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Mechanism of boron-induced embrittlement of Ni3Al in hydrogen
Scripta Materialia, Vol. 37, No. 7, pp. 1023-1029,1997 Elsevier Science Ltd Copyright 8 1997 Acta Metallurgica Inc. F’rinted in the USA. All rights reserved 1359~6462197~$17.00+ .OO
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MECHANISM OF BORON-INDUCED EMBRITTLEMENT OF N&AI IN HYDROGEN Gayle S. Painter Metals and Ceramics Division, Oak Ridge National Laboratory P.O. Box 2008, Oak Ridge, Tennessee 3783 l-6114 (Received November 19,1996) (Accepted April 14,1997) Introduction
Materials thal: have been exposed to HZ or hydrogen-containing gases often show embrittlement due to atomic hydrogen. It is now recognized that hydrogen embrittlement is the source of brittle behavior of a number of intermetallics that were previously thought to be “intrinsically brittle”. A well-known example is polycrystalline NtAl, which is susceptible to environmental embrittlement in atmospheres containing water vapor, and the mechanism of H production is the dissociation of Hz0 at active Al sites (l-3). Most ahuninides are susceptible to HzO-based environmental embrittlement, and a number of LIZ polycrystalline alloys are also embrittled by exposure to gaseous hydrogen atmospheres (4). Studies of N&Al in gaseous hydrogen have reported a susceptibility to embrittlement that is not prevented by boron additions (5). Recently an investigation of the room temperature fracture behavior of N&Al in low pressure dry hydrogen atmospheres found that NixAl showed only a slight loss in ductility on exposure to HZ (6). But when a hot tungsten filament was introduced, it served as a source of atomic H resulting in severe embrittlement. Further investigations to identify how B additions affect the flatture susceptibility of NiJAl in HZ atmospheres discovered a surprising phenomenon: boron-induced embrittlemeni (7). While B-free N&Al remained highly ductile (30% elongation to fracture) and fractured in a predominantly transgranular mode in dry Hz atmospheres at pressures of ld Pa, B-doped Ni3Al displayed brittle intergranular fracture with very limited ductility under the same conditions (see Table 1). The chemisorption of Hz on transition metal surfaces is usually dissociative (8), but experiment (6) indicates that Ni3Al polycrystalline surfaces do not readily dissociate Hz at low pressures, in spite of the high Ni-content of the alloy. Experiment (7) clearly associates boron with a HZ dissociation mechanism, and the atomic hydrogen leads to grain boundary embrittlement. Here, as with environments containing H;:O, the sonrce of the atomic hydrogen is known. The crucial question concerns how atomic H is produced and, in particular, what is the role of boron in the Hz molecular dissociation that releases atomic H. Insufficient data are available to build more than a speculative atomic-level model to explain this counterintuitive connection of boron with embrittlement. One explanation has been proposed in terms of the chemical participation of boron in the Hz dissociation step (7). However the chemistry of boron 1023
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TABLE 1 Effects of Hz Pressure and B-Content on Ductility of Ni-24 at % Al (1.2 x 10” s-’ Strain Rate) from Data of Cohron et al. (7).
does not support a direct dissociation step. And simple electronic structure considerations discussed later suggest boron would not act as a promoter in an indirect dissociation of Hz. The alternate mechanism proposed here, although tentative until further evidence is attained, is based on well-established atomic-level properties of boron in Ni3Al. In the following, the concept of the activity of interfaces with regard to HZ molecular dissociation in the region of intersection of the grain boundary with the gas/solid interface is used to explain how boron can affect atomic hydrogen production. Activity will refer exclusively to the Hz molecular dissociation reaction unless noted otherwise. Mechanism From experiment it is known that there are grain boundaries of different intrinsic strength in Ni3Al (12,13). It is also known that the segregation potential of boron in Ni 3Al depends on grain boundary character (14), and that boron enhances the cohesive strength of grain boundaries (3, 9-11). These properties form a basic part of the mechanism proposed here. Since grain boundaries vary in structure and composition, and surface molecular dissociation is often sensitive to both of these variables, it is reasonable to hypothesize that some grain boundaries are active for the dissociation of adsorbed Hz, if they are penetrated by HZ,while other interfaces are inactive for dissociation, whether exposed or not. These boundaries are labeled type A (active) and NA (inactive) respectively in Figure 1. Structure sensitivity for molecular dissociation occurs frequently in surface chemistry, and it is quite plausible in the present case, especially in light of the observed specificity of different grain boundaries with regard to B-absorption (14). In the mechanism proposed here, the structural characteristics that differentiate grain boundaries according to their strength and potential for B-segregation are associated with differences in their activity for Hz dissociation. It is assumed that type A boundaries remain largely B-free and, if exposed, present surfaces active for Hz-dissociation, by reason of composition and structural character. The reactivity of grain boundaries depends critically on the local stress and atomic-scale opening that leads to HZ adsorption and dissociation. The crack opening stress depends on the intrinsic cohesive strength of the boundary and the effect of segregants. It is reasonable to assume that active boundaries are intrinsically stronger, due to their structural character-perhaps they are more compact (close-packed) than the inactive ones-and that B segregates preferentially to the more open (and relatively weaker) distribution of inactive boundaries. In response to a fixed strain, this Bstrengthening of inactive boundaries increases strain in the distribution of (stronger) active grain boundaries. If the opening exceeds the critical amount, &, that allows Hz penetration, dissociation can occur. In the absence of boron, a fixed strain is accommodated by the distribution of weaker inactive boundaries. This reduces strain on the active boundaries, so that HZ molecules cannot penetrate them, limiting Hz dissociation (Fig. 1(a)). But boron-induced strengthening of the weaker, inactive boundaries requires more strain accommodation in the distribution of active boundaries. This allows HZ penetra-
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Figure.1. Schemlaticof mechanism of boron-inducedembrittlementof NisAI stressed in Hz atmosphere.In B-free N&Al (a) strain is accommodated by weaker boundaries inactive (NA) for H2 dissociation so that active (A) boundaries are not exposed. In B-doped N&Al l(b)NA boundaries arestrengthened by segregatedboron with result that A boundaries(unstrengthened)open for H2 penetration,Iresultingin Hz dissociation and release of H.
tion and disslociation, and atomic H diffusion into the boundary causing embrittlement (Fig. l(b)). The role of boron in promoting embrittlement in this model is then by inducing exposure of active grain boundaries by preferential B-strengthening in the distribution of inactive boundaries. The mechrmism can be illustrated by a simple example with linear local strain response to stress and two distributions of boundaries-a weaker inactive (NA) group and a stronger active (A) set as shown in Figure 2. The distributions are separated only for clarity-in actuality, a gradation in strength (12,13) and activity exists. Overlapping of the distributions is accommodated without loss of argument. In B-free NijAl a given level of strain produces grain-boundary openings accommodated primarily in the weaker inactive set, while the openings of the stronger active distribution are smaller, remaining below the critical level, EE,for HZ penetration (Fig. 2(a)). The local strain distribution in B-free Ni3Al with inactive boundaries that are weaker than the active boundaries is altered in a critical way by the presence of boron. Segregation of B to the weaker distribution of boundaries strengthens them preferentially; as a result the distribution of active boundaries will experience greater local strains for a fixed level of overall strain (Fig. 2(b)). If the increase in strain across some active boundaries exceeds the critical value for penetration of HZ molecules, the resulting onset of Hz dissociation and H-embrittlement leads to the decrease in ductility found in B-doped NisAl. The surprising association of boron with induced embrittlement thus involves no unusual activity beyond the well-known behavior that B strengthens those Ni-enriched boundaries to which it segregates. The proposed mechanism involves only well-understood phenomena and leads to behavior qualitatively consistent with the experimental observations summarized in Table 1. In the absence of atomic hydrogen, the intrinsic grain boundary cohesion is large enough to support appreciable ductility in B-free N&Al. This ductility is largely retained even in an Hz atmosphere of moderate pressure, as long as there is insignificant HZpenetration of active boundaries. But if boron doping results in E > me for a significant population of active interfaces, a pathway for molecular dissociation is established, and H-embriltlement is promoted.
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B-FRE
LOCAL.SlRAIN E (a)
LOCALSTRAIN
E
(b)
Figure 2. Schematic of B-effect on local stress/strain for distributions that are active (A) and inactive (NA) for Hz dissociation at N&AI grain boundary& gas interface. Without boron (a) a fixed strain is accommodated primarily by NA boundaries. With boron (b) NA boundaries are strengthened and show less strain, leading to greater strain in the distribution of active boundaries. For those A boundaries where E > &, Hz penetration and dissociation occur.
This mechanism is also consistent with the effect of boron in eliminating HzO-driven environmental embrittlement of N&Al. In terms of the mechanism presented here, hydrogen is produced along boundaries that are active for Hz0 dissociation. Both the chemical and structural properties of such interfaces are expected to be completely different from those that support Hz dissociation, since Al is the essential element for Hz0 dissociation, and Ni is the active element for HZ dissociation. Molecular size considerations suggest the more open (and weaker) interfaces are the active boundaries for Hz0 dissociation-in contrast to the case with Hz. Experimental results for N&Al in vucuum suggest that boron segregates to the weaker interfaces that fail by intergranular fracture in the absence of boron (6) and strengthens them (to give a 100% transgranular fracture mode (7)). In the mechanism proposed here, this is the action of boron in both Hz and Hz0 environments. However, the results of B-doping in these two cases are entirely different (7), according to the predictions of this mechanism. In the former case, boron breaks the reaction path for Hz0 dissociation by reducing the grain boundary opening under stress so that HZ0 penetration of the active boundaries does not occur. This eliminates embrittlement of NLAI in HzO-containing atmospheres. But in the latter case (HZ environment), conditions for hydrogen embrittlement are initiated by boron’s strengthening of the weaker boundaries, leading to the opening of boundaries that are active for Hz dissociation. The local effect of boron is the same in these two cases-the crucial difference is in the dissociative activity of the interface that becomes exposed by boron’s selective segregation and strengthening. The more open, weaker interfaces active for Hz0 dissociation are strengthened by boron so that Hz0 can not penetrate them, and environmental embrittlement can no longer occur. But boron strengthens the interfaces that are inactive for Hz dissociation, so that active boundaries experience greater local strain under applied stress in Hz atmospheres and allow Hz penetration, dissociation and embrittlement. Boron’s preferential segregation to the more open and presumably weaker distribution of grain boundaries in both cases is understood in terms of first-principles calculations (10,ll).
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Discussion It is well-known that the grain boundaries in off-stoichiometric N&rich N&Al are enriched in Ni, independent of the presence of boron (15). Thus composition is unlikely to be the sole factor for determining the strength of grain boundaries in NiJAl. On the other hand, the large influence of structure on the strength and ductility of grain boundaries is indicated by observations of the fracture resistance, of “special” (low angle and X3) boundaries in NtAl(12,13) and the presence of weak boundaries (such as high angle and certain C boundaries) (13). The importance of these observations with regard to the mechanism proposed here is confirmation of the presence of a distribution of boundaries having different strengths, which is not just related to Ni-content. This supports the assumption that the cohesive strength of Ni-rich grain boundaries of NiJAl is variable (related to the structural character of the boundary). Experiment strongly correlates boron segregation with grain boundary structural character. Atom probe studies (16) and more recent observations with field ion microscopy (17) indicate that the concentration of boron along an interface is highly variable. Miller and Horton also reported that the commonly occurring W twin boundary in NisAl exhibits very little B segregation (16). Recent EELS experimental work of Muller et al. (14) also shows that B segregation depends on grain boundary character and is variable, even along a particular interface. It has been observed that some portions of a Ni-enriched interface are inert for B segregation, and variability in segregation occurs even when the Ni composition along the boundary is uniform. Experiment demonstrates that both large and small angle boundaries in Ni,Al display Ni-enrichment, but exhibit very different B segregation characteristics (14). These EELS studies also confirm that boron segregation along certain (large angle) grain boundaries can vary, although the level of Ni-enrichment is uniform. No segregation of boron was detected at small angle boundaries, even though they were N&enriched. These observations clearly suggest that the structural characteristics of Ni-enriched boundaries can determine the segregation behavior of boron. This sensitivity of B segregation to grain boundary structure supports an essential part of the mechanism proposed here, because it provides an explanation for why boundaries are strengthened in a selective way. These experimental observations thus support the assumption that B segregates only minimally to the strong boundaries (14,16). By exclusion, it can be inferred that B segregation is preferential to the weaker boundaries, as assumed in this model. The EELS ‘data of Muller et al. (14) also support the proposed mechanism regarding the HZ dissociation mechanism. Electron energy loss spectroscopy probes the excitations of core-like electrons to valence levels, thus giving a measure of the local density of states (LDOS) in the volume. The experimental data for B-doped NbAl show that grain boundaries having a high B-content yield spectra (Ni L2,3edge) similar to that from bulk N&Al. On the other hand, spectra from grain boundary regions which are boron-free appear more N&like (14). This associates the electronic structure characteristics of Ni-enriched boundaries of B-free NijAl with those of Ni (an element active for HZ dissociation). This interestbig correlation of boron segregation with characteristics in the LDOS supports the mechanism proposed here, because grain boundaries having Ni-like electronic features (the active, strong boundaries in the present model) are associated with an absence of boron. The segregation of boron (to the non-active, weaker boundaries in this model) produces grain boundary bonding characteristics that are more lNisAl-like than Ni-like. So the boundaries to which B segregates would be expected to be less active for Hz dissociation than the B-free boundaries. These boron-strengthened boundaries display EELS spectral features that resemble those of bulk N&Al (14) and may well be stronger than the B-free boundaries with Ni-like character. In this proposed model, the B-tree strong boundaries become relatively weaker and show greater strain effects as a result of the B-effect on the distribution of weaker boundaries. This increased strain at the active boundaries can exceed the critical value for HZ
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penetration so that interfaces are exposed that have Ni-like character and are active for dissociation of penetrating H molecules. At the same time, however, these EELS observations are problematic for the alternative model of boron-activation/promotion (7), because the B-rich interfaces show bulk NisAl LDOS characteristics (14). More detailed d-hybridization considerations must be included to properly address Hz dissociation at surfaces (18), but to the extent that LDOS character relates to reactivity, dissociation at the Brich surface would not be consistent with the observed inactivity of N&Al surfaces for HZ dissociation (7). Nor do specific electronic structure features of the B-doped surface match the requirements for the spontaneous dissociation of HZ, since the presence of boron decreases the number of d-band holes that serve to lower the entrance channel barrier for Hz chemisorption (19). While it appears difficult to reconcile an active species or promoter role of boron with these electronic structure factors, the mechanism of boron-induced exposure of active grain boundaries is consistently accommodated. The ductile behavior of the B-free samples deserves special comment. At the highest pressure Hz atmospheres, NisAI shows a 30% elongation to fracture (Table l), representing a 25% reduction from the value at lo-’ Pa H2 pressure. Failure in the higher pressure environment ( lo3 Pa) is mixed-mode (about 45% intergranular compared with 30% intergranular at 1U’ Pa H2 pressure), so some environmental effect is manifest without boron. The conditions in this model for an absence of embrittlement requires that stress levels at active boundaries produce strains less than the critical value for Hz penetration. Active boundaries must be intrinsically strong enough that they support the stress levels that cause plastic deformation in some grains without opening themselves for Hz penetration, at least until stress levels are reached that cause the observed mixed-mode failure. At failure the percentage of boundaries showing intergranular fracture is 45% in this case, but the distribution among inactive and active boundaries is not known. Operative modes for failure of the B-free samples at lo3 Pa HZ pressure include: 1) weaker, inactive boundary failure (the principal mode in the proposed mechanism), 2) a small number of “inactive” boundaries show some non-vanishing level of activity, 3) a small percentage of active boundaries exceed the critical strain value. Certainly a combination of these factors can be operative and fall within the scope of the proposed mechanism. On the other hand, the 15% increase in intergranular fracture observed for these samples in going from 10’ to lo3 Pa HZ pressure (7) is unrelated to boron, since there is no boron present in this case. Conclusion The mechanism proposed here is attractive in its conceptual simplicity and is at least qualitatively consistent with a large set of data. This simplicity reflects more our lack of knowledge of atomic-level detail than any inherence of the model. The principal effect (B-induced exposure of active boundaries) should persist with refinements. The embrittlement mechanism as described in this model is plausible, and the role of boron involves only the ordinary strengthening of Ni-rich boundaries to which boron segregates. This mechanism also offers an explanation for the effect of boron in eliminating embrittlement of N&Al in HzO-containing atmospheres. In both cases boron is hypothesized to control the production rate of atomic H indirectly by affecting the exposure of grain boundary interfaces that are active for molecular dissociation (HZ or Hz0 in each case). Further development and verification of this model requires atomic-level information about the distribution of strain over the population of grain boundaries and how active and inactive boundaries differ in their crack nucleation abilities. Extensions of this investigation to address strain effects on the energy barrier for Hz penetration of N&Al grain boundaries and the effect of boron would benefit from experimental study of Hz chemisorption on well-characterized surfaces of NtAl.
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Acknowledgments This work was supported by the Division of Materials Science, Office of Basic Energy Sciences, U. S. Department of Energy under contract DE-AC05-960R22464 with Martin Marietta Energy Systems, Inc. The author thanks J.W. Cohron, E.P. George, C.T. Liu and C.L. Fu for helpful discussions on this problem. References 1. 2. 3. 4. 5. 6. 7. 8. 9. 10. 11. 12. 13. 14. 15. 16. 17. 18. 19.
CT. Liu, Scripta Metall. et Mater. 27,25 (1992). E.P. George, C.T. Liu and D.P. Pope, Scripta Metall. et Mater. 30,37 (1994). E.P. George, C.T. Liu and D.P. Pope, Acta mater. 44, 1757 (1996). N.S. Stoloff rmd CT. Liu, Intermetallics 2,75 (1994). K.H. Lee and1 C.L. White, Scripta Metall. et Mater. 33,129 (1995). J.W. Cohron, E.P. George, L. Heatherly, C.T. Liu and R.H. Zee, Intermetallics 4,497 (1996). J.W. Cohron, E.P. George, L. Heatherly, C.T. Liu and R.H. Zee, Acta mater. (1997) in press. B. Poelsema, L.K. Verheij and G. Cornsa, Surface Science 148, 101 (1985). G.S. Painter rmd F.W. Averill, Phys. Rev. Lett. 58,234 (1987). CL. Fu and G.S. Painter, Acta mater. (1996), in press. G.S. Painter, C.L. Fu and F.W. Averill, J. Appl. Phys. (1996), in press. A. Chiba, S. IHanada, S. Watanabe, T. Abe and T. Obana, Acta Metall. Mater. 42, 1733 (1994). H. Lin and D.P. Pope, Mater. Sci. Eng. A192/193,349 (1995). D.A. Muller, S. Subramanian, P.E. Batson, J. Silcox and S.L. Sass, Acta mater. 44, 1637 (1996). E.P. George, CT. Liu and R.A. Padgett, Scripta Metah. 23,979 (1989). M.K. Miller rmd J.A. Horton, Scripta Metall. 20,789 (1986). S.S. Bremrer and H. MingJian, Scripta Metall. 25, 1271 (1991). B. Hammer and J.K. Norskov, Surface Science 343,211(1995). J. Harris, S. Andersson, C. Holmbcrg and P. Nordlander, Phys. Scripta T13,155 (1986).
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MECHANISM OF BORON-INDUCED EMBRITTLEMENT OF N&AI IN HYDROGEN Gayle S. Painter Metals and Ceramics Division, Oak Ridge National Laboratory P.O. Box 2008, Oak Ridge, Tennessee 3783 l-6114 (Received November 19,1996) (Accepted April 14,1997) Introduction
Materials thal: have been exposed to HZ or hydrogen-containing gases often show embrittlement due to atomic hydrogen. It is now recognized that hydrogen embrittlement is the source of brittle behavior of a number of intermetallics that were previously thought to be “intrinsically brittle”. A well-known example is polycrystalline NtAl, which is susceptible to environmental embrittlement in atmospheres containing water vapor, and the mechanism of H production is the dissociation of Hz0 at active Al sites (l-3). Most ahuninides are susceptible to HzO-based environmental embrittlement, and a number of LIZ polycrystalline alloys are also embrittled by exposure to gaseous hydrogen atmospheres (4). Studies of N&Al in gaseous hydrogen have reported a susceptibility to embrittlement that is not prevented by boron additions (5). Recently an investigation of the room temperature fracture behavior of N&Al in low pressure dry hydrogen atmospheres found that NixAl showed only a slight loss in ductility on exposure to HZ (6). But when a hot tungsten filament was introduced, it served as a source of atomic H resulting in severe embrittlement. Further investigations to identify how B additions affect the flatture susceptibility of NiJAl in HZ atmospheres discovered a surprising phenomenon: boron-induced embrittlemeni (7). While B-free N&Al remained highly ductile (30% elongation to fracture) and fractured in a predominantly transgranular mode in dry Hz atmospheres at pressures of ld Pa, B-doped Ni3Al displayed brittle intergranular fracture with very limited ductility under the same conditions (see Table 1). The chemisorption of Hz on transition metal surfaces is usually dissociative (8), but experiment (6) indicates that Ni3Al polycrystalline surfaces do not readily dissociate Hz at low pressures, in spite of the high Ni-content of the alloy. Experiment (7) clearly associates boron with a HZ dissociation mechanism, and the atomic hydrogen leads to grain boundary embrittlement. Here, as with environments containing H;:O, the sonrce of the atomic hydrogen is known. The crucial question concerns how atomic H is produced and, in particular, what is the role of boron in the Hz molecular dissociation that releases atomic H. Insufficient data are available to build more than a speculative atomic-level model to explain this counterintuitive connection of boron with embrittlement. One explanation has been proposed in terms of the chemical participation of boron in the Hz dissociation step (7). However the chemistry of boron 1023
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TABLE 1 Effects of Hz Pressure and B-Content on Ductility of Ni-24 at % Al (1.2 x 10” s-’ Strain Rate) from Data of Cohron et al. (7).
does not support a direct dissociation step. And simple electronic structure considerations discussed later suggest boron would not act as a promoter in an indirect dissociation of Hz. The alternate mechanism proposed here, although tentative until further evidence is attained, is based on well-established atomic-level properties of boron in Ni3Al. In the following, the concept of the activity of interfaces with regard to HZ molecular dissociation in the region of intersection of the grain boundary with the gas/solid interface is used to explain how boron can affect atomic hydrogen production. Activity will refer exclusively to the Hz molecular dissociation reaction unless noted otherwise. Mechanism From experiment it is known that there are grain boundaries of different intrinsic strength in Ni3Al (12,13). It is also known that the segregation potential of boron in Ni 3Al depends on grain boundary character (14), and that boron enhances the cohesive strength of grain boundaries (3, 9-11). These properties form a basic part of the mechanism proposed here. Since grain boundaries vary in structure and composition, and surface molecular dissociation is often sensitive to both of these variables, it is reasonable to hypothesize that some grain boundaries are active for the dissociation of adsorbed Hz, if they are penetrated by HZ,while other interfaces are inactive for dissociation, whether exposed or not. These boundaries are labeled type A (active) and NA (inactive) respectively in Figure 1. Structure sensitivity for molecular dissociation occurs frequently in surface chemistry, and it is quite plausible in the present case, especially in light of the observed specificity of different grain boundaries with regard to B-absorption (14). In the mechanism proposed here, the structural characteristics that differentiate grain boundaries according to their strength and potential for B-segregation are associated with differences in their activity for Hz dissociation. It is assumed that type A boundaries remain largely B-free and, if exposed, present surfaces active for Hz-dissociation, by reason of composition and structural character. The reactivity of grain boundaries depends critically on the local stress and atomic-scale opening that leads to HZ adsorption and dissociation. The crack opening stress depends on the intrinsic cohesive strength of the boundary and the effect of segregants. It is reasonable to assume that active boundaries are intrinsically stronger, due to their structural character-perhaps they are more compact (close-packed) than the inactive ones-and that B segregates preferentially to the more open (and relatively weaker) distribution of inactive boundaries. In response to a fixed strain, this Bstrengthening of inactive boundaries increases strain in the distribution of (stronger) active grain boundaries. If the opening exceeds the critical amount, &, that allows Hz penetration, dissociation can occur. In the absence of boron, a fixed strain is accommodated by the distribution of weaker inactive boundaries. This reduces strain on the active boundaries, so that HZ molecules cannot penetrate them, limiting Hz dissociation (Fig. 1(a)). But boron-induced strengthening of the weaker, inactive boundaries requires more strain accommodation in the distribution of active boundaries. This allows HZ penetra-
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t
0
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Figure.1. Schemlaticof mechanism of boron-inducedembrittlementof NisAI stressed in Hz atmosphere.In B-free N&Al (a) strain is accommodated by weaker boundaries inactive (NA) for H2 dissociation so that active (A) boundaries are not exposed. In B-doped N&Al l(b)NA boundaries arestrengthened by segregatedboron with result that A boundaries(unstrengthened)open for H2 penetration,Iresultingin Hz dissociation and release of H.
tion and disslociation, and atomic H diffusion into the boundary causing embrittlement (Fig. l(b)). The role of boron in promoting embrittlement in this model is then by inducing exposure of active grain boundaries by preferential B-strengthening in the distribution of inactive boundaries. The mechrmism can be illustrated by a simple example with linear local strain response to stress and two distributions of boundaries-a weaker inactive (NA) group and a stronger active (A) set as shown in Figure 2. The distributions are separated only for clarity-in actuality, a gradation in strength (12,13) and activity exists. Overlapping of the distributions is accommodated without loss of argument. In B-free NijAl a given level of strain produces grain-boundary openings accommodated primarily in the weaker inactive set, while the openings of the stronger active distribution are smaller, remaining below the critical level, EE,for HZ penetration (Fig. 2(a)). The local strain distribution in B-free Ni3Al with inactive boundaries that are weaker than the active boundaries is altered in a critical way by the presence of boron. Segregation of B to the weaker distribution of boundaries strengthens them preferentially; as a result the distribution of active boundaries will experience greater local strains for a fixed level of overall strain (Fig. 2(b)). If the increase in strain across some active boundaries exceeds the critical value for penetration of HZ molecules, the resulting onset of Hz dissociation and H-embrittlement leads to the decrease in ductility found in B-doped NisAl. The surprising association of boron with induced embrittlement thus involves no unusual activity beyond the well-known behavior that B strengthens those Ni-enriched boundaries to which it segregates. The proposed mechanism involves only well-understood phenomena and leads to behavior qualitatively consistent with the experimental observations summarized in Table 1. In the absence of atomic hydrogen, the intrinsic grain boundary cohesion is large enough to support appreciable ductility in B-free N&Al. This ductility is largely retained even in an Hz atmosphere of moderate pressure, as long as there is insignificant HZpenetration of active boundaries. But if boron doping results in E > me for a significant population of active interfaces, a pathway for molecular dissociation is established, and H-embriltlement is promoted.
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B-FRE
LOCAL.SlRAIN E (a)
LOCALSTRAIN
E
(b)
Figure 2. Schematic of B-effect on local stress/strain for distributions that are active (A) and inactive (NA) for Hz dissociation at N&AI grain boundary& gas interface. Without boron (a) a fixed strain is accommodated primarily by NA boundaries. With boron (b) NA boundaries are strengthened and show less strain, leading to greater strain in the distribution of active boundaries. For those A boundaries where E > &, Hz penetration and dissociation occur.
This mechanism is also consistent with the effect of boron in eliminating HzO-driven environmental embrittlement of N&Al. In terms of the mechanism presented here, hydrogen is produced along boundaries that are active for Hz0 dissociation. Both the chemical and structural properties of such interfaces are expected to be completely different from those that support Hz dissociation, since Al is the essential element for Hz0 dissociation, and Ni is the active element for HZ dissociation. Molecular size considerations suggest the more open (and weaker) interfaces are the active boundaries for Hz0 dissociation-in contrast to the case with Hz. Experimental results for N&Al in vucuum suggest that boron segregates to the weaker interfaces that fail by intergranular fracture in the absence of boron (6) and strengthens them (to give a 100% transgranular fracture mode (7)). In the mechanism proposed here, this is the action of boron in both Hz and Hz0 environments. However, the results of B-doping in these two cases are entirely different (7), according to the predictions of this mechanism. In the former case, boron breaks the reaction path for Hz0 dissociation by reducing the grain boundary opening under stress so that HZ0 penetration of the active boundaries does not occur. This eliminates embrittlement of NLAI in HzO-containing atmospheres. But in the latter case (HZ environment), conditions for hydrogen embrittlement are initiated by boron’s strengthening of the weaker boundaries, leading to the opening of boundaries that are active for Hz dissociation. The local effect of boron is the same in these two cases-the crucial difference is in the dissociative activity of the interface that becomes exposed by boron’s selective segregation and strengthening. The more open, weaker interfaces active for Hz0 dissociation are strengthened by boron so that Hz0 can not penetrate them, and environmental embrittlement can no longer occur. But boron strengthens the interfaces that are inactive for Hz dissociation, so that active boundaries experience greater local strain under applied stress in Hz atmospheres and allow Hz penetration, dissociation and embrittlement. Boron’s preferential segregation to the more open and presumably weaker distribution of grain boundaries in both cases is understood in terms of first-principles calculations (10,ll).
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Discussion It is well-known that the grain boundaries in off-stoichiometric N&rich N&Al are enriched in Ni, independent of the presence of boron (15). Thus composition is unlikely to be the sole factor for determining the strength of grain boundaries in NiJAl. On the other hand, the large influence of structure on the strength and ductility of grain boundaries is indicated by observations of the fracture resistance, of “special” (low angle and X3) boundaries in NtAl(12,13) and the presence of weak boundaries (such as high angle and certain C boundaries) (13). The importance of these observations with regard to the mechanism proposed here is confirmation of the presence of a distribution of boundaries having different strengths, which is not just related to Ni-content. This supports the assumption that the cohesive strength of Ni-rich grain boundaries of NiJAl is variable (related to the structural character of the boundary). Experiment strongly correlates boron segregation with grain boundary structural character. Atom probe studies (16) and more recent observations with field ion microscopy (17) indicate that the concentration of boron along an interface is highly variable. Miller and Horton also reported that the commonly occurring W twin boundary in NisAl exhibits very little B segregation (16). Recent EELS experimental work of Muller et al. (14) also shows that B segregation depends on grain boundary character and is variable, even along a particular interface. It has been observed that some portions of a Ni-enriched interface are inert for B segregation, and variability in segregation occurs even when the Ni composition along the boundary is uniform. Experiment demonstrates that both large and small angle boundaries in Ni,Al display Ni-enrichment, but exhibit very different B segregation characteristics (14). These EELS studies also confirm that boron segregation along certain (large angle) grain boundaries can vary, although the level of Ni-enrichment is uniform. No segregation of boron was detected at small angle boundaries, even though they were N&enriched. These observations clearly suggest that the structural characteristics of Ni-enriched boundaries can determine the segregation behavior of boron. This sensitivity of B segregation to grain boundary structure supports an essential part of the mechanism proposed here, because it provides an explanation for why boundaries are strengthened in a selective way. These experimental observations thus support the assumption that B segregates only minimally to the strong boundaries (14,16). By exclusion, it can be inferred that B segregation is preferential to the weaker boundaries, as assumed in this model. The EELS ‘data of Muller et al. (14) also support the proposed mechanism regarding the HZ dissociation mechanism. Electron energy loss spectroscopy probes the excitations of core-like electrons to valence levels, thus giving a measure of the local density of states (LDOS) in the volume. The experimental data for B-doped NbAl show that grain boundaries having a high B-content yield spectra (Ni L2,3edge) similar to that from bulk N&Al. On the other hand, spectra from grain boundary regions which are boron-free appear more N&like (14). This associates the electronic structure characteristics of Ni-enriched boundaries of B-free NijAl with those of Ni (an element active for HZ dissociation). This interestbig correlation of boron segregation with characteristics in the LDOS supports the mechanism proposed here, because grain boundaries having Ni-like electronic features (the active, strong boundaries in the present model) are associated with an absence of boron. The segregation of boron (to the non-active, weaker boundaries in this model) produces grain boundary bonding characteristics that are more lNisAl-like than Ni-like. So the boundaries to which B segregates would be expected to be less active for Hz dissociation than the B-free boundaries. These boron-strengthened boundaries display EELS spectral features that resemble those of bulk N&Al (14) and may well be stronger than the B-free boundaries with Ni-like character. In this proposed model, the B-tree strong boundaries become relatively weaker and show greater strain effects as a result of the B-effect on the distribution of weaker boundaries. This increased strain at the active boundaries can exceed the critical value for HZ
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penetration so that interfaces are exposed that have Ni-like character and are active for dissociation of penetrating H molecules. At the same time, however, these EELS observations are problematic for the alternative model of boron-activation/promotion (7), because the B-rich interfaces show bulk NisAl LDOS characteristics (14). More detailed d-hybridization considerations must be included to properly address Hz dissociation at surfaces (18), but to the extent that LDOS character relates to reactivity, dissociation at the Brich surface would not be consistent with the observed inactivity of N&Al surfaces for HZ dissociation (7). Nor do specific electronic structure features of the B-doped surface match the requirements for the spontaneous dissociation of HZ, since the presence of boron decreases the number of d-band holes that serve to lower the entrance channel barrier for Hz chemisorption (19). While it appears difficult to reconcile an active species or promoter role of boron with these electronic structure factors, the mechanism of boron-induced exposure of active grain boundaries is consistently accommodated. The ductile behavior of the B-free samples deserves special comment. At the highest pressure Hz atmospheres, NisAI shows a 30% elongation to fracture (Table l), representing a 25% reduction from the value at lo-’ Pa H2 pressure. Failure in the higher pressure environment ( lo3 Pa) is mixed-mode (about 45% intergranular compared with 30% intergranular at 1U’ Pa H2 pressure), so some environmental effect is manifest without boron. The conditions in this model for an absence of embrittlement requires that stress levels at active boundaries produce strains less than the critical value for Hz penetration. Active boundaries must be intrinsically strong enough that they support the stress levels that cause plastic deformation in some grains without opening themselves for Hz penetration, at least until stress levels are reached that cause the observed mixed-mode failure. At failure the percentage of boundaries showing intergranular fracture is 45% in this case, but the distribution among inactive and active boundaries is not known. Operative modes for failure of the B-free samples at lo3 Pa HZ pressure include: 1) weaker, inactive boundary failure (the principal mode in the proposed mechanism), 2) a small number of “inactive” boundaries show some non-vanishing level of activity, 3) a small percentage of active boundaries exceed the critical strain value. Certainly a combination of these factors can be operative and fall within the scope of the proposed mechanism. On the other hand, the 15% increase in intergranular fracture observed for these samples in going from 10’ to lo3 Pa HZ pressure (7) is unrelated to boron, since there is no boron present in this case. Conclusion The mechanism proposed here is attractive in its conceptual simplicity and is at least qualitatively consistent with a large set of data. This simplicity reflects more our lack of knowledge of atomic-level detail than any inherence of the model. The principal effect (B-induced exposure of active boundaries) should persist with refinements. The embrittlement mechanism as described in this model is plausible, and the role of boron involves only the ordinary strengthening of Ni-rich boundaries to which boron segregates. This mechanism also offers an explanation for the effect of boron in eliminating embrittlement of N&Al in HzO-containing atmospheres. In both cases boron is hypothesized to control the production rate of atomic H indirectly by affecting the exposure of grain boundary interfaces that are active for molecular dissociation (HZ or Hz0 in each case). Further development and verification of this model requires atomic-level information about the distribution of strain over the population of grain boundaries and how active and inactive boundaries differ in their crack nucleation abilities. Extensions of this investigation to address strain effects on the energy barrier for Hz penetration of N&Al grain boundaries and the effect of boron would benefit from experimental study of Hz chemisorption on well-characterized surfaces of NtAl.
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Acknowledgments This work was supported by the Division of Materials Science, Office of Basic Energy Sciences, U. S. Department of Energy under contract DE-AC05-960R22464 with Martin Marietta Energy Systems, Inc. The author thanks J.W. Cohron, E.P. George, C.T. Liu and C.L. Fu for helpful discussions on this problem. References 1. 2. 3. 4. 5. 6. 7. 8. 9. 10. 11. 12. 13. 14. 15. 16. 17. 18. 19.
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