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The effects of heat treatment and purity on the mechanical properties of monocrystalline NiAl
ELSEVIER
Materials Science and Engineering
A192/193
(1995) 333-339
The effects of heat treatment and purity on the mechanical properties of monocrystalline NiAl Ha K. DeMarcoa, Alan J. Ardell”, Ronald D. Noebeb aDepartment of Materials Science and Engineering, University of California, Los Angeles, CA 90024-1.595, USA ‘NASA Lewis Research Center, MS. 49-3, Cleveland, OH 441353191, USA
Abstract The yield stress, 4, strain to fracture, q, and thickness-compensated fracture load of monocrystalline NiAl were investigated as functions of purity, annealing treatment and cooling rates. A miniaturized disk bend test was employed to test disks 3 mm in diameter and approximately 250 pm thick. Specimens in the soft (110) orientation from commercial purity (CP) and high-purity (HP) alloys were tested. The annealing treatments affected uYmore strongly than the other mechanical properties. The yield stresses of both of the as-received alloys increased significantly after annealing at 1300 “C followed by furnace cooling. Subsequent annealing at 400 “C for 2 h resulted in a reduction of uY.This behavior is attributed to the role of excess vacancies retained during cooling, which annealed out at the lower temperature. The mean values of the yield strength also tended to decrease with increasing cooling rate from 400 “C, but the effect was small. A substantial increase in uY of the CP alloy was found on prolonged aging at 400 “C, whereas aY of the HP alloy was unaffected. We attribute this behavior to the precipitation of very small particles or solute atom clusters in the CP alloy. The ductility of the CP alloy, and the thickness-compensated fracture loads of both alloys, were relatively insensitive to the heat treatments. Keywords: Heat treatment;
Mechanical
properties;
Nickel; Aluminium;
1. Introduction There is considerable current interest in ordered intermetallic compounds as candidate materials for high-temperature structural applications. NiAl, with its
high melting temperature and low density, has received special attention, as evidenced by recent comprehensive reviews of the literature on its physical and mechanical properties [1,2]. However, the low roomtemperature ductility and fracture toughness of monolithic NiAl continue to impede its utilization as a structural material. Consequently, a substantial amount of research has focused on improving the ductility of both polycrystalline and monocrystalline NiAl by microalloying additions [3-91. This effort has been aided by the development of techniques [ 10,l l] for growing high-purity single crystals, which enables investigators to determine the mechanical properties of NiAl containing very low impurity levels. The widely accepted values for the tensile strains to fracture for monocrystalline NiAl in the “soft” (110) orientation range from 1 to 3% [ 1,121, but it has been 0921-5093/95/$9.50 0 1995 - Elsevier Science S.A. All rights reserved SSDZ 0921-5093(94)03215-7
Crystals
demonstrated recently that the inherent tensile ductility of NiAl may be as high as 7-8% strain to failure [ 131. Brzeski et al. [ 131 and Hack et al. [ 141 have presented experimental evidence that the cooling rate from 400 “C strongly affects the fracture toughness. They concluded that a strain-aging effect at low temperatures ( -200 “C), involving either interstitial impurities or constitutional vacancies resulting from slight deviations from stoichiometry, hinders the room-temperature mobility of dislocations. Field et al. [15] have reported the occurrence of serrated yielding in monocrystals tested in compression at several temperatures. Their observations are consistent with dynamic strainaging effects involving interactions between point defects and dislocations [ 161, and support the idea that strain aging influences the mechanical behavior of NiAl. Vacancies and alloy purity also affect the mechanical properties of NiAl. Weaver et al. [ 171 found that annealing commercial-purity (CP) and high-purity (HP) NiAl at 727 “C, followed by a furnace cool of -0.083 K s-l, decreased the compressive yield
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Materials Science and Engineering AI921193 (1995) 333-339
that plastic deformation of discs in the (100) orientation does not occur by bending, but by a kinking process that probably involves slip on { lOO} [20,21]. This is because ( 100) is a hard orientation for the disks, and normal plastic bending would require (111) slip. Results are therefore presented only for the specimens deformed in the ( 110) orientation. Each crystal of NiAl was oriented using the Laue back-reflection method, and slices averaging 0.7 mm in thickness were cut parallel to ( 110) using an electricspark discharge machine. The slices were electropolished after cutting to remove surface damage and contaminants, and subjected to several different heattreatment conditions: (1) as-received material (this is designated as the AR condition); (2) annealed at 400 “C for 70 h, then furnace-cooled at an average rate of 0.023 “C SK’ (AR+ 400 [70 h] + FC); (3) annealed at 1300 “C for 36 h in flowing ultra-highpurity (99.995%) argon and then furnace-cooled at an initial rate of 0.12 “C s- ’( 1300 + FC); (4) as in (3), but then annealed at 400 “C for 2 h and either furnace-cooled at a rate of 0.023 “C s- ‘, aircooled or water-quenched (these are designated as 1300 + 400 [2 h] + FC, AC or WQ respectively); (5) as in 3, but annealed at 400 “C for 70 h and furnace-cooled ( 1300 + 400 [70 h] + FC). Disks 3 mm in diameter were cut from the heattreated slices using an abrasive slurry cutter. Each disk was then ground on each side to a thickness of -250 pm using silicon carbide abrasive papers, then metallographically polished to a final finish using 0.05 pm alumina slurry. The orientation of each individual disk was not verified, but it is assumed that the planes of the disks were within 2” of ( 110) according to the orientations of several disks that were checked using electron channeling patterns or Laue back-reflection X-ray patterns. The MDBT apparatus is described in detail elsewhere [22-241. The tests were conducted in the ringon-ring mode using a table-model Instron testing machine at a cross-head displacement rate of 8.3 pm s- I. Despite the elastic anisotropy exhibited by
strength of both alloys. The resulting yield strength of the HP alloy was lower than that of the CP alloy. Jayaram and Miller [ 181 observed fine carbide precipitates in C-doped NiAl, and Duncan et al. [ 191 observed inhomogeneously distributed ultrafine precipitates in the CP alloy, as well as an occasional precipitate in HP crystals. The existence of these ultrafine preciptiates may affect the yield strength and ductility of NiAl. It seems safe to conclude that the combined effects of alloy purity and thermal maniulation on the mechanical properties of monocrystalline NiAl are complex and incompletely understood. In this paper we report the results of an investigation of the effects of heat treatment, cooling rates, and alloy purity on yield strength, bending strain to fracture and fracture load of NiAl using a miniaturized disk-bend test (MDBT). The advantage of the small specimen size utilized by the MDBT, besides the obvious ability to conduct many tests on a small quantity of material, is that small specimens are rapidly equilibrated thermally, and can therefore be cooled relatively rapidly without acquiring large residual stresses.
2. Experimental Monocrystalline NiAl of two different purities was investigated: a CP crystal grown by GE Aircraft Engines using the Bridgman method and a HP crystal grown at the University of Tennessee using a containerless processing method [lO,ll]. The as-received CP alloy had already been subjected to a homogenizing anneal at 13 16 “C for 48 h and cooled at a rate of -0.213 “C s-‘. The HP alloy was provided in the asgrown condition, with no post-growth processing treatment. Chemical analysis of the CP crystal used in this study was not performed, but the typical impurity levels in identically grown and processed GE NiAl monocrystals are listed in Table 1. Chemical analysis of the HP crystal was performed at NASA Lewis Research Center; details of the analytical techniques used are described by Weaver et al. [ 171. Experiments were performed on disks in both (100) and ( 110) orientations. However, we have discovered
Table 1 Interstitial
and Si content (at.%) of typical NiAl alloys
Alloy
Si
C
0
N
S
Comments
HP CP
0.06 0.15
< 0.0036 0.0306
< 0.0027 0.0155
< 0.003 1 0.0036
< 0.0013 < 0.001
Same ingot as Weaver et al. [ 171 Average C, 0, N values from Hack et al. [ 141; Si and S from[17]
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Materials Science and Engineering A1921193 (1995) 333-339
NiAl, the yield stresses, oY, were calculated using the formula [25], valid for elastically isotropic materials,
oy=g+
(I+.),,;+~$ 1
[
?? HP
CP
1-c II (1)
335
a*
where Py is the applied load at yielding, t is the thickness of the specimen, Y is Poisson’s ratio, a and b are the radii of the support die and loading ring respectively, and R is the radius of the specimen. The values of the parameters used in the calculations were a=0.955 mm, b=0.36 mm, R= 1.5 mm and ~=0.31 [1,2]; t varied from one specimen to another, but was typically 250 pm. In addition to aY, we report the thickness-compensated load at fracture, Pr/t2, where P, is the load at fracture. This quantity provides an estimate of the relative fracture stresses of the various specimens, Representative curves of P/t* vs. displacement of the center of the specimen, W, are shown in Fig. 1 of the paper by DeMarco and Ardell[26]. The strain to failure, sr, was also measured in this investigation, and was estimated using the formula [27]
(2) Since az % t*, a useful measure of the strain is simply yt/a*. In all the measurements, wr was taken as the difference between the displacement at fracture and that at yielding.
Fig. 1. Yield stresses, heat treatments.
u,,, of the HP and CP alloys for the various
CP
H
HP
2.5
0.5 0
3. Results The data on aY, sf and Pr/t2 for both alloys are shown in Figs. l-3. A description of these results follows, organized according to the heat treatment to which the specimens were subjected. As a prelude to the more detailed descriptions that follow, we can state with some generality that the yield strength of the CP alloy was by and large more sensitive to the annealing conditions than that of the HP alloy (Fig. l), that the reverse was true for et (Fig. 2), and that Pf/t2 was essentially insensitive to the annealing treatments (Fig. 3). 3.1. The AR condition The yield strengths of the AR material are comparable to those obtained in previous studies [ 12,17, 28,291 from standard tensile and compression tests. The yield strengths were 241.2 * 28.6 MPa and 232.3 f 24.7 MPa for the HP and CP materials respectively. Even though the scatter in the measured values of cr of the HP samples was extremely large, the
2
Y Y + + “i;fF; @
Fig. 2. Strains to fracture, various heat treatments.
8 +
g +
!
!
i-
$ .-
4 _
i
E,, of the HP and CP alloys for the
strain to fracture of the two as-received alloys were similar, 2.04 + 0.7% for the CP alloy and 1.89 & 1.29% for the HP alloy. The fracture strains measured from the bend samples were within the range of previously reported results on tensile ductility of single crystal NiAl ( -0.5 to 2% [ 1,121). Pf/t2 was also comparable for both alloys. 3.2. 1300 -I-FC Both alloys subjected to this annealing treatment exhibited a considerable increase in yield strength. In contrast, neither material exhibited as significant a change in sf or P, /t*.
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H.K. DeMarco et al. CP
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Materials Science and Engineering A1921193 (1995) 333-339
?? HP
anneal; i.e. the CP alloy was strengthened considerably, while the HP alloy was not. The ductility of the HP alloy was lower than that resulting from the 1300 + 400 [70 h] + FC anneal, but .sf for the CP alloy was only marginally affected. The ductilities of the CP and HP alloys were slightly lower than, but comparable, to those in the AR condition. The values of Pr/tz were comparable to those for the AR alloys.
4. Discussion
Fig. 3. Thickness compensated fracture loads, P,/t’, of the HP and CP alloys for the various heat treatments.
3.3. 1300+400[2h]+FC,ACor
WQ
Regardless of the cooling rate from 400°C the samples subjected to this treatment had lower yield stresses than those annealed only at 1300 “C, suggesting that the microstructural feature(s) responsible for the prior increase in yield strength were eliminated by the low-temperature anneal. Furthermore, the yield stresses of both alloys tended to decrease with increasing cooling rate from 400 “C (Fig. l), though the observed trend may not be statistically significant. The ductility of the CP alloy was not affected by cooling rate, but there was a systematic increase in the ductility of the HP alloy as the cooling rate from 400 “C increased (Fig. 2). Pf/t2 was relatively insensitive to the cooling rate from 400°C except possibly for the furnace-cooled HP alloy, for which the scatter was quite large. 3.4. 1300 + 400(70 h] + FC The properties of the HP alloy were relatively unaffected by this annealing treatment, compared to those in the AR condition. However, the CP alloy had a yield strength of 430.3 _+87.2 MPa, which on average was the highest value of aY found for any group of samples. For the HP material aY was much lower than that produced by the 1300 + FC treatment (291.2 f 6.5, cf. 410.9 + 57.6 MPa). The ductility and Pf/t2 values for the CP alloy were roughly the same as those produced by the other annealing treatments. 3.5. AR+400[70h]+FC The yield strengths of both alloys were comparable to those resulting from the 1300 + 400 [70 h] +FC
The CP and HP alloys have approximately the same values of oYin the as-received condition. This was also the case for the as-received CP and HP alloys investigated by Weaver et al. [ 171, which are nearly equivalent to ours. Nevertheless, the yield strength of HP NiAl is expected to be lower than that of CP NiAl because of its lower concentration of impurities. The as-received CP alloy was already homogenized whereas the asreceived HP alloy was not, but this was the only difference in post-growth processing. The most likely explanation for the equivalent values of oY is that the dislocation density and/or thermal vacancy concentration in the HP alloy is higher than that in the CP alloy. Several factors could account for the substantial increase in uY experienced by NiAl of both purities after the 1300 + FC anneal: ( 1) excess thermal vacancies; (2) strain aging; (3) the formation of ultra-fine precipitates or solute-atom clusters. The concentration of thermal vacancies introduced is the only factor common to both materials. It might be expected that the vacancy concentrations in both alloys would have enough time to equilibrate during the slow initial furnace-cooling rate of 0.12 “C s- ‘. A review [l] of various measurements of the formation and migration energies of vacancies in NiAl indicates that the former is much lower than the latter, which is the reason why it is possible to quench a high concentration of thermal vacancies into NiAl in the first place. Since the dislocation density in annealed and slowly cooled NiAl is small [30], the contribution of dislocation sinks in reducing the vacancy concentration is expected to be minimal. Therefore, a relatively high concentration of thermal vacancies probably remains in NiAl when cooling from elevated temperatures, even at moderate to slow cooling rates. The pinning of dislocations by excess vacancies can produce signifcant increases in oY
[311. Strain aging could also account for the dramatic increase in yield strength because the slow cooling rate from 13OO”C, and, more critically, through 4OO”C, should allow enough time for interstitials to pin dislocations, thereby increasing the yield strength. The principal difficulty with this explanation is that o, of the
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specimens subjected to the 1300 + FC and 1300 + 400 [2 h] + FC heat treatments should be comparable if the cooling rate from 400 “C is the decisive heat-treatment variable. Reference to Fig. 1 shows that 4 is reduced after this heat treatment. The reduction is especially significant for the HP alloy, for which the effect of strain aging is expected to be negligible. The nucleation and growth of ultrafine precipitates or the formation of small clusters of solute atoms could also partially account for the observed increases in strength, but if this were the case we would not expect much of an effect in the HP alloy, and we would not expect such a significant reduction in oY on annealing the alloys for an additional 2 h at 400 “C. Consideration of all these factors leads us to conclude that excess vacancies are responsible for the large values of ur produced by the 1300 + FC anneal, although possible contributions from the other factors discussed cannot be discounted. The low-temperature anneals at 400 “C for 2 h reduced the yield stress and evidently eliminated the effect of the high-temperature homogenizing anneal. This is most probably due to the partial annealing out of the retained excess vacancies. With an activation energy for vacancy motion in NiAl of - 160 kJ mol- ’ [l] and D,, of -4 x lo-10 cm* ss’ [32], we estimate that it would take approximately 9 days for the concentration of vacancies to reach equilibrium at the center of a slab of material 700 km thick. Alternatively, vacancies will be depleted to a depth of -0.34 pm from the surface during a 2 h anneal at 400 “C. These vacancydepleted regions are polished away during specimen preparation; hence the excess vacancies in the interior of the samples are most probably eliminated by coalescence into voids [ 17,33-351. This implies that vacancies in the form of individual point defects are more effective obstacles to the motion of dislocations than small voids, which is possible if vacancies effectively pin mobile dislocations. We note here that Weaver et al. [ 171 found that heat-treating monocrystalline NiAl of both purities at -727 “C for 2 h also decreased aY, but that the effect was more pronounced in the HP alloy. They attributed the reduction in uY to the elimination of the excess thermal vacancies in the lattice by agglomeration into small voids. The strain aging phenomenon proposed by Brzeski et ul. [ 131 and Hack et al. [14] can account for the increase in oY and the decrease in .sr with decreasing cooling rate, because interstitials that are randomly distributed at -400 “C would form Cottrell atmospheres more effectively as the cooling rate is reduced (a faster cooling rate would freeze in randomly distributed interstitials, thereby reducing uY). However, the increase in aY with decreasing cooling rate can also be explained by precipitation on cooling from 400 “C,
337
especially in the CP alloy. Ultrafine precipitates (or small clusters of solute atoms) have been observed in NiAl of both purities, with the CP alloy having a much higher concentration [ 191; this is consistent with the higher strength of the CP alloy. Although the kinetics of precipitate growth are unknown, slower cooling rates will allow nucleation, growth and coarsening to occur to a greater extent than would be possible in rapidly cooled specimens. If precipitation strengthening occurs while the alloy is still underaged, aY will increase with slower cooling rates from the annealing temperature, or longer aging times at a fixed temperature. It is clear that the significant increase in aY of the CP alloy produced by the 1300 + 400 [70 h] + FC heat treatment, compared to the 1300 + 400 [ 2 h] + FC heat treatment (Fig. l), can also be accounted for by precipitation strengthening. We would not expect aYto change significantly with annealing time if strain aging were the strengthening mechanism, because Cottrell atmospheres are established relatively quickly, after which their effectiveness in immobilizing dislocations is independent of annealing time. Longer annealing times will produce higher yield strengths because of precipitation or cluster formation, so long as the alloy is underaged and the peak strength of the alloy has not yet been attained. This explanation also accounts for the large increase in yield strength of the CP alloy subjected to the AR + 400 [70 h] + FC heat treatment, compared to that of the as-received material (Fig. l), and the fact that a,, of the HP alloy subjected to this annealing treatment was unaffected. The various annealing treatments did not produce significant changes in the values of Pr/t2; hence some insight into the effect of annealing on fracture toughness can possibly be obtained from the variations of cl (in this regard we note the correlation between tensile fracture strain and toughness observed by Brzeski et al. [ 131). The systematic increase in ductility of the HP alloy to a maximum value of cf = 3% with increasing cooling rate from 400 “C (notwithstanding the large uncertainty in cf for the furnace-cooled specimens) is consistent with either a strain aging effect or a precipitation-strengthening effect. However, cf was also quite large for the specimen subjected to the 1300+ 400 [70 h] + FC anneal, so that there is no unique correlation between ductility and cooling rate. Additionally, the ductility of the CP alloy was essentially unaffected by the heat treatments. Since strain aging is expected to be a more important process in the CP alloy, with its larger interstitial content, it is difficult to rationalize the observed variations in ductility of the HP alloy in the context of strain aging. We note that even the largest ductilities of the HP alloy were small compared to the maximum tensile ductilities reported by Brzeski et al.
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[ 131 on their CP alloy. Hence the bend test itself may play some role in limiting the ultimate ductility of the samples. We believe that excess vacancies, very small precipitates and/or clusters of atoms play the predominant role in affecting the yield strengths of the CP alloy, while vacancies and perhaps strain aging influence the mechanical behavior of the HP material. We cannot otherwise explain the high strength of specimens of both alloys annealed at 1300 “C and furnace-cooled, the subsequent reduction in oY of both alloys on annealing for an additional 2 h at 400 “C, and the large increase in uY (of the CP alloy uniquely) on annealing for 70 h at 400 “C. The precise details of the relative roles of excess vacancies, strain aging and ultrafine precipitates are not yet clear, although it should be kept in mind that excess vacancies will enable precipitates to grow at low temperatures faster than they would under conditions where vacancies are in thermal equilibrium. 5. Summary Annealing both alloys at 1300 “C, followed by furnace cooling, produces a large increase of the yield strength over that of the as-received alloys. The microstructural factor responsible for this effect has not been positively identified, but it is most probable that excess vacancies participate in an important way. After having been annealed at 1300 “C and furnacecooled, a,, of both alloys decreases dramatically on annealing ‘at 400 “C for 2 h. We attribute this observation to the coalescence of excess vacancies. The yield strengths of both alloys increase slightly with increasing cooling rate from 400 “C. This could be due to strain aging, as suggested by Brzeski et al. [ 131, or to the precipitation of ultrafine particles, especially in the CP alloy. Prolonged aging at 400 “C after a homogenizing anneal of 1300 “C produces a large increase in oYof the CP alloy, while that of the HP alloy is relatively unaffected. This observation is consistent with strengthening by the precipitation of ultrafine particles in the CP alloy. The fracture stresses of both alloys are relatively insensitive to the annealing treatments, within the limits of experimental error, as are the ductilities of the CP alloy. However, the ductility of the HP alloy increases with some of the annealing treatments, but these increases cannot be attributed to variations in the cooling rate from the annealing temperature. Acknowledgments
H.K. DeMarco and A.J. Ardell are grateful to the NASA Lewis Research Center for financial support of
this research under grant No. NAG 3-1325. We also thank Dr. R. Darolia of GE Aircraft Engines for the CP alloy and Professor B. Oliver of the University of Tennessee for the HP material. References and M.V. Nathal, ht. Mater. Rev., 38 (1993) 193. D.B. Miracle, Acta Met. Mater., 41 (1993) 649. t:i R.D. Field, D.F. Lahrman and R. Darolia, Acta Met. Mater., 29(1991)2961. [41 R. Darolia, D.F. Lahrman, R.D. Field and A.J. Freeman, in C.T. Liu et al. (eds.), in High Temperature Ordered Intermetallic Alloys III, Mater. Res. Sot. Symp. Proc., Vol. 133, 1989,~. 113. 151 R. Darolia, D.F. Lahrman and R.D. Field, Ser. Met. Mater., 26(1992) 1007. S.J. Jeon and H.-C. Lee, Mater. Sci. Eng., A153 (1992) 392. jsi R.D. Noebe and M. Behbehani, Ser. Met. Mater., 27 (1992) 1795. E.P. George and CT. Liu, J. Mater. Res., 5 (1990) 754. t”,; R.D. Noebe and A. Garg, Ser. Met. Mater., 30 (1994) 8 15. [lOI R.D. Reviere, B.F. Oliver and D.D. Bruns, Mater. Manuf Proc., 4 (1989) 103. [Ill D.R. Johnson, S.M. Joslin, B.F. Oliver, R.D. Noebe and J.D. Whittenberger, in H. Henein and T. Oki (eds.), Processing Materials for Properties, Proc. 1st Int. Co@, TMS, Warrendale, PA, 1993, p. 865. [I21 D.F. Lahrman, R.D. Field and R. Darolia, in L. Johnson et al. (eds.), High Temperature Ordered Intermetallic Alloys ZV, Mater. Res. Sot. Symp. Proc., Vol. 213,1991, p.603. J.M. Brzeski, J.E. Hack, R. Darolia and R.D. Field, Mater. Sci. Eng., A170 (1993) 11. J.E. Hack, J.M. Brzeski and R. Darolia, Ser. Met. Mater., 27 (1992) 1259. R.D. Field, D.F. Lahrman and R. Darolia, in I. Baker et al. (eds.), High Temperature Ordered Intermetallic Alloys V, Mater. Res. Sot. Symp. Proc., Vol. 288, 1993, p. 423. [lb1 J.P. Hirth and J. Lothe, Theory of Dislocations, McGrawHill, New York, 1968. [I71 M.L. Weaver, M.J. Kaufman and R.D. Noebe, Ser. Met. Mater., 29(1993) 1113. R. Jayaram and M.K. Miller, Appl. Surf Sci., 67 (1993) 3 11. t:;; A.J. Duncan, M.J. Kaufman and M.K. Miller, Appl. Surf Sci., 76/77( 1993) 160. [=I H.L. Fraser, R.E. Smallman and M.H. Loretto, Phil. Mag., 28(1973)651. [211 H.L. Fraser, R.E. Smallman and M.H. Loretto, Phil. Msg., 28 (1973) 667. [221 H. Li, EC. Chen and A.J. Ardell, Met. Trans., 22A (1991) 206. Mater. Res., 6(1991) 1950. [231 J.ZhangandA.J.Ardell,J. 1241 D.L. Meyers, EC. Chen, J. Zhang and A.J. Ardell, J. Test Eval., 21 (1993) 263. [251 W.C. Young, Roark’s Formulas for Stress and Strain, 6th edn., McGraw-Hill, New York, p. 391. I261 H.K. DeMarco and A.J. Ardell, in I. Baker et al. (eds.), High Temperature Ordered Intermetallic Alloys V, Mater. Res. Sot. Symp. Proc., Vol. 288,1993, p. 641. P71 F.H. Huang, M.L. Hamilton and G.L. Wire, Nucl. Technol., 57( 1982) 234. [281 R.J. Wasilewski, S.R. Butler and J.E. Hanlon, Trans. AlME, 239(1967) 1357.
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[29] R.D. Field, D.F. Lahrman and R. Darolia, in L. Johnson et al. (eds.), High Temperature Ordered Intermetallic Alloys IV, Mater. Res. Sot. Symp. Proc., Vol. 213, 1991, p. 255. [30] K. Forbes, Ph.D. Thesis, Stanford (1994). [31] R.R. Bowman, R.D. Noebe, S.V. Raj and I.E. Locci, Mef. Trans., 23A (1992) 1493. 1321 J.H. Westbrook, Intermetallic Compounds, John Wiley &
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Sons, New York, 1967. [33] W.J. Yang and R.A. Dodd, Ser. Met., 8 (1974) 237. [34] J.E. Epperson, K.W. Gerstenberg, D. Berner, G. Kostorz and C. Ortiz, Phi/. Msg. A, 38 (1978) 529. [351 A. Parthasarathi and H.L. Fraser, Phil. Msg. A, 50 ( 1Y84) 89.
Materials Science and Engineering
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(1995) 333-339
The effects of heat treatment and purity on the mechanical properties of monocrystalline NiAl Ha K. DeMarcoa, Alan J. Ardell”, Ronald D. Noebeb aDepartment of Materials Science and Engineering, University of California, Los Angeles, CA 90024-1.595, USA ‘NASA Lewis Research Center, MS. 49-3, Cleveland, OH 441353191, USA
Abstract The yield stress, 4, strain to fracture, q, and thickness-compensated fracture load of monocrystalline NiAl were investigated as functions of purity, annealing treatment and cooling rates. A miniaturized disk bend test was employed to test disks 3 mm in diameter and approximately 250 pm thick. Specimens in the soft (110) orientation from commercial purity (CP) and high-purity (HP) alloys were tested. The annealing treatments affected uYmore strongly than the other mechanical properties. The yield stresses of both of the as-received alloys increased significantly after annealing at 1300 “C followed by furnace cooling. Subsequent annealing at 400 “C for 2 h resulted in a reduction of uY.This behavior is attributed to the role of excess vacancies retained during cooling, which annealed out at the lower temperature. The mean values of the yield strength also tended to decrease with increasing cooling rate from 400 “C, but the effect was small. A substantial increase in uY of the CP alloy was found on prolonged aging at 400 “C, whereas aY of the HP alloy was unaffected. We attribute this behavior to the precipitation of very small particles or solute atom clusters in the CP alloy. The ductility of the CP alloy, and the thickness-compensated fracture loads of both alloys, were relatively insensitive to the heat treatments. Keywords: Heat treatment;
Mechanical
properties;
Nickel; Aluminium;
1. Introduction There is considerable current interest in ordered intermetallic compounds as candidate materials for high-temperature structural applications. NiAl, with its
high melting temperature and low density, has received special attention, as evidenced by recent comprehensive reviews of the literature on its physical and mechanical properties [1,2]. However, the low roomtemperature ductility and fracture toughness of monolithic NiAl continue to impede its utilization as a structural material. Consequently, a substantial amount of research has focused on improving the ductility of both polycrystalline and monocrystalline NiAl by microalloying additions [3-91. This effort has been aided by the development of techniques [ 10,l l] for growing high-purity single crystals, which enables investigators to determine the mechanical properties of NiAl containing very low impurity levels. The widely accepted values for the tensile strains to fracture for monocrystalline NiAl in the “soft” (110) orientation range from 1 to 3% [ 1,121, but it has been 0921-5093/95/$9.50 0 1995 - Elsevier Science S.A. All rights reserved SSDZ 0921-5093(94)03215-7
Crystals
demonstrated recently that the inherent tensile ductility of NiAl may be as high as 7-8% strain to failure [ 131. Brzeski et al. [ 131 and Hack et al. [ 141 have presented experimental evidence that the cooling rate from 400 “C strongly affects the fracture toughness. They concluded that a strain-aging effect at low temperatures ( -200 “C), involving either interstitial impurities or constitutional vacancies resulting from slight deviations from stoichiometry, hinders the room-temperature mobility of dislocations. Field et al. [15] have reported the occurrence of serrated yielding in monocrystals tested in compression at several temperatures. Their observations are consistent with dynamic strainaging effects involving interactions between point defects and dislocations [ 161, and support the idea that strain aging influences the mechanical behavior of NiAl. Vacancies and alloy purity also affect the mechanical properties of NiAl. Weaver et al. [ 171 found that annealing commercial-purity (CP) and high-purity (HP) NiAl at 727 “C, followed by a furnace cool of -0.083 K s-l, decreased the compressive yield
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that plastic deformation of discs in the (100) orientation does not occur by bending, but by a kinking process that probably involves slip on { lOO} [20,21]. This is because ( 100) is a hard orientation for the disks, and normal plastic bending would require (111) slip. Results are therefore presented only for the specimens deformed in the ( 110) orientation. Each crystal of NiAl was oriented using the Laue back-reflection method, and slices averaging 0.7 mm in thickness were cut parallel to ( 110) using an electricspark discharge machine. The slices were electropolished after cutting to remove surface damage and contaminants, and subjected to several different heattreatment conditions: (1) as-received material (this is designated as the AR condition); (2) annealed at 400 “C for 70 h, then furnace-cooled at an average rate of 0.023 “C SK’ (AR+ 400 [70 h] + FC); (3) annealed at 1300 “C for 36 h in flowing ultra-highpurity (99.995%) argon and then furnace-cooled at an initial rate of 0.12 “C s- ’( 1300 + FC); (4) as in (3), but then annealed at 400 “C for 2 h and either furnace-cooled at a rate of 0.023 “C s- ‘, aircooled or water-quenched (these are designated as 1300 + 400 [2 h] + FC, AC or WQ respectively); (5) as in 3, but annealed at 400 “C for 70 h and furnace-cooled ( 1300 + 400 [70 h] + FC). Disks 3 mm in diameter were cut from the heattreated slices using an abrasive slurry cutter. Each disk was then ground on each side to a thickness of -250 pm using silicon carbide abrasive papers, then metallographically polished to a final finish using 0.05 pm alumina slurry. The orientation of each individual disk was not verified, but it is assumed that the planes of the disks were within 2” of ( 110) according to the orientations of several disks that were checked using electron channeling patterns or Laue back-reflection X-ray patterns. The MDBT apparatus is described in detail elsewhere [22-241. The tests were conducted in the ringon-ring mode using a table-model Instron testing machine at a cross-head displacement rate of 8.3 pm s- I. Despite the elastic anisotropy exhibited by
strength of both alloys. The resulting yield strength of the HP alloy was lower than that of the CP alloy. Jayaram and Miller [ 181 observed fine carbide precipitates in C-doped NiAl, and Duncan et al. [ 191 observed inhomogeneously distributed ultrafine precipitates in the CP alloy, as well as an occasional precipitate in HP crystals. The existence of these ultrafine preciptiates may affect the yield strength and ductility of NiAl. It seems safe to conclude that the combined effects of alloy purity and thermal maniulation on the mechanical properties of monocrystalline NiAl are complex and incompletely understood. In this paper we report the results of an investigation of the effects of heat treatment, cooling rates, and alloy purity on yield strength, bending strain to fracture and fracture load of NiAl using a miniaturized disk-bend test (MDBT). The advantage of the small specimen size utilized by the MDBT, besides the obvious ability to conduct many tests on a small quantity of material, is that small specimens are rapidly equilibrated thermally, and can therefore be cooled relatively rapidly without acquiring large residual stresses.
2. Experimental Monocrystalline NiAl of two different purities was investigated: a CP crystal grown by GE Aircraft Engines using the Bridgman method and a HP crystal grown at the University of Tennessee using a containerless processing method [lO,ll]. The as-received CP alloy had already been subjected to a homogenizing anneal at 13 16 “C for 48 h and cooled at a rate of -0.213 “C s-‘. The HP alloy was provided in the asgrown condition, with no post-growth processing treatment. Chemical analysis of the CP crystal used in this study was not performed, but the typical impurity levels in identically grown and processed GE NiAl monocrystals are listed in Table 1. Chemical analysis of the HP crystal was performed at NASA Lewis Research Center; details of the analytical techniques used are described by Weaver et al. [ 171. Experiments were performed on disks in both (100) and ( 110) orientations. However, we have discovered
Table 1 Interstitial
and Si content (at.%) of typical NiAl alloys
Alloy
Si
C
0
N
S
Comments
HP CP
0.06 0.15
< 0.0036 0.0306
< 0.0027 0.0155
< 0.003 1 0.0036
< 0.0013 < 0.001
Same ingot as Weaver et al. [ 171 Average C, 0, N values from Hack et al. [ 141; Si and S from[17]
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Materials Science and Engineering A1921193 (1995) 333-339
NiAl, the yield stresses, oY, were calculated using the formula [25], valid for elastically isotropic materials,
oy=g+
(I+.),,;+~$ 1
[
?? HP
CP
1-c II (1)
335
a*
where Py is the applied load at yielding, t is the thickness of the specimen, Y is Poisson’s ratio, a and b are the radii of the support die and loading ring respectively, and R is the radius of the specimen. The values of the parameters used in the calculations were a=0.955 mm, b=0.36 mm, R= 1.5 mm and ~=0.31 [1,2]; t varied from one specimen to another, but was typically 250 pm. In addition to aY, we report the thickness-compensated load at fracture, Pr/t2, where P, is the load at fracture. This quantity provides an estimate of the relative fracture stresses of the various specimens, Representative curves of P/t* vs. displacement of the center of the specimen, W, are shown in Fig. 1 of the paper by DeMarco and Ardell[26]. The strain to failure, sr, was also measured in this investigation, and was estimated using the formula [27]
(2) Since az % t*, a useful measure of the strain is simply yt/a*. In all the measurements, wr was taken as the difference between the displacement at fracture and that at yielding.
Fig. 1. Yield stresses, heat treatments.
u,,, of the HP and CP alloys for the various
CP
H
HP
2.5
0.5 0
3. Results The data on aY, sf and Pr/t2 for both alloys are shown in Figs. l-3. A description of these results follows, organized according to the heat treatment to which the specimens were subjected. As a prelude to the more detailed descriptions that follow, we can state with some generality that the yield strength of the CP alloy was by and large more sensitive to the annealing conditions than that of the HP alloy (Fig. l), that the reverse was true for et (Fig. 2), and that Pf/t2 was essentially insensitive to the annealing treatments (Fig. 3). 3.1. The AR condition The yield strengths of the AR material are comparable to those obtained in previous studies [ 12,17, 28,291 from standard tensile and compression tests. The yield strengths were 241.2 * 28.6 MPa and 232.3 f 24.7 MPa for the HP and CP materials respectively. Even though the scatter in the measured values of cr of the HP samples was extremely large, the
2
Y Y + + “i;fF; @
Fig. 2. Strains to fracture, various heat treatments.
8 +
g +
!
!
i-
$ .-
4 _
i
E,, of the HP and CP alloys for the
strain to fracture of the two as-received alloys were similar, 2.04 + 0.7% for the CP alloy and 1.89 & 1.29% for the HP alloy. The fracture strains measured from the bend samples were within the range of previously reported results on tensile ductility of single crystal NiAl ( -0.5 to 2% [ 1,121). Pf/t2 was also comparable for both alloys. 3.2. 1300 -I-FC Both alloys subjected to this annealing treatment exhibited a considerable increase in yield strength. In contrast, neither material exhibited as significant a change in sf or P, /t*.
336
H.K. DeMarco et al. CP
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Materials Science and Engineering A1921193 (1995) 333-339
?? HP
anneal; i.e. the CP alloy was strengthened considerably, while the HP alloy was not. The ductility of the HP alloy was lower than that resulting from the 1300 + 400 [70 h] + FC anneal, but .sf for the CP alloy was only marginally affected. The ductilities of the CP and HP alloys were slightly lower than, but comparable, to those in the AR condition. The values of Pr/tz were comparable to those for the AR alloys.
4. Discussion
Fig. 3. Thickness compensated fracture loads, P,/t’, of the HP and CP alloys for the various heat treatments.
3.3. 1300+400[2h]+FC,ACor
WQ
Regardless of the cooling rate from 400°C the samples subjected to this treatment had lower yield stresses than those annealed only at 1300 “C, suggesting that the microstructural feature(s) responsible for the prior increase in yield strength were eliminated by the low-temperature anneal. Furthermore, the yield stresses of both alloys tended to decrease with increasing cooling rate from 400 “C (Fig. l), though the observed trend may not be statistically significant. The ductility of the CP alloy was not affected by cooling rate, but there was a systematic increase in the ductility of the HP alloy as the cooling rate from 400 “C increased (Fig. 2). Pf/t2 was relatively insensitive to the cooling rate from 400°C except possibly for the furnace-cooled HP alloy, for which the scatter was quite large. 3.4. 1300 + 400(70 h] + FC The properties of the HP alloy were relatively unaffected by this annealing treatment, compared to those in the AR condition. However, the CP alloy had a yield strength of 430.3 _+87.2 MPa, which on average was the highest value of aY found for any group of samples. For the HP material aY was much lower than that produced by the 1300 + FC treatment (291.2 f 6.5, cf. 410.9 + 57.6 MPa). The ductility and Pf/t2 values for the CP alloy were roughly the same as those produced by the other annealing treatments. 3.5. AR+400[70h]+FC The yield strengths of both alloys were comparable to those resulting from the 1300 + 400 [70 h] +FC
The CP and HP alloys have approximately the same values of oYin the as-received condition. This was also the case for the as-received CP and HP alloys investigated by Weaver et al. [ 171, which are nearly equivalent to ours. Nevertheless, the yield strength of HP NiAl is expected to be lower than that of CP NiAl because of its lower concentration of impurities. The as-received CP alloy was already homogenized whereas the asreceived HP alloy was not, but this was the only difference in post-growth processing. The most likely explanation for the equivalent values of oY is that the dislocation density and/or thermal vacancy concentration in the HP alloy is higher than that in the CP alloy. Several factors could account for the substantial increase in uY experienced by NiAl of both purities after the 1300 + FC anneal: ( 1) excess thermal vacancies; (2) strain aging; (3) the formation of ultra-fine precipitates or solute-atom clusters. The concentration of thermal vacancies introduced is the only factor common to both materials. It might be expected that the vacancy concentrations in both alloys would have enough time to equilibrate during the slow initial furnace-cooling rate of 0.12 “C s- ‘. A review [l] of various measurements of the formation and migration energies of vacancies in NiAl indicates that the former is much lower than the latter, which is the reason why it is possible to quench a high concentration of thermal vacancies into NiAl in the first place. Since the dislocation density in annealed and slowly cooled NiAl is small [30], the contribution of dislocation sinks in reducing the vacancy concentration is expected to be minimal. Therefore, a relatively high concentration of thermal vacancies probably remains in NiAl when cooling from elevated temperatures, even at moderate to slow cooling rates. The pinning of dislocations by excess vacancies can produce signifcant increases in oY
[311. Strain aging could also account for the dramatic increase in yield strength because the slow cooling rate from 13OO”C, and, more critically, through 4OO”C, should allow enough time for interstitials to pin dislocations, thereby increasing the yield strength. The principal difficulty with this explanation is that o, of the
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specimens subjected to the 1300 + FC and 1300 + 400 [2 h] + FC heat treatments should be comparable if the cooling rate from 400 “C is the decisive heat-treatment variable. Reference to Fig. 1 shows that 4 is reduced after this heat treatment. The reduction is especially significant for the HP alloy, for which the effect of strain aging is expected to be negligible. The nucleation and growth of ultrafine precipitates or the formation of small clusters of solute atoms could also partially account for the observed increases in strength, but if this were the case we would not expect much of an effect in the HP alloy, and we would not expect such a significant reduction in oY on annealing the alloys for an additional 2 h at 400 “C. Consideration of all these factors leads us to conclude that excess vacancies are responsible for the large values of ur produced by the 1300 + FC anneal, although possible contributions from the other factors discussed cannot be discounted. The low-temperature anneals at 400 “C for 2 h reduced the yield stress and evidently eliminated the effect of the high-temperature homogenizing anneal. This is most probably due to the partial annealing out of the retained excess vacancies. With an activation energy for vacancy motion in NiAl of - 160 kJ mol- ’ [l] and D,, of -4 x lo-10 cm* ss’ [32], we estimate that it would take approximately 9 days for the concentration of vacancies to reach equilibrium at the center of a slab of material 700 km thick. Alternatively, vacancies will be depleted to a depth of -0.34 pm from the surface during a 2 h anneal at 400 “C. These vacancydepleted regions are polished away during specimen preparation; hence the excess vacancies in the interior of the samples are most probably eliminated by coalescence into voids [ 17,33-351. This implies that vacancies in the form of individual point defects are more effective obstacles to the motion of dislocations than small voids, which is possible if vacancies effectively pin mobile dislocations. We note here that Weaver et al. [ 171 found that heat-treating monocrystalline NiAl of both purities at -727 “C for 2 h also decreased aY, but that the effect was more pronounced in the HP alloy. They attributed the reduction in uY to the elimination of the excess thermal vacancies in the lattice by agglomeration into small voids. The strain aging phenomenon proposed by Brzeski et ul. [ 131 and Hack et al. [14] can account for the increase in oY and the decrease in .sr with decreasing cooling rate, because interstitials that are randomly distributed at -400 “C would form Cottrell atmospheres more effectively as the cooling rate is reduced (a faster cooling rate would freeze in randomly distributed interstitials, thereby reducing uY). However, the increase in aY with decreasing cooling rate can also be explained by precipitation on cooling from 400 “C,
337
especially in the CP alloy. Ultrafine precipitates (or small clusters of solute atoms) have been observed in NiAl of both purities, with the CP alloy having a much higher concentration [ 191; this is consistent with the higher strength of the CP alloy. Although the kinetics of precipitate growth are unknown, slower cooling rates will allow nucleation, growth and coarsening to occur to a greater extent than would be possible in rapidly cooled specimens. If precipitation strengthening occurs while the alloy is still underaged, aY will increase with slower cooling rates from the annealing temperature, or longer aging times at a fixed temperature. It is clear that the significant increase in aY of the CP alloy produced by the 1300 + 400 [70 h] + FC heat treatment, compared to the 1300 + 400 [ 2 h] + FC heat treatment (Fig. l), can also be accounted for by precipitation strengthening. We would not expect aYto change significantly with annealing time if strain aging were the strengthening mechanism, because Cottrell atmospheres are established relatively quickly, after which their effectiveness in immobilizing dislocations is independent of annealing time. Longer annealing times will produce higher yield strengths because of precipitation or cluster formation, so long as the alloy is underaged and the peak strength of the alloy has not yet been attained. This explanation also accounts for the large increase in yield strength of the CP alloy subjected to the AR + 400 [70 h] + FC heat treatment, compared to that of the as-received material (Fig. l), and the fact that a,, of the HP alloy subjected to this annealing treatment was unaffected. The various annealing treatments did not produce significant changes in the values of Pr/t2; hence some insight into the effect of annealing on fracture toughness can possibly be obtained from the variations of cl (in this regard we note the correlation between tensile fracture strain and toughness observed by Brzeski et al. [ 131). The systematic increase in ductility of the HP alloy to a maximum value of cf = 3% with increasing cooling rate from 400 “C (notwithstanding the large uncertainty in cf for the furnace-cooled specimens) is consistent with either a strain aging effect or a precipitation-strengthening effect. However, cf was also quite large for the specimen subjected to the 1300+ 400 [70 h] + FC anneal, so that there is no unique correlation between ductility and cooling rate. Additionally, the ductility of the CP alloy was essentially unaffected by the heat treatments. Since strain aging is expected to be a more important process in the CP alloy, with its larger interstitial content, it is difficult to rationalize the observed variations in ductility of the HP alloy in the context of strain aging. We note that even the largest ductilities of the HP alloy were small compared to the maximum tensile ductilities reported by Brzeski et al.
338
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Materials S crence and Engineering Al 92/l 93 (1995) 333-339
[ 131 on their CP alloy. Hence the bend test itself may play some role in limiting the ultimate ductility of the samples. We believe that excess vacancies, very small precipitates and/or clusters of atoms play the predominant role in affecting the yield strengths of the CP alloy, while vacancies and perhaps strain aging influence the mechanical behavior of the HP material. We cannot otherwise explain the high strength of specimens of both alloys annealed at 1300 “C and furnace-cooled, the subsequent reduction in oY of both alloys on annealing for an additional 2 h at 400 “C, and the large increase in uY (of the CP alloy uniquely) on annealing for 70 h at 400 “C. The precise details of the relative roles of excess vacancies, strain aging and ultrafine precipitates are not yet clear, although it should be kept in mind that excess vacancies will enable precipitates to grow at low temperatures faster than they would under conditions where vacancies are in thermal equilibrium. 5. Summary Annealing both alloys at 1300 “C, followed by furnace cooling, produces a large increase of the yield strength over that of the as-received alloys. The microstructural factor responsible for this effect has not been positively identified, but it is most probable that excess vacancies participate in an important way. After having been annealed at 1300 “C and furnacecooled, a,, of both alloys decreases dramatically on annealing ‘at 400 “C for 2 h. We attribute this observation to the coalescence of excess vacancies. The yield strengths of both alloys increase slightly with increasing cooling rate from 400 “C. This could be due to strain aging, as suggested by Brzeski et al. [ 131, or to the precipitation of ultrafine particles, especially in the CP alloy. Prolonged aging at 400 “C after a homogenizing anneal of 1300 “C produces a large increase in oYof the CP alloy, while that of the HP alloy is relatively unaffected. This observation is consistent with strengthening by the precipitation of ultrafine particles in the CP alloy. The fracture stresses of both alloys are relatively insensitive to the annealing treatments, within the limits of experimental error, as are the ductilities of the CP alloy. However, the ductility of the HP alloy increases with some of the annealing treatments, but these increases cannot be attributed to variations in the cooling rate from the annealing temperature. Acknowledgments
H.K. DeMarco and A.J. Ardell are grateful to the NASA Lewis Research Center for financial support of
this research under grant No. NAG 3-1325. We also thank Dr. R. Darolia of GE Aircraft Engines for the CP alloy and Professor B. Oliver of the University of Tennessee for the HP material. References and M.V. Nathal, ht. Mater. Rev., 38 (1993) 193. D.B. Miracle, Acta Met. Mater., 41 (1993) 649. t:i R.D. Field, D.F. Lahrman and R. Darolia, Acta Met. Mater., 29(1991)2961. [41 R. Darolia, D.F. Lahrman, R.D. Field and A.J. Freeman, in C.T. Liu et al. (eds.), in High Temperature Ordered Intermetallic Alloys III, Mater. Res. Sot. Symp. Proc., Vol. 133, 1989,~. 113. 151 R. Darolia, D.F. Lahrman and R.D. Field, Ser. Met. Mater., 26(1992) 1007. S.J. Jeon and H.-C. Lee, Mater. Sci. Eng., A153 (1992) 392. jsi R.D. Noebe and M. Behbehani, Ser. Met. Mater., 27 (1992) 1795. E.P. George and CT. Liu, J. Mater. Res., 5 (1990) 754. t”,; R.D. Noebe and A. Garg, Ser. Met. Mater., 30 (1994) 8 15. [lOI R.D. Reviere, B.F. Oliver and D.D. Bruns, Mater. Manuf Proc., 4 (1989) 103. [Ill D.R. Johnson, S.M. Joslin, B.F. Oliver, R.D. Noebe and J.D. Whittenberger, in H. Henein and T. Oki (eds.), Processing Materials for Properties, Proc. 1st Int. Co@, TMS, Warrendale, PA, 1993, p. 865. [I21 D.F. Lahrman, R.D. Field and R. Darolia, in L. Johnson et al. (eds.), High Temperature Ordered Intermetallic Alloys ZV, Mater. Res. Sot. Symp. Proc., Vol. 213,1991, p.603. J.M. Brzeski, J.E. Hack, R. Darolia and R.D. Field, Mater. Sci. Eng., A170 (1993) 11. J.E. Hack, J.M. Brzeski and R. Darolia, Ser. Met. Mater., 27 (1992) 1259. R.D. Field, D.F. Lahrman and R. Darolia, in I. Baker et al. (eds.), High Temperature Ordered Intermetallic Alloys V, Mater. Res. Sot. Symp. Proc., Vol. 288, 1993, p. 423. [lb1 J.P. Hirth and J. Lothe, Theory of Dislocations, McGrawHill, New York, 1968. [I71 M.L. Weaver, M.J. Kaufman and R.D. Noebe, Ser. Met. Mater., 29(1993) 1113. R. Jayaram and M.K. Miller, Appl. Surf Sci., 67 (1993) 3 11. t:;; A.J. Duncan, M.J. Kaufman and M.K. Miller, Appl. Surf Sci., 76/77( 1993) 160. [=I H.L. Fraser, R.E. Smallman and M.H. Loretto, Phil. Mag., 28(1973)651. [211 H.L. Fraser, R.E. Smallman and M.H. Loretto, Phil. Msg., 28 (1973) 667. [221 H. Li, EC. Chen and A.J. Ardell, Met. Trans., 22A (1991) 206. Mater. Res., 6(1991) 1950. [231 J.ZhangandA.J.Ardell,J. 1241 D.L. Meyers, EC. Chen, J. Zhang and A.J. Ardell, J. Test Eval., 21 (1993) 263. [251 W.C. Young, Roark’s Formulas for Stress and Strain, 6th edn., McGraw-Hill, New York, p. 391. I261 H.K. DeMarco and A.J. Ardell, in I. Baker et al. (eds.), High Temperature Ordered Intermetallic Alloys V, Mater. Res. Sot. Symp. Proc., Vol. 288,1993, p. 641. P71 F.H. Huang, M.L. Hamilton and G.L. Wire, Nucl. Technol., 57( 1982) 234. [281 R.J. Wasilewski, S.R. Butler and J.E. Hanlon, Trans. AlME, 239(1967) 1357.
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