ORNL-3593 UC-25 — Metals, Ceramics, and Materials TID-4500 (27th ed.) WASTER MECHANICAL PROPERTIES OF SOME REFRACTORY METALS AND THEIR ALLOYS H. E. McCoy, Jr. R. L. Stephenson I RiWeie; -Jdrs OAK RIDGE NATIONAL LABORATORY operated by UNION CARBIDE CORPORATION for the U.S. ATOMIC ENERGY COMMISSION DISCLAIMER This report was prepared as an account of work sponsored by an agency of the United States Government. Neither the United States Government nor any agency Thereof, nor any of their employees, makes any warranty, express or implied, or assumes any legal liability or responsibility for the accuracy, completeness, or usefulness of any information, apparatus, product, or process disclosed, or represents that its use would not infringe privately owned rights. 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MeCoy, Jr. Contract No. W-7405-eng-26 METALS AND CERAMICS DIVISION MECHANICAL PROPERTIES OF SOME REFRACTORY METALS AND THEIR ALLOYS R. L. Stephenson J. R. Weir, Jr. APRIL 1964 OAK RTNGE NATIONAL LABORATORY Ozk Ridge, Tennessee operated by UNION CARBIDE CORPORATION for the U. S. ATOMIC ENERGY COMMISSION R @ ORNL-3593 . THIS PAGE WAS INTENTIONALLY " LEFT BLANK iii CONTENTS Strengthening Mechanisms in High-Temperature Materials ........ Solid-Solution Strengthening ...... it nneanns Interstitial and Dispersion Strengthening ................ The Mechanical Properties of Nb, Mo, Ta, and W ................ Niobium-Base Alloys ...................................;.. Molybdenum-Base Al OYS ... vivrerrnecasrsncensennsnsonsss Tantalum-Base AlloysS i iieitretnetietenerenesssesennennnsaa Tungsten-Base AlLlOYS .iiviieieeerrnnneanecssoaanannansanas Internal Ericéion Studies of Refractory-Metal Systefis ......... Effects of Irradiation on Refractory Metals ..............c.... )« B .1 Molybdenum ....ieeiiiieiiieneeeeetessenesnannsionnnsosssnes B =0 v - V0 X 33 37 MECHANICAL PROPERTIES OF SOME REFRACTORY METAIS AND THEIR ALLOYS H. E. McCoy, Jr. R. L. Stephenson J. R. Weir, Jr: ABSTRACT A critical evaluation has been made of the available mechanical property data for Nb-, Mo-, Ta-, and W-base alloys. It was found that insufficient data are available to allow the design and construction of complex engineering systems involving these materials. A general evaluation of the potential service temperatures for Nb-, Mo-, Ta-, and W-base alloys was made on the basis that conventional alloys have been used up to two thirds of their absolute melting point. Strengthening mechanisms that have been used to achieve high operating temperatures for conventional alloys and that could be applied to refractory alloys are discussed. A review of the literature on the effects of irradi- ation on the mechanical properties of niobium, molybdenum, tantalum, and tungsten has been made. It has been found ‘that the existing data on this topic are rather scant. The date in gecneral show that the ductility of molybdenum, tantalum, and tungsten is reduced after irradiation at ambient temperatures. The yield and ultimate‘strengths are increased slightly by irradiation. High-temperature tube-burst tests show that the rupture life of the Nb—1% Zr alloy is not 'drastically influenced by irradiation. | THIS PAGE WAS INTENTIONALLY - LEFT BLANK INTRODUCTION Conventional high-temperature alloys, sucfi as the stainless steels and nickel-base alloys, have constantly been improved. These materials have and will continue to be invaluasble structural materials in the nu- clear field. However, proposed future nuclear systems require materials that will operate satisfactorily at temperatures in excess of the melting points of the nickel- and iron-base alloys. A scan of the periodic chart, melting points of the elements, and availability and subsequent costs reveals only four candidate materials: niobium, molybdenum, tantalum, and tungsten. In considering the potential of these metals and their alloys, the physical property data in Table 1 are useful. The data on nickel and iron are tabulated for comparison. The values for one half and two thirds of the absolute melting point are significant because they indicate, re- spectively, the temperature for which creep begins to be a problem and the maximum service temperature to which engineering alloys are commonly sub jected.? | The values of the microscopic thermal neutron absorption cross section for these metals are of interest for nuclear applications. With other parameters remaining constant, the use of niobium or molybdenum as a fuel elefient cladding material would result in better neutron economy than would the use of tantalum or tungsten. In considering a material for engineering application, it is nec- essary that the requirements of the particular application be carefully evaluated and contrasted with the properties of the material. Consider in particular the problem of choosing the structural and fuel cladding materials for a nuclear reactor using a liquid-metal heat-transfer medium. The material must have sufficient strength at the operating temperature, must be capable of fabrication into the desired shapes, and must withstand the corrosive influences of its environment. Although the fabricability and corrosion resistance are of the utmost importance, it is the purpose of this discussion to deal specif- ically with the mechanical property requirements of these materials. The 1D. Mclean, "Point Defects and the Mechanical Properties of Metals and Alloys at High Temperatures,”" p. 179, Vacancies and Other Point Defects in Metals and Alloys, Institute of Metals Monograph No. 23, 1957. Table 1. Physical Property Data Microscopic . Thermal Neutron ' o o 1/2 Absolute 2/3 Absolute Absorption ‘Modulus of Element Density Melting Point~ Melting Point Melting Point Cross Sectio Elasticity (8/cm?) (°c) - (°0) - (°c) (barns/atom) (psi) ‘ | x 108 Nickel 8.90 1453 - 590 878 4.5 30.0% Iron 7.87 1537 632 934 2.4 28.5° Niobium 8.57 2468 1098 ' 1554 | 1.1 17.7° IS Molybdenum 10.22 2610 1169 1649 2.4 47 Tantalun 16.6 2996 1362 1906 21 27 a Tungsten 19.3 - 3410 1569 2183 o 19 , 50 "Phy51cal Propertles of the Elements," Metals Handbook Vol. I, pp. 4451, American Society for Metals, 8th ed., 1961. ‘ ' Samael Glasstone, Principles of Nuclear Reactor Englneerlng, pp. 841;423 Van Nostrafid, . Princeton, N. J., 1955. - 4 . : CoeE °L. p. Jahnke et al., ""Columbium Alloys Today," Metal Progr., 78: 77 (1960). specific properties that must be evaluated include: (1) engineering design data; (2) data concerning the long-time chemical stability of the alloy; (3) the ductility between the minimum and maximum service temperatures; (4) effect of atmosphere on the strength and ductility; (5) the influence.of irradiation; and (6) thermal fatigue properties. - The components of é reactor system fequire materials having con- siderably different properties in Lhese six areas. Fér example, materials used for fuel element cladding or radiators must be consider- ably more dgctile than turbine blade or nozzle materials. Likewise, resistance to damagefiby irradiatibn is of importance for core structural materials but not for radiator materials. | ~ Although considerfible‘information is available on the high- temperature mechanical properties of Nb-, Mo-, Ta-, and W-base alloys, no Single alloy has been sufficiently evaluated in these six areas to make it ready for service in a nuclear system. In this discussion an attempt will be made to assess the state of affairs relative to these four refractory metals. The available data will be reviewed critically. Recommendations as to the choice of alloys for service over specific temperature ranges will be made. Areas in which data are lacking wiil be poihted out. . . STRENGTHENING MECHANISM> IN HIGH-TEMPERATURE MATERIALS - The following discussion of strengthening mechanisms is not in- tended to be a complete "textbook" treatment of the subject, but rather a meane of bringing to the reader's attention the many possibilitiés that must be considered. For example, when 1% Zr is added to niobium, it does not necessarily follow that the strengthening observed is due to solid~solution strengthening. The entire chemistry of the metal is changed and it is quite likely that the major portion of the strength- ening is due to the formation of zirconium-interstitial complexes (clusters or compounds). ' Specific data are presented ifi this discussion only where it illustrates a partiéuléf point. Data on the mechanical behavior of | refractory metals will be given in the ncxt scetion. Solid-Solution Strengthening Solid-solution strengthening may be defined as the increase in . resistance to deformation of a material brought about by dissolving in it another element. The introduction of atoms having a diameter dif- ferent from those of the parent lattice introduces strains. These distorted regions in the lattice interfere with the motion of disloca- tion and increase the resistance of the material to deformation. The amount of strengthening obtained by this mechanism ié proportional to the amount of solute up to the solubility limit. The strengthening is likewise proportional to the size difference in the solute and solvent atoms. However, this is not an independent factor since the degree of solubility decreases as the atomic misfit increases. This picture of strengthening based on a size factor, as was originally proposed by Mott and N'a'barro,2 is somewhat an oversimplification and recént workers37% have shown the valency or electronic effects to be important. ' Another effect of solid-solution alloying is that of lowering the stacking fault energy. This causes the dislocations to split into partials with a faulted region in between. For cross slip to occur, these partials must be forced together. This efféect, however, is confined to face-centered cubic materials and has beénrdbserved in copper and stainless steels. | | Some of the most interesting effécts arise from the tendency of impurity atoms to migrate to dislocations and to grain bofindaries. This tends to anchor the dislocations and to lock the sources./’The.segfega- tion of impurities in the grain boundaries is largely responsible for _the large effects that impurity atoms have on the recrystallization temperature of metals. Vandermeer et 32.5 have shown that alloy addi- tions to high-purity aluminum alter the rate of grain boundary migration in proportion to the diffusion rate of the solute in the solvent. 2N. F. Mott and F. R. N. Nabarro, p. l,:Réport of Conference on Strength of Solids, Physical Society, London, 1948. 3N. P. Allen, T. H. Schofield, and A. E. Tate, Nature, 168: 378 (1951). 4 “W. R. Hibbard, Jr., Trans. Met. Soc. AIME, 212: 1 (1958). °P. Gordon and R. A. Vandermeer, The Mechanism of Boundary Migration in Recrystallization, Tech. Rep. No. 3, Department of Metallurgical Engineering, Illinois Institute of Technology, August 1961. & However, recrystallization in the transition metals may be more complex. Abrahamson® has shown that the effect of alloying élements on the recrys- tallization temperature of niobium can be correlated with the atom per- cent solute and the free atom electron configuration of the solute element. The elements Mn, Fe, Co, Ni, W, Re, and Os lower the recrys- tallization temperature and Ti, V, Cr, Zr, Mo, Ru, Rh, Pd, Hf, Ta, Ir, and Pt raise the recrystallization tefiperature. Darken’ points out that the effect of substitutional alloy additions per se cannot account for the strength realized in materials. Substitutional alloying may be of more importance in conjunction with other strengthening mechanisms. The studies by Darken’ of the oxidation of a silver-aluminum alloy illustrate this point. At the temperatures studied, aluminum has a high affinity for oxygen whereas silver oxide is unstable. It was felt that at low temperatures the oxygen would diffuse to the aluminum atoms and in the limiting case the aluminum atoms would remain stationary. If the aluminum atoms were completely surrounded by oxygen'atoms the oxygen-to-aluminum ratio would be 6. As the aluminum atoms migréte, the cluster size would increase and the oxygen-to-aluminum ratio would decrease. Observations by Wriedt® on the oxidation of a Ag—0.1% Al and a Agr0.48% Al alloy support the proposed model. The oxygen- to-aluminum ratio decreased with increasing temperature and increasing aluminmm content. It was also found that when an alloy was oxidized at one température and subsequently exposed to an oxidizing atmosphere at a higher temperature the oxygen-to-aluminum ratio did not change. This indicates the very high stability of the aluminum-oxygen clusters. These are considerably more effective in strengthening the alloy than would be the strain fields due to the aluminum atoms alone. The effect that a substitutional alloying addition has on theil strength of a metal would be lost if' the solute element were removed; therefore, the allny addition must be compatible with the service '6E. P. Abrahamson II, Trans. Met. Soc. AIME, 221: 1196 (1961). 7L. S. Darken, Am. Soc. Metals Trans. Quart. 54(4): 60042 (1961). 8D. ¥. Wriedt and L. S. Darken, Research Laboratories, U. S.. Steel Corp., unpublished data. enviromment. Two of the processes whereby the solute may be lost are by evaporation -at high temperatures in vacuum and by selective leaching in a corrosive environment. Interstitial and Dispersion Strengthening Althbugh the interstitial atoms are,smaller‘énd diffuse more rapidly than substitutional allpying elements, they can effectiyely:alter the motion of disloéations. They are believed to be reéponsible for the ductile~-to-brittle transition that is characteristic of body-centered y cubic metals. The interstitial atoms become quite immobile at low (L% temperatures and prevent the dislocations from moving. Strain aging is another phenqmenon attributed to interstitial impurity afioms. This process is brought about by interactions between moving dislocations and mobile interstit%al solute atoms. Strain aging may be manifested in a "return of the yiéld point" in a tensile | specimen after interrupting a tensile test and aging the specimen or'by strengthening during a continuofis tefisile test with accompanying reduction in ductility and discontinuous yielding. The empirical relationship that has been determined for the. occurrence of discontinuous yielding is e =10° D, where ¢ is the strain rate, and D is the diffusion rate of the interstitial responsible for the - strain aging. This describes the condition for which the velocities of ‘moving disloca- tions and impurifiy atoms afe comparable. Strain aging is a relatively low-temperature phenomenon. For example, in niobium at strain rates of 1072 to 1074 sec™l, strain aging due to oxygen and the combined effects of nitrogen and carbon is observed over the temperature range of 200 to 450°C. % 9B. Longson and C. Tyzack, The Effect of Hydrogen on the Mechanical Properties of Niobium, TRG. Memo 880 (C), p. 6, March 1962. | The precipitation of a second phase has been used as a strengthen- ing mechanism in metals for some time. The general concept of strength- ening by this mechanism is that the second phase particles introduce strain fields that interfere with the motion of dislocations. In light of this mechanism, the concept of a critical particle size (or spacing) was proposed by Orowan. 10 Particles of sizes greater or smaller than this critical size are relatively ineffective. However, many complica- tions may arise that make this picture a gross oversimplification. The particles formed may or may not produce a strain field; they may or may not be coherent; they may have various shapes; and they may or may not deform plastically under stress. In fact, the critical particle size concept predicted'by Orowan has never been observed. The closest approach has been the observation of Meiklejohn and Skodal?l on the yield strength of solid mercury containing iron particles. However, a particle size effect was noted that canceled out the influence of particle spacing and gave the net result that the yield strength was a function only of the volume fraction of the precipitate. The silver-aluminum alloys referred to in the previous. section likewise showed only a slight dependence of strength upon aluminum-oxygen cluster size but exhibited a marked depend- 2 ence upon the volume fraction of the précipitate.l Recent transmission 2 on iron-gold alloys and by electron microscope studies by Horn’bogen1 Leslie et g;.13 on iron-bismuth alloys show that the second-phase par- ticles can act as copious sources of dislocations. The cell structure of dislocations originating from the particle offers more strengthening than would be predicted by the Orowan concept of strengthening. Another interesting effect is produced by cold working. Garofalol* pretreated 10E. Orowan, Discussion in Symposium on Internal Stresses in Metals and Alloys, p. 451, Institutc of Mctalc, London, 1948. 1ly. H. Meiklejohn and R. E. Skoda, Acta Met., 8: 773 (1960). 121,. S. Darken, Am. Soc. Metals Trans. Quart., 54(4): 600—42 (1961). 1 13W. C. Leslie et al., "Annealing of Cold Worked Iron," paper presented at the Metallurgical Society AIME Conference on High-Purity Iron and Its Dilute Solid Solutions, to be published. 14F. Garofalo, F. Von Gemmingen, and W. F. Domis, Am. Soc. Metals Trans. Quart., 54: 430 (1961). 10 type 316 stainless steel specimens by .solution annealing, cold working, and agifig. It was found that the stréngth was greatly improved by pre; treatments which resulted in fine randomly dispersed carbide formation. The resulting dislocation networks were studied and correlated with the mechanical properties of the steel. The desirable dislocation structure consisted of tangles that had apparently been trapped by the precipitate particles and the most undesirable structure was the stabilized cross grids of dislocations which offered little back force on dislocation fiotion. One particularly important factor concerning the mechanical prop- erties of refractory metals is the. influence of substitutional alloying element on the strength attainablé-throUgh the formation'of a dispersed phase. Thé case of the aluminum alloy addition to silver and the forma- tion of aluminum-oxygen clusters has already been discussed. One further - example is the influence of carbon on the properties of niobium. McCoy1? and Cortes and Feildl® have independently shown that carbon additions up to 0.21% do not result in measurable strengthening nor embrittlement of niobium. The niobium-base alloys F-44 (Nb—15% Mo—1% Zr—C) and F-48 (Nb—15% W—5% Mo—1% Zr—C) are, however,.strengthened by carbide dispersions as illustrated by the data of Chang!? given in Table 2. The formation of Nb,C in the latter alloy is due to the zirconium content being 0.6% instead of the nominal 1%. This illustrates the importance of the precise control of the zirconium-to-carbon ratio in these alioys. Besides being instru- mental in the initial formation of a strengthening dispersion, a solid- solution alloying element can affect the solubility of the precipitated °H. E. McCoy, Jr. , Conference on Corrosion of Reactor Matérials, June 4—8, 1962, Proceedings, Vol. I, pp. 263-94, International Atomic Energy Agency, Vienna, 1962. 16F. R. Cortes and A. L. Feild, Jr., J. Less-Common Metals, 4: 169 (1962). - 17w, H. Chang, p. 105, Refractory Metals and 209 > 189 12-13¢ >63 >56 38-41 B-66 (5% V—5% Mo~1% Zr) Good to excellent? ~1260¢ 38¢ 114 36 Cb-752 (10% W—2.5% Zr)¢ Good? g ~ 12609 26¢ 18-21 18 14 8 B-33 (4% V) Excellent? l 11754 209 D-31 (10% Mo-10% Ti) Pilot producuo}ab 1204° 50° 23-26° 11-20° 148 48 D-14 (5% Zr) fi 13709 550 28¢ 177 124 5d 39 16 D-36 (10% Ti~5% Zr) 34h 17k 140 C-103 (10% Hf-1% Ti) . 13159 187 ' SCb-291 (10% Ta~10% W) ) 1150-1315¢9 374 FS$-85 (27% Ta—10% W—1% Zr) Excellent? | 13709 29-60%-9 22-414.4 20-22° X-110 (10% W—1% Zr—0.1 C) Excellentd 13159 359 17.59 54 Nb—40% V | 980/ 33 (Rupture life of <13 hr at 102 psi and 1090°C) Pure molybdenum Sheet avauable’g 1425-1705% 21-24° 18-28° 10-20° 12-132 9144 Mo-0.5% Ti Sheet available;fr 1340 68 20-452 16-222 29-54¢ 12¢ 7¢ 210° 952 258 TZM (0.5% Ti—0.08% Z1) ’ 1325-17057 85 67-78° 50-552 38-80° 30-51¢ 35° 208 210° 90° 552 Mo—30% W 65° 35¢ Mo—25% W—0.1% Zr—0.05% C | 758 TZC (1.25% Ti-0.15% Zr-0.15 C) 15407 602 452 ~40° 30 20-30° Mo—50% Re K 85k 30% 20k Pure tantalum Tubing avaxlab’I;e 1090¢ 224 9-16° 8-167 ~6.5° 3f 34 -~11 5 5 Ta—10% W . 1370¢ 50-80%f 42-67f 40-45° 448 Ta-20% W g Ta-30% W * Ta-10% Hf-5% W 50-78¢ 46-60° 40-467 Ta—30% Nb-7.5% V° ‘ 1204/ 80° 622 422 Ta—8% W—2% HE i 15408 85¢ P 1204°C 1315°C 1650°C 1090°C 1204°C 1315°C 1650°C 1090°C 1204°C 1315°C 1650°C Pure tungsten [;\ 558 40-50° 20-30° 222 192 44 32 27 5.7 W—3% Mo }‘ ~188. W—30% Mo i ~ 302 W-1% ThO, { ~ 14008 38 W~2% ThO, i > 26908 ~42! >20° ~ 20! W-30% Re 135/ 50/ 4T. E. Tietz and J. W. Wilson, Mechanical, Oxidation, and ihermal Property Data for Seven Refractory Metals and Their Alloys, Lockheed Report, Code 2-36-61-1 (Sept. 15, 1961). bE. S. Bartlett and J. A. Houk, Physical and Mechanical Pmpert:es of Columbium and (‘qumbxum-Base Alloys, DMIC Report 125 (Fcb. 1960). €Alloys selected for study by AEC-NASA-AF Tubing Evaluatxon Committee. . 9AEC-AF-NASA Table on Niobium Alloys. ll €Creep-rupture data on the 0.5% Ti-Moé alloy at 535 to 1315°C from Climax Molybdenum Company, Sept. 1957. fAEC-AF-NASA Table on Tantalum and Vanadium Alloys. !| 8M. Semchyschen and J. J. Harwood, Refractory Metals and A,Iloys, Interscience, New York, 1961. BDu Pont Metal Products, Product Data Sheet No. 1, 1962. | B. R. Rajala and J. R. Van Thine, Improved Vanadxum-Base' ‘j{B R. Rajala and R. J. Van Thine, Improved Vanadmm-Base' kManufacturer s Literature, Chase Brass and Copper Company, Waterbury, Conn. . IB. S. Lement and 1. Perlmutter, ‘‘Mechanical Propertles Ath}mable by Alloying of Refractory Metals,’ b K 1 I I Alloys, ARF 2210-6 (Dec 20, 1961). Alloys, ARF 2191-6 (Dec. 27, 1960). ' p. 316, Niobium, Tantalum, Molybdenum, and Tungsten (ed. by A. G. Quarrell) Elservier, New York, 1961, 15 UNCLASSIFIED ORNL-LR-DWG 77766 o S Eg l e '&J NICKEL - ~TANTALUM @ w / Z 60 , 4 = / —TUNGSTEN | w g 40 ' / z NIOBIUM /A—MOLYBDE‘N‘UM / pzd © 20 : . — . 2 4// u//// a / lé':J 0 -300 -200 -100 0 {00 200 300 400 500 600 700 800 ’ - TEMPERATURE (°C) ’ ' Fig. 1. Effect of Temperature on Ductility. [L. Northcott, "Some Features of the Refractory Metals," p. 8, Niobium, Tantalum, Molybdenum, and Tungsten, (ed. by A. G. Quarrell) Elsevier Publishing Co., New York, 1961. ] : 16 with interstitial impurities. The F-48 and F-50 alloys have been studied by Chang.?? Both alloys were found to be age hardenable and the aging was attributed to carbide precipitation. Studies?3 of the No—1% Zr alloy have also shown it to be age hardenable under specific circumstances. The tensile properties of niobium are improved appreciably by the addition of vanadium. However, the creep properties are not improyed.24’25 This illustrates the fact that a tensile test is not a valuable screening teét for engineering materials. The range of values found in the litera- ture for the tensile strength of pure niobium at 982°C indicates the un- reliability of much of the mechanical property data on refractory metals. Molybdenum-Base Alloys Additions of titanium and zirconium appreciably improve the mechan- ical properties of molybdenum. Although this effect is often attributed to solution strengthening, it seems more reasonable that the strengthening is due to clusterihg.or dispersion strengthening caused by substitutional- interstitial atom interactions. Chang22 has studied the aging response of the MofTZC alloy and has clearly established the precipitation-hardenable nature of the alloy. Three dispersed phases were identified, consisting of TiC, Mo,C, and ZrC. The formation of TiC was prifiarily responsible for the aging, and ZrC was felt to have little inflfience on the strength. Chang suggested that another important role of the titanium was that of enhanéing the high-temperature solubility of carbon. Molybdenum, Mo-TZ, and Mo—0.5% Ti were found not to be age hardenable. ' The Mo—50 wt % Re (35 at. %) alloy has some very unique properties. Figure 2 compares the ductility of this alloy with that of pure molybde- 26 num. Note that the ductile-to-brittle transition temperature is signif- icantly lowered by the rhenium addition. This is due to the onset of 22y, H. Chang, A Study of the Influence of Heat Treatment on Microstructure and Properties of Refractory Alloys, Report No. ASD-TDR-62-211 (1962). 23p. o. Hobson, A Preliminary Study of the Aging Behavior of Wrought Columbium—1% Zirconium Alloys, ORNL-2995 (Jan. 1961). 24B. R. Rajala and R. J. Van Thine, Improved Vanadium-Base Alloys, ARF 2191-6 (Dec. 27, 1960). 25Tpid., ARF 2210-6 (Dec. 20, 1961). 261,. Northcott, "Some Features of the Refractory Metals," p. 17, Nicbium, Tantalum, Molybdenum, and Tungsten, (ed. by A. G. Quarrell) Elsevier Publishing Co., New York, 196l. 17 UNCLASSIFIED ORNL—-LR-DWG 77767 Mo —35 Re (RECRYST) S g L How >_' = / / Mo (RECRYST) / : / / / 3 | Mo (H.C.W.) 2142 wl 16 I I H.C.W.= HOT- COLD - WORKED ’ 20 -200 —100 O {00 TEMPERATURE (°C) Figs 2. Bend-Transition Curves for Molybdenum and the Mo—35% Re Aloy in the Recrystallized and Hot- and Cold-Worked (HCW) Conditions. [L. Northcott, "Some Features of the Refractory Metals," p. 17, Niobium, Tantalum, Molybdenum, and Tungsten, (ed. by A. G. Quarrell) Eleevier Publishing Co., New York, 1961.] 18 twinning in the molybdenum-rhenium alloy at low temperatures. This alloy . is also more resistant to oxygen embrittlement than pure molybdenum. In pure molybdenum the oxide phase accumulates in the grain boundaries, thus forming brittle grain boundary layers. The rhenium addition influences the surface energy of the oxide, and the oxide occurs as globules in the grains as well as at the boundaries rather than as a continuous grain boundary layer. The availability and cost of rhenium make the widespread use of the Mo—50% Re alloy doubtful.. In addition to the composition wariable that influences the proper- ‘ties of molybdenum, fabrication is also an important variable.?? This is illustrated in Fig. 3. Note the very large differences in room tempera- ture ductility depending on whether the final rolling temperature is 1204 or 1648°C. Significant strength differences also result. Tantalum-Base Alloys Pure tantalum is relatively weak at elevated temperatures. Additions of W, Hf, Nb, and V to tantalum result in significant strengthening. The ~influence of relatively low concentrations of oxygen and nitrogen on the elevated temperature behavior of tantalum has been investigated by Schmidt gz'gl.zs Additions of 560 ppm O and 225 ppm N were not effective strengtheners above 1100°C. However, carbon was an effective strengthener up to 1200°C. No systematic study has been made of strengthening due to interstitials when a substitutional alloying element is present. Chang29 has worked with a complex tantalum-base alloy of the nominal composition Ta—20% Nb—10% W5% V—1% Zr—0.08% C. This alloy was found to have a recrystallization temperature of 1704°C. Preliminary studies have shown that severe intergranular cracking occurred when annealed above 1648°C. Studies are continuing on this alloy. 27M. Semachyschen, R. Q. Barr, and G. D. McArdle, Effect of Thermal- Mechanical Variables on the Properties of Molybdenum Alloys, WADD-TR-60- 451 (Nov. 1960). ‘ 28F, F. Schmidt et al., WADD Report 59-13, p. 123 (Dec. 31, 1959). 29W. H. Chang, "A Study of the Influence of Heat Treatment on Microstructure and Properties of Refractory Alloys," Quar. Rep. No. 6, March 1, 1961, to May 31, 1961, DM62-140, pp. 10-13. 19 UNCLASSIFIED ORNL-LR-DWG 77768 % 1o Q O - o Q ~ 100 - ] (O] Z [e] & 90 - . o wn - ] w - ”)(\ 2200 °F = // \ 2 80 % M = L4 N w b3 \\ = N g ° S 5 o S~< 3000°F wl o SO % e — - 60 - a p= 3 @ 50 70 ,’5 (=N (o] S 60 w ) 8 50 ® = z 40 — 2200°F ——T° 5 .._-——-—-—‘—0- o = o . U) 30 r \\ w \ 7 Tee 3000°F =J ’ 4 "“——-___-4_.__-‘-- (%] Vd z 20 w - {0 S ~ 100 - < [ - . Bz oy T S = ® R T' " S ® studied the mechanical properties of commercially pure molybdenum irradiated in the MIR for an estimated exposure of 1.9 to 5.9 X 10°° thermal nvt. The specimen temperature was maintained at 90°C. The material used in this investigation was arc melted by the Climax Molybdenum Company. Two heats of material were used having carbon contents of 0.061 and 0.066 wt %. No other analytical details were given. The material was hot worked to 5/8-in. diam, annealed at 1100°C in hydrogen, and swaged to 1/2-in. diam. The impli- cation is that this last fabrication step was carried out at room temperature and represents a reduction in area of 36%. The test material had an average hardpess of 264 VPN (99.2 RB) and an average of 5000 grain/mm?®. The tensile specimens were rods having a gage section 1.00 in. long and 0.182 in. in diameter. The strain rate used in the tensile tests was 1.3 X 107% per second. The results of tensile tests performed in this program are summarized in Table 4. The unirradiated material was ductile at —20°C but was completely brittle at'—40 and —60°C. The irradiated material was completely brittle in tests conducted at room temperature and 60°C but was ductile at 80°C. Hence, the ductile-to-brittle transition 33y, E. Brundage et al., Solid Slule Div. Ann. Prog. Rcp. August 1961, ORNL-3213, pp. 124—33. 3%W. E. Brundage et al., Solid State Div. Ann. Prog. Rep. August 1962, ORNL-3364, pp. l44—45. 35¢. A. Bruch, W. E. McHugh, and R. W. Hockenbury, "Embrittlement ol Mulybdenum by Neutron Irradiation," Trans. ATME, 203: 281-85 (1955). Table 4. Tensile Properties of Mblybdenum? Integrated Thermel Upper b Neutron Test Yield Tensile Fracture Reduction Material Flux Temperature Point Strength Stress Elongation in Area Condition (nvt) - (°c) (psi) (psi) (psi) (%) | (%) x 102C . x10® x10° x 103 Unirradiated -- +22 102.5 100.8 214.0 45.7 72.4 Unirradiated -- +22 93.8 98.8 193.0 41,7 65.0 Unirradiated -- —20 125.5 120.0 . 243.0 32.8 63.8 Unirradiated -- —40 -- 123.0 123.0. 0 0 Unirradiated -- —60 -- 142.0 142.0 0 0 Aged® . - +24.6 9. b 97.7 182.6 40.8 67.4 Aged® - +24.6 94.0 9.3 181.6 42.5 65.3 Irradiated 5.1 +21.8 151.7 151.7 149.0 0 0.08 Irradiated 5.1 +22.4 -- 109.7 109.7 0 0 Irradiated 5.85 +60 -- 148.5 148.5 0 0 Irradiated 5.85 +80.5 143.5 143.5 185.0 14.7 60.7 Irradiated 5.8 +100 111.5. 111.5 134.0 10 59.7 ?C. A. Bruch, W. E. McHugh, and R. W. Hockenbury, "Embrittlement of Molybdenum by NeutronEIrradiation,” Trans. ATIME, 203: 281-85 (1955). bMaximum load divided by original area. “Unirradiated specimen heated for 30 days at 90°C. K4 25 temperature increased from about —30 to +70°C. Unirradiated and irradiated specimens were not tested at comparable temperatures at which each deformed plastically so that a meaningful comparison of the strength could be made. Bruch et al.?? do not make any comments con- cerning the relative values of the elongation and reduction in area. However, it seems that an important trend exists. 1In the unirradiated specimens, the uniform elongation is over one half the value of the reduction in area. In the irradiated specimens, the elongation is only one fifth to one sixth the wvalue of the reduction in area. A possible explanafion of this observation is that the irradiation-igduced detf'ects pin the dislocations in the metal so that the stress to cause plastic deformation is quite high. When this stress is exceeded, the dis- locations breask away from their pinning defects with such driving force that the normal processes of work hardening are ineffective. Hence, failure occurs with very high local deformation and very small unifdrm elongation. Several metallographic specimens were included in the tests of Bruch et §£.35 They were polished and photographed at points marked wifh hardness impressions before insertion into the experiment. Photo- graphs were made of the same fields after irradiation without further polishing. It was concluded that no visible metéllographic changes occurred. ‘ The results of hardness tests performed before and after irradi- ation are given in Table 5. The hardness increased by approximately 35 BHN as a result of the irradiation. In the last paragraph of their pafier, the authors ingerted some additional data concerning irradiation hardening. Few experimental detaiis are given. Opecimens were irradi- ated at 400°C for an estimated 3 X 10%°0 thermal nvt (3 X 10%° fast nvt) and found to inereage® in hardness froam 169 ta 216 BHN (converted from RB values). These results indicate that the-defects are introduced by the irradiation at a rate greater than they can be annealed out at 400°C. Table 5. Hardness Test Resultsa . b \ b e RC RA "Material Average o Average o Condition Value ° Minimum Maximum BHN Value Minimum Maximum BHN Unirradiated 23.0 - 19.7 24.9 242 62.0 59.9 63.0 243 Irradiated 28.5 24,2 31.7 275 64.9 62.8 66.5 280 Change +5.5 . +3.0 +7.7 433 +2.9 +0.5 +h 4 +37 8c. 5. Bruch, W. E. McHugh, and R. W. Hockenbury, "Embrittlement of Molybdenum - by Neutron Irradiation,"” Trans. AIME, 203: 281-85 (1955). Each number in the table represents the average of test results for 12 specimens. At least three Rockwell hardness measurements of each kind were made per specimen. ‘ CConverted from the Rockwell numbers. 9¢ 27 Makin and Gillies?® inveétigated the effects of neutron irradi- ation on the mechanical properties of molybdenum. The test material was obtained from the Johnson, Matthey and Company, ILtd., in the form of 0.040-in.-diam wire. A comélete spectrographic analysis of the material was given, but no mention was made of interstitial impurities. The wires were given a stress-relief anneal at 1030°C for 30 min. The specimens were irradiated for six months in a Windscale pile at approxi- mately 100°C. The flux was 6 X 10%% thermal nv and the integrated thermal flux was 5 X 10'% nvt. The ratio of fast to thermal neutrons was estimated to be unity. Tensile tests were run on a Hounsfield Tensometer at a strain rate of 8.2 X 1077 per second. The ductile-to- brittle transition temperature was determined by bend tests on the 0.040-in.-diam wires. Specimens were defined as ductile when they could be bent .90° around a pin of 6-mm diam without fracture. The strain rate in the bend tests was designated as "slow." Eight specimens were used to determine each transition temperature to a reported accuracy of *2°C. The results of postirradiation tensile tests over the temperature range of 20 to 200°C are given in Table 6. Four specimens were tested at each condition. The yield stress was increased by irradiation over the entire range of test temperatures, the effect becoming more pro- nounced with increasing temfierature. The ultimate strength changed in a similar mahner. Yield points'were observéd in the irradiated and unirradiated specimens tested at 20°C. However, the drop in stress associated with the yield point of the irradiated material was the greatest and, at 83°C, only the irradiated material showed a yield point. Neither materiai exhibited a yield point at the 200°C test temperature. - The elongation at rupture was, in general, decreased slightly by the irradiation. The 200°C test condition was an exception with slightlytgreater elongation occurring in the irradiated specimen. However, the elongation of both materials was quite low. 36M. J. Makin and E. Gillies, "The Effect of Neutron Irradiation on the Mechanical Properties of Molybdenum and Tungsten," J. Inst. Metals, 86: 108-12 (1958). Table 6. Tensile Tests on Stress-Relieved MOlybdenum? Test 8¢ Temperature ~Material Yield Sbressb Ultimate Strengtfib ‘Elongationb (°C) Condition (psi) | (psi) ' (%) . x10° - x 102 | 20 Irradiated 95.8-102.2 (99.4) 99.5-106.8 (104.3) 20.5-24.3 (22.0) 20 Unirradiated 90.6-95.6 (93.7) = 96.0-102.3 (99.8) 20.0-26.7 (23.6) 83 Irradiated - 93.3 . 93.5 18.5 97 Unirradiated 80.3 ‘ 90.5 23.8 200 Irradiated 85.5 ' 85.9 5.8 200 E Unirradiated 68.5-72.2 (70.4) 7. 6774.6 (74.6) 2.7-2.8 (2.8) ®M. J. Mekin and E. Gillies, "The Effect of Neutron Irradiation on the Mechanical Properties of Molybdenum and Tungsten," J. Inst. Metals, 86: 10812 (1958). bLimits of experimental results given with average values in parentheses. 29 Bend tests showed that the irradiation dose of 5 X 10%° nvt raised the ductile-to-brittle transition temperature from —136 + 1°C to =73 + 1°C, a rise of about 63°C. Attempts were made to study the recovery characteristics of the radiation effect by annealing treat- ments, but the complexity of the process coupled with the small number ' of specimens prevented conclusive results from being obtained. Makin and Gillies3® explained their results in terms of both impurity atoms ~and irradiation-produced defects. It was postulated that the influence of the defects produced by irradiation was not fully observed until the test temperature was increased to about 200°C. Since quite large yield points could be produced by postirradiation annealing at 200°C, it was concluded that the defects were mobile at this temperature. Studies of recovery in cold-worked molybdenum in a radiation 7 revealed a recovery stage at 150°C. field by Kinchin and Thompson3 From the activation energy of 1.3 ev, the process was felt to be vacancy migration. This observation led Makin and Gillies to conclude that vacancies were the important defects in their specimens. Makin and Gillies discuss their results in light of those obtained by Bruch et §£.35 The former authors observed the transition tempera- ture to increase from —136 to —73°C (+63°C) after a dose of 5 X 101° nvt. The latter authors reported an increase from —30 to +70°C (+100°C) after a dose of 1.9 to 5.9 X 1029 nvt. It was concluded by Makin and Gillies that the magnitude of the increase in the transition temperature was not proportional to the neutron dose. This seems a rather dangerous con- clusion in light of the different metallurgical histories of the test material, possible chemical differences, use of 0.04U-in.-diam wires vs 0.182-in.-diam rods as specimens, and uncertainties in flux measurements. " Another conclusion based upon this comparison was that "...it is possible that the greatest effect is produced in materials possessing initially the lowest transition temperature." 1In light of the uncertainties just mentioned, this conclusion is not supported by the available data. Although Makin and Gillies seemed to realize this, this has been passed on through the literature as a general rule. 37G. H. Kinchin and M. W. Thompson, "Irradiation Damage and Recovery in Molybdenum and Tungsten," J. Nucl. Energy, 6: 275-84 (1958). 30 -, Bruch et §£.38 discuss the general problem of irradiation damage in a later papér. In this paper, an attempt was made to assess the relative effects of irradiation on high-purity copper, nickel, zirconium, and iron, commercial grade 75A titanium, commerical-purity molybdenum, and cold-worked and annealed type 347 stainless steel. The data on molybdenum presented in this paper are the same as presented earlier. However, some general statements were made which have bearing upon the properties of molybdenum. Although it is not possible to calculate the number of vacancy interstitial pairs which exist in a metal in a radia- ‘tion field at a given time due to complex annealing and other inter- actions, reasonable models exist for calculating the total number of sfich pairs that have been produced by a given flux. Such calculations for the metals mentioned above showed that only minor differences existed in the number of pairs produced in the metals. Hence, it was concluded that the very large differences in the effect of comparable doses on the properties of these metals were due to factors such as the annealing of defects or variations ambng metals in the property change produced by a given defect. Tantalum Franklin‘gz gl.Bg have run a limited number of postirradiation tests on.tantalum and two tantalum-tungsten alloys to evaluate the effect of irradiation on the mechanical properties. Besides the lattice - defects normally produced by irradiation, tantalum is also converted to tungsten by the thermal neutron reaction Ta8l(n,7)Ta'®2(—8)W182. The - intrinsic effects of the two processeé were evaluated by comparing the properties of the irradiated specimens with those of control specimens containing comparable amounts of tungsten. The analysis of the test material is given in Table 7. Sheet specimens were used which had a 38¢. A. Bruch, W. E. McHugh, and R. W. Hockenbury, "Variations in Radiation Damage to Metals," Trans. AIME, 206: 1362—72 (1956). 39C. K. Franklin, D. Stahl, F. R. Shober, and R. F. Dickerson, Effects of Irradiation on the Mechanical Properties of Tantalum, BMI-1476 (Nov. 18, 1960). 31 Table 7. Chemical Analysis of the gantalum.and Tantalum-Tungsten Alloys®’ Analyses for Indicated Specimen (ppm) Unalloyed Tmpurity Tantalum Ta~l.5 wt % W Ta—3.0 wt % W Aluminum - <5 15 . 20 Chromium - 10 4 5 Copper 10 20 . 15 Iron. 3 6 15 Molybdenum -- 20 _ , 15 Niobium -- 300 100 Nickel ~- 6 3 Silicon -- 30 60 Zirconium -- 15 | 10 Nitrogen < 10 20 35 Carbon 10 20 35 Hydrogen ' 0.3 1 | 1 Oxygen 40 53 22 aA.verage of two analyses taken of the alloys after they had been cold rolled to 0.030-in. strip and vacuum annealed. bC. K. Franklin et al., Effects of Irradiation on the Mechanical Properties of Tantalum, BMI-1476 (Nov. 18, 1960). gage cection 1.00 in. long, 0.75 in. wide, and 0.030 in., thick. The tantalum specimens were annealed at 1371°C and the alloy specimens were annealed 2 hr at 1426°C prior to test. The specimens were irradiated in the MIR at a thermal flux level of 4 to 5 X 10%% nv and a tempera- ture of 93°C. Two capsules were irradiated having total thermal doses of 8.6 X 1029 and 1.8 x 10?! nvt. Chemical analyses showed that the tungsten content of the six irradiated specimens varied from 0.6 to 2.2 wt %. The room-temperature tensile properties of the irradiated specimens are compared with those of unirradiated specimens having comparable chemistry in Table 8. The strain rate was 5 X 107° per minute. Table 8. Room-Temperature Mechanical Properties of Unirradiated and Irradlated Tantalum and Unirradiated Ta—1.5 wt % W and Ta—3.0 wt % W Alloys? Total Integrated Average of Properties Tests Thermal Flux 0.2% Based on Ultimate Offset Elongation Number of Dosimetry Tensile Yield at Maximum Elongation Hard- Specimen Specimens .Analysis Strength Strength Load in 1 in. . ness Description Tested (nvt) (psi) (psi) (%) (%) KHN - Unirradiated 4 -- 42,000 30,000 -- 40 103 tantalum . Unirradiated 3 -- 44, ;900 31,000 -- 39 151 Ta—1.5 wt % W | Unirradiated 3 - 52,400 38, 500 -- 35 170 Ta~3.0 wt % W | Irradiated 4 7.8 x 10°° 69, 500 65,800 -- 16 274 tantalum - (66-day irradiation) , Irradiated 4 1.57 x 10%% 86,300 81,400 -- 7 309 tantalum (109-day irradiation) 8. K. Franklln, D. Stahl, F. R. Shober, and R. F. Dickerson, Effects of Irradiation on. the Mechanical Properties of Tantalum, BMI-1476 (Nov. 18, 1960). ce 33 The yield strength of the irradiated specimens increased nearly three times over that of the unirradiated tantalum. The properties of the arc-melted alloys showed that the increase in strength could not be accounted for in terms of the tungsten content since only minor strength changes resulted from the addition of up to 3 wt % W. Hence, the change in strength was attributed to some mechanism of defect production by neutrons, Table 8 also illustrates the marked reduction in ductility which occurred as a result of irradiation. However, the data indicate that a large portion of the ductility effect can be attributed to the tungsten content of the specimens. Tantalum showed a large increase in hardness after irradiation. The hardness changed from 103 to 309 KHN after an integrated thermal dose of 1.8 X 1021 nvt. The hardnesses of the tantalum alloys containing 1.5 and 3.0 wt % W were 151 and 170 KHN, respectively. 40 reported the results of a limited number of Sutton and Ieeser room—tempefature postirradiation tensile tests on tantalum. No experi- mental details other than neutron dose were given. The results of these tests are given in Table 9. Irradiation appears to raise the ultimate strength and decrease the rupture ductility. These authors also reported that the hardness of tantalum was increased 8 points to 57 RA by a neutron dose of 1 X 101° nvt thermal and 5 X 10'° nvt fast. Tungsten The effects of bombardment by 13.7-Mev deuterons on the internal friction and Young's modulus of polycrystalline tungsten were studied d.4' TIrradiations were carried out at 300°%K using by Muss and Townsen the University of Pittsburgh cyclotron, and the effects of integrated doses of 5 X 1014 to 8 x 101° deuterons/cm2 were studied. Because of 400, R. Sutton and D. O. Leeser, "Radiation Effects on Keactor Materials," Nucleonics, 12: 8-16 (1954). 4lp, R. Muss and J. R. Townsend, "Internal Friction and Young's Modulus in Irradiated Tungsten,"” J. Appl. Phys., 33: 180407 (1962). 34 - Table 9. Postirradiation Tensile Properties of Tantalum at Room Temperature Ultimate Integrated Neutron Dose Tensile (nvt) ‘ Strength Elongation Thermal Fast (1b/in.?) (%) x 1019 X 10%° x 103 O 0 72.0 19 0 0 65.0 23 i 1 5 88.0 17 ) 1 5 g5.0 17 the nature of the bombarding particles, the number of defects produced by deuterons is approximately 1000 times greater than the number pro- duced by a comparable dose of neutrons. The test material was 0.0015-in.-diam tungsten wire supplied by the Sylvania Company. The wire was designated as type NS-50, but no composition was givén. The modulus and internal friction measurements were made using a mechanical resonance system consisting of the tungsten wire mounted as a cantilever. Preirradiation internal friction meaéurements as.a function of temperature showed that a reproducible peak occurred at 140°%K. This peak was found to occur at higher temperatures as the frequency of vibration was increased. This 1s a characteristic of a relaxation phenomenon and it was'fiBStUlatédfthat this peak was due to the thermal - activation of dislocations. The data of this study and those of Chambers and Schultz*? were combined to obtain an activation-enefgy for the process of 0.21 + 0.05 ev. The effects of integrated dose rate-op tungsten were'evaluated at an irradiation temperature-of 295°%K. The internal friction decreased and Young's modulus increased. Both of these effects were found to fit 42R. H. Chambers and J. Schultz, "Dislocations Relaxation Spectra of Cold-Worked Body-Centered Cubic Transition Metals,'" Phys. Rev. Letters, 6: 273 (1961). 35 the model of a dislocation pinning mechanism that was proposed by 43 The pinning in this case was concluded to be Dieckamp and Sosin. due to interstitials. This conclusion was based on the work of Kinchin and Thompson** which showed that radiation-induced interstitials begin to anneal out at 80°K whereas vacancies become mobile at 650°%K. The internal friction effect began to saturate at an integrated flux of about 2 X 1016 deuterons/cmz. The elastic modulus went through an inflection at approximately the safie dose and decreased linearly with incrcasing dose. This decrease was explained in terms of the bulk effect of vacancies being frozen into the lattice. This effect amounted quantitatively to a 0.44% decrease in Young's modulus per atomic percent vacancies. The amplitude of the internal friction peak at 140°%K and the background internal‘frictjon were found to. decrease as a result of irradiation. This was attributed to dislocation pinning. However, after irradiation several small peaks appeared in the internal friction spectrum that were not satisfactorily explained. Makin and Gillies“® studied the effects of neutrofi irradiation on the mechanical properties of tungsten. The material used in this investigation was obtained from the Johnson, Matthey and Company, Ltd., in the form of 0.040-in.-diam wire. A complete spectrographic analysis of the material was given, but no mention was made of interstitial impurities. The specimens were fully recrystallized by annealing 30 min at 1600°C. The specimens were irradiated for six months in a Windscale pile at approximately 100°C. The flux was 6 X 10'? thermal nv and the total integrated thermal flux was 5 X 102 nvt. The ratio of fast to thermal neutrons was not known accfirately but was estimated Lo be unity. Control specimens were annealed at an equivalent temperature and time. “3H. Dicckaipy and A. Sosin, "Effect of Flectran Irradlatlon on Young's Modulus,"” J. Appl. Phys., gz 1416 (1956). 44G. H. Kinchin and M. W. Thompson, "Irradiation Damzge and Recovery in Molybdenum and Tungsten,'" J. Nucl. Energy, 6: 275-84 (1958). 45M, J. Makin and E. Gillies, "The Effect of Neutron Irradiation on the Mechanical Properties of Mblybdenum and Tungsten," J. Inst. Metals, 86: 108-12 (1.958). 36 Tensile tests were run on a Hounsfield Tensometer at a strain rate of 8.2 x 1072 per second. The ductile-to-brittle transition temperature was determined by bend tests on the 0.040-in.-diam wires. Specimens were defined as ductile when they could be bent 90° around a pin of 60-mm diam without fracture. The strain rate was given as "slow." Eight specimens were used to determine each transition temperature and the reported accuracy is *2°C. | _ The results of' tensile postirradiation tests at 100 and 200°C are compared with the results obtained from control specimens in Table 10. Although no statement is made of the exact number of specimens tested, the implication is that these are average values. Both the irradiated and unirradiated specimens were brittle at 100°C. However, the fracture stress was raised by the irradiation. At 200°C the yieid strength was increased by irradiation but the ultimate strength was unaffected. The elongation at rupture and reduction in area at 200°C were increased by irradiation. Smooth stress-strain curves were obtained with no yield points being observed. The ductile-to-brittle transition temperature was increased from 118 # 2°C to 126 * 2°C by the irradiation. \ Table 10. Tensile Tests on Recrystallized Tungstena Ultimate ‘ ' Test Tensile . Material Temperature Yield Stress Strength Elongation Condition (°C) (1b/in.?3) (1b/in.3) (%) x 103 X 103 Irradiated 100 152.0 - 0 (fracture stress) Unirradiated 100 137.0 ' -- 0 (fracture stress) Irradiated 200 | 131.0 ~ 173.0 4.2 Unirradiated 200 ‘ 148.0 ©173.2 2.4 M. J. Makin and E. Gillies, "The Effect of Neutron Irradiation on the Mechanical Properties of Molybdenum and Tungsten," J. Inst. Metals, 86: 108-12 (1958). 37 40 reported that the room-temperature ultimate Sutton and Leeser tensile strength of tungsten increases 20 to 25% after irradiation with 5 X 10'° fast neutrons/cm®. The reported data (given in Table 11) do not seem to support the statement made by Sutton and ILeeser. All . specimens were tested below the ductile-to-brittle transition tempera- ture; therefore, all were brittle. No'additional experimental details were given. Table 11. Postirradiation Tensile Properties of Tungsten at Room Temperature Ultimate Integrated Neutron Dose Tensile (nvt) Strength Elongation Thermal Fast (1b/in.?) (%) x 10%° X 10%° x 103 0 O ' 145.5 O 0 0 160.0 0] 1 5 132.0% 0 1 5 102.0% 0 ®Decrease can be attributed partly to difficulty of alloying brittle specimens by remote control. SUMMARY The available data on the mechanical behavior of niobium, molyb- denum, tantalum, and tungsten have been reviewed critically. Several important conclusions have been reached as a result of this study — the most important one being that insufficient engineering data are available for the design of complex systems using refractory metals and structural materials. It was also found that each of the four metals reviewed has certain unique properties that make it desirable for specific application. Molybdenum and tungsten have low coefficients of thermal expansion which may more nearly match those of cermets and ceramic components. These materials also have high moduli of elasticity which are desirable / 38 from a design standpeoint. -However, both of these materials present problems with respect to fabricatility. Niobium has a very low modulus . of. elasticity — an undesirable feature. Niobium and tantalum are relatifiely easy to fabricate and have good ductility. Niobium and molybdenum have low neutron absorption cross sections, whereas tantalum and tungsten are an order of magnitude higher. Because of the complexity of high-temperature nuclear systems, materials are needed that have a variety of properties. Hence, at this stagé of refractory-metal technology it is important not to limit our studies to those metals that can be fabricated into tubing or those metals that can be welded, since there may be applications for which such materials can be used in different parts of the reactor. Also, service conditions such as stress, temperature, temperature cycle, and desired nuclear properties must be known. _ In order to expedite the development of the technology necessary for the use of refractory metals in engineering systems, it is felt that persons involved in evaluating the mechanical behavior of these metals should consider the following factors. 1. A lot of deception is being injected into this field by the production of small melts of alloys and by the evaluation of these alloys by short-time tensile tests. These small melts often are made under nonrefiroducible conditions and are fabricated by unknown pro- cedures. Unless the application is one requiring a short life, the use of short-time tensile tests for screening purposes cafi be very deceiving. Short creep tests of 10 to 100 hr duration are better screening tests for materials to.be used in long-time applications. 2. More attention needs to be given to defofmation mechanisms in refractory metals. Just as the strength of many superalloys far ‘exceeds that of pure iron .and nickel, so can the properties of refrac- tory superalloys excel those of the pure refractory metals if one learns more about the deformation mechanisms in refractory metals. Fabrication procedure and impurity content can be used to an advantage if understood. Dispersed phases may possibly be found important in these alloys. It may be found that these dispersions improve the 39 strength by impeding dislocation motion as well'as serving as sources for dislocations in normally brittle materials. Hence, the production of ultra-pure alloys by electron-beam melting may not be the most practical approach to ductile molybdenum and tungsten. 3. The more promising alloys need to be evaluated with respect to their strength and metallurgical stability over long periods of time. 4. Experiments should be carried out to evaluate the behavior of refractory metals under neutron irradiation at ambient and elevated temperatures. These studies should be directed toward defining mecha- nisms responeible for the different mechanical behavior under irradia- tion and hence would make it possible to design alloys which are not greatly altered by irradiation. | THIS PAGE WAS INTENTIONALLY - LEFT BLANK . 52. 56—58. 48, 41 INTERNAL DISTRIBUTION Central Research Labrary 68. Reactor Division Library 69, ORNL — Y-12 Technical Iibrary 70, Document Reference Section 71. Laboratory Records Department 72. Iaboratory Records, ORNL R.C, 73. ORNL Patent Office 74—78. R. E. Adams 79. S. E. Beall 80. R. J. Beaver 81. R. L. Bennett 82. R. G. Berggren 83. J. O. Betterton, Jr. 84. E. G. Bohlmann 85. N. H. Briggs 86. J. Burka 87. G. W. Clark 88. R. E. Clausing 89. J. A. Conlin 90. W. H. Cook 91. G. A, Cristy 92. J. E. Cunningham 93, J. H. DeVan 9%4. J. R. DiStefano 9599, R. G. Donnelly 100. W. S. Ernst, Jr. 101. S. T. Bwing 102. J. 1. Federer 103, H. A. Friedman 104. J. H Frye, Jr. 105. W. R. Gall 106, R. G. Gilliland 107, A. Goldman 108. K. W. Haff 109. M. R. Hill 110. N. E. Hinkle 111, D. O. Hobson 112. H. Inouye 113-115, D. H. Jansen 116. G. W. Keilholtz 117. R. B. Korsmeyer 118. C. E. Iarson 119, A. P. Litman 120. R. A. Iorenz ORNL~3593 UC-25 — Metals, Ceramics, and Materials TID-4500 (27th ed.) MOUPTHEPRQEOUIUQI LU IARORNARENRPpHOSQRESSHway Iotts Tundy Tyon MacPherson Manly ' Martin McCoy McDonald McHargue Mixon Mossman Peishel Perry Quist Rabin Reed Roche Rosenthal Samuels L. Senn Sisman Slaughter Stelzmann Stephenson Stiegler Strehlow Swartout Taboada Thurber Tolson Trauger Vandermeer Venard Wantland Watson Wechsler Weinberg Weir Wichner Burr (consultant) Johnson (consultant) Smith (consultant) Smoluchowskl ERrppREfpIoHRIDQRNH PPOoP&ER MAPUIRRIPRA>EEQ (consultant) 42 EXTERNAL DISTRIBUTION _ . -Adams, Jr., Massachusetts Institute of Technology 121. C. M 122. G. M. Anderson, U. S. Atomic Energy Commission, Washington, D.C. 123. D. E. Baker, General Electric Company, Hanford 124. 8. 8. Christopher, U. S. Atomic Energy Commission, Washington, D.C. 125-126. D. F., Cope, Oak Ridge Operations Office _ 127. E. M. Douthett, U. S. Atomic Energy Commission, Washington, D.C. 128. Ersel Evans, General Electric Company, Hanford 129. J. L. Gregg, Cornell University 130. T. W. McIntosh, U. S. Atomic Energy Commission, Washington, D.C. 131. R. G. Oehl, U. S. Atomic Energy Commission, Washington, D.C. 132. F. C. SBchwenk, U. S, Atomic Energy Commission, Washington, D.C. 133. J. Simmons, U. S. Atomic Energy Commission, Washington, D.C. 134. E. E. Stansbury, University of Tennessee 135. D. X. Stevens, U. S. Atomic Energy Commission, Washington, D.C. 136. Research and Development, Oak Ridge Operations Office 137. G. W. Wensch, U.~S. Atomic Energy Commission, Washington, D.C. 138. M. J. Whltman, U. S. Atomic Energy Commission, Washington, D.C. 139—710. Given distribution as shown in TID-4500 (27th ed.) under Metals Ceramics, and Materials category (75 copies — OTS)