Effect of vacuum on room-temperature ductility of Ni3Al

Scripta METALLURGICA et MATERIALIA

Vol. 30, pp. 37-42, 1994 Printed in the U.S.A.

Pergamon Press Ltd.

EFFECT OF VACUUM ON ROOM-TEMPERATURE DUCTILITY OF Ni3A1 E. P. George, C. T. Liu, and D. P. Pope t

Metals and Ceramics Division, Oak Ridge National Laboratory, Oak Ridge, TN 37831-6093. tDepartment of Materials Science and Engineering, University of Pennsylvania, Philadelphia, PA 19104-6272. (Received September 3, 1993) (Revised September 20, 1993) Introduction There is increasing evidence [1-3] that grain boundaries in Ni3A1 are not intrinsicallybrittle as once thought. Rather, the low ductility commonly encountered when polycrystalline Ni3AI is tested in air is due to extrinsic factors such as H20-induced environmental embrittlement, and microcracks introduced during processing. Thus, in a recent paper [1], we showed that the elongation to fracture of Ni3AI can be as high as 16% at room temperature, when defect-free specimens (produced by cold working and recrystallizing single crystals) are tested in relatively dry oxygen, but that it can be as low as 3% when the same material is tested in ordinary ambient air. The embrittling agent in air was assumed to be H20---by analogy with earlier work on iron aluminides [4-6]--and the operative mechanism was thought to involve the reduction of water vapor in air by aluminum atoms in the aluminide, resulting in the generation of atomic hydrogen, which then caused embrittlement. In such a mechanism, H20 serves merely as a source of atomic H. If atomic H can be generated by other means, e.g., by the dissociation of molecular H 2 at Ni3AI surfaces, 1-I20 may not be required for embrittlement. Therefore, until experiments are conducted in dry H2, both H20 and H2 should be considered as ~otential culprits. It is worth noting, though, that the partial pressure ofH 2 in atmospheric air is only about 5 x 10- Pa [7], which is considerably less than the vapor pressure (-10 j Pa) of H20 in air at 50% relative humidity [8]. Therefore, it is possible that, while H 2 at relatively high pressures may well embrittle Ni3A1, the low concentration of H2 in ordinary ambient air is insufficient to cause embrittlement. Regardless of whether it is H20, or both 1--I20and H2, that are responsible for the embrittlement of Ni3A1 in air, it is of interest to measure the ductility of carefully prepared (i.e., defect-free) Ni3A1 specimens in vacuum as a function of pressure (our earlier experiments were performed in oxygen). It is of interest also to determine by how much the ductility increases when the test is performed in ultrahigh vacuum (UHV), i.e., when the moisture content is reduced to extremely low levels. Put another way, conventional vacuums may not be good enough to entirely eliminate environmentalembrittlement, and testing in ultrahigh vacuum may be necessary to determine the intrinsic properties of grain boundaries in Ni3A1. In this paper we report the ductility of polycrystalline, B-free Ni3A1 as a function of pressure, over 12 orders of magnitude, from -10Spa (in air) to -10 -8 Pa (in UHV). Experimental Procedure A single crystal of Ni-23.4AI (at.%) was rolled at room temperature and subsequently recrystallized to get defect-free polycrystalline material with a grain size of approximately 13 lam. Details are given in ref. [1]. For the UHV test, miniature tensile specimens having a gauge section of 3.2 x 1.6 x 0.5 mm were first electro-discharge machined from the rolled plate with their tensile axes parallel to the rolling direction, and then electropolished (at room temperature, in a solution of 13 v/o H2SO4 in methyl alcohol, at approximately 7 V) to remove the surface oxide. For the remainder of the tests (at higher pressures), specimen geometry and preparation were as reported in our earlier paper [1], and specimen orientation was again parallel to the rolling direction. Room temperature tensile tests at pressures down to -10 -4 Pa were performed on a screw-driven Instron machine equipped with a turbo-pumped vacuum chamber. These tests were conducted at a constant cross-head speed of 4.2 x 10-2 ram.s-l, which corresponds to an engineering strain rate of 5.3 x 10-3 s-1. For the tensile test at 1.3 x 10-I Pa, the test chamber was first evacuated to a pressure of 7.3 x 10-4 Pa, and then air was leaked into the test chamber through a Varian leak valve until the pressure reached 1.3 x 10-1 Pa. The UHV tensile test was conducted in a fracture chamber attached to an Auger electron microprobe. This chamber is pumped with a combination of liquid-nitrogen-cooled sorption pumps (down to pressures of ~1 Pa) and an ion pump (at lower pressures). To achieve UHV, the specimen and chamber were first baked out at -180°C for 18 h. During the bake.out, the ion pump was kept running and the chamber pressure fell progressively from -10 -4 Pa to -10 "~ Pa. Following cooldown to room temperature, pumping was carded on for an additional 48 h prior to fracture so that

37 0956-716X/94 $6.00 + .00

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pressure could drop further--into the low 10-8 Pa range. Just before fracture, a cold finger in the vicinity of the specimen was cooled with liquid nitrogen to freeze out any residual moisture. The extension rate used for fracture in UHV was 1.1 x 10-2 mm-s-1, which corresponds to an engineering strain rate of 3.3 x 10-3 s-1. Resul~ Table 1 and Figure 1 summarize the results of room-temperature tensile tests. (Also included for comparison are the results of tensile tests in air from ref. [1].) The measured ductilities are 3.1% in air (105 Pa), 7.9% at 1.3 x 10-1 Pa, 12.8% at 3.2 x 10-4 Pa, and 23.4% at 3.6 x 10-8 Pa. Yield strengths were relatively insensitive to test environment. Ultimate tensile strengths, on the other hand, increased with increasing elongation to fracture. As before [1,2], the yield strengths reported in Table 1 were calculated from either the load corresponding to the fiat (discontinuous yielding) region of the load--displacement curves, or the load at 0.2% offset (if there was no well-def'med fiat region). Loads were not measured during the UHV test, but yield strength is not expected to be a function of vacuum level. TABLE 1. Effect of vacuum on the room temperature tensile properties of Ni-23.4A1. Pressure (Pa)

Elongation to Fracture (%)

Yield Strength (MPa)

Ultimate Strength (MPa)

1.0 x 105 t

3.1

308

392

1.3 x 10-1

7.9

323

519

3.2 x 10-4

12.8

301

602

23.4

:

t

3.6 x 10-8 /±Atmosphenc pressure. TNot measured.

25

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20

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~ W

B...,.... 10

"*~'~..OOo°°''==oo.. 5

! 10"81~0"710"610"510"410"310"210"1

100 101 102 103 104 105 1

Pressure (Pa) Fig. 1. Room temperature tensile ductility of Ni-23.4AI as a function of vacuum level.

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Figure 2 shows the fracture surface appearance in air and in the different vacuums. While fracture is predominantly intergranular in all cases, increasing amounts of transgranular fracture as well as an increasingly ductile appearance are evident as the vacuum level improves.

Fig. 2. Room temperature fracture surfaces of Ni-23.4A1 tensile tested in (a) air (760 torr), (b) 1 x 10-3 torr vacuum, (c) 2.4 × 10-6 tort vacuum, and (d) 2.7 x 10-10 torr vacuum. Discussion Our present results demonstrate clearly the strong dependence of ductility on vacuum level:* In ultrahigh vacuum, B-free Ni3AI can have ductility as high as 23%. This is the highest-ever ductility reported for equiaxed, polycrystalline, B-free Ni3A1, and demonstrates that, contrary to conventional wisdom, Ni3A1 is intrinsically quite ductile. With increasing amounts of air in the test environment, ductility drops systematically (Fig. 1), until in ordinary ambient air it drops to about 3%. Moisture in air is mainly responsible for this embrittlement although, as cautioned earlier, H 2 cannot be ruled out as a contributing factor. This must mean that a significant portion of the so-called boron effect is related to suppression of H20-induced environmental embrittlement. And, indeed, although the detailed mechanism by which it improves ductility is unclear at present, it has been shown that boron, * For purposes of this discussion, we adopt the following descriptions: ultrahigh vacuum (-10 -10 torr), high vacuum (~10 -6 ton'), and low vacuum (-10 -3 torr).

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depending on the amounts added, suppresses environmental embrittlement to varying degrees [9,10]. The interesting question is whether boron does anything else. As shown in Fig. 2, despite its high ductility, undoped Ni3A1 fractures predominantly intergranularly----even in UHV. Boron-doped.Ni3A1, on . t h e o ~ hand, fractures predominantly transgranularly----cvenin air [11,12]. By segregaung to the grmn bounaanes tizJ, mererore, voron apparently enhances grain boundary strength [13]. How important this effect is in improving overall ductility remains an open question. It depends to a large extent on how important environmental embrittlement is in decreasing ductility. As we have pointed out [14], the phenomenon of environmental embrittlement has evolved dramatically over the course of the last two years---from being essentially unrecognizedto being a major cause of brittleness in Ni3A1. While the intrinsic ductility of B-free N i 3 ~" (23%) is still considerably..!ess than.that of ~e. best B-doped alloys (50%, ref. [12]), it remains to be seen whemer mture experiments wau rcsutt m even nlgner ductilities. Clearly, if the intrinsic ductility of B-free Ni3A1 (in the absence of mois.ture) a.pp.roaches ~at of B-doped Ni3AI, it would raean that the principal role of boron is to suppresse.nvtronm..c.ntal emtmttlcmcnt. Mean wn.ue, on the basis of results obtained to date, one has to conclude that part ot me auctilizing effect or t)oron is remtc~a to its beneficial effect on grain boundary cohesion. An interesting result of our present work is that the ductility obtained in high vacuum (12.8%, Table 1) is lower than the ductility (15.8%) obtained in our previous study [1] in 02. In that study, the oxygen tests were conducted at 6.7 x 104 Pa after first evacuating the test chamber into the high vacuum range and then backfilling with 02. To rationalize these different ductilities, we estimated the water-vapor contents of the various test environments, as shown in Table 2. Consider first ordinary ambient air: at room temperature (~22°C), the saturation water vapor pressure is 2.7 x 103 Pa [8]; so, assuming 50% relative humidity, which is typical for laboratory air, the partial pressure of moisture in air is 1.3 × 103 Pa. Next consider oxygen: although its reported purity was 99.993%, such analyses usually ignore H20. So we experimentally determined the moisture content of our oxygen with a hygrometer, and obtained a value of 44.3 ppm. This corresponds to a water vapor pressure of 3 Pa when the total oxygen pressure is 6.7 × 104 Pa. Next consider low vacuum which, in our experiments, was obtained by first evacuating the chamber to - 1 0 -4 Pa and then backfilling with air. It is reasonable to assume that the fraction of H20 in this low-vacuum environment is about the same as that~n air (vi.'z., 1.3%.at 50% relative humidity), yielding a partial pressure of 1-12O for the low-vacuum test of 2 x 10- Pa. (Smular esumates arc given in rcf. [15].) Next consider the high-vacuum test in an unbaked system, where 70-90% of the residual gas is expected to be water vapor [15]. Assuming a middle value of 80%, the partial pressur.c. of H.20 in the.high-vacuum test can be calculated to be 3 x 10"4 Pa. Finally, in the ultrahigh vacuum test, most ot me resmu~, gas ~s expectco to be H 2 (after a bake-out), with H20 contents in the range 3-5% [15]. Picking a middle value of 4%, the partial pressure of H20 in our UI-IV environment is estimated to be 1 x 10-y Pa. TABLE 2. Effect of I-I20 partial pressure on the room-temperature tensile ductility of Ni-23.4A1. Test Environment

Estimated Partial Pressure of H20 (Pa)

Elongation to Fracture (%)

Air (105 Pa)

1 x

103

3.1

Oxygen (6.7 x 104 Pa)

3 x 10°

15.8

L o w vacuum (1.3 × 10 -I Pa)

2 x 10-3

7.9

High vacuum (3.2 × 10 -4 Pa)

3 x 10-4

12.8

Ultrahigh vacuum (3.6 x 10 -8 Pa)

1 x 10-9

23.4

On the basis of the above estimates, therefore, it appears that the amount of residual H20 in the oxygen test is considerably greater than that in the high vacuum test. (Indeed, it appears to be greater than that even in the low vacuum test.) Yet ductility was higher in the oxygen test. This must mean that the beneficial effect of oxygen is not related simply to its low moisture content. Because if that were the case, the high- and low-vacuum tests---because of their lower moisture contents---should have resulted in higher ductilities. This point is further elaborated in Fig. 3, where we have plotted the ductility of Ni3AI as a function of the partial pressure of H20. Note that the data in Fig. 3 can b¢ analyzed so that the results of the vacuum tests all fall on one curve, whereas those from tests

1

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conducted in environments containing significant amounts of oxygen fall on another curve. (The exact shapes of these curves will be determined after additional data are obtained.) The horizontal separation between the curves (-4 orders of magnitude) is the excess amount of 1-120 that the oxygen environment can tolerate relative to the vacuum environment, and still obtain the same ductility. Similarly, the vertical separation between the curves is the improvement in ductility obtainable by testing in oxygen---for a given moisture content.

25

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Oxygen-containing

-

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environments

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Vacuum environments

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4"''''~ 1 0 . 3.1. .0._. 2' 1 .0. . . ... 1 1"0 0. . . .

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1 0• 1. . . . 1" ~0 2" "

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Partial Pressure of Water Vapor (Pa) Fig. 3. Room temperature tensile ductility of Ni-23.4A1 as a function of partial pressure of H20. There are two possible explanations for the beneficial effect of oxygen. One explanation is that the reaction involving oxygen (4A1 + 302 ~ 2A1203) competes with the reaction involving moisture (2AI + 3H20 -~ A1203 + 6H) so that, for a given amount of residual moisture, less atomic H is available during the oxygen test than in either the low- or high-vacuum tests. Another possibility is that the presence of 6.7 × 10'~ Pa 02 in the test environment screens the specimen surfaces from the I-I20 molecules. To separate these two effects, namely the chemical reaction effect and the physical screening effect, we are currently performing additional experiments in which we compare the effects of argon and oxygen at similar water vapor contents. Regardless of the outcome of these experiments, however, it is worth noting that similar results were obtained in FeAI [4], i.e., specimens tested in oxygen exhibited higher ductilities than specimens tested in high vacuum, but not in Fe3AI, where ductilities were found to be about the same in oxygen and high vacuum [5]. These different responses to apparently identical test environments may be related to possibly different A1 activities in the different intermetallics. The other possibility is that the moisture contents in some of the early oxygen tests may not have been well controlled. Finally, in our previous paper [ 1], we examined several possible reasons for the higher ductilities obtained in polycrystalline Ni3A1 produced from single crystals, as compared to those obtained in Ni3A1 produced by conventional casting and forming operations. While the most likely reason appeared to be the presence of fewer flaws (mierocracks) in material produced from single,crystals, we could not rule out the possibility that it might be related also to the significantly higher fraction of ' special" (i.e., low angle or low E) boundaries in material produced from single crystals. Such special boundaries rarely fracture in Ni3A1 [16,17]. Consistent with this, directionally solidified Ni3A1, made by a special floating zone technique (FZ-UDS), and having a preponderance of low angle boundaries, exhibits very good ductility [18]. Interestingly, Nishimura et al. [19] have recently discovered that Ni-24.3A1 produced by FZ-UDS shows high tensile ductilities of 31, 38 and 32% in water, air and vacuum, respectively. This must mean that low angle and low Z boundaries are not very susceptible to moistureinduced environmental embrittlement. In addition, it probably also means that our present Ni3A1 alloy (produced by

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recrystallizing a single crystal) does not have many special boundaries---sinceit is severely embrittled by the water vapor in ordinary ambient air. Experiments aimed at confirming this by determining the grain boundary character distributions in our Ni3AI alloy arc currently under way. Conclusions

Polycrystalline,B-free Ni3AI (23.4 at.% AI), produced by cold working andr.ecr~YStallizinga <100> si.ng.le crystal,exhibitsroom temperature tensileductilityof 23.4% in ultrahighvacuum, inls IS me mgncst ever uuctmty reported for equiaxed, polycrystalline,B-free Ni3AI, and demonstrates that, contrary to conventional wisdom, Ni3AI is intrinsicallyquite ductile. Ductility decreases systematically with increasing amounts of air in the test environment: 23% at I0-~Pa, 13% at 10 -4 Pa, 8 % at I0-I Pa, and 3 % at 105 Pa. Moisture in air has been identified as a major cause of thisenvironmental embrittlement. Based on estimates of the moisture contents of the various test environments, the beneficialeffect of testing in oxygen appears not to bc related just to its (low) moisture content, because vacuum environments with even lower moisture contents resultin lower ductilities.W e speculate that oxygen competes with moisture for the oxidation of aluminum, thereby lowering the amount of atomic H generated for a given amount of moisture in the environment. The other possible effectof oxygen is thatitscreens the specimen surfaces from water vapor molecules, thereby suppressing the embrittlement reaction. Despite the high intrinsicductilityof Ni-23.4AI in U H V , finalfracture is predominantly intergranular,i.e.,grain boundaries remain as weak links even in the absence of moisture. Therefore, even though a significant portion of the ductilizingeffectof boron is relatedto suppression of environmental ernbrittlemcnt,part of the boron effectappears to be relatedto itsbeneficialeffecton grain boundary cohesion. ~cknowled grnents We thank E. H. Lee and P. S. Bishop for technical assistance, and P. F. Tortorelli and C. G. McKarney for reviewing the manuscript. This research was sponsored by the Division of Materials Sciences, U.S. Department of Energy under contract DE-AC05-84OR21400 with Martin Marietta Energy Systems, Inc., and by the LRSM at the University of Pennsylvaniasupported by the National Science Foundation under Grant No. DMR/MRL91-20668. References 1. 2. 3. 4. 5. 6.

7. 8. 9. 10. 11. 12. 13. 14. 15. 16. 17. 18. 19.

E.P. George, C. T. Liu, and D. P. Pope, Scripta MetaU. et Mater. 28, 857 (1993). E.P. George, C. T. Liu, and D. P. Pope, Scripta Metall. et Mater. 27, 365 (1992). C.T. Liu, Scripta Metall. et Mater. 27, 25 (1992). C.T. Liu, E. H. Lee, and C. G. McKarney, Scripta. Metall. 23, 875 (1989). C.T. Liu, C. G. McKamey, and E. H. Lee, Scripta. Metall. 24, 385 (1990). M. Shea, A. Castagna, and N. S. Stoloff, in High-Temperature Ordered lntermetallic Alloys IV, eds. L. A. Johnson, D. P, Pope, and J. O. Stiegler (Materials Research Society, Pittsburgh, PA, 1991), Vol. 213, p. 609. Handbook of Chemistry and Physics, 54th edition, CRC Press, Cleveland, OH (1973), p. F-187. Ibid., p. D-158. X.J. Wan, J. H. Zhu, and K. L. Jing, Scripta MetaU. et Mater. 26, 473 (1992). N. Masahashi, T. Takasugi, and O. Izurni, Aeta MetaU. 36, 1823 (1988). K. Aoki and O. Izumi, Nippon Kinzoku Gakkaishi 43, 1190 (1979). C.T. Liu, C. L. White, and J. A. Horton, Aeta Metall. 33, 213 (1985). S.P. Chen, A. F. Voter, R. C. Albers, A. M. Boring, and P. J. Hay, Scripta Metall. 23, 217 (1989). E.P. George, C. T. Liu, and D. P. Pope, in Proc. Intl. Symp. Structural intermetallics, Seven Springs, Champion, PA (1993), to be published. A. Roth, Vacuum Technology, Nortla-Holland, Amsterdam (1976), p. 4, S. Hanada, T. Ogura, S, Watanabe, O. Izurni, and T. Masumoto, Acta Metall. 34, 13 (1986). H. Lin and D. P. Pope, in High.Temperature Ordered lntermetallic Alloys IV, eds. L. A. Johnson, D. P. Pope, and J. O. Stiegler (Materials Research Society, Pittsburgh, PA, 1991), Vol. 213, p. 391. T. Hirano and T. Mawari, in High-Temperature Ordered lntermetallic Alloys V, eds. I. Baker, R. Darolia, J. D. Whittenberger, and M. H. Yoo (Materials Research Society, Pittsburgh, PA, 1993), Vol. 288, p. 691. C. Nishimura, T. Hirann, and M. Amano, Scripta Metall. et Mater. 29 (1993) in press.

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