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Dislocation structure in γ phase grains of Ti55Al45 alloy
Scripta METALLURGICA et MATERIALIA
Vol. 25, pp. 543-548, 1991 Printed in the U.S.A.
Pergamon Press plc All rights reserved
DISLOCATION STRUCTURE IN 7 PHASE GRAINS OF Tiss?gl4s ALLOY J.Y. Kim, Y.D. Hahn and S.H. Whang Department of Metallurgy and Materials Science Polytechnic University 333 Jay Street, Brooklyn, NY 11201 (Received October 8, 1990) (Revised December 13, 1990)
Introduction In A1 rich, single y phase(Llo) Ti-A1 alloys deformed at room temperature, the dislocations neither decompose nor cross-slip[I,2]. However, upon deformation at high temperatures, the ordinary dislocations00=1/2<110]) become curled[3] and superdislocations(b=<101]) tend to decompose into perfect dislocations[4-6], or undergo cross-slip[4,5]. Recently, the two phase alloys (7+az) in the Ti-A1 system have shown much higher ductilities at room temperature than those of Al-rich single y-phase alloys[7,8]. Furthermore, a y-Ti-A1-Mn alloy poor in AI(< 50 at %) showed an enhanced ductility which was accompanied by the decomposition of superdislocations into perfect dislocations upon deformation at room temperature[9]. For this reason, it is very important to understand the dislocation structure in the Al-poor, ? phase of Ti-Al alloys in connection with the enhanced ductility. However, a study on an Al-poor, single 7 phase of the Ti-A1 system is not possible due to the fact that such a phase is not stable in a single phase, Llo structure. Alternatively, one can investigate the dislocation structure in the ¥ phase of the two phase(,/ +c~) alloys since the y phase in the lamellar structure has an A1 content of less than 49 at % depending upon the solidification rate and the subsequent heat treatment. Experimental Pure titanium (>99.9 %) and aluminum (>99.99 %) were melted together into small alloy buttons in an arc furnace under an argon atmosphere. The alloy buttons were annealed for a week in vacuum at 1273 K, after which the oxygen content is typically 500 ppm. Subsequently, the buttons were machined into rectangular specimens (3x3x7 mm) by a spark cutter and deformed at room temperature to 2 % by uniaxial compression with a strain rate of 3x10a s "I. The deformed specimens were sliced into thin sheets and electropolished into thin foils by a twin jet polisher for TEM study. Results After 2 % deformation, the y phase layers of the lamellar structure in TissAlas shown in Fig. la-c contain predominantly 1/2<110] type dislocations and a few <101] type superdislocations. Dislocations "A","B"and "C" in Fig. la and b were identified as having the Burgers vectors [011], 1/21110] and 1/211]'0], respectively. The 'qY' and "C" dislocations are of the same type, but the "B" dislocations are screw while the "C" dislocations are either mixed(C0 or edge(Cz) in character. In Fig. la, 1/2<110] dislocations are invisible while the [011] superdislocations labeled "A" are visible. One of the "A" superdislocations of "S" shape, indicated with an arrow sign in Fig. lc, is curved when viewed from the [-1-10] beam direction. This indicates that the "A" dislocation (indicated by an arrow) does not lie in any of the planes (111), CI'T1), (110) or (001). Also, the dislocation "A" was viewed from [010] direction as shown in Fig. la and f. The curved portion of the "A" dislocation remains curved, indicating that the curved portion is in neither the (100) nor the (001) plane. Therefore, the 'iS" shaped segment of the superdislocation might have undergone cross-slip.
543 0036-9748/91 $3.00 + .00 Copyright (c) 1991 Pergamon Press plc
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STRUCTURE
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The complete determination of the cross-slip would require an examination from another beam direction, namely [110] whch could not be achieved. The fully annealed TissA145 alloy contains y phase layers of the lamellar structure that are free of dislocations, as shown in Fig. 2a. However, after deformation at RT, the y phase, the bright layers in Fig. 2a, were decorated with dislocations generated during deformation. Dislocation "A" is invisible with g=(1T1)(Fig. 2c); dislocation '13" is invisible with g=('2"20) (Fig. 2d); and dislocation "C" is invisible with g=(002) (Fig. 2b). Dislocations "A" and '23" in Fig. 2b were found to possess the Burgers vectors [10"1"]and 1/21112"], respectively. The "C" dislocation with Burgers vector 1/211"1-0] is connected to the junction of "A" and '13" dislocations. The majority of dislocations in this ~/layer originate from the interface between the c~ and y phases, as indicated with arrows in Fig. 2c. Also, indicated with arrow signs are 1/211T0] dislocations('1~") initiated from the interface, shown on the left side of Fig. 2c. The distribution of dislocations in the y layer in Fig. 2c is rather localized along the interface. Obviously, misfit dislocations with an equal spacing at the interface[10] must exist, but are not shown in these micrographs. The dislocation structure in the primary y grains was examined as shown in Fig. 3a-e. Dislocations A, B and D in Fig. 3a, b, d, and e, identified by the contrast principle, have Burgers vectors: [10T], 1/211-1-2-] and 1/21110], respectively. The "A" and '23" dislocations are edge in character and the "C" and '19" dislocations are mixed. Again, these dislocations A, B, and D meet at a node point "N" shown in Fig. 3b. Most superdislocations appear straight while the ordinary dislocations are extensively curved in shape. Discussion The dislocation dissociation at the node observed in the y phase of the lamellae, and the primary y grains in Ti~Al4s can be written as <101] --* 1/2 <112] + 1/2
Vol. 25, No. 3
DISLOCATION
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545
in shape in contrast with the curled shapes in the y phase with less than 50 at% A1 at RT. On the basis of the morphology of 1/2<110] dislocations in the Al-poor, y TiAl alloy, one can postulate that the Peierls stress is low due to the imbalance in Ti-A1 bonds, which makes kinks move more easily, and thereby lower the flow stress in the ~¢phase. In the lamellar structure of TissA145alloy, dislocations appear to originate from the interface between the ct2 and "/phases upon deformation. A possible mechanism may be that the mismatch at the interface between the two phases increases with increasing applied stress due to the difference in the modulus of phase and that of the y phase; therefore, additional dislocations may be generated at the interface and propagated into the matrix via an appropriate propagating mechanism, in addition to the initial mismatch dislocations along the interface[10]. Conclusions 1.
Dislocations in the y-phase of the lamellar structure in TissA145alloy consist of predominantly 1/2<110] type dislocations and a small portion of superdislocations.
.
A portion of the <101] type superdislocations dissociate into perfect dislocations upon room temperature deformation in Al-poor, ~¢phase grains of TissA145,consistent with observation in an Al-poor, y TiA1-2 wt% Mn alloy.
3.
A large portion of the deformation dislocations in the y-phase of the lamellar structure in Ti~A~ originate from the interface boundary between the y and ~t2 layers, indicating that a large mismatch at the interface during deformation may promote a dislocation source or sink behavior.
4.
Ordinary dislocations in the ,/layers of the lameUae and y grains in T%sA145alloys are curved and bent, similar to those observed in Al-rich, single y phase Ti-Al alloys deformed at high temperatures. References
1. 2,
3. 4. 5. 6. 7. 8. 9.
10. 11.
S.H. Whang and Y.D. Hahn, Proc. High Temperature Aluminides and Intermetallics, S.H. Whang, C.T. Liu, D.P. Pope and J.O. Stiegler eds, pp. 91-110 (1990). G. Hug, A. Loiseau and A. Lasaimonie, Phil. Mag. A, 54, 47-65 (1986). S.H. Whang and Y.D. Hahn, Scripta Met., Vol. 24, pp.485-490(1990). Y.D. Hahn and S.H. Whang, Scripta Met., Vol. 24, pp.139-144 (1990). G. Hug, A. Loiseau and P. Veyssiere, Phil. Mag. A, 57, pp. 499-523(1988). Y.D. Hahn and S.H. Whang, submitted to Met. Trans. A. S. C. Huang and E.L. Hall, MRS Res. Soc. Sym. Proc., C.T. Liu, A.T. Taub, N.S. Stoloff, and C.C. Koch, Vol. 133, pp. 373-383 (1989); Ibid, T. Kawabata, T. Tamura and O. Izumi, pp. 329-334. Y.W. Kim and F.H. Froes, Proc. High Temperature Aluminides and Intermetallics, pp. 465-492, S.H. Whang, C.T. Liu, D.P. Pope, and J.O. Stiegler eds, TMS Inc.(1990). T. Hanamura and M. Tanino, J. Mat. Sci. Lett, 8, 24-28(1989). C.R. Fang, D.J. Michel and C.R. Crowe, Scripa Met., 33, 1707-12(1989). B.A. Greenberg, Scripta Metallurgica, vol. 23, 631-636(1989).
546
Vol. 25, No. 3
DISLOCATION STRUCTURE IN TissAI45
:(d)
Fig. 1: Deformation dislocations in ? phase region of the lamellae in TissAl4salloy: a) BF; B=[010], g=(002); b) BF, B:[010], g=(200); c) BF, B=[T10], g=(002); d) BF, B=[T10], g=(TTT); e) BF, B=[T10], g=(11T); f) BF, B=[010], g=(20~-); "A": [011] dislocations; "B":1/21110] dislocations; "C': 1/211-1-0]dislocations.
Vol. 25, No. 3
DISLOCATION STRUCTURE IN TissAI45
(c)
547
(d)
Fig. 2 Deformation dislocations in y phase layers of lamellar structure in Ti~AI~: a) larnellar structure in fully annealed Ti55A145alloys at 1273 K, 1 week; g=(002); b) BF, B=[ll0], g=(002); c) BF, B=[ll0], g=(1T1); d) BF, B=[ll0], g=!~20); the alloys in (b)-(d) deformed to 2% at RT; "A": [10T] dislocations; "B": 1/21112-] dislocations; 'C": 1/2[1T0] dislocations.
548
DISLOCATION STRUCTURE IN TissAI45
Vol. 25, No. 3
i
(c)
(a)
Deformation dislocations in ~/grains in Ti~Al~ alloys: a) BF, B=[110], g=(O02);
b) BF, B=[110], g=(1-n'); c) BF, B=[110], g=Ci-l-f); d) BF, B=[010], g=(200); e) BF, B=[010], g=~0~3; where dislocations "A", '~", "C" and 'I9" have Burgers vectors [10T], 1/211-1-2-], 1/211-1-0] and 1/21110].
(e)
Vol. 25, pp. 543-548, 1991 Printed in the U.S.A.
Pergamon Press plc All rights reserved
DISLOCATION STRUCTURE IN 7 PHASE GRAINS OF Tiss?gl4s ALLOY J.Y. Kim, Y.D. Hahn and S.H. Whang Department of Metallurgy and Materials Science Polytechnic University 333 Jay Street, Brooklyn, NY 11201 (Received October 8, 1990) (Revised December 13, 1990)
Introduction In A1 rich, single y phase(Llo) Ti-A1 alloys deformed at room temperature, the dislocations neither decompose nor cross-slip[I,2]. However, upon deformation at high temperatures, the ordinary dislocations00=1/2<110]) become curled[3] and superdislocations(b=<101]) tend to decompose into perfect dislocations[4-6], or undergo cross-slip[4,5]. Recently, the two phase alloys (7+az) in the Ti-A1 system have shown much higher ductilities at room temperature than those of Al-rich single y-phase alloys[7,8]. Furthermore, a y-Ti-A1-Mn alloy poor in AI(< 50 at %) showed an enhanced ductility which was accompanied by the decomposition of superdislocations into perfect dislocations upon deformation at room temperature[9]. For this reason, it is very important to understand the dislocation structure in the Al-poor, ? phase of Ti-Al alloys in connection with the enhanced ductility. However, a study on an Al-poor, single 7 phase of the Ti-A1 system is not possible due to the fact that such a phase is not stable in a single phase, Llo structure. Alternatively, one can investigate the dislocation structure in the ¥ phase of the two phase(,/ +c~) alloys since the y phase in the lamellar structure has an A1 content of less than 49 at % depending upon the solidification rate and the subsequent heat treatment. Experimental Pure titanium (>99.9 %) and aluminum (>99.99 %) were melted together into small alloy buttons in an arc furnace under an argon atmosphere. The alloy buttons were annealed for a week in vacuum at 1273 K, after which the oxygen content is typically 500 ppm. Subsequently, the buttons were machined into rectangular specimens (3x3x7 mm) by a spark cutter and deformed at room temperature to 2 % by uniaxial compression with a strain rate of 3x10a s "I. The deformed specimens were sliced into thin sheets and electropolished into thin foils by a twin jet polisher for TEM study. Results After 2 % deformation, the y phase layers of the lamellar structure in TissAlas shown in Fig. la-c contain predominantly 1/2<110] type dislocations and a few <101] type superdislocations. Dislocations "A","B"and "C" in Fig. la and b were identified as having the Burgers vectors [011], 1/21110] and 1/211]'0], respectively. The 'qY' and "C" dislocations are of the same type, but the "B" dislocations are screw while the "C" dislocations are either mixed(C0 or edge(Cz) in character. In Fig. la, 1/2<110] dislocations are invisible while the [011] superdislocations labeled "A" are visible. One of the "A" superdislocations of "S" shape, indicated with an arrow sign in Fig. lc, is curved when viewed from the [-1-10] beam direction. This indicates that the "A" dislocation (indicated by an arrow) does not lie in any of the planes (111), CI'T1), (110) or (001). Also, the dislocation "A" was viewed from [010] direction as shown in Fig. la and f. The curved portion of the "A" dislocation remains curved, indicating that the curved portion is in neither the (100) nor the (001) plane. Therefore, the 'iS" shaped segment of the superdislocation might have undergone cross-slip.
543 0036-9748/91 $3.00 + .00 Copyright (c) 1991 Pergamon Press plc
544
DISLOCATION
STRUCTURE
IN TissAI45
Vol.
25, No. 3
The complete determination of the cross-slip would require an examination from another beam direction, namely [110] whch could not be achieved. The fully annealed TissA145 alloy contains y phase layers of the lamellar structure that are free of dislocations, as shown in Fig. 2a. However, after deformation at RT, the y phase, the bright layers in Fig. 2a, were decorated with dislocations generated during deformation. Dislocation "A" is invisible with g=(1T1)(Fig. 2c); dislocation '13" is invisible with g=('2"20) (Fig. 2d); and dislocation "C" is invisible with g=(002) (Fig. 2b). Dislocations "A" and '23" in Fig. 2b were found to possess the Burgers vectors [10"1"]and 1/21112"], respectively. The "C" dislocation with Burgers vector 1/211"1-0] is connected to the junction of "A" and '13" dislocations. The majority of dislocations in this ~/layer originate from the interface between the c~ and y phases, as indicated with arrows in Fig. 2c. Also, indicated with arrow signs are 1/211T0] dislocations('1~") initiated from the interface, shown on the left side of Fig. 2c. The distribution of dislocations in the y layer in Fig. 2c is rather localized along the interface. Obviously, misfit dislocations with an equal spacing at the interface[10] must exist, but are not shown in these micrographs. The dislocation structure in the primary y grains was examined as shown in Fig. 3a-e. Dislocations A, B and D in Fig. 3a, b, d, and e, identified by the contrast principle, have Burgers vectors: [10T], 1/211-1-2-] and 1/21110], respectively. The "A" and '23" dislocations are edge in character and the "C" and '19" dislocations are mixed. Again, these dislocations A, B, and D meet at a node point "N" shown in Fig. 3b. Most superdislocations appear straight while the ordinary dislocations are extensively curved in shape. Discussion The dislocation dissociation at the node observed in the y phase of the lamellae, and the primary y grains in Ti~Al4s can be written as <101] --* 1/2 <112] + 1/2
Vol. 25, No. 3
DISLOCATION
STRUCTURE
IN Ti55AI45
545
in shape in contrast with the curled shapes in the y phase with less than 50 at% A1 at RT. On the basis of the morphology of 1/2<110] dislocations in the Al-poor, y TiAl alloy, one can postulate that the Peierls stress is low due to the imbalance in Ti-A1 bonds, which makes kinks move more easily, and thereby lower the flow stress in the ~¢phase. In the lamellar structure of TissA145alloy, dislocations appear to originate from the interface between the ct2 and "/phases upon deformation. A possible mechanism may be that the mismatch at the interface between the two phases increases with increasing applied stress due to the difference in the modulus of phase and that of the y phase; therefore, additional dislocations may be generated at the interface and propagated into the matrix via an appropriate propagating mechanism, in addition to the initial mismatch dislocations along the interface[10]. Conclusions 1.
Dislocations in the y-phase of the lamellar structure in TissA145alloy consist of predominantly 1/2<110] type dislocations and a small portion of superdislocations.
.
A portion of the <101] type superdislocations dissociate into perfect dislocations upon room temperature deformation in Al-poor, ~¢phase grains of TissA145,consistent with observation in an Al-poor, y TiA1-2 wt% Mn alloy.
3.
A large portion of the deformation dislocations in the y-phase of the lamellar structure in Ti~A~ originate from the interface boundary between the y and ~t2 layers, indicating that a large mismatch at the interface during deformation may promote a dislocation source or sink behavior.
4.
Ordinary dislocations in the ,/layers of the lameUae and y grains in T%sA145alloys are curved and bent, similar to those observed in Al-rich, single y phase Ti-Al alloys deformed at high temperatures. References
1. 2,
3. 4. 5. 6. 7. 8. 9.
10. 11.
S.H. Whang and Y.D. Hahn, Proc. High Temperature Aluminides and Intermetallics, S.H. Whang, C.T. Liu, D.P. Pope and J.O. Stiegler eds, pp. 91-110 (1990). G. Hug, A. Loiseau and A. Lasaimonie, Phil. Mag. A, 54, 47-65 (1986). S.H. Whang and Y.D. Hahn, Scripta Met., Vol. 24, pp.485-490(1990). Y.D. Hahn and S.H. Whang, Scripta Met., Vol. 24, pp.139-144 (1990). G. Hug, A. Loiseau and P. Veyssiere, Phil. Mag. A, 57, pp. 499-523(1988). Y.D. Hahn and S.H. Whang, submitted to Met. Trans. A. S. C. Huang and E.L. Hall, MRS Res. Soc. Sym. Proc., C.T. Liu, A.T. Taub, N.S. Stoloff, and C.C. Koch, Vol. 133, pp. 373-383 (1989); Ibid, T. Kawabata, T. Tamura and O. Izumi, pp. 329-334. Y.W. Kim and F.H. Froes, Proc. High Temperature Aluminides and Intermetallics, pp. 465-492, S.H. Whang, C.T. Liu, D.P. Pope, and J.O. Stiegler eds, TMS Inc.(1990). T. Hanamura and M. Tanino, J. Mat. Sci. Lett, 8, 24-28(1989). C.R. Fang, D.J. Michel and C.R. Crowe, Scripa Met., 33, 1707-12(1989). B.A. Greenberg, Scripta Metallurgica, vol. 23, 631-636(1989).
546
Vol. 25, No. 3
DISLOCATION STRUCTURE IN TissAI45
:(d)
Fig. 1: Deformation dislocations in ? phase region of the lamellae in TissAl4salloy: a) BF; B=[010], g=(002); b) BF, B:[010], g=(200); c) BF, B=[T10], g=(002); d) BF, B=[T10], g=(TTT); e) BF, B=[T10], g=(11T); f) BF, B=[010], g=(20~-); "A": [011] dislocations; "B":1/21110] dislocations; "C': 1/211-1-0]dislocations.
Vol. 25, No. 3
DISLOCATION STRUCTURE IN TissAI45
(c)
547
(d)
Fig. 2 Deformation dislocations in y phase layers of lamellar structure in Ti~AI~: a) larnellar structure in fully annealed Ti55A145alloys at 1273 K, 1 week; g=(002); b) BF, B=[ll0], g=(002); c) BF, B=[ll0], g=(1T1); d) BF, B=[ll0], g=!~20); the alloys in (b)-(d) deformed to 2% at RT; "A": [10T] dislocations; "B": 1/21112-] dislocations; 'C": 1/2[1T0] dislocations.
548
DISLOCATION STRUCTURE IN TissAI45
Vol. 25, No. 3
i
(c)
(a)
Deformation dislocations in ~/grains in Ti~Al~ alloys: a) BF, B=[110], g=(O02);
b) BF, B=[110], g=(1-n'); c) BF, B=[110], g=Ci-l-f); d) BF, B=[010], g=(200); e) BF, B=[010], g=~0~3; where dislocations "A", '~", "C" and 'I9" have Burgers vectors [10T], 1/211-1-2-], 1/211-1-0] and 1/21110].
(e)