*Aj. ft :?/-/**& NBSIR 81-1645 TECHNICAL REPORTS \ MATERIALS STUDIES FO MAGNETIC FUSION ENERGY APPLICATIONS AT LOW TEMPERATURES - IV STRESS INTENSITY RANGE TEMPERATURE To Department of Energy Office of Fusion Energy Washington, D.C. 20545 By Frac Natk Bou' ERRATUM : Replacement for page 463 of NBSIR 81-1645, Materials Studies for Magnetic Fusion Energy Applications at Low Temperatures - IV o . M<&£0fiEc3&.2 3k ■ \^T>&, Mag 8X Fracture Face of CVN Tested at -320°F Mag 1.5X Tensile Tested at Room Temperature \ j . • 1 • ! 9 / * . 1 > • i i • * • '. ° 1 FIGURE 4 Impact, Tensile and Micro structure for Weld Deposited Using U.S. -■ Armco Electrodes c Mn P S Si C£ Ni N 05 9.0 .007 .009 . 22 18 16.2 .18 Toughness p> -320 F \ CVN 13'# UTS ksi 93 Ten? k E tenant Chromic Mag 500X Grain Boundary Defect? Ferrite-Free Weld r I NBSIR 81-1645 TECHNICAL REPORTS MATERIALS STUDIES FOR MAGNETIC FUSION ENERGY APPLICATIONS AT LOW TEMPERATURES - IV Edited By R.P. Reed and N.J. Simon Fracture and Deformation Division National Bureau of Standards Boulder, CO 80303 April 1981 Sponsored By Department of Energy Office of Fusion Energy Washington, D.C. 20545 o t »T or c V o z I \ ♦< * U.S. DEPARTMENT OF COMMERCE, Malcolm Baldrige, Secretary NATIONAL BUREAU OF STANDARDS. Ernest Ambler. Director Q % •2 *» CONTENTS Page SUMMARY v PROGRAM DESCRIPTION vii HIGHLIGHTS OF RESULTS 1 STRUCTURAL ALLOYS STRUCTURAL ALLOYS PROGRAM 11 STRUCTURAL ALLOYS FOR SUPERCONDUCTING MAGNETS IN FUSION ENERGY SYSTEMS 17 STRENGTH AND TOUGHNESS RELATIONSHIP FOR INTERSTITIALLY STRENGTHENED AISI 304 STAINLESS STEELS AT 4 K 37 TENSILE AND FRACTURE PROPERTIES OF MANGANESE-MODIFIED Fe-18Cr-8Ni TYPE AUSTENITIC STAINLESS STEELS 77 INTERSTITIAL CARBON AND NITROGEN EFFECTS ON THE CRYOGENIC FATIGUE CRACK GROWTH OF AISI 304 TYPE STAINLESS STEELS 101 ANOMALOUS YIELDING, TRANSFORMATION, AND RECOVERY EFFECTS IN AISI 304L STAINLESS STEELS 131 LOW-TEMPERATURE DEPENDENCE OF YIELDING IN AISI 316 STAINLESS STEELS 147 LOW-TEMPERATURE DEFORMATION OF 5083-0 ALUMINUM ALLOY 185 EFFECTS OF CARBON AND NITROGEN ON THE ELASTIC CONSTANTS OF STAINLESS STEEL 304 203 LOW-TEMPERATURE VARIABILITY OF STAINLESS-STEEL-304 ELASTIC CONSTANTS 215 PREDICTED SINGLE-CRYSTAL ELASTIC CONSTANTS OF STAINLESS STEEL 304 227 ELASTIC-CONSTANT/TEMPERATURE BEHAVIOR OF THREE HARDENED MARAGING STEELS 237 TEMPERATURE BEHAVIOR OF YOUNG'S MODULI OF FORTY ENGI- NEERING ALLOYS 257 WELDING AND CASTING EVALUATION OF WELDMENTS AND CASTINGS FOR LIQUID HELIUM SERVICE 273 STRENGTH AND TOUGHNESS OF FULLY AUSTENITIC STAINLESS STEEL FILLER METALS AT 4 K 289 WELD PROCESS STUDY FOR 316 L STAINLESS STEEL WELD METAL FOR LIQUID HELIUM SERVICE 303 FRACTURE TOUGHNESS PROPERTIES OF WELDS IN ALUMINUM ALLOY 5083-0 AT 4 K 323 THE INFLUENCE OF FERRITE AND NITROGEN ON MECHANICAL PROP- ERTIES OF CF8M CAST STAINLESS STEEL 337 METALLOGRAPHY OF 6-FERRITE IN AUSTENITIC STAINLESS STEEL CASTINGS 357 DEFORMATION AND FRACTURE OF STAINLESS STEEL CASTINGS AND WELDMENTS AT 4 K 415 EVALUATION OF WELD FILLERS TO JOIN NITROGEN STRENGHTHENED AUSTENITIC STAINLESS STEELS FOR CRYOGENIC APPLICA- TIONS (FUSION AND MHD ENERGY) 453 m NONMETALLICS NONMETALLICS FOR MAGNET SYSTEMS 467 NONMETALLIC MATERIALS: STANDARDIZATION STANDARDIZING NONMETALLIC COMPOSITE MATERIALS FOR CRYOGENIC APPLICATIONS 477 NONMETALLIC MATERIALS: RESEARCH AND CHARACTERIZATION THERMAL CONDUCTIVITY OF G-IOCR AND G-11CR INDUSTRIAL LAMINATES AT LOW TEMPERATURES (PART III) 495 FRACTURE ANALYSIS OF COMPRESSION DAMAGE IN GLASS FABRIC/ EPOXY, G-IOCR, 295 K to 4 K 507 STATIC AND DYNAMIC YOUNG'S MODULI OF TWO FIBER-REINFORCED COMPOSITES 589 RADIATION EFFECTS ON COMPOSITES AND NONMETALLIC MATERIALS 597 NONMETALLIC MATERIALS: TECHNOLOGY TRANSFER CONTRIBUTIONS TO TECHNOLOGY FORECAST '81 605 COMMENT ON "MECHANICAL PROPERTIES OF PLASTIC COMPOSITES.." 609 THE MANUFACTURE OF RADIATION RESISTANT LAMINATES 613 TECHNOLOGY TRANSFER NBS-DoE WORKSHOP: MATERIALS AT LOW TEMPERATURES 641 ORGANIZATIONAL CONTACTS 646 ACKNOWLEDGMENT 647 IV SUMMARY This report contains results of a research program to produce material property data that will facilitate design and development of cryogenic structures for the superconducting magnets of magnetic fusion energy power plants and prototypes. The program was conceived and developed jointly by the staffs of the National Bureau of Standards and the Office of Fusion Energy of the Department of Energy; it is managed by NBS and sponsored by DoE. Research is conducted at NBS and at various other laboratories through subcontracts with NBS. The reports presented here summarize the fourth year of work on the low temperature materials research program. Highlights of the results are presented first. Research results are given for the four main program areas: structural alloys, weldings and castings, nonmetall ics, and technology transfer. Objectives, approaches, and achievements are summarized in an introduction to each program area. The major portion of the program has been the evaluation of the low temperature mechanical and physical properties of stainless steel base metal and welds, with particular emphasis on the nitrogen-strengthened stainless steels. Aluminum alloys have received some consideration also. Work has been done on the production and standardization of nonmetallics, primarily industrial laminates, for low temperature ap- plications and on the measurement of their properties at cryogenic temperatures. A brief description is given of the fourth NBS/DoE Vail workshop held in October 1980. Note: Certain commercial equipment, instruments, or materials are identified in this paper to specify the experimental procedure ade- quately. In no case does such identification imply recommendation or endorsement by the National Bureau of Standards, nor does it imply that the material or equipment identified is necessarily the best available for the purpose. Papers by non-NBS authors have not been reviewed or edited by NBS . Therefore, the National Bureau of Standards accepts no responsibility for comments or recommendations contained therein. VI PROGRAM DESCRIPTION The overall objective of the program is to assist in the design, construction, and safe operation of low temperature magnetic fusion energy (MFE) systems, especially superconducting magnets, through effec- tive materials research and materials technology transfer. The specific steps taken to achieve this objective are: (1) evaluation of low tem- perature materials research needs specific to MFE devices; (2) develop- ment and monitoring of a research program to acquire the necessary data; and (3) rapid dissemination of the data to potential users through personal contacts, publications, and workshops. Efforts directed at the first specific objective began with the publication of the Survey of Low Temperature Materials for Magnetic Fusion Energy in March 1977. A recent publication updating part of this survey, entitled Structural Alloys for Superconducting Magnets in Fusion Energy Systems , is included in this volume. Through continuous interactions with all low temperature design, construction, and measurement programs, such as the Large Coil Project, we are aware of new problems as they arise. This year's contribution to the second objective is described in Table 1 in the form of an outline of the research projects. The results appear later in this report. The third objective is satisfied, in part, by these annual reports and by the series of NBS-DoE Workshops on Mate- rials at Low Temperatures, which are held each fall in Vail, Colorado. Starting in fiscal year 1981, handbook pages presenting the available data for specific materials will be added to the annual report. vn Table 1. Outline of the NBS/DoE program on material studies for magnetic fusion energy applications at low temperatures Program Area 1 . Characterization Organization A. Structural Alloys NBS 2. Al loy development NBS Characterization and modeling of martensitic start temperatures U. Wisconsin 1. Stainless steel welds B. Welds and Castings NBS Aluminum welds NBS Program Description Measurement of elastic, tensile, fracture toughness and fatigue crack growth rate properties as a function of temperature for selected austenitic stainless steels and 5083 aluminum. Assess- ment of effects of grain size, martensitic transformations and alloying on low temperature be- havior. Assessment of low temperature properties of selected austen- itic alloys containing varying concentrations of C, N and Mn with a Fe-Cr-Ni base. Study to assess, empirically, the relation between alloy con- tent and the temperature at which martensite forms on cooling for commercial grades of AISI 300 series stainless steels. Evaluate the strength and toughness of several fully aus- tenitic weld metals provided by suppliers in the United States, Austria, Sweden, and the Soviet Union at 4 K. (Four alloys with 16-20% Cr, 16-20% Ni , 4-7% Mn and 2-5% refractory elements (Mo, W and/or Nb) had strength and toughness combinations comparable to base metal properties.) Evaluate the strength and fracture toughness of GMA welds using 5183 wire in 51 mm thick 5083-0 plates that were provided by five alu- minum fabricators. (Property variations were minimal despite significant differences in welding procedures.) vm L 3. Stainless steel castings ESCO NBS Produce stainless steel castings (grade CF8M) with varied ferrite and nitrogen contents to study the effects of ferrite and ni- trogen on the strength and tough- ness at 4 K. (This year's ef- forts consisted of manufacture of the castings, tests at room temperature and 77 K, and metal lo- graphic studies to be followed by evaluation in liquid helium next year.) C. Nonmetallics for Magnet Structures 1. Standardization NBS 3. Characterization NBS 1. NBS/DoE Workshop Technology Transfer NBS Continue cooperation with in- dustry to establish and refine cryogenic grade specifications for insulating laminates and coding systems for laminates. Measure thermal conductivity of composite components at cryogenic temperatures to aid in estab- lishing models for predicting the thermal conductivity of a variety of laminates. Investigate failure modes of laminates under compres- sion at cryogenic temperatures. Develop dynamic methods of meas- uring elastic constants of com- posite laminates. An annual workshop to present re- search results to the fusion com- munity, to discuss new problems, and to promote interaction between interested parties. Handbook data pages on stainless steels will be produced, starting in FY '81. IX HIGHLIGHTS Of RESULTS HIGHLIGHTS OF RESULTS STRUCTURAL ALLOYS: Stainless Steel 304--Characterization and Alloy Development The study of this alloy series has been emphasized because of its favorable cryogenic properties, extensive service record, and avail- ability. Stainless 304 has been used in several designs for the Large Coil Project (LCP) . Previous work has shown that adding nitrogen increases the tensile strength and austenitic stability of the 304 series at low temperatures. Carbon has a similar effect, but the addi- tion of carbon is limited owing to sensitization (chromium carbide precipitation). It would be desirable to optimize the properties of the alloy for different requirements by varying the composition, and this objective has motivated the studies described below. 1. Carbon and nitrogen (C+N) weight percent were varied system- atically in a series of heats for which fracture toughness and tensile yield strength were measured at 4 K. Fracture tough- ness dropped linearly with increasing C+N content, but strength increased as C+N content increased. 2. The effects of C+N content and temperature upon fatigue crack growth were also measured. Low C+N content was found to correlate with better crack resistance at 4 K. The improved crack resistance for low C+N contents at 4 K, compared with that at room temperature was shown to be due to a transition to transgranular fracture at cryogenic temperatures from intergranular at room temperature. 3. Although interstitial nitrogen strengthens the 304 alloy series, the resulting alloys are harder to fabricate. Since the addition of manganese is expected to improve fabricabil ity, a series of heats with manganese content varying from 1 to 6 weight percent and nitrogen content of 0.1 to 0.2 weight percent were tested for fracture toughness and tensile yield strength. The results gave much lower fracture toughness than expected, except for the heat with 6 percent manganese. These anomalous results may indicate the presence of previously unrecognized factors in alloy production that affect fracture toughness significantly at cryogenic temperatures. Research to elucidate these factors is continuing. 4. Yield strength of a selected heat of 304L was measured as a function of temperature. The measurements showed that yield strength does not increase monotonically with decreasing temperature, as is the case for the austenitically stable stainless 310. 5. Effects of C+N content on the elastic constants of 304 steels were measured at room temperature, and the variability among the elastic constants of alloys of the 304 series was measured down to 4 K. Stainless Steel 316 The low temperature flow strength of stainless 316 was measured. The formation of bcc martensite that occurs at low temperatures or as the result of strain influences this property. Aluminum Alloys The yield strength and stress-strain curves of the aluminum alloy 5083-0 were measured down to 4 K. Serrated yielding was observed at 4 K and room temperature, though not at 76 K. The mechanism causing this effect is different at 4 K than at room temperatures. The magnesium present in this alloy increases the strength at both room and cryogenic temperatures. Previously unpublished data on alloys 5154 and 5052, which have different magnesium content, are presented to illustrate this trend further. WELDING AND CASTING: Research has focused on the metallurgical factors that determine strength and toughness of weldments and castings in stainless steel and aluminum alloys. 1. Since it is known that ferrite reduces the fracture toughness of weldments at 4 K, although ferrite reduces microfissuring, work on the development of percent ferrite filler materials has been undertaken. A series of experimental filler alloys was tested at 4 K. Some were found to have a good combination of strength and toughness, comparable to that of the base metal, but a few problems with microfissuring remain. 2. The effects of differing weld processes were studied for 31 6L weldments. The cleanliness of the process correlated well with good fracture toughness at 4 K, so it was concluded that the impurity level may have a strong effect on cryogenic properties of welds. 3. To support LCP magnet design and construction, a number of qualifying and test weldments of stainless steel were tested for fracture toughness and tensile strength. 4. The evaluation of 5183 welds in 5083 aluminum alloys was completed during the past fiscal year. The fracture toughness of weldments at 4 K was found to be essentially the same for welds furnished by different suppliers or produced by different processes. 5. Work has begun on a study of the more fundamental aspects of the structure of duplex austenite/ferrite steels and on how this structure is related to fracture and deformation in weld- ments and castings at cryogenic temperatures. Castings are being studied for two reasons: (a) there is interest in their use as structural elements for superconducting magnets and (b) the factors affecting fracture and deformation modes are more easily observed because their ferrite structure is coarser than that of weldments. A series of CF8M castings with varied ferrite and nitrogen content was produced this year, and the production techniques are reported. This series will be used to study the effects of ferrite and nitrogen in a controlled manner in the current fiscal year. Meanwhile, during the past year, available weldments and castings were subjected to a detailed metallographic study. Fracture and deformation modes were determined for these specimens, and the effect of ferrite structure on these modes was analyzed. NONMETALLICS: Standardization Efforts to refine and improve the specifications of cryogenic grades of glass-fabric-reinforced bisphenol A epoxy industrial laminates have continued. These established grades are known as G-10CR and G-11CR. Recently it was found that G-11CR does not have the radiation resistance that was initially expected from its composition and curing process. Therefore, this grade has been dropped. The standardization of glass- mat reinforced laminate products is in progress. Characterization The thermal conductivity of the epoxy matrix material used in G-10CR and G-11CR was measured, and the results revealed considerable inhomogeneity in the bulk resin. Efforts to compare thermal conduc- tivity measurements made at the National Bureau of Standards with those made by CERN on a laminate of European manufacture indicated a sizeable discrepancy at room temperature. Work on this problem will be con- tinuing during the present fiscal year. Modes of failure in compressive fracture of laminates at cryogenic temperatures were studied. Compressive stresses of insulators pre- dominate in superconducting magnets. Dynamic elastic moduli were measured for two fiber-reinforced com- posites, and the results were compared with static methods. Results of studies on radiation effects on composites are presented in a section from a review paper. TECHNOLOGY TRANSFER: The annual NBS/DoE Workshop on Materials at Low Temperatures at Vail, Colorado in October 1980 was attended by about seventy representatives of the materials and magnet design communities. Again, the meeting was viewed constructively; it served as a focal point to identify and dis- cuss magnet construction or design problems related to materials. For a few who have become recently involved in cryogenic applications, the workshop served an educational role. STRUCTURAL ALLOYS A STRUCTURAL ALLOYS PROGRAM R. P. Reed Fracture and Deformation Division National Bureau of Standards Boulder, Colorado ABSTRACT An overview of the structural alloys program is presented. Selec- tion and properties of the alloys under study are discussed, and the objectives of the program in the current fiscal year (1981) are set forth. 11 The objectives of the structural alloys project are twofold: (1) to characterize the physical and mechanical properties of structural alloys at 4 K to obtain typical properties than can be used for material selection and design estimates and (2) to develop, in collaboration with industry, a stronger, tough, weldable austenitic steel for use as a structural alloy for superconducting magnets. The structural -alloys-project results from this past year are sum- marized: Characterization : The tensile yield strength as a function of temperature (4-330 K) was measured for a thick plate of 5083-0 aluminum alloy. The temperature dependence of the flow strength of AISI 316 was measured and the influences of martensitic transformations, adiabatic heating, magnetic transitions, and grain size were assessed. Measurements on fatigue and creep of metastable AISI 304 were initiated. The fatigue crack growth rate of a series of Fe-18Cr-8Ni alloys with varying C and N con- centrations was measured and assessed. Measurements of elastic constants of stainless and maraging steels, of the variability of AISI 304 elastic properties and a review of the temperature dependence of elastic constants of engineering alloys are presented. 13 Alloy Development : The tensile yield and ultimate strengths and the fracture toughness of a series of Fe-18Cr- 7Ni-0.03C alloys with varying Mn and N concen- trations were measured at 295, 76, and 4 K. Fracture toughness results obtained for the low (1-4 wt. %) Mn alloys were anomalous when compared with the Fe-18Cr-8Ni-(C+N) alloy series measured the previous year. This year (fiscal year 1981), the tensile flow strength of 2219-T87 aluminum alloy will be measured at a series of temperatures ranging from 4 to 300 K. Additional tensile measurements will be conducted on 5083-0 aluminum alloy to assess the influence of grain size on strength and to assess, more accurately, the temperature dependence of the flow strength. The study of the role of martensitic transformation on low- temperature fatigue and creep of AISI 304 steel will continue. Finally, to pursue the development of a stronger, but tough, austenitic steel, the causes for the lower toughness at 4 K of Fe -18Cr-7Ni -l-4Mn-N steels of the last alloy series need to be understood. Metallography, electron mi- croscopy, scanning electron microscopy, and scanning transmission electron microscopy will be used to study the fracture surfaces, crack development, and impurity segregation of these alloys. Subsequently, an alloy heat of selected composition will be prepared. The program has concentrated on four alloy series: AISI 304 (Fe- 18Cr-8Ni), AISI 316 ( Fe-1 6Cr-10Ni-2Mo) , 5083-0 (Al-0.7Mn-4.45Mg-0.01 5Cr) and 2219 (Al-6Cu-0.3Mn-0.lV-0.18Zr-0.06Ti). These alloys are most commonly used for 4 K applications and are thought to represent the best base alloy materials. 14 The stainless steels have a relatively high Young's modulus, low thermal and electrical conductivity, and excellent toughness. Also, the major fabricators have good experience in welding stainless steel. Aluminum alloys are light, less expensive, and have been used exten- sively (alloy 5083) in the construction of LNG containment systems. Toughness and strength do not decrease between 110 K and 4 K for these aluminum alloys. AISI 304 is susceptible to martensitic transformation on cooling or on application of applied stress at low temperatures. Martensitic transformation tends to act as a deformation mode, therefore its pres- ence tends to reduce the resistance of the alloy to applied stress at low temperatures, and the result is low initial flow strengths. How- ever, we have not obtained evidence that the martensitic transformations act to degrade fracture toughness or ultimate tensile strength. The alloy AISI 316 is less susceptible to martensitic transforma- tion and the dependence of strength on temperature for this steel is more predictable. The addition of nitrogen to austenitic stainless steel dramatically increases the low-temperature yield strength of the steel. There is an attendant decrease of toughness. The aim of the alloy development program is to establish a composition of Fe-Cr-Ni-Mn-C-N that will provide a stronger and tough structural alloy at 4 K. Two correlations have been established to assist this process: (1) At 4 K the tensile yield strength has been demonstrated to be inversely related to the fracture toughness (J-integral ) . (2) The total concentration of C+N from about 0.10 to 0.30 wt. % is linearly related to the tensile yield strength at 4 K. 15 The aluminum alloys do not exhibit low-temperature phase or mag- netic transitions. Consequently, both their physical and mechanical properties are well behaved at wery low temperatures. The fracture toughness of the aluminum alloys is considerably lower than stainless steel values. However, there is a range of slow, stable crack growth following crack initiation at the critical J-integral fracture toughness and additional research is required to characterize the range of stable crack growth. 16 ^ ^P! STRUCTURAL ALLOYS FOR SUPERCONDUCTING MAGNETS IN FUSION ENERGY SYSTEMS H. I. McHenry and R. P. Reed Fracture and Deformation Division National Bureau of Standards Boulder, Colorado 80303 17 Nuclear Engineering and Design 58 (1980) 219-236 © North-Holland Publishing Company STRUCTURAL ALLOYS FOR SUPERCONDUCTING MAGNETS IN FUSION ENERGY SYSTEMS H.I. McHENRY and R.P. REED Thermophysical Properties Division, National Bureau of Standards, Boulder, Colorado 80302, USA Received 15 March 1979 The behaviour of selected alloys for superconducting magnet structures in fusion energy systems is reviewed with emphasis on the following austenitic stainless steels (AISI grades 304, 310S and 316), nitrogen-strengthened austenitic stainless steels (types 304LN, 316 LN and 21Cr-6Ni-9Mn) and aluminum alloys (grades 5083, 6061 and 2219). The mechanical and physical properties of the selected alloys at 4 K are reviewed. Welding, the properties of weldments, and other fabrication considerations are briefly discussed. The available information suggests that several commercial alloys have adequate properties at 4 K and sufficient fabrication characteristics for the large magnet structures needed for fusion energy systems. 1. Introduction Alloys will be used for the main structural mem- bers of the superconducting magnet systems needed for fusion devices. Designs of tokamak and mirror machines that require superconducting magnets are currently in the conceptual stage (except for MFTF), but many design features that strongly influence materials selection are reasonably well established. The principal features governing alloy selection for structural applications are the enormous size and stored energy of the magnet systems, the extremely high forces exerted by the magnets, the massive struc- tural elements needed to restrain these forces, the limited space available for the structure, and the need for accessibility to install and periodically remove the blanket and shield systems - systems that are completely surrounded by magnets and the support structure. These features result in the need for alloys with sufficient strength and stiffness to perform the required structural functions within the space available, sufficient fatigue and fracture resistance to operate safely, and sufficient fabricab- ility to permit manufacture and assembly of the com- ponents. The choice of materials is further restricted by the adverse operating environment of the magnet system. The alloys are subjected to temperatures down to 4 K and must not be susceptible to low temperature embrittlement. The presence of high magnetic fields of controlled shape cautions against the use of ferro- magnetic materials in certain applications because they would be subjected to higher forces due to the magnetic field and would result in field distortion. The rapid magnetic field changes associated with cyclic operation of the poloidal field coils cause heat inputs to the structure from eddy current losses. Consequently, alloys with high electrical resistivity are usually preferred. The radiation environment eliminates use of alloying elements, such as cobalt, which transmute to long half-life products. With this exception, radiation does not appear to be a signifi- cant factor in the selection of structural alloys because dosage levels sufficient to degrade struc- tural performance probably cause unacceptable deter- ioration of other system materials. Finally, the micro- structure of the material must be stable when exposed to the operating strains, temperature, and radiation; e.g. the martensite transformation in metastable austenitic stainless steels produces a brittle, ferro- magnetic phase and the attendant volume expansion causes localized areas of stress concentration. The contraints on materials selection caused by size requirements, space limitations and operating environment significantly reduce the number of can- didate alloys. Further limitations arise due to con- 19 //./. McHenry, R.P. Reed / Structural alloys for superconducting magnets siderations related to experience, cost and availability. Currently, the leading candidates are the austenitic stainless steels and aluminum alloys. In this paper the properties of the alloys considered most suitable for liquid helium applications are reviewed with emphasis on their properties at 4 K and on their fabrication characteristics relevant to fusion magnet systems. This review draws heavily on research conducted at the National Bureau of Standards; particularly a three-year program on materials for superconducting machinery sponsored by the Defense Advanced Re- search Projects Agency [1], a survey of low tempera- ture materials for magnetic fusion energy [2], and a continuing program on materials for fusion energy magnets [3] sponsored by the Office of Fusion Energy of the US Department of Energy. For background relating to the properties discussed the reader is refer- red to reviews by Fickett [4] on the physical basis for each property and by McHenry [5] on fracture mechanics and its application to cryogenic structures. 2. Annealed austenitic stainless steels Austenitic chromium— nickel stainless steels, includ- ing the nitrogen-strengthened grades discussed in more detail in section 3, are the most widely proposed alloys for structural applications in superconducting magnet systems. These steels have the best combina- tion of strength, stiffness and toughness at 4 K. The physical properties also offer advantages over compet- ing materials. Specifically, the modulus is high, the thermal expansion is close to that of the copper- stabilized conductor, magnetic permeability is low and the electrical and thermal conductivities are low. The principal disadvantages are the high cost of con- struction and problems arising from microstructural stability and control. The AISI 300-series stainless steels, particularly types 304 and 304L, are widely used in cryogenic applications. The designations and compositions of commonly used grades are summarized in table 1 . These grades have moderate strength, excellent thoughness, good fabrication characteristics and are available in a variety of product forms. The ASTM specifications for the various product forms are listed in table 2. Several 300-series grades not listed in table 1 are cover- Table 1 Compositions of austenitic stainless steels suitable for liquid helium service AISI Composition (%) type no. Cr Ni C, max Other 304 18-20 8-10.5 0.08 304L 18-20 8-12 0.03 316 16-18 10-14 0.08 2-3 Mo 316L 16-18 19-22 0.03 2-3 Mo 310 24-26 19-22 0.25 1.5 Si, max 310S 24-26 19-22 0.08 1.5 Si, max For each grade Mn P= 0.045 max. 2.0 max; S = 0.030 max; Si = 1.0 max; ed by ASTM specifications and can meet ASME Pres- sure Vesels Code requirements for cryogenic vessels. These grades are not described in this chapter because they offer no particular advantage for superconduct- ing magnet structures and consequently their service experience and property data at 4 K are limited. The alloys listed in table 1 can be compared with respect to AISI 304, the basic 19Cr-9Ni stainless steel. AISI 316 contains 2-3% Mo and slightly higher nickel and thus has greater austenite stability than AISI 304. AISI 304L and 316L are low carbon modifications of 304 and 316, respectively; low car- bon is desirable to avoid sensitization, the grain bound- ary precipitation of Cr-carbides. Type 31 OS is a 25Cr-20Ni alloy that has lower carbon (0.08% C) than AISI 310 (0.25% C). The high alloy content Table 2 ASTM specifications for various product forms of austenitic stainless steels Product form ASTM Specifications Plates, sheet and strip A240 Forgings A473 Bars and sections A479 Seamless tubes A213 Welded tubes A249 Seamless and welded tubes A269 Small diameter tubing A632 Bolts, screws, studs A320 Nuts A194 20 H.I. McHenry, R.P. Reed / Structural alloys for superconducting magnets provides austenite stability and greater strength and thus 31 OS is useful where dimensional stability, non magnetic behavior and higher strength are desirable. The low temperature properties of the 300-series austenitic stainless steels have been compiled in cryo- genic handbooks L 6,7] , in a review article by Brickner and Defilippi [8], and in suppliers' publications [9,10]. In this section the general trends in strength, toughness and fatigue resistance are summarized as a function of temperature. Care should be taken in using the data presented here or in data compilations because there are significant heat-to-heat variations in the mechanical properties of the 300-series stain- less steels, particularly in the yield strength and toughness. The principal causes of these variations are interstitial content, i.e. carbon and nitrogen, and mechanical deformation history, i.e. resulting grain size. The influence of nitrogen within the limits of 0—0.08 wt% is particularly significant with the higher nitrogen contents causing substantial increases in yield strength. 2.1. Tensile properties The yield and ultimate strengths of AISI grades 304, 310 and 316 are compared in fig. 1 for tempera- tures ranging from 4 to 300 K. Notice that type 304 has the lowest yield strength and the highest ultimate strength at all temperatures. This is due in part to the martensite transformation which occurs more readily in 304 than in 316 or 310. Since the transformation is strain induced, the yield strength is not significantly influenced by martensite, but the ultimate strength is increased. The yield and ultimate strengths of AISI grades 304 and 304L are compared in fig. 2. Type 304 exhibits a slightly greater increase in strength than type 304L as temperature is reduced. The loss of strength associated with reduced carbon content is considered to be characteristic of all the stainless steels, i.e. 304L vs. 304, 316L vs. 316 and 310S vs. 310, and is attri- buted to the strengthening effect of carbon. The ductility of the annealed 300-series stainless steels is generally exce'lent at cryogenic temperatures. Elongation and reduction-of-area tend to drop with decreasing temperatures, but values generally exceed 30%. 1600 1200 - 800 - 400 1 >/304 1 BAR yTensile Strength 310^^\. ,316 ^310 Yield Strength^ 1 f — — = 1 100 300 200 TEMPERATURE, K Fig. 1. Tensile and yield strengths of three austenitic stainless steels - AISI types 304, 310 and 316 - at temperatures between 4 and 300 K [6]. 2.2. Toughness The austenitic stainless steels retain excellent tough- ness at cryogenic temperatures. The Charpy V-notch impact toughness does decrease with temperature but in all cases known to the authors the toughness far exceeds the 0.38 mm lateral expansion requirement of the ASME Boiler and Pressure Vessel Code [1 1 ] and a ductile-to-brittle transition is not exhibited. The minimum temperature for Charpy testing is about 20 K, because adiabatic heating causes specimen tem- peratures in excess of 20 K even when the test tem- perature is 4 K. Slow strain-rate tests using notched tensile speci- mens indicate that notch sensitivity does not develop at 4 K [6] . The data summarized in fig. 3 indicate that the notch-to-unnotched ratio of tensile strengths is approximately equal to 1 and the ratio of notched tensile strength to unhotched yield strength generally exceeds 2. 21 H.I. McHenry, R.P. Reed / Structural alloys for superconducting magnets 1800 1600 1400 1200 1000 800 600 400 200 Ultimate Tensile 304L] Strength SHEET -304 } -304L >0.2% Offset Yield Strennth u 100 200 TEMPERATURE, K Fig. 2. Tensile properties of two austenitic stainless steels - AISI types 304 and 304L - at temperatures between 4 and 300 K [6]. 300 Tobler [12] used /-integral methods to measure th( fracture toughness of AISI grades 310 and 316 at tem- peratures to 4 K. The toughness, K\ C {J), at 4 K is 230 MPaVm for 31 OS and 460 MPa\/m for 316, and in both alloys the toughness at 4 K exceeds the tough- ness at room temperature. The ratio of toughness to yield strength is sufficiently high to ensure gross duc- tile deformation prior to fracture. 2.3. Fatigue Strain cycling fatigue properties at 300, 76 and 4 K have been measured for AISI grades 304L and 310 by Nachtigall [13] and for grades 304L and 316 by Shepic and Schwartzberg [14]. The results for 304L, which were essentially the same in both investi- gations, and for 316 are shown in fig. 4. The fatigue resistance of 310 and 316 were superior to that of 304L, particularly in low cycle fatigue (less than 10 4 cycles). For each alloy the fatigue resistance at low temperatures is superior to the fatigue resistance at room temperature except at the highest strain ranges where failure occurred in less than 1000 cycles. The fatigue crack growth behavior of AISI grades 304, 304L, 310S and 316 has been determined at 295, 76 and 4 K by Tobler and Reed [15] . The data for 304L are shown in fig. 5. Notice that the growth rates are essentially the same at 76 and 4 K; similar behavior was observed in each of the other alloys. Also, the growth rates at 76 and 4 K are slower than those at room temperature. The best-fit lines through the 304L data at 295 K and at 76 and 4 K form the NOTCHED TEN.SILE PROPERTIES OF l°* STAINLESS STEEL AT 4K _ K =3 T 1 NTS 304L ' J ys sheet V 5. ["J NTS 310 304 sheet bar K T =3 310S K=5.2 forqi nr] _ 310 T plate 316 K=fi ■>. slate r T -6.i Fig. 3. Notched tensile properties of five austenitic stainless steels - AISI types 304, 304L, 310, 310S and 316 - at 4 K [6]. Kf is the stress concentration factor of the notch. _, i-o - \ \l 1 \^-316 1 1 1 ^-304 - I 1 1 ,21Cr-6Ni 1 -9Mn 1 _ 10 10 10 10 FATIGUE LIFE, N. cycles Fig. 4. Fatigue behavior of three austenitic stainless steels - AISI types 304 and 316 and a 21Cr-6Ni-9Mn alloy - at 4K [14]. 22 ■I H.I. McHenry, R.P. Reed / Structural alloys for superconducting magnets 10" J - 10 - 1 i i i 1 1 1 l I I I 1 1 ] i AISI 304L (annealed) •f : • T=295K " ~ o T=76K - ~ a T=4K ; - •L ? in - - " - - - - - •, fa - - - - L° - i i i i i i 1 i -r£ t i i i I 1 1 10 100 STRESS INTENSITY FACTOR RANGE, AK, MPa-™' Fig. 5. Faitgue crack growth rates of AISI type 304L austen- itic stainless steel at 295, 76 and 4 K [ 15] . approximate scatter bands for the data on 304, 31 OS and 316 at 295, 76 and 4 K. Thus, it can be concluded that alloy content and temperature have minimal influence on the fatigue crack growth behavior of the 300-series stainless steels. 2.4. Physical properties The physical properties of the 300-series stainless steels at 295, 76 and 4 K are summarized in table 3. The thermal properties have been taken from the best-fit lines used in the LNG Materials Handbook [7] , and the original sources are given for the elastic properties [16], electrical resistivity [17], and magne- tic permeability [18]. 3. Annealed nitrogen-strengthened austenitic stainless steels Nitrogen has a pronounced effect on the yield strength of austenitic stainless steels, particularly at cryogenic temperatures. For example, the yield strength of 304 N (AISI 304 with deliberate additions of nitrogen) is approximately three times greater than the yield strength of 304 at 76 K, the ultimate strengths are approximately equal and the elongation of the 304N is superior at temperatures below 220 K [20]. This remarkable improvement in mechanical proper- ties occurs in many nitrogen-strengthened grades. Consequently, significant interest is developing in these alloys for future applications in superconducting magnet systems. The designations and compositions of the com- mercially available grades that offer promise for liquid helium applications are summarized in table 4. Several other nitrogen-strengthened grades have been developed and are covered by ASTM specifications such as A240 and A41 2. However, the availability and property data on these grades are limited. Several additional grades are commonly used in Europe where Table 3 Physical Propeties of austenitic stainless steels Property Units AISI 304 AISI 310 AISI 316 295 K 4K Ref. 295K 4 K Ref 295 K 4 K Ref. Density g/cm 3 7.9 _ 19 7.9 _ 19 8.0 19 Young's modulus GN/m 2 189 204 16 184 199 16 217 234 16 Shear modulus GN/m 2 740 800 16 700 770 16 830 910 16 Poisson's ratio - 0.290 0.271 16 0.307 0.296 16 0.298 0.281 16 Thermal conductivity W/m -K 13 0.3 7 nearly the same for 304, 310 and 316 Expansion cm/cm KX10 -6 15.8 10.2 7 nearly the same for 304, 310 and 316 Specific heat J/kg • K 480 0.5 7 nearly the same for 304, 310and 316 Electrical resistivity p ncm 70.4 49.6 17 87.3 68.5 17 75.0 53.9 17 Magnetic permeability 1.02 1.09 18 1.003 1.10 18 1.003 1.02 18 23 H.I. McHenry, R.P. Reed I Structural alloys for superconducting magnets Table 4 Compositions of nitrogen-strengthened austenitic stainless steels suitable for liquid helium service Type Nominal composition (%) Cr Ni Mn C, max N Other AISI 304N 18 -20 8 -10.5 2.0, max 0.08 0.10-0.16 _ 304LN 18 -20 8 -12 2.0, max 0.03 0.10-0.16 - AISI 316N 16 -18 10 -14 2.0, max 0.08 0.10-0.16 - 316LN 16 -18 10 -14 2.0, max 0.03 0.10-0.16 - ASTMXM11 * 18 -21 5-7 8-10 0.04 0.15-0.40 - ASTMXM19* 20.5 -23.5 11.5-13.5 4- 6 0.06 0.20-0.40 Cb 0.10-0.30 V 0.10-0.30 Mol.5 -3.0 * ASTM A412. nitrogen is deliberately added to several of the com- mon Cr-Ni 300-series stainless steels. The equivalent AISI designations for these European grades are 304LNand316LN. Two basic types of nitrogen-strengthened stainless 1600 1400 1200 1000 200 ■304LN [21 316LN [21] | 316LN [26] "0 100 200 TEMPERATURE, K Fig. 6. Tensile and yield strengths of nitrogen-strengthened 300-series stainless steels at temperatures between 4 and 300 K [20-27]. 300 steels are suitable for cyrogenic service: the Cr-Ni-N and the Cr-Ni-Mn-N alloys. The Cr-Ni-N alloys are 300-series stainless steels with deliberate additions of nitrogen. In the US specifications, the nitrogen range is 0.10-0.16% for 304N and 316N. The corres- ponding European specifications permit nitrogen levels to 0.25%. The manganese in Cr— Ni-Mn— N alloys provides austenite stability and can thus be used to replace part of the Ni content. Manganese also increases the solubility of nitrogen in austenite and thus permits higher nitrogen levels —0.4% maxi- mum nitrogen is typical. 3.1. Tensile properties The yield and ultimate strengths of the Cr— Ni— N 300-series stainless steels are summarized in fig. 7. The solid lines and the broken lines represent the data of Randak et al. [21] for 304LN and 316LN, respectively. The various symbols depict the data for seven other investigations [20,22—27] , each on one of the following alloys: 304LN, 304N and 316LN. As in the case for 304 and 316, fig. 1, alloys 304N and 304 LN have approximately 20% greater ultimate strength at cryogenic temperatures than alloy 316LN. All three alloys have essentially the same yield strength at and below room temperature. There is relatively little scatter in the yield strength data considering that data on 17 heats from a variety of sources are plotted. The nitrogen contents of the 17 heats range from 0.09 to 0.17%. Of the four nitrogen-strengthened Cr-Ni— Mn stain- 24 H.I. McHenry, R.P. Reed / Structural alloys for superconducting ma&iets 2.0 1.6 1.2 0.8 - 0.6 - 1 21Cr-6Ni-9Mn 1 o Tobler and Reed cj-4^ v A Scardiqno Q Landon T ^^^ v Masteller 4 \ [ Malin _ Tensi le ~- I a ^~-— a — Yield 1 i 300 100 200 TEMPERATURE, K Fig. 7. Tensile and yield strengths of 21Cr— 6Ni-9Mn stain less steel at temperatures between 4 and 300 K [28]. less steels listed in table 4, the 21Cr-6Ni-9Mn alloy has been most thoroughly evaluated [28—32]. The tensile and yield strength data are summarized in Fig. 6. This alloy typically has about 25% greater yield strength than the Cr— Ni— N alloys. Represen- tative tensile data on the other nitrogen-strengthened Cr-Ni-Mn alloys that have been tested at 4 K indicate that yield strength increases and ductility decreases with decreasing temperature; and thus, fracture resis- tance must be considered when selecting these alloys for liquid helium service. 3.2. Toughness The nitrogen-strengthened stainless steels have excellent toughness at room temperature. However, as shown in fig. 8, significant toughness losses gener- ally occur as the temperature is reduced [25]. Read and Reed [25] observed significant variations in the toughness of a single piece of 21Cr-12Ni-5Mn alloy. As-received mill-annealed material having an inter- granular microconstituent had a plane strain fracture toughness at 4 K of 1 1 1 MPa\/m- A high temperature (1177°C, 1.5 h) anneal dissolved the microconstituent and increased the toughness to 176 MPa-s/m- Thus, certain nitrogen-strengthened grades can be used at 4 K, but care should be taken to ensure that the alloy selected has satisfactory toughness at 4 K for the applicable melting practice, product form and heat treatment. 500 400 200 • Fe-18Cr-3Ni-13Mn ■ Fe-21Cr-12Ni-5Mn _ A Fe-19Cr-9Ni-2Mn ♦ Fe-21Cr-12Ni-5Mn-X ▼ Fe-21Cr-6Ni-9Mn I 300 100 200 TEMPERATURE, K Fig. 8. Fracture toughness of four nitrogen-strengthened austenitic stainless steels at temperatures between 4 and 300K [25]. 3.3. Fatigue The strain cycling fatigue behavior of 21Cr-6Ni- 9Mn has been measured at 295, 76 and 4 K by Shepic and Schwartzberg [14] . The results at 4 K, shown in fig. 4, indicate that the 21Cr-6Ni-9Mn alloy has fatigue strength better than 304L but not as good as 316. Thus, the significantly higher strength of the nitrogen-strengthened grade does not result in a comparable improvement in fatigue life. The fatigue crack growth behavior of several nitro- gen-strengthened grades has been studied at cryogenic temperatures [25,28]. The results for the high nitro- gen Cr-Ni-Mn alloys depend largely on the relative austenite stability of the alloys. For the least stable alloy evaluated, 18Cr-3Ni-13Mn, the fatigue crack growth rates at 4 K are 50 times faster than those for the most stable alloy evaluated, 21Cr— 12Ni— 5Mn, and 10 times faster than the growth rates for 21Cr— 6Ni— 9Mn. For the 21Cr-12Ni-5Mn alloy, the fatigue crack growth rates fall within the same scatter band as the 300 series (N < 0.08%) alloys at 4 K. For the Cr-Ni alloys, comparison of 304 and 304L with 304N and 304LN shows that the nitrogen-strengthened 25 //./. McHenry, R.P. Reed / Structural alloys for superconducting magnets alloys have growth rates at 4 K about 4 to 5 times higher than the lower nitrogen grades. In contrast, the 316 alloy, which has greater austenite stability than 304, has essentially the same growth rates at 4Kasthe316LNalloy. 3. 4. Physical properties In general, the physical properties of the nitrogen- strengthened stainless steels are similar to the 300- series alloys, as summarized in table 3. In general, the nitrogen contributes to austenite stability and thus reduces the magnetic permeability induced in some alloys by deformation at low temperatures. Anomalous behavior has been observed in high man- ganese alloys by Ledbetter [33] in the elastic behavior and by Ledbetter and Collings [34] in the magnetic susceptibility. 4. Fabrication characteristics of austenitic stainless steels The fabricability of austenitic stainless steels is certainly sufficient for fusion magnet systems as attested by the widespread use of stainless steels in a variety of industries. Fabrication characteristics that should receive special consideration for fusion magnet applications are briefly reviewed in this sec- tion. 4.1. Welding The austenitic stainless steels, with and without nitrogen strengthening, are readily weldable by all of the common welding processes, providing the appropriate consumables and procedures are used. The service experience with welded assemblies of the alloys has been satisfactory at cryogenic temperatures. The American Welding Society, AWS, designation sys- tem, which is similar to the AISI system for wrought products, is generally used to classify stainless steel filled metals. The AWS designations and chemical compositions of the filler metals commonly used for cryogenic applications are summarized in table 5. The strength of stainless steel welds at 4 K generally ex- ceeds the corresponding base metal strength, but toughness is usually significantly lower. Three pheno- mena that affect the strength and toughness of the as-deposited weld metal take on added significance at cryogenic temperatures: sensitization, ferrite con- tent and nitrogen pickup. In addition, heat treatment Table 5 Compositions of filler metals for welding austenitic stainless steels AWS type Composition (%) Cr Ni C, max Other E308 18 -21 9 -11 0.08 E308L 18 -21 9 -11 0.04 E310 25 -28 20 -22.5 0.20 0.75 Si, max E16-8-2 14.5-16.5 7.5- 9.5 0.10 0.50 Si, max; 1-2 Mo E316 17 -20 11 -14 0.08 2-2.5 Mo E316L 17 -20 11 -14 0.04 2-2.5 Mo E330 14 -17 33 -37 0.25 Rod and bare electrodes ER308 19.5-22 9-11 0.08 ER308L 19.5-22 9-11 0.03 ER310 25 -28 20-22.5 0.08-0.15 ER316 18 -20 11-14 0.08 2-3 Mo ER316L 18 -20 11-14 0.03 2-3 Mo For covered electrodes: Mn = 2.5, max; Si = 0.9, max; P = 0.04, max; S = 0.03 max. For rod and bare electrodes: Mn = 1.0-2.5; Si = 0.25-0.60; P = 0.03; S = 0.03, max. 26 H.I. McHenry, R.P. Reed / Structural alloys for superconducting magnets generally has a stronger influence on the weld metal than the base metal. Sensitization, the grain boundary precipitation of Cr carbides, reduces weld metal toughness at low temperatures [35]. Consequently, weld metals with extra low carbon, such as 308L and 316L, are com- monly used for cryogenic applications. Sensitization is most extreme if the weldment is given prolonged exposure in the temperature range 500— 800°C. How- ever, the underlying weld beads in multipass weld- ments may receive sufficient exposure to these tem- peratures to cause significant toughness reductions in stainless steels with greater than 0.03% carbon. Subsequent annealing i >f the welds at temperatures greater than 950°C dissolves the carbides and im- proves the toughness. The ferrite content of as-deposited weld metals can be estimated by using the Schaeffler constitu- tion diagram. To avoid microfussuring the weld metal chemistry is generally balanced to provide 4—10% ferrite. However, for cryogenic service there is concern that the ferrite, which is a bcc phase and therefore brittle at low temperatures, may embrittle the weld metal. Szumachowski and Reid [36] have shown that ferrite reduces toughness at 76 K for a wide range of stainless steel weld metals. Read, McHenry and Steinmeyer [35] have evaluated the ? 50 316 Weldment 2 4 6 PERCENT FERRITE 10 Fig. 9. Fracture toughness of AWS 316L weld metal at 4 K as a function of ferrite content [35]. effect of ferrite content on the fracture toughness of shielded metal arc welds of AWS 316L; the results at 76 and 4 K are shown in fig. 9. Nitrogen increases the yield strength and decreases the toughness of stainless steel weld metals. When conventional (low N) filler metals are used, nitrogen variations in the range 0.04—0.13% may occur depend- ing on the amount of air which enters the arc shielding medium. The greatest variations in nitrogen pickup are associated with gas metal arc welding where the nitrogen content ranges from 0.04 to 0.13% depend- ing on the type of gas coverage and the electrode stickout. Smaller variations are attributed to electrode coatings; lime coatings generally give better coverage and less nitrogen pickup than titania coatings. 4.2. Heat treatment One of the principal advantages of annealed austen- itic stainless steels is that heat treatment is generally not required. The possible exceptions are annealing at 1000—1 100°C following severe forming operations and stress relief at 420-480°C to reduce peak residual stresses and improve dimensional stability. Heat treatment of stainless steel welds has been studied by Bennett and Dillon [37] for AWS types 308, 308Land 310. Heat treating or slow cooling in the sensitization range causes significant toughness loss at 76 K in types 308 and 310, but does not affect 308L. Stress relieving at 840°C followed by furnace cooling causes about a 50% toughness loss in each alloy (20—32 J as-welded and 8—20 J stress-relieved). Stress relieving at 950°C followed by furnace cooling caused a 50% toughness loss in 310, no change in 308L, and mixed results (±40% toughness change) in 308. Annealing at 1070°C followed by water quench- ing improved the toughness of 308 and 308L, but decreased the toughness of 310. The beneficial effect of a full anneal on toughness at 76 K has also been reported by Krivobok and Thomas [38] , McConnell and Brady [39] and DeLong [40]. 4.3. Castings The austenitic stainless steels are widely used for manufacturing corrosion resistant castings. The casting alloys equivalent to AISI types 304 and 316 are ACI (Alloy Casting Institute) types CF8 and CF8M, respect- 27 H.I. McHenry, R.P. Reed / Structural alloys for superconducting magnets ively. The ferrite content of these alloys may vary from to 40% due to chemistry variations within the allowable ranges. Ferrite levels in this range are likely to be unacceptable because of low fracture toughness and high permeability. Ferrite-free castings require the use of higher nickel content alloys such as ACI grades CK20 (corresponding to AISI grade 310). A 16Cr-20Ni-10Mn-2.2Mo alloy has been used for liquid hydrogen [41] and liquid helium bubble chambers [42,43] ; the alloy retains excellent strength (550 MPa yield strength) and ductility at temperatures to 2.5 K [43] . 4.4. Machining, forming and forging Machining: Annealed austenitic stainless steels are more difficult to machine than carbon steels or aluminum alloys because of their relatively high tensile strength, work hardening rates, and low thermal conductivity. These properties cause poor finishes, heat buildup and excessive tool wear. As a result, it is necessary to use heavier feeds, slower speeds, better quality tool bits, more rigid machines and more power. Forming: The candidate stainless steels have excel- lent formability. Alloys with more austenite stability than types 304 or 316 may be needed if severe form- ing operations are required. The nitrogen-strengthened 21Cr-6Ni-9Mn alloy is reported to have excellent austenite stability; even after 60% reduction at room temperature the permeability is less than 1 .02 [44] . Forging: The austenitic stainless steels can be forged into shapes suitable for many magnet applica- tions, e.g. forged pancake billets for ring rolling coil forms. Stainless steels are more difficult to forge than most common structural alloys because of their high strength at elevated temperatures. Within the family of austenitic stainless steels, the forgeability of types 304 and 316 is highly rated [45]. 4.5. Inspection During fabrication of fusion magnet systems, the stainless steel material and components are likely to be inspected using conventional methods. The only problem peculiar to the proposed designs is the need for heavy-section stainless steel welds. These welds cannot be effectively inspected for the presence of cracks or crack-like defects. Radiography is of limited use because of its inability to reliably detect. cracks. Ultrasonic methods do not work effectively because of the high noise levels and large and variable attenua- tion arising from the large-grain-size welds. Standards are needed to assure that the defect sizes that can be reliably detected do not adversely affect structural integrity. 5. Aluminum alloys Aluminum alloys have been proposed for many structural applications in fusion magnet systems. Their principal advantages are: low as-fabricated cost, light weight, non-magnetic behavior, stable micro- structure, and good retention of strength and tough- ness at cryogenic temperatures. The main disadvan- tages of aluminum alloys are low strength in weld- ments, high electrical and thermal conductivity, and an unfavorable elastic modulus (too low) and thermal expansion (too high) for many applications. For fusion magnet structures the preferred alumi- num alloys are 5083, 2219 and 6061. The composi- tion ranges for these alloys are summarized in table 6. Alloy 5083 is generally used in the annealed condi- tion (5083-0) and has a typical yield strength of 171 MPa (25 ksi) at 4 K. Alloy 2219 is generally used in a heat treated condition (T85 1 , T87, etc.), and in the T851 temper has a typical yield strength of 482 MPa (70 ksi) at 4 K. Alloy 6061 is usually used in the heat treated condition (T6, T651, etc.), and in the T6 temper has a typical yield strength of 380 MPa (55 ksi) at 4 K. Although each of these alloys is read- ily weldable, only 5083 retains full strength in the welded condition and is the best choice for most large welded components. Alloy 2219 has the highest strength of the three alloys and is recommended for bolted construction where space or weight need to be conserved. Several other aluminum alloys have excellent pro- perties at 4 K and may be used for special applica- tions. Commercially pure aluminum (alloy 1 100) and the 3003 alloy have excellent formability. Alloy 3003 has been widely used for cryogenic applica- tions and is the best choice when brazing is used for joining. The high strength alloys such as 2014 and 7005 may be used for nuts and bolts. 28 H.I. McHenry, R.P. Reed / Structural alloys for superconducting magnets Table 6 Compositions of aluminum alloys suitable for liquid helium service [46] Type Composition (%) Si max Fe max Cu Mn Ml! Cr Zn max Ti Others 1100 1.0 SH 2219 0.2 0.3 3003 0.6 0.7 5083 0.4 0.4 6061 0.4- -0.8 0.7 Si + Fe, max 0.05-0.20 0.05, max 5.8 -6. 1 0.20-0.40 0.02, max 0.10 99.00 Al, minimum 0.10 0.02-0.10 0.05-15 V; 0.10-0.25 Zr 0.05-0.20 1.0 -1.5 - - 0.10 - 0.1, max 0.40-1.0 4.0-4.9 0.05-0.25 0.25 0.15, max 0.15-0.40 0.15, max 0.08-1.2 0.04-0.35 0.25 0.15, max The ASTM specifications for aluminum mill pro- ducts of alloys 5083, 2219 and 6061 are summarized in table 7. For a more complete listing of specifications refer to the Aluminum Standards and Data Book [46] . Procurement of heavy sections may require additional specification requirements on through-thickness ductility [47]. The low temperature properties of aluminum alloys have been compiled in cryogenic handbooks [6,7], in a review article by Kaufman and Wanderer [48] and in suppliers' publications [49] . In this sec- tion the general trends in strength, toughness and fatigue resistance are summarized, particularly the data for room remperature and 4 K. The data are taken on relatively few lots of material, but generally the lot-to-lot variations in the mechanical properties of aluminum alloys are relatively small (with respect to stainless steel). 5.7. Tensile properties The mechanical properties of alloys 5083, 6061 and 2219 at room temperature and 4 K have been determined by Kaufman, Nelson and Wanderer [50] and the results are summarized in table 8. Notice that in all cases the tensile strengths, and to a lesser extent the yield strengths, are superior at 4 K and the ductil- ity values are not significantly changed. The alloys retain good notched-tensile properties at 4 K. Aluminum alloys having a wide range of strength levels are suitable for liquid helium applications. The yield strengths of commercially pure (CP) aluminum (types 1 100-0 and 1099-H14) and the preferred alloys are shown as a function of temperature in fig. 10 [42] . The yield strengths of each alloy gradually increase with decreasing temperature. Strength increases in the following order: annealed CP (1 100-0), cold Table 7 ASTM specifications and availability of aluminum mill products [46] Product form Specification ASTM Available alloys 5083 6061 2219 Sheet and plate B209 Wire, rod, bar; rolled or B211 cold finished Wire, rod, bar, shapes, tube B221 Structural shapes B308 Pipe B241 Forgings and forging stock B247 Structural pipe and tube; B249 extruded 29 H.I. McHenry, R.P. Reed / Structural alloys for superconducting magnets Table 8 Tensile properties of aluminum alloys at room temperature and 4 K [50] Alloy and Plate Orienta- Temp. TS YS Elongation Reduction NTS NTS NTS temper thickness tion (K) tensile yield in4D of area notch TS YS (mm) strength (MPa) strength (MPa) (%) (%) tensile strength (MPa) 5083-0 25 L RT 322 141 19.5 26 372 1.16 2.65 4 557 178 32 33 429 0.77 2.42 5083-H321 25 L RT 335 235 15 23 421 1.26 1.80 4 591 279 29 33 508 0.86 1.82 6061-T651 25 L RT 309 291 16.5 50 477 1.54 1.64 4 483 379 25.5 42 619 1.28 1.63 T RT 309 278 15.2 42 467 1.51 1.68 4 485 363 23 33 601 1.24 1.66 2219-T851 25 L RT 466 371 11 27 547 1.12 1.48 4 659 484 15 26 703 1.06 1.48 T RT 457 353 10.2 22 531 1.16 1.50 4 659 481 13 21 665 1.01 1.38 2219-T87 25 L RT 464 387 11.8 28 567 1.22 1.46 4 674 511 15 23 690 1.03 1.37 A356-T61 19 casting RT 287 208 8.8 10 354 1.23 1.70 casting 4 455 331 7.1 9 495 1.09 1.50 worked CP (1099-H14), solution-strengthened Al-Mg alloy (5083-0), cold worked Al-Mg alloy (5083-H321), precipitation hardened Al-Mg-Si alloy (6061-T651) and precipitation hardened Al-Cu alloy (2219-T851). Still higher yield strengths are available in several other 2024-T851 and 7005- 100 200 TEMPERATURE, K Fig. 10. Yield strengths of six aluminum alloys at tempera- tures between 4 and 300 K [50]. 300 alloys including 2014-T651 T5351. 5.2. Toughness Aluminum alloys generally retain excellent tough- ness at cryogenic temperatures. A useful measure of •aluminum alloys toughness at 4 K is the notch-yield ratio, i.e. the ratio of notched tensile strength to un- notched yield strength [51]. Ratios substantially greater than one, say 1 .5, indicate that the alloys retain the ability to plastically deform and resist crack initiation in the presence of sharp stress con- centrations. As shown in table 9, the 5083 and 6061 alloys have notch-yield ratios greater than 1 .6 and the higher-strength 2219-T851 and T87 alloys have notch-yield ratios of about 1.4. Significantly lower ratios (1 .0-1 .25) occur in other high strength alloys such as 2014-T651 and 2024-T851 and in most 7000 series alloys (7005 is an exception), where the ratio is sometimes below 0.4 at cryogenic temperatures [52]. Limited toughness test data are available on alu- minum alloys at 4 K. Tobler and Reed [53] measured the fracture toughness of 5083-0 (4.3 cm thick plate) 30 H.I. McHenry, R.P. Reed / Structural alloys for superconducting magnets Table 9 Physical properties of aluminum alloys Property Units 5083 (annealed) Precipitation hardened 6061 Precipitation hardened 2219 295 K 4K Ref. 295 K 4K Ref. 295 K 4K Ref. Specific gravity g/cm 3 2.66 2.70 2.83 Young's modulus GN/m 2 71.5 80.9 58 70.1 77.7 7 77.4 85.7 59 Shear modulus GN/m 2 26.8 30.7 58 26.4 29.2 7 29.1 32.5 59 Poisson's ratio 0.333 0.318 58 0.338 0.327 7 0.330 0.318 59 Thermal conductivity W/m-K 120 3.3 7 120-160 3-6 50 120-160 3-6 50 Thermal expansion K -l xl0 -6 23 14.1 7 120-160 14.1 7 120-160 14.1 7 Specific heat J/kg • K 900 0.28 7 120-160 0.28 7 0.28 0.28 7 Electrical resistivity M fi cm 5.66 3.03 17 3.94 1.38 17 5.7 2.9 17 using 7-integral procedures and found that toughness increased with decreasing temperature. The K\ C {J) was 47 MPa\/m at 4 K in the transverse (TL) orienta- tion. Shepic [54] measured the toughness of 2219- T87 (3.8 cm thick plate) using the ASTM E399 plane strain fracture toughness testing procedures. The K\ c for 2219-T87 at 4 K was 45.5 MPaVm in the trans- verse (TL), 47.7 MPa\/m in the longitudinal (LT) and 28.6 MPaVm in the through-thickness (SL) orienta- tions. 5.3. Fatigue Strain cycling fatigue tests have been conducted by Nachtigal [13] on 2219-T851 and 2014-T6 alloys at room temperature, 77 and 4 K. For a given strain level, the fatigue lives of both alloys were greater at 4 K than at the higher temperatures for endurances 1 1 1 1 - 2219-T851 I >h Aluminum - ^^ o 300 K - ^< S. D 77 K - - ^V M( - 1 I 1 ^H. - 1 - 10' 10 c 10 J 10 10 J 10" CYCLES TO FAILURE, N. Fig. 11. Fatigue behavior of 221 9-T851 aluminum alloy at temperatures of 300, 77 and 4 K [13]. greater than about 500 cycles. The results for 2219- T851 are shown in fig. 1 1. The Manson-Hirschberg [55] method of universal slopes gave reasonable pre- dictions of fatigue behavior based on the tensile pro- perties at the test temperature. The fatigue crack growth behavior of 5083-0 [53], 2219-T87 [54] and an Al-6Mg [56] alloys have been determined at 4 K. In each case the results indicate that the growth rates at 4 K are slower than the rates at room temperature. 5.4. Physical properties The physical properties of the candidate aluminum alloys are summarized in table 9. The thermal proper- ties have been taken from the best-fit lines used in the LNG Materials Handbook [7] and from the compila- tion on thermal conductivity by Quids, Ericks and Powell [57]. Original sources are given for the elastic properties [58,59] and electrical resistivity [17]. With the exception of magnetic susceptibility, the physical properties of aluminum alloys are not as good as those of stainless steels for fusion magnet applications at 4 K. The elastic modulus is relatively low, and thus when used for coil case applications more of the load must be carried by the conductor assembly. The thermal and electrical conductivities are relatively high, resulting in increased thermal loads particularly from eddy current losses. 6. Fabrication characteristics of aluminum alloys The relative ease of fabricating large aluminum structures is the principal advantage of aluminum 31 H.I. McHenry, R.P. Reed / Structural alloys for superconducting magnets Table 10 Compositions of aluminum welding rods and electrodes for liquid helium service [46] ; composition, % (max unless range given) Type Cu Mg Mn Cr Si Fe Zn Ti Be Other Suggested for welding 1100 0.20 - 05 - Si + Fe=1.0, max 99.0 - -99.5 Al 1100,3003 2319 5.8-6.8 0.02 0.20- -0.40 - 0.20 0.30 0.10 0.10- -0.20 0.0008 0.10- - 0.25 Zr 2219 4043 0.03 0.05 0.05 - 4.5-6.0 0.80 0.10 0.20 0.0008 0.05- - 0.15 V 1100,6061 5183 0.10 4.3-5.2 0.50- -1.0 0.05- -0.25 0.40 0.40 0.25 0.15 0.0008 5083 5356 0.10 4.5-5.5 0.05- -0.20 0.05- -0.20 Si + Fe = 0.5, max 0.10 0.06- -0.20 0.0008 5083,6061 5556 0.10 4.7-5.5 0.5 - -1.0 0.05- -0.20 Si + Fe = 0.4, max 0.25 0.05- -0.20 0.0008 5083,6061 over stainless steels. In addition the lower weight of aluminum structures may facilitate assembly and erection of the large magnet structures anticipated for fusion energy systems. Fabrication characteristics that should received special consideration for fusion magnet systems are briefly reviewed in this section. Table 1 1 6.1. Welding Alloys 5083, 6061 and, to a lesser extent 2219, are readily weldable by the gas metal-arc and gas tungsten-arc processes. The compositions of aluminum welding rods and electrodes for these processes are raoie 1 1 Tensile properties of aluminum alloy weldments at room temperature and 4 K (60) Alloy and Plate Filler Post- Temp. TS JYS Elonga- Reduc- Joint Loca- NTS NTS NTS temper thichness alloy weld (K) tensile joint tion in tion of effi- tion of notch TS JYS (mm) thermal treat- ment strength (MPa) yield strength (MPa) 4D (%) area (%) ciency (%) failure tensile strength (MPa) 5083-0 25 5183 no RT 293 138 21.5 _ 100 WM 308 1.05 2.22 4 381 174 27 37 69 BM 371 0.98 2.14 5083-H321 25 5183 no RT 305 179 14 39 96 HAZ 376 1.23 2.10 4 455 246 9.0 14 80 WM 405 0.89 1.65 5083-H321 25 5356 no RT 286 167 13.5 47 90 WM 371 1.30 2.22 4 455 235 9.0 17 80 HAZ 398 0.87 1.69 5083-H321 25 5556 no RT 306 176 14 36 97 WM 370 1.21 2.10 4 474 238 13 17 83 WM 399 0.84 1.68 6061-T6 25 4043 no RT 214 144 6.0 19 69 HAZ 234 1.10 1.63 4 338 259 4.5 9 63 WM 275 0.81 1.06 yes RT 298 247 11 44 96 BM 396 1.31 1.57 4 452 309 15 16 84 WM 463 1.02 1.50 6061-T6 25 5356 no RT 225 156 8.0 31 73 WM 323 1.44 2.07 4 398 243 13.5 24 74 WM 367 0.92 1.50 yes RT 279 202 9.5 33 90 BM - - - 4 476 307 19 24 89 WM 419 0.88 1.37 2219-T851 32 2319 no RT 225 185 2.0 5 50 HAZ 280 1.24 1.52 4 411 277 2.5 10 62 HAZ 363 0.88 1.31 2219-T62 25 2319 yes RT 395 277 7.5 7 99 HAZ 439 1.11 1.58 4 496 355 3.5 5 80 HAZ 570 1.15 1.61 Note: Location of failures: WM - weld metal; BM - base metal; HAZ - heat affected zone. 32 H.I. McHenry, R.P. Reed / Structural alloys for superconducting magnets summarized in table 10. Annealed 5083 alloy retains full strength in the as-welded condition, but the cold worked and precipitation hardened alloys are weak- ened by welding. Part of the strength of precipitation hardened alloys can be restored by a post-weld heat treatment. The tensile properties of aluminum alloy weldments were determined by Nelson, Kaufman and Wanderer [60] and the results for selected alloys at room tem- perature and 4 K are summarized in table 1 1 . Notice that part of the strength obtained by cold working (5083-H321) and by precipitation hardening (6061- T6 and 22 1 9-T85 1 ) is retained by the welded joint. This strength increase is not reflected in design-code allowable stresses [1 1 j which assume as-welded alu- minum alloys are in the fully annealed condition. The tensile properties of 6061-T6 and 2219-T62 welds that were subsequently solution treated and aged indicate that substantial strengthening can be achieved by post-weld heat treatments. For most fusion magnet structures, the solution treatment and quenching are not considered practical. The data in table 1 1 also indicate that the ductility of 2219 weldments is quite low due to localized deformation in the heat affected zone; the ductility and notch-yield ratio of 5556 and 4043 filler metals (as-welded) are low at 4 K; and the best ductility and notch-yield ratios in the as-welded condition at 4 K are provided by the 5083 and 5356 alloys. 6.2. Heat treatment and temper designations Aluminum alloys are supplied in a wide range of tempers, and for large complex structures, subsequent heat treatments are not generally performed by the fabricator. The tempers are designated by a letter followed by one to three digits. The letter is the basic temper designation: O-annealed, H-strain hardened and T-fhermally treated. For the strain hardened tempers, the first digit indicates the hardening opera- tions: HI -strain hardening only, H2-strain hardened and partially annealed and H3-strain hardened and stabilized. The second digit (e.g. HI 8 is strain-harden- ed by a cold reduction of 75%) indicates the degree of strain hardening and the third digit is sometimes used for specified variations. For the thermally treated tempers, the first digit indicates the specific sequence of basic treatments; for example, T6 designates solu- tion heat treated and then artificially aged and T8 designates solution heat treated, cold worked and then artificially aged. Subsequent digits indicate specific variations, such as X51 stress relieved by stretching as in T651 or T851. The reader is referred to the Aluminum Association Handbook [46] for a complete description of the temper designation sys- tem. 6.3. Castings Aluminum alloy castings can be produced by a variety of processes and generally are satisfactory for nonstructural applications at low temperatures. The ASME Code [11] permits use of Al-5 Si (type 43) and Al-7Si-0.3Mg (type 356-T6) for nonwelded structural applications. Nelson, Kaufman and Wanderer [60] evaluated the tensile properties at 4 K of several combinations of castings welded to wrought alloys; in general, superior properties were obtained when the properties of the cast and wrought components were nearly equal. 6.4. Machining, forming and forging Machining: Aluminum alloys, particularly the heat treated grades, have excellent machinability. Machined and bolted assemblies of high strength aluminum alloys can replace more costly welded stainless steel assemblies. Care should be taken when using steel bolts because the thermal contraction of aluminum alloys on cooling to 4 K exceeds that of steel. Forming: Aluminum alloys have excellent formab- ility in the annealed or the solution treated conditions. Forging: Aluminum alloys have the best forgeabil- ity of all structural alloys because of their very low strength at forging temperatures of 370— 510°C. The size limit on aluminum forgings is significantly greater than that on stainless steels because of the reduced pressures required. 6.5. Inspection Aluminum alloys and their weldments are readily inspectable by conventional methods: dye-penetrant, ultrasonics and radiography. 33 H.I. McHenry, R.P. Reed / Structural alloys for superconducting magnets 7. Alternative alloys Alternative alloy systems do not provide candidate structural materials for superconducting magnets in fusion energy systems. Consider the following sys- tems: Nickel base alloys are ferromagnetic, expensive and frequently difficult to fabricate. However, many of these alloys have excellent combinations of strength and toughness at 4 K and may be suitable for special applications [61]. Titanium alloys are brittle at 4 K, expensive and difficult to fabricate [62]-. Cobalt base superalloys transmute to long half-life reaction products and are expensive and difficult to fabricate. Carbon and alloy steels including the ferritic and martensitic stainless steels are ferromagnetic and brittle at 4 K. Even the cryogenic nickel steels are brittle at 4 K [63] except for a specially processed 12Ni alloy [64]. Copper alloys are acceptable, but generally not competitive with the austenitic stainless steels or aluminum alloys. 8. Summary This review suggests that there are several com- mercially available alloys that are suitable for con- struction of large superconducting magnet structures for fusion energy systems. These alloys include anneal- ed austenitic stainless steels, nitrogen strengthened austenitic stainless steels where greater strength is needed, an annealed Al~ Mg alloy for welded alumi- num construction and precipitation-hardened alumi- num alloys for bolted assemblies. For each of these alloys there are enough experimental data and fabrica- tion experience to provide reasonable assurance that the alloys have sufficient strength, ductility and toughness in the liquid helium environment; that the alloys can be fabricated into support structures for superconducting magnets; and finally, that the fabricated structure will retain adequate properties during operation of the fusion energy device. It should be realized, however, that the depth of understanding, the degree of characterization, and the extent of related experience for these alloys in the fusion-magnet environment are extremely limited. More work is needed on material behavior at 4 K, particularly in the areas of material variability, pro- cessing effects on properties, and the fabrication and mechanical behavior of large, complex, thick-section components. 9. Recommendations The accumulated knowledge of materials for fusion energy systems should eventually result in a code for the design, construction and inspection of fusion mag- nets. Adherence to this code should provide assurance that the magnet system provides the intended field and operates in a safe, reliable and efficient manner. Estab- lishment of such a code for large magnets must await further developments in the technology, particularly the experience gained from prototype systems and greater material characterization. However, parts of the code should be set forth as preliminary guidelines to facilitate the orderly development of materials tech- nology for magnet systems. Design guidelines needed for selection of materials should identify operating loads and conditions and criteria for establishing allowable working stresses for candidate materials. For magnet systems, a descrip- tion of the operating loads should include the anti- cipated load history for normal operation, loads due to credible fault conditions and procedures for treat- ing combined loads. The operating conditions of interest are the thermal environment, radiation dosage, magnetic field profile and design life. Allowable work- ing stresses should account for the materials strength variability, the fatigue strength for a specified scatter factor on fatigue cycles, fracture mechanics analysis and a factor of safety based on uncertainties in the loads and stress analysis. Material procurement and processing should be controlled by a set of specifications which are suffi- cient to preclude the use of materials in the magnet systems which have properties inferior to those as- sumed in design. Inspection specifications and stan- dards are needed to provide a high degree of confi- dence that no defects are present in the raw materials, or introduced during fabrication, which could cause unsatisfactory performance of the magnet system. 34 H.I. McHenry, R.P. Reed / Structural alloys for superconducting magnets References [21 [1] R.P. Reed, A.F. Clark and E.C. van Reuth, in: Advances [22 in Cryogenic Engineering, vol. 22 (Plenum Press, New York, 1976)1. [23 [2] R.P. Reed, F.R. Fickett, M.B. Kasen and H.I. McHenry, in: Materials Studies for Magnetic Fusion Energy Appli- [24 cations at Low Temperatures, vol. 1, NBSIR 78-884 (National Bureau of Standards, Boulder, CO, 1978) [25 243. [3] F.R. Fickett, M.B. Kasen, H.I. McHenry and R.P. Reed, in: Advances in Cryogenic Engineering, vol. 24 (Plenum Press, New York, 1979). [26 [4] F.R. Fickett, Properties of Non-superconducting Techni- cal Solids at Low Temperatures, Proc. Fourth International Conference on Magnet Technology Brookhaven, New York, 1972)498. [27 [5] H.I. McHenry, in: Advances in Cryogenic Engineering, vol. 22 (Plenum Press, New York, 1976) 9. [6] Handbook on Materials for Superconducting Machinery, [28 Metals and Ceramics Information Center, MCIC-HB-04 (Battelle, Columbus, Ohio, 1977). [7] D.B. Mann, LNG Materials and Fluids (National Bureau [29 of Standards, Boulder, CO, 1978). [8] K.G. Brickner and J.D. Defilippi, in: Handbook of [30 Stainless Steels (McGraw-Hill, New York, 1977) 20. [9] Low Temperature and Cryogenic Steels, Materials [31 Manual (U.S. Steel Corp., Pittsburgh, PA, 1966). [10] Materials for Cryogenic Service-Engineering Properties [32 of Austenitic Stainless Steels (International Nickel Limited, London, 1974). [33 [11] ASME Boiler and Pressure Vessel Code, Section VIII, [ 34 Pressure Vessels, (American Society of Mechanical Engineers, New York, 1977). [35 [12] R.L. Tobler, in: Fracture 1977, vol. 3, Proc. Fourth International Conference on Fracture (Waterloo, Canada, 1977)839. [13] A.J. Nachtigall, in: Properties of Materials for Liquefied [36 Natural Gas Tankage, ASTM STP 579 (American Society for Testing and Materials, 1975) 378. [37 [14] J. A. Shepic and F.R. Schwartzberg, in: Materials Studies for Magnetic Fusion Energy Applications at Low Tern- [39 peratures-1, NBSIR 78-884 (National Bureau of Standards, Boulder, CO, 1978) 13. [40 [15] R.L. Tobler and R.P. Reed, in: Advances in Cryogenic [41 Engineering, Vol. 24 (Plenum Press, New York, 1979). [ 16] H.M. Ledbetter, W.F. Weston and E.R. Naimon, J. Appl. [42 Phys. 46 (9) (1975) 3855. [43 [17] A.F. Clark, G.E. Childs and G.H. Wallace, Cryogenics [44 10(1970)295. [ 18] K.R. Efferson and W.J. Leonard, Magnetic Properties [45 of Some Structural Materials Used in Cryogenic Appli- cations; ORNL-4150 (Oak Ridge National Laboratory, [46 Oak Ridge, TN, 1976) 126. [19] Metals Handbook, vol. 1, 8th edn. (American Society [47 for Metals, Metals Park, OH, 1967). [48 [20] G.P. Sanderson and D.T. Llewellyn, J. Iron Steel Inst. Lond. 27(1969) 1129. V.A. Randak, W. Wesseling, H.E. Bock, H. Steinmaurer and L. Faust, Stahl Eisen 91 (1971) 1255. R. Voyer and L. Weil, in: Advances in Cryogenic Engi- neering, vol. 11 (Plenum Press, New York, 1966) 447. D.C. Larbalestier, Private Communication, University of Wisconsin, Madison, WI (1977). C.E. Spaeder, Jr. and W.F. Domis, Project 44.012-062(2), US Steel Corp., Pittsburgh, PA (1969). D.T. Read and R.P. Reed, in: Materials Studies for Mag- netic Fusion Energy Applications at Low Temperatures-1, NBSIR-78-884 (National Bureau of Standards, Boulder, CO, 1978)91. D.B. A. Macmichael, M. Liddle and J.S.H. Ross, Super- conducting AC Generators Trial Rotor Forging Investiga- tion (International Research and Development Co. Ltd., Newcastle upon Tyne, England, 1978). D.T. Read and R.P. Reed, in: Materials Studies for Mag- netic Fusion Energy Low Temperature Applications-II (National Bureau of Standards, Boulder, CO, 1979). R.L. Tobler and R.P. Reed, in: Elastic Plastic Fracture, ASTM STP-000, (American Society for Testing Materials, Philadelphia, PA, forthcoming). P.F. Scardigno, M.S. Thesis, Naval Postgraduate School, Monterey, CA, NTIS AD/A-004555 (1974). P.R. Landon, Unpublished Data, Lawrence Livermore Laboratories, Livermore, Ca (1975). R.D. Masteller, NASACR-82638, Martin Marietta Corp., Denver, CO (1970). CO. Malin, NASA SP-5921 (01), Technology Utiliza- tion Office, NASA, Washington, D.C. (1970). H.M. Ledbetter, Mater. Sci. Eng. 29 (1977) 255. H.M. Ledbetter and E.W. Collings, in: Metal Physics of Stainless Steels (TMS AIME, New York, forthcoming). D.T. Read, H.I. McHenry and P.A. Steinmeyer, in: Materials Studies for Magnetic Fusion Energy Low Temperature Applications, vol. II (National Bureau of Standards, Boulder, CO, 1979). E.R. Szumackowski and H.F. Reid, Weld. J. 57 (1978) 325-s. V.N. Krivobok and R.D. Thomas, Jr., Weld. J. 29 (1950)493-s. J.H. McConnell and R.R. Brady, Chem. Eng. 67 (1960) 125. W.T. DeLong, Weld. J. 53 (1974) 273-s. C.L. Goodzeit, in: Advances in Cryogenic Engineering, vol. 20 (Plenum Press, New York, 1965) 26. D.F. Shaw, Cryogenics 4 (1964) 193. A.B. Miller, Cryogenics 5 (1965) 320. Armco Nitronic 40 Stainless Steel Sheet and Strip (Arm- co Steel Corp., Baltimore, MD, 1976). J.E. Jenson, ed., Forging Industry Handbook (Forging Industry Association, Cleveland, OH, 1966). Aluminum Standards and Data, Aluminum Association, Inc. Washington, D.C. (1976). H.I. McHenry, Met. Progr. 93 (3) (1968) 62. J.G. Kaufman and E.T. Wanderer, in: Aluminum, vol. II: Design and Applications (American Society for Metals, Metals Park, OH, 1967) 297. 35 H.I. McHenry, R.P. Reed I Structural alloys for superconducting magnets [49] Alcoa Aluminum - The Cryogenic Metal, Alcoa, Pitts- burgh, PA (1971). [50] J.G. Kaufman, K.O. Bogardus and E.T. Wanderer, in: Advances in Cryogenic Engineering, vol. 13 (Plenum [58] Press, New York, 1968) 294. [51] J.G. Kaufman and E.W. Johnson, in: Advances in Cry o- [59] genie Engineering, vol. 8 (Plenum Press, New York, 1963)678. [60] [52] M.P. Hanson, G.W. Stickley and H.T. Richards, in: Low Temperature Properties of High-Strength Aircraft and Missile Materials, ASTM STP 287 (American Society [61] for Testing and Materials, Philadelphia, PA, 1961) p. 3. [53] R.L. Tobler and R.P. Reed, J. Eng. Mat. Technol. 99 (1977) 306. [54] J. A. Sphepic, in: Materials Studies for Magnetic Fusion [62] Energy Low Temperature Applications, vol. II (National Bureau of Standards, Boulder, CO, 1979). [55] S.S. Manson, Exp. Mech. 5 (1965) 193. [63] [56] H.I. McHenry, S.E. Naranjo, D.T. Read and R.P. Reed, in: Advances in Cryogenic Engineering, vol. 24 (Plenum Press, New York, 1979). [64] [57] G.E. Childs, L.J. Ericks and R.L. Powell, Thermal Con- ductivity of Solids at Room Temperature and Below, NBS Monograph 131 (National Bureau of Standard, Boulder, CO, 1973). E.R. Naimon, H.M. Ledbetter and W.F. Weston, J. Mater. Sci. 10(1975) 1309. D.T. Read and H.M. Ledbetter, J. Eng. Mater. Technol. 99(1977)181. F.G. Nelson, J.G. Kaufman and E.T. Wanderer, in: Ad- vances in Cryogenic Engineering, vol. 14 (Plenum Press, New York, 1969)71. J.M. Wells, R. Kossowsky, W.A. Logsdon and M.R. Daniel, in: Materials Research for Superconducting Machinery VI, NTIS No. ADA 036919 (National Bureau of Standards, Boulder, CO, 1976). R.L. Tobler, in: Cracks and Fracture, ASTM STP 601 (American Society for Testing and Materials, Philadel- phia, PA, 1976) 346. R.L. Tobler, in: Materials Research for Superconducting Machinery VI, NTIS No. ADA 036919 (National Bureau of Standards, Boulder, CO, 1976). S. Jin, S.K. Hwang and J.W. Morris, Jr., Met. Trans. 6A (1975) 1569. 36 STRENGTH AND TOUGHNESS RELATIONSHIP FOR INTERSTITIALLY STRENGTHENED AISI 304 STAINLESS STEELS AT 4 K R. L. Tobler, D. T. Read, and R. P. Reed National Bureau of Standards Boulder, Colorado ABSTRACT A study was conducted to determine the effects of carbon and nitrogen on the 4-K fracture properties of Fe-18Cr-10Ni austenitic stainless steels having a base composition corresponding to AISI 304. J-integral fracture toughness tests using 24.5-mm-thick compact specimens (TL orientation) were performed in a liquid helium environment on nine steel heats having carbon plus nitrogen (C+N) contents between 0.067 and 0.325 wt.%. The fracture toughness decreased with increasing C+N content. 1/2 The Kx estimates obtained at 4 K ranged from 337 to 123 MPa-m , exhibiting an inverse dependence on tensile yield stress. A computer- aided J-integral test facility was implemented to conduct this study. This new facility improves measurement accuracy, conserves material specimens and testing time, and systematizes test procedures. 37 INTRODUCTION Acceptance of the parameter J, as a measure of the fracture resis- tance of metals has become widespread in recent years. A large volume of J-integral fracture toughnesss data on a variety of candidate materials for low temperature applications has been accumulated. The J-integral test has the advantage of being applicable to smaller specimens (up to 100 times smaller) and tougher materials than the plane strain fracture test (K, test). A disadvantage of the J-integral technique has been its experimental difficulty. Previous test methods have required approximately five specimens and lengthy data reduction procedures [1]. Recently developed single-specimen J-integral test techniques [2] eliminated the multiple-specimen requirement, but the data reduction procedure was still lengthy. In the computer-aided J-integral test, most of the data- reduction procedure is automated; the apparatus produces a plot of the J-integral value versus crack extension from which Jy can be quickly and simply extracted. The AISI 300 series stainless steels are austenitic Fe-Cr-Ni alloys, offering relatively low strength but excellent cryogenic duc- tility and toughness [3]. By virtue of favorable mechanical properties, service history, and availability, AISI 304 is perhaps the most widely used cryogenic alloy in the world. But nitrogen-strengthened grades such as AISI 304 N or AISI 304 LN have attracted attention recently as possible substitutes for AISI 304 in applications demanding higher strength. Nitrogen is a relatively inexpensive and effective strengthener. Nitrogen also stabilizes the austenitic structure and reduces the probability of martensitic transformation when stress is applied at cryogenic temperatures. 39 Carbon has similar effects, but carbon contents in austenitic stainless steels must be held to low levels to prevent sensitization (chromium carbide precipitation) during thermal excursions. Nitrogen-strengthened austenitic stainless steels are gaining popularity, but some applications are hindered by limited availability, limited service experience, lower fracture toughness, and lack of cryogenic design data. Several mechanical property studies were con- ducted on these steels [4-18], but most papers are concerned only with room temperature behavior. In the present study, fracture toughness measurements for nine heats of an Fe-18Cr-10Ni stainless steel are reported at 4 K. The carbon and nitrogen (C+N) contents of the nine heats were varied to enable a systematic study of interstitial concentration The results can be used to predict the strength and toughness of commercial steels of given compositions or to tailor the properties for a particular application by appropriate specification of the steel composition. MATERIALS Nine stainless steel plates were procured from the research lab- oratories of a major steel manufacturer. These plates had a nominal base composition falling within the limits set by ASTM specification A 240 for AISI 304 stainless steel: Fe-18Cr-10Ni-l.5Mn-0.02P-0.02S-0.55Si- 0.2Mo-0.2Cu. The carbon and nitrogen levels varied, however, with carbon at 0.03, 0.06, or 0.09 wt.% and nitrogen at 0.04, 0.12, or 0.24 wt.%. The mill chemical analyses are listed in Table I. The nine plates were produced from three 136-kg (300-1 b) vacuum-induction- melted heats, split with respect to carbon level, and teemed into 76- x 200- x 360-cm hot- topped cast-iron ingot molds. The ingots were then reheated and soaked at 1561 K (2350°F), hot-rolled to 25.4-mm-thick plates, 40 and air cooled. The plates were finally annealed at 1332 ± 7 K (1937° ± 13°F) for 1 h and water quenched. The hardness and grain size measurements are 1 is ted in Table I . METHODS Since elastic-plastic fracture was anticipated in the test materials, the J T parameter was measured. The measurement problem in J T Ic K Ic testing is to obtain the required plot of J-integral, J, versus crack extension, Aa, from the directly measured quantities, which are the load, P, on the specimen and the relative displacement, 6, of two measurement points located on the specimen load line on opposite sides of the notch (Fig. 1). The J value is proportional to the area under the (P, 25J Ic a f In all low temperature fracture tests, partial transformation the paramagnetic austenite to ferromagnetic a' (bcc) martensite was observed (e.g., Ref. 23). The extent of a' transformation was estimated by using a bar-magnet device [24]. Small amounts of e (hep) martensite may also have formed, but this was not measured. RESULTS AND DISCUSSION The computer-aided J-integral methods have now been in use for nine months and have proved quite satisfactory. In comparison with the 46 multiple specimen and non-computer-aided single-specimen techniques pre- viously used, the new methodology has eliminated about 4 h of data reduction per test and has provided a real-time readout of the test progress and quality. The cost of the complete data-acquisition system, including analog-to-digital converters, processor, digital plotter, and CRT terminal is estimated at about $25,000. J T results were obtained for seven of the austenitic stainless Ic steels tested, but the data for the two lower strength heats (with C+N 0.067 and 0.097 \nt.%) were invalid. The maximum J values in these tests exceeded the allowable limits for 24.5-mm-thick specimens. The resistance curves deviated from the blunting line at exceptionally low J values, and the resistance curves obtained could not be analyzed according to conventional techniques because there was no obvious knee at the blunting line-resistance curve intersection. The J T values for the seven steels Ic 2 at 4 K range from 63 to 499 kJ/m , whereas the yield stresses for the same materials at 4 K range from 329 to 1286 MPa (Table II). Higher toughness in these seven steels is associated with lower interstitial C+N contents, and hence lower strength. All of the steels tested exhibited a ductile dimpled fracture mode at 4 K, as shown by the scanning electron fractographs of Figs. 7 and 8. Estimates of the plane strain fracture toughness parameter, K, , which would be observed under linear-elastic conditions, were obtained using the relationship: K Ic (J) =/Jt„ E \l/2 Id - v 2 )j where E is Young's modulus, v is Poisson's ratio, and Kj ( J) denotes an estimate of K- from J, . The values of E and v for the stainless ic lc 47 steels at 4 K were estimated to be 206.8 GPa and 0.3 [25], respectively. The calculated K T estimates are plotted versus C + N content in Fig. 9 and versus yield strength in Fig. 10. As shown, the K T (J) values decrease linearly with interstitial concentration and are inversely related to yield stress. Therefore nitrogen, which is a particularly effective strengthener, also has a strong influence on fracture toughness at cryogenic temperatures. An expression governing the relationship between fracture toughness and yield stress was derived from Fig. 10: K T (J) = 500 - 0.3 a Ic v y 1/2 The data scatter of K y (J) from this expression is ±20 MPa-m The present results compare well with existing data, some of which are also shown for comparison in Fig. 10. The AISI 304 N datum of Read and Reed [5] falls slightly below the present data, reflecting differences owing to metallurgy or measurement techniques. The Fe-Cr-Ni-Mn-N steels tend to have higher nitrogen contents, higher 4-K strength, and lower toughness compared with the Fe-Cr-Ni-N (AISI 304) steels [4]. Yet, the Fe-Cr-Ni-Mn-N steels in Fig. 10 appear to possess higher toughness than would be expected from the extrapolated trend for Fe- Cr-Ni-N steels. It is uncertain whether this is due to metallurgical effects or measurement bias, since the Fe-Cr-Ni-Mn-N data were obtained by other methods including the multispecimen technique, which have greater J T uncertainty. The inverse relationship of a and K T is well known for linear-elastic y Ic fracture toughness data. Now there is appreciable evidence that the inverse relationship holds in the elastic-plastic region as well. 48 Previous 4-K data for a variety of austenitic stainless steels and uncontrolled metallurgical conditions [5] indicated a linear trend of lesser slope than that shown in Fig. 10. It is advantageous that the present data derive from a controlled alloy series in which strength was varied by interstitial content while other compositional and processing variables were constant. The approximate amount of a 1 (bcc) martensite formed during fatigue and fracture at 4 K is plotted as a function of interstitial content in Fig. 11. These results provide only an approximate indication of the extent of phase transformation, because the magnetometer measurements were taken directly from the compact specimen fracture surfaces, which are rough and unmachined. It is clear from Fig. 11 that some martensite forms in all of the steels during testing. The quantity of martensite that forms would be expected to increase with the plastic-zone size. Referring to Fig. 11, more a' is detected in steel specimens of lower strength, higher toughness, and presumably, larger plastic-zone size. It is not clear at this time whether the formation of martensite plays a significant role in the fracture process of these metastable steels. Stress-strain curves of various heats at 295, 195, 76 and 4 K are presented in Figures 12-17. At 4 K, all heats exhibited discontinuous yielding. Both the magnitude and the strain of initiation of the yield drops increased with increasing austenite stability and strength. The inflections of the curves at low temperatures result from the martensitic transformation. Notice in Figure 12 that the effect of C + N at 4 K is to increase the early deformation of austenite and that the ultimate tensile strength remains relatively unchanged. 49 SUMMARY AND CONCLUSIONS 1. New computer-assisted J-integral test methods and facilities have been developed and applied to cryogenic testing. 2. Fracture toughness measurements at 4 K are reported for Fe-18Cr-10Ni austenitic stainless steels having C+N contents ranging from 0.067 to 0.325 wt.%. 3. The fracture toughness at 4 K decreased with increasing C+N content, and the K, values estimated from Jj (ranging from 123 to 336 1 /2 MPa*m ' ) were inversely related to yield stress. ACKNOWLEDGMENTS This research was supported by the Naval Ship Research and Development Center (J. Gudas, coordinator) and by the Office of Fusion Energy, DoE (D. Beard, coordinator). T. Whipple, A. Castillo, D. Burkhalter, and D. Beekman assisted in performing the fracture tests. 50 REFERENCES 1. J. D. Landes and J. A. Begley, "Test Results from J-integral Studies: An Attempt to Establish a J, Testing Procedure," in Fracture Analysis , ASTM STP 560, American Society for Testing and Materials, Philadelphia, PA (1974), pp. 170-186. 2. G. A. Clarke, W. R. Andrews, P. C. Paris, and D. W. Schmidt, "Single Specimen Tests for J\ Determination," in Mechanics of Crack Growth , ASTM STP 590, American Society for Testing and Materials, Philadelphia, PA (1976), pp. 27-42. 3. R. P. Reed, F. R. Fickett, M. B. Kasen, and H. I. McHenry, "Magnetic Fusion Energy Low Temperature Materials Program: A Survey," i n Materials Studies for Magnetic Fusion Energy Appl ications at Low Temperatures-I , NBSIR 78-884, F. R. Fickett and R. P. Reed, Eds., National Bureau of Standards, Boulder, CO (1978), pp. 245-335. 4. D. T. Read and R. P. Reed, "Toughness, Fatigue Crack Growth, and Tensile Properties of Three Nitrogen-Strengthened Stainless Steels Cryogenic Temperatures," in Materials Studies for Magnetic Fusion Energy Applications at Low Temperatures-I , NBSIR 78-884, F. R. Fickett and R. P. Reed, Eds., National Bureau of Standards, Boulder, CO (1978), pp. 91-154. 5. D. T. Read and R. P. Reed, "Fracture and Strength Properties of Selected Austenitic Stainless Steels at Cryogenic Temperatures," i n : Materials Studies for Magnetic Fusion Energy Appl ications at Low Temperatures-II , NBSIR 79-1609, F. R. Fickett and R. P. Reed, Eds., National Bureau of Standards, Boulder, CO (1979), pp. 81-122. 51 9. 10, 11 12, D. T. Read, R. P. Reed, and R. E. Schramm, "Low Temperature Deformation of Fe-18Cr-8Ni Steels," in: Materials Studies for Magnetic Fusion Energy Applications at Low Temperatures-II , NBSIR 79-1609, F. R. Fickett and R. P. Reed, Eds., National Bureau of Standards, Boulder, CO (June 1979), pp. 151-172. D. T. Read and R. P. Reed, "Heating and Strain Rate Effects in AISI 304 L Stainless Steel at 4 K," in: Materials Studies for Magnetic Fusion Energy Applications at Low Temperatures-II , NBSIR 79-1609, F. R. Fickett and R. P. Reed, Eds., National Bureau of Standards, Boulder, CO (1979), pp. 125-147. R. L. Tobler and R. P. Reed, "Tensile and Fracture Behavior of Nitrogen-Strengthened, Chromium-Nickel-Manganese Stainless Steel at Cryogenic Temperatures," in Elastic-Plastic Fracture , ASTM STP 668, American Society for Testing and Materials, Philadelphia, PA (1979), pp. 537-552. P. P. Dessau, "LN 2 and LH 2 Fracture Toughness of Armco Alloy 22-13-5," Report No. TID/SNA-2083, Aerojet Nuclear Systems Company, Sacramento, CA (1971). D. C. Larbalestier and D. Evans, "High Strength Austenitic Stain- less Steels for Cryogenic Use," in Proceedings of Sixth International Cryogenic Engineering Conference , IPC Science and Technology Press Ltd., Guildford, Surrey, England (1976), pp. 345-347. R. B. Gunia and G. R. Woodrow, "Nitrogen Improves Engineering Properties of Chromium-Nickel Stainless Steels," J. Mater. , JMLSA, 5 (1970), pp. 413-430. R. Voyer and L. Weil, "Tensile and Creep Properties of a High Nitrogen Content 18/10 (AISI 304L) Stainless Steel at Cryogenic Temperatures," in Adv. Cryo . Eng. 1 1 , K. D. Timmerhaus, Ed., Plenum, New York (1965), pp. 447-452. 52 13. L.-A. Norstrom, "The Influence of Nitrogen and Grain Size on Yield Strength in Type AISI 316L Austenitic Stainless Steel," Met. Sci_. , 21 (1977), pp. 208-212. 14. M. K. Booker and V. K. Sikka, "Effects of Composition Variables on the Tensile Properties of Type 304 Stainless Steel," ORNL-5353, Oak Ridge National Laboratory, Oak Ridge, TN (1977), pp. 1-38. 15. P. Soo and W. H. Horton, "The Effect of Carbon and Nitrogen on the Short-Term Tensile Behavior of Types 304 and 316 Stainless Steels," Ward NA 3045-2, Westinghouse Electric Corporation, Madison, PA (1973) pp. 1-48. 16. K. J. Irvine, D. T. Lewellyn, and F. B. Pickering, "High Strength Austenitic Stainless Steels," J. Iron Steel Inst ., 199 (1961), pp. 153-175. 17. K. J. Irvine, T. Gladman, and F. B. Pickering, "The Strength of Austenitic Stainless Steels," J. Iron Steel Inst. , 207 (1969), pp. 1017-1028. 18. C. E. Spaeder, Jr. and W. F. Domis, "Cryogenic and Elevated-Temperature Mechanical Properties of a High-Nitrogen Type 304 Stainless Steel," Report No. 44012-062-2, United States Steel Corporation, Pittsburgh, PA (1969). 19. C. W. Fowlkes and R. L. Tobler, "Fracture Testing and Results for a Ti-6A1-4V Alloy at Liquid Helium Temperature," Eng . Fract. Mech. , 8 (1976), pp. 487-500. 20. J. A. Joyce and J. P. Gudas, "Computer Interactive J T Testing of Navy Alloys," in Elastic-Plastic Fracture , ASTM STP 668, American Society for Testing and Materials, Philadelphia, PA (1979), pp. 451-468. 53 21. "Standard Method of Test for Plane Strain Fracture Toughness of Metallic Materials (Designation E399-74)," 1974 Annual Book of AST M Standards , Part 10, American Society for Testing and Materials, Philadelphia, PA (1974), pp. 432-451. 22. J. G. Merkle and H. T. Corten, "A J-Integral Analysis for the Compact Specimen, Considering Axial Force as Well as Bending Effects," J. Pressure Vessel Tech., Trans . ASME , 6 (1974), pp. 1-7. 23. R. P. Reed and C. J. Guntner, "Stress-Induced Martensitic Trans- formations in 18Cr-8Ni Steel," Trans, Metal 1. Soc. AIME 230 (1964), pp. 1713-1720. 24. R. P. Reed and R. P. Mikesell, "The Stability of Austenitic Stain- less Steels at Low Temperatures as Determined by Magnetic Measure- ments," in Adv. Cryo. Eng. , 4, K. D. Timmerhaus, Ed., Plenum, New York (1960), pp. 84-100. 25. H. M. Ledbetter, N. V. Frederick, and M. W. Austin, "Elastic- Constant Variability in Stainless Steel 304," J . Appl . Phys. , 51 (1980), pp. 305-309. 54 LIST OF TABLES Table No. I. Stainless steel compositions (wt.%), hardnesses, and grain sizes II. Yield strength and J-integral toughness results for Fe-18Cr-10Ni steels at 4 K. 55 NiNcorNi^^tOrsin kO cti oj in co i — r-». c^ c\j OOi — i — i— CNJ C\J CNJ CO ooooooooo .— = = C\J : rvjincofoiDai^inco ooooooooo ooooooooo = r— z = C\J : Oz = 0= = 0= = O = = O = O = = C* = = C^ : 00 = = CO = i-csjro^tirnorscocn 56 ■ Table II. Yield strength and J-integral tou ghness results Fe- -18Cr-10Ni steel s at 4 K. Specimen C+N Con tent Yielc 1 Strength, a y Flow Streng th, a* 25 J, c Ic K Ic (J)* (wt.%) (MPa) a (MPa) a f (mm) (kj/m' ■) b (MPa-m 1/2 ) c 1 0.067 329 892 NA NA NA 2 0.097 445 996 NA NA NA 3 0.128 530 1029 12.2 499 337 4 0.157 745 1158 6.7 312 266 5 0.187 876 1288 4.5 230 230 6-A 0.214 896 1264 4.3 218 222 6-B 0.214 896 1264 5.3 272 249 1186 1178 1178 1286 2 2 Calculated from the expression K, = J, *E/(l-v ), assuming E = 206.8 GPa and v = 0.3 at T = 4 K. To convert from MPa to ksi, divide by 6.895. L 9 9 To convert from kJ/m to in-lb/in , multiply by 5.709. C r. — To convert from MPa-Zm to ksi /in, divide by 1.099. 57 LIST OF FIGURES Figure 1. Typical J-integral-versus-crack extension curve. Figure 2. Typical load-versus-specimen displacement curve, indi- cating J-integral area for specimen at point a. Figure 3. Calibration curve of load change versus displacement change. Figure 4. Flow chart of recording and analyses computer system for fracture testing. Figure 5. Specimen geometries used for fracture mechanics tests. Figure 6. Typical J-resistance curve for an austenitic stainless steel tested at 4 K. Figure 7. Scanning electron microscopy photograph of Fe-18Cr-10Ni ; C+N = 0.325 wt.%. The 4-K fracture surface is shown at 200 X and 1000 X. Figure 8. Scanning electron microscopy photograph of Fe-18Cr-10Ni ; C+N = 0.067 wt.%. The 4-K fracture surface is shown at 200 X and 1000 X. Figure 9. Dependence of Ky on C+N content at 4 K. Figure 10. Inverse relationship between fracture toughness and yield stress for austenitic stainless steels at 4 K. Figure 11. Relative amounts of a' (bcc) martensite formed during tests of compact specimens at 4 K. (Arbitrary Units) Figure 12. Engineering stress-strain curves for Fe-18Cr-10Ni steel heat no. 1, representative of low interstitial C+N con- tents. 58 Figure 13. Engineering stress-strain curves for Fe-18Cr-10Ni steel heat no. 4, representative of medium interstitial C+N contents. Figure 14. Engineering stress-strain curves for Fe-18Cr-10Ni steel heat no. 8, representative of high interstitial C+N contents. Figure 15. True stress-strain curves for Fe-19Cr-10Ni steel heat no. 1 (serrations at 4K are not shown). Figure 16. True stress strain curves for Fe-18Cr-10Ni steel heat no. 4 (serrations at 4K are not shown). Figure 17. True stress-strain curves for Fe-18Cr-10Ni steel heat no. 8 (serrations at 4K are not shown). 59 2 S- 3 o c: o ^■"V •r— E ens E ext ^^ -^ o CO 2 < V CO Z3 w* CO z > o 1 ^^ m > t— * o < . DC HI O 13 (ui/N>i) ivaoaiNi-n 60 i •r— -a c: •r~ CD > n3 S- 13 +-> U C ^>» -M c o CL E CD E -M E 03 V-^ a. E 00 • i — 00 •i— cu Cl ms c 00 \- a E S- z •r— O CD o 4- 111 CL CO 03 ai 2 1 CO =3 rt3 LLI 00 n3 O > CD < 1 -a 4-> _l o •r— a. , O CO ■> 03 U CXI cu s- Z3 CD (N>i) d 'avcn 61 CO to 6 o b 10 — lo CO CM (N>i) dV '30NVH0 avcn o < I o LU LU o < Q. CO Q CD oj E n3 E 00 •r— ^^ "° «> 00 :3 < 00 S- CD > t» CD 1 1 l Ol LU c: -a re o o cu > o o 4-> u 62 >% ital tter 4— o CO CO o o l_ Q. O Floppy Disks o 13 a E E CD ^ CO c III O) o E III I— o 05 o o 'c CO > CO Is F -J ^ o <*- O) ^~ 4-> TO CO +-» c >> 00 ?ii < 1° °- ' (D t CD l< cl< c < 1 ?1 ■ | CO o co i. Z5 OV A Chan t Q mV Chan 25V Chan t | Q ! < c > x o c; ■i t i J. < co ; ! 5 O CD C CL o <_> oo ■ O T^ 05 o O) 1 oo ! c O >> III! 1 < 03 i 1 1 i 03 c CD E "O cc ■Di2 C CO CO ca c 03 o « "S _j co "Sr75 CD Disp Sign. o a c o 3 W CO •r— -a i. o o aj CO E cr > _c0 Q. CO 5 >♦— Q. E < L < Analog -Y Plotter Unload and :op Switche sst Syste o 'c o V- -t— < o CD <+- o 4-> S- 03 jC u 2 o ■p 00 QJ SL +3 U X 03 CO LU u_ M- l 1 1 *- i l c . I \ 53- \ \ E OJ s- =5 \ \ O O) CD \ \ CO CO ■r— Ll_ \ aO \ w Q >> V v — \ ading V \/ -rtfr i V. J CO — o OT _, CO < ) < C 63 63.5 Dimensions *n mm Width,W=50.8 Razor Blades COMPACT SPECIMEN Thickness = 24.5 mm 9.53 R TENSILE SPECIMEN Thickness = 2.54 mm Figure 5. Specimen geometries used for fracture mechanics tests G4 1200 000 ^ 800 ; 6oo CD ^ 400 i 200- J- Resistance Curve Fe-l8Cr-l0Ni Stainless Steel (C+N = 0.128 weight percent) Compact Specimen 24.5 mm thick a/W* 0.630 TL Orientation Blunting Line J = 2Aa O Figure 11 . BCC Martensite Formation of CT Specimens T=4K Fracture Surface Fatigue Surface at AK £* 40 MPa-m 1/2 0.10 0.20 0.30 Carbon + Nitrogen (weight percent) Relative amounts of a' (bcc) martensite formed during tests of compact specimens at 4 K. (Arbitrary Units.) 70 1.8 I I I I I I I Fe-18Cr-10Ni Heat No. 1 C+N = 0.067 wt % 1.6 — "V 76* ■ 1.4 — 1.2 CO a. O b 10 A95* — STRESS, o bo | 0.6 295 K 0.4 "" "* 0.2 n I I I I I I I 10 250 20 30 40 50 ENGINEERING STRAIN, € (%) 60 — 200 — 150 o b eo — 100 50 70 Figure 12. Engineering stress-strain curves for Fe-18Cr-10Ni steel heat no. 1, representative of low interstitial C+N con- tents. 71 1.8 1.6 0.2 Figure 13, Fe-18Cr-10Ni Heat No. 4 C + N = 0.157wt % 16 24 32 40 48 ENGINEERING STRAIN, € (%) 56 A 250 200 150 o CO JC CO CO 111 cr i- w 100 — 50 64 72 Engineering stress-strain curves for Fe-13Cr-l ONi steel heat no. 4, representative of medium interstitial C+N contents. 72 2.0 1.8 CO Q- O CO CO LU cr co 0.2 — Figure 14, 10 Fe-18Cr-10Ni Heat No. 8 C + N = 0.297 wt % 20 30 40 50 ENGINEERING STRAIN, € (%) 1 60 250 200 v> 150 ^ CO CO w cr i- co 100 50 70 Engineering stress-strain curves for Fe-18Cr-l ONi steel heat no. 8, representative of high interstitial C+N contents. 73 2.4 2.2 0.2 10 Fe-18Cr-10Ni Heat No. 1 C + N =0.067 wt % 20 50 60 — 250 300 200 jc CO CO LU ec CO LU 150 3 DC 100 50 70 Figure 15, 30 40 TRUE STRAIN, C (%) True stress-strain curves for Fe-19Cr-10Ni steel heat no. 1 (serrations at 4K are not shown). 74 2.4 2.2 0.2 Figure 16, Fe-18Cr-10Ni Heat No. 4 G + N » 0.157 wt % 10 20 30 40 TRUE STRAIN, € {%) 50 60 300 250 ft 05 H 200 w fr- 150 g 100 50 70 True stress strain curves for Fe-18tr-10Ni steel heat no. 4 (serrations at 4K are not shown). 75 2.4 0.2 Figure 17. 10 Fe-18Cr-10Ni Heat No. 8 C+N = 0.297 wt % — 300 250 w 200 5 b co" CO HI DC I- co HI 150 3 100 50 20 30 40 TRUE STRAIN, € (%) 50 60 70 True stress-strain curves for Fe-18Cr-10Ni steel heat no. 8 (serrations at 4K are not shown). TENSILE AND FRACTURE PROPERTIES OF MANGANESE-MODIFIED Fe-18Cr-8Ni TYPE AUSTENITIC STAINLESS STEELS R. L. Tobler and R. P. Reed Fracture and Deformation Division National Bureau of Standards Boulder, Colorado ABSTRACT A series of ten low-carbon AISI-?04-type austenitic stainless steels having 1 to 6% Mn (by weight) and 0.1 to 0.2% N were produced and tested to determine the effect of these elements on properties at 4 K. Tensile tests (at 295, 76, and 4 K) and J-integral fracture toughness tests (at 4 K) were conducted on developmental steels con- taining 18.25 to 19.50% Cr, 7.9 to 8.75% Ni, and 0.02 to 0.03% C. All steels were hot-rolled at 1450 K (2150° F) from ingots to 25.4-mm (1- inch) plates. The 4-K yield strengths ranged from 620 MPa (90 ksi) to 1068 MPa (155 ksi), increasing strongly with nitrogen content. Unac- ceptably low toughness was observed in the low manganese compositions, but the fracture toughnesses of alloys containing 6% Mn were equivalent to those of conventional AISI 304 stainless steels. 77 INTRODUCTION The AISI-300 series stainless steels are austenitic Fe-Cr-Ni alloys offering relatively low strength but excellent cryogenic ductility and toughness. Recently, nitrogen-strengthened grades, such as AISI 304 N or AISI 304 LN, have attracted attention as possible substitutes for appli- cations demanding higher strength. Nitrogen is a relatively inexpensive strengthener, and it stabilizes the austenitic structure with respect to martensitic phase transformation. A disadvantage of the nitrogenated grades is that they are more difficult to produce and fabricate. To overcome this disadvantage, manganese additions have been recommended. The potential advantages of manganese-modified type AISI 304 LN steels are (1) consistent fabricabil ity, (2) consistent weldability, and (3) re- liable supply in all product forms [1]. It was the purpose of this study to investigate the effects of manganese on the 4-K mechanical properties of nitrogenated AISI 304 L type alloys. Therefore, a series of steels containing 1 to 6% Mn and 0.1 to 0.2% N were prepared, tested, and compared with previous results for Fe-18Cr-10Ni stainless steels [2,3] MATERIALS Ten stainless steel plates were obtained from a commercial steel manufacturer. These steels had chemical compositions as listed in table 1 and grain size and hardness values as listed in table 2. All of these steels were produced from 225 N ingots (12.7 x 12.7 x 17.8 cm) which were hot-rolled at 1450 K (2150° F) to plates approxi- mately 25-mm thick. The plates were annealed at 1340 K (1950° F) for one-half hour, then water quenched and acid pickled. After trimming, the average plate size was 25.4 x 140 x 400 mm. 79 Figure 1 shows one example of as-received plate. These steels were free from corner checks (edge tears), which sometimes may occur in nitrogen-alloyed austenitic stainless steels. Figure 2 shows an example of corner checks as they appeared in an Fe-18Cr-10Ni-0.03C-0.024N-l.5Mn stainless steel of a previous study [2]. In the previous study a series of nine Fe-18Cr-10Ni-l .5Mn stainless steels were hot-rolled at 1560 K (2350° F) from ingots to 25.4-mm-thick plates, and corner checks were observed in steels having carbon plus nitrogen (C + N) contents of 0.270 wt.% or greater. EXPERIMENTAL PROCEDURE Tensile Tensile tests were conducted using the cryostat described by Reed [4]. Low temperatures were achieved by immersing the specimen and cryostat in liquid nitrogen at 76 K and liquid helium at 4 K. The specimen geometry is shown in fig. 3. The specimens were machined in the transverse orientation. The load-vs-deflection curves were recorded using the outputs from a commercial load cell and strain-gauge extensom- eter. The specimen yield stress at 0.2% plastic strain, ultimate ten- sile stress, elongation, and reduction of area were measured. Estimated strain sensitivity is 50 ym at 4 K, and low-temperature accuracy is estimated to be ± 500 ym at 4 K. Fracture Toughness J-integral tests were necessary because austenitic stainless steels typically fail in a ductile (elastic-plastic) mode with gradual crack extension at 4 K. Therefore, single-specimen J-integral tests were used to measure fracture toughness. The technique incorporates a computer to enable digital data acquisition and J-resistance curve plotting during 80 the tests, as described previously [2,3]. The fracture toughness parameter measured in these tests is Jt , the critical value of the J- integral at the beginning of crack extension, which is a material property enabling quantitative comparisons of ductile materials. Jt also enables an estimate to be made of K T , which is the plane-strain fracture toughness as conventionally measured by the ASTM E 399 test method. The compact fracture toughness specimens shown in fig. 3 had a 24.6-mm thickness (B) and 50.8-mm width (W). The notch was machined in the TL orientation, as defined in ASTM E-399-78a [5]. The fracture specimens were fatigue cracked at 76 K to a crack length, a, between 27.25 and 31.44 mm (g * 0.535 to 0.620). The maximum load used in fatigue precracking was 22 kN. The specimens were then loaded in displacement control. As the load increased, partial unloadings were performed periodically. The load reduction during unloading was typically ten percent of the maximum load. Using the crack-length-vs- compliance correlation, the crack length at each unloading was inferred and used to obtain the crack extension increment (Aa) from the difference of the initial precrack length. The J value at each unloading point was calculated using the expression: J = gr- where X is the Merkle-Corten factor, A is the area under the load-vs-defection curve at the unloading point, and b is the specimen ligament (b = W - a). J-resistance curves consisting of many sets of J and Aa points were generated. The J T value for each steel was taken at the intersection of the J-resistance curve with the blunting line. The blunting line is given by J = 2a f A where a f is the flow stress, which is equal to the 31 average of the yield and ultimate tensile strengths. It was discovered that the final crack extensions, Aa, measured on the fracture specimens, were 20 to 200% greater than the values inferred by the computer. The error between actual and inferred crack extensions increases as Aa increases, and is lower at the J T measurement point. Owing to this bias, the J T values reported here are estimated to be 15% higher than would be expected from multiple specimen tests. Scatter in the J- resistance curves led to uncertainties of ± 15% in the values of J T . Ic The validity of J. results was determined by the size criterion, which requires that 25J B, b, a > — -£- Previous results have shown that valid results at 295 K are impossible for annealed austenitic stainless steels of 25-mm thickness, so a com- plete temperature dependence of J, was not attempted. Only 4-K values are reported. The 4-K J T values were converted to K T estimates, denoted Kj c (J)» using the following equation: 1/2 K Ic (J) = (S)' where E is Young's modulus and v is Poisson's ratio. The values of E and v at 4 K were estimated as 206.8 GPa and 0.30, respectively. RESULTS AND DISCUSSION Tensile Properties The tensile properties of the steels tested in this study are listed in table 3 and plotted in figs. 4 through 6. As shown in fig. 4, the yield strength (a ) data for the present Fe-18.25 to 19.5Cr-7.90 to 8.75Ni steels (open symbols) agree well with previous data trends for Fe-18Cr-10Ni steels (solid lines). The yield 82 strength is highly sensitive to interstitial concentration, so differ- ences in other metallurgical parameters (such as processing temperature, grain size, and nickel content) can be neglected for this comparison. The 295-K data confirm a previous conclusion: the yield strength is a linear function of C+N content [2]. The 4-K data also confirm the existence of a kink in the a -versus-C+N plot at C+N contents between 0.127 and 0.157 wt.%. In the previous study, this kink was attributed to the structural instability of the austenitic phase at these inter- stitial contents [2]. In the present study there is a clear trend of increasing a values with higher Mn contents at C+N contents from 0.131 to 0.138 wt.%. This effect is probably due to the stabilizing effect of Mn additions at this critical interstitial content. The ultimate tensile strength results of the present study, shown in fig. 5, are also in general agreement with expected trends. The data at room temperature fall on the same trend line that was previously reported for the Fe-18Cr-10Ni stainless steels. The 4-K data show a minor variation owing to the austenite stabilizing effect of Mn modi- fications. As more manganese is added at a given interstitial con- centration level, the austenite is progressively more stable during tests at 4 K; as a result, less martensite is formed during testing and ultimate tensile strength is reduced as Mn concentration is increased to 6%. Figure 6 shows the spread of the elongation and reduction of area measurements for the ten steels. Not all of the ductility values are typical of commercial stainless steels. In particular, the reduction of area values for the developmental steels containing 1 to 4% Mn are lower compared with values previously reported for Fe-18Cr-10Ni steels [2] or 33 commercially available AISI 304 LN steels. Only the 6% Mn-modified steels possess ductility values at the expected levels. Fracture Toughness The Jt results obtained for one specimen of each of the ten dif- ferent steel compositions tested are reported in table 4, and the cor- responding Kr (J.) values are plotted in fig. 7. As shown in fig. 7, the Kj ( J) estimates follow regular trends depending on Mn and N content. Solid lines have been drawn to connect data representing constant Mn contents; for each Mn content the 0.1N composition is represented by the symbol at the lowest yield strength, whereas the 0.2N composition is represented by the higher yield strength, and 0.15N is intermediate. From fig. 7 it is clear that higher N at a given Mn content lowers the K T (J) value. It is also clear that Mn has a favorable effect, raising both the strength and the toughness of these developmental alloys. The major observation from fig. 7 is that the 6%-Mn-modified steels have acceptable fracture toughness, whereas the modifications containing 1 to 4% Mn display inferior properties. The 1 and 2% Mn steels are 1 17 particularly poor, having estimated Kr values as low as 117 MPa-m . Their load-deflection fracture toughness test records showed a nearly linear-elastic behavior, with regions of fast crack advance (pop-ins). This evaluation is based on comparison of the present results with the previous results for a series of nine Fe-18Cr-10Ni stainless steels con- taining variable C+N contents (for all of these steels the same J T test procedures were used). As shown in fig. 7, the 6% Mn-modified steels exhibit an inverse K T (J)/a relationship at 4 K and, within the mea- J-c y surement error indicated by the error bars shown in the figure, the 34 inverse-linear trend falls along the lower bound of the scatter band of data for the Fe-18Cr-10Ni steels. The cause of the unexpected, low K* (J) values of the 1 to 4% Mn grades is not clear, but it may be related to chemical composition or chemical segregation. The nominal compositional variations between the 1 to 2% Mn steels of this study and the Fe-18Cr-10Ni-l .5Mn steels of the previous study include a 1.25 to 2.5% difference in nickel content. Metallurgical factors, such as interstitial segregation, may account for the low fracture toughness of the 1 to 4% Mn steels of the present study. However, light microscopy revealed no evidence of sensitization or other obvious microstructural differences. Evidence for interstitial (nitrogen) segregation to grain boundaries in the 1, 2, and 4% Mn steels was obtained by scanning electron micros- copy performed on the fracture surfaces of the compact specimens subse- quent to Jt testing at 4 K. Figure 8 presents typical results. The specimens containing 4% Mn or less exhibited a significant portion of brittle intercrystalline fracture, whereas the three heats containing 6% Mn (heats 72, 76, 78) all showed a more completely dimpled fracture mode. The lower manganese heats (69, 70, 73, and 74) also show more intercrystalline fracture than was observed in the Fe-18Cr-10Ni steels of our previous study. The only apparent distinction between the thermal -mechanical treatments of the two alloy series was the tempera- ture of hot-working of the billet. The hot-rolling temperature of the initial 304 LN alloy series was 1560 K (2350° F) and the rolling tem- perature of the 1-6 wt.% Mn alloy series was 1450 K (2150° F). Perhaps this higher deformation temperature is required to preclude excessive 85 nitrogen segregation to grain boundaries. Further research is con- tinuing to identify the distinctions between the two alloy series. CONCLUSION The data of this study show that 6% Mn-modified AISI 304 LN stain- less steel can be produced, having conventional tensile and fracture toughness properties equivalent to standard AISI 304 LN stainless steel when tested at 4 K. The Mn-modified stainless steel offers potential processing and fabrication advantages when compared with the standard grade, which contains 1 to 2% Mn. ACKNOWLEDGMENTS The stainless steel plates were produced by Armco, under the di- rection of R. H. Espy of Armco' s Research and Technology division. Dr. E. L. Brown of NBS contributed valuable metal lographic assistance and technical advice, and D. Baxter of NBS conducted many of the tensile and fracture toughness tests. This assistance is gratefully appreciated. REFERENCES [1] Espy, R. H., "Stainless steel alloy development," presented at NBS/DoE Workshop on Materials at Low Temperatures, Vail, CO, October 16-18, 1979. [2] Tobler, R. L., and Reed, R. P., "Interstitial carbon and nitrogen effects on the tensile and fracture parameters of AISI 304 stain- less steels," in Materials Studies for Magnetic Fusion Energy Applications at Low Temperatures- II I . NBSIR 80-1627, R. P. Reed, ed., National Bureau of Standards, Boulder, CO, June 1980, pp. 17-48. 86 [3] Tobler, R. L., Read, D. T., and Reed, R. P., "Strength and tough- ness relationship for interstitially strengthened AISI 304 stain- less steels at 4 K," in Materials Studies for Magnetic Fusion Energy Applications at Low Temperatures-IV . NBSIR 81-1645, R. P. Reed and N. J. Simon, eds . , National Bureau of Standards, Boulder, CO, June 1981 , pp. 37-76. [4] R. P. Reed, "A Cryostat for Tensile Tests in the Temperature Range 300 to 4 K," in Advances in Cryogenic Engineering , Vol. 7, Plenum Press, NY (1961), pp. 448-454. [5] Standard Method of Test for Plane Strain Fracture Toughness of Metallic Materials (Designation E399-74) , in 1974 Annual Book of ASTM Standards , Part 10, American Society for Testing and Mate- rials, Philadelphia, PA (1974), pp. 432-451. 37 cu a> +-> GO CO 00 a> rO +-> CO «3- o oo CD a. 4-> XJ CD ■a o o 00 o -M 00 O Q. 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CO CO in co co co in in CM co CM r^ r— in o CO CD in in o o in r-^ co CD O r> CD in CD O >> • -Q ■— . >> >> in r— jQ CD Q. 00 •i— CD • ■4-J "O in r— •r— 3 > >1 E •i — -Q A T3 CU CM #* T3 c 1 = •i — •r— kr > "^ s* •i— -Q • "O ■"" •r— to •> C jx: •r— 'i— L0 o -Si o 4-> •4-J o 4-> CM Lf E • » t> S- S_ S- — a» CD CU |CM > > > C c II ° o o 11 o o u t> 1— o 1— o 1 — 91 LIST OF FIGURES Figure 1. As-received 25 mm-thick stainless steel plate (heat 76, Fe-19.5Cr-7.83Ni-5.8 Mn), showing freedom from corner checks. Figure 2. Corner checks as they appear in hot-rolled nitrogen-modified Fe-18Cr-10Ni stainless steel of a previous study (see text). Figure 3. Compact fracture and tensile specimens used in this study. Figure 4. Effect of carbon plus nitrogen interstitials on yield strength of stainless steels, with temperature as a parameter. Figure 5. Effect of carbon plus nitrogen interstitials on ultimate strength at two test temperatures. Figure 6. Ductility of stainless steels. Figure 7. Fracture toughness-versus-yield strength results for Mn-N modified steels of the present study, as compared with previous results for AISI 304 type stainless steels. Figure 8. Scanning electron micrographs of 4 K fracture surfaces of low toughness (heat 69, top) and high toughness (heat 72, bottom) austenitic stainless steels. ■••• ■-. '. •'? K. 11 ■ i«; J ^ i >& (; »•>. i i CO r-» s- +-> cu (O e cu s- 4C o — * o co E CD o +-> S- ro «t- ^« Q. E o !— •o CU cu CU cu +-> s_ CO ct * — V -^ c <_> s: •r— -£T CO 4-> • 1 ir> E i E •p* 2: LO ro CNJ CO -o 1^. CU 1 > S- •r— O CO ID o CO cu CD -i<£ s- 1— O 1 1 CU CO CU -C < U- O ,-' CU S- rj CD 93 T3 CD , , •r- +j M- X •r— QJ -o +-> o E OJ 1 OJ c CO OJ * — -- en o >> s- -a +-> 13 •i— +-> c CO "O CO ai 3 p— O r— •r— o > S- OJ 1 S- +J a. o jC fO c M- •r— o S- r __ (O OJ CD OJ CL +-> Q. co fC CO >> CO CD OJ jC r— +J c • 1 — CO fO fO +-> CO CO .^ •r— a -21 CD O .£: r— o | s- s- o 0) co c r— S- I o OJ O u_ CM OJ £_ =3 ai 63.5 Width,W=50.8 28.58 2.70 Dimensions in mm 13.08 Dia. Razor Blades COMPACT SPECIMEN Thickness = 24.5 mm 9.53 R TENSILE SPECIMEN Thickness = 2.54 mm Figure 3. Compact fracture and tensile specimens used in this study 95 1,500 a Q_ b I I- o Ld (T h- C/) Q _l UJ >- ,000 500 - 0.10 0.20 0.30 CARBON + NITROGEN (wt.%) Figure 4. Effect of carbon plus nitrogen interstitials on yield strength of stainless steels, with temperature as a parameter. 2,000 a! 1,800 . 1,600 h x e> lu 1,400 Ql h- UJ z UJ UJ 1,200 1,000 - 1 Fe-l8Cr- I I -8tolONi Steels - — y — 1 Mn-y 2Mn \l — 4Mn- ;: J=U^ — - 6Mn^ Previous Study - - l T=4 K I I ^ 800 3 600 400 0.I0 0.20 0.30 CARBON + NITROGEN (wt.%) Figure 5. Effect of carbon plus nitrogen interstitial s on ultimate strength at two test temperatures. 97 ^ 80 — 70 < £ 60 < u. 50 O Z 40 O P 30 o q 20 LU * 10 Fe-l8Cr-8tolONi Steels — This study 100 200 TEMPERATURE (K) 300 Figure 6. Ductility of stainless steels 93 350 CM ^ 300 E o 250 £ 200 in LU X <£> O UJ cr ID O < 50 100 - 50 - Previous Fe-l8Cr- lONi Results 400 800 1200 YIELD STRENGTH, Oy(MPa) 1600 Figure 7. Fracture toughness-versus-yield strength results for Mn-N modified steels of the present study, as compared with previous results for AISI 304 type stainless steels. 99 (a) 100 ym (b) 100 ym Figure 8. Scanning electron micrographs of 4 K fracture surfaces of low toughness (heat 69, top) and high toughness (heat 72, bottom) austenitic stainless steels. 100 INTERSTITIAL CARBON AND NITROGEN EFFECTS ON THE CRYOGENIC FATIGUE CRACK GROWTH OF AISI 304 TYPE STAINLESS STEELS R. L. Tobler and R. P. Reed Fracture and Deformation Division National Bureau of Standards Boulder, Colorado ABSTRACT Constant-load-amplitude fatigue crack propagation (FCP) rate meas- urements are reported for AISI 304 type stainless steels having variable carbon-plus-nitrogen (C+N) contents. Rates at stress-intensity factors 1/2 from 20 to 70 MPa«m ' were measured using 25-mm-thick compact speci- mens. The Fe-18Cr-10Ni steels tested exhibited partial martensitic phase transformations during tests at 76 and 4 K, but not during tests at 295 K. At high C+N contents (0.187 wt.% or greater), the FCP re- sistance was lower at 4 K than at 295 K. At low C+N contents (0.067 wt.%), the FCP resistance at cryogenic temperature was significantly higher than at room temperature. The improvement in steels of low C+N contents was associated with a transition from partially intergranular failure at 295 K to completely transgranular failure at 76 and 4 K. 101 INTRODUCTION When James [1] reviewed the subject of fatigue crack propagation in austenitic stainless steels in 1976, few cryogenic data were available. Since then, fatigue crack propagation rate measurements (da/dN) at intermediate stress intensity factor ranges (aK) have shown that stable (nontransforming) austenitic stainless steels display improved resist- ance to fatigue cracking at extreme cryogenic temperatures, whereas metastable grades are less predictable [2-5]. A metastable stainless steel is one in which the austenitic phase in response to stress or plastic deformation at low temperatures transforms to some degree to body centered cubic (a 1 ) or hexagonal (e) martensitic phases. Metastable AISI 304 stainless steel (Fe-18 to 20O-8 to 12Ni) is a common cryogenic structural alloy, whose mechanical properties are highly sensitive to composition; a broad range of mechanical property combinations can be achieved by varying the C+N contents, because these elements significantly stabilize and strengthen the austenitic phase. A previous study described the effects of C+N contents on the 4-K fracture toughness properties of a series of Fe-18Cr-10Ni developmental stainless steels [6]. In the present study, the same series of steels was tested to evaluate fatigue crack propagation behavior. MATERIALS Nine stainless steel alloy plates with an Fe-18Cr-10Ni base were tested. These plates had a nominal base composition falling within the limits set by ASTM specification A 240 for AISI 304 stainless steel. The C+N levels varied, however, with carbon at 0.03, 0.06, or 0.09 wt.%, and nitrogen at 0.04, 0.12, or 0.24 wt.%. The nine plates were produced 103 from three 136-kg (300-1 b) vacuum-induction-melted heats, split with respect to carbon level, and teemed into 76 x 200 x 360 cm hot-topped cast-iron ingot molds. The ingots were then reheated and soaked at 1561 K (2350° F), hot-rolled to 25.4-mm-thick plates, and air cooled. The plates were finally annealed at 1332 ± 7 K (1937° ± 13° F) for 1 h and water quenched. The mill chemical analyses, hardnesses, and grain size measurements are listed in Table I. Tensile properties are listed in Table II . SPECIMEN The modified compact specimen geometry of Figure 1 was used (the basic geometry is described in ASTM Method E-399-74 [7]). The modi- fication consisted of mounting knife edges for clip-gage deflection (6) measurements at the load line. The specimen thickness (B) was 25 mm, and the width-to-thickness ratio (W/B) was 2. The notch was machined in the TL orientation [7], and the crack propagated in a direction parallel to the plate rolling direction. PROCEDURES The fatigue tests were performed using a 100-kN servo-hydraulic testing machine and cryostat [8]. Room-temperature tests were conducted in ambient air (295 K) at a relative humidity of 10 to 20%. Cryogenic tests were performed with the specimens submerged in a stainless steel Dewar containing liquid nitrogen (76 K) or liquid helium (4 K) . Specimens tested at 295 and 76 K were precracked at the same tem- peratures, whereas specimens tested at 4 K were precracked at 76 K. During precracking at 76 K, the number of cycles to initiate a 1 .9 -mm crack was recorded. The machine was operated in load control, using a 104 sinusoidal load cycle at frequencies between 20 and 25 Hz. The load, P, was accurate to ± 3%. The load ratio, R (P m j M /P_ aw )» was constant at mi n max 0.1. The maximum load for a given test was held constant at a value between 22 and 24 kN. Crack growth-to-a/W ratios of about 0.65 were monitored by specimen compliance (6/P) measurements. The compliance measurements were correlated with the average crack lengths based on measurements at the specimen edges, quarter points, and center of thickness [5,8]. Crack curvature was such that the average crack length was 93% of the center crack length for a/W ^ 0.6. Robert's crack length-compliance solution was corrected (shifted a few percent to agree with the experimental measurements for the final crack length for each test specimen) and then used to infer all other crack lengths from compliance readings taken during fatigue tests [5,8]. A computer was used to fit the a-versus-N curves with third order polynomial expressions, to differentiate the curves to obtain da/dN, and to calculate the corresponding stress intensity factor ranges, AK: AK = K -K . = (P -P . )B"V 1/2 [f(a/W)] (1) max mm v max min' L w yj v ; where f(a/W) is given by Roberts [9]. Data were obtained from one specimen per material per temperature. Magnetic measurements verified that cyclic loading of the Fe-18Cr- lONi steels during crack growth experiments at 76 K and 4 K caused formation of a 1 martensite. The a 1 martensite is ferromagnetic, whereas the original austenite and e martensite are not. Therefore, by using a torsion bar magnetometer [6] on the compact specimen fatigue surfaces 105 subsequent to testing, the amount of a' martensitic transformation was determined. RESULTS Crack initiation at 76 K The number of load cycles required to produce a 1.9-mm crack at 76 K for the particular specimen geometry and loading conditions used in the present tests are plotted in Figure 2. As C+N contents are in- creased, the resistance to crack initiation decreases rapidly at first, but reaches a plateau at C+N = 0.157 or greater. The effect is large, amounting to a five-fold variation of initiation time. Crack propagation For the stress intensity factor ranges studied, the FCP rates for Fe-18Cr-10Ni stainless steels conformed to power law equations of the form: da/dN = C(AK) n (2) where the parameters C and n are material and temperature dependent. Evidence justifying the applicability of this equation to austenitic stainless steels, and to AISI 304 in particular, is presented by James [1] and Brose and Dowling [10]. In the present study, log-log plots of da/dN versus AK were constructed to obtain the C and n values that correspond to the ordinate intercepts at AK = 1 and the slopes, respectively. Table III lists the C and n values for all material and temperature combinations. A general presentation of results is shown in Figure 3. The room- temperature data (Fig. 3a) for the present steels form a scatter band 106 that fits within the larger scatter band drawn by James [1] for many commercial AISI 304 steels tested at different laboratories using ten different specimen geometries and variations of test procedure. The data for the Fe-18Cr-10Ni steel of our study having 0.067 wt.% C+N was also found to be in close agreement with previous measurements for a commercial AISI 304L steel of similar composition [5]. Therefore, our results for laboratory heats are representative of commercial steels. Figures 3b and 3c show the effect of temperature reductions on the spread of present results. Reducing the temperature to 76 K improves the FCP resistance of many of these steels. At 4 K, however, the high C+N compositions are adversely affected. As indicated in Figure 3c, the 4-K scatter band encompasses the 295-K scatter band for the same steels. The greater variation of fatigue crack growth rates at 4 K is attended by significant differences in austenitic stability and tensile proper- ties (Table II) for these alloys, owing to the effects of C+N at cryo- genic temperatures. Interstitial effects Figure 4 presents the results for each heat at 295, 76, or 4 K. Four heats were tested at 295 K. As shown in Figure 4a, the rates for heats 1 and 3 are equivalent. The rates for heats 5 and 9 are also equivalent, but lower compared with those of heats 1 and 3. Therefore, higher C+N contents tend to improve fatigue propagation resistance at 295 K. Individual datum points are shown only for heats 1 and 9. The specified C+N contents for these two steels (0.067 and 0.325 vit.%) rep- resent the limits studied, and the indicated data spreads are typical. 107 Also shown in Figure 4a are results for an AISI 310 stainless steel (Fe-25Cr-20Ni) [5], which fall among the Fe-18Cr-10Ni results. At 295 K, all of these steels are stable with respect to the austenite-to- martensite phase transformation, and all exhibit similar microfailure modes. As shown in Figure 5, the failure modes at 295 K consist of significant portions of intergranular failure. Scanning electron mi- crographs were taken from the fatigue surfaces where the aK value was 1/2 approximately 40 MPa-m ' . In each case, the failures show a number of features, including secondary cracks, coarse transgranular failure zones, and smooth, highly reflective grain-boundary facets. Occa- sionally, a sample would show delaminations. The exposed grain facets frequently displayed slip-line markings, sometimes twins. At 76 K (Figure 4b), the effect of C+N content on fatigue crack growth rates is reversed from that at 295 K. The rates for heat 1 are considerably lower than for heats 3, 6, 7, and 9, which are nearly equivalent. Thus, the steel with the lowest C+N level has the lowest rates at 76 K. At 76 K, partial martensitic transformations take place, and transformation is most extensive for heat 1. Significantly, there is also an obvious difference in microfracture mode for heat 1. As shown in Figure 6, the 76-K fatigue surface reveals a completely trans- granular microfailure, quite distinct from the partially intergranular mode exhibited by this steel at 295 K. As a result of this transition to a transgranular failure, the 76 K fatigue surfaces for heat 1 are noticeably smoother than those at 295 K. In Figure 6, the tendency to transgranular failure is also somewhat in evidence for heat 3. In fact, as the C+N contents are increased, grain facets gradually become more 103 prominent on the 76-K fatigue surfaces. Heats 6, 7, and 9 at 76 K exhibit the same degree of intergranular cracking observed in tests at 295 K. At 4 K (Figure 4c), where all heats were tested, the results were similar to those at 76 K. Again, heat 1 exhibits the lowest fatigue crack growth rates. As C+N contents are increased, the 4-K rates in- crease rapidly at first 3 but the effect saturates at C+N * 0.128 or more, and heats 4 through 9 have equivalent fatigue crack growth re- sistances at this temperature. The relative amount of a' (bcc) martensite formed during fatigue testing at 4 K is plotted as a function of interstitial content in Figure 7. These data provide only an approximate indication of the extent of phase transformation, because the magnetometer readings were taken directly from the compact specimen fatigue surfaces, which are unmachined. From Figure 7, it is clear that some martensite forms in all of the alloys fatigued at 4 K. But more martensite forms in the low-strength, low-C+N grades, because these grades are the least stable and also have the largest crack-tip plastic-zone sizes. Scanning electron microscopy results for specimens fatigued at 4 K are shown in Figure 8. The findings closely parallel those at 76 K: at low C+N contents, the failure mode is completely transgranular , but at high C+N contents there is a partially intergranular failure mode. Temperature effects The results for three heats representing low, intermediate, and high C+N contents are replotted in Figure 9, illustrating the tempera- ture dependence of fatigue crack growth resistance. 109 Heat 1 (C+N = 0.067 wt.%) represents the lowest interstitial content, lowest yield strength, and lowest austenite stability. For this steel, temperature reductions to 76 or 4 K greatly improve the da/dN values at a given AK value. The fatigue crack growth rates at 76 and 4 K are nearly equivalent and much lower than those at 295 K. The rates for heat 1 are in excellent agreement with rates previously re- ported for a commercial AISI 304L and 310 stainless steels [5] at 76 and 4 K. As indicated previously in Figure 5, the improved fatigue crack growth resistance at cryogenic temperatures is associated with a transi- tion from partially intergranular failure (295 K) to transgranular failure (76 and 4 K) . Heat 3 (C+N = 0.128 wt.%) represents intermediate interstitial content, strength, and austenite stability. As shown in Figure 9b, the rates at a given AK value decrease slightly between 295 and 76 K and then increase between 76 and 4 K. Thus, temperature effects are mini- mized at intermediate interstitial content. This concurs with previous reports of temperature-insensitive fatigue crack growth behavic In commercial AISI 304 grade steel [5]. The scanning electron micros ;y results for heat 3 show a more well-defined intergranular failure mode at 295 and 4 K, compared with that at 76 K. Heat 9 (C+N = 0.325 wt.%) represents the highest interstitial content, highest strength, and highest austenite stability. The effect of a temperature reduction to 4 K on this heat is completely opposite to the trend for heat 1, shown in Figure 9c. The 76-K rates are slightly higher or nearly equivalent to the rates at 295 K, but the 4-K rates are two or three times higher. The failure mode for heat 9 is partially in- tergranular at all test temperatures. 110 DISCUSSION It was concluded from previous work that desirable fatigue crack growth resistance at cryogenic temperatures may be achieved with alloys having the austenitic structure and high elastic modulus [5]. Neverthe- less, as shown in this paper, a broad range of properties may be dis- played within a given austenitic steel family at 4 K, owing to compositional variations. At room temperature, where the da/dN data represent crack propa- gation through stable austenite, increasing the C+N level tends to in- crease fatigue crack growth resistance, in accordance with previous observations [2], although heat-to-heat variations are not very great. The microfailure mode at room temperature always consists of a rough transgranular surface plus a significant portion of intergranular failure. The grain facets corresponding to the intergranular failure are undoubtedly the reflective "cleavage-like" facets observed under low magnification in previous studies of austenitic steels [5]. At cryogenic temperatures, where the strengths of some of the steels were greatly increased and where martensitic transformations occurred in varying degrees, the variation of fatigue-crack-growth resistances was especially great. During tests of these steels at 76 and at 4 K, it can be assumed that the martensitic transformation oc- curred locally in the plastic zone ahead of the crack, so that trans- formation preceded material separation. The parameter measured in these tests was the rate of fatigue crack propagation through partially mar- tensitic material . At low C+N levels, the effects of martensitic transformation are complex: strength is low, strain-hardening capability is high, and a relatively large portion of martensite is formed. Since the martensite 111 transformation is an energy-absorbing process, a decrease of crack propagation rates may result, as compared to rates under conditions where the austenite is more stable. At higher C+N contents, the strength level is high and strain hardening capability is low. The amount of phase transformation is reduced, and its effects may be masked by other factors, such as a growing tendency to intergranular fracture. Intergranular fracture may be encouraged at higher C+N contents if the grain boundary energy is reduced by absorbed interstitials or if the grains are hardened relative to the grain boundaries. These results must be interpreted cautiously, because it is believed that the cyclic stress-strain behavior, not the static tensile properties, govern fatigue behavior [2]. In our work, the effects of C+N contents and martensitic transformation on cyclic stress-strain properties were not studied. The effect of the volumetric expansion accompanying martensitic transformation is also unknown. The volumetric expansion in austenitic stainless steels is about 3%, and this should introduce residual stresses. Separation of the temperature, transformation, strain hardening, and residual stress effects is not easy and is beyond the scope of this study. A primary purpose of the present work is to compare the fatigue crack growth rates of steels that may be selected for large cryogenic machinery. Nine metastable Fe-18Cr-10Ni steels and an AISI 310 steel are compared. The two alloys having superior fatigue crack growth resistances are the fully stable (AISI 310) and the least stable (AISI 304, heat 1). The 76- and 4-K rates of these two (Fig. 4) are 112 equivalent and less than the rates at 295 K. The equivalent perfor- mances at extreme cryogenic temperatures are surprising since the im- provement of the metastable steel is associated with a failure mode transition, whereas the improvement for AISI 310 must be attributed solely to inherent temperature effects on material properties (i.e., there is no failure mode transition for AISI 310). Fatigue crack propagation at 4 K may be an important design pa- rameter for superconducting magnet structures in some magnetic fusion energy devices. The rates reported here were determined primarily for material characterization, but the results may also be useful for life prediction, with precautions. Owing to the complex nature of the steels tested, the data trends of Table III should not be extrapolated beyond the range of actual data accumulation. The rates in the present study were measured at R = 0.1, at constant load amplitude, and at interme- diate aK. These conditions are not representative of many applications. Before the fatigue life of a component can be accurately predicted, the effects of stress ratio, variable amplitude loading, cycle frequency, and deviations from the power law equation at threshold or yery high AK levels need to be evaluated. These effects have never been studied at 4 K. SUMMARY AND CONCLUSION A study was conducted to determine the effects of interstitial C+N on fatigue crack propagation resistance of AISI 304 type stainless steels at room and cryogenic temperatures. Rate measurements were reported for conditions of constant load amplitude, at a stress ratio of 0.1, and at intermediate AK values, using compact specimens of the TL 113 orientation, taken from 25-mm-thick plate. The C+N content determines the temperature dependence of fatigue crack propagation rates in these alloys : (1) At low C+N contents, the rates decrease at cryogenic tempera- tures, relative to those at 295 K, (2) At intermediate C+N contents, the rates are relatively in- sensitive to temperature, and (3) At high C+N contents, the rates increase at cryogenic tem- peratures, relative to those at 295 K. The improvement for low C+N contents is associated with a failure mode transition from partially intergranular failure at 295 K to transgran- ular failure at 76 and 4 K. 114 REFERENCES 1. James, L. A., "Fatigue crack propagation in austenitic stainless steels," At. Energy Rev. 14, 1 (1976), 37-85. 2. Pineau, A. G., and Pelloux, R. M., "Influence of strain-induced martensite transformations on fatigue crack growth rates in stain- less steels," Metal 1. Trans., 5 (1974) 1103-1112. 3. Logsdon, W. A., Wells, J. M., and Kossowsky, R., "Fracture me- chanics properties of austenitic stainless steels for advanced cryogenic applications," in: Proceedings of the Second Inter - national Conference on Mechanical Behavior of Materials , American Society for Metals, Metals Park, Ohio (1976), 1283-1289. 4. Tobler, R. L. and Reed, R. P., "Fatigue crack growth rates of structural alloys at 4 K," in Advances in Cryogenic Engineering , Vol. 22, K. D. Timmerhaus, R. P. Reed, and A. F. Clark, eds., Plenum, New York (1977), 35-46. 5. Tobler, R. L. and Reed, R. P., "Fatigue crack growth resistance of structural alloys at cryogenic temperatures," in: Advances in Cryogenic Engineering , Vol. 24, K. D. Timmerhaus, R. P. Reed, and A. F. Clark, eds., Plenum, New York (1978), 82-90. 6. Tobler, R. L. and Reed, R. P., "Interstitial carbon and nitrogen effects on the tensile and fracture parameters of AISI 304 stain- less steels," in: Materials Studies for Magnetic Fusion Energy Applications at Low Temperatures-III , NBSIR 80-1627, R. P. Reed, ed., Fracture and Deformation Division, National Bureau of Stand- ards, Boulder, Colorado (June 1980), 17-48. 115 7. Standard test method for plane-strain fracture toughness of me- tallic materials (designation ANSI/ASTM E399-78a)," in 1979 An - nual Book of ASTM Standards , Part 10, American Society for Testing and Materials, Philadelphia (1979), 540-561. 8. Fowlkes, C. W. and Tobler, R. L., "Fracture testing and results for a Ti-6A1-4V alloy at liquid helium temperature," Eng. Fract. Mech. 8 (1976), 487-500. 9. Roberts, E. J., Jr., "Elastic crack-edge displacements for the compact tension specimen," Mater. Res. Stands., 9_, No. 2 (1969), 27. 10. Brose, W. R. and Dowling, N. E., "Size effects on the fatigue crack growth rate of type 304 stainless steel," in: Elastic-Plastic Fracture , ASTM STP 668, J. D. Landes, J. A. Begley, and G. A. Clarke, eds., American Society for Testing and Materials, Phila- delphia (1979), 720-735. 116 •a o c E o 1TJ cu N CO ••- CJIOO - S_ c cu -r- • > 10 < s- OJ (1) C 3 U") 00 00 CM co r^ en 00 00 ■— r-~ co co en cm cm en co en r— 00 en O r- ,— CM CM CM CO coeoeni — r-^ *j- o 1^ lt> cMLncocoiocnroLOco 000000000 000000000 o r^ o ■— en o CM r r 1— r r M : z OOO LO ID o en CM CM 1— o " " o o O r^ LT> lt> un LO o " " o o en r^ ^o o= = o= = o= = 000 1— <£> rt CM CM CM 0= = 0= = Or = OOO 1— cm en 117 r— , — id iopho co cm o co co cr> incor- cm o i— ou-r— en cm id rN in . — id 00 OMn 01 n* o in en o lO 2 CO o >1 O +J I— to to to CU 03 +J lO c O) +-> GO 3 03 s- . o CM <4- 1 to E S- • ai 03 +-> a. a> 2: E 03 C S- •r- fO Q. to •r- CM i^ < cr -a LlJ c 1 1 03 c cu o ^~ •^ u +J >> 03 u 3 "**s. cr F cu E cu c 3 •^ en •r— to +J •r- 03 «4- z -a s *> — ra 03 >— -a s- 0) cu 5- 2 ai o -C D_ 5 cu 03 C CM o ->» cn^E cu a: ra a. < — r-» i^ in in in m «3- o o «=j- *fr m lO O «* o to r^ O CO LD ID o o in to o in n3 s- CU Q. 03 03 +J 03 CU O O i— .— O O O O i— o o i— i — I i — i — i— r— Ol i — CT> r— CT> i — r— i — i iii i i i i o ooo oooo X X X in r— t*^ CM X XXX i— «^- in X X X Cn i— X X ^t- 1— X X «d- i— X XXX CTi ^" i — X X X X in r~- r — i — CO CO «vf "- CO co •=*■ — CO i— CO i— CO i— - CO CO i— CO "si" r— t— CfiO<* CTi en r-» co to in CM lO r-*. to r-«. to to in cm r^ to «d- o r-*. i^. CO «* CO CO CO CO CO CO ^l- CO CO CO CO CO CO «3" CO CO «=j- CO CO CO in to *d- cn r*. CM in to *3- cn r-» CM in «* cx> CM to "si- r's. to «* m to «sf cn r-s. CM m in id <* cr> en r*. CM CM CM CO to co cn in i i o r— CO I— I go I— I < CT> CTi O .a cu T3 > ■o !c 00 03 Q. E O s- s- cu > c o o 119 LIST OF FIGURES Figure 1. Compact specimen used for fatigue tests. Figure 2. Crack initiation results for specific conditions at 76 K. Figure 3. General comparison of da/dN vs. AK results. Figure 4. Fatigue crack propagation behavior as a function of inter- stitial C+N contents with temperature as a parameter. Figure 5. SEM views of the partially intergranular failure modes at 295 K. Figure 6. SEM views of the fatigue failure modes at 76 K, showing the transgranular (heat 1) to partially intergranular transition (heat 9). Figure 7. Relative amounts of a' (bcc) martensite formed during fatigue tests at 4 K. (Arbitrary units) Figure 8. SEM views of the fatigue failure modes at 4 K, showing the transgranular to partially intergranular transition. Figure 9. Fatigue crack propagation behavior as a function of tempera- ture, with interstitial C+N contents as a parameter. 120 63.5 28.58 Width,W=50.8 12.70 Dimensions in mm 13.08 Dia. COMPACT SPECIMEN Thickness = 24.5 mm Figure 1. Compact specimen used for fatigue tests 121 1 1 1 i • CO i _ c o ii n X o n or • o 00 u. / • 1 / u. 1 •-* ** 1 i / _ 1 o o o o o o ro O O CJ O O O ro « o Q. O CJ O O u •r— M- •r— o oj CL OO S- O 4- co +J CO a> • s_ ^ c vo o r^ +j -p •r— 4-> 00 •r- C c o +-> ^. -i- o -a re c s- o CJ o o JQ O a 3 ° CD S- ZJ CD >)0DJQ 01016*1 D a * D !i! u l °4 S9|0A00|!>| 122 Fe-l8Cr-IONi STAINLESS STEELS 10 -2 0) o o E E o ■o Ld h- < or o or o o < o UJ CD Li_ 10" I0" 4 10 -5 1 1 1 — | — r 295 (This Study) 77K M il "1 I I | I I I ± J ! I I I 20 40 60 80 20 40 60 80 20 40 60 80 STRESS INTENSITY FACTOR RANGE, AK, MPa • m l/2 Figure 3. General comparison of da/dN vs. AK results 123 Fe-l8Cr-IONi STAINLESS STEELS 10 I i i | i ~i — i — | — i i i i i r T= 4K I I i I _ T = 295K E £ D T3 h- < or o O o < cr o UJ Z) o 10' Hea 10 io 5 I I I I I I 1 I T = 76 K Heats 3 Heatl 1 I I I Heats 4-9 V / - Heat 1 _ 3I0 i I i i i i 20 40 60 80 20 40 60 80 20 40 60 80 STRESS INTENSITY FACTOR RANGE, AK, MPa • m l/2 Figure 4. Fatigue crack propagation behavior as a function of interstitial C+N contents with temperature as a parameter. 124 en cu ro o CO I — 1 ■4-J (1) Mqure b SEf'i views of the partially intergranular failure nodes at 295 K. OJ It! I a XI ** ass WSM I ■ ■ • '• • - tMimC^ Figure 6. SEM views of the fatigue failure modes at 76 K, showing the transgranular (heat 1) to partially intergranular transition (heat 9), 126 o 00 t— I < 7 CD c CD o o GO c o E < > CD 5- BCC Martensite Formation of Fatigue Specimens T=4K, AK^40MPa-m 0.10 0,20 Carbon H-Nitrogen (wt.%) 0.30 Figure 7. Relative amounts of a' (bcc) martensite formed during fatigue tests at 4 K. (Arbitrary units.) 127 CD ■«= 1 cu Figure SEM views of the fatigue failure modes at 4 K, showing the transgranular to partially inter- granular transition. 128 Fe-l8Cr-IONi STAINLESS STEELS O O E E D "D uJ a: x O o *: o < tr o UJ Z> O £ 10 -2 TT Heat 1 C + N= 0.067 wt.% I0 -3 h I0" 4 - I0 -5 1 1 — i — r~r-r Heat 3 C+N = O.I28wt. % j_ j I i i i 1 1 — i — | — i i i i. Heat 9 C + N =0.325 wt.% _L X I I I 20 40 60 80 20 40 60 80 20 40 60 80 STRESS INTENSITY FACTOR RANGE, AK, MPa • m l/2 Figure 9. Fatigue crack propagation behavior as a function of temperature, with interstitial C+N contents as a parameter, 129 ANOMALOUS YIELDING, TRANSFORMATION, AND RECOVERY EFFECTS II AISI 304L STAINLESS STEELS R. L. Tobler, D. H. Beekman, and R. P. Reed Fracture and Deformation Division National Bureau of Standards Boulder, Colorado ABSTRACT Experiments describing the temperature dependence of yielding and anomalous stress-strain behaviors of metastable AISI 304L stainless steels were performed. In these steels transformation to a' bcc mar- tensite can be induced by thermal cycling or creep at stresses sig- nificantly lower than those required to form a' by simple monotonic tensile loading. A recovery phenomenon is observed for tensile speci- mens unloaded and warmed to 295 K after plastic strain at 76 K. 131 INTRODUCTION Austenitic stainless steels have a metastable face-centered cubic structure (y). But at room temperature and below both body-centered cubic (a 1 ) and hexagonal close-packed (e) structures form during cooling, under applied elastic stress, or during plastic deformation. In the commercial grades of Fe-18Cr-8Ni the martensitic phase transformations may occur during service, depending on alloy content, service temper- ature, stress history, and metallurgical condition. Various physical and mechanical property changes occur in conjunc- tion with martensitic transformation in austenitic stainless steels. The most significant consequences of transformation are volumetric expansion (about 2%), magnetic transition from paramagnetic (y) to ferromagnetic (a 1 ), and mechanical property changes. In the Fe-Cr-Ni austenitic steels, the martensitic transformation is believed to be responsible for anomalous temperature dependence of the flow strength. Suzuki et al . [1] have suggested that the transition to e is responsible for flow strength decreases at lower temperatures, whereas 01 sen and Azrin [2] have produced data that suggest a strong correlation of the decline of flow strength at low temperatures with a 1 martensite. There are other ramifications of the role of martensitic transformation in plastic deformation of austenite. This paper suggests that thermal cycling and creep behavior are significantly influenced. In particular, preliminary data are presented that suggest that austen- itic alloys do creep at low temperatures and that thermal cycling (76- 295 K) serves to alter the creep rate. 133 EXPERIMENTAL PROCEDURE The materials studied are AISI 304L stainless steels. One of these steels had been previously tested [3]. The second steel has the fol- lowing composition in weight percent: Fe-18.3Cr-l 0.1 Ni-1 .57Mn-0.021 P- 0.019S-0.50Si-0.20Mo-0.21Cu-0.028C-0.039N. This steel was obtained as a 25-mm-thick plate, having an average grain diameter of 85 urn and a Rockwell hardness of Rg74. All tests were performed on dog-boned specimens 2.54 x 4.77 mm in cross section, 50.8 mm in gage length, and transversely oriented with respect to the plate rolling direction. Conventional tensile tests used a universal test machine at a crosshead rate of 8.47 ym/s. Creep and thermal cycling tests were conducted using a servo-hydraulic test machine in load control. Clip gage extensometers and inked' scribe lines were used to measure strain. Test temperatures were achieved using liquid helium (4 K) , liquid nitrogen (76 K) , or a vapor cooling system [3]. In thermal cycling and stress tests, the specimens were loaded to a specified stress at 295 K, cooled to 76 K in a Dewar of nitrogen, and then warmed to 295 K in a jet of air after removing the Dewar. This procedure was repeated 10 times in about 0.5 h. The amount of body-centered cubic (a 1 ) martensite formed after various tests was measured using a spring-loaded magnetometer. Five magnetometer readings spaced across the gage length were taken on op- posite sides of the specimen, and averaged. The accuracy of the a' martensite determination is estimated at ± 0.02%. The magnetic readings had been previously associated with volume percent of the a', a ferro- magnetic phase [2]. 134 RESULTS AND DISCUSSION Tensile yield strength Yield strength data for AISI 304L stainless steels are plotted as a function of temperature in Fig. 1. A regular trend is expected for stable austenitic stainless steels, whereas the data for AISI 304L demonstrate an irregular temperature dependence: the yield strength increases between 295 and 200 K, then decreases between 200 and 100 K, and finally increases again between 100 and 4 K. The present results follow the trend of Read, Reed, and Schramm [3], who tested a slightly stronger heat (higher carbon content). Owing to the anomalous behavior demonstrated in Fig. 1, the 76-K yield strength of AISI 304L is only 40 MPa higher than at 295 K. For the thermal cycling and creep experiments described below, this means that for a given applied stress, a , there is not a great difference between the applied stress to yield strength ratio (a /a ) at 295 and 76 K. Behavior during the initial stage of plastic deformation is il- lustrated by the stress-strain curves shown in Fig. 2 for the steel of reference 3. During the first few percent of plastic strain, the AISI 304L curves at different test temperatures cross over one another. This cross-over behavior is not evident in stress-strain curves for stable stainless steels, such as AISI 310. Stable stainless steels remain completely austenitic in structure, and the elastic moduli and flow strengths at a given strain progressively increase as test tem- perature is lowered. In contrast, the metastable AISI 304L steel stress-strain curves exhibit early departures from linear elasticity at 100 K (and also at 76 K) , so that the proportional limit is lower at 135 these temperatures than at 199 K. In a study of Fe-18Cr-10Ni stainless steel, an observed decrease of the proportional limit between 243 and 76 K was attributed to e martensite transformation [1]. In these steels, e forms first, followed by a'. Yet strain-induced a 1 phase transformation was confirmed in the present tests of AISI 304L (see Table 1), and Olsen and Azrin [2] have shown a strong correlation be- tween the stress to initiate a' and a . Furthermore, the higher strain- hardening rates at larger strains have always been correlated with a' formation; in this range of the stress-strain curve the volume fraction of a 1 is linearly related to strain. Hence at strains above 0.011 for this alloy, there is the gradual transition to a monotonic temperature dependence of flow strengths. Clearly, there is currently conflicting evidence regarding the roles of e and a' on flow strength. The evidence suggests a dual role or possible dual product form for a': at low strains a' forms at slipband intersections and acts as a "window" for cross-slip, and at higher strains the a' forms within slipbands and acts as a dislocation or deformation barrier. Creep Application of a constant stress at constant temperature (76 K) resulted in isothermal martensitic transformation and elongation of AISI 304L specimens. It has been reported that creep occurs in AISI 300 series steels at temperatures as low as 20 K [5]. The purpose of our creep tests was to obtain observations at 295 and 76 K to aid inter- pretations of the thermal cycling tests described later. Figure 3 shows a series of creep tests performed at three tempera- tures. No creep occurred at 4 K at the stress levels indicated, but 136 significant creep occurred at 76 and 295 K. The creep at 76 K was ac- companied by a' martensite formation, a indicated in Table 1. Thermal cycling at constant stress One AISI 304L specimen was thermally cycled 10 times between 295 and 76 K with no externally applied stress, but no measurable amount of a' formed. Two additional specimens were "hemally cycled 10 times while constant stresses or 100 or 200 MPa were applied A stress of 100 MPa falls in the elastic region, whereas a stress of 200 MPa causes a small plastic strain of 0.0005. Measurable a' martensite formed in both specimens and there was also a significant elongation of the speci- mens. As shown in Table 1, greater elongation and martensitic trans- formation occurred in the specimen at 200 MPa, but the ratio of perxent a' martensite ic elongation wao a factor of three higher for the specimen held at 100 MPa. The thermal cycling test at 200 MPa took 0.5 h to complete. In comparison, more a 1 phase transformation and specimen extension resulted from thermal cycling (295 to 76 K) at constant stress than from the 0.5-h, 76 K creep test at the same stress level. However, the ratio of a' to specimen strain was reduced when thermal cycling was combined with creep at 200 MPa (Table 1). Recovery at room temperature Figure 4 illustrates a novelty discovered in our tests. The usual strain-hardening behavior for an austenitic stainless steel is depicted by data from specimen 1, which was monotonically loaded in tension at 76 K. Specimen 2 was (1) cooled to 76 K, (2) strained into the plastic range, (3) completely unloaded, (4) warmed to 295 K, (5) cooled to 76 K, 137 and (6) reloaded at 76 K. The stress-strain curve on reloading at 76 K diverges from the previous unloading line, and plastic flow resumes at a lower stress than expected. In essence, the strain-hardening of the previous loading cycle is partially erased by the intermediate thermal cycle (76 to 295 to 76 K) at zero stress. The proportional limit, observed on each reloading, remains nearly equivalent with each thermal cycle and this can be repeated for a number of cycles. The effect is apparently linked with microstructural instability. When the same experiments were repeated using metastable AISI 316 and stable AISI 310 steels, the anomalous behavior was observed for meta- stable AISI 316, but not for stable AISI 310. Therefore, to observe the effect it is necessary (1) to have a metastable steel and (2) to include the thermal cycle to 295 K and back to low temperature. After a total elongation of 5%, specimen 1 (monotonic loading) contained 5% a' martensite, whereas specimen 2 (repeated stress and thermal cycling) contained 6% a 1 martensite. This tends to confirm the suspicion that a 1 martensite forms during each loading cycle at about the same stress level and that warming to room temperature provides sufficient internal recovery to permit additional a' transformation. However, this explanation does complicate the earlier proposed concept that early a' transformation occurs at cross-slip sites, where suffi- cient dislocation pile-up stresses are presumably generated to induce transformation. Under such circumstances, recovery would not be ex- pected to encourage transformation. SUMMARY The results of experiments on metastable AISI 304L stainless steels are : 138 1. Significant creep occurs at 295 and 76 K, but not at 4 K, at the stress ranges studied. 2. Thermal cycling between 295 and 76 K serves to increase the creep rate. 3. Warming to room temperature produces limited recovery after tensile deformation at 76 K. 4. The role of a 1 and e martensite transformations in the de- formation processes remains unclear. ACKNOWLEDGMENT The assistance of D. Baxter of NBS in performing tensile tests is gratefully acknowledged. REFERENCES 1. Suzuki, T., Kojima, H., Suzuki, K., Hashimoto, T., Koike, S., and Ichihara, M., Plastic deformation and martensitic transformation in an iron-base alloy, Scr. Metall. lp_ (1976), pp. 3 r o-358. 2. Olsen, G. B. and Azrin, M., Transformation behavior of TRIP steels, Metall. Trans. A , 9A (1978) 713-721. 3. Read, D. T., Reed, R. P., and Schramm, R. E. (1979), Low tempera- ture deformation of Fe-18Cr-8Ni steel, Metall. Trans, (to be pub- lished) . 4. Reed, R. P. and Mikesell, R. P., The stability of austenitic stainless steels at low temperatures as determined by magnetic measurements, in Advances in Cryogenic Engineering , Vol. 4, K. D. Timmerhaus, ed. (Plenum, New York, 1960), pp. 84-100. 139 5. Voyer, R. and Weil, L., Creep properties of 18-10 and 25-20 stain- less steels at cryogenic temperatures, in: Advances in Cryogenic Engineering , Vol. 10, K. D. Timmerhaus , ed. (Plenum, New York, 1965), pp. 110-116. 140 03 +-> o c: o cn o 03 +-> ' -M 03 < O cn* 1— C O UJ CD CO c . CD S- 03 CD Z3 00 CD S_ -t-> 00 CD <4- O S- rt3 OO s: CO CD 03 . — - S- b^ CU » — Q. CD O CD Q. CO -a CD E s- o it- S- CD D_ +-> 00 CD CO LO r*. CM CM o LO o CM o «=J- I— CO r— o o o CM o o CD O o o o o cn o o o o «=c <=c CT> CM CM CM CM CM CO CD CD cn C c c: •r" *r— •r— -a -a T3 rO 03 03 o o O 1 — 1 — r ~ a o a •r— •r- ■r~ C cz C o o o +-> +-> +J o o o c c c: o o o >> >> c 03 +-> 00 o u 00 CD O 03 +-> 00 e o cj 00 CD CJ a c c o o o_ Q. CD CD CD CD S- s_ C_> C_> 03 03 -t-> 03 00 o o 00 CD O 03 s_ S- S- CD CD CD -C -C -C LO cn CO >> .a CD ■a "r— > ■a 00 o 03 O S- Cf- S- CD > o o 141 LIST OF FIGURES Figure 1. Yield strength of AISI 304L stainless steels, showing ir- regular temperature dependence. Figure 2. Stress-strain curves for AISI 304L steel, showing cross-over behavior. Figure 3. Creep test results for AISI 304L at three temperatures. Figure 4. Stress-strain curves for AISI 304L specimens illustrating recovery effect after unloading and warming to 295 K. 142 AISI 304 L STAINLESS STEELS 500 400 o Q_ 300 x h- UJ (T h- (/) Q _l UJ > 200- 100- 1 1 1 1 1 X. Fe-l8.5Cr-8.75Ni-0.03C -0.061 N [3] ^^o Fe-l8.3Cr-IO.INi-0.028C-0.039N [this study] i 1 i 1 I 100 200 TEMPERATURE, K 300 Figure 1. Yield strength of AISI 304L stainless steels, showing ir- regular temperature dependence. 143 700 AISI 304L Fe -18.5 Cr -8.75 Ni-0.03C-0.061N 0.6 0.8 STRAIN, €, % 1.0 1.4 Figure 2. Stress-strain curves for AISI 304L steel, showing cross-over behavior. 144 0.18 0.16 CREEP TEST Fe -l8.3Cr-IO.INi -0.028C-0.039N 8 76 K C7=200 MPa, 81% of cr y elongation at start of creep - 0.095% 4 K, no measurable creep from act as windows for dislocation pileup release; the result is an "easy glide" region of the stress-strain curve. This "easy glide" region had previously been reported by Reed and Guntner [5], who argued that it was caused by strain-induced e martensite, since their x-ray and transmission electron microscopy data for an Fe-18Cr- 8Ni alloy indicated an increasing concentration of hep during easy glide. However, Olson and Azrin [6], studying an Fe-9Cr-8Ni alloy, found an excellent correlation between flow strength reduction at lower temperatures and the initiation of a' martensite. No "easy glide" region of the stress-strain curve was observed. The effect of grain size on low-temperature flow strength of aus- tenitic stainless steels has not been documented. The characterization of grain size effects is significant in the design of strong and tough steels for low-temperature applications. Soviet authors [7,8] report the existence of three major anomalies in the temperature dependence of the yield strength for Fe-18Cr-xNi austenitic stainless steels containing 8, 10, or 20% Ni . These anomalies amount to sharp decreases of the yield strength on the order of 100 MPa . The first anomaly occurs between 20 and 4 K and is unexplained. The second anomaly occurs at temperatures between 60 and 35 K and is attrib- uted to the Neel transition (magnetic ordering). The third anomaly occurs in metastable alloys (lower Ni) at temperatures just above the NL temperature and is attributed to strain-induced martensitic phase trans- formations. Clearly, continuous flow-strength measurements as a func- tion of temperature should be made to assess these results. Therefore, tensile tests were conducted on a grade of AISI 316 austenitic stainless steel to assess the influence of martensitic transformation, grain size, magnetic transitions, and \/ery low temperature c n flow strength. 150 MATERIALS The AISI 316 stainless steel tested in this study was purchased in the form of a 38-mm-thick plate. Tensile specimens were machined in the longitudinal orientation in the form of flat coupons having a 50.8-mm gage length and a 2.54- x 4.77-mm-gage cross section. After the specimens were annealed at 1338 K for 0.5 h, they had a Rockwell B (R g ) hardness of 71 and an average grain diameter of 55 ym (ASTM grain size number 5.5). The chemical compositions of this steel (denoted heat 2) and other AISI 316 steels referred in this paper [9-11] are listed in Table 1. PROCEDURE Tensile tests were performed on (1) annealed and (2) strained and recrystall ized material. The strained and recrystall ized material was used to determine grain size effects. All tensile tests were performed at a crosshead rate of 8.5 ym/s using a screw-driven machine and cryostat [1,12]. Test temperatures of 4 and 76 K were achieved with baths of liquid helium and liquid nitrogen. Other temperatures were achieved by controlled thermal conduction using liquid nitrogen reservoirs and heating wires attached to the specimen grips and load train. Current to the heating wires was varied until thermal equilibrium was reached at the specified test temperature, which was accurate to ± 1.5 K. Load- versus-deflection curves were recorded with a commercial load cell and clip-gage extensometer. Clip-gage extensometers magnify the specimen strain and permit measurement of only the specimen strain within the (2,5-cm) gage length. Estimated flow strength measurement inaccuracies are ± 2%, but low- temperature specimen-to-specimen variations (including both material and 151 measurement variations) may reach ± 7.5%. After plastic strains of 0.2 to 3.65% the specimens were unloaded, warmed to 295 K, and examined for the presence of ferromagnetic a' martensite with a torsion-bar mag- netometer [2]. X-ray diffractometer measurements were used to try to detect small concentrations of e martensite. No e martensite was de- tected and our estimated sensitivity is 2%. The stress-strain curves were calculated in the conventional manner, and the measurement un- certainty of flow strength values is not greater than 3%. As an indi- cation of data scatter, seven tests were performed on the AISI 316 steel at 76 K, resulting in an average yield strength (a ) of 448 MPa and an overall spread of ± 7 MPa. Grain size variations were obtained by recrystall ization treat- ments. Tensile specimens were plastically strained 20 to 30% at 295 or 76 K and then annealed for 0.25 to 0.5 h in an argon atmosphere at 1233 to 1573 K. These treatments produced grain sizes ranging from 22 to 140 ym. Subsequent to recrystall ization treatments, the specimen sur- faces were cleaned by electropol ishing. The recrystall ization treatments increased the scatter of yield strength measurements to about ± 15 MPa. TENSILE YIELD STRENGTH RESULTS The tensile yield strength (a ) and initial stress-strain behavior of the Fe-17Cr-12Ni-2Mo alloy were measured as a function of tempera- ture, grain size, martensitic transformations, and magnetic transitions. These effects are described in this section. Temperature : As listed in Table 2, twenty-six a measurements were obtained for annealed AISI 316 stainless steel (heat 2) at seventeen temperatures between 323 and 4 K. Magnetometer inspections of the 152 specimens indicated no a' martensitic transformation for this heat after 0.2% plastic strain, regardless of test temperature. The a data between 295 and 4 K are plotted and compared with previously published data [9,10, 13-16] in Fig. 1. This comparison reveals significant heat-to-heat variations of o , with considerably greater differences at cryogenic temperatures than at room temperature. Several references cited in Fig. 1 report data at only a few test temperatures, whereas the functional dependence of o on temperature can be identified only by extensive measurements performed over a broad temperature range. Therefore, data for three heats of AISI 316 are replotted in Fig. 2. These data show that o increases progressively as temperature is lowered. Despite minor variations of grain size, strain rate, and composition, the a values for the three heats are nearly equivalent between 973 and 175 K, where the test temperature intervals from different studies overlap. Below 175 K, significant variations of strength begin to emerge. At still lower temperatures, the variations become more obvious until, at 4 K, a correlates with composition: Fe- 17.25Cr-13.48Ni (heat 1) exhibits the highest strength, Fe-17.34Cr- 12.17Ni (heat 2) exhibits intermediate strength, and Fe-1 6.720-11 .27Ni (heat 3) exhibits the lowest strength. This apparent association of higher strength with higher nickel content may relate to crystal structure stability as discussed below. Thermally activated plastic flow theory predicts that a should exhibit some dependence upon strain rate. A few experiments were per- formed on the annealed AISI 316 stainless steel to evaluate this effect. Increasing the crosshead rate by a factor of 100 from 0.85 ym/s to 85 ym/s increased the flow strength at strains lower than 1% by 153 factors of 6% at 295 K and 3% at 76 K. The effect of crosshead rate on the a at 4 K was negligible. Therefore, strain rate effects do not play a significant part in explaining the variability noted in Fig. 1. Strength variations may arise between specimens from a single heat of steel owing to inhomogeneity or thermomechanical processing effects. As shown in Fig. 1, Read and Reed's data [13] for a 316 stainless steel at 295, 76, and 4 K are 5 to 17% higher compared with the results of Tobler, Reed, and Burkhalter [9], who tested the same heat. In the former study, the specimens were tested in the as-machined condition, whereas the specimens used in the second study were annealed (1289 K, 1 h) after machining. This annealing step reduced the hardness at 295 K from R R 79 to R R 75 and lowered a by increasing the grain size slightly and removing any residual cold work from the fabrication or machining processes. An example of within-heat variability is also noted in the present study: the annealed AISI 316 steel specimens had a 76-K a of 448 ± 7 MPa, whereas reprocessed (strained and recrystallized) specimens having a similar grain size (55 ym) had a 76-K a of about 480 ± 15 MPa. Thus, process history influences strength. Larger scatter for the recrystallized specimens is to be expected since dislocation density and substructure within the grains depend on prior history and recrystalliza- tion temperature. Grain Size : Figure 3 illustrates the effect of grain size on a at 295, 76, and 4 K. The data at each temperature fit an inverse square root dependence on the average grain diameter, d, which is in reasonable accord with the familiar Hall-Petch equation: 154 a y = a. + kd" 1/2 (1) wh ere a. refers to the flow strength in the absence of grain boundaries -1/2 (the "friction stress") and k is the slope of the a -versus-d ' plot and indicates the grain boundary dependence of the yield strength. The values of a. and k are given in the figure. Martensitic transformation : Specimens from this study were screened using a magnetometer, but no a' martensite was detected after 0.2% plastic strain. To force the phase transformation to occur, a series of specimens was loaded to higher strains. Figure 4 presents the strain dependence of a' at 76 K. Notice that a plastic strain of about 2% is necessary for a 1 to form at 76 K. The dependence of a 1 on plastic deformation is linear, and this trend continues at least to a strain of 5%. At larger strains, the dependence apparently increases since fractured specimens show between 44 and 52% a' at uniform strains of 40 to 50%. The stress-strain curves at higher strain increments are presented in Figs. 5 and 6. The quantity of a 1 martensite formed after either 3.65 or 2.75% plastic strain is negligible for test temperatures between 300 and 190 K, but below 199 K, a' is formed in increasing amounts. In fact, as illustrated in Fig. 7, the percent a' martensite divided by the strain is linearly dependent on test temperature. The curves of Fig. 5 and 6 represent specimens in various states of transformation, containing from to 0.1% a' martensite per unit strain. But the stress-strain curves are regular and free from discontinuities, indicating that the martensitic transformation occurs gradually with no anomalous effects. These findings contrast with the behavior of Fe-18Cr-8Ni and Fe-18Cr- lONi stainless steels that exhibit: (1) irregular yield and flow strength 155 temperature dependences owing to the initiation of e and a' phase transformations [7], and (2) suppression of a' martensite at 4 and 20 K, compared with that at 76 K [5]. DISCUSSION Thermally activated flow : According to classical theory (e.g., Seeger [17]), a may be separated into two components: °y = a A + a T (2) Here a» is an athermal stress component determined by metallurgical structure, and a T is a thermal stress component governed by lattice vibrations and rate processes. By estimating from Fig. 2, the value of a. (uncorrected for shear modulus) for the AISI 316 steel heats 1 and 2 is taken to be 100 MPa, which is the asymptote approached by a at elevated temperatures. From previous work involving thermodynamic and thermally activated plastic flow theories, the following expression was developed for a T : a y = (a Q - a A ) exp (- CT) (3) Here a n is the extrapolated value of the yield strength at absolute zero and C is a materials constant comprised of activation energy, entropy, and strain-rate terms. This expression derived from the work of Yaroshevich and Ryvkina, who did not correct the a. value for the effect of shear modulus [18]. The applicability of Eq. (3) for AISI 316 stainless steels was tested by plotting log (a - a^) versus T. Data agreeing with Eq. (3) will produce a single straight line when graphed this way, assuming that 156 a single mechanism of plastic flow is operative. As demonstrated in Fig. 8, Eq. (3) successfully predicts the experimental results for the three AISI 316 steels down to 175 K, but divergences from the main trend occur for each of the three heats at 175, 150, or 30 K, depending on alloy composition. A graphical solution leads to the following ex- pression for the AISI 316 data (heats 1 and 2) between 975 K and the temperature of divergence: a = 100 + 700 exp (-6.2 x 10' 3 T) (4) As a consequence of divergences from the main trend, the experimentally measured 4-K a values for all heats are lower than predicted by Eq. (4) If the trend lines of Fig. 8 are extrapolated, then the measured o values for heats 1 and 2 at 4 K (near absolute zero) are 15 to 45% lower than the extrapolated value of a Q , which is 780 MPa. The data for heat 3 suggest an even greater difference. The reasons for deviations from Eq. (4) are not completely under- stood. It was suggested previously that dislocation tunneling and magnetic ordering may affect c at temperatures approaching absolute zero [9]. Likewise, adiabatic heating becomes significant as absolute zero is approached. But these processes alone are unable to account for the deviant behavior of AISI 316 heats 2 and 3 at temperatures in the 150- to 175-K range. Moreover, these processes do not account for the apparent sensitivity to alloy composition. As noted in Fig. 1, diver- gences from Eq. (4) for heats 1, 2, and 3 occur at 30, 150, and 175 K, respectively. This sequence follows the order of decreasing alloy content (Table 1), decreasing stacking fault energy (Table 3), and 157 decreasing crystal structure stability, as indicated by the NL and M D temperatures (Table 3). The possible martensite roles are discussed in the next section. It is also possible that more than one dislocation mechanism is operative, depending on temperature and composition. If a second thermally activated dislocation process becomes rate controlling at sufficiently low temperatures, the change of slopes in Fig. 5 would result. In Figure 9, the a data for the three AISI 316 stainless steels have been normalized by using the shear modulus and plotted versus test temperature, which is normalized by using the melting point, taken as L, = 1658 K. Values of the shear modulus at cryogenic temperature were taken from Ledbetter [11]. These data indicate that the thermal component of a becomes active at a temperature of 1/3 the melting point. A linear extrapolation of the thermal portion of the yield strength gives a value of about 300 K for the intersection of the two components. The divergences from the thermal components at low temperatures, shown in Figure 8, is also observed here. Role of martensitic transformation : The nonlinear dependence of In Oj (Fig. 8) or a on (Fig. 9) on temperature at yery low temperatures may result from strain-induced martensitic transformation. The NL and M n temperatures in Table 3 provide only a relative indication of stability, having been calculated on the basis of the Cr, Ni, Mn, Si, C, and Mo contents. The calculated M s values are negative, indicating complete stability on cooling to absolute zero, which is consistent with the observation that no a' martensite formed on cooling 158 to 4 K. Beyond this, a quantitative comparison of the measured and calculated transformation temperatures is impractical because the cal- culated NL and M D temperatures are influenced by trace elements which are neither controlled nor analyzed during steel manufacture. It is unlikely that the a' transformation can account for the observed deviations since no a 1 martensite was detected at 0.2% plastic strain. Furthermore, from Fig. 4 it is apparent that no a' is expected to form until a plastic strain of about 2% is reached. However, it is possible that martensitic transformations producing small amounts of e phase may be responsible. From Table 3, the calculated stacking fault energies decrease, but remain positive, with increased deviation of a from linearity. This possibility could not be confirmed, since our metallographic and x-ray inspection methods lacked the required sensi- tivity to detect a small amount (< 2 %) of e . In the commercial Fe-18Cr-8Ni austenitic stainless steel grade (AISI 304), the a 1 martensitic product and, perhaps, the e martensitic product apparently contribute to the anomalous temperature dependence of a and to the "easy glide" portion of the stress-strain curves at tempera- tures near 77 K [4,5,23]. However, the data represented in Figs. 4-7 suggest that in the alloy of this study (Fe-17Cr-12Ni-2Mo) the tempera- ture dependence of the flow strength (including a ) is well behaved and the a 1 martensitic product, therefore, does not contribute to anomalous behavior. The nature of the linear relationship between the "normalized" percent bcc martensite (percent magnetic martensite divided by engi- neering plastic strain) and temperature deserves additional experimental 159 clarification. It is implied that the temperature of about 190 K repre- sents the onset of strain-induced martensite, and at higher strains and temperatures little magnetic product phase is detected. This behavior has been substantiated by the tests to fracture at 195 and 295 K. How- ever, the nature of the probable family of curves, generated by taking data at larger strains, is not clear. Similar data for AISI 304 steel indicate a gradual deviation from linearity at larger strains, but the existence of a constant temperature of a' initiation [24]. With respect to the role of martensitic products in affecting the temperature dependence of the flow strength of Fe-Cr-Ni alloys at low temperatures, there are two schools of thought: Suzuki et al . [4] argue that the e phase must be responsible for the reduction of a at low temperatures since in their study a 1 martensite was detected at strengths considerably greater than a . Additionally, they have argued from optical metallography evidence that e martensite forms at strength levels near the measured a . Conversely, Olson and Azrin [6] studied a less stable Fe-9Cr-8Ni alloy and obtained excellent correlation between the temperature dependence of o and the flow strength at which a 1 mar- tensite was first detected. From the results of this study, it is apparent that the onset of a' strain-induced martensite in austenitic Fe-Cr-Ni alloys does not lead to a reduction of the flow strength at low temperatures. This finding, again, raises the question of the nature of the a 1 , ferromagnetic martensitic product or products that form(s) during plastic deformation. Conflicting parametric studies and lack of conclusive microscopy preclude assessment of the role of e and a' mar- tensites in affecting low temperature flow strength. 160 Role of grain size : Data on the effect of grain size for aus- tenitic stainless steels are rare. The present data at room temperature indicate a very slight effect, as evidenced by the low-k value of -1/2 -1/2 191 MPa«ym ' . This value is in agreement with the 221 MPa»um ' value obtained by Norstrom [25] for a 31 6L stainless steel containing no nitrogen. Low-k values are typical for face-centered cubic alloys at room temperature. At cryogenic temperatures, however, the grain size effect increases significantly. As shown in Fig. 3, the values of a. and k at 4 K are greater by factors of 2 and 5, respectively, than the values at 295 K. In Fig. 10, the temperature dependences of k and a. from Eq. (1) are plotted, using the 4- to 295-K data from this study and the 295- to 875-K data of Norstrom [25]. The k is strongly temperature-dependent at low temperatures, but insensitive to temperature above room temperature. The frictional stress, a., increases with decreasing temperature and the In o. is linearly dependent on T(K) for temperatures lower than about 500 K. These data clearly imply that grain size has a significant effect on low-temperature flow strength. From Fig. 10, the quantity a - o. represents the grain boundary contribution to flow strength and about 30% of a at 4 K is contributed by grain-boundary strengthening. Role of magnetic transition : Similar to other Fe-Cr-Ni austenitic stainless steels, AISI 316 undergoes a magnetic ordering phenomenon characterized by a change from the paramagnetic to the antiferromagnetic state at low temperatures. Accordingly, a peak occurs in the tempera- ture dependence of magnetic susceptibility at T N , the Neel transition temperature, which corresponds to the onset of antiferromagnetic or- dering. Since a decrease of elastic constants occurs at T N [26,27], it 161 was suggested that the Neel transition likewise affects dislocation motion and hence tensile properties [7,8]. On the basis of Soviet experimental data [7,8], a sharp decrease of a of about 100 MPa should be expected for our Fe-18Cr-12Ni alloy at temperatures between 60 and 48 K. As shown in Fig. 11, elastic property measurements for Fe-17.5Cr- 12.9Ni stainless steel do reveal a noticeable drop in Young's modulus (E), identifying the Neel transition between 40 and 20 K. From Table 3, the calculated predicted value of T N is 26 K, but the data of Fig. 11 indicate no decrease of a at temperatures corresponding to T... From Table 3, T N values for the three AISI 316 steels of Fig. 8 are nearly equivalent, despite minor compositional differences. Care must be exercised in these arguments, however. The approxi- mate reduction of E after magnetic transition is about 2%, which is within the possible a data variation. For better assessment, the large grain size contribution should be subtracted from a , and a. be examined 3 y i carefully as a function of temperature. SUMMARY Tensile tests were performed on AISI 316 stainless steel (Fe- 17.34Cr-12.17Ni ) at many temperatures between 323 and 4 K, and the results were compared with data for other AISI 316 heats at temperatures to 973 K. The findings are: 1) The existing data for AISI 316 show a monotonic temperature dependence, with a increasing continuously as temperature is lowered between 973 and 4 K. 162 2) Between 973 and 175 K, a data for AISI 316 follow a single equation derived from the theory of thermally activated plastic flow. Below 175 K, deviations occur, apparently depending on chemistry variations which in turn affect phase stability and dislocation dynamics. 3) Deformation beyond 0.2% plastic strain is needed to initiate transformation to a' martensite at temperatures below 180 K. 4) Reduction of grain size contributes significantly (30% at 4 K) to low-temperature flow strength and at low temperatures the grain size dependence of a follows a Hall-Petch relationship. 5) Magnetic transitions do not noticeably affect the deformation of AISI 316. ACKNOWLEDGMENTS We thank D. Beard, Office of Fusion Energy, DoE for his support and continued encouragement. D. Beekman and D. Burkhalter of NBS contributed significantly to this study by conducting most of the tensile tests. Dis- cussions with Dr. E. L. Brown of NBS, who contributed valuable metallographic assistance and reviews of the manuscript, are also gratefully acknowledged. REFERENCES 1. Read, D. T. , Reed, R. P., and Schramm, R. E., Low temperature deformation of Fe-18Cr-8Ni steels, in: Materials Studies for Magnetic Fusion Energy Applications at Low Temperatures-II , NBSIR 79- 1609, F. R. Fickett and R. P. Reed, eds., National Bureau of Standards, Boulder, CO (June 1979), pp. 151-172. 2. Reed, R. P. and Mikesell, R. P., The stability of austenitic stain- less steels at low temperatures as determined by magnetic measure- ments, in: Adv. Cryo. Eng. , _4, K. D. Timmerhaus, ed. Plenum Press, New York (1960), pp. 84-100. 163 6. 7. 10, 11 Larbalestier, D. C. and King, H. W., Austenitic stainless steels at cryogenic temperatures I, structural stability and magnetic proper- ties, Cryogenics , Tj3 (1973) 160-168. Suzuki, T., Kojima, H., Suzuki, K., Hashimoto, T, and Ichichara, M., An experimental study of the martensite nucleation and growth in 18/8 stainless steel, Acta Metal 1 . , 23 (1972). Reed, R. P. and Guntner, C. J., Stress-induced martensitic trans- formations in 18Cr-8Ni steel, Trans. AIME , 230 (1969) 1713-1720. Olson, G. B. and Azrin, M., Transformation behavior of TRIP steels, Metal! . Trans. A , 9A (1978) 713-721. Ilyichev, V. Ya., Medvedev, Ya . M., Shapovaliev, I. A., and Klimenko, I. N., Low temperature anomaly of the temperature dependence at the flow stresses in iron-chromium-nickel alloys, Phys. Met. Metal! ogr. , 44, No. 1 (1978) 173-176. Verkin, B. I., Ilyichev, V. Ya., and Klimenko, I. N., The low- temperature change of the magnetic structure and plastic properties of Fe-Cr-Ni alloys, in: Adv. Cryo. Eng. , 26 Plenum Press, New York (1980), pp. 120-125. Tobler, R. L., Reed, R. P., and Burkhalter, D. S., Temperature dependence of yielding in austenitic stainless steels, in: Adv . Cryo. Eng. , 26^, Plenum Press, New York (1980), pp. 107-113. Sanderson, G. P. and Llewellyn, D. T., Mechanical properties of standard austenitic stainless steels in the temperature range -196 to +800C, J. Iron Steel Inst. , 207 (1969) 1129-1140. Ledbetter, H. M., Sound velocities and elastic constants of aus- tenitic stainless steels, to be published. 164 12. Reed, R. P., A cryostat for tensile tests in the temperature range 300 to 4 K, in: Adv. Cryo. Eng. , 7_, K. D. Timmerhaus , ed . , Plenum Press, New York (1961), pp. 448-454. 13. Read, D. T. and Reed, R. P., Fracture and strength properties of selected austenitic stainless steels at cryogenic temperatures, to be published, Cryogenics (1980). 14. Baughman, R. A., Gas atmosphere effects on materials, Progress Report #2, AF 33(616), General Electric Co. (November 1958). 15. Tobler R. L., Mikesell, R. P., Durcholz, R. L., Fowlkes, C. W., and Reed, R. P., Fatigue and fracture toughness testing at cryogenic temperatures, NBSIR 74-359, National Bureau of Standards, Boulder, CO (March 1974), pp. 182-308. 16. Hoke, J. H., Mabus, P. G., and Goller, G. N., Mechanical properties of stainless steels at subzero temperatures, Met. Prog. (May 1949) 643-648. 17. Seeger, A., The generation of lattice defects by moving disloca- tions, and its application to the temperature dependence of the flow-stress of FCC crystals, Phil. Mag., 46 (1955) 1194-1217. 18. Yaroshevich, V. D. and Ryvkina, D. G., Thermal-activation nature of plastic deformation in metals, Sov. Phys.-Solid State (Engl. Transl.) 12, No. 2 (1970) 363-370. 19. Schramm, R. E. and Reed, R. P., Stacking fault energies of seven commercial austenitic stainless steels, Metal! . Trans . , 6A (1975) 1345-1351. 165 20. Larbalestier, D. C. and King, H. W., Prediction of the low tempera- ture stability of type 304 stainless steel from a room temperature deformation test, in: Proceedings of the Fourth International Cryo - genic Engineering Conference , K. Mendelssohn, ed . , IPC Science and Technology Press, Guildford, Surrey, England (1975) pp. 338-340. 21. Williams, I., Williams, R. G., and Capellaro, R. C, Stability of austenitic stainless steels between 4 K and 373 K, in: Proceedings of the Sixth International Cryogenic Engineering Conference , ed . K. Mendelssohn, IPC Science and Technology Press, Guildford, Surrey, England (1976), pp. 337-341. 22. Warnes, L. A. and King, H. W., The low temperature magnetic proper- ties of austenitic Fe-Cr-Ni alloys 2, the prediction of Neel tem- peratures and maximum susceptibilities, Cryogenics , 16 , No. 11 (1976) 659-667. 23. Guntner, C. J. and Reed, R. P., The effect of experimental vari- ables including the martensitic transformation on the low tempera- ture mechanical properties of austenitic stainless steels, Trans. Am. Soc. Met. , 55 (1962) 399-419. 24. Starr, C. D., Notes on the plastic critical temperature in strain- induced martensite reaction, Trans. AIME , 197 (1953) 654. 25. Norstrom, L. A., The influence of nitrogen and grain size on yield strength in Type 316L austenitic stainless steel, Met. Sci . , 11 (1977) 208-212. 26. Collings, E. W. and Ledbetter, H. M., Sound velocity anomalies near the spin glass transition in an austenitic stainless steel alloy, Phys. Lett. , 72A , No. 1, 11 (1979) 53-56. 166 27. Ledbetter, H. M. and Col lings, E. W., Low-temperature magnetic and elastic-constant anomalies in three manganese stainless steels, in The Metal Science of Stainless Steels , E. W. Col lings and H. W. King, eds. 167 o o T3 ■»-> 00 oo ■M O ■a aj s- s- QJ 4- CU S- 00 +-> <1J co CO o a> e>3 CD • U O C s- +-> aj A3 4- d) O) DC en co o «3- CO C\J 00 LO cr> o o cm o CO CO CO CO LO cm CD u fO en CT> o CO CM CO CO O O CO CM o LO LO CO +-> CM CO LO «3- co I O 1 CM CO CM LO i— CM CM CM CD o CM CO LO O CM O O CO o r-^ r— • CM LO LO LO CO o O O O CO cr> CO +-> I 1 I 1 O0 O r— r— r— 00 I 1 I 1 CO 168 Table 2. Yield stress measurements at 0.2% plastic strain obtained for AISI 316 (heat 2) of this study, Test Temperature, T, Yield Stress, o K MPa* : 323 185 295 204 295 198 295 1 98 240 244 215 281 194 298 192 306 190 307 164 349 154 368 143 380 125 399 100 428 76 456 76 451 76 445 76 450 76 441 76 447 76 448 60 456 39 487 23 528 4 537 4 527 *To convert from MPa to ksi , divide by 6.895 169 Table 3. Calculated stacking fault energies (SFE) and NL, NL, and T N temperatures for AISI 316 steel heats of this study. Heat No. & Reference SFE a , M S ' M C 'n ' mJ/M 2 K K K 1 [9] 70.4 -226 190 24.7 2 this study 59.6 -126 232 25.5 3 [10] 57.4 - 47 246 7.4 4 [11] NA NA NA 26.4 a. Calculated from equation (9) of Schramm and Reed [19]. b. NL is the temperature at which martensite starts forming spontaneously on cooling, as calculated from the equation of Larbalestier and King [20]. c. NL is the temperature above which no deformation induced martensite is possible, as calculated from the equation of Williams, Williams, and Capellaro [21]. d. T., is the temperature of the onset of anti ferromagnetic ordering, as calculated from the equation of Warnes and King [22]. 170 LIST OF FIGURES Fig. 1. Yield strength of AISI 316 stainless steels at low tempera- tures, illustrating heat-to-heat variability. Heat 1 is from Tobler, Reed, and Burkhalter [9]; heat 3 is from Sanderson and Llewellyn [10]. Fig. 2. Yield strength data for AISI 316 stainless steels in the elevated-to-cryogenic temperature range, illustrating overall temperature dependence. Fig. 3. Effect of grain size on the yield strength at three test temperatures for the Fe-17Cr-12Ni-2Mo alloy. Fig. 4. Dependence of bcc (a 1 ) martensite on tensile plastic strain. Martensite was detected by using a bar magnet and torsion balance beam. Fig. 5. Temperature dependence of the yield and flow stresses for AISI 316 stainless steels, showing no evidence of anomalies despite martensitic phase transformation occurring at higher strain increments and lower test temperatures. Fig. 6. Stress-strain curves for AISI 316 stainless steels at various test temperatures, showing no evidence of anomalies despite martensitic phase transformations occurring at higher strain increments and low temperatures. Fig. 7. Amount of bcc martensite per unit strain in AISI 316 specimens deformed to plastic strains of 2.75 or 3.65%, indicating clear evidence of martensitic phase transformation at cryogenic temperatures. Extrapolated temperature of 0% bcc martensite is about 180 K. 171 Fig. 8. Temperature dependence of the thermal stress component, illus- trating applicability of thermally activated plastic flow theory at temperatures between 900 and 175 K. Fig. 9. Normalized yield strength versus normalized temperature plot for three AISI 316 stainless steel heats, showing variations in thermal and athermal stress components. Fig. 10. Dependence of the tensile yield strength, a ; the Hall-Petch parameter, K; and the flow strength, a., for AISI 316 steel. High temperature data from Norstrom 25 and a. is adjusted to 0.04 weight percent nitrogen. Fig. 11. Temperature dependence of yield stress and Young's modulus for AISI 316 stainless steel heats, showing no drop in yield strength at the Nee! transition temperature. 172 800 co a. i- O z LU CO C0 Q UJ Fig. 1 . 600 400 = 200- T 1 AISI 316 Steel (annealed) •Heat 1 Heat 2 _Jthis study) Heat 3 * Read and Reed [13] O Baughman [14] v Tobler, Mikesell, Durcholz, Fowlkes, and Reed [15] n Hoke, Mabus, and Goller [1 6] o Sanderson and Llewellyn [10] I i 100 200 300 TEMPERATURE, K Yield strength of AISI 316 stainless steels at low tempera- tures, illustrating heat-to-heat variability. Heat 1 is from Tobler, Reed, and Burkhalter [9]; heat 3 is from Sanderson and Llewellyn [10]. 173 CO b" z UJ DC I- co Q _l UJ >- 600 400- 200 °A AISI 316 Steels (annealed) o Heat 1 Tobler, Reed, and Burk halter [9] • Heat 2 (this study) A Heat 3 Sanderson and Llewellyn [10] ca □ a i_L _L ± ± 200 400 600 800 1000 TEMPERATURE, K Fig. 2. Yield strength data for AISI 316 stainless steels in the elevated-to-cryogenic temperature range, illustrating overall temperature dependence. 174 CM +-> CD -t-> CD CD S- +-> en CD S_ 4-J m -a CD >> CD o CD M l/> C fO S_ cn 4- O +-> O CD <+- O 03 O CNJ I I i CD CD -c ■P S- o n- c/i CD S_ 4-> s_ CD Q. E CD +-> BdW lA -0 'H10N3«1S QT3IA 175 UJ H 55 z UJ fr- < u o A 0.04 TENSILE STRAIN 0.06 Fig. 4. Dependence of bec (a 1 ) martensite on tensile plastic strain, Martensite was detected by using a bar magnet and torsion balance beam. 176 1000 CO 0. CO CO UJ ■ 0. 100 Li. 2 O „ 1- b« Z 1 LU Z b* II o 1- 0. b 2 o 10 O -1 < 2 tr LU X 1- 1 ^A-^O AISI 316 Steels (annealed) o Heat 1 • Heat 2 (this study) a Heat 3 200 400 600 800 TEMPERATURE, K Fig. 8. Temperature dependence of the thermal stress component, illus- trating applicability of thermally activated plastic flow theory at temperatures between 900 and 175 K. 180 so O X o UJ GC CO UJ >- CO o o < UJ CO 0.1 0.2 0.3 0.4 0.5 TEST TEMPERATURE, K MELTING TEMPERATURE, K 0.6 T 0.7 M Fig. 9. Normalized yield strength versus normalized temperature plot for three AISI 316 stainless steel heats, showing variations in thermal and athermal stress components. 181 500 400 o Q_ I- O -z. CC h- C/) 300 200- 100 1 1 1 1 1 1 1 1 1 \ °"y *\ \ \* \ — \ \\ - \ \ X — \ V — \ \ — \. - • : 1 1 1 1 1 1 1 1 1600 1400 0J E 1200 <| HIOOO .. h- -z. UJ 800 o Ld 600 8 LlI N -400 w < 200 § °6 100 200 300 400 500 600 700 800 900 1000 TEMPERATURE, K Fig. 10. Dependence of the tensile yield strength, a ; the Hall-Petch parameter, K; and the flow strength, a., for AISI 316 steel. High temperature data from Norstrom 25 and a. is adjusted to 0.04 weight percent nitrogen. 182 900 CO a. b" O z HI CO h- co UJ 700 500 ^ 300 100 T 1 AISI 316 Steels (annealed) Heat 2 (this study) ± ± 100 200 TEMPERATURE, K - 205 210 - 200 - 195 300 (0 O m LU CO -J O o CO b z o Fig. 11. Temperature dependence of yield stress and Young's modulus for AISI 316 stainless steel heats, showing no drop in yield strength at the Neel transition temperature. 183 LOW -TEMPERATURE DEFORMATION OF 5083-0 ALUMINUM ALLOY R. L. Tobler and R. P. Reed Fracture and Deformation Division National Bureau of Standards Boulder, Colorado ABSTRACT This paper describes the temperature dependence of the yield strength and stress-strain characteristics for 5083-0 aluminum alloy. New experimental data in the temperature range 4 to 340 K are compared with previously published data extending to the melting point. Ther- mally activated plastic deformation apparently governs the behavior between 4 and 400 K, but the athermal component is not simply a function of the shear modulus. Aspects of serrated yielding and load-versus- displacement control tensile tests are discussed. 185 INTRODUCTION Aluminum alloys have been proposed for structural applications in superconducting magnets for magnetic fusion energy systems. The most widely used aluminum alloy in cryogenic service today is 5083-0 alu- minum, which is used extensively in thick sections for the storage of liquefied natural gas. This alloy is strengthened by magnesium in solid solution and is relatively soft, but tough. This provides the advantage, however, that strength can be maintained in the welded joints of struc- tures made with this alloy. The low-temperature properties of relatively pure aluminum have been reviewed by Reed [1]. The purpose of this study was to charac- terize the plastic deformation of 5083-0 alloy over a broad temperature range. Although scatter in the studied material prevented determination of the exact form of the temperature dependence of the yield strength, several observations are recorded concerning the effects of temperature, magnesium concentration, and method of testing on the low-temperature strength of aluminum alloys. This work provides a basis for future studies. PROCEDURE Material A 3.8-cm-thick plate of 5083-0 aluminum was obtained from a com- mercial manufacturer. The ASTM specification B 209-66 gives the chemical composition in weight percent as follows: 4.0-4.9 Mg, 0.3-1.0 Mn, 0.05- 0.25 Cr, 0.4 max Si, 0.4 max Fe, 0.10 max Cu, 0.25 max Zn, 0.15 max Ti , and 0.15 max of other impurities. The microstructure consists of elon- gated grains in the rolling direction, as illustrated in Figure 1. Typical Rockwell hardness of the alloy is R g 61. 187 Specimen Tensile tests were performed using unnotched, longitudinally oriented specimens having a 50.8-mm gage length and 2.54 x 4.80 ym cross section. Specimens were cut from the 3.8-cm-thick plate (at various thickness depths), then annealed at 686 K and furnace cooled. Apparatus Most tests were performed using the screw-driven machine and vari- able temperature cryostat described by Read and Reed [2], The crosshead rate during these tests was 8.5 ym/s. The tensile yield strength was recorded at 0.002 offset specimen strain. The specified test tempera- tures ranged from 4 to 340 K and were accurate to better than ± 1.5 K. A few tests were also performed using a servo-hydraulic test machine in load control at a loading rate of 0.1 kN/s. Tests were performed at temperatures between 4 and 76 K in a gaseous helium environment and between 76 and 295 K in a gaseous nitrogen atmosphere. RESULTS AND DISCUSSION Tensile Yield Strength In Figure 2, the tensile yield strength between 4 and 300 K is presented. The data scatter is large, with six measurements of the room-temperature yield strength ranging from 137 MPa to 158 MPa. It is thought that this scatter is real and reflects inhomogeneity in the 3.8-cm-thick 5083-0 plate. Unfortunately, the scatter is too large to permit accurate interpretation of the temperature dependence at low temperatures. The trend shown in Figure 2 follows the minimum values measured in this study. 188 The present results are combined with high temperature results [3] in Figure 3. In the temperature range 150 to 400 K, the yield strength (a ) is constant, if one relies on the handbook values in this tempera- ture interval. Normalizing the yield strength by dividing by the shear modulus (G), which is reported by Weston, Naimon, and Ledbetter [4], results in a negative temperature dependence of a /G in this temperature range (Fig. 3). The rather strong thermal contribution to the flow strength is most likely associated with dynamic solute (Mg) segregation to dislocations. Evidence for this effect is the discontinuous yielding observed in the stress-versus-strain behavior of the alloy at room temperature that disappears at 76 K. Owing to the rather strong dynamic strain aging contribution, it is not possible from these data to sepa- rate the athermal from the thermal components. Stress-Strain Behavior True stress-true strain curves for the 5083-0 aluminum alloy at 295, 76, and 4 K are presented in Figure 4. Here true stress is defined L-L L-L as P/A (1 + — r — ) and true strain equals In (1 + — r — ) where P is load, o o A is the original cross-sectional area, L is the original gage length, and L is the instantaneous gage length. Serrated yielding phenomena were observed at 4 K and at 295 K, but not at 76 K. In the case of 5083-0 alloy at 4 K, the stress-strain curve was smooth up to a strain of about 4%, at which point the flow became discontinuous. The stress- strain curve at 295 K showed single serrations at 1.8 and 2.5% strain, then continuous serrations from 2.1% strain to fracture. The serrated yielding at 295 K is well known in aluminum-magnesium alloys as the Portevain-LeChatelier effect, which is accompanied by 189 L'uders band propagation and is attributed to dynamic strain aging [5]. In the study reported by Lloyd [6], dynamic strain aging was observed at temperatures between 200 and 350 K. In our study, the stress-strain curve at 295 K showed continuous serrations from 3% plastic strain to fracture. The strain at which the serrations begin and the magnitude of the load drops associated with each serration have been demonstrated to be dependent on strain rate, magnesium concentration, and temperature [7] The serrated yielding at 4 K is attributed to adiabatic heating, a completely different mechanism from that of dynamic strain aging at 295 K. Adiabatic heating is possible because of the low specific heat of the alloy at 4 K [8]. Thus the amount of heat necessary to produce an appreciable temperature rise is relatively low. As a result of heating during plastic deformation, appreciable softening occurs and the deformation becomes unstable. Low thermal conductivity at 4 K (compared with that at 295 K) contributes to a localization of the heat produced until flow occurs. During testing at 4 K, the load drops suddenly and the flow becomes localized when thermal softening, due to heat generated in plastic straining, is greater than the strain-rate hardening. This instability criterion is more easily satisfied at the lowest temperatures because 1) the volume specific heat is reduced so that relatively large tempera- ture increases are possible in straining and 2) the flow stress is more strongly temperature dependent. Specimens tested using load control at 4 K were compared with those tested using stroke control . The stress-strain curves obtained are compared in Figure 5. The adiabatic instabilities corresponded to specimen strains of the order of 0.04%. The load-controlled tests did 190 not produce catastrophic deformation during the discontinuous yields. Instead, each local serration is arrested at some strain and temperature where the flow process becomes stable, owing to local specimen strain hardening which distributes the load over an ever-increasing area that precludes continued temperature increases in the local area in which the process initiated. Magnesium Strengthening Tensile yield strength data at low temperatures for three com- mercial aluminum-magnesium alloys (5083, 5154, and 5052) are presented in Figure 6. The 5083-0 results represent the average data at 295, 76 and 4 K from the present study, whereas the 5154 and 5052 data were previously unpublished [7]. In general, the temperature dependences are \/ery similar and little low- temperature strengthening is apparent. Clearly, the addition of magnesium serves to increase the room-temperature yield strength and this increase is maintained at low temperatures. In Figure 7 the effects from magnesium additions to aluminum are plotted at two test temperatures, 295 and 20 K. Again, the dependence of yield strength on magnesium content at these two temperatures appears wery similar. This indicates that, contrary to nitrogen and carbon inter- stitial contributions to austenitic stainless steels, the addition of magnesium serves to increase the long-range stress fields through solid- solution strengthening and does not affect the thermal contribution to the flow strength. Data scatter in Figure 7 may represent the effects of grain size or relate to lack of solute homogeneity. 191 SUMMARY Tensile yield strength was measured for the 3.8-cm-thick plate of aluminum alloy 5083-0 at temperatures ranging from 4 to 340 K. The scatter of yield strength measurements was sizable, preventing an exact determination of the temperature dependence. However, the following results were obtained: 1. The stress-strain behavior of 5083-0 is complex: at 295 K discontinuous yielding from dynamic strain aging occurs; at 4 K discontinuous yielding from adiabatic specimen heating effects occurs. 2. Tensile yield strength is enhanced at room temperature by dynamic strain aging from magnesium solute. 3. The strengthening contribution of magnesium essentially results from its influence on long-range stress fields; little strength- ening from short-range solute clustering or precipitation was measurable from comparison of 295 and 20 K data. ACKNOWLEDGMENTS We thank D. Beard, Office of Fusion Energy, Department of Energy, for his support and encouragement. D. Beekman and D. Baxter of the National Bureau of Standards, Boulder, Colorado contributed by per- forming tensile tests, reducing data, and conducting metallographic analyses. Their assistance is gratefully acknowledged. 192 REFERENCES 1. Reed, R. P., Aluminum: A review of deformation properties of high purity aluminum and dilute aluminum alloys. Cryogenics, 12 (1972) 259-291. 2. Read, D. T. and Reed, R. P., Toughness, fatigue crack growth, and tensile properties of three nitrogen-strengthened stainless steels at cryogenic temperatures. The Metal Science of Stainless Steels , E. W. Collings and H. W. King, eds., American Institute of Mining, Metallurgical, and Petroleum Engineers, New York (1979) pp. 92-121. 3. Metals Handbook , Vol. 2, Properties and Selection of Nonferrous Alloys and Pure Metals, American Society for Metals, Metals Park, Ohio (1979). 4. Weston, W. F., Naimon, E. R., and Ledbetter, H. M., Low temperature elastic properties of aluminum 5083-0 and four ferritic nickel steels, in Properties of Materials for Liquefied Natural Gas Tankage , ASTM STP 579, American Society for Testing and Materials, Philadelphia (1975), pp. 297-420. 5. Morris, J. G., Dynamic strain aging in aluminum alloys, Mater. Sci . Eng. 13 (1974), 101-108. 6. Lloyd, D. J., The deformation of commercial aluminum-magnesium alloys, Metall. Trans. 11A (1980), 1287-1294. 7. Reed, R. P., Unpublished data, National Bureau of Standards, Boulder, Colorado (1980). 8. Read, D. T. and Reed, R. P., Heating effects during tensile tests of AISI 304L stainless steel at 4 K, Advances in Cryogenic Engineering , Vol. 26, Plenum Press, New York (1980), pp. 91-101. 193 LIST OF FIGURES Figure 1. Three-dimensional metal lographic representation of aluminum alloy 5083-0. Figure 2. Tensile yield strength of 5083-0 aluminum alloy at low temper- atures. Figure 3. Variation of temperature dependence of tensile yield strength of 5083-0 aluminum alloy from 4 K to the melting temperature. Figure 4. Tensile true stress-true strain curves of 5083-0 aluminum alloy at 295, 76, and 4 K. Figure 5. Load and displacement controlled tensile stress-strain be- havior at 4 K. Figure 6. Tensile yield strength at low temperatures from aluminum- magnesium alloys 5083, 5154, and 5052 in the annealed condi- tion. Figure 7. Influence of magnesium alloying additions in aluminum on tensile yield strength at 295 and 20 K. 194 •A \ ^r% L \ // ' ' / ¥" Figure 1. Three-dimensional metal! ographic representation of aluminum alloy 5083-0. 195 200 LU 120 5083-0 (annealed) 100 200 (0 300 TEMPERATURE, K Figure 2. Tensile yield strength of 5083-0 aluminum alloy at low temper atures. 196 200 50 £ -z. LU cr h- co Q _l LU >- 100 50 5083-0 Aluminum Annealed 77/1 This Study, «avg. o Handbook 1 1 Solidus, 847 K O O if) Q O or < LU X if) -z. LU or i- (/) Q _l LU >- 100 200 300 400 500 600 700 800 900 TEMPERATURE (K ) Figure 3. Variation of temperature dependence of tensile yield strength of 5083-0 aluminum alloy from 4 K to the melting temperature. 197 800 Figure 4. Tensile true stress-true strain curves of 5083-0 aluminum alloy at 295, 76, and 4 K. 198 Load Control Aluminum Alloy 5083-0 4K Displacement Control Figure 5. Load and displacement controlled tensile stress-strain be- havior at 4 K. 199 (0 Q. 2 200 X H O z LU OC h- (/> LU - 160- 120- 28 24 (0 20 -16 O z UJ cc H 0) UJ -112 100 200 300 TEMPERATURE, K Figure 6. Tensile yield strength at low temperatures from aluminum- magnesium alloys 5083, 5154, and 5052 in the annealed condi- tion. 200 200 £ o Z UJ or h- co Q _l UJ > UJ _J UJ 150 100 TS" Q This study D USSR alloys I , •, X Unpublished NBS wt % Mg Figure 7. Influence of magnesium alloying additions in aluminum on tensile yield strength at 295 and 20 K. 201 EFFECTS OF CARBON AND NITROGEN ON THE ELASTIC CONSTANTS OF STAINLESS STEEL 304 H. M. Ledbetter Fracture and Deformation Division National Bureau of Standards Boulder, Colorado ABSTRACT Nine stainless-steel 304-type alloys were studied at room tempera- ture. Carbon-plus-nitrogen contents of these alloys ranged from 0.067 to 0.325 weight percent (0.3 to 1.3 atomic percent). Five elastic constants—longitudinal modulus, Young's modulus, shear modulus, bulk modulus, and Poisson's ratio—were determined by a pulse-echo ultrasonic method. Within an experimental uncertainty of about 0.1 percent, all the elastic constants remain unaffected by carbon and nitrogen addi- tions. 203 Introduction This study sought to determine how carbon and nitrogen alloying ad- ditions affect the elastic constants of stainless-steel 304-type alloys. Both carbon and nitrogen increase strength of austenitic Fe-Cr-Ni alloys. Nitrogen stabilizes the austenite, reducing the tendency at low temperatures toward martensitic transformations. Carbon functions simi- larly, but higher r . rs ?' rations cause welding sensitization. The in- creasing popularity of these alloys together with a dearth of reliable physical -property measurements prompted the present study. Being physical properties, elastic constants are relatively in- sensitive to details of internal structure related to thermomechanical treatment. However, when determined with sufficient sensitivity, they provide important parameters for indicating interatomic-force changes, which may occur during alloying. Indeed, small changes in elastic constants may indicate large changes in the nature of the interatomic bonding . Experiment Materials Nine stainless-steel plates were purchased from the research labo- ratory of a U.S. steel manufacturer. All plates had nearly the same base composition, which fell within the limits of ASTM specification A240 for AISI-304 stainless steel, nominally Fe-18Cr-10Ni-l.5Mn-0.02P- 0.02S-0.55Si-0.2Mo-0.2Cu weight percent. The carbon and nitrogen con- tents varied as shown in the mill chemical analyses in Table I. The nine plates were produced from three 1 36- kg vacuum-induction-melted heats, split with respect to carbon level, and teemed into 7.6 x 20 x 36 cm hot-topped cast-iron ingot molds. The ingots were reheated and 205 soaked at 1561 K (2350°F), hot-rolled into 2.54-cm plates, and air cooled. The plates were finally annealed at 1332 ± 7 K (1937 ± 13°F) for 1 h and water quenched. Hardness and grain size are also listed in Table I, and mass densities are given in Table II. Methods Room-temperature sound velocities were determined by a pulse-echo- overlap technique. Quartz piezoelectric crystals with fundamental resonances between 4 and 7 MHz were cemented with phenyl salicylate to one end of a specimen with flat and parallel surfaces. An x-cut trans- ducer was used for longitudinal waves and an ac-cut for transverse waves. The specimen in this case was a 1.59 x 1.59 x 1.27 cm paral- lelepiped with opposite faces ground flat and parallel within 2.5 urn. Ultrasonic pulses 1 to 2 cycles long were launched into the specimen by electrically exciting the transducer. The pulses propagated through the specimen, reflected from the opposite face, and propagated back and forth. The pulse echoes were detected by the transducer and displayed on an oscilloscope equipped with a time delay and a microprocessor for time-interval measurements. The sound velocity was computed by v = 2£/t (1) where i denotes specimen length, and t the round-trip transit time. On the oscilloscope, t was the time between adjacent echoes, the first and second echoes usually being measured, and within these the time between leading cycles. Elastic constants were computed from the general rela- tionship C = P v 2 (2) 206 where p denotes mass density. The usual engineering elastic constants are related to the longitudinal and transverse sound velocities, v and v t , by 2 longitudinal modulus = C = pv (3) X/ A/ 2 shear modulus = G = pv. bulk modulus = B = C £ -(4/3)G Young's modulus = E = 3GB/(C £ -G) (6) Poisson's ratio = v = (E/2G)-1 = (l/2)(C r 2G)/(C r G) (7) Results Table II gives the principal results of the study: sound velocities and elastic constants for nine stainless-steel type-304 alloys with dif- ferent carbon and nitrogen contents. Table III shows for one case, alloy 1, the directional variation of the ultrasonic wave velocities. Table III implies also that ultrasonic velocities can be measured in these alloys with an uncertainty of 0.05 percent or less. Discussion Results in Table II lead to the following conclusion: within our measurement capability (0.1 percent or less) all elastic constants of stainless-steel type-304 alloys remain unaffected by carbon-plus-nitrogen additions up to 1.3 atomic percent. Various statistical analyses of the results support this conclusion. And it can be seen quite simply by comparing the results for specimens 1 and 9, which have the lowest and 207 highest carbon-plus-nitrogen additions: the two sets of elastic con- stants are nearly identical; neither volume-deformation nor shear- deformation elastic constants change significantly nor show any trend versus composition. With the exception of specimen 4, all sound- velocity results fall within the one-standard-deviation range estab- lished for 304-type alloys by Ledbetter, Frederick, and Austin. And specimen four is within 1.2 standard deviations. Another strong reason for finding no alloying effect comes from comparing the variation of B and G among the nine alloys. Interstitial atoms, especially those in isotropic, octahedral positions, should affect B more than G because they cannot move under a dilation force. Yet, among the nine alloys B and G show the same variations from the mean. Present results contrast sharply with observations reported for 2 ferritic iron. Speich, Schwoeble, and Leslie reported that carbon added to alpha iron reduces elastic stiffnesses such as E and G by about five percent per atomic percent carbon! Indeed, carbon is the most ef- fective known alloying element in altering alpha iron's elastic con- stants. Quantitative results for nitrogen remain unreported, but one would expect nitrogen to behave similarly to carbon because their atomic diameters are similar. The essential distinction between the ferritic and austenitic cases is the crystal structure. Body-centered-cubic crystals are less accommodating to interstitials than are face-centered- cubic crystals. Carbon alloyed into bcc iron occupies octahedral inter- stices that have tetragonal symmetry, which produce a strong tetragonal distortion; no similar distortion occurs in fee iron, where carbon occupies larger octahedral interstices and the lattice distortion is 208 isotropic. This distortion reflects also in the much higher solubility of carbon and nitrogen in fee iron compared with bcc. Present results dispute those of a previous study. Laddach and 3 Schuler studied the influence of nitrogen on the elastic properties of an Fe-21Cr-20Ni alloy containing 0.02 and 0.29 weight percent nitrogen. They concluded that "nitrogen raises the modulus of elasticity as well as the modulus of rigidity, regardless of the nitrogen being either in interstitial solution or precipitated in the form of nitrides." No reconciliation of f Ms conclusion and that of the present study seems apparent. Conclusion In the studied composition range, elastic constants are determined by interactions among atoms occupying lattice sites; carbon and nitrogen interstitial atoms play no significant part. Since elastic constants are determined by the interatomic force constants, they also must be independent of carbon and nitrogen content. Thus, the present result has importance also for other force-constant-determined physical proper- ties, such as thermal expansivity and specific heat. Acknowledgments The DoE supported part of this study. R. L. Tobler of NBS con- tributed the materials and the grain-size measurements. J. Dahnke of NBS assisted in measurements. References 1. H. M. Ledbetter, N. V. Frederick, and M. W. Austin, J. Appl . Phys. 51_, 305-309 (1980). 209 2. G. R. Speich, A. J. Schwoeble, and W. C. Leslie, Metall. Trans. 3_ 2031-2037 (1972). 3. H. Laddach and P. Schuler, DEW-Techn. Ber. 13, 85-88 (1973). 210 CD M in *~* E e p. fO CJ3 CO CO i — CD i— C CD ■O 2 S- -^ (O o n: o or rt5 S- Cn -a 03 CO ai c -o 5- n3 CD u aj CL CO o Q. E O o cu x: o JO + c_) ificoomLficocoNO con ai o coi — cr> oo cn ^i-i — inoiinNnOw ro -=3- CO ■ — oo^d-r—oooo r--r--i--.cococococooo NNCONN o LT> LT) cn r»» CO o O o o o o CM o CNJ CM o "5fr CNJ o o cn CM cn m o oo s_ • • • o oo oo o • — ■ — CM i — CMro^d-cncor^cocn 211 Table II. Mass density, sound velocities, and elastic constants of nine carbon-plus-nitrogen-alloyed stainless-steel 304 alloys. Density (g/cm 3 ) (cm/ys) v t (cm/us) f S (GPa) E (GPa) G (GPa) B (GPa) V 1 7.944 0.5751 0.3140 262.7 201 .7 78.3 158.3 0.288 2 7.937 0.5757 0.3139 263.1 201 .5 78.2 158.8 0.288 3 7.936 0.5751 0.3142 262.5 201 .7 78.3 158.0 0.287 4 7.935 0.5743 0.3155 261.7 202.8 79.0 156.4 0.284 5 7.927 0.5747 0.3150 261 .8 202.2 78.7 156.9 0.285 6 7.923 0.5748 0.3144 261 .8 201 .5 78.3 157.3 0.287 7 7.869 0.5769 0.3141 261 .9 200.2 77.6 158.4 0.289 8 7.883 0.5764 0.3135 261 .9 199.9 77.5 158.6 0.290 9 7.900 0.5766 0.3146 262.6 201 .4 78.2 158.4 0.288 Mean 0.5755 0.3144 262.2 201 .4 78.2 157.9 0.287 Std. dev. 0.0009 0.0006 0.5 0.9 0.5 0.8 0.002 Pet. std. error 0.2 0.2 0.2 0.4 0.6 0.5 0.7 212 I Table III. Directional variation of ultrasonic wave velocities in specimen 1, where z direction is perpendicular to rolling plane. Propagation direction Polarization direction v £ (cm/ys) v t (cm/ys) X X 0.5746 X y 0.3140 X z 0.3143 y y 0.5753 y X 0.3139 y z 0.3138 z z 0.5753 z X 0.3143 z y 0.3140 Mean Std. dev. Pet. std. error 0.5751 0.3141 0.0004 0.0002 0.07 0.07 213 LOW-TEMPERATURE VARIABILITY OF STAINLESS-STEEL-304 ELASTIC CONSTANTS H. M. Ledbetter Fracture and Deformation Division National Bureau of Standards Boulder, Colorado ABSTRACT By measuring ultrasonic velocities, the elastic-constant behavior of four stainless-steel 304- type alloys was determined between ambient and liquid- helium temperatures. Despite their equivalence near room temperature, results show that cooling through the magnetic transition produces slightly different elastic constants in different alloys at low temperatures. 215 INTRODUCTION Recently the author 1 collaborated in studying the room- temperature variability of stainless-steel 304- type elastic constants: Young's modulus, shear modulus, bulk modulus, Poisson's ratio. For twenty samples acquired randomly, the lot-to-lot variation was surprisingly small: 0.9 percent for Young's modulus, for example. Thus, for this commercial alloy, chemical -compo- sition differences and thermomechanical -processing differences affect the elas- tic constants practically negligibly. Because of many cryogenic applications for 300-series austenitic stainless steels, the author^ recently determined accurately the low- temperature elas- tic constants of three such alloys: 304, 310, 316. At low temperatures all three alloys undergo a magnetic transition, which affects the elastic con- stants. The 304-alloy results given in that study were representative rather than comprehensive. Different 304 alloys actually behave slightly differently at low temperatures. The present study determined approximately the low- temperature variability of stainless-steel -304 elastic constants. The longitudinal sound velocity was measured between 295 and 4 K for four steels: one regular 304 alloy, two low- carbon alloys, and one nitrogen-enriched alloy. The author-^ showed recently that interstitial alloying elements such as carbon and nitrogen do not affect the room- temperature elastic constants of 304- type stainless steels. The longitudinal modulus manifests both volume and shear effects and is therefore sensitive to any change in the elastic constants. 217 EXPERIMENTS Low- temperature ultrasonic velocities were determined similarly to room- temperature measurements made in our laboratory by a pulse-echo-superposition method. 1 * 3 The notable difference was that for low temperatures the trans- ducer-specimen bond consisted of stopcock grease rather than phenyl salicylate. The specimen was inserted into a probe, 4 which was inserted into the ullage of a helium dewar. Measurements were made at 5-K intervals during continuous cooling or warming at rates of 1 to 2 K per minute. Temperature was monitored with a chromel-constantan thermocouple placed close to the specimen. Table I shows the chemical compositions, hardnesses, and mass densities of the studied alloys. RESULTS Figure 1 shows the experimental results graphically: relative longitudi- nal modulus versus temperature for six alloys. To determine the slope dC/dT more accurately, phenyl salicylate bonds were used in two cases and cooled only to 245 K. Table II gives quantitative characteristics of the cooling curves (which reproduced during subsequent reheating, showing no hysteresis). "Max" indicates the maximum value of the C(T)/C(295) curve, "min" indicates the minimum value at low temperatures, and A the difference between the two. The transition temperature, T t > corresponds to A/2. The slope dC/dT was deter- mined by a linear-regression fit to eleven points between 245 and 295 K, and r^ is the coefficient of determination. 218 DISCUSSION Elastic constants reflect the magnetic-interaction energy of a system according to C.. k , = 3 2 U/3n 1j 3n kl (1) where C-fj^i denotes the fourth-rank elastic-stiffness tensor, U denotes internal -energy density (including magnetic energy), and n-jj denotes Brugger strains. Four types of magnetic energy were discussed by Ledbetter and Collings 5 in a study of Fe-Cr-Ni-Mn steels. A priori , one expects strong magnetic interactions in these steels because they contain four magnetic elements: iron, chromium, nickel, and manganese. Fundamental calculations by Moruzzi , Janak, and Williams 6 show that for these elements, especially Mn, magnetic energy affects properties: cohesive energy is reduced, bulk modulus reduced, and volume increased. At lower temperatures, because lattice vibrations are reduced and because magnetic atoms move closer together because of thermal contraction, magnetic interactions become stronger. Stronger magnetic interactions manifest themselves in physical -property changes. The existence and complexity of low- temperature magnetic transitions in 304-type alloys have been discussed especially by Collings and Hart.' They view austenitic stainless steels as completely paramagnetic at high tempera- tures. At some lower temperature, clustered ferromagnetic ordering leads to 219 superparamagnetism; at still lower temperatures, the paramagnetic matrix under- goes spin-glass condensation. Interactions among spin clusters and with the matrix lead to further possible complexities. In principle, all these magnetic transitions affect the elastic constants. The reversible two-percent elastic-constant softening during cooling shown in Fig. 1 near T = 47 K results from a magnetic transition, probably the spin- glass condensation. Why should spin-glass condensation cause elastic softening? Landau and Lifshitz^ proved for second-order phase transitions that "compressibility drops for a transition from an unsymmetrical to a symmetrical phase." Thus, the bulk modulus drops for a transition from a symmetrical (paramagnetic) phase to an unsymmetrical (magnetic) phase. Magnetic in this case means spin-glass. Spin-glass properties arise, apparently, from magnetic-moment interac- tions. Principal among these is the s-d or Rudermann-Kittel-Kasuya-Yosida (RKKY) indirect-exchange interaction. The energy function for this spin system is E- -(1/2)^^1.-1. (2) where J^ = jCft^) = J(^. - ft.) (3) J(R) = A cos 2k f R/(2k f R) 3 (4) 220 where the spin positions Rj are random variables, kf is the fermi-level wavevector of the host conduction electrons, and J(0) = 0. Interaction (4) alternates in sign and is long range, representing alternatingly ferromagnetic and anti ferromagnetic spin coupling. This type of magnetic interaction can become important when the usual S -,• • S j direct exchange becomes insignificant, that is, when the magnetic ions are separated so widely that their overlap integral is zero. The basic idea in RKKY exchange is that an s conduction electron interacts with a localized d electron through a direct-exchange inter- action S s »sy, travels through the lattice, and then has a second such inter- action S s *Sj. This amounts to an effective (indirect) interaction tj*3j between the two localized d electrons at Rj and Rj. The anomalous elastic-constant behavior shown in Fig. 1 relates in the most general sense to magneto-elastic energy, to a mechanism that couples a crystal's elastic vibrations to its magnetic modes. Here we are especially concerned with the magneto-elastic energy arising from the strain dependence of the J-jjS-j'Sj indirect exchange interaction. From Hausch and Warlimont, y who used a molecular-field approximation, and for simplicity using Voigt averaging, it follows that the magnetic-exchange-energy contributions to the bulk modulus, B, and the shear modulus, G, for the face-centered-cubic first neighbor-only case are: B = -4N(1 - L) (r* J" - Zr J') (5) ex T. o o 221 and e ex " I** 1 " T t > {r o J " + 4r o J,) (6) where N = atoms/volume, T t = transition temperature, r = equilibrium atomic spacing, and J' and J" are the first and second derivatives of J with respect to r evaluated at r = r . Since J(r) is not known a priori for 3d metals and alloys, J' and J" must be determined experimentally. In principle, J' and J" can be estimated from measurements made in the present study. How- ever, the many approximations and the alloy's complexity render this approach dubious. Equations (5) and (6) are set out mainly to give the flavor of a final explanation, which awaits an ab initio relationship for J(r), discussions of which DehlingerlO gave for alloys. ACKNOWLEDGMENT The Department of Energy supported this study. Several people contributed experimentally: M. W. Austin, J. Dahnke, and G. Maerz. REFERENCES 1. H. M. Ledbetter, N. F. Frederick, and M. W. Austin, J. Appl . Phys. 5JU 305-309 (1980). 2. H. M. Ledbetter, J. Appl. Phys., forthcoming. 222 3. H. M. Ledbetter, unpublished research. 4. E. R. Naimon, W. F. Weston, and H. M. Ledbetter, Cryogenics J4, 246-249 (1974). 5. H. M. Ledbetter and E. W. Collings, in The Metal Science of Stainless Steel s (Met. Soc. AIME, New York, 1979), pp. 22-40. 6. V. L. Moruzzi , J. F. Janak, and A. R. Williams, Calculated Electronic Properties of Metals (Pergamon, New York, 1978), p. 6. 7. E. W. Collings and S. C. Hart, Cryogenics ^9, 521-530 (1960). 8. L. D. Landau and E. M. Lifshitz, Statistical Physics (Pergamon, London, 1959), p. 438. 9. G. Hausch and H. Warlimont, Z. Metallkd. 63, 547-552 (1972). 10. U. Dehlinger, Z. Metallkd. 28, 116 (1936); 28, 194 (1936); 29, 388 (1937). Reviewed in F. Seitz, The Modern Theory of Solids (McGraw Hill, New York, 1940), pp. 623-626. 223 to o> ■p s_ 0) Q. o i- Q- S- o ■o c +-> e •r— 0) o to O O- E o a to o •^ E o a) *0 >> 4->CO to •r- E to CO (J fO c ^s» ^ a» cr> Q ,_^ CO (/) oo r— QJ r— c O CTi O LO CO 1^ r-H CO 1^ cvj CM CO LO r^ CO lo O 1^- r- r~~. LO o LO o CO o CM LO I — o CM i— I o 1— 1 o O i— ( o 1^. 1—1 o CO o CM o o o o o o CM O CM CM O o CM o LO CM o CI 1— 1 o O O o o o CM O LO CM o 1—1 LO o o CO o CO o lo CM CO CTi LO r-l CO 00 CO o o o o o *d- LO C 44 --of stainless steel 304 are estimated from the polycrystall ine bulk and shear moduli together with an empirical C, ? /C,, value, which is dis- cussed theoretically. The estimate involves a reverse Kroner method for relating single crystal and polycrystall ine elastic constants. 227 A solid's engineering elastic constants, such as Young's modulus and Poisson's ratio, serve many practical purposes. Among these are calculations of load-displacement, thermoelastic stress, elastic sta- bility (buckling), residual stress, and fracture-toughness plane-stress- to-plane-strain conversions. A solid's single crystal (anisotropic) elastic constants, the C..'s and the S..'s, provide additional and more fundamental information than do the engineering (quasi-isotropic) elastic constants. The C.-'s are essential, for example, for theory of point defects, anisotropic dislo- cation interactions, elastic-energy calculations, force-constant-related properties such as lattice vibrations, accurate determination of Debye characteristic temperatures, correlations with electronic theory of solids, and correlations with other physical properties. Stainless steels form a key component of materials technology, and their basic properties are now being studied by methods applied formerly to elements or to simple binary alloys. Elastic constants comprise one subset of basic mechanical-physical properties, which must be known and understood before the metal physics of stainless steels can be carried yery far. Three existing experimental studies [1-3] of the single-crystal elastic constants of Fe-Cr-Ni alloys omit the Fe-18Cr-8Ni alloy, which is of most interest because of large tonnages and multifarious applica- tions and because it provides a base material for developing many other special purpose stainless steels. In principle, these three studies on four alloys--Fe-12Cr-18Ni, Fe-18Cr-12Ni , Fe-18Cr-14Ni , Fe-18Cr-19Ni — should permit by linear-regression analysis the determination of the single-crystal elastic constants of Fe-18Cr-8Ni. But, this analysis 229 fails. Reasons for this probably include: a small sample set (four alloys for three independent coefficients), poor distribution of sample chemical compositions, experimental errors, nonlinear dependence on composition, and possible nonadditive, interactive composition effects. The present study predicts the single-crystal elastic constants of Fe-18Cr-8Ni. The prediction depends on three ingredients: one observa- tional, one theoretical, and one empirical, which has some theoretical basis. Observationally, Ledbetter, Frederick, and Austin [4] measured the polycrystall ine elastic constants of twenty randomly acquired Fe- 18Cr-8Ni alloys with the results that the bulk modulus, B, and the shear modulus, G, are well defined: B = 1.582 and G = 0.774-10 11 N/m 2 . Theo- retically, Kroner [5] showed how the polycrystall ine elastic constants (two independent, e.g., B and G) can be calculated from the single- crystal elastic constants (three independent for the cubic-symmetry case: C,,, C, 2 > and C**). Empirically, as shown in table 1 for Fe, Ni, and the four Fe-Cr-Ni alloys described above, the dimensionless ratio y = C^/C-i-i is surprisingly constant for these materials: 0.635 ± 0.031. In theory, y varies between zero and one. Both the upper and lower limits correspond to a mechanical instability. For the cubic elements, in practice y varies between 0.12 (diamond) and 0.85 (gold). Before considering the above case, we consider a simple, familiar model: central forces between atoms, which require that C, ? = C,.. Table 1 shows this assumption is not too bad for these materials. We will treat the results as approximate guidelines to a better calcula- tion. Since we are computing three elastic constants (C,,, C, ? , C«.) from two (B and G), we need another relationship among the C.'s. For central interatomic forces, this relationship is 230 ^1 2 = 44 ' ' For all cubic crystals, the bulk modulus relates to the C.'s according to B = j (C n + 2C 12 ) (2) Following Kroner, the shear modulus relates to the C.'s by G 3 + aG 2 + 6G + y = (3) where a = (5C n + 4C ]2 )/8 (4) P = " C 44 (7C 11 " 4C 12 )/8 (5) and Y = -C 44 (C T1 - C 12 )(C n + 2C ]2 )/8 (6) Solving eqs. (1) through (3) simultaneously gives the central-force (cf) results for Fe-18Cr-8Ni shown in table 1. All three C.'s and the bulk modulus fall within the ranges established by the four alloys. Both C = ( C, -j -C-. 2 ) /2 and A = C../C are outside these bounds. Thus, when fitted to observed values of B and G, a central -force model for stain- less steel is not bad and fails mainly for the C elastic constant. We turn now to a more realistic noncentral -force model where C, ? /C,, 0.635. (The basis for this is described below.) Simultaneously solving eqs. (2) and (3) subject to this condition yields the results shown in table 1 in the bottom line. When combined with previous observations, these new theoretical results show how Ni affects elastic constants in the Fe-18Cr-8Ni region. Increasing Ni decreases C,-,, C, 2 » B > and C, 231 but leaves C,, relatively unaffected, perhaps slightly increased; Zener anisotropy increases slightly, indicating a possible tendency toward instability. Except for C and A, results of Salmutter and Stangler [1] suggest that Cr produces opposite effects from Ni . Finally, it remains to try to understand why y = 0.635 for these alloys. If the two-body, central -force interatomic potential is denoted cj>(r), then the Brugger elastic constants, which are fourth-rank tensors, are given at T = K by: C, 8 2 U ijkl 3r >ij 3r i| < i = ^ [£D 2 <»(r)r0r0r° k r0] r = r0 (7) where D4>(r) denotes (l/r)[3(r)/3r]. The energy density U of the crystal is obtained by a sum over two-body atom-atom interaction en- ergies : U = (l/2V°)E*(r), (8) where V° is the undeformed atomic volume. The n-- are components of the Lagrangean elastic strain matrix n = Jd^J-I), (9) where I is the 3x3 identity matrix, t denotes transposition, and the deformation matrix J is defined by r = Jr°, (10) 232 where r° and r are the interatomic spacings in the initial state and in the deformed state, respectively. Details concerning this approach to elastic constants can be found elsewhere [9], Face-centered cubic crystals have twelve nearest neighbors at <^±1 , ±1, 0/a/2, and six second-nearest neighbors at <±1 , 0, ($>a, where a is the unit-cell dimension. In this case one obtains from eq. (7): C ll = 2^ ^(|) 4d2 * (—) + 2aV(r) is already negative at r = a. If so, then y is increased because C,, is decreased while C,,, remains unaffected. Equation (7) represents both structure-dependent and volume-dependent contributions to the C. .'s. But there may be other, purely volume-dependent terms in the energy, often represented as U(V) = AV n (13) where V denotes volume and A and n are positive adjustable parameters. It follows simply that these terms (denoted by asterisks) contribute to the C. .'s as follows: C n * = nA(n-2) = n 2 A-2nA (14) C ]2 * = n 2 A (15) C 44 * = nA (16) 233 Thomas [10] showed that for f.c.c. lattices -nA = /2a(3c}>/3r) (17) evaluated at the nearest-neighbor spacing. Since 3<}>/9r is negative there, nA is positive. (If the volume-dependent energy arises, for example, entirely from the electron gas, then n = -2/3 in a Sommerfeld model.) Thus, C-, 2 * exceeds C,,* and y is increased above 0.50 by volume-dependent energy terms. For cubic metals, only four (Cr, Ir, Mo, W) have C^/C-,, ratios less than 0.5, and only one of these is f.c.c. Values exceeding 0.50 are expected for metals because their electron gas contributes a volume-dependent energy. The near constancy of y among Fe, Ni, and these Fe-Cr-Ni alloys suggests that for some purposes the potential-energy function of d-band transition metals can be treated similarly to the simpler sp-band metals: a sum of pair potentials and a purely volume-dependent energy term. ACKNOWLEDGMENT This study was supported partly by the DoE Office of Fusion Energy and partly by the NBS Office of Standard Reference Data. REFERENCES [1] K. Salmutter and F. Stangler, Z. Metallkd. 51_, 1-4 (1960). [2] M. Kikuchi, Trans. Jap. Inst. Met. 1_2, 417-421 (1971). [3] M. C. Mangalick and N. F. Fiore, Trans. Metall. Soc. AIME 242, 363- 2364 (1968). [4] H. M. Ledbetter, N. V. Frederick, and M. F. Austin, J. Appl . Phys . 51_, 305-309 (1980). 234 [5] E. Kroner, Z. Phys. 151_ s 504-518 (1958). [6] R. A. Johnson, Phys. Rev. V45, 423-433 (1966). [7] D. Diesburg, Ph.D. Thesis, Iowa State University (1971). [8] G. A. Alers, J. R. Neighbors, and H. Sato, J. Phys. Chem. Solids, T3 S 40-55 (1960). [9] R. A. Johnson, Phys. Rev. B 6, 2094-2100 (1972). [10] J. F. Thomas, Scripta Metal 1. 5, 787-790 (1971). 235 Table 1 Single-Crystal Elastic Constants of Several Face-Centered-Cubic Materials Material C ll C 12 C 44 11 2 10 M N/m B C A y Ref. Fe 2.51 1 .57 1.12 1 .88 0.47 2.38 0.625 6 Fe 2.760 1.735 1.363 2.077 0.513 2.41 0.629 7 Ni 2.516 1.544 1.220 1 .868 0.486 2.51 0.614 8 Fe-12Cr- ■18Ni 2.332 1.626 1.225 1.862 0.353 3.48 0.697 1 Fe-18Cr- ■12Ni 1.912 1 .179 1.386 1 .423 0.367 3.78 0.617 2 Fe-18Cr- •14Ni 1.98 1.25 1.22 1.49 0.365 3.34 0.631 3 Fe-18Cr- ■19Ni 1.91 1.19 1 .24 1.43 0.360 3.44 0.623 3 Fe-18Cr-8Ni (cf) 1.99 1.38 1.38 1.58 0.306 4.50 0.692 Present Fe-18Cr-8Ni 2.09 1.33 1.21 1.58 0.382 3.17 0.635 Present 236 ELASTIC-CONSTANT/TEMPERATURE BEHAVIOR OF THREE HARDENED MARAGING STEELS H. M. Ledbetter and M. W. Austin Fracture and Deformation Division National Bureau of Standards Boulder, Colorado ABSTRACT Elastic constants of three maraging steels were determined by measuring ultrasonic velocities. Annealed steels show slightly lower bulk moduli and considerably lower shear moduli than hardened steels. All the elastic constants -- Young's modulus, shear modulus, bulk modulus, and Poisson's ratio -- show regular temperature behavior between 76 K and 400 K. Young's and shear moduli increase with in- creasing yield strength, but the bulk modulus and Poisson's ratio are relatively unchanged. Elastic anisotropy is quite small. 237 INTRODUCTION Elastic constants find many uses. In technological materials they assume importance whenever problems arise concerning deflections, elastic stability (buckling), excessive elastic deformation (jamming), thermoelastic stress, residual stress, fracture, and so on. Ratios of elastic modulus to density or to thermal conductivity provide well-known structural- design parameters. Elastic constants correlate empirically with many practical properties such as strength, hardness, and wear. Because they relate intimately to both the interatomic potential and the lattice- vibration spectrum, elastic constants relate in turn to many fundamental solid-state phenomena such as diffusion, theoretical strength, and lattice-defect energies. Elastic constants appear as coefficients in thermodynamic equations of state where their relationships to free energy, specific heat, and thermal expansivity become especially clear. Since they can be measured accurately (typically within 1 percent or 5 less) and small changes can be detected easily (typically 1 in 10 or better), elastic constants are useful physical properties for characterizing and comparing materials and for detecting changes in material properties. In the present case, the variation of elastic constants with composition and microstructure is of special interest. Ledbetter and Read recently reported the elastic-constant behavior of an annealed 300-grade maraging steel between 4 and 300 K. [The 300- grade steel has a yield stress near 2.07 GPa (300,000 psi).] They reported four elastic constants -- Young's modulus, shear modulus, bulk modulus, and Poisson's ratio. Besides corresponding closely to the elastic constants of an Fe-18Ni alloy, all these constants showed regular temperature behavior. 239 ■ The present report describes a similar, extended study. Three maraging steels -- 200-grade, 250-grade, and 300-grade, both annealed and hardened -- were studied between 76 and 400 K. Elastic anisotropy was studied at room temperature. Despite their relatively poor low- temperature fracture toughness in the hardened condition, these steels occur in many cryogenic applications because of their exceptionally high strength. Because they are relatively new structural alloys, their elastic properties have received little attention. EXPERIMENT Tables I and II characterize the alloys obtained from a commercial source as 2.7-cm and 4.2-cm diameter bars. The pycnometric mass-density determination used distilled water as a standard. Sound velocities were determined on two sizes of cubes -- 1.3 cm and 1.6 cm -- whose opposite faces were flat and parallel within 3 urn. Pulse-echo and pulse-echo-superposition techniques for room-temperature 2 3 and low-temperature measurements, respectively, are described elsewhere. ' Measurements between 300 and 400 K were done in a stirred, heated mineral- oil bath using the oil as the transducer-specimen couplant; a mercury- in-glass thermometer indicated the temperature. Elastic constants were obtained from longitudinal and transverse sound velocities, v and v., by the relationships: 240 2 longitudinal modulus = C £ = pv^ (1) 2 shear modulus = G = pv. (2) bulk modulus = B = C £ - yG (3) Young's modulus = 9GB/(G + 3B) (4) and Poisson's ratio = (E/2G) - 1 (5) where p denotes mass density. RESULTS Table III gives, for the hardened steels, the room-temperature elastic constants and their directional variation. Table IV gives the averaged (over three orthogonal directions) room-temperature elastic constants for the annealed specimens. Table V and Figs. 1-5 give the temperature dependence of the elastic constants. Curves in these figures 4 represent best fits of the data to Varshni s relationship: C = C° - s/(e t/T - 1) (6) which contains three adjustable parameters: C° is the zero-temperature elastic constant, -s/t is the high-temperature slope dC/dT, and t is an approximate Einstein temperature. Figures 1-5 contain only the 76 to 300 K results; the 300 to 400 K data form linear extensions of the 76 to 300 K curves. 241 .5 DISCUSSION In their annealed (Fe-Ni-martensite) condition, these alloys show decreasing E and G and increasing B and v upon progressing from lower to higher yield strength (see Table IV). Since only Mo content increases systematically in this series and its bulk modulus exceeds considerably that of Fe, Ni, and Co, it may cause these property variations. Takeuchi's elastic-constant results confirm that Mo additions decrease Fe's shear modulus. Molybdenum effects on the bulk modulus remain uncertain. Ledbetter and Read found that, except for the bulk modulus, an annealed 300-grade maraging steel has elastic constants similar to those of Fe-18Ni. The present study confirms this. Hardening increases Young's modulus, shear modulus, longitudinal modulus, and bulk modulus but decreases Poisson's ratio (see Tables III and IV). These changes increase upon progressing from 200-grade to 250-grade to 300-grade. These elastic-constant changes suggest their possible use as an indicator of the precipitation process. The elastic-constant temperature behavior is hardly remarkable. It shows all the features of regular behavior: tendency toward zero slope at zero temperature, as required by the third law of thermodynamics; continuous decrease with increasing temperature, consistent with the softening of interatomic bonding forces caused by increased thermal vibrations; and a linear slope at high temperatures. High-temperature linearity is consistent with the theoretical result of Leibfried and Ludwig that the elastic constant depends on temperature according to C(T) = C(0) (1-DU.) (7) where U denotes mean energy and D is a parameter that depends on the type of crystal and the type of oscillator model used (Einstein, Debye, etc.). At high temperatures U varies as kT. 242 ACKNOWLEDGMENT This study was supported by NASA Langley Research Center and by the DoE Office of Fusion Energy. 243 REFERENCES 1. H. M. Ledbetter and D. T. Read, Metal 1. Trans. A 8A, 1805-1808 (1977). 2. E. R. Naimon, W. F. Weston, and H. M. Ledbetter, Cryogenics J_4, 246-250 (1974). 3. H. M. Ledbetter, N. F. Frederick, and M. W. Austin, J. Appl . Phys 51, 305-309 (1980). 4. Y. P. Varshni, Phys. Rev. 2, 3952-3955, (1970). 5. S. Takeuchi, J. Phys. Soc. Jpn. 2_7, 929-940 (1969). 6.. G. Leibfried and W. Ludwig, in Solid State Physics, Volume 12 (Academic, New York, 1961), pp. 275-444. 244 03 C_3 s- r^J on dj Li- ■^ CU l— U c rO r— re XI n™ M JC C71 •r— OJ 2 00 A 00 r— O) CL) -t-> to C s: T3 O) •i— -a =j ■(-> c o •r— -l-> o •r" to O Q. E O u o r— 2: id u E O) .c c_> o o 1 — 1 CO ^— XI ro r— •r - •ZL LT> O O o O I— r— o o o CO CM CO o o o o o o o CM o 00 to m o o ro o o CO o c LD ID o o o o <=d- CM o o o o IX> o CO ro 00 Lf> o LD cr> o o cu CU OJ ■o -o "O r0 (O r0 S- s- s- CD o> cn O o O O LD O CM CM ro 245 1 Table II.. Metallurgical characterization of studied alloys. i Density Hardness Hardening Heat Treatment j (g/cm 3 ) (Rockwell C) Annealed Hardened Annealed Hardened 200-grade 8.068 8.030 28 43 489°C, 4.5 h, air cool 250-grade 8.076 8.066 29 51 same 300-grade 8.087 8.087 31 54 same 246 Table III. Room-temperature elastic constants of three hardened maraging steels B E G V (10 11 N/m 2 ) (10 11 N/m 2 ) (10 11 N/m 2 ) 200-grade x a 1 .654 1.853 0.706 0.313 y 1 .657 1 .848 0.703 0.314 z 1.656 1 .852 0.705 0.314 250-grade X 1 .666 1 .885 0.719 0.311 y 1 .664 1 .884 0.718 0.311 z 1 .663 1 .888 0.720 0.310 300-grade X 1 .680 1 .902 0.725 0.311 y 1 .681 1 .899 0.724 0.312 z 1 .686 1 .899 0.724 0.312 z = bar axis; x and y are perpendicular to z. 247 Table IV. Room-temperature elastic constants of three annealed maraging steels. (10 11 N/m 2 ) (10 11 N/m 2 ) (10 11 N/m 2 ) 200-grade 1.655 250-grade 1.656 300-grade 1.664 1.773 0.671 0.321 1.749 0.661 0.324 1.729 0.652 0.327 248 Table V. Elastic constants of three hardened maraging steels at selected temperatures TOO 6 E G V (10 11 N/m 2 ) (10 11 N/m 2 ) (10 11 N/m 2 ) 200-grade 400 1.63 1.794 0.681 0.317 350 1.643 1.822 0.693 0.315 300 1.655 1.849 0.704 0.314 250 1.667 1.875 0.714 0.313 200 1.678 1.900 0.724 0.311 150 1.687 1.920 0.733 0.310 100 1.694 1.935 0.739 0.310 75 1.696 1.938 0.740 0.309 250-grade 400 1.645 1.826 0.695 0.315 350 1.654 1.855 0.706 0.313 300 1.664 1.883 0.718 0.311 250 1.673 1.910 0.729 0.310 200 1.682 1.935 0.740 0.308 150 1.691 1.956 0.748 0.307 100 1.698 1.971 0.754 0.307 75 1.701 1.976 0.756 0.306 300-grade 400 1.664 1.841 0.700 0.316 350 1.674 1.870 0.712 0.314 300 1.684 1.898 0.723 0.312 250 1.694 1.925 0.734 0.311 200 1.704 1.949 0.744 0.309 150 1.713 1.971 0.753 0.308 75 1.72 1.986 0.759 0.308 75 1.722 1.990 0.761 0.307 249 LIST OF FIGURES Fig. 1. Temperature variation of longitudinal modulus of three hardened maraging steels. Fig. 2. Temperature variation of shear modulus of three hardened maraging steels. Fig. 3. Temperature variation of Young's modulus of three hardened maraging steels. Fig. 4. Temperature variation of bulk modulus of three hardened maraging steels. Fig. 5. Temperature variation of Poisson's ratio of three hardened maraging steels. 250 -a a> (2^w/n iw0i) snnnaoN idNianiiDNcn 251 -a s- O) CD S- o s> Q CO CM s-\ Z3 ^1 13 \— ' O E Lxi a: n3 Z) a) (X CO Ct 4- Ld O CL E 21 O U +-> tS h- a3 Q 2. Temperature vari maraging steels. G> en GO IS. CD in T m Ol rs. IV rs. rs. rs. |S- rs. is. 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Ratios of dynamic (adiabatic) and static (isothermal) Young's moduli, E./E . Temp. (K) Nickel Iron Copper Aluminum 100 1 .000 1 .000 1 .001 1.001 200 1.002 1.001 1 .001 1 .003 300 1 .003 1 .003 1.003 1 .005 400 1.004 1 .004 1 .004 1.007 500 1 .005 1 .005 1.006 1 .008 600 1.006 1.006 1.007 1.010 268 (spupsnoqi ui) ejw A1DI1SV13 dO SmndOW < t— < to o CO 5- CD en c CD +-> s- o 4- S_ o =3 o 1/1 en O o +-> fO •I — S- > CD s_ 13 +J res S- CD o. E CD CD S- =3 en (spuesnom ui) |s>| A1DI1SV13 dO SniTiaOW 269 WELDING AND CASTING 271 EVALUATION OF WELDMENTS AND CASTINGS FOR LIQUID HELIUM SERVICE T. A. Whipple and H. I. McHenry Fracture and Deformation Division National Bureau of Standards Boulder, Colorado ABSTRACT In this paper the FY80 results of the NBS program on the evaluation of weldments and castings for liquid helium service are reviewed. The welding process study for stainless steel castings and the study on the toughness of 5183 aluminum weldments were completed during FY80, and studies on the toughness of stainless steel castings and the fracture and deformation phenomena in duplex austenite/ferrite structures were begun. The research planned for FY81 is also reviewed. 273 Introduction The program to evaluate weldments for structural applications in superconducting magnet systems [1,2,3] was continued in FY80. Except for the addition of the evaluation of stainless steel castings for liquid helium service, the objectives for the program have remained unchanged from those of FY79: 1. To investigate the metallurgical factors that affect the toughness of stainless steel weldments and castings at cryogenic temperatures. 2. To contribute to the development of improved filler metals for welding stainless steels for liquid helium service. 3. To evaluate the strength and toughness of stainless steel weldments and castings and aluminum weldments at 4 K. Research began this year on a study that will take a more fundamental approach to the first objective, and work has continued on the evaluation of experimental stainless steel filler metals, evaluation of aluminum weldments and support of superconducting magnet construction programs. This report summarizes the results obtained in FY80 and outlines the research to be conducted in FY81 . Stainless Steel Weldments Since previous studies [4,5] have shown that ferrite in weldments substantially reduces the fracture toughness at cryogenic temperatures, substantial work is being conducted on the development of 0% ferrite filler materials. A number of weldments made with experimental filler alloys have been tested at 4 K over the past year. The results of 275 these tests are presented later in this report by Elmer, McHenry and Whipple. Some of these experimental weldments have shown excellent combinations of strength and fracture toughness. However, some of the alloys have considerable problems with microfissuring and none are as yet commercially available. The study on the effects of the welding process on the cryogenic properties of 316L weldments, which was begun in FY79 [6], was completed during the past year. The results of 4 K mechanical property testing are presented later in this report by Whipple and Kotecki. The weldment made by gas tungsten arc welding (GTAW) was found to have the highest toughness followed by gas metal arc welding (GMAW) and submerged arc welding (SAW). This order of weld metal toughness is the same as the expected order of weld metal cleanliness, indicating that impurities in the weld metal may have a significant influence on cryogenic properties. As in past years, a number of qualification and test weldments were tested for tensile and fracture toughness at 4 K in support of the LCP magnet design and construction programs. These weldments included a 316L shielded metal arc (SMA) weldment, which was annealed, and 308L flux core metal arc (FCMA) weldments. The results of these tests are presented in Tables 1 and 2. In addition, these tables contain data obtained on a 31 6L electroslag (ES) weldment and a 308 SMA weldment (in the as-welded and annealed conditions), which were tested under separate contracts with the MFTF and the ANL-MHD projects, respectively. These data are included here because the weldments are being used as part of the overall study of 4 K fracture and deformation of stainless steel weldments at 4 K, which is included in this program. 276 As mentioned previously, a major fundamental study of fracture and deformation in duplex austenite/ferrite stainless steels at cryogenic temperatures has been undertaken. The objective of this investigation is to relate the structure of these materials to the cryogenic properties. During FY80 two initial parts of this investigation were completed and are presented later in this report. The first, by Brown, presents a detailed literature review, critical analysis and observations concerning structure development in duplex austenite/ferrite stainless steel weldments and castings. This study is necessary to define structural parameters to be used in deformation modeling and to understand effects that may arise due to such factors as microsegregation. This knowledge will also be necessary for the eventual modification of structural parameters to optimize mechanical properties. The second paper, by Whipple and Brown, presents the results of a metallographic investigation of deformation and fracture mechanisms in duplex austenite/ferrite stainless steel weldments and castings. In this study, the role of ferrite in 4 K deformation and fracture is examined with particular emphasis on possible fracture modes. The fracture-mode study has pointed out the advantages to be gained in studying castings as a model material that will give indications of the behavior of weldments. Because castings have a much coarser ferrite structure than weldments, their deformation and fracture phenomena are much more easily observed. Stainless Steel Castings The increased interest in the use of stainless steel castings for structural components of superconducting magnets has resulted in a program to evaluate castings for cryogenic service. Ferrite content and 277 nitrogen concentration have been found to have a significant influence on the cryogenic properties of stainless steel weldments [4,5] and it seems reasonable to expect that the same will be true of castings because of the similarities in their microstructures. Therefore, the program has been designed to test both of these effects. Two series of CF8M castings have been produced, one with a constant nitrogen content and varied ferrite content and one with a constant ferrite content and a varied nitrogen content. The production and initial testing of these castings is presented later in this report by Finch. In FY80, 4 K mechanical property tests were conducted on some CF8M and CF8 castings under separate contracts. The data from these tests are presented in Tables 3 and 4. These castings have been used as part of the deformation and fracture mechanism study being conducted in this program, the results of which are presented later in this report by Whipple and Brown. Aluminum Weldments The program on the evaluation of 5183 welds in 5083 aluminum, which was begun in FY79 [7], was completed this year with 4 K fracture toughness testing. It was found that the toughness of aluminum weldments is not sensitive to welding procedure or supplier. The results of this program are presented later in this report by Elmer and McHenry. Test-Method Development In FY79 work had begun on the evaluation of a simple test method for estimating J. [8]. The aim of that work was to develop a simple and inexpensive test method for weld qualification for 4 K service. The initial results indicated that the test worked well for the base metals tested. In the past year the test method was evaluated for use on weld- ments. The weldments tested were those from the welding process study, 278 mentioned previously. The results are shown graphically in figure 1, which is a plot of the J, estimated by this procedure versus the J, measured by the single-specimen J-integral procedure. From this figure it is evident that, in the case of weldments, the correlation between the estimated Jj and the actual Jj is very poor. On the basis of these results, it was determined that the test method is not useful for weld property evaluation. This is probably due to the large grain size of weldments compared with the small ligament of the specimens used for this test. Future Work The major portion of the research effort in FY81 will be devoted to the previously mentioned deformation and fracture study on duplex austenite/ ferrite stainless steel weldments and castings at 4 K. The objective of this work is to gain a much more comprehensive understanding of factors influencing the mechanical properties, including the causes of the observed large variability in weld metal properties. Specifically, this research will include the following efforts: 1. A theoretical model of duplex austenite/ferrite stainless steel deformation at 4 K will be developed from a two-phase deformation modeling approach. This model is intended to define quantitatively the effects of ferrite content, ferrite morphology, and phase properties on deformation. 2. The CF8M castings produced in FY80 will be used to verify and refine the above model and to make additional fracture and deformation mechanism observations. 279 3. Weldments of 31 6L stainless steel weldments will be made and tested with different welding processes and ferrite contents to extend the model to the finer structure in weldments. 4. Extensive metal lographic investigations will be conducted in support of the above program. Optical microscopy will be used to define and measure microstructural parameters. Transmission electron microscopy will be used extensively to investigate factors such as orientation relationships, interface structures and dislocation substructures or deformed and undeformed specimens. Scanning electron microscopy will continue to be used to correlate microstructure to fracture and deformation micromechanisms. 5. Microstrain tensile testing, using high-sensitivity strain measuring equipment, will be conducted to establish microdeformation phenomena further. In addition to and as part of the above program, a COD test methodology will be developed. This test method will use three-point bend specimens and will be capable of measuring J, . It is expected that this form of testing will significantly reduce the cost of 4 K fracture toughness testing. The 4 K mechanical property testing in support of superconducting magnet design and construction programs will continue in FY81 . 280 REFERENCES 1. H. I. McHenry, "Evaluation of Stainless Steel Weld Metals at Cryogenic Temperatures," in Materials Studies for Magnetic Fusion Energy Applications at Low Temperatures-I , NBSIR 78-884 (1978), pp. 159-166. 2. H. I. McHenry, D. T. Read, and P. A. Steinmeyer, "Evaluation of Stainless Steel Weld Metals at Cryogenic Temperatures," in Materials Studies for Magnetic Fusion Energy Applications at Low Temperatures U, NBSIR 79-1609 (1979), pp. 299-308. 3. H. I. McHenry and T. A. Whipple, "Weldments for Liquid Helium Service," in Materials Studies for Magnetic Fusion Energy Appli- cations at Low Temperatures-III , NBSIR 80-1627 (1980), pp. 155-165. 4. D. T. Read, H. I. McHenry, P. A. Steinmeyer, and R. D. Thomas, Jr., "Metallurgical Factors Affecting the Toughness of 316L SMA Weldments at Cryogenic Temperatures," Material Studies for Magnetic Fusion Energy Applications at Low Temperatures-II , NBSIR 79-1609 (1979) pp. 313-352. 5. E. R. Szumachowski and H. F. Reid, "Cryogenic Toughness of SMA Austenitic Stainless Steel Weld Metals: Part I - Role of Ferrite," Weld. J., Weld. Res. Suppl . 57, (1978) pp. 325-s-333-s. 6. D. J. Kotecki, "Weld Process Study for 31 6L Stainless Steel Weld Metal at Cryogenic Temperatures," in Materials Studies for Magnetic Fusion Energy Applications at Low Temperatures-III , NBSIR 80-1627 (1980), pp. 197-204. 7. R. A. Kelsey and L. N. Mueller, "Tensile and Notch-Tensile Properties at Room Temperature and 4K of Welds in Aluminum Alloy 5083-0 Plate," in Material Studies for Magnetic Fusion Energy Applications at Low Temperatures-III , NBSIR 80-1627, (1980) pp. 217-244. 281 8. T. A. Whipple and H. I. McHenry, "Note on a Simple Test Method for Estimating Jj c >" in Materials Studies for Magnetic Fusion Energy Applications at Low Temperatures-III , NBSIR 80-1627 (1980), pp. 245-256. : 282 Table 1. Tensile properties of qualification weldments tested at 4 K. Weld Type Ferrite, % Yield Ultimate Elongation, Reduction Strength, Strength, % in Area, MPa (ksi) MPa (ksi) % 636 939 13 11 685 1349 40 24 650 919 13 20 727 1315 46 43 674 (97.8) 1130 (164.0) 28 25 700 1436 34 23 680 1294 38 25 690 (100.0) 1365 (197.4) 36 24 483 1168 22 16 431 1395 29 19 460 1374 32 21 458 (66.4) 1312 (190.3) 28 19 541 1418 31 22 574 1219 21 16 558 (80.9) 1319 (191.3) 26 19 31 6L SMA |16L Electrosl ag 308L FCMA 0.60N 4.5 308L FCMA 0.8N 283 Table 2. Fracture properties of qualification weldments at 4 K. Weld Ferrite, % kj/m p (in* lb/ in ) K Ic (J), MPa/m (ksi/Tn) Fatigue Crack Growth C n 2 metric units (Engl ish units)' 31 6L SMA (annealed) 241 288 272 267 (1526) 247 (224) 0.501 x 10q 2.83 (0.263 x 10" y ) 31 6L Electroslag 4.5 174 140 157 ( 897) 189 (172) 308L FCMA 0.8N 80.3 77 80.3 79.2 ( 453) 134 (122) 0.489 x 10"? n 3.69 (0.281 x 10 U ) 308 SMA 12 46.0 32.2 39TT ( 224) 94.4 (85.9) 0.254 x 10" 3.84 (0.163 x 10 ) 3, 308 SMA Annealed 328 350 339 (1938) 278 (253) 0.196 x 10 q 3.78 (0.846 x 10" y ) 3.34 1. Constants fit the equation: da/dN = CAK 2. Metric units: AK in MPavfiTT da/dN in mm/cycle English units: AK in ksi/i!7, da/dN in in/cycle 284 Table 3. Tensile properties of stainless steel castings at 4 K. Material Ferrite, Y ield L)l timate Elongation, % Strength, Strength, % MPa (ksi) MPa (ksi) CF8 413 (60) 125 (181) 46.1 8 689 (100) 920 (133) 20.2 14.5 470 (68) 1512 (219) 47.6 12 578 (84) 753 (109) 10.7 CF8M 9.0 621 (90) 1317 (191) 41.0 469 (68) 1338 (194) 44.0 285 Table 4. Fracture properties of stainless-steel castings at 4 K. Material Ferrite, J IC 2 kj/m 2 ( i n • 1 b/ i n ) K Ic (J). Fatigue Crack Growth h MPa^T C 2 n metric units (ksi/in) CF8 481 (2750) 331 (301) 8 175 (1000) 199 (181) __ 14.5 69.3 (396) 125 (114) __ 12 31.5 (180) 84.5 (76.9) __ CF8M 9.0 302 403 280 372 339 (1938) 2.35 x 10' 12 5.21 278 (253) 1. Constants fit the equation: da/dN = CAK 2. Metric units; AK in MPai/rrTT da/dN in mm/cycle 286 i: \ 1 1 1 1 \ \ \ \ __ \ \ \ \ (0 t \ _ c 9 \ O — E \ r >el Weld < 1- < 5 < o < CO \ \ \ < V CO CD a> \ e D ^ 0) '5 i i \ c 1 0) i a> _C0 3 \ 75 CO CD CO OQ 1- 3 \ \ _i < O e D (0 \ *™ co 1 1 1 1 \ \ o o o o * CM E 1 — 1 ""3 \ T3 ■^ J* "3 E o ■t +-> o CO Q U1 LU -a 1- fO < o s ro E 1HM "O o H <4- r— O 0) CM 0) s O _l Hi I/) •r- l£> S- r— fO CO o a. E i- o "D o o O <4- o 0) =3 en o o ID O O o o CO o o CM o o ui/r* 'aaansvaifl °'r 287 STRENGTH AND TOUGHNESS OF FULLY AUSTENITIC STAINLESS STEEL FILLER METALS AT 4 K J. W. Elmer, H. I. McHenry, and T. A. Whipple Fracture and Deformation Division National Bureau of Standards Boulder, Colorado ABSTRACT Several experimental fully austenitic stainless steel welds have been prepared by filler metal suppliers in the United States, Austria, Sweden and the Soviet Union. The strength and toughness of these welds have been tested and are reported at 4 K. Some Charpy impact toughness and tensile data at temperatures other than 4 K are presented. These data were reported by the suppliers. Four of the eight welds that were tested had strength-toughness combinations comparable to that of aus- tenitic base metals. Of these four welds, three did not exhibit micro- fissuring that is generally present in fully austenitic welds. 289 INTRODUCTION Stainless steel weldments are used for the main structural members of the superconducting magnet systems for fusion energy devices. An- nealed austenitic stainless steels of the AISI 300 series are currently being used in weldments with thicknesses up to 125 mm. The weldments operate at high stress levels in liquid helium at 4 K. The weld metal that has been most thoroughly evaluated and is most frequently used for liquid helium service is AWS 31 6L deposited by the shielded metal -arc (SMA) process. Fracture toughness tests on 31 6L welds at 4 K indicate that toughness decreases with increasing ferrite content and with 2 increasing grain size in ferrite-free welds. In general, the toughness 3 of 31 6L welds varies significantly for the different welding processes, and the best assurance of satisfactory toughness is obtained with a ferrite-free weld deposited with a low-heat input. Problems arise 4 5 because ferrite-free welds are susceptible to microfissuring ' and low-heat inputs result in low deposition rates. The purpose of this program is to conduct a preliminary evaluation of the strength and toughness at 4 K of several proprietary fully aus- tenitic weld metals provided in the form of as-welded test plates by filler metal suppliers in the United States, Austria, Sweden, and the Soviet Union. The evaluation is limited to the specific welding process and procedures chosen by the suppliers to prepare the test plates. Promising grades will be identified, but further evaluation will be needed to assess their suitability for structural members of super- conducting magnets. 291 MATERIALS Test weldments were provided for evaluation by five filler metal suppliers. The chemistry of the weld metals is summarized in Table I. The welds are identified by their chemistry: the numbers indicate the Cr-Ni-Mn content and any additional characteristic alloying elements follow, e.g., alloy 18-20-6 Nb is 18% Cr, 20% Ni , 6% Mn, and contains Nb. The welding process and procedures used to prepare the test weldments were selected by the suppliers. The available information is summarized in Table II. Tensile and Charpy impact data provided by suppliers 4 and 5 are summarized in Table III. Microfissuring was observed on the fracture surface of all the welds except for the 20-16-7 Mo-W submerged arc weld (SAW), 17-16-4 Mo (SMA), and 18-20-6 Nb gas metal-arc (GMA) welds. PROCEDURES Tensile Testing Tensile testing was conducted on all-weld-metal specimens by means of a screw-driven testing machine equipped with a cryostat to perform the tests in liquid helium. A 0.635-cm (0.250-in)-diameter tensile bar was used, and the strain was measured over a 25-mm-gage length with a clip-gage extensometer mounted on the tensile specimen. A constant crosshead displacement rate of 0.008 mm/s was used for all specimens. Fracture Toughness Testing All fracture toughness testing was performed using the computer- aided, single-specimen, J-integral test facility described by Read. The testing was done in liquid He at a temperature of 4.0 K using a standard 2.54-cm (1 .00-in)-thick ASTM E399 compact tension specimen. The specimen was removed with the plane of the notch perpendicular to 292 the surface of the plate and centered along the fusion zone. In some instances, the weldment was not thick enough to remove a 2.54-cm-thick compact tension specimen. In these cases, the planar dimensions were kept the same as for a 2.54-cm-thick specimen and only the thickness was reduced. The specimens were precracked in tension-tension fatigue to a depth of 3.05 cm (1.20 in) prior to testing. Displacements were measured with a clip-gage extensometer mounted at the load line. RESULTS AND DISCUSSION The tensile properties of the fully austenitic welds at 4 K are summarized in Table IV. The yield strength ranged from 606 to 1110 MPa. Much of the variation in yield strength can be attributed to variations in the C-and-N contents. These interstitials are effective strengtheners at low temperatures. However, the correlation of yield strength with interstitial content reported by Tobler and Reed for AISI 304 base metal was not observed. Apparently, other factors such as grain size, oxygen content, and overall chemistry are important contributers to strength. For example, four welds with the same carbon plus nitrogen (0.08%) had yield strengths ranging from 606 to 814 MPa. Similarly, the weld ductility, as measured by reduction of area, generally decreased with increasing C plus N. For example, the three welds with less than 20% reduction of area had the highest C-plus-N contents. The fracture toughness values of the fully austenitic welds at 4 K are summarized in Table V. As a group, these welds had high fracture toughness relative to previous results on AWS 31 6L welds with and with- 1 2 out ferrite. ' A comparison of the present results with the trend line o of Read and Reed is shown in Figure 1, a plot of fracture toughness versus yield strength. Note that the strength and toughness combinations 293 for 20-16-7 Mo-W (SAW), 17-16-4 Mo (SMA) ? 18-20-6 Nb (GMA) and 16-16-4 Mo (FCAW) each are close to the base-metal trend line. This is noteworthy because four different welding processes are used, a wide range of strengths is covered, and three of the alloys (all except 16-16-4 Mo) did not exhibit microfissures . The similarity in chemistry of these alloys suggests that a fully austenitic alloy with 16 to 20% each of Cr and Ni, 4 to 1% Mn and refractory elements may provide the properties needed for liquid helium service. SUMMARY AND CONCLUSION A preliminary evaluation of the strength and toughness of several proprietary fully austenitic weld metals has been conducted. The re- sults for 20-16-7 Mo-W, 17-16-4 Mo, 18-20-6 Nb and 16-16-4 Mo indicate that strength and toughness combinations comparable to the properties of austenitic stainless steel base metals are possible in fully austenitic weld deposits. Further investigation is needed to identify the optimum chemistry and welding process for achieving microfissure-free welds that have strength and toughness properties needed for structural applica- tions in superconducting magnets. 294 REFERENCES 1. D. T. Read, H. I. McHenry, P. A. Steinmeyer, and R. D. Thomas, Jr., "Metallurgical Factors Affecting the Toughness of 316L SMA Weld- ments at Cryogenic Temperatures," Weld. J. Res. Suppl . , Vol. 59, No. 4, pp. 104s-113s (1980). 2. T. A. Whipple, H. I. McHenry, and D. T. Read, "Fracture Behavior of Ferrite Free Stainless Steel Welds in Liquid Helium," Materials Studies for Magnetic Fusion Energy Applications at Low Temper - atures-III , NBSIR 80-1627, National Bureau of Standards, Boulder, Colorado, pp. 167-194 (1980). 3. T. A. Whipple and D. J. Kotecki , "Weld Process Study for 31 6L Stainless Steel Weld Metal for Liquid Helium Service," this volume, pp. 303-321 (1981). 4. C. D. Lundin, W. T. Delong, and D. F. Spond, "Ferrite-Fissuring Relationship in Austenitic Stainless Steel Weld Metals," Weld. J. Res. Suppl., Vol. 54, No. 8, pp. 241s-246s (1975). 5. C. D. Lundin and D. F. Spond, "The Nature and Morphology of Fis- sures in Austenitic Stainless Steel Weld Metals," Weld. J., Vol. 55, No. 11, pp. 3565-3675 (1976). 6. D. T. Read, "The Computer-Aided J-Integral Test Facility at NBS," Materials Studies For Magnetic Fusion Energy Applications at Low Temperatures-III , NBSIR 80-1627, National Bureau of Standards, Boulder, Colorado pp. 205-216 (1980). 7. R. L. Tobler and R. P. Reed, "Interstitial Carbon and Nitrogen Effects on the Tensile and Fracture Parameters of AISI 304 Stain- less Steels," Materials Studies for Magnetic Fusion Energy Applica - tions at Low Temperature-III , NBSIR 80-1627, National Bureau of Standards, Boulder, Colorado pp. 17-48 (1980). 295 8. D. T. Read and R. P. Reed, "Fracture and Strength Properties of Selected Austenitic Stainless Steels at Cryogenic Temperatures," Materials Studies for Magnetic Fusion Energy Applications at Low Temperature-II , NBSIR 79-1609, National Bureau of Standards, Boulder, Colorado pp. 79-122 (1979). 296 * co -a i — CD 4-> CO CD I— 4- O s: CO CO >> c o 'r— E cu O (_3 cu fC cu CO co CD O o S- o •r— +J 03 C CD CO =3- r~» CO ■=!• r^> co — tn CO «=j- CO co CTi O co LO O o o LO LO LO ■"" •"" CM CM CM •— •"" r— LO CO o 1— CM i— r- r^ CTi r~- CO CO CO CO CO LO < s: o «=c crs C3 CD r^ *d" CO LO CO CD "=j- CO CO o CM o CM O CM CO CO o CM r-. 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Alloy Process 'Ic' lc' J/m' . 2 (in* lb/ in ) MPav^ (ksi/Tn) 1 20-16-7 Mo-N SAW AVERAGE 116 77 97 (660) 440) (550) 163 132 148 (148) 120) (148) 2 17-16-4 Mo SMA AVERAGE 240 306 273 (1370) 1748 (1560) 234 264 250 (213) (240 (227) 3 18-20-6 Nb GMA AVERAGE 130 350 240 (740) (2000) (1370) 171 282 227 (156) (257) (207) 3 18-20-6 Ti GMA 362 (2070) 287 (267) 4 18-16-9 GMA 65 (370) 121 (no) 5 16-16-4 Mo FCMA 238 (1362) 239 (217) 5 16-15-4 Mo GMA 114 (650) 170 (155) 301 1 1 ' 1 • o £ • 2 i 0) * o 1 - i 2 <0 - (0 h- T" ^ m i 1 1 (0 CO h» r- 1— T- » • • o •MM c CO 2 i o CM — • 4) +■* (0 3 (0 (0 C 4) • O z < D H 2 • (0 *■ o 2 1 o i CM ^» *■ 1 i i 00 c CU I— +-> rO CO +-> 3 CU to E >> cu c r — CO O t— fO T- 13 -O +J 4- ro o en s_ •!- ■■- O 4-> +-> M- t- -c cu +J +-> CD 00 c :3 CU rd CU S- +-> ^ to ■o i— CU rO Q. 1— CO 4-> E 3 •r- o (/> CO 3 S- CU -a > CU T3 Q CO rO CU j CO CU CL s- £ 4-> LU .c: o O CO CD Z3 CO CU CU >- (O sz CD CU rO 3 s- +-> O 3 rO +-> 4-> Q O CU fd 5- S- • 3 4- CO +-> -a a ^ i— ro CU S- ■=d- 3 M- i—" CU s- CD w/vedw '(r) 0, >i 'ssaNHonoi aumovud 302 WELD PROCESS STUDY FOR 31 6L STAINLESS STEEL WELD METAL FOR LIQUID HELIUM SERVICE T. A. Whipple Fracture and Deformation Division National Bureau of Standards Boulder, Colorado D. J. Kotecki Teledyne McKay York, Pennsylvania ABSTRACT This study was conducted to determine the effects of welding process choice on the cryogenic properties of 316L stainless steel welds. Six weldments were produced using two welding wires and three welding processes The weldments were impact tested down to 77 K and tensile and fracture toughness tested down to 4 K. The best properties obtained were from a GTA weld, followed by GMA welds; SA welds had the poorest properties. This variation in properties was attributed to the cleanliness of the weld metal, which is dependent on the welding process and parameters selected. 303 INTRODUCTION To date, most structures intended for liquid helium service have been welded with the shielded metal arc (SMA) welding process. This is mainly due to the lack of cryogenic mechanical property data on stainless steel weldments using higher deposition rate processes. The purpose of this program was to assess the effect of welding process choice and, to a lesser extent, the effect of welding parameter choice on the cryogenic toughness of type 31 6L weld metal at two nominal ferrite 1 evels--approximately 5 ferrite number (FN) and FN. This was intended to add to the data base on different welding processes and to assess the factors which cause differences between welding processes. MATERIALS Type 31 6L bare wire is fairly easy to obtain with ferrite calculated by the DeLong Diagram [1] at about 5 FN. But unless high-nitrogen content is acceptable, type 31 6L bare wire calculated to have FN is rare. From previous work [2,3], it is apparent that high-nitrogen content has a detrimental effect on cryogenic toughness. In order to begin with a low nitrogen wire of calculated FN, a tubular wire of special composition was produced for this work. For the weldments intended to contain 5 FN, a heat of 0.24 cm (3/32 in) diameter 31 6L stainless steel bare wire with the following composition was selected: C Mn Si Cr Ni Mo J? _S _N Co Cu 0.021 1.58 0.50 18.61 12.79 2.17 0.019 0.002 0.033 0.08 0.11 305 The ferrite number calculated from the mill analysis using the DeLong Diagram is 5.2 FN. A tubular wire was produced deliberately aiming to exceed the 11.0 to 14.0 percent nickel range for type 3]6L wire in AWS A 5.9, so that FN might be obtained. The aim composition of the wire was: _C _Mn _Si Cr Ni Mo _J 0.02 2.3 0.4 18.5 15.5 2.3 0.02 The calculated ferrite by the DeLong Diagram for this composition is 0.2 FN. Samples were taken from four locations within this lot. Gas tungsten arc (GTA) melts were prepared and checked for nickel, chromium, and molybdenum homogeneity, which proved to be satisfactory. PROCEDURES Welding Six weldments were produced for cryogenic testing. All welds were made in 2.54 cm (1.0 in) thick 304L stainless steel plate in the flat position. The basic joint preparation was a 90° included angle double vee which was symmetrical about the plate midthickness. Three weldments were made with each of the welding wires described above. The bare wire was used for a GTA weldment, a gas metal arc (GMA) weldment, and a submerged arc (SA) weldment. The tubular wire was used for a GMA weldment and two SA weldments, one with normal heat input and one with high-heat input. Details of the welding procedures, ferrite measurements, and fissure bend testing have been reported previously [4]. 306 Mechanical Testing The weldments were evaluated using tensile tests, Charpy V-Notch (CVN) impact tests and fracture toughness tests. The tensile tests were conducted at 295 K and 4 K; CVN tests were conducted at 77 K, 116 K, 144 K, and 172 K, and fracture toughness tests were done at 4 K. Tensile tests were conducted at 4 K using 2.5 cm (1 in) gage length and 0.64 cm (0.25 in) diameter specimens. The specimens were oriented along the length of the weldment so that the entire sample was welded metal. The test procedures and the cryostat for the low temperature tests have been described by Reed [5]. The 295 K tensile tests were done on standard 1.28 cm (0.505 in) diameter all weld metal specimens. CVN impact tests were conducted on 12 specimens from each weldment. The specimens were machined such that the notch was perpendicular to the plate surface and the crack propagated along the weld center! ine. Fracture toughness tests were conducted at 4 K using the ASTM E399 compact specimen with a 5.1 cm (2 in) width and a 2.5 cm (1 in) thickness. The specimens were oriented so that loading was perpendicular to the weldment and crack extension was along the weldment centerline. The single-specimen J-integral procedure [6] was used to determine J. , the value of J at the onset of crack extension. The plane strain fracture toughness, K, (J), was determined from J. using the relationship proposed by Begley and Landes [7]: 1/2 : IC (J) - Vi-v7 K i where E is Young's modulus and v is Poisson's ratio. Tests were conducted in an electrohydraulic fatigue machine equipped with the cryostat described by Fowl kes and Tobler [8]. 307 All tests were conducted using the computer aided J-integral testing facility described by Read [9]. RESULTS AND DISCUSSION Chemical composition and ferrite measurements were made on all weld metal deposits. The results of these measurements are given in table 1. Note that the SA weldments generally have higher concentrations of residual elements; of particular interest are silicon and oxygen. The higher concentration of these elements results from the flux used in SA welding. Also, note that the ferrite numbers do not correlate well with either the ferrite numbers calculated from the all-weld-metal deposit composition or the ferrite numbers calculated from the chemical composition of the welding wire. As a result of this unpredictability in ferrite contents, the original objective of obtaining three weldments with 5 FN and three with FN was not achieved. Ferrite measurements made on the actual double-vee weldments, which were tensile and fracture toughness tested at 4 K, are shown in table 2. These measurements were made with a ferrite meter that samples only a small area for ferrite content. (Note: These measurements were made in percent ferrite, which at these ferrite levels is approximately equal to FN.) By using this instrument, it is possible to see the variation in ferrite content through the thickness of the weldment. These results show that, although the average ferrite content of the double-vee weldment is wery close to that of the all weld-metal pad, there is a significant variation in ferrite content through the thickness of the weldment. Generally, higher ferrite contents are found in the center of the weldment, as would be expected from the higher degree of base metal 308 dilution in this area. However, two of the weldments deviated from this expected trend. The 31 6L bare wire GMA weldment had higher ferrite on one side of the double-vee than on the other side. The 316L tubular wire SA weldment with high heat input had higher ferrite near the plate surfaces than in the center of the weldment. In the case of the high heat input SA weldment this effect can be explained by the fact that this weldment had ^jery large passes, which resulted in considerable remelting of the previous pass and good mixing of the fused base metal. As a result of the primary austenite solidification mode, ferrite stabilizing elements were segregated to the top of the weldment. No explanation is offered for the curious result seen in the GMA weldment. The above results on ferrite distribution were qualitatively confirmed by metallographic examination. Figures la and lb show micrographs from locations C and E, respectively (see table 2). This sample was prepared using an electrolytic 10 N KOH etch, which selectively deposits a film on the ferrite phase. The weldments were also examined after color etching with a solution of 125 ml FLO, 50 ml HC1 , and 2 g Na 2 S0 4 , which was used to show whether the solidification mode was primary austenite or primary ferrite. The solidification mode is generally considered to be important with regard to microfissuring in the weldment [10] and it may also have an important effect on cryogenic toughness because it affects the continuity of the ferrite phase [11]. In general, the ex- amination with the color etch showed that the areas that contained greater than two percent ferrite solidified as primary ferrite and those areas with less than two percent ferrite solidified as primary austenite. The results of CVN impact tests are shown in table 3. From the yery low lateral expansion it is evident that the two normal heat input 309 SA weldments have a low toughness. Although CVN impact results at 77 K are not a good indication of 4 K fracture toughness [12], it was decided that the lateral expansions of these two SA weldments were so poor that 4 K testing was unwarranted. The 295 K and 4 K tensile properties of all weld metal specimens are shown in table 4. The 316L bare wire weldments had higher 4 K yield and ultimate strengths than the tubular wire weldments. This difference was large at 4 K and was probably due to the extremely low nitrogen content of the tubular wire weldments (see table 1). The results of 4 K fracture toughness tests are shown in table 5. The highest toughness was achieved with the 316L bare wire GTA weldment. The two GMA weldments had almost the same fracture toughness even though the bare wire GMA weldment had a significantly higher strength. The tubular wire high heat input SA weldment had poor fracture toughness as well as a low yield strength. As noted previously, the 4 K fracture toughness of stainless steel weldments has been correlated with both ferrite content and nitrogen concentration. The weldments tested in this study do not have signifi- cantly different ferrite contents, and the weldments with lower nitrogen concentrations also had lower fracture toughness, which is against the expected trend. The major difference in the weldments is the welding process used and it seems likely that this is the cause of the wide variation in fracture toughness. Of the three processes used, GTA, GMA, and SA, it would be expected that the GTA process would produce a "clean" weld and the SA process would result in a "dirty" weld. In this case, clean and dirty refer to the amount of nonmetallic inclusions in the weld metal. The SA welding 310 process uses a flux that should result in a higher inclusion count than processes that use gas shielding. Micrographs on unetched samples from the bare wire GTA and the tubular wire SA weldment are shown in figures 2a and 2b, respectively. This figure qualitatively demonstrates that the SA weldment had a much higher inclusion count. It is well known that a high concentration of hard second phase particles can significantly reduce fracture toughness, and this appears to be the cause of the observed variation in fracture toughness of the welds tested in this study. SUMMARY Six 316L stainless steel weldments, which were produced by three welding processes with two welding wires, were evaluated for cryogenic mechanical properties. The best combination of 4 K yield strength and fracture toughness was achieved with the GTA process using 31 6L bare wire. This weldment had both the highest strength and highest fracture toughness of the welds tested. SA welding was found to result in both low strength and low toughness, attributed to the significantly "dirtier" weld produced by this process. 311 REFERENCES 1. W. T. DeLong, "Ferrite in Austenitic Stainless Steel Weld Metal," Weld. J., Res. Suppl . 53, 273s-275s (1974). 2. E. R. Szumachowski and H. F. Reid, "Cryogenic Toughness of SMA Austenitic Stainless Steel Weld Metals: Part I: Role of Ferrite," Weld. J., Res. Suppl. 58, 325s-333s (1978). 3. E. R. Szumachowski and H. F. Reid, "Cryogenic Toughness of SMA Austenitic stainless Steel Weld Metals: Part II: Role of Nitrogen," Weld. J., Res. Suppl. 58, 34s-44s (1979). 4. D. J. Kotecki, "Weld Process Study for 316L Stainless Steel Weld Metals at Cryogenic Temperatures," in Materials Studies for Magnetic Fusion Energy Appl ications at Low Temperatures-III , NBSIR 80-1 627, National Bureau of Standards, Boulder, Colorado, pp. 195-204 (1980). 5. R. P. Reed, "A Cryostat for Tensile Tests in the Temperature Range 300 to 4 K," in Advances in Cryogenic Engineering , Vol. 7, Plenum Press, New York pp. 448-454 (1962). 6. G. A. Clarke, W. R. Andrews, P. C. Paris, and D. W. Schmidt, in Mechanics of Crack Growth , Proceedings of 1974 National Symposium on Fracture Mechanics, ASTM STP 590, American Society for Testing and Materials, Philadelphia, Pennsylvania, pp. 27-42 (1976). 7. J. A. Begley and J. D. Landes, in Fracture Toughness , Proceedings of 1971 National Symposium on Fracture Mechanics, Part II, ASTM STP 51, American Society for Testing and Materials, Philadelphia, Pennsylvania, pp. 1-20 (1972). 8. C. W. Fowl kes and R. L. Tobler, "Fracture Testing and Results for a Ti 6A1-4V Alloy at Liquid Helium Temperatures," Eng. Fract. Mech., 8, 487-500 (1976). 9. D. T. Read, "The Computer-aided J-integral Test Facility at NBS," i n Materials Studies for Magnetic Fusion Energy Appl ications at Low Temperatures-III , NBSIR 80-1627, National Bureau of Standards, Boulder Colorado, pp. 205-216 (1980). 10. T. Takalo, N. Suntala, and T. Moisio, "Austenitic Solidification Mode in Austenitic Stainless Steel Welds," Met. Trans. A, 10A , pp. 1173-1181 (1979). 11. T. A. Whipple and E. L. Brown, "Deformation and Fracture of Stainless Steel Weldments and Castings at Low Temperatures," in Materials Studies for Magnetic Fusion Energy Appl ications at Low Temperatures- J_V, NBSIR 81-1645, National Bureau of Standards, Boulder, Colorado pp. 415-451 (1981). 12. T. A. Whipple, H. I. McHenry, and D. T. Read, "Fracture Behavior of Ferrite- free Stainless Steel Welds in Liquid Helium," in Materials Studies for Magnetic Fusion Energy Appl ications at Low Temperatures- III, NBSIR 80-1627, National Bureau of Standards, Boulder, Colorado, pp. 167-194 (1980). 312 LIST OF TABLES Table 1. All -weld metal composition and ferrite content. Table 2 Delta- ferrite measurements on double-vee 31 6L stainless steel weldments. Table 3. CVN impact test results on 31 6L weldments. Table 4. Tensile properties of 316L stainless steel weldments. Table 5. Fracture toughness properties of 316L stainless steel weldments at 4 K. 313 Table 1 All -Weld-Metal Composition and Ferrite Element GTA 31 6L Bare 316L Tubul ar GMA SA GMA SA SA High-Heat Input C 0.017 0.017 0.023 0.014 0.016 0.017 Mn 1.39 1.40 1.97 2.17 2.50 2.43 Si 0.45 0.47 0.84 0.46 1.03 0.84 Cr 18.02 17.95 18.95 18.68 19.61 19.07 Ni 11.89 11.76 12.75 15.15 15.32 14.59 Mo 2.12 2.12 2.08 2.31 2.21 2.22 S 0.003 0.003 0.005 0.013 0.011 0.011 P 0.020 0.020 0.031 0.010 0.021 0.021 N 0.069 0.069 0.065 0.030 0.016 0.018 0.053 0.066 0.110 0.072 0.106 0.087 Co 0.092 0.090 0.095 0.012 0.025 0.020 V 0.063 0.063 0.06 0.013 0.017 0.016 Nb 0.02 0.02 0.02 0.02 0.02 0.02 Ti 0.021 0.022 0.024 0.15 0.092 0.104 Cu 0.089 0.086 0.10 0.037 0.043 0.042 FN measured 1.5 3.0 7.4 0.9 2.2 1.8 FN calculated by diagram 4.4 4.1 5.7 0.8 6.8 5.2 314 Table 2 Del ta-ferrite Measurements on Double-vee 316L Stainless Steel Weldments Wire Process A Percent Del ta-ferrite* C D E 'Measurements by ferrite meter at indicated locations. Avg, 316L GTA 0.0 0.6 1.9 0.7 0.0 0.6 Bare Wire GMA 3.5 3.0 2.8 0.6 0.0 2.0 316L GMA 0.5 3.2 2.3 3.2 0.6 2.0 Tubular Wire SA 4.7 0.2 0.7 0.6 2.7 1.8 (high-heat input) 315 Table 3 CVN Impact Test Results 1,2 on 316L Weldments Temperature , l< 316L Bare Wi GTA GMA re SA 316L Tubular GMA SA Wire SA High Heat Input 172 102 (0.91) 114 (1.17) 49 (0.41) 58 (0.91) 24 (0.30) 45 (0.53) 144 89 (0.99) 109 (1.19) 43 (0.41) 47 (0.69) 19 (0.18) 40 (0.53) 116 75 (0.69) 92 (0.99) 37 (0.28) 39 (0.43) 18 (0.18) 33 (0.36) 77 58 (0.46) 77 (0.81) 29 (0.05) 31 (0.43) 16 (0.08) 39 (0.38) 1 Results given are energy absorbed in joules and lateral expansion in (mm) "All values are the average of three tests. 315 Table 4 Tensile Properties of 316L Stainless Steel Weldments Yield Ultimate Reduction Temperature, Strength, Strength, Elongation, in Area, Process K MPa (ksi) MPa (ksi) % % Wire 316L Bare Wire GTA 316L Tubular Wire GMA SA GMA SA SA (high heat input) 295 430 (62.4 602 ) (87.3! 32 55 4 807 (117. o; 1549 ) (224.6! 34 27 295 355 (52.0 554 I (80. 3; 45 55 4 779 (113. o; 1510 ) (219.0; 37 27 295 441 (63.9] 587 ) (85. i; 36 54 295 332 (48. i; 503 ) (73.0; 30 34 4 608 (88.2; 1162 ) (168.6; 29 20 295 387 (56. I! 542 » (78.6; 32 41 295 322 (46.7] 555 1 (80. 5; 34 38 4 606 (87.9] 1122 1 (162. 7; 21 18 317 Table 5 Fracture Toughness Properties of 316L Stainless Steel Weldments at 4 K J lc, K Ic< J > kJ/m^ MPa/nT Wire Process (in«lb/in ) (ksi/Tn) 316L Bare GTA 138 Wire 148 143 (818) 181 (165) GMA 103 130 116 (666) 163 (148) 316L Tubular GMA 111 (635) 159 (145) Wire SA (High 47 (270) 103 (94) heat input) 318 LIST OF FIGURES Figure 1 Micrographs of locations C and E on the 316L bare-wire GTA weldment, showing the variation in ferrite content through the thickness of the weldment. KOH etch, 400X. Figure 2 Micrographs showing the difference in the cleanliness of the weldments (a) 316L bare-wire GTA weldment (b) 316L tubular-wire SA weldment. Unetched, 200 X. 319 J- " ■ \ >. , V-** . » _# % v { 4% * , *■ x :>fi< ■ a* *■ ■" **•>" s\~ (a) % t *%•«. ■— I (b) Figure 1. Micrographs of locations C and E on the 31 6L bare wire GTA weldment, showing the variation in ferrite content through the thickness of the weldment. KOH etch, 400X. 320 » _ » fc^^ * r » r (a) r • i ' * * ?* #* ;.« i -. • ■ • • * ' * i ' **- (b) Figure 2 Micrographs showing the difference in the cleanliness of the weldments (a) 316L bare wire GTA weldment (b) 316L tubular wire SA weldment. Unetched, 200X. 321 FRACTURE TOUGHNESS PROPERTIES OF WELDS IN ALUMINUM ALLOY 5083-0 AT 4 K J. W. Elmer and H. I. McHenry Fracture and Deformation Division National Bureau of Standards Boulder, Colorado ABSTRACT A series of weldments was prepared by five cooperating companies in 51-mm thick 5083-0 aluminum plate using 5183 filler metal and the gas metal arc (GMA) welding processes. The J-integral fracture toughness and fatigue crack growth rates of the weldments were measured at 4 K and found to be similar to base metal properties. In a related study, the tensile and notched tensile properties of the same weldments were ob- tained at room temperature and 4 K. Data for each of the mechanical properties did not vary significantly from weld to weld. It is con- cluded that the mechanical properties of 5183 welds at 4 K do not depend on the GMA welding parameters and procedures. 323 INTRODUCTION Aluminum alloy 5083-0 is a solid-solution-strengthened alloy fre- quently used for cryogenic storage tanks for service to 76 K. Alloy 5083-0 derives its limited strength from the solid-solution-strengthening contributions of 4.4% Mg and 0.7% Mn. It is readily weldable with alloy 5183 filler metal and has outstanding toughness at low temperatures. New applications for 5083-0 aluminum in magnetic fusion energy systems are being examined; these place additional requirements on the alloy since it is necessary to retain favorable mechanical properties down to 4 K. To date, few mechanical property data are available for 5083-0 weldments at 4 K.^ 1 ' 2 -' In this study, fracture toughness and fatigue crack growth rate data were obtained for 5083-0 weldments at 4 K. This study was preceded by an evaluation of the tensile and notched tensile properties of the m same weldments at room temperature and 4 K by Kelsey and Mueller. The objectives were as follows: 1. Develop fracture toughness and fatigue crack growth rate data for 5083 aluminum weldments at 4 K. 2. Compare the 4-K toughness of 5183 aluminum weldments produced by different companies using different welding parameters and processes. WELDING PROCEDURES Weldments were prepared and donated by five different companies. Each company was asked to weld 51 -mm (2-in) thick 5083-0 aluminum plate using 5183 filler wire and the gas metal arc (GMA) welding process. All welding was performed in the flat position using a double-vee joint 325 preparation; however, specific welding parameters were the choice of the individual company. In addition, two companies supplied high deposition rate (HDR) GMA weldments of the same material to determine if high heat input influences the mechanical properties. The welding parameters used are listed in Table 1. The welding parameters varied considerably; for example, the number of passes varied from 2 for an HDR weld to 55 for one of the conventional welds. A total of seven welds were made, five by the conventional GMA welding process and two by the HDR-GMA welding process . TESTING PROCEDURE All fracture toughness testing was performed in liquid helium at 4 K using a standard 2.54-cm (1-in) thick ASTM E399 compact tension specimen. The specimen was removed from the center of the 51 -mm welded plate with the plane of the notch perpendicular to the plate and cen- tered in the fusion zone, such that the 5183 weld metal was tested. The specimens were precracked in tension-tension fatigue to a depth of 3.05 cm (1.20 in) prior to testing. The fracture toughness was measured using a computer-aided, single-specimen J-integral test, described by T3l Read. J Fatigue crack growth rate measurements were made using the same specimen configuration and were tested in liquid helium with an R ratio of 0.1 . RESULTS Tensile properties of the 5083 weldments at room temperature (Table 2) and at 4 K (Table 3) are provided to include additional in- formation about the weldments. 1 - - 1 Table 3 also shows the 4-K fracture toughness properties for each of the seven weldments as measured by the 326 J-integral technique. In Table 3, K Ic (J) refers to the plane strain fracture toughness as estimated from the J-integral test results. Fatigue crack growth rate data at 4 K (Table 4) compares conventional and HDR weld metal properties with 5083-0 base metal properties measured [41 by Tobler and Reed. L J The empirical constants n and C refer to the Paris equation of fatigue crack growth rate: ^=C(AK)" C has metric units corresponding to da/dN in mm/cycle and aK in MPav^. DISCUSSION A comprehensive summary of the tensile data indicates that welding parameters do not influence the yield strength or tensile strength to any significant extent. Except for the generally lower tensile prop- erties of the GMA weldment of supplier 2, the tensile properties of the GMA weldments were quite consistent. The HDR weldments proved to be only slightly lower in strength than the conventional weldments. The notch yield ratios at 4 K and room temperature are all greater than 1 .6, which suggests that the weld metal is very tough. Results of the fracture toughness testing at 4 K show the average fracture toughness, K, (J), determined by the J-integral method, to be 45.1 MPa/m. The K T (J) data are consistent for the seven weldments and Ic all the data lie within one standard deviation of the mean except for that of the HDR weldment from supplier 4. The tensile data are also consistent indicating that the mechanical properties of weldments in 5083-0 aluminum are relatively insensitive to welding parameters. 327 The average toughness of 45.1 MPa^m at 4 K is low when compared with those of other tough materials used at low temperatures and seems to be in conflict with the notch-yield-ratio results. Actually the K, (J) values for the welds are the same as the 5083-0 base plate (TL T3l orientation) results at 4 K reported by Tobler and Read. J The low values are due to the testing technique used and do not reflect the true fracture resistance of the material. The problem with measuring the toughness of the 5183 aluminum weldments using the J-integral technique is that the aluminum tears at low J-levels where K T (J) is measured and Ic thus the K T (J) value is low. However, 5183 aluminum weldments are able Ic to sustain much higher J-levels under stable crack propagation condi- tions, which is an indication of toughness. This is shown in Figure 1, which is a load-displacement curve recorded during one of the tests. The point indicated is the point of initial tearing, determined from the J-Aa curve. This occurs at a low load, but the specimen was able to sustain loads past the peak load without fracturing. Thus, 5183 alu- minum weldments are tough even though their K, (J) values are low. Fatigue crack growth rates for 5183 weld metal deposited by con- ventional and HDR-GMA processes are compared with the fatigue crack growth rate of 5083-0 base metal at 4 K in Figure 2. The fatigue crack growth rates for the conventional and HDR welds are similar and gener- ally faster than the base metal rates; however, without further study the differences between them cannot be considered significant. SUMMARY AND CONCLUSION Tensile properties at room temperature and at 4 K for 5183 aluminum weld metal were shown to be consistent for welds made with large varia- tions in welding parameters. - 1 The K, (J) values and fatigue crack 328 growth rates at 4 K for the welds were also shown to be insensitive to welding process and parameters. Therefore, it can be concluded that the mechanical properties of 5183 weld metal do not depend upon the GMA welding parameters to any great extent. 329 REFERENCES R. A. Kelsey and L. N. Mueller, "Tensile and Notch-Tensile Proper- ties at Room Temperature and 4 K of Welds in Aluminum Alloy 5083-0 Plate," in Material Studies for Magnetic Fusion Energy Applica - tions at Low Temperatures-III , NBSIR 80-1627, 219-244 (1980).* F. G. Nelson, J. G. Kaufman, and E. T. Wanderer, "Tensile Proper- ties and Notch Toughness of Groove Welds in Wrought and Cast Alu- minum Alloys at Cryogenic Temperatures," Advances in Cryogenic Engineering, Vol. 14, 1969, pp. 71-82. D. T. Read, "The Computer-Aided J-Integral Test Facility at NBS," in Materials Studies for Magnetic Fusion Energy Applications at Low Temperatures-III , NBSIR 80-1627, 205-216 (1980). R. L. Tobler and R. P. Reed, "Fracture Mechanics Parameters for a 5083-0 Aluminum Alloy at Low Temperatures," Journal of Engineering Materials and Technology , Vol. 99, (October 1977). *N0TE: A slightly revised version of Reference 1 (same authors and title) has been published: Report Number 57-80-02, Alcoa Laboratories, Alcoa Center, PA 15069, March 14, 1980. 330 Table 1 Welding Procedures [1 ] c ,. Process Current Voltage Gas Su PP lier Weld (A) (V) (L/min) ° f Number of Passes GMA 230-250 26-29 15.5 Ar 31.0 He GMA 250-300 26-29 11.7 Ar 35.2 He GMA 240-300 24-26 28.2 Ar GMA 335-365 33-34 9.4 Ar 18.8 He GMA 380-390 28 — HDR* 580 34 70.5 Ar 47.0 He HDR 740-750 33 11.7 Ar 35.2 He 55 3 GMA 240-300 24-26 28.2 Ar 25 4 GMA 335-365 33-34 9.4 Ar 16 r High Deposition Rate 331 Table 2 Room-Temperature Mechanical Properties of 5083 Welds [1] Suppl ier Process 0.2% Yield Strength (MPa) Notch Tensile Strength (MPa) Notch* Yield Ratio 1 GMA 166 312 1 .88 2 GMA 148 315 2.13 3 GMA 166 314 1 .89 4 GMA 167 327 1.96 5 GMA 178 356 2.00 4 HDR 152 312 2.06 5 HDR 146 306 2.09 *NYR _ Tensile Yield Str ength Notch Tensile Strength Table 3 Mechanical Properties of 5083 Welds at 4 K [1] Suppl ier Process 0.2% Yield Strength (MPa) Notch Tensile Strength (MPa) Notch Yield Ratio K Ic (J) (MPa^n) 1 GMA 202 363 1 .80 41 .7 2 GMA 177 344 1 .94 46.9 3 GMA 213 347 1 .63 41 .3 4 GMA 200 362 1.81 42.8 5 GMA 214 391 1 .83 49.0 4 HDR 202 343 1.69 51 .9 5 HDR 200 337 1 .69 42.0 332 Table 4 Fatigue Crack Growth Rate at 4 K AK Material Condition n C Range MPa/m 5083-0 annealed 5.2 1.3 xlO" 10 10-30 5183 HDR weld 6.5 3.8 xlO" 12 15-31 5183 GMA weld 4.6 7.3 x 10" 10 14-32 333 E E n* 'avon QJ o aj o ai o en s- •i— o -(-> M- u ' a> •1— > -a s- c 3 »l— u • CO CD 4-> -a c: sr i — •i — cu a s- f= S rd cu CU o i — 4-> fO •S -a > fO c: o o 4- _j o o CD S- CD TENSILE AXIS 50 ym Figure 1 Micrograph of a 31 6L SA weld metal tensile specimen deformed at room temperature, taken near the fracture surface and showing the elongation of the ferrite. KOH etch. 435 , THE INFLUENCE OF FERRITE AND NITROGEN ON MECHANICAL PROPERTIES OF CF8M CAST STAINLESS STEEL Leroy Finch ESCO Corporation Portland, Oregon ABSTRACT The production and mechanical -property testing of a cast molybdenum stainless steel are described. 337 THE INFLUENCE OF VARIABLE FERRITE AT A CONSTANT NITROGEN LEVEL AND OF VARIABLE NITROGEN AT A CONSTANT FERRITE LEVEL ON MECHANICAL PROPERTIES OF CF8M CAST STAINLESS STEEL INTRODUCTION The National Bureau of Standards in Boulder desired to study the ef- fect of a) variable ferrite content at a constant nitrogen level and b) a variable nitrogen content at a constant ferrite level in cast molyb- denum stainless steel, CF8M. These materials were to be tested in tension at room temperature and 77°K (liquid nitrogen) at ESCO and at 4°K (liquid helium) at NBS in Boulder. Impact resistance values were also to be determined at room temperature and 77°K at ESCO and at 4°K at NBS. Tables 1 and 2 note the aim chemical analysis of the ten compositions decided on, five with variable ferrite content from 0%, fully austentic, to 28% with a nitrogen content of 0.05% and five with a calculated fer- rite content of 10% with variable nitrogen contents from 0.02 to 0.20%. The ferrite contents were calculated using the data and curve reported at the recent ASTM symposium on cast stainless steels, "Ferrite Measure- ment and Control in Cast Duplex Stainless Steels" by L. S. Aubrey et al. Figure 1 is the curve used to estimate the ferrite content from the chemical compositions. Pattern X3829 was used to make the molds for the casting to be used by NBS for the tests to be made at 4°K. The final casting is 4"kl2"x24". The test bars tested at ESCO were from a conventional keel bar which has a diameter of 1**" at the bottom. PRODUCTION Ten one thousand pound heats were melted in an induction furnace with a monolithic magnesia lining and powered by a motor generator having a 960 Hertz output. The sand mix used to form the molds was 33% zirconium silicate, 66% silica bonded with bentonite. This was washed with an alcohol base silica wash and torch dried. When the metal was all melted a preliminary sample was sent to the laboratory for spectrographic analysis. A "suck-up* sample was also sent for anlaysis of nitrogen content. When the results of these analyses were received back at the melting location, additions were calculated to bring the composition to the de- sired level. These additions were carefully weighed out and added to the molten metal. When the proper temperature was reached as measured with a platinum rhodium thermocouple, the metal was tapped into a magnesia lined ladle and transported to the mold. Final chemical analysis samples were taken from the ladles full of metal. When the proper pouring temper- ature was reached the metal was ooured into the mold. 339 After the metal had cooled down the castings were removed from the molds, heated to 2000 f, water quenched and the risers burned off using an oxyacetylane-powder torch. After the risers were removed the castings were again heated to 2000°F and water quenched. This double heat treat- ment was used because it has been found that CF8M is susceptible to sigma formation during cooling in the mold and if the risers are cut from'mate- rial with sigma present in the microstructure cracking is likely to occur in the areas subjected to thermal shock. By heating to 2000 F the sigma is converted to ferrite and this ferrite is retained by the rapid water quench. After the risers are removed the material is again heat treated to redissolve any carbide which may have precipitated during the removal of the riser and also to dissolve any sigma that may have formed. Conventional tension test bars were machined from keel bars poured with metal from each heat. These were heat treated with the castings. Charpy V notch impact test specimens were also machined from the keel bars. The ferrite content of each heat was calculated using the Schoefer ratio and curve, Figure 1. It was also measured with a calibrated Magna gage and a Ferritescope. A portion of the end of each of the tension test bars was polished, etched and a point count made of the ferrite present using the procedure given in ASTM-E 562-76. The spectrographic chemical analyses of the heats produced are listed in Tables 3 and 4. The nitrogen contents included in the tables were obtained by vacuum fusion using a Leco rapid nitrogen analyzer. Also included in the tables are the ferrite contents determined by the methods noted above. The heat which was designed to have 0.02% nitrogen was melted in an argon atmosphere in order that the metal would not absorb nitrogen from the air and finish with a nitrogen content higher than 0.02%. The ferrite contents of four of the original heats were lower than desired. They were not used. The nitrogen content of another was higher than de- sired so it also was not used. Five more heats were produced in a similar manner as described above. In order to insure that the ferrite contents of these were satisfactory before the molten metal was tapped the spectrographic sample poured into the copper chill mold was measured with the calibrated Ferritescope. When the ferrite content was that desired the metal was tapped and poured as before. 340 MECHANICAL PROPERTIES Table 5 notes the room temperature mechanical properties obtained at ESCO for the ten heats which will be tested by the NBS. Table 6 lists the impact properties obtained to date. The tensile properties of the heats have not been obtained at 77 K at this time nor have the impact tests been completed for the last four heats. They will be reported when obtained. Figure 2 is a plot of the tensile and yield strengths (at 0.2% offset) at room temperature of the variable ferrite content heats versus the ferrite content. The ferrite content is expressed as the Cr /Ni ratio. Figure 3 is the plot of the tensile and yield strengths (at 0.2% offset) at room temperature of the variable nitrogen content heats versus the nitrogen contents. In Figure 3 the numbers above each point plotted show the ferrite con- tent of the material expressed by its ratio. The variable nitrogen heats did not have exactly the same ferrite contents. Heat 48987 was not used in these plots since the ferrite was 14% while the others were 8-10%. Heat 48988 with 8% ferrite and 0.05% nitrogen was used. Figure 4 is a plot of the average impact strength at room temperature and 77 K versus the ferrite content. Figure 4 is a similar plot of impact strength versus nitrogen contents. DISCUSSION From Figure 2 and Table 5 it can be said that the ultimate tensile strength increases 940 psi per 1% ferrite when the ferrite varies from 0%, fully austentic, to 29%. The yield strength increases 890 psi per 1% ferrite over the same range of ferrite. From Figure 3 and Table 5 it can be said that increasing the nitrogen content from 0.02 to 0.05% has little or no effect on mechanical proper- ties. When the nitrogen content increases from 0.05 to 0.14% the tensile strength increases 855 psi per 0.01% of nitrogen. When it increases from 0.14 to 0.20% the increase in tensile strength averages 1250 psi per 0.01% of nitrogen. From Figure 2 and Table 6 it will be noted that the increase in tensi le •strength between 8 and 10% ferrite, 1.14 and 1.19 ratio is 2100 psi so if 2100 is subtracted from 93.7 the tensile strength of the heat with 1.19 ratio and 0.20% nitrogen the tensile strength of the material if it had had 8% ferrite and 0.20% nitrogen could have been 91.6 ksi, so 91.6 - 84.1 = 7500 psi, which could have been due to the increase in nitrogen content or 1250 psi per point of nitrogen. 341 It will be noted with similar calculation that the tensile strength of the heat with 0.10% nitrogen would have fallen on the line if its fer- rite content had been 8%, 1.14 ratio, instead of the 10%, 1.17 ratio. The effect of the nitrogen content increase on yield strength appears to be linear, ,1160 psi per 0101% nitrogen. It will be noted in Figure 4 that the impact strength decreases as ferrite content increases. This is true at room temperature as well as at 77 K. Figure 5 shows that the impact strength at room temperature and 77 K is also lowered by increasing the nitrogen content. Of course this would be expected with the tensile strengths increased. When the mechanical property data is completed the results will be reported. 342 7 14 21 28 0.06 0.06 0.06 0.06 0.06 1.20 0.40 0.40 0.40 0.40 0.50 1.25 1.25 1.25 1.25 18.25 18.35 18.65 19.20 20.50 13.00 10.75 9.05 8.20 8.20 2.15 2.15 2.15 2.15 2.15 0.05 0.05 0.05 0.05 0.05 Table 1 - Aim Chemical Analysis in Percent of Five Heats of CF8M with Variable Ferrite and Constant Carbon and Nitrogen Contents Ferrite % C Mn Si Cr Ni Mo N 2 CrVNi 1 0.90 1.12 1.27 1.39 1.485 Table 2 - Aim Chemical Analysis in Percent of Five Heats of CF8M with Variable Nitrogen Contents at a Constant Ferrite Level, 10% C Mn Si Cr Ni Mo N 2 Cre'/Nie' 1.19 1.19 1.19 1.19 1.19 0.06 0.06 0.06 0.06 0.06 0.40 0.40 0.40 0.40 0.40 0.75 1.25 1.25 1.25 1.50 18.55 18.90 19.25 21.20 21.00 10.10 10.00 9.25 9.45 9.25 2.15 2.15 2.15 2.15 2.85 0.02 0.06 0.10 0.15 0.20 343 Table 3 - Final Spectrographic Analysis in Percent of Five Heats of CF8M with Variable Ferrite Content and a Constant Carbon and Nitrogen N Heat No. C Mn Si Cr Ni Mo 2 Ferrite Content * C M Pt Ct 47640 0.06 1.10 0.59 18.04 13.24 2.12 .05 71T8 48988 0.06 0.38 1.16 19.62 10.61 2,12 .05 1.21 6 8 47521 0.06 0.36 1.13 19.23 8.25 2.16 .05 1.39 17 15 47552 0.06 0.35 0.95 20.40 8.22 2.13 .05 1.45 22 22 49021 0.05 0.39 1.11 22.54 9.84 2.26 .05 1.48 27 29 Table 4 - Final Spectrographic Analysis in Percent of Five Heats with Variable Ni trogen at a Constant Ferrite Content Heat No. C 0.07 Mn 0. 39 Si 1.09 Cr Ni Mo 2.07 .02 Ferrite C M 1.25 8 Content* 49020 19.32 10.00 Pt Ct 8 48987 0.07 0.51 1.17 20.13 10.23 2.49 .06 1.27 11 14 47364 0.06 0.40 1.30 19.17 9.14 2.00 .10 1.18 8 9 47390 0.07 0.41 0.98 21.00 9.46 2.13 .144 1.16 7 8 47399 0.06 0. 33 1. 35 21.66 9.68 2.83 .20 1.19 9 10 *Ferrite Content C - Calculated using Schoefer formula for ratio Cr =Cr+l. 5Si+l. 4Mo-4.99 Ni 1 =30xC+. 5Mn+Ni+2 6(N-0 2)+2. 7 7 M - Measured with calibrated Ferritescope Pt Ct - Metallographic point count, ASTM E 562 The Cre'/Nie' ratios used in the plotting of the figures were obtained from the ferrite content determined by point count and taken from Figure 1. 344 Table 5 - Room Temperature Mechanical Properties of Heats Shipped to NBS Tested at ESCO UTS YS ksi E RA Heat No. ksi (0.20% offset ksi ) Variable Ferrite, 0.05% N 47640 65.3 29.6 48988 76.4 36.4 47521 87.2 48.5 47552 89.0 50.2 49021 92.6 55.4 Variable Nitrogen Content, Constant Ferrite 49020 76.4 36.2 48987 82.9 41.0 47364 81.3 44.3 47390 84.1 47.6 47399 93.7 53.8 Table 6 - Impact Properties Obtained to date % % Ferri Ratio te % 55. 5 65 .9 56 71 1. 14 8 48 64 1. 30 15 46 5 63 1. 43 22 42 64 1. 53 29 58. 65 1. 14 8 47. 69 1. 28 14 48. 5 70 1. 17 9 51 66 1. 14 8 42 59 1 19 10 Impact Data, Ft-Lbs ,o. 156-140-151 149 74- 78- 76 76 149-122-145 139 63- 66- 61 63 Heat No. R.T. Avg . 77 K Avg . 47640 177.-240-240 219 136-122-147 135 48988 47521 47552 49021 49020 48987 47364 167-232-223 47390 170-175-169 47399 190-175-172 207 66-101-110 92 171 65- 68- 61 65 179 38- 39- 22 33 345 ADDENDUM The Influence of Variable Ferrite at a Constant Nitrogen Level and of Variable Nitrogen at a Constant Ferrite Level on Mechanical Properties of CF8M Cast Stainless Steel This is an addendum to the original report which did not contain all of the data. This report includes the missing data: impact test results at room temperature and 77 K and tensile properties at 77 K. Table 6 reports the impact strength at room temperature and 77°K obtained on all of the alloys. Some of these were included in the original report. Figure 4 is a plot of the impact strengths of the heats with variable ferrite content vs. the ferrite content at room temperature and 77°K. Figure 5 is a plot of impact strength vs. the nitrogen contents for the variable nitrogen series at room temperature and 77 K. The impact strength drops with increase in ferrite content at both room temperature and at 77 K. At room temperature the loss per percent of ferrite is as follows: 0- 8% 3.9 ft-lbs 0-15% 4.7 ft-lbs 0-22% 3.6 ft-lbs 0-29% 3.4 ft-lbs At 77°K the drop is: 0- 8% 3.9 ft-lbs 0-15% 3.9 ft-lbs 0-22% 3.3 ft-lbs 0-29% 3.1 ft-lbs The rate of loss is nearly equal at both temperatures. Table 7 reports the tensile properties obtained at 77 K. 346 Figure 6 is the plot of the tensile strengths at 77°K plotted against the ferrite content expressed as Cr /Ni ratio and the nitrogen contents. It will be noted that the tensile strength increases both with ferrite content and with nitrogen contents to maxima and drops off with further increases: 22% ferrite with a ratio of 1.43 has the greatest strength of the ferrite series and the nitrogen content of 0.10 percent was the strongest of the nitrogen series of heats. LFinch /nmo 2/25/81 347 T able 5 - Room Temperature Mechanical Properties of Heats Shipped to NBS Tested at ESCO Heat No. UTS ksi (0 20% YS ksi offset ksi) E % RA % Ferrite Variable Ferrite, 0.05% N Ratio _%_ 47640 65.3 29.6 55.5 65 .9 48988 76.4 36.4 56.0 71 1.14 8 47521 87.2 48.5 48.0 64 1.30 15 47552 89.0 50.2 46.5 63 1.43 22 49021 92.6 55.4 42.0 64 1.53 29 Variable Nit: rogen Content, Constant Ferrite 58.0 65 1.14 49020 76.4 36.2 8 48987 82.9 41.0 47.0 69 1.28 14 47364 81. 3 44. 3 48.5 70 1.17 9 47390 84. 1 47.6 51.0 66 1.14 8 47399 93.7 53.8 42.0 59 1.19 10 Table 6 ■ Heat No. 47640 48988 47521 47552 49021 49020 48987 47364 47390 47399 Impact Properties Impact Data, Ft-Lbs R.T. Avg . 177-240-240 219 203-162-199 188 156-140-151 149 149-122-145 139 125-118-115 119. 3 162-137-139 146 184-198-205 195. 7 167-232-223 207 170-175-169 171 190-175-172 179 7 7°K 136-122- 147 98-108- 107 74- 78- 76 63- 66- 61 40- 45- 47 73- 87- 85 118-102- 93 66-101- 110 65- 68- 61 38- 39- 22 Avg. 135 104.3 76 63 44 81.7 104.3 92 65 33 348 Table 7 - Mechanical Properties at 77 K of Heats Shipped to NBS Tested at ESCO Heat No. UTS, psi E, % Variable Ferrite Content, 0.05% N 47640 142,300 38 48988 190,800 53 47521 211,200 30 47552 214,100 37 49021 169,200 30 Variable Nitrogen Content, Constant Ferrite 49020 183,100 35 48988 190,800 53 48987 186,000 53 47364 202,700 35 47390 173,300 39 47399 170,500 25 LFinch /nmo 1/23/81 Ferrite RA, % Ratio _%_ 47 .9 54 1.14 8 20 1.30 15 28 • 1.43 22 56 1.53 29 !k 55 1.14 8 .02 54 1.14 8 .05 61 1.28 14 .06 25 1.17 8 .10 52 1.14 8 .144 36 1.19 10 .20 349 Cr 1 = %CR + 1.5%Si = 1 .4%Mo - 4.99 Ni 1 = 30(%C) + ^ + %Ni + 26(%N - .02) + 2.77 o FERRITE, % Figure 1 Relationship between Cr /Ni and percent Ferrite in Cast Duplex Corrosion Resistant Stainless Steels 350 •s>| 'H10N3H1S CH3IA O a 9 Ifl "* CO CO I 1 1 1 1 1 • -C r- CD cu r»» T" i- CM z + ^» CD CM • k. >- • O -o OJ C 2 CD IT) CM O) c 0) O) bet Cont r~ c 0) LU Q. •r- O) + O CO CO K onsh rrit <*S CO "35 "O q QC •r- CD +-> Ll_ fl3 II II c 0) r- QC 1— -a c •r- z 1 < 1 1 00 ■ LU LL Figure 2 R a c 0) 00 1^ « > !S>| 'H19N3U1S 31ISN31 !51 !S>| 'HlONaaiS (H3IA Z LU h- Z o o z LU o o DC +-> en oo T3 'a; >- -u c ra a> •r- to a; E -t-> +-> c: CD o -Q C_J Q. E •r- a> -E C7i to O c s- o +-> •^ • ^ +■> ^ (O r— T3 O) C a: 03 s- cn !S>1 'H10N3H1S 31ISN31 * LL 'sq|-u *H±0N3U±S lOVdIAII o 00 o o o o <0 o u~> r-- C\J + C\J o C\J C\J o 1 1 1 1 1 1 4> 1 1 1 y im S *-* •/ (0 Q) a E 0) E o * / ° /% *™ o ■"■ o • o/ 1 1 1 <0 _*■ 1- _CM "O o- fD •<- ^ 1— \ ■■■ Q&- z \ to ™ -C ro 4-> o CD 00 S- LO *\ 4-> ID Z CO S- Q- +-> X O O) LU Q-+-> E c 1- i — i o CD -M -O -i- LU •r- CD -C Li. H CO ■^ O C OC •r- 03 4-> oc nD ^ LU U. Sj" o ■* fC H Z CD c III CD H Z +-> O Q. C 9 o E CO i— i 4-> c E O • o z O) UJ o betv roger o DC H ionship and Nit 9 m o o LT) cu i- CD Id 'sqi-U 'H19N3U1S lOVdlAII 354 CO 0) > c 4) 0) •■■■ a> o ■>■» u. z o o o o CM o CO <0 CM CM '■ ■*■ o CM r-> r^. hs m +-> - r-. *■» OJ o H Z +-> ■vj- Z i o o O 1— X O) c + T" o «3" zi • o z betw onte + + UJ o c/1 s: c o •r- O) LT> + o tions Fern* •- cc + CXL O cH3 1- 03 r— -O O CO *■**■ y) transformation on the resultant 6-ferrite microstructure in welds and castings. An optical metallographic study was performed on three ASTM A351 Grade CF8M alloy castings (e.g. 18-21Cr, 9-12Ni, 2-3Mo)that possessed different 6-ferrite contents. Metallographic techniques permitted distinction between the solidification and 6-^-y transformation behavior in these castings. All of the alloys solidify in a primary 6 mode in which y precipitates before solidification is complete. The residual 6-ferrite microstructure, owing to the 6->y transformation, is generally characterized by curved segments of 6/y interface connected by straight or faceted segments. Frequently occurring interfacet angles are consistent with the facets being parallel to a {110}6 crystallo- graphic plane and an orientation relationship between 6 and y across the faceted 6/y interface. The metallographic features of the 6-ferrite microstructure can be rationalized with the aid of a 6+y transformation model analogous to that proposed for the proeutectoid y-n< transformation. 357 J INTRODUCTION The solidification behavior of austenitic stainless steels is of concern in both welding and casting technology. The term "austenitic" stainless steel is actually a misnomer with respect to weldment and casting microstructures, since they may contain significant amounts of delta-ferrite (6-ferrite), the high temperature bcc phase of iron, as well as the fee austenite phase (y). The presence of 6-ferrite can significantly alter the physical and mechanical properties of austenite. Property variations due to 6-ferrite can be beneficial or deleterious depending upon the nature of the material application and its temperature range. For example, in the case of weldments the presence of 5 to 10% 1 2 6-ferrite has been reported to reduce the incidence of hot cracking ' and microfissuring . At high temperatures (>973 K) 6-ferrite can transform to the brittle sigma (a) phase, which results in a degradation of mechanical properties . In highly corrosive environments a-phase can provide a path for selective corrosion attack of welds and castings. Delta-ferrite in the austenite matrix also generally increases the strength and lowers o fracture toughness of welds and castings , especially in cryogenic g applications . Since 6-ferrite profoundly affects the properties of austenitic stainless steel, it would obviously be advantageous to control its volume fraction and morphology in welds and castings. This report will consist of a review of 6-ferrite in austenitic stainless steel followed by a description and discussion of metallographic studies of 6-ferrite in stainless steel castings. 359 REVIEW OF 6-FERRITE IN AUSTENITIC STAINLESS STEEL Composition Effects Most research has been conducted in the area of welding austenitic stainless steels and, more specifically, on the effects of various filler metal compositions on 6-ferrite content. Some additional work has been performed to extend the concepts developed for weld metals to castings. Quantification of the effects of composition on 6-ferrite content are conveniently accomplished by Cr- and Ni-equivalents . The Cr-equivalents integrate the effects of ferrite formers (Cr, Mo, Si) and the Ni-equivalents, the effects of austenite formers (Ni, Mn, C, N) . Delong reviewed ferrite in austenitic stainless steel weld metal and 11 12 traced the development of the Schaeffler and DeLong Diagrams (Figure 1) Both of these diagrams relate 6-ferrite content, in volume percent (v/o), to Cr- and Ni-equivalents. The Schaeffler Diagram was determined by quantitative metal lographic techniques to within a ±4 v/o accuracy. The 6-ferrite contents of the DeLong Diagram were determined magnetically to within ±3 v/o, and a factor to account for the strong austenite forming effect of nitrogen was added to the Ni-equivalent. The DeLong Diagram predicts higher 6-ferrite contents than the Schaeffler Diagram for more highly alloyed materials. A recent modification of these diagrams has been made to account for the effects of precipitation in stabilized austenitic stainless steels, since precipitation removes carbon or nitrogen or both from solid solution, thus effectively reducing 1 3 austenite stability. Despite the fact that cooling rates encountered during common welding practice do not exert a significant effect on total ferrite content, DeLong cautions against a literal translation of 360 data on welds to castings. Similar empirical correlations between 14 15 6-ferrite content and steel composition have been established for castings ' and are contained in the Schoefer Diagram ( Figure 2). This diagram is actually derived from the Schaeffler Diagram and is supported by measurements 14 made on castings. All of these diagrams refer to room temperature 6-ferrite formed in a nonequil ibrium manner. In fact, the equilibrium microstructures of these austenitic stainless steels at room temperature do not contain any 6-ferrite. The departure from equilibrium can be effected by altering the solidification rate, the cooling rate subsequent to solidification, or the chemical composition. The following sections discuss the specific ways in which this departure from equilibrium affects the extent and morphology of 6-ferrite in austenitic stainless steels. Solidification Effects Numerous investigations have been conducted to elucidate the solidifi- 1 "7 _ nc I c OJ cation behavior in austenitic stainless steel welds and castings. ' Basically, the solidification behavior of this class of stainless steels can be divided into primary 6 and primary y modes according to the ratio of Cr- and Ni-equivalents. This division is frequently illustrated by construction of appropriate constant Fe pseudobinaries (isoplethes) of 17 18 27 the Fe-Cr-Ni ternary. ' ' It is important to realize that pseudo- binary sections through ternaries are only qualitative, inasmuch as tie lines drawn across multiphase fields do not represent true equilibrium 3G1 28 tie-lines. They are useful, however, for qualitatively illustrating solidification modes as a function of composition. A typical pseudobinary of this sort is depicted in Figure 3. The two-phase, y+L and 6+L regions are separated by a eutectic triangle of Y+6+L region. Below the triangle, a y + 6 phase field exists, enclosed by the y and 6 solvus lines. The y + 6 region broadens with decreasing iron content. Chromium- and Ni-equivalents are substituted for Cr and Ni to increase the generality of the ensuing review. If the nominal composition of the steel possesses a low Cr-equivalent (low proportion of ferrite formers) the primary solid phase is austenite. Similarly, in a steel with a high Cr-equivalent (high proportion of ferrite formers), ferrite is the primary solid phase. Therefore, solidification can be divided generally into primary y and primary 6 modes. The solidification mode becomes more complex in the vicinity of the eutectic triangle. For nominal compositions C 3 > Cr > C-, , the primary solid phase is still austenite; however, ferrite precipitates before solidification is complete. Likewise, for nominal compositions C 3 > Ni > C~, the primary solid phase is ferrite but austenite precipitates prior to complete solidification. The microstructures that result from solidification and cooling in the vicinity of the eutectic triangle are also more complex. The classification 1 fi 1 Q-2? 71 ? of solidification modes has been variously described in the literature. ' ■--.- / ) Table I summarizes these classification schemes. Composition ranges and application (welding or casting) are included. The ternary eutectic behavior described above is applicable to the majority of commercial austenitic stainless steel grades. In the materials 362 under consideration the eutectic has a divorced morphology since the 23 amount of liquid at the eutectic composition is generally small. In a divorced eutectic mixture one of the solidifying phases is discontinuous and present in smaller volume fraction (i.e. embedded in the continuous Of) Ol phase). ' In primary y alloys the divorced phase at the subgrain and grain boundaries is 6. In primary 6 solidified alloys the divorced 23 phase is y. Studies of solidification in conventional ingots and by a directional solidification and quenching technique have identified three possible 1 fi 1 ft 71 ?Q "3? stages of solidification in an initially primary 6 alloy. ' ' ' ' . Initial precipitation of 6 is interrupted by precipitation of y, and at the end of solidification, a transition back to precipitation of 6 is 32 observed. This corresponds to the Type B solidification mode depicted in Table I, L*L+6+L+6+y+6+y. In a conventional casting, the L+6-*l_f5+Y transition has been associated with the transition from a columnar or branched columnar solidification structure to an equiaxed structure. Segregation of alloying elements ahead of the solidification front as well as differences in cooling rate have been proposed as reasons for the interruption of initial solidification of 5 by precipitation of y from the melt and the transition back to 6 precipitation at the end of 1 fi ?7 solidification. The classification scheme of McTighe and Beech ' (see: Table I) will be used in the following descriptions. Schematic diagrams of Type A and Type B solidification extracted from the work of these investigators are shown in Figure 4. For Type A alloys, the primary 6 dendrites grow out into the liquid and 6 occupies close to 100 v/o of the as-solidified structre. With falling temperature y forms by the solid state &+y transformation, first in the primary grain boundaries 363 (y-i) and then intragranularly at larger undercoolings (y? and Y 3 )- The intragranular y can possess more than one crystal lographic orientation (e.g., difference between y 2 ancl Yo)- For Type B alloys, the primary 6 dendrites grow into the liquid and a y envelope precipitates directly onto the 6 from the liquid (y-i). The 6^-y transformation results in the growth of y (yo) ^ rom tne envelope into the 6 dendrite. Some 6 remains in the dendrite cores. The orientation of the 6 is the same within a given grain or dendrite. Similar behavior has been observed by other 29 investigators. The columnar-to-equiaxed grain morphology transition often observed in conventional castings takes place in Types A and Type B solidification. However, in Type A alloys it is observable in both the primary and secondary structures, whereas in Type B alloys it is only observable in 27 the primary structure ' (Figure 8). Primary structure refers to the structure produced by solidification; secondary structure refers to that 33 produced by the subsequent 5->7 transformation. The primary and secondary structures in Type B alloys can be differentiated by metal lographic etching techniques. 16 ' 27 ' 32 ' 34 Quenching techniques provide a powerful tool for analyzing the course of solidification by allowing the isolation of a particular stage 1 r i q 07 O Q of solidification. ' ' ' ' By quantitative metallography of specimens quenched from various points on a cooling curve, the volume fraction of primary 6 has been determined as a function of time or 29 temperature and composition for a cooling rate of 20 K/min. At this cooling rate the maximum volume fraction of primary 6 was 0.5 prior to the start of austenite growth in Type B alloys, and decreased monotonically with increasing C- equivalent (C + 0.65 N) until the onset of primary y freezing. The maximum volume fraction of primary 6 could be increased 354 further by alloying in the Type A alloy region (high Ni-equivalent) . Other investigators have reported the ability to retain increasingly larger amounts of residual 6 (subsequent to 6-^y transformation) by 1 r r\n OO OC increasing the cooling rate. ' ' ' The observed increase in residual 6-ferrite may reflect, at least partially, an increase in primary 6 with increasing cooling rate. This has been attributed to an increase in 1 fi 71 departure from equilibrium conditions with increasing cooling rate. ' In contrast, work done on austenitic stainless steel arc welds has demonstrated only a small effect of cooling rate on total ferrite content, 10 24 but a significant effect on ferrite scale and morphology. ' Concomitant with the evolution of a solidification structure is the redistribution of solute elements, commonly referred to as solute segregation. Reference to the equilibrium phase diagram indicates the various possibilities, which have been described in detail elsewhere. " Under conditions of equilibrium (slow cooling), diffusion in the solid and liquid and complete mixing in the liquid allow for complete solute redistribution, so that the composition of the solid. formed is always that defined by the tie-line intersection with the solidus curve. The solid composition at the solidus temperature is the nominal composition. A more realistic situation is one in which solid-state diffusion is minimal so that coring or 37 38 microsegregation occurs during solidification. ' Perfect mixing in the liquid and mixing in the liquid by only diffusion represent extreme situations. Real cases probably exist between these two extremes. For the cases described above, the final distribution of solute after normal freezing is depicted in Figure 5. The theoretical analyses of segregation are easily applied to planar interfaces, but it has already been implied that more complex growth forms, such as cellular, cellular-dendritic, and dendritic, are the case in most castings and welds. These solid- liquid interface types are mainly due to the constitutional supercooling produced by solute redistribution. However, they do not affect the generality of the statements made with respect to planar interfaces since the cell or dendrite can be viewed as composed of a series of small volume elements that grow into the liquid as a planar interface, A generalized theory of solidification of austenitic stainless 17 steel as primary 5-ferrite has been formulated for weldments by assuming mixing in the liquid by diffusion only. A schematic diagram in Figure 6, extracted from this work, depicts the solute distributions of Ni and Cr across a subgrain following solidification. The gradients in composition are associated with the subgrain (dendrite) core and the subgrain boundary ( interdendritic) region. Attempts to analyze the solute distribution in these alloys have relied on electron microprobe and scanning transmission electron micro- scopy (STEM) analyses. The STEM possesses a higher spatial resolution u-T4. a i u * A + i ia 1 7 , 1 8 , 2 0~ 22 , 24 , 2 5 capability. Analyses have been performed on actual welds 18 29 32 and specimens in which solidification has been interrupted by quenching. ' ' In conventional welds and castings difficulties arise in separating the parts of the total solute redistribution due to solidification and the 6-^y transformation. Assumptions regarding the nature of solidification and the 6+y transformation are generally necessary; e.g., no solute 17 24 25 redistribution during the 6-*j transformation. ' ' Reasonably indepen- dent assessments of the solute distribution coefficients and the solute profiles due to soldification have only been made in interrupted solidifi- cation studies. Solid State 6->y Transformation and Residual S-ferrite Morphologies As-cast and as-welded austenitic stainless steel microstructures do not generally represent an equilibrium state. Treatments that promote equilibrium tend to dissolve 5-ferrite, e.g. wrought grades possess virtually no 6-ferrite at room temperature and the 6-ferrite levels in castings and welds can be substantially reduced by annealing at tempera- 2fi "39 tures of approximately HOOK. ' A primary 6 alloy may contain in excess of 50 v/o 6 upon solidification. Since the room temperature microstructure of such an alloy contains substantially less 6-ferrite, it is obvious that the 6-*y transformation plays a major role in the evolution of microstructure in duplex austenitic stainless steel welds and castings. A 5+y transformation only occurs for alloys that undergo primary 6 solidification, and the nature of the transformation appears to vary with composition, cooling rate, and perhaps with solidification behavior. It is important to note that a number of labels have been associated with the various observed residual ferrite morphologies that are descriptive in nature but may not describe the source of the morphology accurately. This section attempts to describe the various ferrite morphologies, and associate them with labels commonly employed in the literature. Primary 6 alloys can be classified further into alloys remote from the eutectic triangle that solidify with almost 100 v/o 6 (forming y interdendritically) and alloys that intersect the eutectic triangle, in which solidification to y occurs directly from the melt after some 6 has solidified. In the former type of alloy, the resultant room temperature microstructure is completely determined by the solid state transformation. The character of the phase transformation depends on the degree of 1 fi 17 undercooling, which in turn, depends upon the cooling rate. ' For 357 relatively rapid cooling, the 6-^7 transformation has been referred to as 18 23 24 32 1 9 22 either a Widmanstatten ' ' ' or an acicular ' transformation, but the former is probably a more accurate description. The Widmanstatten phase transformation can be described as a thermally activated and diffusion controlled reaction in which an orientation relationship 40 exists between the parent and product phases. In this case, the 1 9 20 22 residual 6-ferrite has been termed 'lathy 1 . ' ' Although the residual 6-ferrite morphology has been termed 'lathy 1 , it is actually the austenite that forms laths by the 6+y transformation. Microstructural investigation 20 of the lathy morphology in welds has indicated the presence of several differently oriented lath bundles in a given large growth unit. The model for Type A alloys, depicted in Figure 4a, describes this case. The association of lathy 6-ferrite morphology with a Widmanstatten transformation on quenching allows for the metal lographic delineation of 1 8 32 regions in a specimen that have solidified as primary 6. ' '' At slower 1 9 22 cooling rates the mode of 6->y transformation has been termed 'equiaxial'. ' It is not exactly clear what 6-ferrite morphology this type of 6+y transformation produces. Many primary 6 alloys of interest solidify so that y envelops the primary 6 prior to completion of solidification. This mode is depicted in the model of Figure 4b and refers to the sequence L+L+6->l_+6+y. In this case there is already solid y present in the microstructure in the form of the y envelope, as a result of a L->y transformation. The solid y envelope allows for heterogeneous nucleation of y during the 6-vy transformation with less undercooling (higher temperature) than for the Widmanstatten 6->-y transformation. Therefore, the residual 6 exhibits 368 1 C\ 1 Q 1 Q oo q"7 oO less of the Widmanstatten type morphology. ' ' * * ' The morphology of the residual 6-ferrite in this case has been termed 'vermicular' 17 22 23 because of its roundish appearance. ' ' The residual 6-ferrite morphology in primary 6 alloys has also been described as a cell-like 22-24 27 network. ' It is unclear whether this morphology is associated with the vermicular or lathy 6 morphologies or is a transition morphology between the two. A continuous spectrum of ferrite morphology, from 22 vermicular to lathy, has been observed with increasing Cr /Ni . Regardless of the observed morphology, the residual 5-ferrite in primary 6 alloys is definitely located at the cores or spines of dendrite / . . \ 16,18.19,22,23,27 Tu . .. , . . .. arms (subgrains). The progressive growth of y into the primary 6 (see Figure 4) results in segregation of ' ferritizers' to the 6 at the dendrite or subgrain core, thus stabilizing it against trans- 1 o 23 24 32 formation. ' ' ' The formation of a y envelope occurs with attendant segregation of alloying elements. Since diffusion, and hence homogenization, of alloying elements is more difficult in austenite the segregation persists. The segregation of alloying elements permits etching and 27 hence, observation of the prior dendritic form at room temperature. In alloys that solidify as virtually 100 v/o 6, the ease of diffusion in the bcc lattice minimizes segregation at high temperature, and therefore the prior dendritic structure is not observable. Transmission electron microscopy (TEM) observations have shown that lathy 6 possesses a Kurdjumov-Sachs (K-S) type orientation relationship with austenite, whereas vermicular 6 possesses no crystallographic relationship with 22 austenite. Angular faceting of 6/y interfaces in primary 6 alloys has 23 also been observed and noted in castings. 369 The observations of primary 6 alloys described above have all 18 22 implied a diffusion-controlled mechanism for the 6-^y transformation. ' Other investigations, in welds, suggest that the rapid cooling rates attendant with relatively low heat input welding precludes the high temperature residence time necessary for a diffusion controlled trans- 1 -j pope formation. ' " w They propose, instead, a massive phase transformation that would allow the bcc->fcc crystal structure transition to occur at constant composition. The massive transformation is a thermally activated, 42 17 parti tionl ess phase transformation. Lippold and Savage have proposed that most of the 5 dendrite solidifies at constant composition (Figure 6). These workers are suggesting that this region of constant composition 17 23 29 6 transforms to y by a massive transformation. ' ' The assertion that the Ni and Cr segregation profiles characteristic of solidification are preserved when a massive 6^y transformation occurs is supported by 25 quantitative chemical microanalysis (STEM) across a dendrite. The concentration profiles measured are also cited as evidence for short-range diffusion across the interphase boundary as the mechanism of transformation pr (i.e., local depletion of Cr in y ana* excess in 6). Other investigations have shown that a significant amount of microsegregation results from 18 29 32 the solid state 6-+y transformation. ' ' Thus, it is uncertain whether all the microstructural observations of residual 6-ferrite in these investigations are consistent with or peculiar to the occurrence of a massive transformation. Perhaps a massive &->y transformation occurs only under conditions of rapid cooling. It is possible that an alternative explanation of the observed concentration profiles in welds lies in the enhanced homogenization associated with finer dendrite spacings. 16,27 370 Some of the observations described above have been used to explain the behavior of larger castings and ingots. In Type A alloys the first Y precipitates in the primary grain boundaries and as the temperature drops, intragranularly as well. All of these regions of y have different orientations and the columnar-to-equiaxed transition is thus observable 27 in the primary (solidification) and secondary (solid state) microstructures. In Type B alloys, where y begins to solidify from the melt after initial 6 solidification and prior to completion of solidification, columnar y crystals nucleate at high temperature and grow under steep temperature gradients across the entire ingot cross section. With the high driving force for y growth produced by the temperature gradient, changes in the 21 32 33 original 6 orientation present no barrier to that growth. ' ' Mixtures of Type A and Type B behavior are frequently observed in the same ingot. Once 6 has solidified, the y envelope is difficult to form 27 at high freezing rates. Therefore, Type A solidification might be expected in the vicinity of the mold wall and Type B solidification in the ingot interior. Variations in 6-ferrite level across an ingot cross section must be rationalized in light of the fact that higher 6-ferrite levels, associated with higher cooling rates, can be offset by rapid homogenization in the finer solidification structures, which lowers the 26 39 6-ferrite content. ' As the 6-ferrite levels in welds and castings 43 44 increase, the 6 phase becomes more continuous. ' Spheroidization 2fi of 6 can also take place during the 6-^y transformation. Measurement of 6-ferrite Content There are a number of methods that have been employed to determine the 6-ferrite levels in duplex austenitic stainless steels and most have 371 10 45 been reviewed recently. ' The commonly employed methods are magnetic measurements and quantitative metallography. Quantitative metallography is not as useful for measurements of the fine 6-ferrite structures of welds, and extensive counting is necessary to compensate for the wide intrinsic variations in 6-ferrite content within a given microstructure. Several more exotic techniques for 6-ferrite measurement have been 15 attempted. Of these, integrated Mossbauer spectra and electrolytic 43 separation ' may prove particularly useful because they are not dependent on the shape, size, orientation or distribution of the 6-ferrite. RESULTS OF METALLOGRAPHIC INVESTIGATIONS OF CASTINGS AT NBS The bulk of metallographic investigations on castings conducted at 46 NBS to date has been performed on CF8M material (ASTM A351 ) . These alloys are a cast version of AISI 316 stainless steel. Alloy designation, ferrite level, and chemical composition of the CF8M alloys are contained in Table II. These castings were keel blocks, shown in Figure 7. All keel blocks were annealed at approximately 1400 K. The macrostructure of these alloy castings is typified by the CF8M-9 alloy. The macrostructure, delineated by a light color etch,* is shown in Figure 8 and is characterized by a region of columnar dendritic grain growth perpendicular to the casting surface, followed by a transition to an equiaxed dendritic growth form. The color etch delineates the solidification structure by depositing a blue interference film on the regions of the original 6 dendrites (high Cr) and a gold film on the interdendritic regions (low Cr) . When viewed in color, the columnar region to the left of the transition boundary possesses a blue cast *C^lor Etch - A deposition etch composed of: 100 cc distilled water, 20 cc hydrochloric acid, and 2 g potassium metabisul phite. 372 whereas the equiax ed region to the right possesses a gold cast, indicative of a greater volume of interdendritic liquid in the latter region. A series of higher magnification micrographs depicts the nature of the transition in solidification growth morphology (Figure 9a, b, and c). Note that the dark gray areas correspond to the blue original 6-ferrite dentrites and the light gray areas correspond to the gold interdendritic liquid regions. The region adjacent to the casting surface is composed almost completely of 6-ferrite dendrites, aligned perpendicular to the casting surface (Figure 9a). The microstructure in the transition region is characterized by 6-ferrite dendrites that grow at angles to the heat flow direction, effectively blocking columnar dendritic growth in this direction (Figure 9b). The transition also appears to signal an increase in the volume of interdendritic liquid in the as-solidified structure. This correlates with the gold appearance of the central zone of the keel block casting when viewed at lower magnification. The structure in the central portion of the casting is characterized by randomly oriented columnar dendrites as well as equiaxed dendrites (Figure 9c). The finer network structure observed inside the original 6-ferrite dendrites is the residual 6-ferrite present at room temperature subsequent to the 6->y transformation. When a heavier color etch is applied to the same casting section a coarse columnar grain structure parallel to the direction of heat flow is observed (Figure 10). The columnar grain structure appears to grow across the previously described columnar-to-equiaxed grain morphology transition. A higher magnification macrograph of the columnar crystals 373 shows that the dendritic structure, as defined by the residual 6-ferrite, possesses the same orientation across many of the columnar grain boundaries (Figure 11). The columnar grain boundaries also appear faceted on a microscale. A 10% oxalic acid electro-etch is employed to examine the residual 6-ferrite morphology since this etch preferentially attacks the 6/y interphase boundaries. Some characteristic features of the 6-ferrite morphology are immediately apparent from observations of the CF8M-9 alloy (Figure 12). On a given plane of polish the 6-ferrite appears composed of either relatively long, filamentary particles (e.g., at A in Figure 12A) or equiaxed cell-like networks (e.g., at B in Figure 12a). The 6-ferrite at this volume fraction also appears to be composed of discrete particles that are faceted in many instances. A facet is defined as a straight segment of 6/y interphase boundary. There are many regions of apparently fine, equiaxed 6 particles (e.g. at C in Figure 12a) and other regions where 6/y interfaces are tangent to one another (e.g., at D in Figure 12a). The latter regions would seem to be locations of incipient spheroidization of larger 6-ferrite particles. A less prevalent feature of the micro- structure is the apparent twinning of the y matrix between the regions of resiudal 6-ferrite (Figure 12b). The location and morphology of the 6- ferrite is obscured in these regions when the oxalic acid etch is employed. The faceted 6-ferrite particles also seem to be characterized by definite angles that frequently occur between different facet variants (Figure 13). The frequently occurring intervariant angles are 60° and 90°. Although no independent assessment was made of the crystal lographic 374 plane in the 6-ferrite parallel to the plane of polish, these intervariant angles are consistent with the interplanar angles of different 110 6-ferrite planes ({110}6. This means that faceted 6/y interfaces could be parallel to {110)6. Comparison of the residual 5-ferrite microstructure with the original 6-ferrite dendrite morphology provides some independent confirmation of this notion. A micrograph of a color etched region of the CF8M-9 alloy illustrates this relationship for facet traces with a 60° angle between them (Figure 14). The primary dendrite axis has been 47 established parallel to a <100>6 easy growth direction by other investigators. The measured angle between the primary dendrite axis and the facet trace is 45° (see the schematic diagram in Figure 14). The assumption of the facet plane parallel to {110}6 is then consistent with the measured angle of 45° since this is one of the tabulated angles between {110}6 and {100)6 in a bcc crystal. X-ray and TEM determination of the crystal lo- graphic orientations of 6 and y are planned to confirm the accuracy of this analysis. The residual 6-ferrite distribution was employed to define the orientation of the original 6-ferrite dendrite spines in Figure 11. As shown in Figure 15, the 6-ferrite located at the dendrite spines is actually composed of a network of 6-ferrite. 375 As the 6-ferrite content of the CF8M alloys is increased, the 6-ferrite microstructure on a given plane of polish becomes coarser while still possessing many of the characteristics observed in the CF8M-9 alloy, i.e., facets, points of interface tangency, networks (Figure 16). A ION KOH electro-deposition etch is employed to deposit a light blue to light brown film on the 5-ferrite. This provides greater contrast between the 6 and y phases while eliminating the etching of y twins. The coarser 6-ferrite microstructures are also characterized by a greater apparent continuity. The 6/y interface morphology is quite intricate and can generally be described as straight facets connecting curved segments of interface. Closer scrutiny of the <5/y interfaces indicates the presence of ledges on both faceted and curved segments (Figure 17) . DISCUSSION The metallography of the CF8M alloys is consistent with the Type B solidification sequence: L->L+6->L+6+y->-6+y. In addition to residual 6-ferrite located at dendrite spines, this sequence is distinguished metallographically by a secondary structure, due to solid state transformation, in which the columnar-to-equiaxed grain morphology transition of solidifi- cation is not observed (Figures 8 and 10). Figure 8 shows the primary (solidification) structure with its columnar-to-equiaxed transition; 176 Figure 10 shows the secondary (5^y) structure composed of completely columnar crystals that bend parallel to the direction of heat flow. The fact that dendrites are observed to grow across the columnar grain boundaries of the secondary structure also supports the assertion that the columnar y structure is a secondary structure produced by the 6-+y transformation (Figure 11). In the case of a solidification sequence where a y envelope begins to solidify from the melt, columnar y crystals can nucleate heterogeneously at high temperature and grow with the aid of very steep temperature gradients. Once nucleated under these conditions, a columnar y crystal will grow right across misorientations between S-ferrite dendrites and parallel to the path of maximum heat flow (Figure 10). These metallographic observations are also consistent with respect to the various Ni- and Cr-equivalent ranges specified for primary 6 alloys in Table I. The larger volume fraction of primary S-ferrite near the casting surface is partly attributable to the higher cooling rate in this region and the difficulty in forming the y envelope at high solidification 27 rates. The interruption of 6 solidification by y solidification has been associated with the columnar-to-equiaxed transition, in which large 32 pockets of liquid exist between settling 6 crystals (Figure 9). The level of residual 6-ferrite does not appear to vary significantly with distance from the casting surface and neither does the dendrite arm spacing (Figure 9). The latter is a monitor of local solidification . 48 time and determines the time available for a diffusion controlled transformation to occur. Finer spacings might be expected to promote J77 the 6->7 transformation and thereby decrease the level of residual 6-ferrite. Therefore, the constancy of both residual 6-ferrite content and 6 dendrite arm spacing implies that the cooling rate does not vary sufficiently within the casting to affect these parameters significantly, although it still might affect the interruption of 6 solidification by y solidification. A general model for the 6->7 transformation is proposed to rationalize the observed microstructural features of the residual 6-ferrite in primary 6 alloys. The model assumes that the 6+y transformation is thermally activated and diffusion controlled. A nucleus of y forms either at solidification subgrain boundaries or on the y envelope formed by direct y solidification from the melt. These alternative modes of solid state y nucleation would affect the nucleation rate and distribution of y nucleii but not the basic morphological evolution. The shape of a nucleus and the implications of faceted nucleation and growth have been 49-52 considered generally and with specific regard to proeutectoid a. Analogous nucleation and growth is assumed to occur during the 6-+y transformation. The occurence of a facet on a nucleus is associated with formation of a low-energy interface between parent and product phase. The low energy interface, furthermore, is associated with a high degree of lattice matching (coherency) across the interface and a crystal lographic orientation relationship between the abutting phases. The segment or segments of low energy interface minimize the total surface energy of the nucleus and thereby minimize the activation barrier for nucleation. 370 This type of nucleus has been used in models of Widmanstatten-type 49-52 proeutectoid a formation (saw teeth and side plates). In this kind of proeutectoid y-*y transformation, the nucleus and precipitate faceting is associated with a Kurdjumov-Sachs (K-S) type of orientation relationship. A K-S orientation relationship, between bcc and fee phases, involves a parallelism between planes and directions of densest atomic packing in the product and parent phases. For the K-S orientation relationship the following parallelisms exist: {111>Y // {H0}a <110>y // a Faceting is definitely an aspect of 6/y interface morphology in the castings of the present investigation. Furthermore, optical metal lographic evidence suggests that the facet plane in the 6-ferrite parent phase is a {110}6 (Figures 13 and 14). This is a plane of densest packing in a bcc phase and is highly suggestive of a K-S orientation relationship existing between y and the 5 it is growing into. As mentioned in the previous section, X-ray and TEM will be employed to confirm whether or not a K-S orientation relationship exists between 6 and y in these alloys. This orientation relationship has been observed for so-called 'lathy' 6. 20 Once y nucleus formation has been considered, it remains to demonstrate the y growth process necessary for the observed evolution of 6-ferrite microstructure. Assuming a distribution of y nucleii and facet morphologies, growth is accomplished mainly through the movement of the curved, incoherent segments of <5/y interface. These are high energy and high mobility 379 boundaries. The faceted, coherent segments possess low mobility since 49-52 atom jumping across the interface is energetically infeasible. Movement of a faceted boundary perpendicular to itself can be envisioned 49-52 as proceeding by the movement of ledges or steps parallel to the boundary. These steps are actually incoherent boundary segments that possess high mobility. The presence of ledges in a macroscopical ly curved interface moving along specific crystal lographic planes can also provide a means for maintaining an orientation relationship between the parent and product phases separated by the interface. Therefore, a faceted interface is not necessary for observation of a crystallographic orientation relationship between parent and product phases. Ledges have been observed on both faceted and curved 5/y interfaces in this investigation (Figure 17). The features of the residual 6-ferrite microstructure can then be viewed as evolving through the rapid movement of incoherent segments of 8/y interface with coherent segments being dragged along behind. The coherent segments move outwards as well, by a ledge mechanism, but more slowly. Points of tangency between transformation fronts after impingement are commonplace and accelerate spheroidization processes. A schematic diagram showing the evolution of the residual 6-ferrite morphology is shown in Figure 18. It is important to note that for both the model and the actual 6-ferrite microstructures observed in this study, observations on a given plane of polish can be deceiving. These types of observation might lead one to the erroneous conclusion that discrete 6-ferrite particles are formed, even at high volume fraction. In fact, observations 380 43 in the author's laboratory and by other investigators suggest that a continuous 5-phase exists at contents greater than about 7 v/o. Therefore, although spheroidization of 6-ferrite certainly does take place at high temperature, the "discrete" particles and points of interfacial tangency observed may be peculiar to only that plane of polish (Figure 18). The characteristic filamentary 6-and cell-like 6-ferrite morphologies observed (Figure 16B) could then be considered as two views of a continuous 6-phase (Figure 18). The ferrite network can be more equiaxed if impinge- ment of incoherent segments occurs more uniformly, and this will certainly occur if the solidification structure is finer and the y nucleation rate higher. These conditions might be expected with welding and other higher cooling rate processes. Observations of the residual 6-ferrite microstructures in primary 6 weld metals show them to be more equiaxed, i.e., to exhibit a higher volume of cell-like than filamentary 6-ferrite. The cell-like networks still possess faceted 6/y interfaces (Figure 19). The similarities between residual 6-ferrite microstructures in welds and castings indicate potential similarities in their solidification and 6->y transformation behavior. RECOMMENDATIONS FOR FURTHER WORK The following are recommendations for further work: 1) TEM, selected area electron diffraction and X-ray diffraction studies to determine the fine structure of 6/y interfaces and the crystallo- graphic orientation relationships that exist between 6 and y. 2) Scanning transmission electron microscopy (STEM) studies to determine the concentration gradients of Ni and Cr that exist within the micro- structures of duplex austenitic stainless steels. Knowledge of the 381 Ni and Cr concentration profiles across 6/y interphase boundaries may assist in elucidating the nature of the 6-*y transformation. 3) Quantification of the proposed model for y transformation to confirm its ability to predict observed volume fractions of residual 6-ferrite within the constraints of normally encountered cooling rates. 332 CONCLUSIONS 1. The CF8M castings in this study are primary 6 solidification alloys with the following transformation sequence: L-H-+6-*L+6+y-*6+Y 2. The primary structure of these alloys is characterized by a columnar- to-equiaxed grain morphology transition in the casting interior. 3. The secondary structure, owing to the 6->y transformation, is characterized by columnar crystals that extend through the entire casting cross section, bending in the direction of heat flow. The secondary structure completely masks the presence of the primary structure, as noted in previous investigations. 4. The residual 6-ferrite microstructure is generally characterized by curved segments of 6/y interface connected by straight or faceted segments of 6/y interface. The 6-ferrite becomes coarser and appears to possess greater continuity as the 6-ferrite content increases. 5. Frequently occurring interfacet angles of 60° and 90° are consistent with a {110H crystallographic plane parallel to the facet plane and the establishment of a Kurdjumov-Sachs orientation relationship between 6 and y across the faceted 6/y interface. 6. The observation of ledges on curved and faceted 6/y interfaces demonstrates the possibility of interface movement via a ledge mechanism that could maintain an orientation relationship across any type of 6/y interface. 7. The features of the residual 6-ferrite morphology in this type of alloy can be rationalized with the aid of a transformation model analogous to that proposed for the proeutectoid y+a transformation. 383 REFERENCES 1. H. F. Reid and W. T. Delong: "Making Sense Out of Ferrite Requirements in Welding Stainless Steel," Met. Prog. , Vol 84 (1973), p. 73. 2. H. Fredriks and L. J. VanderToorn: "Hot Cracking in Austenitic Stainless Steel Weld Deposits," Br. Weld. J. , Vol. 15 (1968), p. 178. 3. C. D. Lundin and D. F. Spond: "The Nature and Morphology of Fissures in Austenitic Stainless Steel Weld Metals," Weld. J. Res. Suppl . Vol. 55 (1976), p. 356-s. 4. R. T. King, N. C. Cole, and G. M. Goodwin, "Creep Behavior of SMA Type 316 Weld Metal with Controlled Residual Elements," J. Pressure Vessel Techno!. , Vol. 98 (1976), p. 194. 5. L. R. Poole: "An Unwanted Constituent in Stainless Weld Metal," Met. Prog. , Vol. 65 (1954), p. 108. 6. 0. H. Henry, M. A. Cordovi, and G. J. Fischer, "Sigma Phase in Austenitic Stainless Steel Weldments," Weld. J. Res. Suppl . , Vol. 34 (1955), p. 75-s. 7. K. E. Pinnow and A. Moskowitz: "The Corrosion Resistance of Stainless Steel Weldments," Weld. J. Res. Suppl. , Vol 49 (1970), p. 278-s. 8. V. N. Zemzin and G. L. Petrov: "Influence of Ferrite Content on the Properties of Austenitic Weld Metal," Svar. Proizvod ., No. 5 (1967), pp. 6-8. 9. D. T. Read, H. I. McHenry, P. A. Steinmeyer, and R. D. Thomas, Jr. "Metallurgical Factors Affecting the Toughness of 31 6L Weldments at Cryogenic Temperatures," Weld. J. Res. Suppl . , Vol. 59 (1980), p. 104-s 384 10. W. T. Delong: "Ferrite in Austenitic Stainless Steel Weld Metal," Weld. J. Res. Suppl . , Vol. 53, (7), (1974), p. 273-s. 11. A. Schaeffler: "Constitution Diagram for Stainless Steel Weld Metal", Metal Progress , Vol. 56, (5), (1949), pp. 680-680B. 12. C. Long and W. T. DeLong: "The Ferrite Content of Austenitic Stainless Steel Weld Metal," Weld. J., Res. Suppl. , Vol. 52, (7), (1973), pp. 281-s-297-s. 13. G. Molinder: "Ni Requirements for Equivalent Structures in Austenitic Stainless Steels", Met. Prog. , Vol. 86, (11), (1979), pp. 56-59. 14. F. H. Beck, E. A. Schoefer, J. W. Flowers, and M. G. Fontana, "New Cast High Strength Alloy Grades by Structure Control," in Advances in the Technology of Stainless Steels and Related Alloys, ASTM STP 369, American Society for Testing and Materials, Philadelphia (1965), pp. 159-174. 15. L. J. Schwartzendruber, L. H. Bennett, E. A. Schoefer, W. T. DeLong, and H. C. Campbell, "Mossbauer-Effect Examination of Ferrite in Stainless Steel Welds and Castings", Weld. J. Res. Suppl. Vol. 53, (1), (1974), p. 1-s. 16. J. Beech: Personal Communication. 17. J. C. Lippold and W. F. Savage: "Solidification of Austenitic Stainless Steel Weldments: Part 1-A Proposed Mechanism," Weld. J. Res. Suppl. , Vol. 58, (11), (1979), p. 330-s. 18. S. A. David, G. M. Goodwin, and D. N. Braski, "Solidification Behavior of Austenitic Stainless Steel Filler Metals," Weld. J. Res. Suppl . , Vol. 58, (11), (1979), p. 330. 385 19. N. Suutala, T. Takalo, and T. Moisio, "The Relationship Between Solidification and Microstructure in Austenitic and Austenitic-Ferritic Stainless Steel Welds," Met. Trans. A , Vol. 10A, (4), (1979), p. 512. 20. N. Suutala, T. Takola, and T. Moisio, "Single-Phase Ferritic Solidification Mode in Austenitic-Ferritic Stainless Steel Welds," Met. Trans. A , Vol. 10A, (8), (1979), p. 1183. 21. T. Takalo, N. Suutala, and T. Moisio, "Austenitic Solidification Mode in Austenitic Stainless Steel Welds," Met. Trans. A , Vol. 10A, (8), (1979), p. 1173. 22. N. Suutala, T. Takato, and T. Moisio, "Ferritic - Austenitic Solidification Mode in Austenitic Stainless Steel Welds", Met. Trans. A , Vol. 11A, (5), (1980) p. 717. 23. M. J. Cieslak and W. F. Savage: "Weldability and Solidification Phenomena of Cast Stainless Steel," Weld. J. Res. Suppl . , Vol. 59, (5), (1980), p. 136-s. 24. J. C. Lippold and W. F. Savage: "Solidification of Austenitic Stainless Steel Weldments: Part 2-The Effect of Alloy Composition on Ferrite Morphology," Weld. J. Res. Suppl ., Vol. 59, (2), (1980), p. 48-s. 25. C. E. Lyman: "Analytical Electron Microscopy of Stainless Steel Weld Metal," Weld. J. Res. Suppl ., Vol. 58, (7), (1979), p. 189-s. 26. V. S. Raghunathan, V. Seetharaman, S. Venkadesan, and P. Rodriquez, "The Influence of Post Weld Heat Treatments on the Structure, Composition and the Amount of Ferrite in Type 316 Stainless Steel Welds," Met. Trans A, Vol. 10A, (11), (1979), p. 1683. 386 27. J. M. McTighe and J. Beech: "The Structure of As-Cast Austenitic Stainless Steels," J. Res. , SCRATA, No. 37, (June 1977), p. 43. 28. G. Masing: Ternary Systems: Introduction to The Theory of Three Com- ponent Systems , Dover, New York, (1960). 29. 0. Hammar and U. Svensson: "Influence of Steel Composition on Segregation and Microstructure During Solidification of Austenitic Stainless Steels," in Solidification and Casting of Metals , Proceedings of an International Conference on Solidification , University of Sheffield, England (July 1977), p. 401. 30. G. A. Chadwick: "Eutectic Alloy Solidification," Progress in Materials Science , Vol. 12, No. 2 (1963) Pergamon Press. 31. W. Kraft: "Solidification Structures of Eutectic Alloys, "ASM Metals Handbook , Vol. 8, 8th Edition, p. 157. 32. H. Fredriksson: "The Solidification Sequence in an 18-8 Stainless Steel, Investigated by Directional Solidification," Met. Trans. , Vol. 3, (11), (1972), p. 2989. 33. A. Hultgren: "Crystallization and Segregation Phenomena in 1.10% Carbon Steel Ingots of Smaller Sizes," J.I.S.I ., Vol. 70 (1929), pp. 69-113. 34. P. Lichtenegger and R. Bloch: "Color Etching of High Alloy Steels," Prak. Metal 1. Vol. 12, (1975), p. 567. 35. J. C. Lippold: PhD Thesis, Rensselaer Polytechnic Institute, Troy, NY, (1978). 36. J. W. Rutter and B. Chalmers: "A Prismatic Substructure Formed During Solidification of Metals," Can. J. Phys , Vol. 31 (1973) p. 15. 307 37. J. W. Christian: "The Theory of Transformations in Metals and Alloys," Pergamon Press, New York (1965), p. 566. 38. W. A. Tiller, K. A. Jackson, J. W. Rutter, and B. Chalmers, "The Redistribution of Solute Atoms during the Solidification of Metals," Acta Metal!. , Vol. 1 (1953), p. 428. 39. Y. Kinoshita, S. Tadeda, and H. Yoshimura, "On the Dissolution of 6-ferrite into Austenite in Continuously Cast Slabs of 18-8 Stain- less Steel," Tetsu-to-Hagane , Vol. 65, (8), (1979), pp. 1176-1185. 40. J. B. Newkirk: "General Theory, Mechanism and Kinetics" in Precipitation from Solid Solution ASM Seminar, American Society for Metals, (Nov. 1957). 41. G. Kurdjumov and G. Sachs: "Crystallographic Orientation Relationship between a-and y -Fe," Ann. Phys. , Vol. 64 (1930), p. 325. 42. T. B. Massalski: "Massive Transformation, Chapter 10" in Phase Transformations , ASM Seminar, American Society for Metals, Metals Park, Ohio, October 1968, p. 433. 43. T. P. S. Gill, R. K. Dayal , and J. B. Gnanamoorthy, "Estimation of Delta Ferrite in Austenitic Stainless Steel Weldments by an Electro- chemical Technique," Weld. J. Res. Suppl . , Vol. 58, (12), (1979), p. 375-s. 44. R. G. Berggren, N. C. Cole, G. M. Goodwin, J. 0. Stiegler, G. M. Slaughter, R. J. Gray, and R. T. King, "Structure and Elevated Temperature Properties of Type 308 Stainless Steel Weld Metal with Varying Ferrite Contents," Weld. J. Res. Suppl ., Vol. 57, (6), (1979), p. 167-s. 388 45. R. B. Gunia and G. A. Ratz: "The Measurement of Delta Ferrite in Austenitic Stainless Steel," Weld. Res. Counc. Bull. 132, (1968). 46. L. Genens, S. T. Wang, S. H. Kim, and E. L. Brown, "Design and Fabrication of the Cold Bore Tube for a Large Superconducting CFFF MHD Magnet", paper presented at the 1980 Superconducting MHD Magnet Design Conference, MIT, Cambridge, Massachusetts, March 26-27, 1980. 47. J. W. Christian: "Solidification and Melting," Chapter 14 in The Theory of Transformations in Metals and Alloys , Pergamon Press, New York (1965), pp. 562-566. 48. M. C. Flemings: "Solidification Processing," Met. Trans. , Vol. 5, (2), (1974), p. 121. 49. W. C. Johnson, C. L. White, P. E. Marth, P. K. Ruf, S. M. Tuominen, K. D. Wade, K. C. Russell, and H. I. Aaronson, "Influence of Crystallography on Aspects of Solid-Solid Nucleation Theory," Met. Trans. A , Vol. 6A, (4), (1975), p. 911. 50. H. I. Aaronson: "The Pro-Eutectoid Ferrite and the Pro-Eutectoid Cementite Reactions" in Decomposition of Austenite by Diffusional Processes , The Metal Society, American Institute of Mining, Met- allurgical, and Petroleum Engineers, New York (1962), p. 387. 51. H. I. Aaronson, C. Laird, and K. R. Kinsman, "Mechanisms of Diffusional Growth of Precipitate Crystals, Chapter 8", in Phase Transformations , ASM Seminar, American Society for Metals, Metals Park, Ohio, (October 1968) p. 313. 52. H. I. Aaronson, J. K. Lee, and K. C. Russell, "Diffusional Nucleation and Growth," in Precipitation Processes in Solids , Proceedings of a Symposium sponsored by The Metal Society, American Institute, Fall Meeting, Niagara Falls, New York (1976). 389 Table I. Classification of Solidification Modes in Austenitic Stainless Steels Investigation Classification Classification Composition Appl ication References 16, 27 References 19-22 L-*-L+y+L+y+6->6+y L-*L+y->T !_•*!_ +Y+Y Type A Type B Type C Type D Type A Cr = 18.2-24.2 Ni !n = 8.6-11.2 eq Cr = 18.1-21.7 Ni„ q = 10.1-13.6 eq Cr = 18.9-20.3 Nif q = 13.9-26.6 eq Cr /Ni * 1.48 eq eq Castings Welds L+L+6+6 Type C Cr /NI n z 1 eq eq .95 L+L+6+L+6+y+6+y Type B 1.48 $ Cr /Ni s eq eq 1.95 Reference 29 L^L+6->L+6+y^y+«5 $ < Castings L-M_+y+6- > Y+<5 $ = $ = Ni* ■ r eq • 75Cr eq L->L+y-*y $ > + 0.257 i 390 Table II. Alloy Designation, Chemical Composition, 6-ferrite Level Chemical Composition (w/o) Alloy Designation C Mn Si Cr Ni Mo v/o 6-ferrite CF8M-9 .02 1.06 1.10 19.33 11.67 2.14 9 CF8M-14 .05 .83 1.25 19.50 12.20 2.69 14 CF8M-24 .07 .52 1.32 21.48 9.48 3.18 24 391 List of Figures Figure 1 (a) Schaeffler Diagram (b) DeLong Diagram. Figure 2 Schoefer Diagram. Figure 3 Constant Fe pseudobinary phase diagram. Figure 4 Schematic diagrams of primary 6-ferrite solidification modes. Figure 5 Schematic diagram of solute distribution after normal freezing. Figure 6 Schematic diagram of solute distribution across a subgrain, following primary 6 solidification of a weld Figure 7 Schematic diagram of keel block casting. Figure 8 Keel block cross section showing columnar-to-equiaxed grain morphology transition in solidification structure, light color etch. . Figure 9(a) Keel block cross section in columnar grain region near edge, light color etch. Figure 9(b) Keel block cross section near columnar-to-equiaxed grain morphology transition in solidification structure, light color etch. Figure 9(c) Keel block cross section in equiaxed grain region, light color etch. Figure 10 Keel block cross section showing columnar y crystals formed during &->y phase transformation, heavy color etch. Figure 11 Columnar y crystal structure in relation to solidification structure delineated by residual 6-ferrite, heavy color etch. Figure 12 Residual 5-ferrite morphology in CF8M-9 alloy (a) filamentary and cell-like 6-ferrite, (b) twinning in y, oxalic acid electro-etch. Figure 13 Residual 6-ferrite faceting, oxalic acid electro-etch. Figure 14 Residual 6-ferrite in relation to 6-ferrite dendritic solidification structure, light color etch. Figure 15 Residual 6-ferrite networks at 6-ferrite dendrite spines, light color etch. 392 Figure 16 Residual 6-ferrite morphology (a) CF8M-14, (b) CF8M-24, ION KOH electro- deposition etch. Figure 17 Residual 6-ferrite in CF8M-24 alloy showing ledge morphology, ION KOH electro-deposition etch. Figure 18 Schematic diagram of evolution of residual 6-ferrite morphology via 6^ phase transformation. Figure 19 Residual 6-ferrite in weld metal; ~3v/o 6-ferrite, 31 6L GMA, oxalic acid electro-etch. 393 II fc£ - » c o 2 o °" 2 UJ CO ._ + 2 Z 20 21 22 23 24 25 Cr. Equivalent = (%Cr+%Mo+1.5x%Si+0.5x%Cb) ^ 18 ii x in ■ o + > # "5 o O" CO UJ + •- z z * o CO + 17- 16-, 15 I I y i / ~ V 7 -6+y) Figure 4 Schematic diagrams of primary 6-ferrite solidification modes, 397 z o (75 o Q. o o Equilibrium Perfect mixing in liquid DISTANCE Figure 5 Schematic diagram of solute distribution after normal freezing. 398 LU 5 LU LU C5 Z > o < Subgrain Boundary LU 5 LU LU o z o < Subgrain Core Nickel PRIMARY AUSTENITE Chromium (b) Boundary kC Q Cr Figure 6 Schematic diagram of solute distribution across a subgrain, following primary 6 solidification of a weld^ ' . 399 ■< — ~8 cm — h ~1 4 cm Figure 7 Schematic diagram of keel block casting. 400 Transition Transition 3 cm Figure 8 Keel block cross section showing columnar-to- equiaxed grain morphology transition in solidification structure, light color etch. 401 Columnar Grain Growth Dir 500 urn Figure 9(a) Keel block cross section in columnar grain region near edge, light color etch. 40; Columnar Region Transition Equiaxed Region 500 ym Figure 9(b) Keel block cross section near columnar-to- equiaxed grain morphology transition in solidification structure, light color etch. 403 500 urn Figure 9(c) Keel block cross section in equiaxed grain region, light color etch. 404 2 cm Figure 10 Keel block cross section showing columnar y crystals formed during 6->y phase transformation, heavy color etch. 405 20 cm Figure 11 Columnar y crystal structure in relation to solidification structure delineated by residual 6-ferrite, heavy color etch. 406 2& r (a) r-0 X^k " V\| i w (b) 250 urn Figure 12 Residual S-ferrite morphology in CF8M-9 alloy (a) filamentary and cell-like G-ferrite, (b) twinning in y ■> oxalic acid electro-etch. "Jbi i .% 7" ■>-• \ > ^ ns rl \ 4gpg v.-^ \ • ¥ , 1 . ' • -^ • /*» ■ ^. <"\ c v> ' • / X. *A „ • • /. Figure 13 Residual 6-ferrite faceting, oxalic acid electro-etch. 250 vim 408 250 urn Facet variants of residual 8-ferrite Secondary dendrite arm <100> 8 growth direction of 8 dendrite (primary dendrite arm) Figure 14 Residual 5-ferrite in relation to 6-ferrite dendritic solidification structure, light color etch. 409 500 urn Residual 8-ferrite network at dendrite spines 8-ferrite dendrite Figure 15 Residual 6-ferrite networks at 6-ferrite dendrite spines, light color etch. 410 J ' (a) (b) 250 urn Figure 16 Residual 6-ferrite morphology (a) CF8M-14, (b) CF8M-24, ION KOH electro- deposition etch. 411 (a) (b) 31 yiii Figure 17 Residual 6-ferrite in CF8M-24 alloy showing ledge morphology, ION KOH electro-deposition etch. 412 CO c E 60 ■o 0) X (0 3 M c ' / (0 M *o m o>i: i ' ent rfac< .q o 3 o v2 «- c • t: a>£ +- a> £ c ©~ o — o.E o (0 c u. o 3 C 0) a U. +-> ra E S- 1 — o ra u_ ^ 1/1 -a c •i — 03 to S_ oj +-> s- cu M- 00 O ra -C c CL o •r- >- +-> + 3 • i — CD > 4- >> O cr> O s i — ro o S- -£Z cn Q. ra ^ •i — o -a E o CD •i— +-> +-> •i— ra S- F s_ O) O) sz 4- a 1 oo ** >>* -♦" v » r 1 ■ J*" rsr^wc r i * *~ ->>^\^.^ .-; s i 1 / ;\V /./Jr c#C- • **> r ^ 125ym Figure 19 Residual 6-ferrite in weld metal; ~3v/o 6-ferrite, :16L 6MA, oxalic acid electro-etch, 41 DEFORMATION AND FRACTURE OF STAINLESS STEEL CASTINGS AND WELDMENTS AT 4 K T. A. Whipple and E. L. Brown Fracture and Deformation Division National Bureau of Standards Boulder, Colorado ABSTRACT A number of stainless steel weldments and castings that were de- formed at 4 K have been examined with optical and electron microscopy. The purpose of this investigation was to assess qualitatively the effects of residual del ta-ferrite on the deformation and fracture mechanisms of these materials at cryogenic temperatures. Weldments and castings have \/ery similar microstructures, but the ferrite in castings is much coarser as a result of slower solidification and cooling rates. This wide variation in the scale of the structure provides a good opportunity for the study of structure and property relationships in these materials. It has been found that deformation and fracture are greatly influ- enced by the ferrite morphology. The study of the coarser structure in castings has allowed more direct observation of fracture, which has led to some preliminary conclusions concerning fracture mechanisms. Cleavage fractures of ferrite particles have been observed in some castings and in an electraslag weldment. It appears that these fractures, although not forming a continuous crack path, assist in the void formation in the ductile austenite matrix. Another observed crack propagation mechanism is that of void nucleation near the austenite/ferrite interface. Although interface separation has not been observed, the inhomogeneous deformation near the interface appears to cause premature void nucleation and growth. The occurrence of one of the above mechanisms rather than the other depends on a number of factors, including ferrite morphology and the properties of the two phases. The morphology and properties, in turn, depend on both the alloy composition and solidification conditions. 415 INTRODUCTION The structural components for most superconducting magnets are being made of stainless steels. Because of the large structures, high stresses and low operating temperature, welded construction is the most attractive fabrication method. Since the structures are being made of thick plate, fracture toughness is an important consideration. It has been found that stainless steel weldments generally have a much lower 1 2 fracture toughness than the stainless steel base metal. ' To date, three factors have been established that have an effect on the cryogenic fracture toughness of stainless steel weldments: the 1 3 delta-ferrite content of the weldment, ' the nitrogen concentration in 4 2 the weld, and the fusion zone grain size of fully austenitic weldments. Nevertheless, it still is not possible to predict the toughness of a stainless steel weldment on the basis of these factors. Apparently there are factors beyond those stated above that cause the observed wide 9 C C variations in fracture toughness. ' ' This paper contains the results of the initial stage of an investigation that is intended to define more precisely the role of ferrite and to identify other factors that have an influence on the cryogenic toughness of stainless steel weldments and castings. In many structural applications it would be advantageous to use stainless steel castings to facilitate fabrication and to eliminate the need for welding. There are presently wery few 4 K mechanical property data on castings, but a few initial tests have shown a wide variation in 5 fracture toughness similar to weldments. This is not suprising 417 because weldments are actually rapidly solidified castings. Stainless steel castings also generally contain a certain percentage of delta-ferrite, although on a much coarser scale than that of weldments, owing to the slower solidification of castings. This paper presents the results of a metal lographic and fractographic investigation of deformation and fracture mechanisms in duplex austenite/ ferrite structures at 4 K. MATERIALS The materials examined in this study are stainless steel weldments and castings that have been tensile and fracture toughness tested at 4 K. These materials did not come from a controlled series, however, they all fall into the AISI type 304 or 316 stainless steel composition ranges. These variations in composition, as well as variations in the casting or welding process, are appropriate since the intent of this study was to investigate the possible deformation and fracture mechanisms in duplex austenite/ferrite structures at 4 K and not just the mechanism for a specific material. The available chemical composition and mechanical property data for the materials used in this study are presented in Tables 1 and 2. PROCEDURES For this investigation a number of sectioning, metallographic and fractographic techniques were employed. Each technique allows the observation of different phenomena. Macroscopic examination, optical microscopy, and scanning electron microscopy (SEM) were used for making observations. 418 Sectioning Many of the figures presented in this paper are taken from samples that were sectioned perpendicular to the fracture plane of tensile and compact tension specimens. This sectioning method allows the observa- tion of crack propagation paths and deformation around the crack tip and provides insight into the crack propagation mechanisms. Metallographic Specimen Preparation For the most part metallographic specimens were prepared by using standard mechanical grinding and polishing techniques, but in a few cases electrolytic polishing was used. A 10% oxalic acid electrolitic etch was used as a general purpose etch to reveal ferrite morphology and grain boundaries. This etch can also reveal strain-induced martensite if the etching time is increased; however, if the specimen contains ferrite, this procedure leads to con- siderable surface relief and pitting, which make metallography difficult. A 10 N KOH electrolytic etch was used to reveal ferrite structure only. This is a deposition etch that colors the ferrite light blue to brown. The main attribute of this etchant is that it does not remove any material; therefore, it will not round the edges of cracks or produce pitting and other etching artifacts. The KOH etch is sometimes used prior to oxalic acid etching because it tends to passivate the ferrite and allow deeper oxalic acid etching of the austenite with less overall damage. A solution of 50 ml HC1 , 125 ml H^O and 2g Na-SO* was used to color etch stainless steel weldments and castings. This etchant is useful in 419 revealing the original solidification pattern after the solid-state ferrite-to-austenite transformation has occurred. A magnetic etching procedure was used to reveal the magnetic phases, ferrite, and bcc martensite. For this procedure a suspension of fine magnetic particles is spread on the surface of the specimen and the specimen is placed in a magnetic field; the magnetic particles will then migrate to the magnetic phases. This etch was used mainly in conjunction with the KOH etch to reveal bcc martensite which is an indication of plastic deformation in the austenite matrix. DISCUSSION Low Temperature Deformation It is generally true that bcc materials, such as del ta-ferrite, go through a ductile- to-brittle transition at low temperatures, whereas fee materials, such as austenite, do not. This is due to the difficulty of operating non close-packed slip systems at low temperatures. The tem- perature at which this transition occurs in steels is influenced mainly by alloy content and microstructure; a higher nickel content and a smaller grain size lead to a lower transition temperature. Another factor to consider is the increase in strength with a reduction in temperature; bcc materials have a much more dramatic increase in yield strength than fee materials at low temperatures. These general observations lead to some expectations for the behavior of duplex stainless steels. At higher temperatures, both austenite and ferrite are ductile and have similar strengths; in this case the amount 420 or morphology of the ferrite is not expected to significantly affect the strength or toughness of the material. At temperatures below the transition temperature of the ferrite, it becomes a strong, brittle second phase. Factors that influence the inhomogeneity of deformation in the ductile austenite matrix will influence the strength and toughness of the ma- terial. This means that the amount, morphology and properties of the ferrite can all have significant influences. Observations that support these hypotheses have been made by Read et al. Figures 1 and 2 are micrographs of 316L stainless steel weld metal tensile specimens that were tested at room temperature and 4 K, respec- tively. Note that at room temperature (Figure 1) the ferrite is highly elongated in the direction of the tensile axis and at 4 K (Figure 2) the ferrite maintains its original dendritic morphology and shows no indica- tion of elongating with the matrix. During deformation at low temperatures, austenite can go through a strain induced martensitic transformation. The degree of transformation and the critical strain at which it occurs are a function of the alloy content. Although the stability of the austenite matrix will quanti- tatively effect the strength and toughness of duplex austenitic stain- less steel castings and weldments, it is only of interest in this qualitative discussion in that it provides a method of observing defor- mation in the austenite. Figures 3 and 4 are micrographs of areas near the crack tips in two CF8* stainless steel castings tested at 4 K. The intent of these micro- *CF8 and CF8M are ASTM A351 grades of cast stainless steel and are analogous to AISI 304 and 316, respectively. 421 graphs is to show the inhomogeneity in deformation caused by the pres- ence of ferrite. Figure 3 is from a 0% ferrite casting; the straight lines near the crack tip are strain-induced martensite. Note that the lines are straight and fairly uniform, only changing direction at grain boundaries. This can be contrasted with Figure 4, which is a similar region in a 14.5% ferrite casting. In this case, the martensite forma- tion is extremely nonuniform and seems to be concentrated in association with the ferrite morphology. Another example of the effect of the presence of ferrite on deformation patterns in the austenite matrix is shown in Figures 5 and 6. In Figure 5, bending of martensite bands near the ferrite particles can be seen in a CF8M* casting with 9% ferrite, indicating that the ferrite is nondeforming and the matrix is flowing around it. Figure 6 is from a 316L electroslag weldment; this micro- graph shows the reinforcement of the austenite matrix by the ferrite in a low strain area. Note that there is little martensite formation between the parallel ferrite plates and more extensive martensite formation where the matrix is not as well reinforced. All of the above evidence lends support to the observations cited at the beginning of this discussion: (1) at low temperatures ferrite is a \/ery strong nondeforming phase and (2) the ferrite content and morphology are important influences on how the ductile austenite matrix deforms. Low Temperature Fracture As a result of the above discussion on low-temperature deformation, it appears appropriate to classify duplex austenitic stainless steel, at *CF8 and CF8M are ASTM A351 grades of cast stainless steel and are analogous to AISI 304 and 316, respectively. 422 low temperatures, as a ductile matrix composite with hard, brittle particles or discontinuous fibers. This is not an exact classification, since the ferrite is very seldom in the form of spherical particles and can be either continuous or discontinuous depending on solidification and cooling conditions, but this classification does provide a good starting point for the discussion of fracture mechanisms. There are three possible mechanisms for the fracture of ductile matrix/hard- discontinuous fiber composites. 1. Brittle fracture of the fibers. 2. Matrix/fiber interface separation. 3. Fiber pull-out. Brittle fracture of the fibers occurs when the local stress in the fiber exeeds its fracture strength. The local stress in the fiber is not necessarily the same as the macroscopic stress on the sample, but can be considerably higher owing to inhomogeneous plastic deformation in the matrix. Interface separation is an important consideration with artifical composites and with metal-matrix/particulate composites; however, the interface between austenite and ferrite is a yery strong one and no evidence of interfacial separation has been observed. Fiber pull-out is sometimes associated with interface separation, but for the purpose of this discussion it means a ductile failure of the matrix near the matrix/fiber interface. This occurs when the fiber does not fracture, but the highly inhomogeneous deformation of the matrix near the fiber results in void nucleation at this location. The occurrence of one of these mechanisms over the others depends on fibre strength, content and morphology. 423 In this study evidence was found for both fiber fracture and fiber pull-out. Until this investigation there had been no direct observation of the brittle fracture of ferrite in weldments and castings at 4 K. Although this had long been assumed to be the cause of the reduction in toughness of stainless steel weldments with increasing ferrite content, direct observation was difficult because of the extremely small size of the ferrite in weldments. In this investigation observations were made on both castings and weldments made with high heat input processes. The much slower solidification and cooling rates of these materials results in a coarser structure, which is more easily observed. With these materials extensive evidence was found of brittle fracture of ferrite in weldments and castings deformed at 4 K. Figure 7 is a micrograph of a 316L electroslag weldment taken near the fracture surface of a tensile specimen. In this specimen, numerous brittle fractures of ferrite particles are evident. The fact that the fractures are at 90° to the tensile axis lends further support to the conclusion that they are brittle cleavage fractures. Additional evidence for cleavage fracture of ferrite is shown in Figure 8 (a graph from the 14.5% ferrite CF8 casting). This micrograph shows consistant angular relationships between the fractures, indicating that fracture is occurring on a specific set of crystallographic planes, i.e., cleavage planes. Figure 9 and 10 are crack profiles for two stainless steel cast- ings. Figure 9 is a CF8M casting that contains 9% ferrite and has a very high toughness at 4 K, similar to that of wrought base metals. Figure 10 is a CF8 casting containing 14.5% ferrite; this casting has a low toughness comparable to that of a weldment with a similar ferrite content. 424 In both cases the crack path is associated with the ferrite morphology, but Figure 10 has a much more brittle appearance and significant cracking not associated with the ultimate crack path. Also, it can be seen in Figure 4 (a higher magnification of the crack tip) that there is significant ferrite cracking in advance of the crack tip, which was not observed in the CF8M casting of Figure 9. A definitive explanation of the difference in these two castings cannot be offered at this time; however, two factors appear to have contributed the difference in the amount of ferrite cracking. First, the ferrite morphology is different; the tougher CF8M casting was a sand casting with a slower solidification and cooling rate than the centrifugal ly cast CF8 casting. This resulted in a coarser and discontinuous ferrite morphology, which provided less restriction to deformation in the austenite matrix than the finer more continuous ferrite in the CF8 casting. The second contributing factor may be the toughness of the ferrite phase: CF8M has a higher nickel content than CF8, and it is well known that nickel increases the cryogenic toughness of ferritic materials. Weldments differ from castings mainly in that the solidification and cooling rates are much greater. This results in a more directional solidification pattern and a much finer structure. The greater cooling rate results in less diffusion during the ferrite-to-austenite transformation causing the ferrite to be more closely associated with the original dendritic solidification pattern. Therefore, the residual ferrite of weldments is more likely to be continuous. When fracture is associated with cracking of the residual ferrite, having continuous ferrite can provide 425 a preferred crack path and lower toughness than occurs for a material of the same ferrite content in a discontinuous morphology. An example of this preferred crack path is shown in Figure 11 in the crack profile of a fracture toughness specimen from a 316L electroslag weldment. An electroslag weldment experiences a cooling rate greater than that of a casting, but much less than that of other welding processes. This results in a relatively coarse but continuous ferrite morphology. Figure 11(a) shows that the dendritic growth direction determined the crack propagation direction even though the stress state was such that the crack should propagate straight. The closer view, Figure 11(b), shows the crack was propagating down dendrite cores and secondary dendrite arms. The crack periodically moved to an adjacent dendrite owing to the stress state. Figure 12(a) is a micrograph of the crack tip of this weldment, showing that the crack is indeed propagating in a brittle manner through the ferrite. Although this fracture does have a very brittle appearance, the measured fracture toughness was above average for a weldment of this ferrite content. Figure 12(b) shows the same area using the magnetic etch to reveal strain-induced martensite. Even though the ferrite fractures in a brittle manner, there was still considerable plastic deformation in the austenite matrix. This matrix deformation and the fact that energy was consumed in the crack propagating off-line accounts for the relatively high toughness of this weldment. Further evidence that the ferrite in this weldment fractured in a brittle manner is provided by the SEM fractographs in Figure 13. These fractographs demonstrate the mixture of brittle ferrite fracture and ductile matrix failure. 426 Behavior similar to that described above can also be seen in the rapidly solidified 308L flux-cored metal arc (FCMA) weldment shown in Figures 14 and 15. Because of the rapid solidification, the ferrite is extremely fine and it was not possible to observe brittle fracture of the ferrite directly. Figure 14 shows that the crack propagation again followed the dendrite growth direction. The closer view of Figure 15 shows that the crack path was determined by the local ferrite morphology, which varies considerably from the bottom to the top of a weld pass. As mentioned previously, two possible discontinuous fiber composite fracture mechanisms are of interest: fiber fracture and fiber pull-out. The evidence already presented has shown that fiber fracture does occur in coarse duplex structures. In the case of weldments with much finer structures, limited evidence has been found for both fracture mecha- nisms. Figure 16 is a stereo fractograph of the fracture surface of the 308L FCMA weldment described previously. The area indicated by the arrow appears to be a subsurface ferrite dendrite, where the inhomo- genous deformation nucleated a ductile failure near the ferrite/austenite interface. Composite fracture theory predicts that, for a discontinuous fiber composite, there is a critical aspect ratio (fiber length-to- diameter ratio) at which a transition from fiber fracture to fiber pull-out fracture mechanisms occurs; a high aspect ratio favors fiber fracture. The aspect ratio of a fiber is dependent on orientation to the tensile axis: the same fiber will have a high aspect ratio if it is aligned with the tensile axis and a low one if it is perpendicular to the tensile axis. In Figure 16 it is apparent that the dendrite was perpendicular to the tensile axis; therefore fiber pull-out occurred. 427 Figure 17 shows a case where the fibers were apparently parallel to the tensile axis and fiber fracture occurred. SUMMARY Samples of stainless steel castings and weldments that had been deformed and fractured at 4 K were examined. The thrust of the investigation was to determine the mechanisms by which residual del ta-ferrite affects the deformation and fracture of duplex austenite/ferrite structures at cryogenic temperatures. Evidence was presented that indicates that at low temperatures the ferrite acts as a hard, brittle second phase particle. The manner in which these particles affect the mechanical properties of the material depends on the amount, properties and morphology of the second-phase particles. It was observed that the ferrite reinforces the austenite matrix and causes the deformation to be inhomogeneous. Evidence was found for two fracture mechanisms: (1) brittle fracture of the ferrite mixed with ductile failure of the matrix between ferrite fractures and (2) ductile failure in the matrix near the austenite/ferrite interface, which was caused by the highly inhomogeneous deformation in this region. The occurrence of one of these mechanisms over the other depends on the morphology and properties of the ferrite. Future work in this area will concentrate on obtaining a more quantitative description of the role of ferrite in the deformation and fracture of duplex austenite/ferrite materials at liquid helium temperatures. 428 REFERENCES 1. D. T. Read, H. I. McHenry, P. A. Steinmeyer, and R. D. Thomas, "Metallurgical Factors Affecting the Toughness of 31 6L SMA Weld- ments at Cryogenic Temperatures," Weld. J. Res. Suppl . 59 , 104s- 113s (1980). 2. T. A. Whipple, H. I. McHenry, and D. T. Read, "Fracture Behavior of Ferrite-Free Stainless Steel Welds in Liquid Helium," to appear in Weld. J. Res. Suppl. 60 (4), 1981. 3. E. R. Szumachowski and H. F. Reid, "Cryogenic Toughness of SMA Austenitic Stainless Steel Weld Metals: Part I - Role of Ferrite," Weld. J. Res. Suppl. 57, 325s-333s (1978). 4. E. R. Szumachowski and H. F. Reid, "Cryogenic Toughness of SMA Austenitic Stainless Steel Weld Metals: Part II - Role of Nitrogen," Weld. J. Res. Suppl. 58, 34s-44s (1979). 5. T. A. Whipple and H. I. McHenry, "Evaluation of Weldments and Castings for Liquid Helium Service," in Materials Studies for Magnetic Fusion Energy Applications at Low Temperatures-IV , NBSIR 81-1645, pp. 273-287 (1981). 6. T. A. Whipple and D. J. Kotecki, "Weld Process Study for 316L Stainless Steel Weld Metal for Liquid Helium Service," in Materials Studies for Magnetic Fusion Energy Application at Low Temperatures- LV, NBSIR 81-1645, pp. 303-321 (1981). 7. A. Kelly, "Particle and Fibre Reinforcement," in Strengthening Methods in Crystals , Halsted Press Division, John Wiley & Sons, New York, pp. 433-484 (1971). 429 Table 1 Table 2 LIST OF TABLES Process and chemical composition data on stainless steel weldments and castings. Mechanical properties of stainless steel weldments and castings 430 s- o • 14- to en T- (O -t-> -a to c o o •r- -O +J c •r- fO CO O tO Q-4-> E s= O CO a e -o ra co ■*■• E 1- co a) -C co o +j to -a c to c co T- o S- a. to o Q. E o •>- co o o co •r- S- s- ai o o «3- o "Sj" o Ln "* O O , — «=J- CM CO CM O O O o o o 00 CM CM o 00 o «3- 00' to en CM to 00 o o •vj- CM to CO co CM o o o to to o CO CO I — CO LO I — LO "vl- cr> LO "5f ■3- 00 o 00 U_ o 00 Ll. CM O to O O CM O 025 • • • 1 1 CM 00 (U I — r— en 3 (0 (0 B3 "O x> C3> D> en en en p— •a ■— CO 3 C 3 c E to co cu S- i. 4_ .,_ <4- •!- •r - O en 2 (O +-> (0 ■a fO "a E S- x CO ai fO r— 3 t— CO C_) 00 LU CO U- E 00 00 CD CO co 431 CO +-> c co 00 CO a> c •r- • fO Si +-> *3- (/> 4-> «»- +-> 1_ CO c o » •i- fO +j CU a s- 3<5« 1 "O 1 CU c CrC -r- en? o CU -C CO +j +-> -*s ro CD E S= •r- CU +-> s- fO r— ■!-> CL. =3 oo s: • CO -C .^ T3 4-> r— CO CU C •r- CU (13 >- S- Q- •m s: oo cu E o o CM 10 00 oo CO CO 00 en CM co CT1 o en CM CM en 00 CM co Lf) CD CO O CO 00 CO o o co 5 o LO LO o CTl CO LO *3- o LO •si- CO U_ O 00 Li_ O 00 C_5 CO CO CM i— CM CM 00 «=t i— CM CM co r-~ CT> r ~ ,— 00 «=J- CO LO r— CO co co i — •=1- CM 00 LO i— LO cr> oo CM oo" CU r— r - " CD 3 (0 CT» n3 OT en to -o o> c D> C c i — -a cu cj 13 -r- 3 -r- •r— CO CU s- s- M- +-> 4- +-> 4-> o x> CT O rtJ •i— CO •r- CO CO S_ r- 1_ o S- CO S_ (0 (O +J CU CU 1 1— +j a -i-> o -o O O 3 e x ■* a-. •§'!•• (?/%&$£ r.wJ A & xf r "' TENSILE AXIS I 200 y m Figure 2 Micrograph of a 31 6L SA weld metal tensile specimen deformed at 4 K, showing that the ferrite does not deform at low temperatures Oxalic acid etch. 436 ■MujejByiM^ ~^r^aMfi^Ji||(jMTLj»n| * r. : V ^ * •.•*• *' J: 7 : t <■ it o o ■St a) +-> •1 — s_ i. OJ <+~ _ C"*i o JC o re +J <3J 4- O -o •i™ a. CJ "r— fO +J O J* *i — o i — fO a (0 a (0 OJ DO s_ u_ ^ * '""■--*. / - ■■■ *'*' LD "O t < * : " ■" ' v 'w **< • *r— ; -" * h <* o ' ' ■'■■ \ !„, *V -^ i— fO .'. * %/*%,': 1 " » ' • re O 7' ~ ^ »l— • ' ■.- 1 -*- J q_ ,_ O ro .•- -:V^ ; X a. o •r— • .. * , ■*» * ~3s^s*~ » .%*•«** -iii fO u *'■•■- •'•• *.".■■ .***"; rC 3= S- O a ^ '/■ 1 «*- \ V ( 438 cu +-> en c: S- -r- rO -t-> O) to C ra o rO CD CO s- u_ cu s_ CD 50 pm Figure 5 Micrograph of a 9% ferrite CF8M casting deformed at 4 K, showing the deformation of the austenite matrix around a ferrite particle, Color etch. 439 TOO u m Figure 6 Micrograph of a 316L electroslag weldment deformed at 4 K, showing the reinforcement of the austenite matrix by the ferrite particles. KOH and magnetic etch. 440 • r t I TENSILE AXIS & • 50 ym Figure 7 Area near the fracture surface of a tensile specimen from a 31 6L electroslag weldment. Oxal ic acid etch. 441 . Ml ! W 50 ym Figure 8 Area near the crack of a 14.5% ferrite CF8 casting. KOH and oxalic acid etch. 442 E 3- O o 00 +-> y; u CO Li- &- 4- Ol tO 4- O J^! ^ O o O r— £- o cr> CD S- C7> 443 « J « CO •I — +-> rD <-> CO u_ O CD +-> ■i — S- i- O) 4- ft-S .c: U~> (_> • 4-5 <* CU i — "O IT3 •r— a 4- ra O o o> •1 — 1 — 1 — •r— (T3 4- X o o S- Q.-0 cr ^ fTD o fO n: S- o <_) i^ O) 13 444 5 mm (a) 4 LOADING DIRECTION ♦ Figure 11 Crack profile of a 31 6L electroslag weldment Oxal ic acid etch. 445 &- \> (a 50 urn vj* ' • ' .'-^n.^: v. '<' ^r f jf* , y (b) 50 ym Figure 12 Micrographs of the crack tip of a 31 6L electroslag weldment. a) KOH etch, b) KOH and magnetic etch. 446 (a) 400 urn (b! 50 ym Figure 13 SEM fractographs of the fracture surface of a 31 6L electroslag weldment. 447 I LOADING DIRECTION I (a) 3 mm « ^wv - H 'J-sd (b) 1 mm Figure 14 Macroscopic crack profile of a 308L FCMA weldment. Oxalic acid etch. 448 *£&* \\ I m E 3- O O CO O $fe* ifilf <+- o c o • •r- JZ CD O O) -t-> S_ CO Q.-0 •I — »l — ■*-> o ro o o (O •«- t- r— O rO X CD O .jr +-> (+_ 4_> o c cu JC E cl-o ro r— CD S o s- < O S •r- CJ) 01 1 449 \ 50 um Figure 16 SEM stereo fractograph of a 308L FCMA weldment. 450 25 ym Figure 17 SEM fractograph of a 308L FCMA weldment. 451 EVALUATION OF WELD FILLERS TO JOIN NITROGEN STRENGTHENED STAINLESS STEELS FOR CRYOGENIC APPLICATIONS (FUSION AND MHD ENERGY) R. H. Espy Armco Inc., Research and Technology Middletown, Ohio ABSTRACT Tests are reported of austenitic stainless steel weldments that use filler materials manufactured outside the United States. 453 Evaluation of Weld Fillers to Join Nitrogen Strengthened Austenitic Stainless Steels for Cryogenic Applications (Fusion and MHD Energy) R. H. Espy - Armco Research & Technology US/USSR Technical Exchange Program on Welding Stainless Steels BACKGROUND INFORMATION During the exchange visits of the Russian and United States Cryogenic Welding Groups both in the USA and Russia, information on weld fillers for joining the Nitrogen Strengthened Austenitic Stainless Steels was presented. While the weld compositions studied in the U.S.A. had acceptable strengths and toughnesses, there was some concern about the presence of micro fissure defects that were almost always present. The concern was mainly that of the effect of defects under cyclic loading conditions. Initially the cooperative effort between the U.S. and USSR was to exchange sample weldments having defects typical of those encountered in welding fully austenitic stainless steels. During our first ex- change it became evident that the defects we were talking about were metallurgical in nature (micro fissure) and the defects the Russian delegation was referring to were of the workmanship type like incom- plete penetration, unfused or unwet areas, cold laps, entrapped slag, etc. Actually, we did not consider workmanship defects as being important because of our nondestructive test and repair practices which are very common and effective throughout the USA where high quality weldments are desired. On the other hand, the Russian group did not consider metallurgical defects like micro fissures because they did not occur with the weld fillers they were using. In general one might say that Russian weld metallurgy in this area appeared more advanced than in the U.S., but on the other hand that U.S. manufac- turing technology appeared more advanced than the Russian manufac- turing capability. The Russian approach of applying large numbers of people to technological problems appears to be the reason for this. In Russia's Paton Welding Institute alone there were, in 1979, 34 engineers and 17 technicians assigned to welding the stainless steels. As a first step in resolving this major difference we inquired con- cerning the typical composition of the Russian weld filler. This was received along with a coil of bare wire for our laboratory evaluation. Just prior to our exposure to Russian technology in the exchange program, we at Armco decided to examine two foreign made weld fillers touted to be good for weldments used for 455 ., Cryogenics. One designated as Grinox 41 (stick electrodes) and Grinox 5 (bare wire) is made by Messer. Griesheim in Germany (distributed by MG Weld Products - Menomonee Falls, Wisconsin) . The second designated as P6 (stick electrodes & bare wire) is made by Avesta in Sweden (distributed by A. Johnson & Co . - Lionville, Pennsylvania) . The results of our work with these two wires are included with those of the Russian wire in this report. COMPOSITION OF WIRES & WELD DEPOSITS Analyses of the filler wires and weld deposits (for covered electrodes) showed them to have the following compositions. A typical composition of fully austenitic weld fillers studied in the U.S. is shown for comparison. - Russian Bare Wire CMnP_ £2i_C_E Ni Mo W N .027 6.7 .01 .003 .09 19.45 16.1 2.8 2.1 .23 German Grinox 41 .17 9.9 .012 .004 .80 15.6 12.4 - 2.3 Weld Deposit Swedish Avesta P6 .03 4.9 .015 .001 .25 18.5 17.5 2.6 - .17 Weld Deposit U.S. Armco .05 9.0 .007 .009 .22 18.0 16.2 - - .18 Weld Deposit PROCEDURE USED IN PREPARATION OF ALL-WELD TEST SPECIMENS The weld geometry employed was as specified in AWS 5.4-79 except the base plate was 1" thick, the root opening was 3/4" and the buttered overlay was deposited as one layer using 5/32" Type 309 Stainless covered electrodes . The parameters for welding with the bare wire were as shown. Weld Process - Automatic Gas Metal Arc (Pulse) Filler Wire - .062" <|> X Coil Shield Gas - Argon + 2% 2 at 30 CFH Weld Current - 160 amperes Volts - Average, 19.5 Peak, 66.0 Cycles - 120 Weld Travel Speed - 15 IPM Stringer Bead Practice Used (No Oscillation) Weld Interpass Temperature RT/300°F 456 3. The parameters for welding with covered electrodes were as shown. Weld Process - Shielded Metal Arc (Manual) Filler - 5/32" $ Weld Current - 100 amperes Weld Voltage - DCRP 24 volts Stringer Bead Practice Used Weld Interpass Temperature RT/300 F The .505" tensile and standard Charpy V-notch specimens were sectioned from the top center of each deposit. The notch in each Charpy specimen was oriented top to bottom of the weld deposit. TEST RESULTS All of the items welded in a satisfactory manner with the processes employed . The room temperature and -320 F tensile and Charpy data are shown below. Included again for comparison are data obtained with the U.S. weld filler. Tests at Room Temperature Tensile Toughness Defects Defects UTS .2% YS % El on Fracture CVN on Fracture ksi 104 ksi 73 2" 41 % RA 56 Face 114 Face Russian None 1 ppt. German Grinox 41 95 73 35 65 Some 11 (1) Separations 116 None Swedish Avesta P6 95 66 41 63 None 95 None U.S. Armco 93 71 25 26 Tests 15% <2) M.C. at -320°F 44 25% M.C. Russian 186 133 19 26 5%( 2 ) M.C. 68 3 small ppts German Grinox 41 173 118 31 31 Some 11 Separations 62 None Swedish Avesta P6 179 123 40 40 None 37 None U.S. Armco 162 134 10 12 io% (2) M.C. 13 20% M.C. (1) Appear to be heat affected area grain boundary separations parallel to direction of welding. (2) M.C. (Micro Condition) appear to be areas of grain boundary separations and weakness associated with the freezing pattern, 457 Me tallo graphic examination of the weld structures to determine reasons for soundness or conversely lack of soundness , showed the molybdenum and tungsten bearing weld metals to contain coring or a phase that reacted sufficiently different to electrolytic etches so as to stand out significantly. These weld deposits were also free of grain boundary breaks in the high temperature heat affected areas of the multipass welds. The sensitized or lower temperature affected areas of the German Grinox weld contained many grain boundary precipitates of carbides. The tungsten and molybdenum free welds (U.S. - Armco) showed numer- ous grain boundary breaks in the high temperature heat affected zones. These structural conditions and their affect on tensile and impact tests are shown in Figures 1 to 4. X-ray diffraction examination of the areas in the tungsten and molybdenum containing welds that showed a differential reaction to the etchants showed all areas to have the same crystal structures. The SEM analytical efforts were not successful in showing molybdenum and tungsten segregation, the elements believed responsible for the etching phenomena. Microprobe efforts with a Phillips AMI/3 Electron Microprobe to identify element separation were not successful. A new electron micro- probe unit, Cameca CAMEBAS at Armco, with which it is believed that proper identification of the highlighted areas can be made will be in service shortly. The results from this effort will be reported separately when available. DISCUSSION The use of tungsten and molybdenum together as in the Russian wire or separately as in the German and Swedish electrodes is effective in controlling the high temperature hot short con- dition found to occur in the welds made with the U.S. - Armco electrodes . All of the welds showed increased yield strengths at cryogenic temperatures confirming the fact that both N&C influence yield strengths as temperatures are lowered. The German Grinox electrode contained residual nitrogen (low) with C near .20% while all other Russian, Swedish and U.S. fillers contained N at about .20% with C at a low level. While all of tungsten and molybdenum containing welds were acceptably sound and exhibited good properties at all temper- atures there was a trend in the high carbon German alloy towards a tensile fracture pattern that reflected the presence of grain boundary carbides in lower temperature or sensitized areas of the multipass welds. While the Russian filler showed low indi- cations of porosity and micro condition, both it and the Swedish alloy appeared overall quite good. 458 CONCLUSIONS AND COMMENTS The use of tungsten and molybdenum in ferrite free stainless weld metal is effective in controlling the grain boundary breaks that have been found to occur in high N containing welds . Normally ferrite in high nitrogen stainless welds is an effective means for producing sound defect- free welds, but ferrite when present has an adverse effect on toughness at cryogenic temperatures when associated with the high strength imparted by nitrogen. While all ferrite-free stainless steel welds are prone to form some grain boundary breaks , the use of nitrogen for strength caused significant increases. The data reported shows that both tungsten and molybdenum are effective in controlling cracking with the pro- cesses used. The German Grinox 41 (stick electrode) and Grinox S (bare wire) and Swedish Avesta P6 (stick electrode and bare wire) are commercially available for any added weld process work that may be desired. Armco is currently making an effort to produce the Russian wire so as to make it also available for additional weld process test work. Overall the US/USSR exchange program has been beneficial in that the higher nitrogen stronger stainlesses like 304 LN (.15 N) and NITRONIC 40 ( . 30 N) can be considered for weldments at the liquid helium temperatures associated with fusion and MHD energy projects. The initial objective of determining the effect of weld defects still remains to be addressed and should be considered in future exchanges. While the use of tungsten and molybdenum were effective in controlling the nitrogen induced grain boundary defect, there were indications that a small amount of the defect still occurred and this may vary with the weld processes employed. It is suggested that the program to determine the effect of defects be continued in future exchange visits. R.H.Espy 10/20/80 459 Mag 8X Fracture Face of CVN Tested at -320°F Mag 1.5X Tensile Tested at Room Temperature I i I / '* * < ;■•':'■■, « v if l h * , ■ 4 \ * Etchant NaOH Mag 500X Coring or Phase in Ferrite-Free Austenite FIGURE 1 : Impact, Tensile and Microstructure for Weld Deposited Using Russian Wire c Mn P S Si Cr Ni 027 6.7 .01 N .003 Mo .09 W 19.5 16.1 .23 2.8 2.1 Toughness at -320 F CVN Defects 68' # 3 small ppts Tensile at Room Temperature UTS .2% YS ksi ksi %E1 2" %RA Defects 104 73 41 56 None 460 Fracture Face of CVN Tested at -320°F Mag 1.5X Tensile Tested at Room Temperature I . ' *& ■»tT hf- ■ § Etchant NaOH Mag 500X Coring or 2nd Phase in Ferrite-Free Austenite FIGURE 2 : Impact, Tensile and Micro structure for Weld Deposited Using German Grinox 41 Electrodes CMn P S Si Cr Ni W .17 9.9 .012 .004 .80 15.6 12.4 2.3 Toughness at -320 F CVN Defects 62 '# None Tensile at Room Temperature UTS .2% YS ksi ksi 95 73 %E1 2" 35 %RA Defects 65 Some 11 Separations 461 Mag 8X Fracture Face of CVN Tested at -320°F Mag 1.5X Tensile Tested at Room Temperature FIGURE 3 : Impact, Tensile and Microstructur for Weld Deposited Using Swedish Avesta P6 Electrodes C ,03 Mn P 4.9 .015 Si Cr Ni .25 18.5 17.5 UTS ksi S_ .001 N Mo ,17 2.6 Toughness at -320 F CVN Defects 37' # None Tensile at Room Temperature %RA Defects ,2% YS %E1 ksi 2" 95 66 41 63 None Etchant NaOH Mag 500X Coring or 2nd Phase in Ferrite-Free Austenite 462 Mag 8X Fracture Face of CVN Tested at -320°F Mag 1.5X Tensile Tested at Room Temperature Etchant Chromic Mag 500X Grain Boundary Defects ir Ferrite-Free Weld Metal FIGURE 4 : Impact, Tensile and Microstructure for Weld Deposited Using U.S. - Armco Electrodes c Mn P S Si Cr Ni N 05 9.0 .007 .009 .22 18 16.2 .18 Toughness at -320 F CVN Defects 13'# 20% Micro Condition UTS ksi 93 Tensile at Room Temperature .2% YS ksi 71 %E1 2" 25 ;RA 26 Defects 15% Micro Condition 463 HONMETAlllCS 465 NONMETALLICS FOR MAGNET SYSTEMS M. B. Kasen Fracture and Deformation Division National Bureau of Standards Boulder, Colorado ABSTRACT An introduction to the nonmetallics program is presented. Topics discussed are: standard grades, radiation resistance, materials characterization, and technology transfer. Plans for FY 81 are presented. 467 The nonmetal 1 ics program has continued to respond to the needs of magnetic fusion energy (MFE) designers and fabricators by working with industry to provide commercial sources of standard grades of insulating laminates, emphasizing those grades that are most widely used in critical parts of superconducting magnets. The Annual Report for FY 79 (NBSIR 80-1627) reported the establishment of specifications for cryogenic grades of glass-fabric reinforced bisphenol A epoxy industrial laminates identified as G-10CR and G-11CR. The initial mechanical, elastic, electrical, and thermal property characterization data were reported to 4 K. This work has continued during FY 80, concentrating on refining the specification to achieve a balance between minimum performance variability at 4 K and minimum cost to the purchaser. It was recognized at the outset that radiation resistance of the insulating laminates would be an important performance requirement. The G-11CR product was incorporated as an initial effort to achieve better radiation resistance than that of the more common and lower cost G-10 type laminates. Available data had indicated that the aromatic amine cure of the G-ll product should be beneficial. However, tests performed in the Oak Ridge reactor failed to confirm this expectation. Because this left the G-11CR, a fairly expensive material, with no overall performance advantage over the less expensive G-10CR, the G-11CR product was dropped from the list of standard laminate grades. G-10CR is, therefore, the sole standard cryogenic grade at this time. Reception of this product by the cryogenic industry has been good, and it is presently being manufactured by three major U.S. laminating firms 469 As originally conceived, the standard grades were to provide the basis for systematic material development in response to fabricators' needs, in addition to providing baseline materials having predictable cryogenic properties. Thus, a variant of the basic G-l OCR product incorporating a lightweight glass cloth has been developed by one of the manufacturers. This variant, designated G-10CR-L, was required to increase the number of reinforcement layers in a thin wall thermal standoff for a Lawrence Livermore Laboratory magnet design. Magnet fabricators also require a uniaxial glass-reinforced laminate product and would find it advantageous to have available a glass-mat- reinforced epoxy laminate having proven performance at 4 K. The uniaxial product would primarily serve in the fabrication of thermal standoff support straps, and the mat product would be, in many instances, a lower cost substitute for G-10CR, having the added advantage of isotropic properties in the plane of the laminate. Substantial progress has been made in working with industry to establish standard specifications for such materials. Consideration has also been given to the possibility of establishing a laminate coding system that would permit the designer to specify the least expensive commercial material required to meet his needs without having to be intimately familiar with all of the parameters affecting insulating and structural laminate performance. This is long-range effort which, if successful, would encompass both standard cryogenic grades and the proprietary products of various manufacturers. These considerations are further expanded in a publication "Standardization Of Composites for Cryogenic Service" reproduced in this section of the report. 470 An example of industry interaction and participation is given by the discussions between NBS representatives and two industrial laminators during the NBS-DoE Vail Workshop that addressed refinement of the G-10CR specification. A copy of the Vail presentation by J. R. Benzinger entitled "The Manufacture of Radiation Resistant Laminates" is reproduced with the permission of Mr. Benzinger to provide background on the capability and limitations of industry in meeting the performance requirements of insulators for MFE magnets. The materials characterization portion of the nonmetallics effort was continued in FY 80 with additional studies of the effects of cryogenic temperatures on the thermal conductivity of G-10CR and G-11CR laminates. An effort to measure conductivity of the epoxy matrices separately was only partly successful owing to inhomogeneity in the bulk resin. Such data would be of value in establishing models for predicting the conductivity of laminates of differing construction. A glass-fabric-reinforced industrial laminate of European manufacture that had previously been evaluated at CERN was included in the NBS thermal conductivity program to provide a comparison of results between the two laboratories. For reasons not yet clarified, the NBS measurements indicated about a 45% higher conductivity at room temperature than did the experiment at CERN. It appears that the differences were due to experimental discrepancies between the laboratories rather than to specimen variabilities. Work to define the source of the differences is continuing. 471 The effect of temperature on the failure mode of insulating laminates must be understood if accurate predictions of the life-cycle performance of magnets are to be forthcoming. The document "Fracture Analysis of Compression Damage in Glass Fabric/Epoxy , G-10CR, 295 K to 4 K" included in this report reflects an initial effort in this direction. Compression failure was selected for study, since such stresses predominate in operating magnets. The high anisotropy of composite laminates creates problems in measuring elastic constants by conventional strain gage methods. The effort to develop dynamic methods of obtaining such measurements has therefore been continued during FY 80. Such methods are attractive because the small specimen size facilitates the evaluation of anisotropic effects. But the necessary conditions for the dynamic moduli to be the technical equivalent of engineering moduli obtained by conventional techniques has not been established. The progress made during the past year in clarifying the equivalency is documented in the contribution entitled "Static and Dynamic Moduli of Two Fiber-Reinforced Composites". The review paper entitled "Cryogenic Properties of Filamentary-Reinforced Composites: An Update" that appeared in the FY 79 report has been revised in part and submitted for publication to the journal Cryogenics . The major change was addition of a section on the radiation effects on composites, a copy of which is included in the present report. Two miscellaneous contributions were made to the literature during FY80. A letter to the editor of the journal Composites comments on a contribution by K. Kadotani relating to the effect of cryogenic temperatures 472 on the performance of composite materials. Work at NBS has indicated that the authors' conclusions could lead to serious misjudgments in materials selection if they are misinterpreted. We also replied to an invitation from the American Society for Metals to contribute to their forthcoming issue "Technology Forecast '81'. Copies of both these documents appear in this report. We anticipate that the nonmetallics program for FY 81 will continue to emphasize the areas of materials standardization, materials research and characterization, and technology transfer. The standardization effort will pursue the introduction of standard cryogneic grades of glass mat and uniaxial glass-reinforced epoxy laminates. The characteri- zation program will screen the performance of a G-10CR product from a third producer as well as that of a lightweight glass variant of this product. The performance of a polyimide-matrix variant of G-10CR will also be screened, because this laminate has shown substantially better radiation resistance than the epoxy grades. The research efforts will involve a cooperative program with industry to produce a cycloal iphatic epoxy resin variant of the basic bisphenol A G-10CR for future studies of radiation resistance. In-house research will attempt to resolve the noted discrepancy in laminate thermal conductivity and will continue to assess the validity of the dynamic method for determination of laminate elastic constants. Work will also begin on a system for precise measurement of viscoelastic effects in laminates as a function of temperature. 473 NONMETALLIC MATERIALS STANDARDIZATION 475 STANDARDIZING NONMETALLIC COMPOSITE MATERIALS FOR CRYOGENIC APPLICATIONS M. B. Kasen Fracture and Deformation Division National Bureau of Standards Boulder, Colorado ABSTRACT The designer of superconducting magnet systems is confronted with a difficult task in specifying the nonmetallic insulating and structural support materials required for proper system functioning. He finds it necessary to select from a profusion of products identified by trade names, most of which were developed for other than cryogenic use and for which reliable cryogenic performance data are nonexistent. This paper describes the ongoing efforts by NBS and industry to provide a com- mercial supply of standardized cryogenic grades of the most widely used insulating laminates having well -characterized cryogenic performance. These laminates will also provide the basis for a systematic materials development program in response to user needs. The elements required for a more generalized laminate coding system incorporating all materials of interest to magnet designers are also considered. 477 INTRODUCTION The technology of composite materials is presently at the state that metals technology was before being systematized by the establish- ment of industry standards. It is difficult to imagine working in metals technology in the United States without the ability to associate performance with designations such as 304L for stainless steel or 2024-T4 for aluminum. It would be equally difficult for individuals in other countries to perform effectively without similar national metal and alloy coding systems that transcend proprietary designations, but this is the situation presently confronting those working in the com- posite materials area, in the United States and elsewhere. Up to the present time, the problem has not been of major concern because the consumption of high-performance composites has been rela- tively low. However, consumption is expected to increase greatly as escalating energy costs place premiums on lightweight, fuel -efficient transportation systems and the development of alternative energy sources It will become increasingly difficult to meet the future needs of com- posite technology with proprietary systems or with laminates fabricated from the few formulations published in the literature. The need for composite materials standards is therefore common to many industries. As reflected in the forum chaired by Evans at the 1979 ICMC/CEC conference, the need of the cryogenic industry is par- ticularly acute. This industry is already at the point where many of the nonmetallic materials needs of magnetic fusion energy (MFE), magneto- hydrodynamic (MHD) , and rotating cryogenic machinery systems cannot be adequately met by existing commercial products. 479 A primary problem is the inability to associate reliably the cryo- genic performance required for diverse applications with the profusion of products identified by trade names, most of which were developed for other purposes. This paper outlines the current thinking of the author on this overall problem and describes the efforts currently underway to meet the immediate industrial cryogenic requirements. The author hopes that these considerations may be combined with similar considerations based on European experience to contribute to the eventual establishment of workable international standards for nonmetallic materials. The author will concentrate on nonmetallic composite materials, reflecting the current need for such materials in the fabrication of the cryogenic portions of new energy systems. It is recognized, however, that unreinforced polymers pose a similar standardization problem. IMMEDIATE CONCERNS The immediate task is to assist designers and fabricators in the selection of dependable insulating and structural nonmetallic laminates for MFE and MHD systems presently being built in the United States. The high costs of these systems and the substantial performance demands made on the laminates require that the selected materials be well character- ized at cryogenic temperatures and the variability in performance be minimized. Because the present systems will provide experience that can be extrapolated to more sophisticated, larger systems, the selected materials must be commercially available over a period of years. Fur- thermore, they must lend themselves to a systematic materials develop- ment program in response to industry needs. These boundary conditions 480 imply that control over the constitutive elements, manufacturing pro- cesses, and performance characterization will be necessary. Current construction will necessarily draw on state of the art, commercially available nonmetallic laminate materials developed for other than cryogenic service. In the past, many proprietary products have been used. However, this approach is inefficient for demanding applications since it requires that each group responsible for a design conduct an expensive in-house materials characterization program. It must also be assumed that the selected proprietary product will not be substantially altered in response to market forces outside the cryogenic area. The alternative of purchasing laminates to specified cryogenic performance criteria either requires the manufacturers to develop cryo- genic quality-control procedures or assumes that room-temperature data can be confidently extrapolated into the cryogenic range. Unfortu- nately, the latter concept is rudimentary for most conventional prop- erties and nonexistent for special -performance requirements under ir- radiation at cryogenic temperatures. A third alternative is to establish uniform industry specifications for manufacturing a series of well -characterized nonmetallic laminates having the greatest application to cryogenic technology. This approach recognizes that many of the inexpensive commercial laminates developed for room-temperature applications are useful cryogenic materials, whose primary deficiencies are the lack of a cryogenic data base and an excessive variability when used in a temperature region for which they were not intended. Establishment and characterization of uniform speci- fication laminates addresses these deficiencies and provides designers with materials having greatly improved cryogenic reliability at minimum 481 cost. Furthermore, the uniform specifications assure that the critical material factors affecting performance will be controlled by the cryo- genic industry, providing the basis for a systematic materials develop- ment program that reflects the needs of the industry. We have been following this approach in working with the United States laminating industry to develop a series of standard, high- pressure industrial laminates for cryogenic service. Both the industry and the coordinating standards group, the National Electrical Manu- facturers Association (NEMA) have been most cooperative. Thus far, two grades of glass-fabric-reinforced epoxy laminates have been established, characterized, and placed in commercial production. These materials are designated G-10CR and G-11CR to distinguish them from the conventional NEMA G-10 and G-ll products. Performance data for these cryogenic 2 grades have been published in the literature. Acceptance by the cryo- genic industry in the United States has been gratifying. Work is con- tinuing at NBS in cooperation with industry to develop additional standard specifications for other laminate systems of interest to cryo- genics, possibly including random-mat and uniaxial glass-reinforced epoxy-matrix laminates. These developments are meeting the immediate needs of industry while providing the basis for systematic alteration in the performance of standard materials in response to developing industrial needs. The concept of the overall program is outlined by the flow chart of Fig. 1. Built-in iteration steps provide the open-ended type of program required to be responsive to continually developing needs. The sequence depicted for development of the CR grades of G-10 and G-ll illustrates the pro- gress up to the present time. The characterization data indicate that these products will perform very well in most applications. 482 Studies at the Oak Ridge National Laboratory, however, have shown that the mechanical properties of these laminates are severely degraded by 4-K neutron and gamma irradiation at fluences comparable to those 3 4 expected in a functioning MFE reactor. ' This deficiency is being addressed by a subprogram to develop laminates having improved radiation resistance. The first step will be to determine if radiation resistance will be significantly improved by replacing the epoxy matrix of the G-10CR laminate with a polyimide matrix. It appears to the author that such a program would also entail an iterative cycle like the one illustrated in Fig. 2. Because the cryo- genic radiation resistance of a polymer is likely to be a strong func- 5 tion of the exact molecular structure, an essential part of the cycle is a correlation of molecular-level damage with radiation level and property change. Only in this way will the required scientific basis be established. The program envisions the development of methods for preliminary screening of 4-K radiation resistance in the laboratory prior to confirmatory testing in a reactor. It is evident that con- siderable new ground will have to be broken in solving the radiation problems in nonmetallic materials. A major deficiency in the approach thus far discussed is the ab- sence of an industry standards group having the responsibility for coordinating, updating, and disseminating information on developed standards. The NEMA organization is performing the disseminating function for the CR-grade laminates through their Industrial Laminates Subcommittee, but NEMA interest is confined to high-pressure, bulk laminates. An organization responsive to the broader interests of the composite materials industry is required. 483 LONG-RANGE STANDARDIZATION NEEDS The above approach lacks the comprehensiveness required to assure orderly overall composite materials development and implementation in the long run. The author agrees with Evans that the most effective way of alleviating the present chaotic state of nonmetallic materials tech- nology would be to introduce a numerical classification system based upon key elements of a material that affect performance. Here, the cryogenic industry has much in common with other industries exploiting composite technology. There seems no a priori reason why such an approach would not be feasible for composite laminates, when one recognizes that the signifi- cant engineering properties of a laminate at any temperature are largely defined by relatively few parameters. For example, the intrinsic me- chanical, elastic, thermal, and electrical properties of a laminate should be defined within relatively narrow limits by a numerical classi- fication system that establishes the type and flexibility of the matrix, the type and configuration of the reinforcement, and the reinforcement volume fraction. At the yery least, it should be possible to establish meaningful lower bounds on such properties. Special properties, such as radiation resistance, could continue to be associated with controlled- specification products that could be incorporated into the basic coding system. It is beyond the scope of this paper to suggest a definitive system, but it is worthwhile to consider the general form that such a coding or classification system might take. One possibility, paral- leling metals practice, is illustrated in Fig. 3. Here, the first 484 letter(s) and number define the matrix and reinforcement class, iden- tifying the composite as being in a general category, such as epoxy- glass or polyimide-graphite. A following letter defines the rein- forcement configuration, e.g., El U would define a uniaxially reinforced, epoxy-glass composite. The next two digits define the nominal fiber volume fraction. A ± 3% range would be sufficiently precise and would be compatible with existing manufacturing practice. An E1U60 coding thus defines a 60 ± 3% volume fraction uniaxial glass-epoxy laminate. The performance of such a laminate could be quite definitively estab- lished by adding general information on the degree of matrix flexibility and specific information on the type of reinforcement. Thus, a desig- nation E1U60U-E would not only define a 60% volume fraction uniaxial E- glass reinforced, fully reacted epoxy-matrix laminate but would simul- taneously establish meaningful intrinsic performance bounds on the laminate performance. Extrinsic factors related to laminate quality, such as void content and interply bond integrity, will affect perform- ance, but such factors must be independently assessed for all laminates. The fabric used to reinforce laminates is frequently unbalanced, usually having a lower fiber content in the fill direction of the weave than in the warp direction. For example, a type of 7628 glass weave commonly used in the United States has a fill/warp ratio of 32/42. In the scheme of Fig. 3, information on this unbalance is provided by a two-digit number defining the fill/warp fraction, or 76 in the cited case. This type of system offers flexibility in establishing coding at the specific level required for the application of interest. The cryo- genic industry would probably be reasonably well served by the cate- gories listed in Table I. The list could be expanded or individual 485 designations could be made more or less specific as need arises. Some cryogenic data already exist for most of the materials on this list, although in no case are the data sufficient for establishment of sta- tistically significant lower performance bounds. A standardization scheme of this type would simplify materials selection and purchase; furthermore, it would minimize costs by pro- viding the maximum flexibility to the laminating industry in meeting the requirements . Consider, for example, the purchase of a G-10 type of laminate widely used in superconducting magnet construction in the United States. Under the present system, the cryogenic performance of this type of laminate purchased on the open market can vary by 30%, if produced by a reliable laminator, or more, if produced by marginal fabricators. Specifying the CR grade would solve the problem for critical applica- tions; however, the premium paid for this grade may not be justified for many applications. Here, the illustrated classification system provides a solution by allowing the purchaser to specify an E1F48U-E/76 laminate if the generic equivalent to the CR grade was desired or to specify an E1F40U-E/76 laminate as a minimum requirement if greater property vari- ability is acceptable. An E1F40U-E/XX specification would result in still further economy by removing the requirement for a specific fill/ warp ratio. SUMMARY The author has described the ongoing efforts in the United States to provide a commercial supply of specific control led-performance lamin- ates required to meet the short-term needs of the cryogenic industry. 486 This approach also provides the basis for a systematic materials de- velopment program required to meet future needs, but the current effort addresses only the most demanding usage of nonmetallic laminates in cryogenic construction. The overall long-range requirements of the cryogenic industry justify consideration of a more comprehensive system of nonmetallic laminate classification that would simplify materials se- lection while minimizing costs. An example of such a system is provided to illustrate its potential benefits to both the consumer and the pro- ducer. ACKNOWLEDGMENT This work was supported by the Office of Magnetic Fusion Energy, United States Department of Energy. REFERENCES 1. D. Evans, in Advances in Cryogenic Engineering , Vol. 26, Plenum Press, New York (1980), p. 83 (in publ meermg ication) 2. M. B. Kasen, G. R. MacDonald, D. H. Beekman, and R. E. Schramm, in Advances in Cryogenic Engineering , Vol. 26, Plenum Press, New York (1980), p. 235 (in publication). 3. R. H. Kernohan, R. R. Coltman, and C. J. Long, Radiation Effects on Insulators for Superconducting Magnets , 0RNL/TM-61 93, Oak Ridge National Laboratory, Oak Ridge, TN (1978). 4. R. H. Kernohan, R. R. Coltman, and C. J. Long, Radiation Effects on Organic Insulators for Superconducting Magnets , 0RNL/TM-7077, Oak Ridge National Laboratory, Oak Ridge, TN (1979). 5. B. S. Brown, Radiation Effects in Superconducting Fusion Magnet Materials, submitted to J. Nucl . Mater. (1980). 487 Table 1. Laminate Classifications of Primary Interest to Cryogenic Technology Laminate Category Suggested Coding Uniaxial Glass-Epoxy Glass Fabric-Epoxy Glass Mat-Epoxy Glass Fabric-Polyester Glass Mat-Polyester Glass-Polyester Pultrusion Cotton Fabric-Phenolic Glass Fabric-Melamine Uniaxial Aramid-Epoxy Aramid Fabric-Epoxy Uniaxial , High-Strength Graphite-Epoxy Fabric, High-Strength Graphite-Epoxy Uniaxial, Medium-Modulus Graphite-Epoxy Uniaxial, High-Modulus Graphite-Epoxy Uniaxial , 5.6-mil Boron-Epoxy E1U55U-E* E1U55U-S2 E1U55U-S E1F48U-E/76* E1F40U-E/76 E1F40U-E/XX E1M35U-E* PE1F45U-E/76 PE1F40U-E/76 PE1F40U-E/XX PE1M35U-E PE1F45U-E/11(P)* PH5XXU M1F45U-E/76 M1F40U-E/XX E2U60U-49 E2F50U-49/94 E3U60U-HS E3F60U-HS/100 E3U60U-M E3U60U-HM E4U50U-56 For illustration purposes only. Classifications meriting development of special cryogenic grades 488 List of Figures Fig. 1. Flow chart depicting a proposed approach to meeting immediate needs of nonmetallic laminate development for the MFE, MHD, and rotating cryogenic machinery industries. Dashed lines depict the present state of development of the G-l OCR and G-11CR grades of industrial laminates. Fig. 2. Suggested flow chart for development of radiation-resistant nonmetallic laminate insulators. Fig. 3. Possible coding system for classifying nonmetallic composite laminates . List of Tables Table 1. Laminate Classifications of Primary Interest to Cryogenic Technology. 489 Q. CO a. (0 3 •a c * c o +* o (0 (0 k. 0) *■> CO 5 (0 ■♦-» c o *"; E a 0) 4) k. 3 C CT o 0) 4M QC (0 O *o k. (0 3 l_ •t-* ^m\* 3 Li. O) c *D 'E O 4> (0 a o o (/> +* c < c ^ o * 3 o O 0) a 0) qc a (0 .2 o <*- 4) Q .. (0 o c o a> c E 4) k. a> a> a 3 CT a> X oc (0 ■M o C •*•* a> a k. k. 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CD s_ CO +-> ITS 1 1 1 1 m 1 — LU Q.I UJ s: CJ> a_ a_ CO =5 x: c CO , — +-> =5 3 CD J_ 4-> ,— -a c O C" =3 o a> cd o> -a CD 2: CD S_ CD F to CO O +-> 1 +->+->+-> r— ai (O CO co s: •1- E •I- 00 •!- -i- ( > CO ra ra 1 JC zj J= 1 -c E <;. m r— J= Q-'i- CL-C Q. 1 • o CJ3 tD en rO T3 ro cnroiD 14- r ■> 1 1 *r- S_ CD i- -r- S- • +-> C UJ OO DC O 21 O X CDLf) CD (1) 1 i i 1 1 1 a: UJ oo s: =c 21 OO CO DC LO CD CO O Q. E o o ra 4-> CD o •1— 4- to co ra S- o 4- CD +-> CO CO CD -o o o CD CO CO O D_ OO en ■1 — u_ CD ra 492 NONMETALLIC MATERIALS: RESEARCH AND CHARACTERIZATION 493 THERMAL CONDUCTIVITY OF G-IOCR AND G-11CR INDUSTRIAL LAMINATES AT LOW TEMPERATURES (PART III) J. G. Hust and R. Boscardin Thermophysical Properties Division Continuous Process Technology Programs National Bureau of Standards Boulder, Colorado ABSTRACT Thermal conductivity measurements from 2 to 300 K were made on the resins used in two cryogenic grades of fiberglass/epoxy composites, G-IOCR and G-11CR. Effects of specimen thickness and inhomogeneity of the bulk sample were studied. Also, the thermal conductivity of a European glass-fiber/epoxy composite was measured to assess the variation between measurements at CERN and at NBS. The accuracy of the NBS measurements was checked by measurements on a standard reference material. 495 INTRODUCTION Extensive low- temperature electrical, thermal, and mechanical properties characterization of cryogenic grades of G-10 and G-ll fiberglass epoxy com- posites are being performed at the National Bureau of Standards. These special laminate grades, referred to as G-10CR and G-11CR, meet NEMA Grades G-10 and G-ll, MIL-P-18177C, and federal specification LP-509. In a previous report (Part I), 5 - thermal conductivity data were presented for G-10CR and G-11CR at temperatures from 2 to 300 K for heat flow in both the normal and warp directions. These data indicated that the two laminates, pro- duced by a single manufacturer (manufacturer A), had the same conductivity within about 10%. Also, it was found that the warp direction conductivity was about 20% higher than that in the normal direction. In Part II of these reports, * data were presented showing the differ- ences in thermal conductivity on specimens obtained from another manufacturer (B) and the effect of large strain. Preliminary data were also presented for thermal conductivity of the resins used in G-10CR and G-11CR. Further measure- ments on these resins and the conclusions are presented here. In addition, this report (Part III) contains data on a European glass fiber/epoxy composite (Vetronite* 432. 81). 3 The purpose of these measure- ments is to clarify the cause of apparently large differences between NBS and CERN data on materials expected to be quite similar. The Vetronite was mea- sured at both laboratories for direct comparison on a single sheet of material. The NBS measurements were also performed on various thicknesses to illustrate the absence of significant systematic errors due to interfaces and radial heat losses. tradenames are used to describe adequately the materials studied. No endorse- ment of a particular product by NBS is implied. 497 MATERIAL AND SPECIMEN CHARACTERIZATION G-IOCR resin is a heat-activated amine- catalyzed bisphenol A, solid-type, epoxy resin. G-11CR resin is an aromatic amine bisphenol A liquid- type epoxy resin. The G-IOCR resin has an uncured density of 1.10 + 0.01 g/cm 3 and the G-11CR resin has an uncured density of 1.04 + 0.01 g/cm 3 . The resin specimens were cut from blocks supplied by the manufacturer. The dimensions, masses, and calculated densities of these specimens are given in Table 1. The resin density calculated from data for the composites (report- ed earlier) is 1.25 + 0.01 g/cm 3 for both resins. Note that the measured densities of the G-IOCR and G-11CR resins are 1.180 g/cm 3 and 1.227 g/cm 3 , respectively. The specimens designated as G-IOCR resin (a) and (b) are the two obtained by cutting the original specimen in two, to study the effect of thick- ness and inhomogeneity. A sheet of the Vetronite material, measured by K. D. Peterson of CERN, was obtained and twelve square specimens were cut from it. These were fabricated by cutting two 1.9 cm wide strips from the supplied sheet. Each of the two strips were further sectioned into six equal pieces. The twelve resulting pieces were machined to their final dimensions. Their dimensions, masses, and densities are given in Table 1. Measurements were performed on selected groups of these specimens to obtain thermal conductivity values as a function of thickness. EXPERIMENTAL PROCEDURE AND DATA ANALYSIS The thermal conductivity of the test specimens is determined using a ther- mal conductivity integral apparatus. This apparatus is described in detail by Hust and Arvidson. 4 The estimated uncertainty of the results is 5%, unless 498 otherwise noted, as verified by measurements on NBS standard reference material SRM 735. The data range is normally 2 to 300 K and the data are normally spaced approximately logarithmically within this range. The measured thermal conduc- tivity integrals are fit by a special least squares technique (see Hust and Arvidson [1]) to yield the parameters for the following equation: n i K(T) = z A.[ln(T + l)] 1 Eq. 1 i = l n RESULTS AND DISCUSSION To facilitate future analysis of composite data, measurements were per- formed on specimens fabricated from the pure resins used in G-10CR and G-11CR (see Table 1). These specimens were prepared from cured pieces supplied by the manufacturer. Twenty-one data points between 2 and 304 K were obtained on the original G-10CR resin specimen and sixteen data points were obtained from 4 to 304 K on the G-11CR resin specimen. The results, essentially the same with- in the precision of the data (+4%) for both resins, are illustrated in Figure 1. These values, however, are more uncertain and imprecise than usual because the specimen thicknesses were too large for optimum accuracy. Because of this potential problem with these first sets of data and because of the possibility of thermal radiation effects in these nearly transparent specimens, it was decided to section the G-10CR resin specimen to produce two thinner specimens. The thinnest specimen, about 1/3 the original thickness (see Table 1), was measured at 18 temperatures from 4 to 300 K. The results, appreciably lower (25%) than the original G-10CR resin data, are illustrated in Figure 1. This large difference was not understood and led to a series of measurements near 499 room- temperature. The results of the room temperature measurements are given in Table 2. The conclusion reached from these measurements is that there is a signifi- cant inhomogeneity in the G-10CR resin specimen. The source of this inhomogen- eity is uncertain but it is likely to be found in the original resin-mixing and block-curing procedure. The results of the measurements on specimens #1, (a), (b) and (a + b) suggest that the thermal conductivity is greater near the middle of the specimen, since both the results of the original measurement on the whole specimen and the later measurement on the reassembled specimen (a + b) were higher than (a) or (b) measured separately. It is also noted that the reassembled specimen resulted in an intermediate value, which is consistent with the above conclusion, since the center section was partly removed by cutting. The average thermal conductivity of nonhomogeneous specimens can be direction dependent, i.e., the results can depend on whether the hot face is on the low-conductivity side or the high-conductivity side. This effect is not quantified, because the orientation of the specimen in the chamber was not recorded in all cases. Whether the G-11CR resin specimen contains similar inhomogeneities is unknown because it was not sectioned. Since the percentage difference in thermal conductivity between the small and large G-10CR resin specimens is essentially independent of temperature, the thermal radiation component to total heat transfer is small compared with the conduction component. One Vetronite specimen was measured near room temperature for comparison with the CERN data. In addition, as a check on the possibility of systematic errors due to interface resistance, radial heat losses, and the calibration constant of the probe, measurements were performed on randomly selected pieces 500 in the form of stacks (2 thick, 4 thick, and 12 thick). The results of the measurements are given in Table 3. These data reveal only a very small systematic deviation as a function of thickness. The constancy of the results over the extremes of thickness used here illustrate that boundary layer resistance (at the smallest thickness) and radial heat losses (at the largest thickness) are not significant experimental errors. In all cases, these results are about 45% higher than the results from CERN. This confirms that the differences are truly experimental in nature and not specimen differences. The possibility of systematic errors in the NBS thermocouple calibrations and other instrumentation calibrations has been ruled out by measurements on SRM 735. 4 501 REFERENCES 1. J. G. Hust, "Thermal Conductivity of G-10CR and G-11CR Industrial Lami- nates at Low Temperatures," in Materials Studies for Magnetic Fusion Energy Applications at Low Temperatures - H^ NBSIR 79-1609, F. R. Fickett and R. P. Reed, Eds., National Bureau of Standards, Boulder, CO (1979), pp. 417-436. 2. J. G. Hust, "Thermal Conductivity of G-10CR and G-11CR Industrial Lami- nates at Low Temperatures (Part II)," in Materials Studies for Magnetic Fusion Energy Applications at Low Temperatures - III , NBSIR 80-1027, R. P. Reed, Ed., National Bureau of Standards, Boulder, CO (1980), pp. 395-404. 3. K. Dahlerup-Petersen and A. Perrot, "Properties of Organic Composite Materials at Cryogenic Temperatures: CERN ISR-BOM/79-39." (November 1979). 4. J. G. Hust and J. Arvidson, "Thermal Conductivity of Glass Fiber/EPOXY Composite Support Bands for Cryogenic Dewars," unpublished report 275.03-78-2, National Bureau of Standards, Boulder, CO. 502 TABLE 1 DIMENSIONS, VOLUMES, WEIGHTS, AND DENSITIES OF THE THERMAL CONDUCTIVITY SPECIMENS DIMENSIONS (cm) SPECIMEN X a X b X a VOLUME WEIGHT DENSITY 1 2 3 3 (cm ) (g) 3 (g/cm ) G-IOCR RESIN #1 ■- 1.915 1.913 1.781 6.524 7.698 1.180 (a) 1.915 0.652 1.781 2.224 2.615 1.176 (b) 1.915 1.041 1.781 3.550 4.175 1.176 (a+b) 1.915 1.693 1.781 5.774 6.790 1.176 G-11CR RESIN 1.915 1.915 1.915 7.023 8.620 1.227 VETRONITE 432.81 (Al) 1.786 0.199 1.770 0.629 1.200 1.91 (single pieces) (A2) 1.780 0.198 1.768 0.623 1.194 1.92 (A3) 1.783 0.196 1.768 0.618 1.189 1.92 (A4) 1.780 0.196 1.765 0.616 1.186 1.92 (A5) 1.780 0.198 1.765 0.622 1,190 1.91 (A6) 1.780 0.199 1.765 0.625 1.193 1.91 (Bl) 1.780 0.201 1.768 0.632 1.202 1.90 (B2) 1.775 0.200 1.768 0.628 1.200 1.91 (B3) 1.778 0.197 1.765 0.618 1.189 1.92 (B4) 1.775 0.196 1.763 0.613 1.183 1.93 (B5) 1.770 0.198 1.760 0.617 1.180 1.91 (B6) 1.768 0.200 1.753 0.620 1.176 1.90 (stacked groups) A3 & B3 1.780 0.393 1.766 1.235 2.328 1.92 A3, A5, Bl, a rid B3 1.780 0.792 1.766 2.490 4.770 1.91 all 12 spec imens 1.778 2.378 1.765 7.462 14.282 1.91 a - Values of X\ and X3 for the stacked groups are averages of the dimensions of the individual pieces. b - Values of X2 for the stacked groups are the sums of the thicknesses of the individual pieces. 503 TABLE 2 MEASURED THERMAL CONDUCTIVITY VALUES AT 280 K OF THE G-10CR RESIN SPECIMENS [SPECIMEN #1 REFERS TO THE ORIGINAL SPECIMEN, (a) AND (b) ARE THE SPECIMENS OBTAINED AFTER SECTIONING SPECIMEN #1] Specimen Identification and Thickness (cm) Thermal Conductivity (W m" 1 K" 1 ) G-10CR Resin #1 1.913 G-10CR Resin (a) 0.652 G-10CR Annealed 15 h P 100°C (a) 0.630 G-10CR Resin (b) 1.041 G-10CR Annealed 15 h 100°C (b) 1.026 After Anneal ina (a + b) 1.656 0.296 0.230 + 0.008 a 0.232 0.246 0.238 0.252 G-11CR Resin G-11CR Annealed 15 h @ 100°C 1.915 1.915 0.282 + 0.005 b 0.278 a Range of 4 measured values b Range of 2 measured values 504 TABLE 3 THERMAL CONDUCTIVITY VALUES FOR VETRONITE 432.81 NORMAL TO THE FIBERS AS A FUNCTION OF THICKNESS AT 284 K. Specimen Identification Thermal Conductivity (see Table 1) W m" 1 K" 1 B3 0.447 B3, A3 0.450 B3, A3, Bl, A5 0.451 All 12 Specimens 0.480 505 ■11 Z. ■*■* O c c <0 U c o u E c S_ 'I- -c a; l— i- t _>i - ^ui-mui 'AiiAiionaNOO iviau3h± ai s- =3 CD 50C j FRACTURE ANALYSIS OF COMPRESSION DAMAGE IN GLASS FABRIC/EPOXY, G-IOCR, 295 K to 4 K Betty A. Beck Fracture and Deformation Division National Bureau of Standards Boulder, Colorado 80303 ABSTRACT This report describes scanning electron and optical microscope fracto- graphic and modeling studies of G-IOCR compression fracture at room and cryo- genic temperatures. The primary objective of the investigation has been to assess the effect of cryogenic temperature on the failure behavior of these composite materials. Macromechanical analysis in terms of gross composite fracture has been designated as primary. Micromechanical (or constituent) analysis has been specified as secondary. This identification has been used for descriptive and comparative purposes rather than to indicate the level of importance in the fracture process. Results of the fractography have shown that the failure events included three kinds of primary fracture patterns which, to varying degrees, are temperature-related. These fractures have been identified as: 1) a 55° slant mode; 2) a chisel or V-notch mode; and, 3) a complex or mixed mode. The influence of constituent properties and behavior have been integral considerations in assessing the fracture mechanism and in analytical interpretations of the fracture behavior. As in other materials, the damage zone in these composites is the result of either discrete or inter- active constitutent failure processes (or both) occurring in a highly strained region of the specimen. Fiber microcracking or fracture, matrix yielding or cracking, and interface failure are examples of such events. Nucleation and propagation of fracture, as in other fiber-reinforced composites, is a function of microstructure, microdeformation, and the structural interaction of the constituents. Compressive strength data variability has been graphically correlated with primary and secondary fracture characteristics in order to understand the low temperature effects better. 507 Contents Page Abstract A. Introduction 511 B. Previous Studies 512 1 . Introduction 512 2. Constituent Properties and Behavior 515 (a) matrix resin (b) glass reinforcement (c) reinforcement/matrix interface 3. Composite Room Temperature Behavior and Failure 518 (a) general (b) toughness (c) compression 4. Studies of Glass Reinforced Nonmetallics at Low Temperature 526 5. In Summary and Concluding Remarks 531 C. Low Temperature Effects on Fracture Modes of G-10CR Tested in Compression at 295 K, 76 K, and 4 K. 532 1 . Introduction 532 2. Material Specifications, G-10CR 534 3. Test Procedures 535 4. Fracture Analysis 535 5. Results and Discussion 540 (a) general (b) fractographic observations and analysis (c) conclusions 509 Contents (continued) Pa^e 6. Recommendations for Future Studies 551 7. References 552 8. Tables 557 9. Illustrations 561 510 A. Introduction In general, composite failure is complex because of the interrelated behavior and properties of the constituents, i.e., the matrix, the reinforce- ment, the reinforcement/matrix interface as well as the variability resulting from the fabrication process. Fracture analysis experience has shown that no single failure mechanism stands out as a primary cause of composite fracture. Rather, there seems to be a sequence of failure events that lead to total composite failure (1-7). Studies have shown that the predominant primary mechanism of composite fracture as well as the sequence of constituent failure events varies signi- ficantly for different composite types (6). Included with this failure pattern variability are the energy dissipative mechanisms at fracture initiation and growth which also differ from one material system to another. This lack of failure behavior continuity and the inherent complexity of fabric-reinforced nonmetallics emphasize the need for systematically building a theory of failure based in part and initially, on detailed observations of fractured specimens. Numerous investigators have studied and described failure and fracture behavior in both discontinuous and uniaxial ly continuous glass-fiber reinforced nonmetallic composite materials (1-7). For the most part analytical interpre- tations, methods, and models, where they exist, have been developed around some theoretical or experimental aspect of a composite's room temperature behavior. Low temperature studies have had a characterization orientation primarily. Fabric-reinforced nonmetallic composites possess a yery complex geometry and behavior, and as such, are more difficult to describe both qualitatively and analytically. These materials have, to date, eluded conventional analysis 511 and modeling. Most of the difficulty seems to be in the inherent complexity of modeling the reinforcement. Because of the dearth of analytical studies for these materials at both room and low temperatures, there is essentially no analytical base on which to build a theory of failure. In this study and through exploratory data analysis, a detailed descrip- tion of primary and secondary fracture patterns has been developed to elucidate basic mechanisms of fracture and to determine temperature-related modes of failure. The results of this investigation will provide a foundation for developing failure analysis techniques based on correlating fractographic features with fracture properties and stress analysis at fracture. B. Previous Studies 1 . Introduction A summary of information from the published literature has been included in this report. In order to identify, interpret, and analyze failure patterns in the G-10CR specimens, the literature review has provided a basis for developing a fractographic and an analytic frame of reference. Some of the published results have been modified or extended in part to this composite system which involved a different reinforcement geometry and type and loading condition. In utilizing and synthesizing the information on the behavior of matrices, reinforcements, and to some extent, interfaces, a rationale has been established for the sytematic analysis of failure in the G-l OCR specimens. The state of the present study would be signficantly less complete without the direction provided by the earlier studies. Because there are so few published quanti- tative or even qualitative analyses about failure behavior in fabric reinforced nonmetall ics, this approach has seemed necessary. 512 Most investigators agree, that, to formulate a better analytical under- standing of the fracture process in all composite materials, knowledge, not only of the predominent mode of failure but of the sequence of failure events and their probability of occurrence is a primary consideration in the difficult task of reconstructing the fracture process from initiation through propagation to final failure (1-3, 6-14). Similarly there is a consensus about the impor- tance of fractographic analysis in advancing a comprehensive theory of failure and failure analysis techniques (14). There seems to be a lack of a well-ordered knowledge base that can be used to develop low-temperature strength theories of failure as well as pre- dictions for the onset of failure. This has been particularly true for the glass-fabric-reinforced nonmetallics. There have been only a few studies of the low-temperature behavior of glass-reinforced nonmetallic composites. Constituent behavior, the matrix resins, and glass reinforcement have received somewhat more attention. The structural dependence of Young's modulus, Poisson's ratio, and strength, for example, has been shown to be rather small in epoxy resins. On the other hand, their fracture energy shows significant room- and low-temperature dependence on structure (15). For the composites made up of these materials, the low-temperature changes in strength, moduli, and other properties are incompletely understood structurally. Correlation of changes in static mechanical properties of glass-reinforced-epoxy composites with macrolevel fracture characteristics at cryogenic temperatures is almost nonexistent A prevailing idea among the many studies is the need to relate fracture modes to structure and states of stress in the development of valid failure criteria. Morgan et al . , for example, point out that the network structure and void characteristics of epoxy resins control their mechanical response (16). 513 The network structure, which varies with chemical composition and curing conditions, imparts rigidity to the resin. Voids act as stress concentrators and, as such, degrade mechanical response by adversely affecting the effective shear strength. The structure and properties of the resin are significant to the load transfer function of the matrix. Low temperatures can adversely affect this load-transfer mechanism. Fracture analysis of all composite systems is a function of the deterior- ation of the unit structure. Glass-reinforced materials are no exception. The constituents impart a major effect to the breakdown under an applied load. Rosen, as well as other investigators, has shown that there is no unique mechanism that will adequately describe failure in all fiber-reinforced non- metallics (1-3, 4, 5, 9). The individual properties and characteristics of the constituents (the matrix, reinforcement, and interface) and their inter- actions in the process of degradation and failure yield a complex, coupled behavior in total composite fracture that is difficult to trace to its origins. The contributions of the variety of interacting mechanisms are not easily discernible in the complexity of general composite failure (10). To a great extent, for example, composite strength, which is related to how effectively the stress can be transmitted from one fiber to the next, is a function of the individual constituents (17). Kelly, in discussing uniaxial ly reinforced composites subjected to an applied load parallel to the fiber, emphasizes a modulus mismatch that causes displacements in fiber and matrix to differ and leads to stress concentrations proportional to the difference in moduli. Shear stresses and strains resulting from the stress concentrations are the mechanism for load distribution between matrix and fiber via the interface (17, 18). The major part of the load is 514 carried by the reinforcement if the matrix modulus is such that a matrix displacement is greater than that of the fiber reinforcement. Matrix and interface characteristics (void content and shear properties) are primary to an effective load-transfer mechanism (19). Cracking in the resin matrix can result in a poorer distribution of the load or abrasion of the glass. Most investigators agree that the load is transferred by shear stresses at the fiber/matrix interface and that the rate of load transfer is limited by the strength of either the interface, the matrix, or the reinforce- ment, whichever is least (3, 17). A strong interface (low void and crack content) provides a more effective load-transfer effect than a weak interface (high void and crack content that degrade bonding and shear strength) (17). The load-transfer effectiveness of the matrix is adversely affected by increased resin brittleness, increased Young's modulus, and the existence of residual stresses in the composite because of resin contraction and shrinkage during cure (3, 20). The matrix resin protects the fibers and provides a shear rigidity between fibers and lamina, strength and stiffness to the reinforcement, and a coupling effect to the constituents in response to tensile, compressive, bending, and shear loads. Additionally, it sustains a high level of strain concentration that occurs as the result of the high-modulus, brittle fibers in the glass-reinforced-epoxy composites. 2. Constituent Properties and Behavior (a) Matrix Resin In general, polymer failure occurs by excessive deformation, crazing or cracking, and tensile or shear fracture. Cross-linked polymers fail by cracking 515 and fracture (21). Structurally, there is evidence to indicate that stress concentrations can lead to cracking at the cross-links if the cross-link length is too short for plastic flow. Shear fracture for polymers can occur on either a single plane or simultaneously along two planes of maximum resolved shear stress lying at 45° to the longitudinal axis of testing. Fracture results when deformation, due to a concentration of large shear strains, is restricted to the one or two planes of shear failure. Strength and elasticity of epoxy resins show little or no dependence on cross-link distance for temperatures in the range above 50 K to less than 20 K. Energy at fracture, however, for these same resins, is strongly dependent on cross-link distance, increasing as the distance increases (15). The response mechanism of a polymer to an applied mechanical stress is a function of the ability of the molecules to move past or over one another. At cryogenic temperatures there is an increase in the rigidity of these materials owing to an increased rigidity in the molecular structure. Epoxy resins, in general, exhibit room-temperature glasslike textures typical of brittle materials that have characteristic low-energy-absorbing pro- perties, low viscosity, and low shrinkage when cured (5, 22, 23). They don't form bubbles, which minimizes buildup of internal stresses in the cured resin. (b) Glass Reinforcement The high-strength, stiff glass fibers, in the form of a rigid molecular network with no mechanism for ductility, are subject to a statistical distri- bution of flaws and defects characteristic of brittle materials. As such, fiber failure is by random microcracking and fracture inherent to brittle 516 failure. An accumulation of fiber breaks is considered by many investigators to be a major mechanism in total composite failure. A strong, stiff matrix contributes to a greater probability of a statistical failure mode of the glass fibers (1-3, 7, 8, 11). (c) Reinforcement/Matrix Interface Debonding of the fiber from the matrix and delamination of one lamina from another are interface failure modes. Most composite studies have shown that the fiber/matrix interface is primary to load distribution from matrix to fibers. Strength, mode of failure, and work at fracture are a function of the behavior of the interface. Interface shear strength is a function of the density of voids formed during fabrication. High void content results in low shear strength and leads to a weak interface that can be split by stresses set up ahead of an advancing crack (17). There is some evidence to indicate that fiber/matrix debonding and matrix cracking can interact to reduce load concentrations and result in an increased composite stress. Studies have also indicated that, where there is an increase in composite strain at fracture, it may be due to deterioration of the fiber/matrix interface. In synthesizing fracture information for glass-reinforced-plastic com- posite systems, constituent characteristics and their highly coupled nature and behavior relative to the failure process must be integrated with the analysis. The degradation and failure mechanisms may occur independently, simultaneously, or consecutively in total failure. Energy-dissipative mechanisms and levels thereof are integrally important in sequence analyses of failure events (4, 24, 25). Crack branching, for 517 example, in fiber-reinforced-composites extends the size of the fracture surface and results in a large dissipation of energy. Debonding, for brittle fibers enclosed by a semi brittle matrix resin, seems to occur first along an interface, and the debonding work done is stored in the fiber as elastic strain energy. When a debonding sequence is completed, the fiber pulls out of the matrix. This occurs because the load on the fiber, in order to split it from the matrix, is higher than the load for pullout. As a result, the energy dissipated at fiber pullout is higher than the work done in debonding (4). 3. Composite Room Temperature Behavior and Failure (a) General Primary, considerations in composite fracture and fractographic analyses include the identification and description of: (1) the predominant failure mode of the composite; (2) the failure modes of the constituents; (3) their probability of occurrence in the failure of a specific composite material under a well-defined loading condition; and, (4) a probable sequence of failure events from initiation to propagation and final fracture. Collins et al . , for example, conclude that, in determining which mode dominates failure behavior, fractographic analysis is a function of which model is assumed (11). Their study identifies and summarizes two approaches: (1) a cumulative fracture propagation model and (2) a noncumulative model. These models have also been described and expanded upon in other investigations (2, 9). The cumulative model assumes that the total applied load is borne by the fibers. The matrix acts only as a mechanism for redistribution of the stresses when a fiber breaks locally. This model is based on a statistical strength distribution due to inherent defects characteristic of the brittle glass 518 fibers. Total composite fracture occurs after enough fiber fractures have resulted in a weakened cross section so that it can no longer support the applied load. This model is valid if an assumption is made either that (1) the load is distributed uniformly over the remaining fibers after one breaks or (2) surviving fibers are subjected to static and dynamic load concentrations. The noncumulative mode identifies composite failure as a function of crack propagation through the matrix, interface, and reinforcement. As a result, the failure loads are considerably below those which would cause a significant number of fibers to break. According to these investigators, fractographic analysis will be approached differently depending on which mode is assumed. The cumulative mode implies, for example, that interface debonding of fiber and matrix and fiber pullout result in an ineffective fiber length, which carries little of a tensile load in a failed area. The noncumulative mode, on the other hand, implies that interface debonding and fiber pullout determine an optimum path of crack propagation through the matrix, including crack branching at the interface. Zweben has determined that such factors as fiber/matrix debonding, interface cracking, matrix plasticity and cracking, and the three-dimensional nature of the material act as energy-absorbing mechanisms that tend to reduce load concentrations, effectively increasing composite stress. On the other hand, a lack of uniformity in the spacing as well as cracks in the fibers tends to increase stress concentrations and decrease composite failure stress (2, 9, 11 , 24). Early studies of the tensile failure of a uniaxially reinforced matrix identified a simple failure model based on a uniform strain throughout the composite. As a result, fracture occurred at the failure strain of the fibers alone (2,7). More recently most investigators 519 have postulated a nonuniform strain distribution, which has directed consideration of the influence of flaws in the brittle reinforcement on composite failure (4, 5, 7, 9). Rosen, and others, have developed a rather comprehensive description of failure in fiber-reinforced composites (2, 6, 7). This model assumes that because of flaws inherent to the brittle fibers, initial fracture will most likely occur in the fibers causing a local perturbation of the stress field. The way the load is redistributed near the broken fiber ends and the resulting stress concentrations may result in: (1) separation along an interface (i.e., fiber debonding or lamina-to-lamina delamination) ; (2) fracture of nearby fibers (since they are subjected to a greater stress intensity than those at some distance); (3) propagation of transverse cracks; (4) immediate composite fracture; or (5) no further fracture. Because the strengths of the brittle fibers are not uniform, they will break randomly throughout the material as the applied stress is increased. The fractured fibers are then reloaded to some level before being pulled out of the matrix. There is, however, a critical length at which the fibers will no longer support the load. As fractures accumulate, the end result is a number of fibers of ineffective lengths interacting to produce a weakened cross section of the composite, which leads to total failure. Since the shortened fibers can no longer effectively transmit the increasing load, composite fracture occurs as the maximum shear strength of the matrix is exceeded. A strong, stiff matrix, the nonuniform strength of the fibers, and a statistical accumulation of flaw-related fiber fractures are the controlling factors in the failure characterization of this model. Because the fibers do not have uniform strengths, strong fibers will carry an increasingly greater load as weaker fibers break. Additionally, the average strength of some of the individual fibers may 520 be greater than the coupled strength of fiber bundles, which can signi- ficantly affect the fracture sequence. Because of the lack of uniformity in the glass-fiber strengths, investigators have utilized the Weibull distribution as statistically descriptive of the strength variability (4, 5, 7, 9, 26). There is a distinction between the failure processes (including stress distribution and level) that control the onset or initiation of failure and those that control the fast propagation of a crack. The former is more difficult to reconstruct fractographically and has thus received less study. Because of the variety of micromechanisms of failure occurring and interacting, failure mechanisms at the onset of damage must be distinguished from those close to fracture in terms of stress levels and states. Rosen feels that, in developing a failure criterion, it's initially important to understand internal irregularities in the state of stress (7). Tetelman has determined that composite fracture initiates either by: (1) fiber fracture, (2) interface shear fracture (debonding or delamination) , or (3) matrix tensile or shear fracture, depending on which of the above occurs first at a lower value of applied strain. Final fracture may be the result of some or all of the above mechanisms (3). Mull in has analyzed energy-absorbing mechanisms in high-modulus fiber composites (24). His analysis indicates that, at a fiber fracture site, strain energy will be increased in adjacent fibers by a stress-transfer mech- anism acting through the interface and the matrix. The localized redistri- bution of load can produce either matrix cracking perpendicular to the fiber or debonding at the interface. The probability of the former mechanism occurring reflects a strong interface bond; the latter, a weak one. Debonding can 521 result in a localized sliding of the fiber in the interface socket accompanied by fiber fracture and pullout. Beaumont has identified postdebond sliding, rather than fiber fracture and pullout, as a major toughening mechanism in fiber-reinforced composites (27). Vinson and Chow associate fiber pullout with the flaws in the brittle fibers that lead to nonuniform strengths (4). They assume that the matrix has cracked completely. As a result, fibers in the cracked regions become stressed more than the remaining ones, which are firmly fixed in the matrix. Zweben cites a study of composite failure that associates fracture sequence with a propagating stress wave that fractures the fibers destroying their load-carrying capability (25). The matrix, as a result, cannot sustain the applied load level without the strengthening effects of the reinforcement. This leads to total composite failure with most of the energy being dissipated by the matrix. Additional models for describing failure initiation and propagation with associated energy dissipative mechanisms have been proposed. One of these indicates initiation is the result of debonding brought about by increasing fiber fracture strength or decreasing matrix or interface shear strength (3). Voids introduced during fabrication degrade interfacial shear strength and significantly influence debonding. Deformations that develop during initial impairment of a fiber-reinforced material can lead to a tensile fracture of the reinforcement, which is most probably due to bending followed by a final shear mode of composite failure (5, 27). Materials with a high void content can fail by crack extension parallel to the fibers, whereas those with low void content and high fracture toughness have a characteristic shear final - failure mode. Residual stresses and strains introduced during fabrication are coo 22 the result of differential thermal expansion characteristics between reinforce- ment and matrix. The existence of these stresses can complicate the analysis of failure initiation. If, as some investigations have assumed, failure first occurs in the brittle matrix and the crack propagates toward the fibers, it can propagate across the fibers if the interface is strongly bonded and con- tinue to brittle failure. If, however, the propagating crack approaches a weak interface, debonding of fiber/matrix interface occurs. Debonded fibers then fracture as a function of the defect distribution in the imperfect fibers. Stress concentrations ahead of an advancing crack can cause the interface to fail. When a propagating crack reaches a fiber, it is diverted along the cracked interface and blunted. In order for the crack to continue to propa- gate, it must be reinitiated by increasing the load. Some investigators have emphasized that this sequence of events is the primary toughening mechanism in fiber-reinforced nonmetallics (4, 5, 27). (b) Toughness Toughness in a metal is the ability to absorb energy during plastic deformation without fracturing. In a static tensile test this energy can be measured by the area under a stress-strain curve, which represents the work required to fracture a specimen. In glass-reinforced epoxy composites there is no direct toughness analogy with a metallic system because of the plastic nature of the metal. If, how- ever, toughness is considered to be the ability to redistribute stresses in the presence of a notch to reduce the effects of high stress concentrations by dissipating the energy, then the work required to fracture a composite specimen should represent toughness properties analogous to a metal. Toughness in a 523 glass-fiber-reinforced plastic seems to be a function of the interfaces and their ability to deflect cracks propagating normal to them (4, 5, 12, 17, 22, 27). If there is a strong interfacial bond, there is a high probability that stress concentrations set up by failure of one fiber will nucleate cracks in nearby fibers leading to progressive fiber failure. If the bond is weak, interfacial shear strength will be degraded and stress concentrations ahead of a crack advancing normal to the fibers can split the interface. As the crack reaches a flawed boundary, it will be blunted or diverted parallel to the interface. For the crack to continue to propagate, it must be reinitiated by a sufficient stress buildup as the applied load is increased. If this process can be repeated as the cracking spreads, large amounts of energy can be absorbed Some investigators have indicated that toughness relates to the energy dissipated as fiber/matrix debonding occurs. The toughness mechanism becomes a sliding action between fiber and socket, accompanied by friction effects and eventual fiber pullout (24, 27). Fiber-reinforced nonmetallics are considered to be tough and not notch sensitive. (c) Compression Compressive stresses are the result of either an externally applied load or are due to shrinkage stresses developed during cooling because of differ- ential thermal expansion characteristics between matrix and fiber. Rosen and others have developed models for uniaxial continuous-filament composites subjected to an externally applied compression load parallel to the fibers (1, 7, 28-30). The model postulates that failure occurs by microbuckl ing of the fibers on an elastic foundation. The analysis identifies two fiber buckling ;24 modes, a shear mode and extension mode, each descriptive of matrix deformation. For this model the matrix shear modulus controls composite compression strength. Buckling can be prevented by a stiff matrix and strong bonding at the interface (5, 7, 20, 31, 32). There seems to be little evidence to show whether, under compressive loading, local buckling precedes and causes crack propagation or follows crack initiation and propagation (5). There is a study cited in Broutman that investigated a linear relation between interlaminar shear strength and compressive strength for flat filament-wound panels. That analysis showed compression strength as a function of resin properties and either high interlaminar or horizontal shear strengths (5) . For his investigation on compressive fracture Chaplin has expanded the development of the shear mode to a shear instability in a volume of material rather than failure in a plane as a function of a resolved shear stress. The shear instability produces significant shear deformation in the matrix and bending of the reinforcement (32). Externally applied compressive loads parallel to the fibers in a unidirectional fiber-reinforced composite can result in failure by: fiber microbuckl ing in a shear mode; tensile or compressive fiber fracture; shear failure of the matrix; interfacial failure due to delamination or debonding or both; or compression-induced bending accompanied by secondary tensile and shear stresses (4, 5, 7, 31-33). Fiber pullout and fracture are failure events in the debonding sequence. A major characteristic of composite shear failure is lack of damage to the fibers in the direction of shear. A compressive stress will not propagate internal cracks nor will it produce crazing, a failure phenomenon characteristic of many matrix resins. Rather, a compressive stress normal to a crack tends to close the crack. 525 Secondary tensile and shear stresses, however, will propagate internal cracks and are significant to the resulting fracture behavior. Ryder and Black, in discussing types of compression testing, indicate that there are several possible compression failure modes in fiber-reinforced composites. Each has a different ultimate strength. For each testing pro- cedure, failure occurs at the lowest strength with a corresponding failure mode. The range of possibilities suggests that compressive strength for these materials is not a precise term, but rather, is one of definition (34). Extensive work on static compressive behavior of glass-epoxy composites has shown that test results and the failure mode are sensitive to the method of gripping (misalignment causing the specimen to bend), type of reinforcement, and fiber volume fraction. Three major compressive failure modes have been identified: (1) fiber buckling, (2) shear, and (3) interfacial, with the shear mode predominating (7, 28, 33). Total brittle compression fracture of a composite specimen results in separation of the specimen into two or more i parts with little or no measurable inelastic deformation. If deformations develop initially, they can lead to either (1) tensile fracture of the reinforcement due to bending with final fracture of the composite in a shear mode or (2) shear displacement in the matrix, which can cause stress concentrations and crack extension along the fibers if voids are present in a contiguous area. 4. Studies of Glass Reinforced Nonmetallics at Low Temperature Fiber and matrix characteristics have a significant influence upon the mechanics of deformation and fracture of fiber reinforced nonmetallics at both room and cryogenic temperatures. As such, temperature-related behavior and 526 properties of the constituents are important in the analysis of temperature effects on failure modes for these composites. Polymer behavior is a function of a molecular chain response mechanism, which can experience a temperature-related change as temperatures are decreased to those representing cryogenic environments. Such change can have strong effects on material properties and fracture behavior. At room temperature the molecular chain is in thermal motion around an equilibrium position. At this temperature the molecular response to an external stress (that is, the coiling and uncoiling of the chains and their moving over and past one another) is easier because there is no fixed molecular position. In this configuration, the polymer chain provides an energy-absorbing mechanism for stress waves transmitted through the material. As the temperature is decreased, the thermal energy of the molecules decreases and the elastic mobility of the polymer chain slows down. As a result, the chain becomes more rigid, the energy- absorbing alternatives become more restricted, and attractive forces become more effective in the resistance of materials to deformation. Under an applied load the decreasing ability to dissipate energy through the molecular structure leads to an increased resin rigidity until the cohesive strength of the polymer has been exceeded, ultimately leading to fracture as a dissipative mechanism. This molecular rigidity results in a stronger, stiffer material at low temperatures because the polymer chains are unable to slide over one another effectively (3, 15, 16, 23. 35). For epoxy resins, there are cross-linking effects in the molecular structure as well as the effects of molecular anistropy of the binding forces owing to the directional nature of the primary covalent bonds and secondary 527 Van der Waals forces, which act along and across the molecular chains. The cross-linked structure imparts some degree of rigidity to the polymer. Epoxy resins do not seem to be subject to changes in structure with decreasing temperature. They exhibit elastic-plastic stress-strain behavior at room temperature but have nearly linear stress-strain characteristics at low temperature. Strength (tensile, compressive, and shear) and modulus increase with decreasing temperature. Fracture strain, a function of molecu- lar cross-linking, approximates 2% at 77 K and 4 K. Softer and Molho tested a Bisphenol-A epoxy resin at cryogenic tempera- tures (77 K and 20 K) . Their results indicate that, in comparison with room- temperature properties, (1) there is an increase in strength and modulus in tension and compression; (2) there is a change in elastic-plastic behavior at room temperature to essentially linear elastic behavior at 77 K and 20 K; and (3) there is a decrease in fracture strain and resin toughness (as measured by the area under the stress-strain curve) with decreasing temperatures. The rate of decrease for (2) and (3) moderates between 77 K and 4 K. The resin exhibited a maximum notch toughness (the capacity of the resin system to sus- tain loads in spite of defects in the matrix) value at 77 K and decreased to below the room temperature value at 20 K. Their assessment of the fracture- surface examination of this particular resin specimen indicated a smooth, almost glasslike texture typical of low-energy, absorbing, brittle materials (23) Hartwig's low-temperature analysis indicates, in general, that the chemi- cal structure of epoxy resin chains, including cross-link density, does not have a marked effect on elastic mechanical parameters. Poisson's ratio seems to be entirely independent of structure. Energy at fracture, on the other hand, decreases with higher cross-link density at cryogenic as well as higher 52; temperatures (15). Other investigators have noted a relation between the degree of cross-linking and void density and their detrimental effects on matrix and interface shear strength. This effect, however, was not investi- gated below room temperature, probably because low-temperature factors were not a part of the study (21). Kadotani determined in his low- temperature studies that the lower the glass transition temperature, T , of a resin, the larger its tensile elonga- tion at 77 K. The larger the elongation at 77 K, the higher the strength and elongation of the composite at 77 K (36). Glass materials used for reinforcement are brittle at room temperature and below. Their rigid silica network structure and lack of molecular order is such that there is no inherent structural mechanism for ductile behavior characteristics. Because of their brittle nature, strengths are not uniform and their failure modes at room and cryogenic temperatures reflect these brittle characteristics. Glass-fabric-reinforced epoxy composites exhibit significant increases in tensile, compressive and shear strength, tensile modulus, Poisson's ratio, and tensile strain to failure with decreasing temperature compared with room- temperature values (20, 36, 37). There is considerable evidence to indicate that: 1. A stiffer resin holds the reinforcement more firmly fixed for loading parallel to the principal direction of fabric-reinforced composite. As a result, there is a greater resistance of the fibers to dis- placement or deformation. 2. Fiber-to-fiber progressive fracture leading to final composite fracture provides an enhanced energy-absorbing mechanism, the 529 characteristics of which reflect the sum of the energies of the individual fibers. As a result, more energy can be absorbed at room and low temperatures. 3. The increase in strength at low temperatures is the result of inter- face failure and crack blunting, a progressive mechanism that imparts toughness to the material. 4. The high strain of the composite at fracture at the lower tempera- tures is a direct function of the low-temperature, high-fracture strain characterisitics of the glass reinforcement. 5. The increase in composite fracture strain may reflect a deteriora- tion of the fiber/matrix interface. 6. The glass-reinforced plastics are the only composites that exhibit increased elongation at 77 K. 7. Tear-type fractures seem to be more evident at low temperature for epoxy matrix materials. For glass-fabric-reinforced epoxy composites stressed parallel to the principal fiber direction, there is a transition in the stress-strain (20, 22, 36, 37) behavior below room temperature that results in both an initial and secondary modulus characteristic, termed a "knee" of the stress-strain curve. Figure 1 is a typical illustration of this effect. Initially the knee seems to be due to interface failure and matrix cracking. There is some evidence to indicate that the decrease in secondary modulus with decreasing temperature, as illustrated in Figure 1, reflects a decrease in the shear strength of the resin matrix or fiber/matrix interface. 530 5. In Summary and Concluding Remarks For glass-fabric-reinforced nonmetallics this review has demonstrated, not only a dearth of adequate test and analysis techniques for predicting strengths and determining mechanisms of deformation and failure, but also, a need for such analyses and tests. The "rule of mixtures" was an early and relatively simple model for relating fiber and matrix strength and stiffness to estimates of composite strength. Implicit in the model was the assumption that, under an applied load, there was a constant strain throughout the composite. The interpreta- tion based on that assumption indicated that the first fiber to fracture led to crack propagation to failure for the composite with fracture occurring at the failure strain of the fibers. More recently some investigators (1, 5, 26) have pointed out that the "rule of mixtures" analysis is inadequate for composites reinforced with wery brittle materials such as glass fibers. Failure for these materials has statistical features that reflect the nonuniform strength distribution of the fibers' inherent flaws and imperfections. Cryogenic testing of these composite materials has been limited to charac- terization studies which show strength and stiffness increases with decreasing temperatures. These effects seem to be more a reflection of changes in the reinforcement than the matrix. There are, however, significant structural changes going on in the resin at cryogenic temperatures that may or may not be important in trying to deduce how a glass fabric reinforced epoxy will fail. Since the fracture strengths of the individual fibers or fiber bundles - covers a range of values, determination of an appropriate stress for describing the failure stress of the reinforcement - and ultimately the composite - is a 531 complex problem still to be addressed. Adding to this complexity are important considerations relative to the sequence of constituent failure events, their interaction, and contributions to total composite failure. The initial part of this report has been directed toward a comprehensive synthesis of behavior and properties information as well as failure models derived from published investigations of glass-reinforced plastic composites. The analytical and descriptive material and models included herein have served as an excellent basis for the low-temperature fractographic analysis of G-l OCR that follows in this report. C. Low-Temperature Effects on Fracture Modes of G-l OCR Tested in Compression at 295 K, 76 K, and 4 K 1 . Introduction To provide an initial data base for engineering applications a cryogenic- temperature-characterization program was conducted at the National Bureau of Standards (1978-79) for the cryogenic-grade glass-fabric/epoxy composite, G-l OCR (37). These experimental studies included electrical, thermal, and mechanical property characterizations of CR-grade materials over the 295 K to 4 K temperature range. The objective of the present study, described in this paper, is to assess the effect of cryogenic temperatures on the failure behavior of the G-l OCR specimens tested in compression under the program described above. The fol- lowing specific activities were directed toward achieving this objective: 1. Examination and analysis of fracture surfaces of G-10CR specimens previously tested in compression at 295 K, 76 K, and 4 K with frac- tographic methods. 532 2. Identification and characterization of primary composite failure and secondary constituent failure as well as the interactive effects of the constituents that led to gross composite failure. 3. Correlation of failure morphology with room and cryogenic compres- sive strength characteristics. As a result, the development of an interpretation of the experimental results, which shows compressive strength and data variability increasing with decreasing temperature. 4. Comprehensive review of the literature of existing failure analysis information and models that can be used as a frame of reference for identifying and elucidating G-10CR compression failure behavior at room and cryogenic temperatures. Generically failure was considered to be the occurrence of any discernible discontinuity in the response of the composite and constituent material to a mechanically applied external compressive load. Gross composite failure was determined to be primary, constituent failure to be secondary for comparative and descriptive purposes rather than to designate order of importance. Several kinds of failure mechanisms within this classification were used to identify failure in the G-l OCR specimens. Breakdown of the unit structure leading to failure of the constituents or the bond between them; reduction in the load-carrying capacity due to localized damage or separation; significant loss of strength due to crack growth; or total composite fracture were considered to be identifiable failure processes in this analysis. This study involved an after-the-fact critical examination and analysis of the fractured specimens. In the analysis it was important to determine: 1. What constituted failure. 2. What was the predominant composite failure mode. 533 3. What were the modes of constituent failure and the probable occur- rence of each. 4. What was the probable sequence of failure events in total composite failure. 5. How do the above (i.e., 1-4) correlate with decreasing temperature, 295 K to 4 K. 2. Material Specifications, G-10CR Specimen configuration with adhesively bonded end caps is shown in Figure 2 (37). The end caps were designed for and used to eliminate premature specimen failure and to assure uniform load transfer as well as good vertical alignment of specimen ends to avoid column buckling. Figure 2 also shows diagrams of laminate configurations. The reinforcement is a silane-finish, E-glass, 7628 woven glass fabric, plain weave with interlaced warp (length) threads (43/in) and fill (width) threads (32/in). This type of reinforcement results in a high-volume fraction of fibers in the resin with a correspondingly high strength and modulus and different strength properties in directions parallel and perpendicular to the warp. The matrix is a heat-activated, amine-catalyzed bisphenol A solid-type epoxy resin, a cross-linked molecular system. Resin content ranges from 32 to 35%. The composite specimen is made up of layers of similar parallel lamina, as illustrated in Figure 2. Material for the specimens was obtained from two producers. Tables 1 and 2 contain typical property values for the E-glass reinforcement and an epoxy resin (9, 22, 23). 534 3. Test Procedures Static compression testing was done by applying a compressive load in the plane of the cloth. The compression test fixture is shown in Figure 3. The testing environment included air at 295 K, liquid nitrogen at 76 K, and helium at 4 K. Tests were performed parallel to both warp and fill, and compressive strength values were determined for both directions at the above temperatures. 4. Fracture Analysis Fractured specimens of G-10CR tested in compression at 295 K, 76 K, and 4 K were subjected to a critical examination using optical microscope and scanning electron microscope (SEM) fractography to assess damage and determine failure modes, both primary and secondary, for each specimen. Failure was defined in the preceding part of this report. Both macrolevel and microlevel observations and analysis were considered to be important to the results of the study. The former assumed the composite to be a homogeneous solid and an averaging process was utilized to assess damage and failure. The latter was based on the assumption that the material is a heterogeneous solid and that constituent (fiber, matrix, and interface) characteristics are integral con- siderations in composite fracture. For both primary and secondary failure mode analyses, location and direc- tion of failure and failure-surface morphology were stages considered important to fracture characterization. Determining individual constitutive failure events as well as their relation to total composite fracture were subsets in the major focus of the study. Although there seemed to be a significant degree of consistency in the G-10CR fracture morphology, it still became an 535 extremely difficult task to establish the path of fracture and trace it to initiation. In general, damage modes of importance to this analysis were considered to be cohesive failure of the matrix and reinforcement and adhesive failure of the fiber/matrix interface. Constituent damage and failure types, which combine in a variety of ways as they affect the response of a material to the applied compression load, are consistent with the numerous published studies summarized in an earlier part of this report. Those failure modes identified as important to this study were: 1. Matrix (epoxy) cracking and fracture 2. Reinforcement (glass fabric) microcracking and fracture 3. Fiber/matrix interface, debonding, delamination, and fiber pullout. Primary compressive failure modes were either (1) an approximate 55° slant type; (2) a V-notch or chisel type, or (3) a mixed mode or complex type. Types (1) and (2) have been diagrammed and are illustrated in Figure 4. Figure 5 correlates each of these three modes [i.e., (a), (b) or (c)] with temperature for warp and fill directions. Tables 3 and 4 are temperature-related summaries of significant factors in the characterization and fracture analysis of these materials. These tables represent the same information, except that one is a summary for the warp direction and the other, for the fill direction. In the first part of the table, failure modes at each testing temperature have been correlated with the lowest and highest values of compressive strength. The second part is a summary of the percentages of each primary failure mode 536 for each testing temperature. Several factors became apparent in organizing the summaries. There is more variability in the primary failure mode in the warp direction than in the fill direction. In the fill direction, the single- plane ^55° slant-type mode was the predominant mode. This mode represented 100% of the lowest compression strength-vs. -temperature values in the fill direction, but only 33% of those values in the warp direction. Highest com- pressive strength-vs. -temperature values were identified as either a single- or a double-slant (shear-type) failure for the fill tests. The warp-test fracture modes, in this case, were not representative of any highest value mode except a complex type indicating considerable variability. The third part of the table identifies the fracture edge for the primary mode when that mode is either the ^55° slant or V-notch. This edge is a transverse crack across the specimen with some evidence of tearing at the lower temperatures (Figs. 4, 6-20). Part four of the table is a summary of constituent damage or failure-vs.- temperature. The next part of this report emphasizes in more detail some of the fractographic analysis that integrates with and documents certain aspects of the fracture interpretations. Property measurements determined during the initial characterization tests are graphically presented for inclusion in this part of the report in Figures 21 through 23 (35). These diagrams serve as a basis for analysis in the development of this failure-analysis study. Also, results determined from these graphs have directed construction of additional graphical comparisons considered to be significant and illustrated in Figures 24 through 26. Figure 1, as previously indicated, illustrates a typical tensile stress- strain diagram for some of the G-10CR specimens tested in the warp direction 537 at room temperature and 76 K (37). The diagram, even though it represents a typical tensile property, has been included with the report because it is representative of the stress-strain behavior for this material. The work required to fracture a specimen can be related to the area under the stress- strain curve and is an indication of the toughness of a material. Using this area as a measure, the work at fracture shows an increase when the temperature is lowered to 76 K. At 76 K there is a discontinuity in the configuration of the curve, considered to be a "knee" at about 300 MPa with an accompanying change in the modulus. Initial and secondary moduli reflect this transi- tion at low temperatures but not at room temperature. This type of behavior can be correlated with specific failure occurrences i.e., beginning matrix cracking at the knee and a degradation in interface shear strength reflected by the lower secondary modulus. Figure 21 has been included to provide a graphic comparison of the range of measured test values for compressive strength, Young's modulus, and tensile strain to failure for room temperature, 76 K, and 4 K in both warp and fill directions. The number of specimens for each test condition are shown in Tables 3 and 4. The bars at each temperature in Figure 21 are representative of the highest and lowest test value and thus illustrate the variability in the data for each of the properties shown. Figure 22 diagrams all measured values for Young's modulus vs. temperature and Figure 23 is a graph of all the measured data for compressive strength vs. temperature. The outer straight lines have no physical meaning other than to indicate upper and lower bounds for the data and to present a graphical comparison of the scatter in the data at each temperature; the center straight line connects the average value at 538 each temperature. The temperature-related vertical line delineates the calculated standard deviation for the range of values. It's interesting to note the significant increase in strengths for warp and fill directions compared with those at room temperature. Figure 24 illus- trates a graphical comparison for warp and fill tests of the low temperature- room temperature percentage increase in compressive strength, Young's modulus, and tensile strain to failure. Another interesting factor to be noted in Figure 23 is that the low temperature variability (i.e., the scatter in the data) is reversed for measured compressive strength values in the warp and fill directions. At 76 K the scatter is yery large for the warp direction and is significantly less in the fill direction. At 4 K the reverse is true. Variability (as expressed in the standard deviation and coefficient of variation for compressive strength, Young's modulus, and tensile strain to failure) has been addressed in Figures 25 and 26. These graphs summarize and compare variability as a function of temperature. The graph illustrating Young's modulus vs. temperature (Figure 22) was determined and presented in a manner similar to the one for compressive strength Earlier studies have indicated that the low-temperature related increase in composite strength and modulus primarily reflect similar changes in the glass reinforcement. Plotted values for Young's modulus in warp and fill directions indicate, as was also illustrated for compressive strength, that there is significantly greater scatter for the data in the warp direction. Variability at 4 K does not appear to be a significant factor for either warp or fill directions. 539 Optical and scanning electron microscopes were used to examine details of the fracture surfaces to identify and evaluate constituent failure. Figures 6 through 15 are typical photographs obtained with the optical microscope. Figures 16 through 20 are typical SEM photographs of fracture details for specimens tested in the warp direction at various temperatures. Significant details have been identified and delineated on each of the photographs. These illustrations can, in turn, be correlated with the summaries in Tables 3 and 4. 5. Results and Discussion (a) General As indicated in an earlier part of this report, previous studies have shown that the response of a glass-reinforced laminated nonmetallic tested in static compression is a function of the type, orientation, and volume fraction of the reinforcing material; matrix characteristics; precision of alignment of the test specimen; and direction of loading (e.g., load-path distribution as a function of the anisotropy of the material). A woven glass-reinforced epoxy- matrix composite subjected to an externally applied compressive load parallel to the fibers (either warp or fill) can undergo in-plane shear deformation (affecting both the lamina and the interlaminer areas), longitudinal and transverse bending of the reinforcement, and possible bending of the specimen. The matrix undergoes shear deformation. Bending of the fibers results in secondary tensile and shear stresses that may be highly significant to the failure mechanism. Published studies on the compressive strength of laminates have identi- fied shear and interface (debonding, delamination, and pullout) failure and fiber buckling as primary failure modes. The shear mode, however, has seemed 540 to predominate. A propagating shear mode of failure across a specimen can be the result of compression-induced bending or shear stresses (4, 5, 28, 31, 33). There are several constituent structurally related responses to the low- temperature environment that are significant to the behavior of glass-reinforced nonmetall ics . Failure modes of the composite are in some way related to interacting modecular and microstructural (constituent) events. An increased rigidity in the polymer chain degrades its ability to dissipate energy. Chemical structure, in terms of cross-linking and directional characteristics inherent in primary covalent bonds and secondary Van der Waals forces, characterizes a molecular anisotropy in the matrix resin, which can be significant to composite behavior. Hartwig has shown that for an epoxy resin: (1) Young's modulus and Poisson's ratio show little or no dependence on chemical structure at low temperature; (2) tensile strength shows some dependence on cross-link distance at temperatures below 76 K; and (3) the energy of fracture shows significant structure sensitivity at both normal and cryogenic temperatures. The energy of fracture is released locally and can lead to a significant temperature rise in the specimen. Hartwig points out that epoxy resins behave viscoelastically at or near the transition temperature and linearly elastic to fracture with an accompanying low fracture strain (approximately 2%) at cryogenic temperatures. Increased strength, brittleness, and resin rigidity characterize the matrix at the low temperatures (15, 35). Increased strength and fracture strain are the cryogenic effects on the glass reinforcement. The topography of a fracture surface is largely a function of the stress conditions at fracture. An analogy with metals would indicate that lower stress levels are generally identified with a lower level of energy being dissipated at fracture and a smooth, regular fracture surface. 541 (b) Fractographic Observations and Analyses With the above as an introduction, the following relates to the specific compression tests of G-10CR specimens at 295 K, 76 K, and 4 K and the fracto- graphic examination and analysis of those specimens that comprised this study. There were two significantly different testing conditions for the G-10CR specimens compression tested at each of the indicated temperatures. Specimens were oriented with the compression load parallel to warp threads for one series of tests and to the fill threads for the other group. (There was a third test orientation normal to the fabric reinforcement. These specimens however, were not included in this fractographic assessment.) The ratio of fill-to-warp threads is approximately 0.75. Measured values of Young's modulus, cu E, and compressive strength, a , for the various testing temperatures cor- related well with this ratio. As the summary in Tables 3 and 4 indicates a single-plane, slant-type primary fracture configuration occurred most frequently for both warp and fill tests at all temperatures but more frequently for the fill tests. This inves- tigation has identified that mode as an approximate 55° slant-type mode (^55°), see Figure 4. A second slant-type configuration occurred and was identified as two intersecting planes showing a "V" or chisel appearance. This configu- ration was designated as either a V-notch or chisel type (Figure 4). There were more V-notch failures for the warp tests than the fill tests. For the former, all occurred at 295 K and 76 K. At 295 K these failures generally occurred (1) within the gage section but near the face of one of the grips and (2) with considerable longitudinal delamination. At 76 K the fractures occurred (1) at the face of the grip, (2) with delamination, and (3) with considerable general damage and surface roughness. 542 There were more complex or mixed-mode fractures in the warp direction than in the fill direction and more evidence for this at 4 K. There was significantly more variability in the primary fracture mode for tests per- formed in the warp direction than in the fill direction, and this variability increased at the lower temperatures. High values of compressive strength for warp tests generally correlated with complex fracture modes at cryogenic temperatures; low values of compressive strength for these tests correlated with either ^55° slant type or V-notch at 295 K and 76 K and was mixed at 4 K. For the fill tests at all temperatures, the lowest values of compressive strength correlated 100% with the ^55° slant mode. Whereas high values of compressive strength at all temperatures correlated with either the ^55° slant or V-notch mode. For all tests, both warp and fill, damage and surface roughness generally increased at the lower temperatures. This damage was manifest in matrix cracking and greater transverse fiber damage. Individual fiber and fiber- bundle fracture, debonding, and pullout were characteristic of the latter. Increased damage and surface roughness at low temperatures has been correlated with energy or work to fracture at 76 K, a measure of which is the area under the stress-strain curve in Figure 1. Published studies have indicated in- creased work to fracture is a function of the interface and its ability to deflect or blunt a crack propagating normal to the reinforcement (4, 5, 12, 17, 22, 27). For the present study, the low-temperature significance of this effect, in part, can be associated with changes taking place at the interface, which probably intensify the mechanism for occurrence. Another characteristic that has been identified with primary fracture in both warp and fill directions for slant-type failures at all temperatures is a 543 major transverse crack across the specimen, which has been determined to be the primary fracture edge. This crack, illustrated in Figure 4, shows evi- dence of being initially a transverse crack in the matrix of the outer lamina (Figures 6-10 and 14). Fractography shows this edge is deflected \/ery slightly at fiber interfaces (for those fibers parallel to the applied load) but re- initiates in a transverse direction, propagating and deflecting intermittently across the specimen face. Reinitiation and propagation are such that the pattern of the fracture edge is primarily at right angles to the applied load. At cryogenic temperatures the fracture edge is similarly transverse but there is evidence of tearing (Figures 8 and 13). Several factors complicate fractographic analysis in general and the determination of an initiating failure mode in particular. Composite ani- sotropy contributes significantly to variability at the interfaces. As this study shows, there has been more interface damage variability transverse to the applied load than there has been parallel to it. In conjunction with this effect is the fact that load-path and stress distributions are a function of the inherent anisotropy of the composite material. Moreover, the determina- tion of an initiating failure mode is complicated by the differential stiffness, expansion and contraction characteristics, and subsequent residual curing stresses developed in fabrication. And, the load-bearing capabilities of matrix and reinforcement, although integral in forming the unit structure of the composite, represent a significantly important differential in the load- support effect. Thus, the reinforcement carries more of the load than the matrix because of its higher strength and modulus. The high volume fraction of the glass reinforcement represents another parameter that makes a signifi- cant contribution to the failure mode in compression. 544 Constituent behavior, properties, and interactions have received equal emphasis in this fractographic study to describe G-10CR fracture behavior. In examining the fracture surfaces, the analysis was based first on discrete damage events (i.e., matrix, fiber, and interface failure) and second on the constituent interaction, which determined a most probable sequence of failure events. The former was developed in considerable detail and has been typically illustrated in Figures 12 through 20. The latter is more complex and requires detailed techniques of analysis, which will direct efforts in future studies. It has received initial coverage here. (c) Conclusions When glass-fabric-reinforced systems such as the G-10CR materials are subjected to an externally applied compressive load parallel to the fibers at either room or cryogenic temperatures, the stiff epoxy resin holds the fibers firmly fixed, constraining deformation which leads to strong stress and strain concentrations (5, 31, 34). As in other composites under an applied stress, different displacements occur in the matrix resin and glass reinforcement owing to an inequality in the moduli (18). The stiffness differential results in shear strains and stresses, a primary mechanism for distributing the load between matrix and reinforcement. Because the matrix will deform more than the fibers and because of the difference in elastic moduli, the latter support most of the load at all levels. Previous studies (3-15, 18) have shown that an axial load, parallel to the fibers, is transmitted directly to the rein- forcement for continuous-filament composites and via shear stresses at the interface for discontinuous-fiber laminates. The epoxy matrix acts primarily as a mechanism for load transfer and stress redistribution. Although the G-10CR materials are considered to be continuous, in actuality, as with other continuous glass-fiber-reinforced composites, continuity characteristics can be disrupted by fabrication flaws or early fiber fracture at low levels of applied stress. The latter is the result of the lack of uniformity in glass strengths owing to defects inherent in all very brittle materials. The short- ened fibers, the weaker ones that fracture at low applied-stress levels, characterize a discontinuous effect in what is normally considered to be a continuous-filament-reinforced material. Stress perturbations and stress concentrations ultimately result in weakened zones in the composite that lead to total composite fracture'. An increase in strength and strain properties of the G-10CR specimens with decreasing temperatures seems to be more a reflection of changes taking place in the reinforcement (i.e., an increase in strength and modulus) than in the matrix. A fractographic assessment of the specimens has shown some rather specific consistencies as well as differences in the appearance of the glass fabric with decreasing temperature. First and most general, the evidence indicates that the reinforcment fractures by tension or bending. Second and more speci- fically, there is a brush-type fracture of fiber bundles at the crossovers and parallel to the applied compressive load for both warp and fill tests (with little or no debonding of individual fibers); see Figures 14 through 20. There are more of these fractures visible on the fracture surface at the higher temperatures than at 4 K. Alternatively, there are less visible brush fractures at low temperatures for tests performed in the fill direction. A significant number of the brush fractures are accompanied with delamination outermost of the fiber crossover area (Figure 24-26). To a lesser extent, some of the bundle fractures remain embedded (little or no delamination) with 546 only the brush ends showing (Figures 22 and 24). This effect is apparent at low temperatures when there is a transverse matrix crack contiguous to a transverse fiber bundle. For both warp and fill tests, fractographic evidence points to an elastic bending of the fibers with resulting fracture at the crossovers for those fibers parallel to the applied load. There is a charac- teristic consistency to this type of fracture mode at all temperatures. The only temperature-related difference seems to be that fewer of these are visible on the fracture surface at the lower temperatures, particularly for the fill tests. A low-temperature-related effect is that more matrix shows on the fracture surface at 4 K. The fractography has provided strong evidence that longitudinal bending fracture (i.e., for fibers parallel to the applied load) occurs as a result of stress concentrations at the fiber crossovers, where the dissipative mechanism with the highest probability for relieving the stress buildup is the bending fracture of the bundles. Published fracture studies reviewed for this report have shown no evidence of the brush fracture mechan- ism observed in the G-10CR specimens. For both warp and fill tests, examination of the fracture surfaces of the specimens showed a significantly greater degree of randomness for damage and fracture (individual fiber and fiber bundle, debond, and pullout) transverse to the applied load, at cryogenic temperatures in particular. Room-tempera- ture random effects were significantly less. At room temperature, evidence of shear failure was apparent for transverse fibers in that damage was primarily limited to delamination with little or no damage to the fibers in the shear direction. The undulating nature of the transverse fibers was not destroyed at the fracture edge (Figures 16 and 17). 547 Finally, as in other published studies, fiber cracking and fracture appears to be due, in part, to cracks at the surface as well as to intrinsic defects. As the fibers fractured in the composite, the increasing load was redistributed by the matrix via the interface to the shortened, weaker as well as the remaining, stronger, long fibers. Initial fractographic analysis indicates that this process continued until a significant number of fibers were reduced to ineffective lengths, causing a weakened region in the com- posite which, in part, led to total composite fracture. Matrix fracture was characteristically either tensile or shear, with the latter predominating. Since the matrix resins serve primarily as a mechanism for load transfer and stress redistribution, there are certain temperature-related factors that directly or indirectly determine the effectiveness of those functions. On the molecular level, for example, the energy dissipative mechanisms inherent in the mobility of the polymer chain were more constrained with decreasing temperatures. Changes that occurred in the resin at low temperatures affected the load transfer function, a parameter that is sensitive to residual stresses developed during fabrication, to resin brittleness, and to a change in resin modulus. The ductility of the resin matrix decreased rapidly with cryogenic cooling. Resin strength and modulus increased; strain capability was reduced; toughness decreased; and, stress-strain behavior changed from viscoelastic to linearly elastic at 4 K. At cryogenic temperatures there was evidence that the epoxy resin matrix of the G-10CR materials cracked before the glass reinforcement (Figures 1, 9b, 10b). Although there has been extensive consideration in the published liter- ature about which constituent (the matrix or the reinforcement), fails first, this study indicates that the most probable first failure event for these 548 testing conditions was the fracture of the outer layer of the epoxy matrix. As the matrix crack in the outer lamina developed and propagated transverse to the applied load and across the specimen, a less effective distribution of the load resulted. Fibers contiguous to the matrix crack became more highly stressed than those firmly embedded, and some bending of the fibers occurred in the vicinity of the cracked matrix. This process continued as the load in- creased and contiguous regions were weakened. Several factors lend support to matrix cracking preceding significant fiber fracture: First, epoxy strength is lower than that of the glass reinforcement. Second, the nonuniform strengths characteristic of glass fibers signify a randomness in their total fracture pattern for which there is little or no evidence in the constituent failure modes parallel to the applied load at any temperature. There was, however, a distinct randomness in failures transverse to the applied load with a more irregular effect becoming apparent at the lower temperatures. Finally, typical tensile stress-strain curves at 295 K and 76 K for these specimens show a transition in the configuration occurring at 76 K, a "knee" with an accompanying secondary modulus (Figure 1).- The appearance of a "knee" has been correlated with initial matrix cracking and the secondary modulus, with interfacial failure, There were several fractographic features that provided evidence of changing matrix damage at the lower temperatures. The increased brittleness of the matrix resulted in opaqueness. Tear features were apparent at the lamina interface and tear ridges showed in the matrix (Figure 19). A number of matrix cracks, both external and internal were visible. There was more matrix on the fracture surface for the low temperature specimens and there was some evidence of hacklelike marks of the matrix at the fracture surface (21). 549 Finally, interface failures at all temperatures for these specimens were identified as delamination (of the lamina) and debonding and pullout of the fibers and fiber bundles. Debonding and pullout were most characteristic of failures transverse to the applied load. There was significantly more of this type of damage at the lower temperatures. Delamination was identified in several ways. First, there were the short delamination steps associated with brush fracture of the longitudinal bundles (Figure 17). Second, there was almost total delamination of the lamina in a direction parallel to the applied load for some of the mixed mode failures. Third, there was the transverse delamination at the primary fracture edge for the warp tests at 295 K, which left the undulating surface of the fill bundles intact (Figure 17). There is some evidence in the published literature to indicate that an increased fiber fracture strength or a decreased matrix or interface shear strength lead to debonding rather than fracture in the debond/fracture sub- sequence of events. There seems to be some support for this sequence at the lower temperatures of the present study. In conclusion, examination and analysis of the fracture surfaces of the G-10CR specimens fractured in compression at 295 K, 76 K and 4 K have provided a basis for the identification and characterization of primary composite failure and secondary constituent failure. As the summaries in Tables 3 and 4 indicate there are identifiable temperature-related failure effects. Figures 6-20 correlated with Tables 3 and 4 emphasize the major effects. Figures 21- 26 illustrate temperature related increases in strength stiffness, and varia- bility or scatter in the data (in terms of standard deviation and coefficient of variation). This information is also reflected in Tables 3 and 4 and can be correlated with Figures 6-20. 550 6. Recommendations for Future Studies There has been little or no fractographic or modeling studies of the low- temperature compression failure behavior for glass-fabric-reinforced epoxies. The in-depth review of published studies that comprised the initial part of this report has been synthesized to provide a comprehensive survey of relevant information that can be used to support and direct the investigative efforts of this study. The literature review has been immeasurably helpful in the development and organization of this report and the failure studies that preceded it. The explanations, interpretations, and extrapolations have provided an invaluable documentary information base. A logical follow-up of this temperature fracture-mode study should include reconstruction of the fracture process from initiation, through pro- pagation, and on to final fracture developed in terms of a most probable sequence of failure events by using a Fault-free Analysis, which has been described in other publications (10). Concurrent with a sequence-of-events analysis should be a stress analysis describing material behavior just prior to fracture. The fractography and modeling developed in the present report will provide support information for the analyses suggested above. These stages, in turn, can serve as a logical supplement to the syste- matic development of a theory of failure for analyzing the low-temperatuare behavior of glass-fabric-reinforced nonmetallics under a static compression load. 551 REFERENCES Rosen, B. W., "Strength of Uniaxial Fibrous Composites," Mechanics of Composite Materials , Wendt, F. W., H. Liebowitz, and Perrone, N., Eds., Pergamon, New York (1970) 621-49. Zweben, C, "Tensile Strength of Fiber-Reinforced Composites: Basic Concepts and Recent Developments," Composite Materials: Testing and Design , ASTM, STP 460, American Society for Testing and Materials, Philadelphia (1969) 528-39. Tetelman, A. S., "Fracture Processes in Fiber Composite Materials," Composite Materials: Testing and Design , ASTM, STP 460, American Society for Testing and Materials, Philadelphia (1969) 473-502. Vinson, J. R., and Chow, T. -W., Composite Materials and Their Use in Structures , John Wiley & Sons, New York (1973). Broutman, L. J., and Krock, R. H., Eds. Modern Composite Materials , Addison-Wesley Publishing Co., Reading, MA (1967). Rosen, B. W., Kalkaini, S. V., and McLaughlin, P. V., Jr., "Failure and Fatigue Mechanisms in Composite Materials," Inelastic Behavior of Composite Materials , American Society of Mechanical Engineers, Applied Mechanics Division, (AMD, Vol. 13), New York (1975) 17-38. Rosen, B. W., "Mechanics of Composite Strengthening," Fiber Composite Materials , American Society for Metals, Metals Park, OH (1965) 37-75. Williams, R. S., and Reifsnider, K. L., "Fracture Analysis of Fatigue Damaged Mechanisms in Fiber Reinforced Composite Materials Using Scanning Electron Microscopy," Report to the Air Force Office of Scientific Research, College of Engineering, Virginia Polytechnic Institute and State University. 552 9. Zweben, C, "Tensile Failure of Fiber Composites," AIAA J. , Vol. 6, No. 12, (December 1968) 2325-31. 10. Masters, J. W., Yeow, Y. T., Louthan, M. R. , Jr., Reifsnider, K. L., and Brinson, H. F., "A Qualitative Fault Tree Analysis for the Tensile Failure of Fibrous Laminated Composites," Composites , (April 1977) 111-17. 11. Collins, B. R., Brentnall, W. D., and Toth, I. J., "Properties and Fracture Modes of Borsic-Titanium," Failure Modes in Composites , I. J. Toth, Ed., the Metallurgical Society of AIME, New York (1972) 103-29. 12. Kanninen, M. F., Rybicki, E. F., and Griffith, W. I., "Preliminary Development of a Fundamental Analysis Model for Crack Growth in a Fiber Reinforced Composite Material," Composite Materials: Testing and Design (fourth conference), ASTM, STP 617, American Society for Testing and Materials, Philadelphia (1977) 53-69. 13. Schramm, R. E., and Kasen, M. B., "Cryogenic Mechanical Properties of Boron-, Graphite-, and Glass-Reinforced Composites," Mater. Sci . Eng. , 30, (1977) 197-204. 14. Jones, R. C, "Fractography of Aluminum-Boron Composites," Composite Materials Testing and Design , ASTM, STP 460, American Society for Testing and Materials, Philadelphia (1969) 513-27. 15. Hartwig, G., "Mechanical and Electrical Low Temperature Properties of High Polymers," Nonmetallic Materials and Composites at Low Temperatures , A. F. Clark, R. P. Reed, and G. Hartwig, Eds., Plenum Press, New York (1979) 33-50. 16. Morgan, R. J., O'Neal, J. E., and Miller, D. B., "The Structure Modes of Deformation and Failure, and Mechanical Properties of Diaminodiphenyl Sulphone - Cured Tetraglycidyl 4.4' Diaminodiphenyl Methane Epoxy," J. Mater. Sci., 14, (1979) 109-24. 553 17. Salkind, M. J., "The Role of Interfaces in Fiber Composites," Surfaces and Interfaces II, Physical and Mechanical Properties , 14th Sagamore Army Materials Research Conference, J. J. Burke, N. L. Reed, and V. Weiss, Eds., Syracuse University Press (1968) 417-45. 18. Kelly, A., Strong Solids , Second Edition, Clarendon Press, Oxford (1973) 157-236. 19. Si h, G. C, Hilton, P. D., Badaliance, R., Shenberger, P. S., and Villarreal, G., "Fracture Mechanics for Fibrous Composites," The Test Methods for High Modulus Fibers and Composites , ASTM, STP 521 , American Society for Testing and Materials, Philadelphia (1973) 98-132. 20. Hertz, J., "The Effect of Cryogenic Temperatures on the Mechanical Properties of Reinforced Plastic Laminates," Soc. Plast. Eng. J. , 21, (February 1965) 181-89. 21. Hull, D., "Nucleation and Propagation Processes in Fracture," Polymeric Materials , American Society for Metals, Metals Park, OH (1975) 487-550. 22. Wigley, D. A., "The Properties of Nonmetals," Mechanical Properties of Materials at Low Temperatures , Plenum Press, New York (1971). 23. Soffer, L. M., and Molho, R., "Mechanical Properties of Epoxy Resins and Glass/Epoxy Composites at Cryogenic Temperatures," NASA-CR-84451 , NTIS, Springfield, VA (1967). 24. Mull in, J. V., "Influence of Fiber Property Variations on Composite Failure Mechanisms," Analysis of the Test Methods for High Modulus Fibers and Composites , ASTM, STP 521, American Society for Testing and Materials, Philadelphia (1973) 349-66. 25. Zweben, C, "Fracture Mechanics and Composite Materials: A Critical Analysis," Analysis of the Test Methods for High Modulus Fibers and Composites , ASTM, STP 521, 1973. 65-97. 554 26. Harlow, D. G. 5 and Phoenix, S. L., "Bounds on the Probability of Failure of Composite Materials," Int. J. Fract. , Vol. 15, Nov. 4, (August, 1979) 321-36. 27. Beaumont, P. W. R., and Tetelman, A. S., "Fracture Strength and Toughness of Fibrous Composites," Failure Modes in Composites , Istavan Toth, Ed., AIME, New York (1972) 49-80. 28. Greszczuk, L. B., "Compressive Strength and Failure, Modes of Unidirectional Composites," Analysis of the Test Methods for High Modulus Fibers and Composites , ASTM, STP 521, American Society for Testing and Materials, Philadelphia (1973) 192-217. 29. Lanir, Y., and Fung, Y. C. B., "Fiber Composite Columns Under Compression," J. Compos. Mater. , Vol. 6, (July 1972)387-95. 30. Davis, J. G., Jr., "Compressive Strength of Fiber Reinforced Composite Materials," Composite Reliability , ASTM, STP 580, American Society for Testing and Materials, Philadelphia (1975) 364-77. 31. Jones, R. M., Mechanics of Composite Materials , McGraw-Hill, New York (1975) 134-45. 32. Chaplin, C. R., "Compressive Fracture in Unidirectional Glass-Reinforced Plastics," J. Mater. Sci . , 12, (1977) 347-52. „ 33. Chou, T.-W., and Kelly, A., "The Effect of Transverse Shear on the Lon- gitudinal Compressive Strength of Fiber Composites," J. Mater. Sci . , 15, (1980) 327-31. 34. Ryder, J. T., and Black, E. D., "Compression Testing of Large Gage Length Composite Coupons," Composite Materials: Testing and Design (fourth conference), ASTM, STP 617, American Society for Testing and Materials, Philadelphia (1977) 170-89. 555 35. Hartwig, G., "Low Temperature Properties of Epoxy Resins and Composites," Advances in Cryogenic Engineering , Vol. 24, K. D. Timmerhaus, R. P. Reed, and A. F. Clark, Eds., Plenum Press, New York (1978) 17-36. 36. Kasen, M. B., "Mechanical and Thermal Properties of Filamentary-Reinforced Structural Composites at Cryogenic Temperatures, 1: Glass-Reinforced Composites," Cryogenics , Vol. 15, No. 6, (June 1975) 327-49. 37. Kasen, M. B., MacDonald, G. R., Beekman, D. H., and Schramm, R. E., "Mechanical, Electrical, and Thermal Characterization of G-10CR and G-llCR Glass Cloth/ Epoxy Laminates Between Room Temperature and 4 K," Advances in Cryogenic Engineering , Vol. 26, A. F. Clark and R. P. Reed, Eds., Plenum Press, New York (1980) 235-44. 556 List of Tables Table 1. Typical Values, E-glass Table 2. Typical Values, General Purpose Epoxy Resin Table 3. G-l OCR Fracture Analysis Summary, Warp Tests (Compression) Table 4. G-l OCR Fracture Analysis Summary, Fill Tests (Compression) 557 Table 1. Typical Values, E-Glass (9, 22, 33) E-glass Properties Temperature 295 K -230 K ^205 K ~83 K Tens. St., MPa 3447 5171 5343 5654 Young's Modulus, E, GPa 72.4 Modulus increases approximately 3% between 295 K and 76 K. % Fracture Elongation ~2.5% Table 2. Typical Values, General Purpose Epoxy Resin Epoxy Properties Temperature 295 K 77 K 20 K Tens. St. , MPa 41 103 100 Compress. St. , MPa 90-138 Young's Modulus, E, GPa 21 47 66 % Fracture Elongation 4.5% 2% 1.8% 558 Table 3 G-10CR Fracture Analysis Summary, Warp Tests (Compression) 295 K (8 specimens] j vs. Failure Mode: * (a) low a cu : ~55° slant (b) high o : V-notch (chisel) Primary Mode : 37=, "55° slant 25% V-notch 38% Miscellaneous (broke in gage section, mixed or missing) Primary Fracture Edge Major transverse crack across specimen. Initially cracked through matrix of outer lamina. Constituent Failure Til Longitudinal Direction (warp) Brush fracture of fiber bundles with some slight delamination at fabric crossovers (b) Transverse Direction (fill) some fiber microcracks, breaks and pullout. Many fiber bundles show delaminated outer surface but imbedded in inner matrix resulting in an undulating surface. (c) Interface Damage Primarily delamination id) Matrix Glassy and transparent. Smooth at fracture edge. No evidence of matrix cracks outside the fractured area 76 K (7 specimens] vs. Fail ure Mode :* a) low a u : V-notch (chisel (b) high a : mixed mode Primary Mode : 14% -55° slant 43% V-notch at face of grip 43% other (mixed and saw-tooth) Primary Fracture Edge Major transverse crack across specimen; same as 295 K. Some tearing-type damage at edge. Constituent Failure More general damage and more random in nature than 295 K. Rough fracture surface (a) Longitudinal Direction (warp) some brush fracture and delamination of fiber bundles (b) Transverse Direction (fill) Considerable damage to transverse bundles and individual fibers fracture, debond, and pullout. Bundles pulled out--not imbedded in inner matrix as they were at 295 K 4 K (3 specimens) cu vs. Failure Mode:* (a) low o : mixed (Failed at grip--no gage section separation) (b) high o u : mixed (55° slant, delamination, and bending.) Primary Mode : 67% mixed 33% failed at grip--no gage section separation Primary Fracture Edge 2 specimens show major transverse crack similar to 295 K and 76 K. No gage-section separation of third specimen. Interface Damage Debonding, delamination and pullout (d) Matrix Opaque. Microcracks in gage section on reverse side of fracture surface. Constituent Failure Considerable damage at primary transverse crack. Longitudinal Direction (warp) some brush fracture longi- tudinal fiber bundles with some microcracks in fibers and delamination. Transverse Direction (fill) microcracks in fibers. Some debonding, fiber fracture and pul lout. (c) Interface Damage Considerable delamination some fiber debonding (d) Matrix Opaque. Significant matrix damage--cracking and rough- ness. More matrix showing on fracture surface than at higher temperature. Cracks showing on external specimen surface at warp and fill crossovers. *Note: c refers to ultimate :ompressive strength. 559 Table 4. 6-1 OCR Fracture Analysis Summary, Fill Tests (Compression 295 K (3 specimens) 76 K (3 specimens) 4 K (3 specimens) cu a vs. Failure Mode: cu a vs. Failure Mode: cu vs. Failure Mode: (a) low a CU : -55° slant (a) cu low a : ~55° slant (a) low a cu : -55° slant (b) high a CU : ~55° slant (b) cu high a : V-notch (b) high o CU : "55° slant Primary Mode: Primary Mode: Primary Mode: 67% -55° slant 67% -55° slant 100% -55° slant 33% Missing Spec. 33% V-notch at grip Primary Fracture Edge Primary Fracture Edqe Primary Fracture Edqe Major transverse crack across Major transverse crack across Majo r transverse across spec lmen spec imen; see 295 K. More tearing- specimen; see 295 K. Initially cracked through type damage matrix of outer lamina Cons tituent Failure Constituent Failure Cons tituent Failure: (more surface damage than at 295 K) (Not as much visible fiber damaqe racture face; more matrix " on f show ing compared to that at 295 K (a) Longitudinal Direction (fill) (a) Longitudinal Direction (fill) (a) Longitudinal Direction (fill Brush fracture with some Brush fracture of fiber bundles Few bundle fractures delamination at crossovers with more delamination than room with some delamination- of fiber bundles. temperature. Microcracks in fibers. mostly at fracture edge. (b) Transverse Direction (warp) (b) Transverse Direction (Warp) (b) Transverse Direction (warp) Hore random fiber More transverse fiber bundle Considerable fiber damage; damage than warp tests. Con- debonding, fracture and pullout. irregular and random siderable damage of fibers Less individual fiber debonding damage of individual debond, fracture and pullout. fracture and pullout. Considerable transverse damage at fracture edge. fibers and bundles-.debonding fracture and pullout. (c) Interface Damage (c) Interface Damage (c) Interface Damage Delamination. Debonding and Delamination, debond, and pullout Some longitudinal delami- pullout of bundles and individual nation. Transverse fiber fibers. and bundle debond and pullout. (d) Matrix (d) Matrix (d) Matrix Glassy and translucent Opaque. More broken up and cracked than at 295 K. Significant matrix cracks undercutting transverse fibers. Trans- verse crack of matrix on reverse side of fracture surface. Opaque. More matrix showing on fracture surface Small number of glassy regions. Irregular surface Some internal cracks showing on reverse side. 560 LIST OF FIGURES Figure 1. Typical tensile stress-strain diagram G-10CR, Warp (37). Figure 2. Laminate and specimen configurations used in determining static compressive properties of G-10CR at 295 K, 76 K, 4 K. Figure 3. Fixture for static compression testing at cryogenic tem- peratures (13). Figure 4. Typical primary fracture modes in compression. G-10CR at 295 K, 76 K, 4 K. Figure 5. G-10CR. Primary compression fracture mode vs. temperature. See Tables 3 and 4 for information on the number of specimens in each machine mode. Figure 6. G-10CR, warp test, 295 K, ^55° slant mode fracture (37). Figure 7. G-10CR, fill test, 295 K. %55° slant mode fracture (37). Figure 8. G-10CR, warp test, 76 K. (a) ^55° slant mode (b) and (c) mixed-mode fracture, (a), (b), and (c) are different specimens (37). Figure 9. G-10CR, warp test, 76 K. (a), (b), and (c) are different specimens (37). Figure 10. G-10CR, fill test, 76 K. (a) and (b) are different speci- mens (37). Figure 11. G-10CR, warp test, 4 K mixed mode failure (37). Figure 12. G-10CR, warp test, 4 K. (a) and (b) are different examples of mixed mode failure (37). Figure 13. G-10CR, fill test, 4 K. ^55° slant mode failure (37). Figure 14. G-10CR, warp test, 295 K. (a) and (b) are different speci- mens. 561 Figure 15. G-10CR, warp test, 295 K. Fracture surface (35). Figure 16. G-10CR, warp test, 295 K. SEM of fracture surface (37). Figure 17. G-1QCR, warp test, 295 K. SEM of fracture surface (37). Figure 18. G-10CR, warp test, 76 K. SEM of fracture surface (37). Figure 19. G-10CR, warp test, 4 K. SEM of fracture surface (37). Figure 20. G-10CR, warp test, 4 K. SEM of fracture surface (37). Figure 21. G-10CR. Compressive strength, Young's modulus, tensile strain to failure vs. temperature (37). Figure 22. G-10CR. Young's modulus vs. temperature (37). Figure 23. G-10CR. Compressive strength vs. temperature (37). Figure 24. G-10CR. Percent increase in tensile strain to failure, Young's modulus, and compressive strength at cryogenic temperatures relative to room temperature. Calculated from results in reference 37. Figure 25. G-10CR. Standard deviation vs. temperature (37). Figure 26. G-10CR. Coefficient of variation in percent vs. tempera- ture (37). 562 1000r 800 - (0 0L - 600 - H 0) 400 - 200- STRAIN, €, % Figure 1. Typical tensile stress-strain diagram, G-10CR, warp (37) 563 a — w 01 = -*- F ill HI 5 I * 1 k TOP Typical Fiber Crossover Warp fiber bundles Fill fiber bundles (a) FRONT Diagram showing top and front views of woven fabric (with fill and warp bundles) (b) Parallel layers of lamina Delamination surface showing fiber weave Diagram of warp and fill plain weave in a G-10 compression specimen. (c) G-10Cr compression specimen Specimen 0.15x0.3 (.38x.76) Sec A-A 4.7 (11.9) / 0.75 (1.9) fi t 1.0 (2.5) t / 6061-T6 Al grip Al spacer 2.75 Al(7.0) (d) G-10Cr compression specimen with aluminum grips. Units: in. (cm) (37) Figure 2. Laminate and speciman configurations used in determining static compressive properties of G-10CR at 295K, 76K and 4K. 564 SPECIMEN Figure 3 Fixture for static compression testing at cryogenic temperatures (13). 565 Typical fiber crossover Primary fracture edge Matrix missing. Exposed fabric weave showing (a) Primary composite^ fracture ~55°slant Matrix missing. Exposed fabric weave showing Primary fracture edge (b) Primary composite fracture, V-notch or chisel Matrix missing. Exposed fabric weave showing. Primary fracture edge Major transverse crack across specimen. Initially cracked through matrix of outer lamina Primary fracture edge (c) Primary ~55° slant fracture surface Figure 4. Typical primary fracture modes in compression G-10CR at 295K, 76K, 4K. 566 Q O Ul oc ID O < OC LU > OC < oc 0. 100r 80 60 40 20 Warp Test 100 80 60- 40- 20 Mixed mode r/ / Failed at grip V-notch(chisel) Other 55° slant \ \l 4K ~55° slant 76 K FIN Test 295 K ~55° slant \ Missing specimen^ V-notch (chisel) 4K 76 K 295 K Figure 5. G-10CR. Primary compression fracture mode vs. temperature. See Tables 3 and 4 for information on the number of specimens in each fracture mode. 567 aj ■a o m ^_ a. > +j c i — IT3 r— r-» CO LSI oo LD II 3 U e D ai o qj 03 00 (4- s_ s_ CO CO c Qj ra S- S- 3 4-> +-> O QJ ra 1 — 5- +-> M- +J ■r— QJ r— +■> >> r— * — * <+_ a> r-~ > ro »r- * — -* 3 ■*-> QJ CD 3 s_ 4-> ■a O QJ CD IB CD C QJ i- S_ -r- a> <+- « 3 n3 r— O E QJ C -C (O •O UJ CO T3 O E -a c fO CO to LO en CM 4-> CO QJ +-> O. S- •r— -M , — Ol QJ c •.- i- •— s- 3 3 -a CT> •■— CO > -I— i- -a qj aj ■+-> 01 o S- rt3 T3 S_ r— ra LU (_) b68 OJ oi +-> o i. 4- o a) in CO ■ ■ •"*• < QJ CD S- r^ -t-> O cu o o LT> CM rv' O o QJ ra Q- 00 en qj n s_ CO -(-> 11 o rO 3 S- O Li- O .<* 569 CD C 5- n3 CD O c CU -a > CU CU QJ E C7> O TD t/> CU OJ • i. CU 3 T3 +-> o o E 03 4-> <4- c r— i_ co o3 O cn LD o. u S- CU -a o E i XJ CU X -a c CU ro XJ ^ o E t/i I c -(-> cu s= E 03 •<- r— <_) oo CU s/> n 4-> ? c cu — . i- fO CU — <+- co r-» cu s_ " 03 +-> wi - — cu u i- c fO o3 Or jQ CO CU 570 •^ - ■£ utf. tflr CD CO o CD CL oo ro oo CL O s: i- (_) VD ro en ro c oo o rors S-U rt3 CO CO CL S- +-> o rt3 s- CL o CTl O QJ CL •r- ro ■a s a> s- f0 o — ' -a TCS c CL 03 s: •» LO ^— *. -Q ro — ^ O r^ •» -- — x i T3 — -* 3 r-~ <_> i^ ro D — r» OJ oo i- - c: 3 +-> CD i — to E • r— OJ t- fa -l-> o U. CD CL CL CD i- on TD 13 O 2 +-> s: c: - CD T3 i- s- ai C_) CD X O t- ., — r— <+_ £ 1 •<- C5 -O . — o > - CTi CD i~ ex. CO o -o c ai CD a res CL 3 >> <-> test direction (fill) primary fracture edge a) ^55° slant mode. Considerable transverse cu transverse (warp) direction fill) damage, j = 536.6 Mpa transverse matrix crackinq Test (fill: direction face of end cap b) Matrix cracking across specimen gage section face o = 561 .6 Mpa Figure 10. G-10CR, Fill Test, 7 >K (a) and (b) are different specimens. (37). 572 , — c rO o C T- ■ — «. .,_ -M "O ra 3 C . +-> -r- CL ■i- E T3 en ra u C c — O CU -o —i -a c cu 4-> <+- cu Q. re u -o c cu CD -o o E -t-J re 5 ■a cu en i- re td -t- 1 "+- 3 i/i c +J cu •r— h- l/> cr. •r— C X o re r — - CD £_ o ro * •i — 4-> Q- ■M T3 21 C c r— ro Qj cu ld f — jO i — sf X> ro lo O C S- CO ID re re n i_o ars «* u ? c c/> D o •f— a> •i — T3 +j c O •a o (/) E c •r- Q. • r— +-> ro 13 E <_> a ai re ai X i — s_ -o ■ r— ^ in 1-^ cu CO •*-> -— ' a. cu s- s_ TJ =3 X i — •r— «* rO i"L 4- >_) o ai i — -o i c> CD E CU en 573 test direction bendi ng a) Mixed mode failure: del ami nation, bending of the specimen, and slant mode fracture. a cu = 881.9 Mpa test (warp, direction b) Mixed Mode Failure. No gage section separation. - 849.9 Mpa Figure 12. G-10CR, Warp Test, 4K. (a) and (b) are different examples of mixed mode failure (37). 574 »«*i\ •f*>',< "«f oO Z. OJ CO i. rs ro +-> en i- I! *..\ ai in e '■■1 o ai 575 r undulating geometry of woven fabric individual lamina delamination fracture face Fill Warp fill (transverse) fiber bundle fracture and pullout. Crush fracture warp ( longi tudinal fiber bundles Warp fill (b) primary (transverse) fracture edge. Figure 14. G- 1 OCR, warp test, 295 K (a) and (b) are different specimens (37) 576 Figure 15. G-l OCR, warp test, 295K Fracture surface ( 35) . fill fiber bundle debond, fracture, and pul lout. brush fracture of warp fi ber bundles - with some del ami nation primary (transverse) fracture edge ill (transverse) I Warp (test) direction undulating fill fibers '^h^S^J^ 1 - aw typical brush fracture of w^arp fiber bundles Fractured fiber ends are showing. ransverse) warp (test direction) 400HM Figure 16 G-10CR, warp test, 295 K. SEM of fracture surface (37) 577 primary ( trans vers* > fracture ('dye matrix primary transverse) fracture edge fill (transverse^ Matrix steps. Del ami na ted regions showing missing warp bundles. End views of fractured warp bundles are visible at rear boundaries of matrix steps. -\ ■ i ng fi 1 fi transverse, primai ■■ ■. tun i dg< • • ; u r i . . • ••• of frai * 578 fill f i be i debori':, f ra> t ■■ and pul loul brush f ractui e warp f i bet bund! e ; fill (transverse) fracture edge (transverse matn x warp fiber fracture and pul loir b) Figure I! G-10CR, warp test , ?6! . Si M o1 fracl un n ! i 579 fill fibers end view of fractured warp fiber bundles brush fracture. fill (transverse] Primary fracture edge matrix cleavage marks warp fiber bundle brush fracture fill (transverse, warp (test direction ■ ire 19 I0( R narp test, 4K. SI M ol fr icture surface ( 17) 580 nia tri x c 1 g a v a ; marks fill fiber debond , fracture and pu 1 lout warp fiber bundle brush fracture - with del ami na t i on warp (test direction fiber crossover Del ami na ted lamina, showing woven glass rei nforcement . Brush fracture warp fiber bundles. Some del ami nation, Inclined fracture surface of specimen. Shows undulating geometry of fill fiber bundles Some fiber fracture and pullout warp (test direction' Fi Figure 20. G-10CR, warp test, 4K. SEM of fracture surface (37; 581 CO 10 8 Compressive Strength o 6 X O 4K 76 K 295 K Q. O UJ Young's Modulus E TU 40 — G9 | m I ^ 4K 76 K ™ 20 295 K Warp Fill 4- ^ 3- 3 m 2- 1- Tensile Strain to Failure m 4K 76K 295 K Figure 21. G-10CR. Compressive strength, Young's modulus, tensile strain to failure vs. temperature (37). 582 LU 00 I LU Q. s LU CO CD +J rd 5- cu CL E QJ 4-> 00 > Z3 •a o 00 en c: =3 o >- a: o i CD C\J CM CD S- Bdo 'smnaoiAi s.onhoa 583 aww i i i \\ J ;\ O 800 — b z h- \ \ Warp o 700 - z yj DC H \ \ \ X^Average 600 — US > 55 500 — cc Q. 2 Fill\\\ \\\ O o 400 Average \ \\ 300 i i n — 100 200 300 400 TEMPERATURE, K Figure 23. G-l OCR Compressive strength vs. temperature (37) 584 o l^i^^to^^ II ■ l [ivS ,UJ 'i^M^m^ 3 •- wwvwww K^^lll , , : v -tip^ 1 1 1 1 1 1 o o CM O o o o o c 00 (0 Kf CM * *■ asvaaoNi % CD U S_ •r - =3 C CU •i — CD IT1 o M- o <_> 4-> 4-> 1^ c rt3 co •1— (O -C CU u s_ +-> S- +-> cx> rs 1— CD s- CU <4- 00 e +-> cu TO (1) s- S- , — +J cu • i— 00 Q. 5- oo E C CU cu cu > 4-1 i_ 4-> • I— ■' to E c 00 O 00 •1 — CI) n ^ S- 4-» CD Q. 1 00 F o oo CU 5- rn O 4-> -o > E c: c •r— •i — 4-> , — S- Z3 00 -M Ol T3 CU 3 • t/i IT( 1 rr ~ "^ iU C_> CD CU o_) o £Z CL 1 — =5 F 1 O CU CD >- +-> *d- C\J CU i- Z3 D> 585 100 80 60 U ) MPa 4 20 S (cr cu ) Compressive Strength Warp Fill 3r Young's Modulus S (E TU ) GPa , 4K 76K 295K 4K 76K 295K Tensile Strain To Failure S(c TU )x10" 1 2 % El Warp M Fill 4K 76 K 295 K Figure 25. G-10CR. Standard deviation vs. temperature (37) 586 (0 20 15 o lb 10 > ■ Compression Strength Hj Warp ■ Fill —15 lal V) ? 10 IUJ ■ > 5 1- g M 6 5 Young's Modulus 4K 76 K 295 K 4K 76K 295K 20 1- Tensile Strain to Failure w 15- m 3 v 10 5 > ■ 4K J2L m Warp 11 Fill 76 K 295 K Figure 26. G-10CR. Coefficient of variation in percent vs. temperature (37) 587 STATIC AND DYNAMIC YOUNG'S MODULI OF TWO FIBER-REINFORCED COMPOSITES H. M. Ledbetter Fracture and Deformation Division National Bureau of Standards Boulder, Colorado ABSTRACT Two uniaxially fiber-reinforced composites were studied: graphite/ epoxy and E-glass/epoxy . Dynamic Young's modulus was determined at kHz frequencies using a Marx three-component oscillator. Comparison with the static Young's modulus determined in compression shows agreement within experimental error. 589 Elastic constants find many uses. In technological materials they assume importance whenever problems arise concerning deflections, elastic stability (buckling), excessive elastic deformation (jamming), thermoelastic stress, residual stress, fracture, and so on. Ratios of elastic modulus to density or to thermal conductivity provide well-known structural design parameters. Elastic constants correlate empirically with many practical properties such as strength, hardness, and wear. Because they relate intimately to both the interatomic potential and the lattice- vibration spectrum, elastic constants relate in turn to many fundamental solid-state phenomena such as diffusion, theoretical strength, and lattice-defect energies. Elastic constants appear as coefficients in thermodynamic equations of state where their relationships to free energy, specific heat, and thermal expansivity become especially clear. Since they can be measured accurately (typically within one percent or 5 less) and small changes can be detected easily (typically 1 in 10 or better), elastic constants are useful physical properties for characterizing and comparing materials and for detecting changes in material properties. Elastic constants are especially important for composites. These materials exhibit elastic deformation at all stress levels up to rupture. Methods for measuring elastic constants fall into two categories: static and dynamic. Static measurements tend to be made in engineering laboratories at high stress levels. Dynamic measurements tend to be made in physics laboratories at yery low stress levels. The dynamic test frequency varies considerably: near 1 Hz for torsional-pendulum and forced-vibration measurements; up to several hundred kHz for resonance measurements; up to 100 MHz for pulse-echo measurements. 591 Dynamic and static elastic-constant measurements remain unreconciled among many engineers and scientists who use these properties. Engineers tend to be more familiar with static measurements and tend to avoid using results determined dynamically. This situation is unfortunate for several reasons. First, dynamic results are more accurate by approximately an order of magnitude. Second, dynamic measurements are quicker and less expensive. Third, dynamic measurements on small specimens, including sheet and wire, are relatively simple. Fourth, complete sets of elastic constants, including anisotropy effects, are determined with little difficulty. Theoretically, dynamic-static elastic-constant differences arise from three sources: the adiabatic-isothermal difference, J the stress- level (strain-level) difference, and the role of higher-order elastic constants at high-strain levels. However, all these differences are relatively small and relatively well understood so that the existing static-dynamic controversy among practitioners surprises the theorist. One feels the controversy will never be resolved except through comparisons of carefully obtained static and dynamic results on the same materials. Previously, the principal problem has been the dearth of reliable static measurements. This situation changed recently in round-robin Young 's-modul us T3l measurements made at six laboratories on two composites. Table 1 shows the static Young's moduli, which are averages of measurements on five specimens. Table 1 also shows the moduli determined dynamically at 60 and 30 kHz, respectively, by a three-component Marx-oscillator method.'- ^ Three static methods were used: ASTM D695, ASTM D3410, and sandwich- 592 beam. The six laboratories participating in the round-robin testing are designated by letters A through F in table 1. Two composites were studied in this series of tests: (1) a unidirectional graphite/epoxy (AS/3501) of high strength and medium modulus, and (2) a unidirectional glass/epoxy laminate (E-glass/1002) . Panels of both composites were prepared by autoclave curing according to the manufacturer's recommendations at the University of Delaware. The panels were cut into blocks that were sent to participants who prepared specimens. For both composites, results in table 1 show that the static (determined compressively) and dynamic results agree within experimental and/or material uncertainties. Thus, in these cases no basis exists for distinguishing between static and dynamic Young's moduli. All three theoretical factors listed above predict a dynamic elastic stiffness that slightly exceeds the static. Although the graphite/epoxy results show this relationship, the glass/epoxy results do not; there, the seven-percent error for the static value masks this relationship. Within our laboratory, we expect several further studies, theoretical and experimental, on the static-dynamic elastic-constant question for composites. This will include design and construction of apparatus for accurate static values as a function of strain (stress). Acknowledgment This study was supported partly by the DoE Office of Fusion Energy and partly by the NBS Office of Standard Reference Data. 593 REFERENCES [1] L. D. Landau and E. M. Lifshitz, Theory of Elasticity (Pergamon. London, 1959), p. 19. [2] F. D. Murnaghan, Amer. J. Math. 59, 235 (1937). [3] N. R. Adsit, Compos. Tech. Rev. 2 (No. 2), 32 (1980). [4] H. M. Ledbetter, in Nonmetallic Materials and Composites at Low Temperatures , (Plenum, New York, 1979), pp. 277-281. 594 Table 1 Static and Dynamic Young's Modulus, in Units of GPa, for Two Fiber-reinforced Composites AS/3501 [0°] E-glass/1002 [0°] Static A 115 41 B 130 45 C 125 43 D 119 50 E 128 48 F 123 46 Static Avg. 123 ± 6 46 ± 3 Dynamic 130 4 ± 1.3 44.3 ± 0.5 595 RADIATION EFFECTS ON COMPOSITES AND NONMETALLIC MATERIALS M. B. Kasen Fracture and Deformation Division National Bureau of Standards Boulder, Colorado ABSTRACT This discussion of radiation effects is part of a review paper on composites. The other parts of the review appeared 1 ast year, in Volume III of this series. 597 The resistance of nonmetallic materials and composites to damage by neutron and gamma irradiation of 4 K is an important consideration in selecting insulating and structural support materials for superconducting magnets of magnetic fusion energy systems. Current designs anticipate that glass-reinforced polymers and other nonmetallics could be exposed to lifetime dose rates equivalent to about 10 rads in some critical areas. 1 The ability of conventional organic-matrix materials to withstand such doses is significant concern. A number of studies have been conducted on the effect of irradiation on the room-temperature properties of nonmetallic insulators, but few 2 studies have extended into the cryogenic range. In addition to degrading mechanical, thermal, and electrical properties, the low thermal conductivities of the nonmetallics might also result in nuclear heating, which places an added burden on the refrigeration system. Extrapolating from room- temperature experience suggests that bis- phenol A epoxy systems cured with aromatic amines should have better radiation resistance than systems cured with dicyandi amide. Cycloaliphatic epoxies may be superior to bishpenol A, polyimides should be better than epoxies, and filled resin systems superior to unfilled resins, but con- firmation is lacking. Complexities of the processes involved do not permit a theoretical estimate to be made of the effect of irradiation at cryogenic temperature on the various properties of nonmetallic insulators-- an experimental data base must be acquired. Because the radiation resistance of a polymer is strongly affected by the exact molecular structure, such experiments must be carried on with materials characterized on the 2 molecular level . 599 The Oak Ridge National Laboratory has recently screened the effect of 4 K gamma and neutron irradiation on the stability of eight types of organic insulation. The volume resistivities of a bisphenol A epoxy, an inorganic-filled epoxy, an aromatic-polyimide sheet material, a polyvinyl formal wire coating and an FR-5 type glass-epoxy laminate o were found to be very little affected by a dose of 2 X 10 rads (y). The effectiveness of a filled epoxy and of a B-staged glass-epoxy adhesive for joining Cu-Be lap-shear tensile specimens was also relatively unaffected, as was the flexural strength of a filled epoxy, a G-10 laminate, and an FR-5 laminate. When doses were increased to higher levels, the mechanical and electrical properties of glass-reinforced epoxy laminates and filled epoxy resins were found to be significantly deteriorated. Residual flexural and compressive strengths of G-10CR and G-11CR laminates were 9 reduced to 10 to 15% of original values after exposure to 2.4 X 10 rads (y) and to 6 to 10% after exposure to 1.0 X 10 rads (y). These doses 20 2 20 were accompanied by fast neutron doses of 2.2 X 10 n/m and 8.7 X 10 n/m , respectively. The strength retention of the aromatic-ami ne-cured G-11CR product appeared better than that of the dicyandiamide-cured G-10CR after the lower dose, but no difference is noted after the higher dose. A variant of the G-10CR product containing boron-free E-glass reinforcement performed much like the G-11CR laminate. The relatively high fraction of initial flexural modulus retained by the laminates indicates that the glass reinforcement did not sustain major damage. The low strain to failure indicates that deterioration of the matrix- 600 and possibly of the Fiber-matrix bond—accounts for most of the property degradation. The mechanical properties of the filled resins were relatively unaffected by the lower level of radiation, but were substantially degraded at the higher level. These materials appear to be essentially at the end of their useful life after the higher level of radiation. Despite the larger percentage of strength reduction, the high initial strength of the glass/epoxy laminates permitted retention of sufficient residual strength (100 MPa, 15 X 10 3 psi flexural strength; 55 MPa, 8 X 10 psi compressive strength) for some magnet applications. Except for one of the filled epoxies, no significant dimensional changes were observed as a result of the irradiation. Warming of the specimens to room temperature after 4-K irradiation simulated periodic warming of the magnet structures throughout their lifetime and incorporated additional damage anticipated from defect recombination at higher temperatures--the Wigner energy. It is apparent that much work remains to be done before reliable radiation-resistant insulators are developed. Current programs are examining the type of molecular damage incurred during irradiation of bisphenol A epoxy-matrix laminates and testing is being extended to include glass-reinforced laminates having other resin matrices. 601 REFERENCES 1. Coltman, R. R., Klabunde, C. E., Kernohan, R. H. and Long, C. J. "Radiation effects on organic insulators for superconducting magnets," Oak Ridge National Laboratory Report ORNL/TM-7077 (1979). 2. Brown, B., "Low Temperature radiation effects on superconducting magnets" submitted to Journal of Nuclear Materials (1980). 3. Kernohan, R. H., Coltman, R. R., and Long, C. J. "Radiation effects on organic insulators for superconducting magnets," Oak Ridge National Laboratory Report ORNL/TM-6708 (1979). 602 ,*>. NONMETALLIC MATERIALS: TECHNOLOGY TRANSFER 603 CONTRIBUTIONS TO TECHNOLOGY FORECAST '81 + PLASTICS, CERAMICS AND COMPOSITES M. B. Kasen Fracture and Deformation Division National Bureau of Standards Boulder, Colorado + Submitted to Metal Progress 605 Nonmetallic composites perform critical structural, thermal, and electrical insulation functions in large superconducting magnets re- quired for magnetic fusion energy (MFE), magnetohydrodynamic (MHD) , rotating cryogenic machinery and energy storage systems required to meet future energy requirements of the United States and other countries. Large quantities of these material s--up to a thousand tons for a storage magnet--must perform reliably over a thirty-year lifetime while sub- jected to large electromagnetic stresses, thermal stresses, and intense radiation in some applications. Studies at NBS have shown that the variability in cryogenic proper- ties of conventional nonmetallic composites is a primary problem. This variability reflects both the intrinsic embrittlement of the polymer matrix at cryogenic temperatures and the extrinsic variability of cryo- genic properties among laminates produced to similar commercial desig- nations by a variety of producers. NBS personnel are working with magnet designers, industrial lami- nators, and the polymer industry to establish standard grades of non- metallic laminates for high technology cryogenic applications. The principle is to establish specifications for components and manufac- turing procedures that will minimize the property variability while pro- viding a basis for systematic development of improved laminates in response to demonstrated needs. This has resulted in the establish- ment of two commercial glass-fabric-reinforced epoxy laminate types identified as NEMA G-10CR and G-11CR. The specifications have been distributed to all U.S. laminators through the NEMA organization and are presently being manufactured by three concerns. 607 Organic-matrix laminates are unable to retain required mechanical and electrical properties after exposure to radiation levels on the order of 10 rads (y) plus a significantly fast neutron fluence at 4 K. Therefore, ceramics may be chosen over nonmetallic composites for insulation in MHD superconducting magnets, but ceramic use entails sub- stantially higher magnet construction cost. On behalf of the Department of Energy, NBS is cooperating with industry and with other federal lab- oratories to define the extent of the problem and to pursue development of improved radiation resistant laminates. Studies at the Oak Ridge National Laboratory have shown that the anticipated radiation levels reduce the compressive and flexural strengths of conventional glass- fabric epoxy laminates to 6 to 10% of their original values. Polyimide matrices are presently being evaluated. Cycloal iphatic epoxy systems and filled-resin laminates are primary candidates for future studies. A significant detriment to the expanded use of composite mate- rials in all areas of consumer technology is the difficulty of associ- ating performance with the profusion of products identified by trade names. Composite technology would greatly benefit from the establish- ment of an industry standards group who would establish a numerical composite classification system based on the key elements affecting performance. Such a system has been proposed by NBS, in which lower bound mechanical, electrical and thermal properties are defined by a number-letter combination defining type and flexibility of the matrix, the type and configuration of the reinforcement and the reinforcement volume fraction. The AISI designations for steel and the Aluminum Association designations for aluminum alloys serve as models. 608 + COMMENT ON "MECHANICAL PROPERTIES OF PLASTIC COMPOSITES M. B. Kasen Fracture and Deformation Division National Bureau of Standards Boulder, Colorado To appear as a letter to the editor, Composites 609 ■ ^ Comment on "Mechanical Properties of Plastic Composites Under Low Tem- perature Conditions" by K. Kaotani , Composites , V. 11, No. 2, April 1980, pp. 87-94 The author's conclusion that composites fabricated with a highly flexibilized epoxy matrix provide better cryogenic properties than those fabricated with nonflexibil ized systems may be valid for laminates con- taining 32% by volume reinforcement, but is not valid for the higher volume fractions typical of laminates in industrial usage. Literature 1-3 reviews and work in our laboratory indicate that residual stresses created by the large thermal contraction of highly flexibilized systems negate any advantage provided by the intrinsically higher strain capa- bility of the matrix when fiber contents approach 50 to 60% by volume in glass, graphite, or aramid laminates. We find, for example, that 4 the flexible Resin 2 formulation decreases the 4-K uniaxial longitu- dinal tensile strength of an S-901 glass-reinforced laminate to a level about equal to that of a nonflexibil ized matrix laminate at 295 K. The transverse strength of a uniaxial graphite-reinforced laminate at 4 K is reduced by about 60% by a highly flexibilized matrix, while the trans- verse failure strain is reduced about 90%. When used with aramid fiber, the flexibilized matrix reduces the uniaxial transverse strength to an extent rendering the laminate almost useless for structural applications. Additionally, the intrinsic weakness of highly flexibilized resin systems near room temperature has resulted in the catastrophic failure of several structures designed for cryogenic use but necessarily subjected to moderate stresses at room or slightly elevated temperatures. 611 The author quite correctly points out the limitation imposed on the effective cryogenic strength of glass-reinforced laminates owing to the strain limitations of conventional bisphenol A epoxies at such tempera- ture (the "knee point" in the stress-strain curve). But the solution requires a resin system providing reasonably high room-temperature strength and reasonably low thermal contraction as well as cryogenic toughness. Possible approaches have been discussed in the recent litera- cy ture. Until such systems are developed and proven, it appears most prudent to utilize conventional resin systems for cryogenic laminates. REFERENCES 1. Kasen, M. B., "Mechanical and thermal properties of filamentary- reinforced structural composites at cryogenic temperatures" Part I: glass-reinforced structural composites, Cryogenics 15, No. 6 (June 1975), pp. 327-349; Part II: advanced composites, Cryogenics , 15, No. 12 (December 1975), pp. 701-719. 2. Kasen, M. B., "Cryogenic properties of filamentary-reinforced com- posites: An update," National Bureau of Standards, Boulder, Colorado (in preparation). 3. Schramm, R. E. and Kasen, M. B., "Cryogenic mechanical properties of boron-, graphite-, and glass-reinforced composites," Materials Science and Engineering , 30 (1977), pp. 197-204. 4. Soffer, L. N. and Molho, R., "Cryogenic resins for glass-filament- wound composites," NASA CR-72114 (Final). 5. Nonmetallic Materials and Composites at Low Temperatures , A. F. Clark, R. P. Reed, and G. Hartwig, eds. (Plenum Press, 1979). 612 THE MANUFACTURE OF RADIATION RESISTANT LAMINATES J. R. BENZINGER CORPORATE TECHNOLOGY SPAULDING FIBRE COMPANY, INC. TONAWANDA, NEW YORK 14150 PRESENTED AT THE NBS/DOE WORKSHOP ON MATERIALS AT LOW TEMPERATURE, OCTOBER 8, 1980, VAIL, COLORADO 613 THE MANUFACTURE OF RADIATION RESISTANT LAMINATES J. R. BENZINGER Corporate Technology Spaulding Fibre Company, Inc. Tonawanda, New York 14150 INTRODUCTION There have been many studies on the irradiation effects of polymers and insulators, but only a few on high pressure thermo- setting laminates . There was virtually no published information on radiation at cryogenic temperatures on polymers, let alone laminates, until the recent work at Oak Ridge on the Spauldite© CR grade laminates, some particle filled epoxies and films of 2 Mylar and Kapton . One conclusion of the 1979 ORNL report \s T as i n tnat, at a dose o± 1 x 10*" rads, particle filled epoxies are at "end of life" in terms of mechanical strength while CR grade laminates, although showing severe degradation, still retained usable strength. Further testing was suggested to determine margins of safety. In some locations of the magnet, such as adjacent to the injector of neutron beams, higher doses would require additional shielding or redesign at considerable expense . A more prudent approach it is claimed would be to improve the radiation resistance of the insulation. Inorganic insulators like ceramics are much more resistant to radiation but tend to be very brittle. This rules them out 615 for most fusion reactor applications since these magnets are subject to extreme mechanical and magnetically induced forces . Consequently, the critical applications can only be satisfied by the use of tough, high pressure thermosetting laminates . It is the purpose of this paper to demonstrate the capability of industry to modify high pressure thermosetting laminates for improved radiation resistance. Mechanical strength testing of some of the experimental G-10CR variants designed for improved radiation resistance yielded some surprising results. Also to be reviewed are the design rationale, materials of construction, and method of manufacture of some polyimide and epoxy laminates prepared for MIT and ORNL irradiation effect studies. In conclusion, cost comparisons, availability and anticipated manufacturing problems associated with some raw materials and laminates will be dis- cussed. MIT AND ORNL PROGRAMS In response to requests from H. Becker and E. Erez of MIT and J. Long of ORNL, several experimental laminates were prepared in the laboratory with the intent to improve radiation resistance Some of them were variants in some manner to Spauldite® G-10CR or G-11CR. The program at ORNL, sponsored by DOE is an ongoing special purpose materials program on "Radiation Effects On Organic 616 Insulators For Superconducting Magnets". The MIT work was performed in the MIT Plasma Fusion Center also under DOE aus- pices. This work was part of an investigation of material « properties for an ignition test reactor (ITR) subject to high neutron fluence and gamma radiation . The radiation effect results on the laminates of these programs will be reported by their respective authors at the NBS/DOE Workshop On Materials At Low Temperature at Vail, October 7-9, 1980. MATERIALS OF CONSTRUCTION Fabric Reinforcements Listed in Table I is a description and cost of the five woven fabrics used as reinforcements in the experimental lami- nates. The low cost plain weave E glass 7628 A-1100 was selected as a control based en an extensive test data base and a long 7 history of successful use in laminates . S glass was selected not only because of its high strength, but also because it is boron free. Table II lists the chemical composition of the various glasses. E glass which contains boron may present a problem in reactors in that it retains residual radioactivity for longer periods of time than boron free glasses. Quartz was chosen because it tested the best of several inorganic insulators for resistance to irradiation. The standard finish available on the commercial S and quartz fabrics was A-1100, the same finish as used in the CR 617 grade laminates. At least one investigator has concluded that o A-1100 was the best of a variety of coupling agents evaluated . It is suggested that some future program include an update study of coupling agents when time permits ordering and preparing special samples. Although all of the five fabrics used in this program were commercial grades, availability and cost of quartz and S glasses could present problems in delivery. As for manufacturing, there are less problems associated with the manufacture of high pressure laminates with plain weaves. However, the satin weaves tend to perform better in low pressure molding where curvatures are involved. See Figure 1 for a schematic of the comparisons of construction. Resins and Curing Agents O " o It has been reported by several early researchers that aromatic compounds such as polyimides and aromatic epoxies have g good radiation resistance . They are many times more stable than aliphatic compounds to radiation owing primarily to the greater bond strength and the resonance energy of the aromatic structure. Substituted aromatics are more resistant than non- substituted aromatics and saturated aliphatics are more stable than unsaturated aliphatics. Other investigators have reported that the aromatic amine curing agent diaminodiphenyl sulphone was best when used with epoxies of the diglycidyl ethers of bisphenol A type and that diaminodiphenyl methane gave the best 618 results when used with multi-functional glycidyl amines . It was also observed that the more highly functional and cross- linked polymers that perform well at high temperatures also tend to fair best at high levels of radiation. Listed in Table III are the epoxies, curing agents and polyimides selected for the program. The polyimide, Kerimid 601, was selected for its ease of processing, modest cost as far as polyimides are concerned and because it has been reported to be unaltered after a 10 rad exposure. The other polyimide used in the program was DuPont NR-150B2 polyimide precursor solution. Shortly after the experiments were underway, DuPont withdrew this resin from the market . LAMINATE PREPARATION The laminates prepared for the MIT and ORNL programs are listed in Tables IV and V. Most of the laminates were very thin and consisted of only two or three sheets of prepreg. For this reason the laminate specimens were prepared by labora- tory "hand treating" rather than using the pilot facility. This process involved hand dipping of the reinforcements in solutions consisting of the resin polymer, its curing agent and compatible solvents. The resin treated webs were then passed through motor driven squeeze rolls. The rolls were preset so that pressure was applied to the treated web forcing the resin into the interstices of the fabric. The force of the G19 rolls also squeezed off excess resin, leaving a uniform appli- cation. After saturation, the treated webs were placed in an air- circulating oven at 280-300° F. The purpose of the drying operation for the polyimides was to remove the solvent. In the case of epoxies and phenolics, additional polymerization takes place in the drying oven in addition to the removal of solvent. All of the prepregs were subsequently pressed between steel plates at a pressure of 500 psi and a temperature of 360°F with the exception of one of the Kerimid laminates which was pressed at 440°F. Two of the polyimide laminates supplied to ORNL were post cured at 450° F for five hours to insure adequate cure. The epoxy based laminates were not post cured. Results of some radiation effects by R. Sheldon on one wet lay-up epoxy laminate indicated that higher temperature post cure tended to improve g radiation resistance . Subsequent designed experiments at different post cure temperatures by D. Evans indicated no signi- 9 ficant differences could be detected in radiation resistance . Therefore, for this reason and because the process is irrevers- ible, the epoxy laminates were not post cured. Based on initial radiation results, a post cure study, if deemed advisable, could be undertaken simply by post curing replicate samples of the laminates. Also included for controls were the commercially available laminate, Spauldite® G-11CR and a phenolic cured epoxy novolac. 620 G-10CR VARIANTS GEM-CR, Quartz Filled As part of an ongoing research program at Spaulding Fibre Company, Inc. several variants of G-10CR and G-11CR were investi- gated with a view toward improvements in radiation resistance, mechanical strength and cost. One of these was a continuous filament E glass epoxy mat laminate (GEM-CR) which was constructed using the G-10CR resin system. Last year Dr. dinger of OCF reported on the good results he had achieved at liquid nitrogen temperatures using mat epoxy laminates . Also several investigators have reported improved radiation stability by incorporating particle fillers in epoxy 1 2 resin castings and moldings . Other improved properties reported for filled resins are less thermal contraction, reduced cracking or crazing, increased or reduced thermal conductivity 13 based on the type and amount of filler and cost . Therefore, a combination of reinforcement, resin and filler seemed like a good approach to improve radiation resistance and other proper- ties of laminates as was the case with castings. The initial experiment was conducted with glass mat rather than fabric for three reasons: (a) lower cost, (b) random and wide spacing of the fibers for reception of the filler, and (c) prior success in GPO laminates. Although several fillers were considered as possible candidates, see Figure 2, the filler selected for experimentation was Quartz (Si0 2 ) because, as 621 stated previously, it tested the best of several inorganic 14 insulators for resistance to irradiation . Shown in Figure V are some mechanical strength properties of both filled and unfilled GEM-CR. The properties of the filled GEM-CR laminate were not as good as the unfilled GEM-CR, but this is believed due to the high loading of Si0 2 , 80 parts to 100 parts of resin. A high loading was used to reduce costs as well as improve radiation stability. Dr. Kasen has suggested that while the optimum amount is unknown, it is possible that a small amount, well distributed, could be effective. The bar graph in Figure 3 compares the compressive strength edgewise at ambient (293 K) and liquid nitrogen (77 K) tempera- tures of the unfilled GEM-CR laminate versus a NEMA GPO-1. (glass mat polyester). The GEM-CR variant like G-10CR increases considerably in strength at cryogenic temperatures. The bar graph in Figure 4 is similar to Figure 3 in all respects except that it compares interlaminar shear (unsupported) as a function of temperature. S Glass And High Glass Content Some mechanical properties of the S glass variant of G-10CR are shown in Table V compared with the average test data of Spauldite® G-10CR. This laminate was made identical to G-10CR with the only variant being S in place of E glass. The weave, finish and resin were the same as G-10CR. The two properties that stand out are compressive strength normal to laminations 622 and bond strength. Both tests when repeated gave the same identical results. The high E glass content variant of G-10CR demonstrates another way to achieve higher strengths. Both of the above laminates, one because of the absence of boron oxide and other because of less organic content, would be expected to provide improved radiation resistance. Polyimide And Phenolic-Epoxy Also listed in Table V are the properties of the polyimide laminate submitted to Joe Long at Oak Ridge for irradiation effects. This laminate is grouped with an aniline phenolic cured epoxy novolac epoxy. Both laminates contain high tempera- ture, highly aromatic polymers which like G-11CR are considered key elements for resistance to radiation. These resins are harder and tend to. be a little brittle as evidenced by the reduced interlaminar shear and bond strengths when compared with G-10CR (epoxy resin) laminates. G-10CR Quartz Filled Based on the good results achieved with the particle filled glass epoxy mat laminate, a quartz filled variant of G-10CR was planned. Several changes were made, one of which was to reduce the amount of quartz (SiCU) added to the epoxy resin from 80 to 40 parts by weight to 100 parts epoxy. The selection of a finer particle size Si0 2 filler, half the size used in the mat, was considered another improvement. In order to improve the bond of the filler to the epoxy with resultant expected increase in 523 strength a filler was obtained that had been pretreated with A-1100 coupling agent. Another reason to explain the lower strength of the filled GEM-CR laminate is that the percentage of glass fiber mat rein- forcement was reduced by the exact amount of the loading of filler while the amount of resin remained the same as in the unfilled laminate. In the case of the Si0 2 filled fabric laminate, the resin, rather than the reinforcement was displaced by the filler, and the amount of glass fabric reinforcement was not reduced but kept the same as the unfilled G-10CR. Test data on mechanical strength properties of this latest variant of G-10CR is not available for inclusion in this report, but will be reported verbally at the workshop. It should be emphasized that all of the laminates reviewed in this paper, with the exception of G-10CR and G-11CR, are laboratory develop- ments and not commercially available at this time. Also, most of the reported data are based on one time testing of laboratory specimens . CONCLUSION It has been demonstrated that laminates can be modified in many ways to make them more resistant to radiation. The success of these endeavors will be largely unknown until and when the laminates can be properly tested and evaluated. Higher tempera- ture irradiations can serve as a guide, but as concluded by B. Brown only irradiations performed near the temperature of the 624 magnet are realistic . If it is deemed necessary that a more comprehensive program be established, it seems more than likely that a significant improvement in radiation resistance can be successfully achieved with high pressure laminates. ACKNOWLEDGEMENT The author wishes to gratefully thank E. Erez of MIT for the test data used to generate the graphs in Figures 3 and 4. REFERENCES 1. M. G. Young, "Radiation Effects On Laminated Plastic Materials", UD-NEMA Facility, April 1965 2. R. H. Kernohan, R. R. Coltman and C. J. Long, "Radiation Effects On Organic Insulators For Superconducting Magnets", Oak Ridge National Laboratory Report ORNL/TM-670S (1979) 3. R. T. Santaro. J. S. Tang. R, G. Alsmiller, J_ . and J. M. Barnes, Nucl . Tech. .37, 65 (1978) 4. R. P. Reed, F. R. Fickett, M. B. Kasen and H. I. McHenry, "Materials Studies For Magnetic Fusion Energy Applications At Low Temperature - 1", Report No. NBSIR 78-884, National Bureau of Standards, Boulder, Colorado (1978), p. 243 5. M. B. Kasen, G. R. MacDonald, D. H. Beekman and R. E. Schramm, "Mechanical, Electrical And Thermal Characterization Of G-10CR And G-11CR Glass Cloth/Epoxy Laminates Between Room Temperature And 4 K", Advances In Cryogenic Engineering, V 26 Plenum Press N.Y. (1980) 625 6. E. A. Erez, H. Becker, "Radiation Damage In Thin Sheet Fiberglass Insulators", Non-metallic Materials And Com- posites At Low Temperature - II, Plenum Press 1981 7. J. R. Benzinger, "Advances In Cryogenic Engineering", Volume 26 Plenum Press, New York 1980 pp. 252-258 8. R. Sheldon and G. B. Stapleton, "The Effect Of High Energy Radiation On The Mechanical Properties Of Epoxy Resin Systems Used For Particle Accelerator Construction", Rutherford Laboratory, Chilton, Didcot, England (1968) 9. M. H. Van de Voorde, IEEE Trans. Nucl. Sci. 18 , 784 (1970), and IBID 20_, 693 (1973) 10. D. Evans, J. T. Morgan, R. Sheldon, G. B. Stapleton, "Post- Irradiation Mechanical Properties Of Epoxy/Glass Composites RHEL/R200 Chilton, Berkshire, England, 1970 11. J. Olinger, "Cryogenic Grade Glass Mat Fnnxy Laminates", NBS/DOE Workshop On Materials At Low Temperature", Vail, Colorado, October 1979 12. H. Brechna, "Effect Of Nuclear Radiation On Organic Materials; Specifically Magnet Insulations In High-Energy Accelerators", SLAC Report No. 40, 1965 13. G. Claudet, F. Disdier and M. Locatelli, "Interesting Low Temperature Thermal And Mechanical Properties Of A Particular Powder-Filled Polyimide", Non-Metallic Materials And Composites At Low Temperatures, Plenum Press pp. 131- 140, 1979 626 14. C. L. Hanks and D. J. Hamman, "Report On The Effect Of Radiation On Electrical Insulating Materials", REIC Report No. 46, Battelle Institute, Columbus, Ohio (1969) 15. B. S. Brown, "Radiation Effects In Superconducting Fusion Magnet Materials", May 1980, submitted to Journal Of Nuclear Materials LIST OF TABLES TABLE I TABLE II TABLE III TABLE IV TABLE V TABLE VI Description Of The Woven Fabric Reinforcements Chemical Composition Of The Reinforcements Polymeric Materials Of Construction Fabric Reinforced Polyimide Laminates Prepared For Irradiation Fabric Reinforced Epoxies Prepared For Irrsdiation Mechanical Strength And Physical Properties Of Some G-10CR Variants TABLE OF FIGURES Figure 1. Comparison Of Weaves, Style 7628 And Style 181 Figure 2. Relative Radiation Resistance Of Inorganic Insulating Materials, Based Upon Changes In Physical Properties Figure 3. Compressive Strength Edgewise As A Function Of Temperature Spauldite® GEM-CR Versus NEMA GPO-1 527 Figure 4. Interlaminar Shear Strength ASTM Guillotine (Unsupported) As A Function Of Temperature GEM-CR Versus GPO-1 628 CO Z o u < > o w o z o u w < H « o 00 o CO O CO i-J • • . . • o\ CM r-- CD CM CO c_>«=r> CO fc Q o ? en CT> rH NO CT> en •— ' . • • • • . W Kl LD. vO u~t CO CO 3= o H CM cm CM ^3- •^f z ro to hO LO LTj :=> X X X X X o -3- n- CM r^ I-- c_> "«t ** rr in Ln w c P P C C > •H •H •H •H •H < ca rt d +-> +-> w i— i rH t— i nr ca S rt 2= P. cx P- CO to CO to K) tq ■M (X CM u rg >H I i H w CO W co Z o o CD O O t— i p d £ (—■ 1— ( i— i r-i i— I s -1-1 ■ H •H •H T3 to 10 10 to CD 1 •H CD o o p O p fe Wh X C •H X CO PC •i— ' •H •H o t; oh § E co 1 CD O CU-H E CD to w CD CD as < 1 O i i O i O ' s 1— 1 z C-H o P -H o C -H CD I i-H •H +-> cc •H U p •H V-, p •H $M rj rH >% 10 n CO -3 E -t-> CO rH E +-> i-H E +J 5 i-H bO O & Z C_) i i 10 1 1 •H 10 i i •H J-. o -t-> o o u t— 1 t— 1 t— ( CO X & i-H S3 x & T— 1 S3 X fe rt X a • H O P. 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UJ o CM CC UJ h- UJ cc D o 2 D u_ < < BdW 638 TECHNOLOGY TRAHSHR 639 ^'2 NBS-DoE WORKSHOP: MATERIALS AT LOW TEMPERATURES October 7-9, 1980, Vail, Colorado N. J. Simon Fracture and Deformation Division National Bureau of Standards Boulder, Colorado ABSTRACT The subject matter and format of this annual workshop are discussed During fiscal year 1980, this workshop constituted the chief formal ac- tivity of the technology transfer program. However, starting with the present fiscal year, a second major activity will be undertaken: the production of handbook pages for low-temperature mechanical and physical properties of materials for magnets in fusion energy systems. 641 k. J The workshop discussions concerned structural and insulating mate- rials to be used at low temperatures in superconducting magnets for fusion reactors and MHD systems. Each year, a major purpose of the workshop is to bring together representatives from the community of users and suppliers of materials for low temperature magnet applications with those engaged in development and characterization of these materials. Participants this year came from industry, the national laboratories, the magnetic fusion energy programs, NBS, DoE and several universities. A list of attendees is given in the following pages. The technical pre- sentations were informal, with many opportunities for comments and ques- tions. Initially, overviews of the large superconducting magnet programs for fusion energy and MHD were given. Presentations regarding the Large Coil Projects and the ORNL test facility followed. Then, recent measurements of properties of structural alloys, welds and castings were presented. In the next session, fracture, fatigue and radiation damage in composites were considered. In the last session, codes and standards for fusion energy magnets were discussed. A detailed list of all the technical papers presented is attached. The proceedings of the workshop are not published, but abstracts or short summaries of all the presentations are provided to those who attend. Meetings of several advisory groups were held during the workshop. For example, the CCP conductor sheath advisory committee met, as did the joint NBS-industry group which is working on specifications for CR-grade laminates. 643 PROGRAM NBS/DoE Workshop on Materials at Low Temperatures October 7-9, 1980 Vail, Colorado Tuesday, October 7 Session 1 (9:00-12:00): Large Superconducting Magnet Programs R. P. Reed, Chairman DoE Office of Fusion Energy - Overview NBS Program Magnetic Fusion Test Facilities, A and B Engineering Test Facility MHD Magnets Session 2 (2:00-5:00): Large Coil Projects D. Beard, Chairman D. Beard, DoE R. P. Reed, NBS E. N. C. Dalder, LLL R. J. Hooper, ORNL V. Der, DoE Oak Ridge National Laboratory General Dynamics General Electric Westinghouse Japan Atomic Energy Research Institute Kernforschungszentrum Karlsruhe Wednesday, October 8 Session 3 (8:30-12:00): Structural Alloys, Welds and Castings H. I. McHenry, Chairman Cr-Ni-Mn-N Stainless Steels C. J. Long, ORNL J. L. Christian, GD C. L. Linkinhoker, GE C. J. Heyne, W E. Tada, JAER1 A. Ulbricht, KfK Conductor Sheath Materials Welding Research at Lawrence Berkeley Stainless Steel Weldments Aluminum Weldments Stainless Steel Castings Session 4 (1:30-5:00): Insulating and Structural Composites M. B. Kasen, Chairman R. L. Tobler, NBS R. Ogawa, KOBE R. Gold, W J. W. Morris, LBL T. A. Whipple, NBS J. W. Elmer, NBS R. Finch, £SC0 Fracture Modes of G-10CR in Compression Fatigue of Fibrous Glass/Epoxy Laminates in LN, Multiaxial Fatigue Damage in Fiberglass Radiation Damage in Fiberglass Insulators Radiation Effects in Organics at 4 K Manufacturing of Radiation Resistant Laminates Low Temperature Irradiation Studies B. A. Beck, NBS J. L. Olinger, Owens Corning S. Wang, UI H. Becker, MIT C. J. Long, ORNL J. Benzinger, Spaulding B. Brown, ANL 6:00 - 8:00 P.M. COCKTAIL PARTY Thursday, October 9 Session 5 (8:30-12:00) Codes and Standards E. N. C. Dalder, Chairman Structural Design Basis for Superconducting Magnets The Development of Codes and Standards Cooperative Research for Codes and Standards Code Approaches to Fracture Control Material Suppliers Viewpoint Filler Metal Suppliers Viewpoint 644 H. Becker, MIT W. E. Cooper, TES M. Prager, MPC H. I. McHenry, NBS E. H. Espy, Armco D. J. Kotecki, McKay . j CO cn -o • CO c s- co J3 T3 CO •»■ rO -a ro J- T3 J- -o • CO s- — 1 ro t. 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CD CO at ui c o -o o co at •r— •»-> _l 2: ^ CO CO >> to a> <_) 4-> r— S- UJ $- C3 +J ra ro =c s- c ra -M -!-> _l i— to >>T3 C +j jo a» a at . at ra o >> ra >> CO • co 5- =c enze ra ro ^ o J- _l • ro at 1— TO ro JO c -u S- E 3 o ^^ UJ c Q ■*-> c s_ JZ o •«- _i at at at c t>- J= b^ *£ at l_ •D E ^~ 1- s- o o a> a* ro 5- 3 O c o JZ 1- o s- s- E r- JO C 3 c tt- u T3 JO J- c o c 1- > U_ CO Q CO n: rj CO 1- ro . CV T3 •t— o «o •r- at j= • ro >> o o >> £ at f- >> = • o ra 5 • JZ •^ ra ra CO "~3 3 ^ UJ > -3 3: oc jo CO occu CJ r- u_ =C Q i— CM CO Tl- cn to r»* co cn o ^~ CM cn ^i- cn to r^ co cn o r— CM ro ^i- to to r-* CO cn o CM co ■**- LO to r». co CM CM CM CM CM CM CM CM CM CO CO ro CO CO CO CO CO CO 645 «, $ 646 ORGANIZATIONAL CONTACTS The people responsible for the various major aspects of the program are listed below so that specific technical questions may be directed to them. Department of Energy, Office of Fusion Energy, Washington, D.C. 20545 Program Monitor D. Beard (301) 353-4958 (FTS) 233-4958 National Bureau of Standards, Boulder, Colorado 80303 Program Manager R. P. Reed (303) 497-3870 (FTS) 320-3870 Welding and Casting H. I. McHenry (303) 497-3268 T. A. Whipple 497-5795 E. L. Brown 497-5790 Elastic Properties H. M. Ledbetter (303) 497-3443 Nonmetallics M. B. Kasen (303) 497-3558 Structural Alloys R. P. Reed, R. L. Tobler (303) 497-3421 Technology Transfer N. J. Simon (303) 497-3687 ESC0 Corporation, 2141 N.W. 25th Avenue, P. 0. Box 10123, Portland, Oregon 97210 Program Manager L. Finch (503)228-2141 ,X520 University of Wisconsin, Engineering Research Building, Madison, Wisconsin 53706 Program Manager D. Larbalestier (608)262-2194 647 ACKNOWLEDGMENT The editors wish to express their appreciation to JoAnne Wilken for preparing much of the manuscript. They are also grateful for the assist- ance of Carole Montgomery and Vickie Grove. Marilyn Stieg removed many of the inconsistencies, grammatical errors, and awkward usages of language. The editors are responsible for those errors that remain. 648 k. j NBS-114A (rev. 2-80) U.S. DEPT. OF COMM. BIBLIOGRAPHIC DATA SHEET (See instructions) 1. PUBLICATION OR REPORT NO. NBSIR 81-1645 2. Performing Organ. Report No. 3. Publication Date April 1981 4. TITLE AND SUBTITLE Materials Studies for Magnetic Fusion Energy Applications at Low Temperatures-IV 5. AUTHOR(S) R. P. Reed and N. J. Simon, Editors 6. PERFORMING ORGANIZATION (If joint or other than NBS, see instructions) NATIONAL BUREAU OF STANDARDS DEPARTMENT OF COMMERCE WASHINGTON, D.C. 20234 7. Contract/Grant No. 8. Type of Report & Period Covered 9. SPONSORING ORGANIZATION NAME AND COMPLETE ADDRESS (Street. City, State, ZIP) Department of Energy Office of Fusion Energy Washington, D.C. 20545 10. SUPPLEMENTARY NOTES Related earlier reports: NBSIR 80-1627, NBSIR 79-1609, NBSIR 78-884 J Document describes a computer program; SF-185, FlPS Software Summary, is attached. 11. ABSTRACT (A 200-word or less factual summary of most significant information. If document includes a significant bibliography or literature survey, mention it here) This report contains results of a research program to produce material property data that will facilitate design and development of cryogenic structures for the super- conducting magnets of magnetic fusion energy power plants and prototypes. The program was conceived and developed jointly by the staffs of the National Bureau of Standards and the Office of Fusion Energy of the Department of Energy; it is managed by NBS and sponsored by DoE. Research is conducted at NBS and at various other laboratories througf subcontracts with NBS. The reports presented here summarize the fourth year of work on the low temperature materials research program. Highlights of the results are presented first. Research results are given for the four main program areas: structural alloys, weldings and castings, nonmetallics , and technology transfer. Objectives, approaches, and achievements are summarized in an introduction to each program area. The major portion of the program has been the evaluation of the low temperature mechanical and physical properties of stainless steel base metal and welds, with particular emphasis on the nitrogen-strengthened stainless steels. Aluminum alloys have received some consideration also. Work has been done on the production and standardization of nonmetallics, primarily industrial laminates, for low temperature applications and on th*i measurement of their properties at cryogenic temperatures. A brief description is given of the fourth NBS/DoE Vail workshop held in October 1980. 12. KEY WORDS (Six to twelve entries; alphabetical order; capitalize only proper names; and separate key words by semicolons) Alloys; aluminum alloys; castings; composites; cryogenic properties; low temperature; mechanical properties; nonmetallics; physical properties; stainless steels; structural alloys; welding. 13. AVAILABILITY fX~) Unlimited I I For Official Distribution. Do Not Release to NTIS ~] Order From Superintendent of Documents, U.S. Government Printing Office, Washington, D.C. 20402. rXl Order From National Technical Information Service (NTIS), Springfield, VA. 22161 14. NO. OF PRINTED PAGES 604 15. Price $41.00 * U.S. GOVERNMENT PRINTING OFFICE:1981— 780-205 / 386 USCOMM-DC 6043-PBfl \ PENN STATE UNIVERSITY LIBRARIES ADDDD7 AA3120 ,