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Compressive flow stress of a binary stoichiometric Ni3Al single crystal
ScriptaMaterialia, Vol. 37, No. 11, pp. 1777-1782,1997 Elsevier Science Ltd Copyright 0 1997 Acta Metahrgica Inc. Printed in the.USA. All rights reserved 1359~6462197$17.00 + .OO
PI1 S1359-6462(97)00339-4
COMPRESSIVE FLOW STRESS OF A BINARY STOICHIOMETRIC N&Al SINGLE CRYSTAL D. Golberg*, M. Demura, and T. Hirano Nat ional Research Institute for Metals, l-2-1 Sengen , Tsukuba, Ibaraki 305, Japan *Present address: National Institute for Research in Inorganic Materials, l- 1 Namiki, Tsukuba, Ibaraki 305, Japan (Received June 17, 1997) (Accepted August 5, 1997) Introduction N4Al exhibits the yield stress anomaly, i.e. its flow stress increases with increasing temperature (1,2). The yield stress anomaly has been extensively studied using single crystals (3-8). These studies have been carried out on Ni-rich (3) or ternary Ni3Al (4-8), whereas the mechanical properties of a binary stoichiometric Ni3Al single crystal that do not contain ternary additions have never been evaluated. Since deviations from stoichiometry or ternary elements essentially produce point defects, the nature of the yield stress anomaly can be affected by these. Therefore it is worth studying N&Al deformation mechanisms using a binary stoichiometric single crystal. However, there is a difftculty in growth of these crystals due to the peritectic reaction liquid + NiAl + NisAl (l), by which the intermetallic forms in a binary alloy. The peritectic solidification process usually leads to copious grain nucleation. Thus all the single crystals grown to date have contained ternary additions or have deviated towards the Ni-rich side of stoichiometry in order to avoid the peritectic reaction (9). Very recently we have found that large binary stoichiometric Ni3Al single crystals can occasionally be grown using a floating zone (FZ) method (lo- 12). The present paper reports the orientation dependence of the compressive flow stress of such crystals at 293-1273 K for testing axis between [OOl] and [Oll]. Experimental Procedure A binary stoichiometric Ni-25 at.%Al alloy was arc-melted using 99.9999%Al and 99.97% Ni. The seed and feed ‘crystals were prepared from the same arc-melted rod. The procedure for crystal growth has been previously reported by us (lo- 12). Single crystals of length 160- 170 mm and diameter IO- 12 mm were grown. Al-content of the as-grown crystal was measured by using wet chemical method and was found to be slightly Al-rich, i.e., 25.2 at.%Al, with respect to stoichiometry.
1777
1778
FLOW STRESS OF NiJAI
Vol. 37, No. 11
Optical microscope observations were carried out on the transverse and longitudinal crystal sections. Marble reagent (2.5 g CuSO4 + 25 cm3 Hcl + 30 cm3 H20) was used for etching. Crystal orientation was determined by the Laue X-ray back reflection method. Compression specimens 3.7 x 3.7 x 7.7 mm with a long dimension oriented between [OOI] and [01 1] were cut from the as-grown crystals by electrical discharge machining. The specimen surface was mechanically polished on 600 grit Sic paper until an approximately 100 pm thick layer had been removed, followed by final AlzOs lapping for a “mirror” finish. Compression tests were carried out at 293-1273 K in air at an initial strain rate of 2.3 x 104/s with a Tensilon 5000 testing machine. Specimens were heated to the testing temperature in a split furnace within 2-10 min and kept at this temperature for 5-10 min before loading to ensure uniform temperature distribution. Plastic strain was calibrated from the change in specimen height after the compression test by assuming a constant strain rate. Results
Figure la shows the longitudinal microstructure of the as-grown single crystal. Figure lb shows the Laue X-ray back reflection pattern of the crystal. The X-ray spots are slightly split, implying the existence of small angle boundaries, but the misorientation angle was measured as less than lo. The same patterns were obtained over the entire transverse section.
Figure 1. a) Optical micrograph of the longitudinal section of the binary stoichiometric N&Al single crystal; and b) corresponding Laue X-ray back reflection pattern.
Vol. 37, No. 11
FLOW STRESSOF NixAl
1779
co111
[OOII
Figure 2. Orientation of the compression axes.
The compression axes studied are plotted on the standard [OOl]-[ i 1l]-[Ol l] unit triangle in Fig. 2. Calculated S&aid factors and Schmid factor ratio N are listed in Table 1. Hereafter the crystal orientations are referred to orientation A (close to [OOl], orientation B (between [OOl] and [Ol l]), and orientation C (close to [Ol 11). Figure 3 shows the stress-strain curves of the crystal for orientations A and C. All the orientations exhibit positive temperature dependence of the flow stress, that is, flow stress increases with increasing temperature below the peak temperature, 1073 K. Above 1073 K, the yield stress decreases with increasing temperature. The stress-strain curves show a smooth yielding followed by work hardening below 1073 K. At 1073 K and above they show a yield drop followed by almost zero or negative workhardening. This yield drop is thought to be due to rapid dislocation multiplication, i.e. JonhstonGilman type yielding (13). The extent of the yield drop is larger for orientation A compared to orientation C. which is consistent with the trend observed in Nis(A1,Ti) single crystals (13). 1ccc 1
Orientation A
1000
1
Orientation C
600-
.p* 600I
2 600i 1400-
I UJ 4001273 K
200 -
0
2
4
6
8
Plastic Strain, %
10
0
2 4 6 8 Plastic Strain, %
10
Figure 3. Stress-strain curves for orientations A and C as a function of temperature.
FLOW STRESSOF N&AI
1780
I
400 -
I
I
I
I
I
I
Vol. 37, No. 11
I
A -A 0 -0 0 -c
m B 2 300 #E E E 200 g ?I 8 1000 200
I 400
600 600 1000 Temperature, K
Figure 4. Critical resolved shear stress (CRSS) on (11 I)[ iOl]
I 120Q
1400
as a function of temperature.
There is an agreement in the literature that below the peak temperature plastic flow in Ni3Al takes place on the octahedral slip system {111 } (1, 2, 9). Above the peak temperature, the cube slip system (OOl}
A
0.46
:
0.49 0.44
ON; S&mid Factor ratio ior nmalip
(OlO)[iOl]
(OOl)[ilO]
N’
0.08
0.08
0.17
0.35 0.29
0.35 0.28
0.80 0.59
system (OlO)[iOl] ml odahedml slip sysmn (lll)[iOl]
FLOW STRESS OF N&Al
Vol. 37, No. 1I
I
I
I
1781
573 K
I
NB(AI,Ta)
o stoichiometric NBAI, present work A Ni-rich Ni3AI (3)
NiS(AI.Nb) [5]
200 I ; ;; IF f
(6)
i
Ni3(AI.W) (4)
150
c ; 8
’ O”
100
vometric
Ni3A present work
t
0.0
4
0.0
0.2
0.4
0.6
0.8
1.0
W
0.2
0.4
0.6
0.8
1.0
N, (0 io)[ioi]/(iii)[io
1.2
11
The CR:SS on (11 l)[ 7011 as a function of Schmid factor ratio N for the binary stoichiometric N&Al single crystal compared to: a) Ni-rich single crystal [3] at 573 K and 773 K; and b) ternary N&Al single crystals [4-61 at 573 K.
Figure 5.
alloys exhibit a strong positive N-dependence (4-6). The results are that between [OOl] and [Ol l] S&mid’s law is almost obeyed in binary stoichiometric NijAl, while it is violated in off-stoichiometric and ternary N&Al. This means that the orientation dependence of the CRSS originates in point defects such as anti-site defects arising from off-stoichiometric compositions or ternary elements. In other words, Schmid’s law generally holds in stoichiometric and ordered N&Al. However, there is still an ambiguity regarding the general validity of S&mid’s law in our alloy since our data are lacking near [i 1I] orientation where the previously observed N-dependence is most pronounced. Violation of S&mid’s law is one of the important characteristic features of the stress anomaly in L 1z-type intermetallic compounds (1,2). Takeuchi and Kuramoto (14) proposed their model to explain the stress anomaly based on this feature in Ni,Ga: the CRSS for (11 l)[ iOl] slip is dependent on the orientation. Paidar, Pope, and Vitek (15) also proposed their model based on more detailed facts in Ni,(Al,Nb) (7). In these models the yield stress is determined by thermally activated cross-slip of (11 l)[ i0 I] dislocations onto (100) cube planes and the cube cross-slip is aided by the CRSS on (100) leading to the violation of Schmid’s law. However, the present results show that violation of Schmid’s law is characteristic of non-stoichiometric and non-binary NbAl, and is not intrinsic to the compound itself. It is, therefore, necessary to reconsider the deformation model in N&Al by taking the effect of anti-site defects or ternary elements into consideration. Although the reason for the orientation dependence of the CRSS on (11 l)[ iOl] in nonstoichiometric or ternary NbAl is not yet understood, it can be related to the extent of hardening by ternary elements or anti-site defects as follows. The ternary elements Ta, Nb, and W, have a strong hardening effect in Ni3Al (16). The ternary NiJAl containing these elements exhibits a strong N-dependence (4-6), as shown in Fig. 5b. In contrast, Ti has a very weak hardening effect (16) and ternary Ni3Al containing Ti has no N-dependence at room temperature (8). In the case of Ni-rich NisAl, Ni anti-site defects have a weak hardening effect (17) and their N-dependence is not strong, as shown in Fig. 5a. Then, we suppose that interaction between the point defects and mobile dislocations
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FLOW STRESS OF Ni3AI
Vol. 37, No. 11
causes the orientation dependence of the CRSS. Without this interaction, no orientation dependence is observed. Conclusions For the first time compressive flow stress of binary stoichiometric single-crystalline NijAl has been studied for orientations between [OOl] and [Ol l]. All orientations exhibit anomalous positive temperature dependence of the flow stress with a weakly orientation-dependent peak temperature between 973- 1073 K. The CRSS on (11 l)[ 70 l] had almost no orientation dependence in the temperature range of the yield stress anomaly. This is considered to be the nature of binary stoichiometric N&Al. Previously reported orientation-dependent CRSS in Ni-rich and ternary NisAl is thought to be due to the point defects induced by off-stoichiometry and ternary elements. Acknowledgments This research has been carried out under the Japanese Science and Technology Agency (STA) Fellowship Program. D. Golberg expresses his gratitude for the award of an STA post-doctoral Fellowship, tenable at the National Research Institute for Metals, Tsukuba, Japan. The authors wish to thank E.P. George at Oak Ridge National Laboratory for many helpful discussions. References 1. 2. 3. 4. 5. 6. 7. 8. 9. 10. 11. 12. 13. 14. 15. 16. 17.
D.P. Pope and S.S. Ezz, Int. Mat. Rev., 29, 136 (1984). N.S. Stoloff, Int. Mat. Rev. 34, 153 (1989). F.E. Heredia and D.P. Pope, Acta Metall., 39,2027 (1991). E. Kuramotc and D.P. Pope, Acta Metall., 26,207 (1978). S.S. Ezz, D.P. Pope and V. Paidar, ActaMetall., 30,921(1982). Y. Umakoshi, D.P. Pope and V. Paidar, Acta Metall., 32,449 (1984). C. Lall, S. Chin and D.P. Pope, Metall. Trans. A, lOA, 1323 (1979). M.S. Kim, S. Hanada, S. Watanabe and 0. Izumi, ActaMetall., 36,2615 (1988). T. Suzuki, Y. Mishima and S. Miura, ISIJ International, 29,1 (1989). M. Demura and T. Hirano, Phil. Msg. Letters, 75, 143 (1997). M. Demura and T. Hirano, in Mat.Res. Sot. Proc. (Boston, 2-6 Dec., 1996, USA) in press. D. Golberg, M. Demura and T. Hirano, in Proc. 2nd Inter. Symp. on Structural Intermetallics, (Sept. 1997, Seven Springs, Pennsylvania, USA) in press. S. Ochiai, S. Miura, Y. Miihima and T. Suzuki, J. Jpn. Inst. of Metals, 51,608 (1987). S. Takeuchi and E. Kuramoto, Acta metall., 21,415 (1973). V. Paidar, D.P. Pope and V. Vitek, Acta MetaIl., 32,449 (1984). Y. Mishiia, S. Ochiai, N. Hamao, M. Yodogawa and T. Suzuki, Trans. Jpn. Inst. Metals, 27,648 (1986). 0. Noguchi, Y. Gya and T. Suzuki, Metal. Trans., 12A, 1647 (1981).
PI1 S1359-6462(97)00339-4
COMPRESSIVE FLOW STRESS OF A BINARY STOICHIOMETRIC N&Al SINGLE CRYSTAL D. Golberg*, M. Demura, and T. Hirano Nat ional Research Institute for Metals, l-2-1 Sengen , Tsukuba, Ibaraki 305, Japan *Present address: National Institute for Research in Inorganic Materials, l- 1 Namiki, Tsukuba, Ibaraki 305, Japan (Received June 17, 1997) (Accepted August 5, 1997) Introduction N4Al exhibits the yield stress anomaly, i.e. its flow stress increases with increasing temperature (1,2). The yield stress anomaly has been extensively studied using single crystals (3-8). These studies have been carried out on Ni-rich (3) or ternary Ni3Al (4-8), whereas the mechanical properties of a binary stoichiometric Ni3Al single crystal that do not contain ternary additions have never been evaluated. Since deviations from stoichiometry or ternary elements essentially produce point defects, the nature of the yield stress anomaly can be affected by these. Therefore it is worth studying N&Al deformation mechanisms using a binary stoichiometric single crystal. However, there is a difftculty in growth of these crystals due to the peritectic reaction liquid + NiAl + NisAl (l), by which the intermetallic forms in a binary alloy. The peritectic solidification process usually leads to copious grain nucleation. Thus all the single crystals grown to date have contained ternary additions or have deviated towards the Ni-rich side of stoichiometry in order to avoid the peritectic reaction (9). Very recently we have found that large binary stoichiometric Ni3Al single crystals can occasionally be grown using a floating zone (FZ) method (lo- 12). The present paper reports the orientation dependence of the compressive flow stress of such crystals at 293-1273 K for testing axis between [OOl] and [Oll]. Experimental Procedure A binary stoichiometric Ni-25 at.%Al alloy was arc-melted using 99.9999%Al and 99.97% Ni. The seed and feed ‘crystals were prepared from the same arc-melted rod. The procedure for crystal growth has been previously reported by us (lo- 12). Single crystals of length 160- 170 mm and diameter IO- 12 mm were grown. Al-content of the as-grown crystal was measured by using wet chemical method and was found to be slightly Al-rich, i.e., 25.2 at.%Al, with respect to stoichiometry.
1777
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FLOW STRESS OF NiJAI
Vol. 37, No. 11
Optical microscope observations were carried out on the transverse and longitudinal crystal sections. Marble reagent (2.5 g CuSO4 + 25 cm3 Hcl + 30 cm3 H20) was used for etching. Crystal orientation was determined by the Laue X-ray back reflection method. Compression specimens 3.7 x 3.7 x 7.7 mm with a long dimension oriented between [OOI] and [01 1] were cut from the as-grown crystals by electrical discharge machining. The specimen surface was mechanically polished on 600 grit Sic paper until an approximately 100 pm thick layer had been removed, followed by final AlzOs lapping for a “mirror” finish. Compression tests were carried out at 293-1273 K in air at an initial strain rate of 2.3 x 104/s with a Tensilon 5000 testing machine. Specimens were heated to the testing temperature in a split furnace within 2-10 min and kept at this temperature for 5-10 min before loading to ensure uniform temperature distribution. Plastic strain was calibrated from the change in specimen height after the compression test by assuming a constant strain rate. Results
Figure la shows the longitudinal microstructure of the as-grown single crystal. Figure lb shows the Laue X-ray back reflection pattern of the crystal. The X-ray spots are slightly split, implying the existence of small angle boundaries, but the misorientation angle was measured as less than lo. The same patterns were obtained over the entire transverse section.
Figure 1. a) Optical micrograph of the longitudinal section of the binary stoichiometric N&Al single crystal; and b) corresponding Laue X-ray back reflection pattern.
Vol. 37, No. 11
FLOW STRESSOF NixAl
1779
co111
[OOII
Figure 2. Orientation of the compression axes.
The compression axes studied are plotted on the standard [OOl]-[ i 1l]-[Ol l] unit triangle in Fig. 2. Calculated S&aid factors and Schmid factor ratio N are listed in Table 1. Hereafter the crystal orientations are referred to orientation A (close to [OOl], orientation B (between [OOl] and [Ol l]), and orientation C (close to [Ol 11). Figure 3 shows the stress-strain curves of the crystal for orientations A and C. All the orientations exhibit positive temperature dependence of the flow stress, that is, flow stress increases with increasing temperature below the peak temperature, 1073 K. Above 1073 K, the yield stress decreases with increasing temperature. The stress-strain curves show a smooth yielding followed by work hardening below 1073 K. At 1073 K and above they show a yield drop followed by almost zero or negative workhardening. This yield drop is thought to be due to rapid dislocation multiplication, i.e. JonhstonGilman type yielding (13). The extent of the yield drop is larger for orientation A compared to orientation C. which is consistent with the trend observed in Nis(A1,Ti) single crystals (13). 1ccc 1
Orientation A
1000
1
Orientation C
600-
.p* 600I
2 600i 1400-
I UJ 4001273 K
200 -
0
2
4
6
8
Plastic Strain, %
10
0
2 4 6 8 Plastic Strain, %
10
Figure 3. Stress-strain curves for orientations A and C as a function of temperature.
FLOW STRESSOF N&AI
1780
I
400 -
I
I
I
I
I
I
Vol. 37, No. 11
I
A -A 0 -0 0 -c
m B 2 300 #E E E 200 g ?I 8 1000 200
I 400
600 600 1000 Temperature, K
Figure 4. Critical resolved shear stress (CRSS) on (11 I)[ iOl]
I 120Q
1400
as a function of temperature.
There is an agreement in the literature that below the peak temperature plastic flow in Ni3Al takes place on the octahedral slip system {111 }
A
0.46
:
0.49 0.44
ON; S&mid Factor ratio ior nmalip
(OlO)[iOl]
(OOl)[ilO]
N’
0.08
0.08
0.17
0.35 0.29
0.35 0.28
0.80 0.59
system (OlO)[iOl] ml odahedml slip sysmn (lll)[iOl]
FLOW STRESS OF N&Al
Vol. 37, No. 1I
I
I
I
1781
573 K
I
NB(AI,Ta)
o stoichiometric NBAI, present work A Ni-rich Ni3AI (3)
NiS(AI.Nb) [5]
200 I ; ;; IF f
(6)
i
Ni3(AI.W) (4)
150
c ; 8
’ O”
100
vometric
Ni3A present work
t
0.0
4
0.0
0.2
0.4
0.6
0.8
1.0
W
0.2
0.4
0.6
0.8
1.0
N, (0 io)[ioi]/(iii)[io
1.2
11
The CR:SS on (11 l)[ 7011 as a function of Schmid factor ratio N for the binary stoichiometric N&Al single crystal compared to: a) Ni-rich single crystal [3] at 573 K and 773 K; and b) ternary N&Al single crystals [4-61 at 573 K.
Figure 5.
alloys exhibit a strong positive N-dependence (4-6). The results are that between [OOl] and [Ol l] S&mid’s law is almost obeyed in binary stoichiometric NijAl, while it is violated in off-stoichiometric and ternary N&Al. This means that the orientation dependence of the CRSS originates in point defects such as anti-site defects arising from off-stoichiometric compositions or ternary elements. In other words, Schmid’s law generally holds in stoichiometric and ordered N&Al. However, there is still an ambiguity regarding the general validity of S&mid’s law in our alloy since our data are lacking near [i 1I] orientation where the previously observed N-dependence is most pronounced. Violation of S&mid’s law is one of the important characteristic features of the stress anomaly in L 1z-type intermetallic compounds (1,2). Takeuchi and Kuramoto (14) proposed their model to explain the stress anomaly based on this feature in Ni,Ga: the CRSS for (11 l)[ iOl] slip is dependent on the orientation. Paidar, Pope, and Vitek (15) also proposed their model based on more detailed facts in Ni,(Al,Nb) (7). In these models the yield stress is determined by thermally activated cross-slip of (11 l)[ i0 I] dislocations onto (100) cube planes and the cube cross-slip is aided by the CRSS on (100) leading to the violation of Schmid’s law. However, the present results show that violation of Schmid’s law is characteristic of non-stoichiometric and non-binary NbAl, and is not intrinsic to the compound itself. It is, therefore, necessary to reconsider the deformation model in N&Al by taking the effect of anti-site defects or ternary elements into consideration. Although the reason for the orientation dependence of the CRSS on (11 l)[ iOl] in nonstoichiometric or ternary NbAl is not yet understood, it can be related to the extent of hardening by ternary elements or anti-site defects as follows. The ternary elements Ta, Nb, and W, have a strong hardening effect in Ni3Al (16). The ternary NiJAl containing these elements exhibits a strong N-dependence (4-6), as shown in Fig. 5b. In contrast, Ti has a very weak hardening effect (16) and ternary Ni3Al containing Ti has no N-dependence at room temperature (8). In the case of Ni-rich NisAl, Ni anti-site defects have a weak hardening effect (17) and their N-dependence is not strong, as shown in Fig. 5a. Then, we suppose that interaction between the point defects and mobile dislocations
1782
FLOW STRESS OF Ni3AI
Vol. 37, No. 11
causes the orientation dependence of the CRSS. Without this interaction, no orientation dependence is observed. Conclusions For the first time compressive flow stress of binary stoichiometric single-crystalline NijAl has been studied for orientations between [OOl] and [Ol l]. All orientations exhibit anomalous positive temperature dependence of the flow stress with a weakly orientation-dependent peak temperature between 973- 1073 K. The CRSS on (11 l)[ 70 l] had almost no orientation dependence in the temperature range of the yield stress anomaly. This is considered to be the nature of binary stoichiometric N&Al. Previously reported orientation-dependent CRSS in Ni-rich and ternary NisAl is thought to be due to the point defects induced by off-stoichiometry and ternary elements. Acknowledgments This research has been carried out under the Japanese Science and Technology Agency (STA) Fellowship Program. D. Golberg expresses his gratitude for the award of an STA post-doctoral Fellowship, tenable at the National Research Institute for Metals, Tsukuba, Japan. The authors wish to thank E.P. George at Oak Ridge National Laboratory for many helpful discussions. References 1. 2. 3. 4. 5. 6. 7. 8. 9. 10. 11. 12. 13. 14. 15. 16. 17.
D.P. Pope and S.S. Ezz, Int. Mat. Rev., 29, 136 (1984). N.S. Stoloff, Int. Mat. Rev. 34, 153 (1989). F.E. Heredia and D.P. Pope, Acta Metall., 39,2027 (1991). E. Kuramotc and D.P. Pope, Acta Metall., 26,207 (1978). S.S. Ezz, D.P. Pope and V. Paidar, ActaMetall., 30,921(1982). Y. Umakoshi, D.P. Pope and V. Paidar, Acta Metall., 32,449 (1984). C. Lall, S. Chin and D.P. Pope, Metall. Trans. A, lOA, 1323 (1979). M.S. Kim, S. Hanada, S. Watanabe and 0. Izumi, ActaMetall., 36,2615 (1988). T. Suzuki, Y. Mishima and S. Miura, ISIJ International, 29,1 (1989). M. Demura and T. Hirano, Phil. Msg. Letters, 75, 143 (1997). M. Demura and T. Hirano, in Mat.Res. Sot. Proc. (Boston, 2-6 Dec., 1996, USA) in press. D. Golberg, M. Demura and T. Hirano, in Proc. 2nd Inter. Symp. on Structural Intermetallics, (Sept. 1997, Seven Springs, Pennsylvania, USA) in press. S. Ochiai, S. Miura, Y. Miihima and T. Suzuki, J. Jpn. Inst. of Metals, 51,608 (1987). S. Takeuchi and E. Kuramoto, Acta metall., 21,415 (1973). V. Paidar, D.P. Pope and V. Vitek, Acta MetaIl., 32,449 (1984). Y. Mishiia, S. Ochiai, N. Hamao, M. Yodogawa and T. Suzuki, Trans. Jpn. Inst. Metals, 27,648 (1986). 0. Noguchi, Y. Gya and T. Suzuki, Metal. Trans., 12A, 1647 (1981).