- Email: [email protected]
The influence of boron segregation on the structure and mechanical properties of boundaries in bicrystals of Ni3Al
Scripta METALLURGICA et MATERIALIA
Vol. 2S, pp. 1283-1288, 1991 Printed in the U.S.A.
Pergamon
VIEWPOINT
Press plc
SET No. 17
THE INFLUENCE OF BORON SEGREGATION ON THE STRUCTURE AND MECHANICAL PROPERTIES OF BOUNDARIES IN BICRYSTALS OF Ni3Al M. J. Mills, S. H. Goods, S. M. Foiles and J. R. Whetstonet Sandia National Laboratories, Livermore, CA 94551-0969 tAllison Gas Turbines Division, General Motors Corporation, Indianapolis, IN 46206-0420
(Received A p r i l 8, 1991} InQxx]uction In contrast to the high ductilities reported for single crystals of ordered (y) Ni3AI, the ductility in polycrystalline form is extremely limited. The discovery by Aoki and Izumi [1] that microalioying with boron greatly increases the ductility of polycrystallineNi3A1 has initiated spirited scientific investigation into the structure and pmperdes of grain boundaries in this material. One consequence of the addition of boron could be an increase in the cohesive properties of the grain boundaries. However, at present the most widely accepted explanation for the improved ductility of the grain boundaries is that boron causes compositional disordering of the grain boundary region. As first suggested by King and Yon [2], the total number of possible dislocation reactions is generally much larger at a grain boundary in disordered (T) material (the FCC structure) than in the ordered L12 phase. Consequently, for the disordered case, the absorption or transmission of dislocations across the boundary should be easier, large stress concentrations should not develop as readily at the boundary, and the propensity for intergranular fracture should be reduced. Although several high resolution transmission electron microscopy (HRTEM) studies have been performed to date, direct experimental evidence concerning disorder at grain boundaries in horon-doped Ni3AI is presently inconclusive. Mackenzie and Sass [3] have reported the presence of a 1.5 to 2.5 nm thick disordered region at a grain boundary in boron-doped and directionally solidifu:d Ni-24 at.% AI. However, observations by Mills [4] on the same material have failed to reveal such a region. In the latter study, results of image simulations indicated that fine, FCC fringe spacings can be produced at completely ordered grain boundaries if they are slightly inclined to the electron beam, creating the false impression that the boundary region is disordered. These results emphasize the need for well-oriented and carefully controlled boundary geomeiries for HRTEM studies. Finally, Krzanowski [5] performed HRTEM imasing and chemical analysis on chill-cast Ni-24 at.% AI with and without boron. In this nonequilibrium condition, it was found that the boundaries in both alloys were nickel-rich. Nevertheless, superlattice fringe spacings were observed to extend close to the boundary in both the binary and horon-doped material. A fundamental shortcoming of these previous investigations is that the observed boundaries are not necessarily representative of all the grain boundaries in the sample, and may not be the ones responsible for determining the macroscopic behavior. This is particularly true of HRTEM studies in which only limited grain boundary areas are examined. In order to directly link grain boundary structure with macroscopic properties, it is necessary to isolate a single boundary for study. To enable this direct correlation of microscopic structure with ductility, bicrystais have been used in the present investigation. The use of bicrystals with suitable crystallographic orientations is also advantageous with respect to the HRTEM observations since (a) lardce fringes may be seen in the crystals on either side of the boundary and Co) the boundary is assured to be close to an "edge-on" configuration. These factors should limit the possible imaging artifacts which can arise in the HRTEM studies and should also allow the degree of disordering to he evaluated very close to ~ e boundary itself. In this investigation, boron-doped bicrystals of Ni3AI (24 at.% AI + 1 at.% Ta) have been studied using HRTEM, while companion specimens were mechanically deformed in tension. The results of these experimental studies will also be discussed in light of Monte Carlo atomistic calculations which have been performed to evaluate the effect of homn on grain boundary structure and ordering. Ex~*rimental Procedu~ The bicrystal used in this study was produced via a modified single crystal investment casting process. Single crystal seeds were oriented in mold fixtures with a common [001] parallel to the growth direction. While the two seeds were oriented in order to obtain the desired ~,5 (120) tilt botmdary geometry, the resulting bicrystal was found to have an additional tilt misorientafion of about 7 ° about the [100] in one of the crystals. The bulk composition of the sample is 23 at% AI, 1 at% Ta, 0.27 at% B, and the balance Ni. The biorystal is therefore slightly Ni-rich relative to the stoichiometric composition. The Ta is substitutional on AI sites [6] and was ~dd~ to assist the growth
1283 0036-9748/91 $3.00 + .00
1284
BORON
SEGREGATION
Vol.
25, No.
of the crystals. Although the ternary addition of Ta causes an increase in the flow stress at low temperatures [7], it does not strongly affect the f~-ture behavior of polycrystalline material. Tensile specimens with reduced gage sections were spark cut from the as-grown bicrystaL In order to maximize the tensile stresses across the interface, the specimens were cut so that the tensile axis was normal to the nominal orientation of the grain boundary plane. Prior to mechanical testing, one specimen was vacuum annealed (< 1 x 10-4 Pa) at 1000°C for 48 hours, then slow cooled at 50°C/hour. In previous work, Auger electron spectroscopy has shown that this beat treatment maximizes the boron concentration at the grain boundary [8]. A second specimen was sealed in an evacuated quartz ampule and annealed at the same temperature for the same length of time. However, rather than being slow cooled, it was rapidly quenched in a water bath. This heat treatment should freeze in a homogeneous distribution of boron throughout the specimen with minimum segregation to the boundary. A final specimen was tested in the as-grown condition which is expected to have an intermediate cooling rate and boron segregation. All three specimens were tensile tested at room temperature at a constant extension rate of = 4 x 10-3 mm/sec. Foils for Iransmission electron microscopy were prepared from slices cut normal to the growth direction of the bicrystai and anneaied in an identical manner to the tensile test specimens. Observations were conducted using a JEOL 4000EX operating at 400 kV. Image simulations have been performed using the software package of Stadelmann [9].
Results
MechanicalTesting The tensile curves for the three different specimens are shown in Figure I. It is clear that the different thermal histories designed to vary the boron concentrations at the interface had a significanteffect on ductility. The annealed and slow-cooled (SC) specimen exhibited the largest strain to failure, 45%, while the as-cast (AC) specimen failed at a total swain of 28%. In contrast, the annealed and quenched (Q) specimen failed after a total strain of only 2.5%. It should be further noted that at low strains, the curves for the AC- and the SC-specimens superimpose, indicating that the anneal had no effect on the flow characteristics for these two conditions. In concert with these changes in ductility, the different heat treatments also affected the fracture morphologies of the samples. In Figure 2, the specimen marked as "a" is undeformed. The interface is located near the center of the gauge section, and approximately normal to the tensile axis. For the high ductility SC-specimen (marked as "c" in Figure 2), failure actually occurred in the grip region, away from the grain boundary-presumably at a strain concentration arising from the specimen geometry. For the AC-specimen, marked as "b", fracture occurred at the interface. This specimen clearly exhibited far less axial strain prior to failure. The complex path of the fracture surface reveals that the interface was not planar. Rather, it varied considerably from the intended (120) boundary plane. The Q-specimen is not shown. However, it failed in a manner identical to that for the AC-specimen, albeit with even less axial extension. These results are consistent with those reported by Cboudhury, et al. [8] for polycrystalline Ni-24 at% Al doped with boron. In their work, relatively small increases in the grain boundary boron concentration were apparently responsible for the transition from brittle, intergranular fracture to ductile failure. The behavior of the Q-specimen also indicates that, in the absence of boron, this bicrystal boundary would be intrinsically brittle and was not a "special" boundary across which slip was favored. From the observed distortions, it is evident that the amount of strain is not the same in both crystals. In fact, the reduction in area in the upper crystal was precisely half that of the lower crystal. This difference in strain is due to the fact that the biorystal is misoriented from the ideal Y.,5(120) geometry, and therefore the tensile axis is not of the same.type in the two crystals. The tensile direction in the lower crystal is close to the expected [120], so that both the a[011] ( I I I ) and a[011](ll]) slip systems _havelarge resolved shear stresses. The upper crystal is rotated about 7° from the [120], so that only the a[011](III) system is active, as confirmed by slip trace analysis. The large and abrupt slrain discontinuity apparent in Figure 2d indicates that slip transfer across the grain boundary was difficult. This observation further emphasizes the fact that the boundary was not in a uniquely favorable orientation for dislocation transmission. In fact, even for an ideal Y~5(120) boundary under a tensile load, the mmsmission of dislocations from one crystal to the other is complicated since none of the active dislocations (or their partials) can be readily decomposed into grain boundary dislocations, based on the considerations of King and Yoo [2]. In spite of the large stress concentrations that must have developed at the boundary during the deformation of the annealed sample, die boundary remained intact throughout the test. Using HRTEM, the validity of correlating this high ductility with the possible presence of a compositionally disordered region at the grain boundary will now be evaluated.
6
Vol. 25, No. 6
BORON SEGREGATION
800
"
|
!
1285
t
i
Ni3(AI, Ta) + 0.27% B Bicrystal A m
a.
600 J
v
P
400 (AC)
Q
2
I-
SlowCooled (so)
200
~ 0 u.O
nnealed,~ e n e l ' ~ •
(q)
0.1
,
,
0.2 True
,
0.3 Strain
,
,
,
0.4
0.I
Figure 1: Flow curves for tensile tests of Ni3(A1, 1 at% Ta) + 0.27 at% B bicrystals for three heat treatment conditions.
a
b
c
d
Figure 2: Mechanical test specimens (a) before tesdng, (b) LCtcrtesting in the as-cast condition, and (c) after testing in the annealed and slow cooled condition. In (d) a slightly magnified side view of the annealed and slow cooled sample is shown, and the location of the grain boundary (GB) is indicated. Note the difference in specimen thickness indicated by arrows in (d).
1286
BORON SEGREGATION
Vol.
25, No. 6
HRTEM Observations An HRTEM image of the bicrystal boundary in the SC condition is shown in Figure 3. The image reflects the fact the bicrystal is not in the ideal Y5 (120) geometry. As a result, only [100] lattice fringes are observed in the lower crystal, while [100] and [010] cross-fringes can be seen in the upper crystal. The inset simulated images have been calculated for both the ordered and the disordered structures, and for the estimated orientation of both crystals. Comparison of these simulations with the observed image shows that the slructaxe of the bulk crystals are consistent with ordered ? Ni3A1. Although the deviation of this boundary from the pure tilt geometry precluded a determination of the atomic structure of the boundary, local variations in the orientation of the interface plane can be identified. Since both crystals exhibited at least one set of lattice fringes, any overlap between the crystals due to an inclination of the boundary from the "edge-on" orientation would result in the appearance of Mdir~ fi'inges. These interference fringes can be seen in the image by sighting along the direction indicated by the arrow in Figure 3. For the evaluation of local compositional disorder, it is extremely important that the grain boundary be as close to the "edge-on" orientation as possible [4]. This image was obtained after orienting the bicrystal in the microscope so that the extent of the Moir~ fringes was minimized while remaining as close as possible to the [001] zone axis in both crystals. The significant result of these HRTEM observations is that the ordering present in the bulk crystals extended up to within a distance of 0.5 nm from the boundary with no detectable loss of ordering. These results demonstrate that the high ductility exhibited by this hicrystal in the SC condition was not due to the presence of a distinct disordered region at the boundary. In fact, evidence for such a disordered region was not observed in either the AC or the SC condition. While the dramatic reduction in the ductility of the specimen in the Q condition suggests that this bicrystal boundary was inherently brittle and that the segregation of boron was responsible for the increased ductility, it is clear that this ducfilization was not brought about by the formation of a broad disordered region at the boundary. Atomistic Calculations of Boron Effects on Grain Boundary Disordering, Structure and Energy If disorder is present at these boundaries, the HRTEM results indicate that it must be localized to the atomic planes immediately adjacent w the boundary. Monte Carlo atomisfic calculations [I0] have been performed in order to provide an estimate of the extent of boron-induced disorder. A ]:5 (130) symmetric tilt boundary has been considered, with a bulk composition of 76 at.% Ni and two boron atoms per boundary period for the boron-doped case. While this is not precisely the same boundary geometry as that studied experimentally, the results are expected to be qualitatively similar. The calculations utilized the Embedded Atom Method (EAM) potentials for the nickel, aluminum and boron system developed by Voter and Chen [11]. These calculations allow for the compositional rearrangement of the system as well as the spatial displacement of the atoms. They were performed for both binary and boron-doped boundaries and at temperatures of 227 and 727oC. In the simulations, the boron is not allowed to diffuse from the boundary, even though it might do so in the real system at the higher temperatures considered. For calculations both with and without boron, a localized region of reduced ordering is predicted to occur at the boundary. In all cases, a partially disordered region, characterized by an increased concenuation of anti-site defects, extends 2 to 3 (130) atomic planes on either side of the boundary. This results from the reduction in anti-site defect energies for many of the sites near the boundary. The amount of disorder produced in the absence of boron was relatively small; about 8 % of the Ni sites and 30% of the AI sites within 0.5 nm of the boundary have anti-site defects in the simulations at 727°C. (In the bulk at this temperature, 1% of the Ni sites and 7.5% of the Al sites contain anti-site defects.) The presence of boron increases the amount of disorder seen; about 15% of the Ni sites and 34% of the AI sites near the boundary have anti-site defects. The increase in disorder caused by the boron is due to a competition between two factors. First, the atoms want to arrange so as to maximize the number of Ni-A1bonds and minimize the number of AI-A1 bonds since the former are stronger. This is the driving force for the creation of the ordered crystal lattice. On the other hand, it is energetically preferable to place Ni atoms next to the B atoms instead of AI atoms. Since it is not possible to achieve both of these conditions simultaneously, it is easier to induce local disorder in the presence of the boron. The degree of ordering at the boundary is reflected in the image simulations for both swactures shown in Figure 4a and 4b. The strong intensities in the images are due to the ordered superlattice. These intensities are clearly reduced near the boundary for the boron-doped case, while they axe stronger at the boron-free boundary. The effect of partial disordering on these images can be conWastedwith the image simulation in Figure 4c. This simulation is based on a structure which has been statically relaxed so that only spatial displacements of the atoms are allowed. Hence, the region near the boundary remains fully ordered and no reduction in the superlattice intensifies is observed at the boundary. These simulations suggest that it should be possible using HRTEM to distinguish between a fully ordered grain boundary and one in which there is only a narrow region of partial disorder. The addition of boron could also cause an increase in the cohesive strength of the boundaries. Chen, et al. [12] have used the EAM to test this possibility and found that for completely ordered boundaries the addition of
Vol.
25, No. 6
BORON S E G R E G A T I O N
1287
Figure 3: HRTEM micrograph of the bierystal boundary obtained for an objective lens defocus of about 60 nm and a crystal thickness of about 18 nm. The upper crystal is oriented close to the [001] zone axis while the lower crystal is near the [0 1 10] zone axis. Inset image simulations have been calculated for both an ordered (T ') and a disordered ( T ) crystal.
Figure 4: Simulated HRTEM images for the Y5 (130) symmetric tilt boundary based on (a) Monte Carlo calculation with boron, (b) Monte Carlo calculation without boron, and (c) statically minimized structures. The simulations have been conducted for a defocus of 60 nm and a crystal thickness of 18 nm.
1288
BORON SEGREGATION
Vol.
25,
No. 6
boron increases their cohesive strength. Based on the present Monte Carlo results, we have performed some preliminary calculations to determine whether the total cohesive energy of the boundaries is affected when the influence of both boron and the reduction in local ordering are included. Tlie cohesive energy is defined as the difference between the 8rain boundary and surface euefsies associated with cleaving the boundary. It is found that the cohesive energy in the absence of boron is about 3100 mJ/m 2 while in the presence of boron the cohesive energy increases to about 3550 mJ/m2. Discussion and Conclusions The Ixesence cf absence of a disordered region at grain boundaries is impormm in identifying the mechanism by which boron enhances ductility. The experimental findings presented here clearly demonstrate that the enhanced ductility does not result bom the formation of a wide disordered region at the boundary. The absence of such a region suggests that processes which may enhance plasticity and reduce stress concentrations near the boundary (e. g. eross-sfip) are not likely to be responsible for the observed ducall,Jfion. Nevertheless, if only a narrow region near the grain boundary were completely d i ~ slip mmmisslon as proposedby King and Yoo [2] might still be enhanced since the pomible number and mobility of residual grain boundary dislocations should be increased. The Monte C~in cakulations discussed above indicate that boron causes an increase in the local disorder at the Y-,5 (130) boundary. This result is consistent with TEM observations by Kung and Sess [13] which show that the APB energy for dislocations in low angie twist boundaries is reduced in the presence of bororL However, both these observations and the Monte Carlo calculations demonstrate that a fully disordered state is not achieved. Indeed, a high degree of order is retained for the calculations both with and without boron. Since the addition of boron apparently has such a small influence on the local ordering at grain boundaries, it is unlikely that this effect alone could lead to tbe dramatic chanse in tbe ductility. The Monte Carlo caic~,lm/onsindicate that the ~ddltion of boron causes an ~ of about 15% in the cohesive energy of the Y.5 (130) boundary. However, the calculations do not include the effect of boron on the flow strength of Ni3AI which, in turn, will influence crack tip plasticity. In fact, boron increases the flow strength in this material [7]. Such an increase in strength should actually lead to a reduction in total work of fracture by reducing the plastic zone size associated with a propagating crack. As a consequmu:e, it remains unclear what role the small increase in cohesive energy plays in enlauwing the ductility of boron-doped Ni3AI. These theoretical results seem to indicate that neither the disorder-induced enhancement of dislocation transmission across grain boundaries nor the increase in their cohesive energy alone can explain the dramatic influence of boron on ductility. It is possible that the increase in ductility results from a combination of both effects. The experimental results of this study demonslrate conclusively that large tensile ductilities are obtained in the absence of a wide region of compositional disorder at the interface. Since the calculations indicate that the region of disorder is likely to exist only very near the boundary,more precisely oriented bicrystels are needed. Because the orientation of the present set of bicrystals is not close enough to a pure tilt geometry, it has not been possible to verify the presence of such a narrow region of pmliul disorder. Additionul castings are planned to produce bicrystals that are clos~ to pure tilt geometries and should enable a determirunion of the local ordering at the boundaries.
A~md~dstm~ T. J. Sage and J. P. Scola (SNL) are acknowledged for their assistance in the testing of the ma~'ial. Research at SNL has been sponsored by the DOE, Office of Basic E n m ~ Sciences und~ contract DE-ACO4-76DP00789. R
~
1. K. Aoki and O. Izunfi, N~op~n ginzaku Gakkaishi, 43, 1190 (1979). 2. A.H. King and M. H. Yon, Scripta Met., .21, 1115 (1987). 3. R.A.D. Mackenzie and S.L. Snss, Scr/pta Met., 22, 1807 (1988). 4. M. J. Mills, $ c ~ t a Met., 23, 2061 (1989). 5. J. E. Krzanows~, $cripta Met., 23, 1219 (1989). 6. H. ~ and D. P. pope, J. Mater. Reg.,S, 763 (1990). 7. F. E. Heredia and D. P. pope, MRS Proceedings, 133, 287 (1989). 8. A. Chouc~aury, C. L. White and C. R. Brooks, MRS Proceedings, 122, 261 (1988). 9. P. S t ~ l m a n n , Ultramicroscopy,21, 131 (1987). 10. S. M. Foiles, Phys. Rev. B, 32, 7685 (1985). 11. A. F. Voter and S.P. Chen, MRS Proceedings, 82, 175 (1987). 12. S. P. Cben, A. F. Voter, R. C. Albers, A. M. Boring and P. J. Hay, J. Mater. Res., 5, 955 (1990). 13. H. Kang and S. L. Sass, to be published in the MRS Proceedings, Symposium on High Temperature lntemaciailic Alloys (199D.
Vol. 2S, pp. 1283-1288, 1991 Printed in the U.S.A.
Pergamon
VIEWPOINT
Press plc
SET No. 17
THE INFLUENCE OF BORON SEGREGATION ON THE STRUCTURE AND MECHANICAL PROPERTIES OF BOUNDARIES IN BICRYSTALS OF Ni3Al M. J. Mills, S. H. Goods, S. M. Foiles and J. R. Whetstonet Sandia National Laboratories, Livermore, CA 94551-0969 tAllison Gas Turbines Division, General Motors Corporation, Indianapolis, IN 46206-0420
(Received A p r i l 8, 1991} InQxx]uction In contrast to the high ductilities reported for single crystals of ordered (y) Ni3AI, the ductility in polycrystalline form is extremely limited. The discovery by Aoki and Izumi [1] that microalioying with boron greatly increases the ductility of polycrystallineNi3A1 has initiated spirited scientific investigation into the structure and pmperdes of grain boundaries in this material. One consequence of the addition of boron could be an increase in the cohesive properties of the grain boundaries. However, at present the most widely accepted explanation for the improved ductility of the grain boundaries is that boron causes compositional disordering of the grain boundary region. As first suggested by King and Yon [2], the total number of possible dislocation reactions is generally much larger at a grain boundary in disordered (T) material (the FCC structure) than in the ordered L12 phase. Consequently, for the disordered case, the absorption or transmission of dislocations across the boundary should be easier, large stress concentrations should not develop as readily at the boundary, and the propensity for intergranular fracture should be reduced. Although several high resolution transmission electron microscopy (HRTEM) studies have been performed to date, direct experimental evidence concerning disorder at grain boundaries in horon-doped Ni3AI is presently inconclusive. Mackenzie and Sass [3] have reported the presence of a 1.5 to 2.5 nm thick disordered region at a grain boundary in boron-doped and directionally solidifu:d Ni-24 at.% AI. However, observations by Mills [4] on the same material have failed to reveal such a region. In the latter study, results of image simulations indicated that fine, FCC fringe spacings can be produced at completely ordered grain boundaries if they are slightly inclined to the electron beam, creating the false impression that the boundary region is disordered. These results emphasize the need for well-oriented and carefully controlled boundary geomeiries for HRTEM studies. Finally, Krzanowski [5] performed HRTEM imasing and chemical analysis on chill-cast Ni-24 at.% AI with and without boron. In this nonequilibrium condition, it was found that the boundaries in both alloys were nickel-rich. Nevertheless, superlattice fringe spacings were observed to extend close to the boundary in both the binary and horon-doped material. A fundamental shortcoming of these previous investigations is that the observed boundaries are not necessarily representative of all the grain boundaries in the sample, and may not be the ones responsible for determining the macroscopic behavior. This is particularly true of HRTEM studies in which only limited grain boundary areas are examined. In order to directly link grain boundary structure with macroscopic properties, it is necessary to isolate a single boundary for study. To enable this direct correlation of microscopic structure with ductility, bicrystais have been used in the present investigation. The use of bicrystals with suitable crystallographic orientations is also advantageous with respect to the HRTEM observations since (a) lardce fringes may be seen in the crystals on either side of the boundary and Co) the boundary is assured to be close to an "edge-on" configuration. These factors should limit the possible imaging artifacts which can arise in the HRTEM studies and should also allow the degree of disordering to he evaluated very close to ~ e boundary itself. In this investigation, boron-doped bicrystals of Ni3AI (24 at.% AI + 1 at.% Ta) have been studied using HRTEM, while companion specimens were mechanically deformed in tension. The results of these experimental studies will also be discussed in light of Monte Carlo atomistic calculations which have been performed to evaluate the effect of homn on grain boundary structure and ordering. Ex~*rimental Procedu~ The bicrystal used in this study was produced via a modified single crystal investment casting process. Single crystal seeds were oriented in mold fixtures with a common [001] parallel to the growth direction. While the two seeds were oriented in order to obtain the desired ~,5 (120) tilt botmdary geometry, the resulting bicrystal was found to have an additional tilt misorientafion of about 7 ° about the [100] in one of the crystals. The bulk composition of the sample is 23 at% AI, 1 at% Ta, 0.27 at% B, and the balance Ni. The biorystal is therefore slightly Ni-rich relative to the stoichiometric composition. The Ta is substitutional on AI sites [6] and was ~dd~ to assist the growth
1283 0036-9748/91 $3.00 + .00
1284
BORON
SEGREGATION
Vol.
25, No.
of the crystals. Although the ternary addition of Ta causes an increase in the flow stress at low temperatures [7], it does not strongly affect the f~-ture behavior of polycrystalline material. Tensile specimens with reduced gage sections were spark cut from the as-grown bicrystaL In order to maximize the tensile stresses across the interface, the specimens were cut so that the tensile axis was normal to the nominal orientation of the grain boundary plane. Prior to mechanical testing, one specimen was vacuum annealed (< 1 x 10-4 Pa) at 1000°C for 48 hours, then slow cooled at 50°C/hour. In previous work, Auger electron spectroscopy has shown that this beat treatment maximizes the boron concentration at the grain boundary [8]. A second specimen was sealed in an evacuated quartz ampule and annealed at the same temperature for the same length of time. However, rather than being slow cooled, it was rapidly quenched in a water bath. This heat treatment should freeze in a homogeneous distribution of boron throughout the specimen with minimum segregation to the boundary. A final specimen was tested in the as-grown condition which is expected to have an intermediate cooling rate and boron segregation. All three specimens were tensile tested at room temperature at a constant extension rate of = 4 x 10-3 mm/sec. Foils for Iransmission electron microscopy were prepared from slices cut normal to the growth direction of the bicrystai and anneaied in an identical manner to the tensile test specimens. Observations were conducted using a JEOL 4000EX operating at 400 kV. Image simulations have been performed using the software package of Stadelmann [9].
Results
MechanicalTesting The tensile curves for the three different specimens are shown in Figure I. It is clear that the different thermal histories designed to vary the boron concentrations at the interface had a significanteffect on ductility. The annealed and slow-cooled (SC) specimen exhibited the largest strain to failure, 45%, while the as-cast (AC) specimen failed at a total swain of 28%. In contrast, the annealed and quenched (Q) specimen failed after a total strain of only 2.5%. It should be further noted that at low strains, the curves for the AC- and the SC-specimens superimpose, indicating that the anneal had no effect on the flow characteristics for these two conditions. In concert with these changes in ductility, the different heat treatments also affected the fracture morphologies of the samples. In Figure 2, the specimen marked as "a" is undeformed. The interface is located near the center of the gauge section, and approximately normal to the tensile axis. For the high ductility SC-specimen (marked as "c" in Figure 2), failure actually occurred in the grip region, away from the grain boundary-presumably at a strain concentration arising from the specimen geometry. For the AC-specimen, marked as "b", fracture occurred at the interface. This specimen clearly exhibited far less axial strain prior to failure. The complex path of the fracture surface reveals that the interface was not planar. Rather, it varied considerably from the intended (120) boundary plane. The Q-specimen is not shown. However, it failed in a manner identical to that for the AC-specimen, albeit with even less axial extension. These results are consistent with those reported by Cboudhury, et al. [8] for polycrystalline Ni-24 at% Al doped with boron. In their work, relatively small increases in the grain boundary boron concentration were apparently responsible for the transition from brittle, intergranular fracture to ductile failure. The behavior of the Q-specimen also indicates that, in the absence of boron, this bicrystal boundary would be intrinsically brittle and was not a "special" boundary across which slip was favored. From the observed distortions, it is evident that the amount of strain is not the same in both crystals. In fact, the reduction in area in the upper crystal was precisely half that of the lower crystal. This difference in strain is due to the fact that the biorystal is misoriented from the ideal Y.,5(120) geometry, and therefore the tensile axis is not of the same.type in the two crystals. The tensile direction in the lower crystal is close to the expected [120], so that both the a[011] ( I I I ) and a[011](ll]) slip systems _havelarge resolved shear stresses. The upper crystal is rotated about 7° from the [120], so that only the a[011](III) system is active, as confirmed by slip trace analysis. The large and abrupt slrain discontinuity apparent in Figure 2d indicates that slip transfer across the grain boundary was difficult. This observation further emphasizes the fact that the boundary was not in a uniquely favorable orientation for dislocation transmission. In fact, even for an ideal Y~5(120) boundary under a tensile load, the mmsmission of dislocations from one crystal to the other is complicated since none of the active dislocations (or their partials) can be readily decomposed into grain boundary dislocations, based on the considerations of King and Yoo [2]. In spite of the large stress concentrations that must have developed at the boundary during the deformation of the annealed sample, die boundary remained intact throughout the test. Using HRTEM, the validity of correlating this high ductility with the possible presence of a compositionally disordered region at the grain boundary will now be evaluated.
6
Vol. 25, No. 6
BORON SEGREGATION
800
"
|
!
1285
t
i
Ni3(AI, Ta) + 0.27% B Bicrystal A m
a.
600 J
v
P
400 (AC)
Q
2
I-
SlowCooled (so)
200
~ 0 u.O
nnealed,~ e n e l ' ~ •
(q)
0.1
,
,
0.2 True
,
0.3 Strain
,
,
,
0.4
0.I
Figure 1: Flow curves for tensile tests of Ni3(A1, 1 at% Ta) + 0.27 at% B bicrystals for three heat treatment conditions.
a
b
c
d
Figure 2: Mechanical test specimens (a) before tesdng, (b) LCtcrtesting in the as-cast condition, and (c) after testing in the annealed and slow cooled condition. In (d) a slightly magnified side view of the annealed and slow cooled sample is shown, and the location of the grain boundary (GB) is indicated. Note the difference in specimen thickness indicated by arrows in (d).
1286
BORON SEGREGATION
Vol.
25, No. 6
HRTEM Observations An HRTEM image of the bicrystal boundary in the SC condition is shown in Figure 3. The image reflects the fact the bicrystal is not in the ideal Y5 (120) geometry. As a result, only [100] lattice fringes are observed in the lower crystal, while [100] and [010] cross-fringes can be seen in the upper crystal. The inset simulated images have been calculated for both the ordered and the disordered structures, and for the estimated orientation of both crystals. Comparison of these simulations with the observed image shows that the slructaxe of the bulk crystals are consistent with ordered ? Ni3A1. Although the deviation of this boundary from the pure tilt geometry precluded a determination of the atomic structure of the boundary, local variations in the orientation of the interface plane can be identified. Since both crystals exhibited at least one set of lattice fringes, any overlap between the crystals due to an inclination of the boundary from the "edge-on" orientation would result in the appearance of Mdir~ fi'inges. These interference fringes can be seen in the image by sighting along the direction indicated by the arrow in Figure 3. For the evaluation of local compositional disorder, it is extremely important that the grain boundary be as close to the "edge-on" orientation as possible [4]. This image was obtained after orienting the bicrystal in the microscope so that the extent of the Moir~ fringes was minimized while remaining as close as possible to the [001] zone axis in both crystals. The significant result of these HRTEM observations is that the ordering present in the bulk crystals extended up to within a distance of 0.5 nm from the boundary with no detectable loss of ordering. These results demonstrate that the high ductility exhibited by this hicrystal in the SC condition was not due to the presence of a distinct disordered region at the boundary. In fact, evidence for such a disordered region was not observed in either the AC or the SC condition. While the dramatic reduction in the ductility of the specimen in the Q condition suggests that this bicrystal boundary was inherently brittle and that the segregation of boron was responsible for the increased ductility, it is clear that this ducfilization was not brought about by the formation of a broad disordered region at the boundary. Atomistic Calculations of Boron Effects on Grain Boundary Disordering, Structure and Energy If disorder is present at these boundaries, the HRTEM results indicate that it must be localized to the atomic planes immediately adjacent w the boundary. Monte Carlo atomisfic calculations [I0] have been performed in order to provide an estimate of the extent of boron-induced disorder. A ]:5 (130) symmetric tilt boundary has been considered, with a bulk composition of 76 at.% Ni and two boron atoms per boundary period for the boron-doped case. While this is not precisely the same boundary geometry as that studied experimentally, the results are expected to be qualitatively similar. The calculations utilized the Embedded Atom Method (EAM) potentials for the nickel, aluminum and boron system developed by Voter and Chen [11]. These calculations allow for the compositional rearrangement of the system as well as the spatial displacement of the atoms. They were performed for both binary and boron-doped boundaries and at temperatures of 227 and 727oC. In the simulations, the boron is not allowed to diffuse from the boundary, even though it might do so in the real system at the higher temperatures considered. For calculations both with and without boron, a localized region of reduced ordering is predicted to occur at the boundary. In all cases, a partially disordered region, characterized by an increased concenuation of anti-site defects, extends 2 to 3 (130) atomic planes on either side of the boundary. This results from the reduction in anti-site defect energies for many of the sites near the boundary. The amount of disorder produced in the absence of boron was relatively small; about 8 % of the Ni sites and 30% of the AI sites within 0.5 nm of the boundary have anti-site defects in the simulations at 727°C. (In the bulk at this temperature, 1% of the Ni sites and 7.5% of the Al sites contain anti-site defects.) The presence of boron increases the amount of disorder seen; about 15% of the Ni sites and 34% of the AI sites near the boundary have anti-site defects. The increase in disorder caused by the boron is due to a competition between two factors. First, the atoms want to arrange so as to maximize the number of Ni-A1bonds and minimize the number of AI-A1 bonds since the former are stronger. This is the driving force for the creation of the ordered crystal lattice. On the other hand, it is energetically preferable to place Ni atoms next to the B atoms instead of AI atoms. Since it is not possible to achieve both of these conditions simultaneously, it is easier to induce local disorder in the presence of the boron. The degree of ordering at the boundary is reflected in the image simulations for both swactures shown in Figure 4a and 4b. The strong intensities in the images are due to the ordered superlattice. These intensities are clearly reduced near the boundary for the boron-doped case, while they axe stronger at the boron-free boundary. The effect of partial disordering on these images can be conWastedwith the image simulation in Figure 4c. This simulation is based on a structure which has been statically relaxed so that only spatial displacements of the atoms are allowed. Hence, the region near the boundary remains fully ordered and no reduction in the superlattice intensifies is observed at the boundary. These simulations suggest that it should be possible using HRTEM to distinguish between a fully ordered grain boundary and one in which there is only a narrow region of partial disorder. The addition of boron could also cause an increase in the cohesive strength of the boundaries. Chen, et al. [12] have used the EAM to test this possibility and found that for completely ordered boundaries the addition of
Vol.
25, No. 6
BORON S E G R E G A T I O N
1287
Figure 3: HRTEM micrograph of the bierystal boundary obtained for an objective lens defocus of about 60 nm and a crystal thickness of about 18 nm. The upper crystal is oriented close to the [001] zone axis while the lower crystal is near the [0 1 10] zone axis. Inset image simulations have been calculated for both an ordered (T ') and a disordered ( T ) crystal.
Figure 4: Simulated HRTEM images for the Y5 (130) symmetric tilt boundary based on (a) Monte Carlo calculation with boron, (b) Monte Carlo calculation without boron, and (c) statically minimized structures. The simulations have been conducted for a defocus of 60 nm and a crystal thickness of 18 nm.
1288
BORON SEGREGATION
Vol.
25,
No. 6
boron increases their cohesive strength. Based on the present Monte Carlo results, we have performed some preliminary calculations to determine whether the total cohesive energy of the boundaries is affected when the influence of both boron and the reduction in local ordering are included. Tlie cohesive energy is defined as the difference between the 8rain boundary and surface euefsies associated with cleaving the boundary. It is found that the cohesive energy in the absence of boron is about 3100 mJ/m 2 while in the presence of boron the cohesive energy increases to about 3550 mJ/m2. Discussion and Conclusions The Ixesence cf absence of a disordered region at grain boundaries is impormm in identifying the mechanism by which boron enhances ductility. The experimental findings presented here clearly demonstrate that the enhanced ductility does not result bom the formation of a wide disordered region at the boundary. The absence of such a region suggests that processes which may enhance plasticity and reduce stress concentrations near the boundary (e. g. eross-sfip) are not likely to be responsible for the observed ducall,Jfion. Nevertheless, if only a narrow region near the grain boundary were completely d i ~ slip mmmisslon as proposedby King and Yoo [2] might still be enhanced since the pomible number and mobility of residual grain boundary dislocations should be increased. The Monte C~in cakulations discussed above indicate that boron causes an increase in the local disorder at the Y-,5 (130) boundary. This result is consistent with TEM observations by Kung and Sess [13] which show that the APB energy for dislocations in low angie twist boundaries is reduced in the presence of bororL However, both these observations and the Monte Carlo calculations demonstrate that a fully disordered state is not achieved. Indeed, a high degree of order is retained for the calculations both with and without boron. Since the addition of boron apparently has such a small influence on the local ordering at grain boundaries, it is unlikely that this effect alone could lead to tbe dramatic chanse in tbe ductility. The Monte Carlo caic~,lm/onsindicate that the ~ddltion of boron causes an ~ of about 15% in the cohesive energy of the Y.5 (130) boundary. However, the calculations do not include the effect of boron on the flow strength of Ni3AI which, in turn, will influence crack tip plasticity. In fact, boron increases the flow strength in this material [7]. Such an increase in strength should actually lead to a reduction in total work of fracture by reducing the plastic zone size associated with a propagating crack. As a consequmu:e, it remains unclear what role the small increase in cohesive energy plays in enlauwing the ductility of boron-doped Ni3AI. These theoretical results seem to indicate that neither the disorder-induced enhancement of dislocation transmission across grain boundaries nor the increase in their cohesive energy alone can explain the dramatic influence of boron on ductility. It is possible that the increase in ductility results from a combination of both effects. The experimental results of this study demonslrate conclusively that large tensile ductilities are obtained in the absence of a wide region of compositional disorder at the interface. Since the calculations indicate that the region of disorder is likely to exist only very near the boundary,more precisely oriented bicrystels are needed. Because the orientation of the present set of bicrystals is not close enough to a pure tilt geometry, it has not been possible to verify the presence of such a narrow region of pmliul disorder. Additionul castings are planned to produce bicrystals that are clos~ to pure tilt geometries and should enable a determirunion of the local ordering at the boundaries.
A~md~dstm~ T. J. Sage and J. P. Scola (SNL) are acknowledged for their assistance in the testing of the ma~'ial. Research at SNL has been sponsored by the DOE, Office of Basic E n m ~ Sciences und~ contract DE-ACO4-76DP00789. R
~
1. K. Aoki and O. Izunfi, N~op~n ginzaku Gakkaishi, 43, 1190 (1979). 2. A.H. King and M. H. Yon, Scripta Met., .21, 1115 (1987). 3. R.A.D. Mackenzie and S.L. Snss, Scr/pta Met., 22, 1807 (1988). 4. M. J. Mills, $ c ~ t a Met., 23, 2061 (1989). 5. J. E. Krzanows~, $cripta Met., 23, 1219 (1989). 6. H. ~ and D. P. pope, J. Mater. Reg.,S, 763 (1990). 7. F. E. Heredia and D. P. pope, MRS Proceedings, 133, 287 (1989). 8. A. Chouc~aury, C. L. White and C. R. Brooks, MRS Proceedings, 122, 261 (1988). 9. P. S t ~ l m a n n , Ultramicroscopy,21, 131 (1987). 10. S. M. Foiles, Phys. Rev. B, 32, 7685 (1985). 11. A. F. Voter and S.P. Chen, MRS Proceedings, 82, 175 (1987). 12. S. P. Cben, A. F. Voter, R. C. Albers, A. M. Boring and P. J. Hay, J. Mater. Res., 5, 955 (1990). 13. H. Kang and S. L. Sass, to be published in the MRS Proceedings, Symposium on High Temperature lntemaciailic Alloys (199D.