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Phase transformation and microstructures in Ti–Al–Nb–Ta system
March 2002
Materials Letters 53 Ž2002. 57–62 www.elsevier.comrlocatermatlet
Phase transformation and microstructures in Ti–Al–Nb–Ta system Jihua Peng a,b,) , Shiqiong Li b, Yong Mao b, Xunfang Sun a a
Department of Engineering Mechanics, Southwest Jiaotong UniÕersity, Chengdu 610031, People’s Republic of China b Department of Superalloy, Central Iron and Steel Research Institute, Beijing 100081, People’s Republic of China Received 4 April 2000; received in revised form 30 May 2001; accepted 1 June 2001
Abstract A variety of heat treatment has been carried out to explore the transformation and microstructure in a Ti–22Al–20Nb–7Ta. Special emphasis focused on the effect of solution andror aging in O q B2 Žorthorhombicq bcc. phase field. Two forms of the transformation of B2 to O phase were observed. They are grain size change that is controlled by grain boundary diffusion and Widmanstatten precipitation of O phase from the B2 matrix. It has been found that Ta addition in Ti 2 AlNb-based alloy ¨ could increase the Widmanstatten transformation starting temperature Ws , which determines the transformation mode of B2 ¨ to O phase. During recrystallization of O grain, it evolves into lath microstructure. The primary a 2 phase is relatively stable when Ti–22Al–20Nb–7Ta was heat treated in O q B2 phase field, while the transformation of a 2 to O phase occurred in this study. q 2002 Elsevier Science B.V. All rights reserved. Keywords: Intermetallics; Ti 2 AlNb-based alloy; Ti–Al–Nb–Ta system; Phase transformation; Microstructure
1. Introduction A promising new route for increasing the specific strength and rupture life of Ti 3 Al-based alloys, at room temperature, is the addition of a third refractory element such as Nb, Mo, V. Among them, Nb is an effective b stabilizer element. When Nb content exceeds 12 at.%, Ti 3 Al q Nb system alloys exist in various constituent phase mixtures, including a 2 ,
)
Corresponding author. National Key Laboratory for Nuclear Fuel and Materials, Nuclear Power Institute of China, Chengdu 610041, People’s Republic of China. E-mail address: [email protected] ŽJ. Peng..
B2rb, O phase. The orthorhombic phase ŽO. Žcmcm system based on Ti 2 AlNb. has similarities with the hexagonal close-packed Žhcp. a 2 phase ŽDO19 structure base on Ti 3 Al., yet differs by the lattice arrangement of Nb with respect to Ti. bror B2 indicates the bcc phase, which may either be disordered Žb . or ordered structure ŽB2. w1–4x. Transformation in these Ti 3 Al q Nb system alloys has been studied in detail. A composition invariant transformation from B2 to O has been established when Ti–24Al– 15Nb alloy was aged in B2 phase field followed by quenching in water w3,4x. The transformation from B2 to a 2 has found to follow the route of B2 ™ B19 ™ OX ™ O0 ™ a 2 in Ti–24Al–11Nb w5x. Coherent precipitation of O from a 2 has been observed in Ti–24Al–14Nb–3Mo–1V w6x. Compositions near
00167-577Xr02r$ - see front matter q 2002 Elsevier Science B.V. All rights reserved. PII: S 0 1 6 7 - 5 7 7 X Ž 0 1 . 0 0 4 5 3 - 0
J. Peng et al.r Materials Letters 53 (2002) 57–62
58
Table 1 Heat treatments for Ti22Al–27Nb and Ti–22Al–24Nb–3Ta alloys Solution treatment
Aging treatment
Temperature Ž8C.
Hold time Žh.
Cooling
Temperature Ž8C.
Hold time Žh.
Cooling
970 950 900
1 1 1
WQ WQ WQ
800 X X
24 X X
WQ X X
WQ—water quenching.
Ti–25Al–25Nb that have received recent attention are Ti–22Al–23Nb w7x, Ti–22Al–27Nb w8x, Ti– 22Al–25Nb w9x, and Ti–23Al–27Nb w10x. Transformation in these alloys has been studied thoroughly. The recently emerged class of alloys based on orthorhombic phase Ti 2 AlNb seems to represent a more appropriate alloy exploration avenue for higher strength and low density titanium aluminides alloy. Our prior studies have shown that Ta addition is beneficial to optimize microstructure and obtain good combination of yield strength and ductility at room temperature for Ti 2 AlNb-based alloy w11–13x. In this work, Ti–22Al–20Nb–7Ta was employed to study the effect of Ta on transformation in Ti 2 AlNbbased alloy, and special emphasis was focused on the solution andror aging in B2 q O phase field.
Samples for the phase transformation and microstructure study were cut from the rolled sheet with EDM. The heat treatment was scheduled as Table 1. The microstructures of the alloy were examined using X-ray diffraction ŽXRD. ŽCuK a ., scanning electron microscope ŽSEM. with back scattered electron ŽBSE., and transmission electron microscope ŽTEM.. According to the Nb content in constituent phases, in BSE images, the bright contrast,
2. Experimental procedure Cast ingots with the nominal composition Ti– 22Al–24Nb–7Ta Žat.%. were prepared by induction skull melting ŽISM.. The chemical analysis showed that the cast ingots composition was in agreement with nominal composition, and the gas impurity contents were lower Že.g., oxygen ( 500 ppm, hydrogen ( 30 ppm, nitrogen ( 80 ppm.. The cast ingots canned with stainless steel were forged to a rod with a diameter of 25 mm at 1100–1200 8C in B2 phase field after homogenizing at 1200 8C for 24 h. The rod with oxidation proof coating was resolutionized at 1100 8C for 30 min followed by aging in the O q B2 phase field for 2 h, and then rolled to 4-mm-thick sheet by multiple passes at the temperature in O q B2 phase field. Unidirectional multipass rolling was carried out with interpass re-heating at the rolling temperature for 15 min.
Fig. 1. As-rolled microstructure of Ti–22Al–20Nb–7Ta. Ža. SEM image. Žb. X-ray diffraction spectrum.
J. Peng et al.r Materials Letters 53 (2002) 57–62
gray contrast, and dark contrast are designated as B2, O and a 2 phase, respectively.
3. Results and discussion The as-rolled microstructure of Ti–22Al–20Nb– 7Ta is shown in Fig. 1a. The as-rolled microstructure contained almost equiaxed OrB2 phase grains elongated slightly, together with a little amount of a 2 phase. Fig. 1b presents XRD of the as rolled Ti– 22Al–20Nb–7Ta, and it proves the existence of a small amount of a 2 phase. Fig. 2a–c shows the microstructures of as solutionized alloy at different temperature in O q B2 phase field, 970, 950, and 900 8C, respectively. The figures show those solution treatments above 900 8C in B2 q O phase field produce equiaxed microstructures containing B2 along with O andror a 2 phase. Decreasing solution temperature, the volume fraction of O phase increases. Near 900 8C, solution treatment provides the microstructure containing equiaxed primary O and
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B2 phase with Widmanstatten O precipitation. Fig. ¨ 2d shows the TEM bright field images of the Widmanstatten. The orientation relationship between B2 ¨ matrix and O phase precipitation is ²ı11:B2´²1ı0:O, Ž110.B2´Ž001.O, which was reported in the literatures w1–3x. Thus, changes in the volume fraction of the B2 and O phase, as dictated by equilibrium requirements, are brought about by changes in the sizes of equiaxed B2 and O phases, when solutionized above 900 8C, Below 900 8C, solution treatment leads to larger O phases at the expenses of B2 grain size. Fig. 3a and b shows the microstructure of Ti– 22Al–20Nb–7Ta after solution and aging in O q B2 phase field. The microstructure contains fine lath mixture of B2 and O phase and a small amount of equiaxed a 2 phase, as shown in Fig. 3a. When aged at 800 8C Žthis temperature is located in B2 q O phase field., three kinds of phase evolution have been observed for the alloy of Ti–22Al–20Nb–7Ta. The first is that B2 transforms to B2 q O, by Widmanstatten precipitation of O phase in B2. The sec¨
Fig. 2. As solutionized microstructures of Ti–22Al–20Nb–7Ta BSE: Ža. 970 8C = 1 h; Žb. 950 8C = 1 h; Žc. 900 8C = 1 h; Žd. TEM BF image for 900 8C = 1 h.
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J. Peng et al.r Materials Letters 53 (2002) 57–62
Fig. 3. The microstructure of Ti–22Al–20Nb–7Ta treated by 970 8Cr1 hrWQr800 8Cr24 hrWQ. Ža. BSE image. Žb. TEM BF image.
ond is the prior equiaxed O grain evolves into lath microstructure. The third is that the prior a 2 phase transforms to B2 andror O. Fig. 3b shows bright field image of microphology of the alloy as aged. In Fig. 4a, there are many precipitates in prior a 2 phase. Fig. 4b shows a typical moire pattern when double beam condition is nearly met. In Fig. 4a, the inset SAD from the prior a 2 phase shows that there is an orientation between the O phase precipitation and the matrix a 2 phase, i.e., w001xO´w0001xa 2 , Ž110.O´Ž1010. a 2 . If the solution temperature is above 900 8C Žwhich is located in the B2 q O phase field., the decrease of B2 content is completed by a reduction in B2 grain size and a corresponding increase in O grain size. In these processes, the volume fractions of the O and
B2 phase are controlled by grain boundary movement. When solution-treated below 900 8C, however, volume fraction changes of B2 and O phases are no longer grain-growth controlled. Solution-treated above 900 8C, grain boundary diffusion and grain growth kinetics are high enough to allow volume fraction changes by grain boundary movement. Widmanstatten transformation requires lattice reconstruc¨ tion involving dislocations and ledges, as in martenstic transformation from a bcc to hcp phase of the same composition. Widmanstatten transformation ¨ would start at the temperature Ws , which is dictated by the necessary driving force DG B2 ™ O s D SŽT0 y Ws . w15x. T0 is designated as the temperature at which the Gibbs free energy of B2 is equal to that of O phase, and D S means the entropy change for the transformation from B2 to O phase. DT is defined as Ws y T0 , which refers to the overcooling need to initiate the Widmanstatten transformation. When so¨ lutionized andror aged in B2 q O phase field, two types of competing process occur, which are controlled by solution andror aging temperature. Only when solutionized below 900 8C, this DT requirement is met. When solutionized above 900 8C, the DT needed by driving force for Widmanstatten ¨ transformation is not met; however, the diffusion rate is high enough to initiate grain growth by grain boundary movement. Compared with Ti–23Al–27Nb with Widmanstatten O phase precipitation when so¨ lutionized below 875 8C w14x, Ta substitution with a part of Nb in Ti 2 AlNb based alloy increases the Ws . From 1025 to 875 8C isotherm in Ti–Al–NB system, the O phase-border is relatively unchanged, the equiaxed O grain evolving into lath microstructure results from recrystallization. The mechanism for lath microstructure formation by means of recrystallization needs further study. Although aged in B2 q O phase field for a long time, the transformation of a 2 phase operated slightly. This is in agreement with report in Ref. w3x, where aged for 720 h, the primary a 2 phase in Ti–24Al–15Nb has little amount of precipitation of O phase with the form of plate. This may be from the sluggish diffusion of Ti in a 2 phase. The fixed crystalline orientation relationship ŽOR. between precipitation of O phase and a 2 has also been reported in Refs. w3,6x. The lattice cell parameters of O and a 2 phase are as following: a s 0.605 nm, b s 0.98 nm, c s 0.473 nm for O
J. Peng et al.r Materials Letters 53 (2002) 57–62
61
Fig. 4. TEM images for the precipitation of O phase from a 2 phase in aged microstructure. Ža. TEF BF image and SAD. Žb. Moire pattern.
phase; a s 0.58 nm, c s 0.465 nm for a 2 phase. The difference of size of lattice cell is very small, so it is easy to explain the moire pattern in dark field image when double beam condition nearly met.
4. Conclusion The transformation of Ti–22Al–20Nb–7Ta alloy solutionized andror aged in O q B2 phase field is studied carefully. The transformation of B2 phase to O phase can take place in either the size changes of O and B2 phase, or Widemanstatten precipitation of ¨ O phase in B2 matrix. When heat-treated in B2 q O phase field, the type of B2 transformation is determined by overcooling DT, which is related to the driving force for Widmanstatten transformation. Re¨ crystallization of primary O grain produces lath microstructure. a 2 phase is relatively stable when solutionized andror aged in O q B2 field, although the phase transformation of a 2 to O occurred for Ti– 22Al–20Nb–7Ta alloy.
References w1x H.T. Kestner-Weykamp, C.H. Ward, T.F. Broderick, M.J. Kaufman, Scr. Metall. 23 Ž1989. 1697–1702. w2x A.K. Gogia, D. Banerjee, T.K. Nandy, Metall. Trans. 21A Ž1990. 609–625. w3x K. Muraleedharan, A.K. Gogia, T.K. Nandy, D. Banerjee, S. Lele, Metall. Trans. 23A Ž1992. 401–415. w4x K. Muraleedharan, T.K. Nandy, D. Banerjee, S. Lele, Metall. Trans. 23A Ž1992. 415–431. w5x L.H. Hsiung, H.N.G. Wadley, Mater. Sci. Eng. A 192r193 Ž1995. 908–913. w6x Y. Wu, Mater. Lett. 32 Ž1997. 319–323. w7x F.C. Dary, T.M. Pollock, Mater. Sci. Eng., A 208 Ž1996. 188–302. w8x R.G. Rowe, in: Y.W. Kim, R.R. Boyer ŽEds.., MicrostructurerProperty Relationship in Titanium Aluminide Alloys, The Minerals Metal and Materials Society, 1991, pp. 387– 397. w9x J. Kumpfert, C. Leyens, in: M.V. Nathal, C.T. Liu, P.L. Martin, D.B. Miracle, R. Wagner, M. Yamaguchi ŽEds.., Stuctural Intermetallics, TMS, Warrendale, PA, 1997, pp. 895–904. w10x C.J. Boehlert, B.S. Majumdar, V. Seetharman, D.B. Miracle, R. Wheeler, in: M.V. Nathal, C.T. Liu, P.L. Martin, D.B. Miracle, R. Wagner, M. Yamaguchi ŽEds.., Stuctural Intermetallics, TMS, Warrendale, PA, pp. 795–804.
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J. Peng et al.r Materials Letters 53 (2002) 57–62
w11x J.H. Peng, Y. Mao, S.Q. Li, X.F. Sun, Trans. Nonferrous Met. Soc. China 10 Ž3. Ž2000. 378–381. w12x Y. Mao, S.Q. Li, J.W. Zhang, J.H. Peng, D.X. Zou, Z.Y. Zhong, Intermetallics 8 Ž2000. 659–662. w13x J.H. Peng, Y. Mao, S.Q. Li, X.F. Sun, Microstructure controlling by heat treatment and complex processing for
Ti 2 AlNb based alloys, Mater. Sci. Eng., A 299 Ž2001. 75–80. w14x C.J. Boehlert, B.S. Majumdar, V. Seetharman, D.B. Miracle, Metall. Mater. Trans. 30A Ž9. Ž1999. 2305–2323. w15x C.H.R. Rao, K.J. Rao, Phase Transformation in Solids, McGraw-Hill, New York, NY, 1978.
Materials Letters 53 Ž2002. 57–62 www.elsevier.comrlocatermatlet
Phase transformation and microstructures in Ti–Al–Nb–Ta system Jihua Peng a,b,) , Shiqiong Li b, Yong Mao b, Xunfang Sun a a
Department of Engineering Mechanics, Southwest Jiaotong UniÕersity, Chengdu 610031, People’s Republic of China b Department of Superalloy, Central Iron and Steel Research Institute, Beijing 100081, People’s Republic of China Received 4 April 2000; received in revised form 30 May 2001; accepted 1 June 2001
Abstract A variety of heat treatment has been carried out to explore the transformation and microstructure in a Ti–22Al–20Nb–7Ta. Special emphasis focused on the effect of solution andror aging in O q B2 Žorthorhombicq bcc. phase field. Two forms of the transformation of B2 to O phase were observed. They are grain size change that is controlled by grain boundary diffusion and Widmanstatten precipitation of O phase from the B2 matrix. It has been found that Ta addition in Ti 2 AlNb-based alloy ¨ could increase the Widmanstatten transformation starting temperature Ws , which determines the transformation mode of B2 ¨ to O phase. During recrystallization of O grain, it evolves into lath microstructure. The primary a 2 phase is relatively stable when Ti–22Al–20Nb–7Ta was heat treated in O q B2 phase field, while the transformation of a 2 to O phase occurred in this study. q 2002 Elsevier Science B.V. All rights reserved. Keywords: Intermetallics; Ti 2 AlNb-based alloy; Ti–Al–Nb–Ta system; Phase transformation; Microstructure
1. Introduction A promising new route for increasing the specific strength and rupture life of Ti 3 Al-based alloys, at room temperature, is the addition of a third refractory element such as Nb, Mo, V. Among them, Nb is an effective b stabilizer element. When Nb content exceeds 12 at.%, Ti 3 Al q Nb system alloys exist in various constituent phase mixtures, including a 2 ,
)
Corresponding author. National Key Laboratory for Nuclear Fuel and Materials, Nuclear Power Institute of China, Chengdu 610041, People’s Republic of China. E-mail address: [email protected] ŽJ. Peng..
B2rb, O phase. The orthorhombic phase ŽO. Žcmcm system based on Ti 2 AlNb. has similarities with the hexagonal close-packed Žhcp. a 2 phase ŽDO19 structure base on Ti 3 Al., yet differs by the lattice arrangement of Nb with respect to Ti. bror B2 indicates the bcc phase, which may either be disordered Žb . or ordered structure ŽB2. w1–4x. Transformation in these Ti 3 Al q Nb system alloys has been studied in detail. A composition invariant transformation from B2 to O has been established when Ti–24Al– 15Nb alloy was aged in B2 phase field followed by quenching in water w3,4x. The transformation from B2 to a 2 has found to follow the route of B2 ™ B19 ™ OX ™ O0 ™ a 2 in Ti–24Al–11Nb w5x. Coherent precipitation of O from a 2 has been observed in Ti–24Al–14Nb–3Mo–1V w6x. Compositions near
00167-577Xr02r$ - see front matter q 2002 Elsevier Science B.V. All rights reserved. PII: S 0 1 6 7 - 5 7 7 X Ž 0 1 . 0 0 4 5 3 - 0
J. Peng et al.r Materials Letters 53 (2002) 57–62
58
Table 1 Heat treatments for Ti22Al–27Nb and Ti–22Al–24Nb–3Ta alloys Solution treatment
Aging treatment
Temperature Ž8C.
Hold time Žh.
Cooling
Temperature Ž8C.
Hold time Žh.
Cooling
970 950 900
1 1 1
WQ WQ WQ
800 X X
24 X X
WQ X X
WQ—water quenching.
Ti–25Al–25Nb that have received recent attention are Ti–22Al–23Nb w7x, Ti–22Al–27Nb w8x, Ti– 22Al–25Nb w9x, and Ti–23Al–27Nb w10x. Transformation in these alloys has been studied thoroughly. The recently emerged class of alloys based on orthorhombic phase Ti 2 AlNb seems to represent a more appropriate alloy exploration avenue for higher strength and low density titanium aluminides alloy. Our prior studies have shown that Ta addition is beneficial to optimize microstructure and obtain good combination of yield strength and ductility at room temperature for Ti 2 AlNb-based alloy w11–13x. In this work, Ti–22Al–20Nb–7Ta was employed to study the effect of Ta on transformation in Ti 2 AlNbbased alloy, and special emphasis was focused on the solution andror aging in B2 q O phase field.
Samples for the phase transformation and microstructure study were cut from the rolled sheet with EDM. The heat treatment was scheduled as Table 1. The microstructures of the alloy were examined using X-ray diffraction ŽXRD. ŽCuK a ., scanning electron microscope ŽSEM. with back scattered electron ŽBSE., and transmission electron microscope ŽTEM.. According to the Nb content in constituent phases, in BSE images, the bright contrast,
2. Experimental procedure Cast ingots with the nominal composition Ti– 22Al–24Nb–7Ta Žat.%. were prepared by induction skull melting ŽISM.. The chemical analysis showed that the cast ingots composition was in agreement with nominal composition, and the gas impurity contents were lower Že.g., oxygen ( 500 ppm, hydrogen ( 30 ppm, nitrogen ( 80 ppm.. The cast ingots canned with stainless steel were forged to a rod with a diameter of 25 mm at 1100–1200 8C in B2 phase field after homogenizing at 1200 8C for 24 h. The rod with oxidation proof coating was resolutionized at 1100 8C for 30 min followed by aging in the O q B2 phase field for 2 h, and then rolled to 4-mm-thick sheet by multiple passes at the temperature in O q B2 phase field. Unidirectional multipass rolling was carried out with interpass re-heating at the rolling temperature for 15 min.
Fig. 1. As-rolled microstructure of Ti–22Al–20Nb–7Ta. Ža. SEM image. Žb. X-ray diffraction spectrum.
J. Peng et al.r Materials Letters 53 (2002) 57–62
gray contrast, and dark contrast are designated as B2, O and a 2 phase, respectively.
3. Results and discussion The as-rolled microstructure of Ti–22Al–20Nb– 7Ta is shown in Fig. 1a. The as-rolled microstructure contained almost equiaxed OrB2 phase grains elongated slightly, together with a little amount of a 2 phase. Fig. 1b presents XRD of the as rolled Ti– 22Al–20Nb–7Ta, and it proves the existence of a small amount of a 2 phase. Fig. 2a–c shows the microstructures of as solutionized alloy at different temperature in O q B2 phase field, 970, 950, and 900 8C, respectively. The figures show those solution treatments above 900 8C in B2 q O phase field produce equiaxed microstructures containing B2 along with O andror a 2 phase. Decreasing solution temperature, the volume fraction of O phase increases. Near 900 8C, solution treatment provides the microstructure containing equiaxed primary O and
59
B2 phase with Widmanstatten O precipitation. Fig. ¨ 2d shows the TEM bright field images of the Widmanstatten. The orientation relationship between B2 ¨ matrix and O phase precipitation is ²ı11:B2´²1ı0:O, Ž110.B2´Ž001.O, which was reported in the literatures w1–3x. Thus, changes in the volume fraction of the B2 and O phase, as dictated by equilibrium requirements, are brought about by changes in the sizes of equiaxed B2 and O phases, when solutionized above 900 8C, Below 900 8C, solution treatment leads to larger O phases at the expenses of B2 grain size. Fig. 3a and b shows the microstructure of Ti– 22Al–20Nb–7Ta after solution and aging in O q B2 phase field. The microstructure contains fine lath mixture of B2 and O phase and a small amount of equiaxed a 2 phase, as shown in Fig. 3a. When aged at 800 8C Žthis temperature is located in B2 q O phase field., three kinds of phase evolution have been observed for the alloy of Ti–22Al–20Nb–7Ta. The first is that B2 transforms to B2 q O, by Widmanstatten precipitation of O phase in B2. The sec¨
Fig. 2. As solutionized microstructures of Ti–22Al–20Nb–7Ta BSE: Ža. 970 8C = 1 h; Žb. 950 8C = 1 h; Žc. 900 8C = 1 h; Žd. TEM BF image for 900 8C = 1 h.
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J. Peng et al.r Materials Letters 53 (2002) 57–62
Fig. 3. The microstructure of Ti–22Al–20Nb–7Ta treated by 970 8Cr1 hrWQr800 8Cr24 hrWQ. Ža. BSE image. Žb. TEM BF image.
ond is the prior equiaxed O grain evolves into lath microstructure. The third is that the prior a 2 phase transforms to B2 andror O. Fig. 3b shows bright field image of microphology of the alloy as aged. In Fig. 4a, there are many precipitates in prior a 2 phase. Fig. 4b shows a typical moire pattern when double beam condition is nearly met. In Fig. 4a, the inset SAD from the prior a 2 phase shows that there is an orientation between the O phase precipitation and the matrix a 2 phase, i.e., w001xO´w0001xa 2 , Ž110.O´Ž1010. a 2 . If the solution temperature is above 900 8C Žwhich is located in the B2 q O phase field., the decrease of B2 content is completed by a reduction in B2 grain size and a corresponding increase in O grain size. In these processes, the volume fractions of the O and
B2 phase are controlled by grain boundary movement. When solution-treated below 900 8C, however, volume fraction changes of B2 and O phases are no longer grain-growth controlled. Solution-treated above 900 8C, grain boundary diffusion and grain growth kinetics are high enough to allow volume fraction changes by grain boundary movement. Widmanstatten transformation requires lattice reconstruc¨ tion involving dislocations and ledges, as in martenstic transformation from a bcc to hcp phase of the same composition. Widmanstatten transformation ¨ would start at the temperature Ws , which is dictated by the necessary driving force DG B2 ™ O s D SŽT0 y Ws . w15x. T0 is designated as the temperature at which the Gibbs free energy of B2 is equal to that of O phase, and D S means the entropy change for the transformation from B2 to O phase. DT is defined as Ws y T0 , which refers to the overcooling need to initiate the Widmanstatten transformation. When so¨ lutionized andror aged in B2 q O phase field, two types of competing process occur, which are controlled by solution andror aging temperature. Only when solutionized below 900 8C, this DT requirement is met. When solutionized above 900 8C, the DT needed by driving force for Widmanstatten ¨ transformation is not met; however, the diffusion rate is high enough to initiate grain growth by grain boundary movement. Compared with Ti–23Al–27Nb with Widmanstatten O phase precipitation when so¨ lutionized below 875 8C w14x, Ta substitution with a part of Nb in Ti 2 AlNb based alloy increases the Ws . From 1025 to 875 8C isotherm in Ti–Al–NB system, the O phase-border is relatively unchanged, the equiaxed O grain evolving into lath microstructure results from recrystallization. The mechanism for lath microstructure formation by means of recrystallization needs further study. Although aged in B2 q O phase field for a long time, the transformation of a 2 phase operated slightly. This is in agreement with report in Ref. w3x, where aged for 720 h, the primary a 2 phase in Ti–24Al–15Nb has little amount of precipitation of O phase with the form of plate. This may be from the sluggish diffusion of Ti in a 2 phase. The fixed crystalline orientation relationship ŽOR. between precipitation of O phase and a 2 has also been reported in Refs. w3,6x. The lattice cell parameters of O and a 2 phase are as following: a s 0.605 nm, b s 0.98 nm, c s 0.473 nm for O
J. Peng et al.r Materials Letters 53 (2002) 57–62
61
Fig. 4. TEM images for the precipitation of O phase from a 2 phase in aged microstructure. Ža. TEF BF image and SAD. Žb. Moire pattern.
phase; a s 0.58 nm, c s 0.465 nm for a 2 phase. The difference of size of lattice cell is very small, so it is easy to explain the moire pattern in dark field image when double beam condition nearly met.
4. Conclusion The transformation of Ti–22Al–20Nb–7Ta alloy solutionized andror aged in O q B2 phase field is studied carefully. The transformation of B2 phase to O phase can take place in either the size changes of O and B2 phase, or Widemanstatten precipitation of ¨ O phase in B2 matrix. When heat-treated in B2 q O phase field, the type of B2 transformation is determined by overcooling DT, which is related to the driving force for Widmanstatten transformation. Re¨ crystallization of primary O grain produces lath microstructure. a 2 phase is relatively stable when solutionized andror aged in O q B2 field, although the phase transformation of a 2 to O occurred for Ti– 22Al–20Nb–7Ta alloy.
References w1x H.T. Kestner-Weykamp, C.H. Ward, T.F. Broderick, M.J. Kaufman, Scr. Metall. 23 Ž1989. 1697–1702. w2x A.K. Gogia, D. Banerjee, T.K. Nandy, Metall. Trans. 21A Ž1990. 609–625. w3x K. Muraleedharan, A.K. Gogia, T.K. Nandy, D. Banerjee, S. Lele, Metall. Trans. 23A Ž1992. 401–415. w4x K. Muraleedharan, T.K. Nandy, D. Banerjee, S. Lele, Metall. Trans. 23A Ž1992. 415–431. w5x L.H. Hsiung, H.N.G. Wadley, Mater. Sci. Eng. A 192r193 Ž1995. 908–913. w6x Y. Wu, Mater. Lett. 32 Ž1997. 319–323. w7x F.C. Dary, T.M. Pollock, Mater. Sci. Eng., A 208 Ž1996. 188–302. w8x R.G. Rowe, in: Y.W. Kim, R.R. Boyer ŽEds.., MicrostructurerProperty Relationship in Titanium Aluminide Alloys, The Minerals Metal and Materials Society, 1991, pp. 387– 397. w9x J. Kumpfert, C. Leyens, in: M.V. Nathal, C.T. Liu, P.L. Martin, D.B. Miracle, R. Wagner, M. Yamaguchi ŽEds.., Stuctural Intermetallics, TMS, Warrendale, PA, 1997, pp. 895–904. w10x C.J. Boehlert, B.S. Majumdar, V. Seetharman, D.B. Miracle, R. Wheeler, in: M.V. Nathal, C.T. Liu, P.L. Martin, D.B. Miracle, R. Wagner, M. Yamaguchi ŽEds.., Stuctural Intermetallics, TMS, Warrendale, PA, pp. 795–804.
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J. Peng et al.r Materials Letters 53 (2002) 57–62
w11x J.H. Peng, Y. Mao, S.Q. Li, X.F. Sun, Trans. Nonferrous Met. Soc. China 10 Ž3. Ž2000. 378–381. w12x Y. Mao, S.Q. Li, J.W. Zhang, J.H. Peng, D.X. Zou, Z.Y. Zhong, Intermetallics 8 Ž2000. 659–662. w13x J.H. Peng, Y. Mao, S.Q. Li, X.F. Sun, Microstructure controlling by heat treatment and complex processing for
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