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Phase equilibrium and martensitic transformation in near equiatomic TiPd alloys
Materials Science and Engineering A 438–440 (2006) 327–331
Phase equilibrium and martensitic transformation in near equiatomic Ti Pd alloys T. Yamamuro ∗ , Y. Morizono, J. Honjyo, M. Nishida Department of Materials Science and Engineering, Kumamoto University, 2-39-1 Kurokami, Kumamoto 860-8555, Japan Received 27 May 2005; received in revised form 11 November 2005; accepted 10 January 2006
Abstract The homogeneity range of the TiPd compound is experimentally determined from martensitic transformation behaviors of solution treated and aged specimens with various compositions and diffusion couple method, and then the phase diagram of Ti Pd binary system in Ti side is reinvestigated. The phase boundary of TiPd compound on the Pd side is almost constant at about 50.5 at.% Pd. On the other hand, that on the Ti side extends from 50.6 at.% Ti at around 873 K to about 51.8 at.% Ti at 1073 K. From the diffusion couple experiments, it is recognized that there is an A2(bcc) to B2 order–disorder transformation around 30 at.% Pd instead of a -Ti and TiPd duplex phase region. It is also confirmed that the Ti3 Pd compound is formed by a peritectoid reaction between -Ti and Ti2 Pd. From these results, we propose a new phase diagram of the Ti Pd binary system. © 2006 Elsevier B.V. All rights reserved. Keywords: High temperature shape memory alloy; Homogeneity range of TiPd compound; A2/B2 transformation; Peritecroid reaction; Ti3 Pd compound
1. Introduction Ti Pd shape memory alloys of near equiatomic compositions are highlighted as potential materials for high temperature applications, since the intermetallic phase of TiPd undergoes a thermoelastic martensitic transformation around 800 K [1]. However, there is no sufficient information about the phase equilibrium around the TiPd compound, i.e., near-equiatomic composition. In the present study, the homogeneity range of the TiPd compound is experimentally determined from martensitic transformation behaviors of solution treated and aged specimens with various compositions. Subsequently, the diffusion couple method is applied to examine the phase equilibrium. Based on these results, the phase diagram of the Ti Pd binary system on the Ti side is reconstructed. The formation of Ti3 Pd compound by a peritectoid reaction of -Ti and Ti2 Pd is also described. 2. Experimental procedure Ti–15 to 60 at.% Pd alloys were prepared from 99.7 mass% Ti and 99.8 mass% Pd by arc melting in an argon atmosphere. The ∗
Corresponding author. Tel.: +81 96 342 3707; fax: +81 96 342 3710. E-mail address: [email protected] (T. Yamamuro).
0921-5093/$ – see front matter © 2006 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2006.01.120
ingots were homogenized at 1273 K for 36 ks in vacuum. Rods of 3 mm in diameter were spark cut from the ingots. Disks of 1.0 mm in thickness for differential scanning calorimetry (DSC) were cut from the rods. They were solution-treated in vacuum at 1273 K for 3.6 ks, and then quenched into ice water. Some of the disks were aged from 873 to 1273 K with 100 K intervals for 36–360 ks. Diffusion couples with combinations of appropriate alloys were prepared by partial arc melting technique. They were annealed from 948 to 1273 K for 36–1800 ks, and then quenched into ice water. Concentration–penetration profiles in the diffusion zone were obtained by electron probe microanalysis. Transmission electron microscopy (TEM) was also applied to the identification of diffusion products in the couple. Microstructure examination of some alloys after heat treatment was carried out by scanning electron microscope (SEM). 3. Results and discussion Fig. 1 shows the summary of experimental results plotted at the Ti-rich portion in Ti–Pd binary phase diagrams assessed by Okamoto [2]. Major modifications are summarized as follows. The phase boundary of TiPd compound in Pd side is almost constant about 50.5 at.% Pd. On the other hand, that of in Ti side extends from 50.6 at.% Ti at around 873 K to about 51.8 at.% Ti at 1073 K. There is A2/B2 order–disorder transformation as
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Fig. 1. Summary of experimental results plotted at the Ti-rich portion in Ti–Pd binary phase diagrams assessed by Okamoto [2]. Open circles are obtained from diffusion couple experiments and closed ones are deduced from transformation temperature of quenched and aged specimens measured by DSC.
plotted by the broken line around 30 at.% Pd above 1180 K, instead of the two phase region consisting of -Ti and TiPd. The other new finding is that Ti3 Pd compound is formed below 963 ± 10 K with peritectoid reaction between -Ti and Ti2 Pd. The evidence of these modifications is discussed below. The homogeneity range of TiPd compound is estimated from transformation behaviors of quenched and aged near-equiatomic alloys. Fig. 2 shows two typical DSC curves of quenched alloys after solution treatment. In the Ti–45 at.% Pd alloy, two endothermic and exothermic peaks are clearly observed upon heating and cooling. In other words, there is successive transformation in the Ti-rich alloys. We define the first and second peaks on heating as A2* and A1* , respectively. In the same way, the first and second peaks on cooling are denoted as M1* and M2* . The inversion of A2* and M2* is clearly seen. However, it must be also noted that A1* lies very close to the reverse transformation peak temperature of the Ti–50 at.% Pd alloy. In the Ti–50 at.% Pd alloy, only a single sharp peak is observed upon heating and cooling and there is no inversion between reverse and forward transformations. The origin of these phenomena has been described in our previous reports [3,4] and is briefly summarized as follows. There is no decomposition in the Ti–50 at.% Pd alloy, during the DSC measurement. On the other hand, the precipitation of fine Ti2 Pd particles takes place in the Ti-rich alloys during the DSC measurement. Therefore, the peak of A1* in the DSC heating curve represents the reverse martensitic transformation of the TiPd matrix. The peak temperature increases with increasing Pd content and reaches a maximum at 50 at.% Pd as shown later in Fig. 3. In other word, the first peak overlaps the second peak at 50 at.% Pd. The peak of A2* is due to the reverse martensitic transformation in the local neighboring areas around the precipitates in which the Pd concentration is considered to be equiatomic, since the peak temperature is almost constant with Pd content and is identical to that of the Ti–50 at.% Pd alloy. Fig. 3(a) shows the compositional dependence of transformation peak temperatures during heating process in the quenched and furnace cooled alloys. There is no difference of transformation behavior and temperature in the quenched and the furnace
Fig. 2. Typical DSC curves of near-equiatomic Ti Pd alloys quenched from 1273 K.
cooled Ti–50 at.% Pd alloy. In the furnace cooled Ti-rich alloys, only one peak was observed in heating curve. The temperature of this peak for all the compositions lies very close to the peak temperature of Ti–50 at.% Pd alloy. On the other hand, two peaks appear on the heating curves of quenched Ti-rich alloys. The temperature of A2* increases with increasing Pd content and reaches the maximum at 50 at.% Pd. The temperature of A1* is almost constant with Pd content and is identical to the transformation peak temperatures A* of the furnace cooled specimens. The change of transformation behavior with Pd content and thermal history suggests that the phase boundary of TiPd compound in Ti-rich side has significant variation with temperature and concentration. Fig. 3(b) shows the aging time and temperature dependence of transformation peak temperatures in the Ti–47 at.% Pd alloys. Here, we plot A1* and A2* in the specimens aged at various temperature for 3.6 ks and relatively prolonged period. The relatively prolonged period is defined as the aging time by which the transformation behavior has become stable. All the specimens aged for 3.6 ks show two endothermic peaks independently of aging temperature as described above. In this case, the A1* presents no significant variation with the aging temperature, while A2* firstly increases and then decreases with decreasing aging temperature below 1073 K. In the specimens aged above 1123 K for 3.6 ks, the A2* shows almost constant temperature. On the other hand, after relatively long time aging
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Fig. 3. (a) Compositional dependence of transformation peak temperatures during heating in near-equiatomic Ti Pd alloys quenched from 1273 K. (b) Effect of aging temperature on the transformation peaks in Ti–47 at.% Pd.
only one endothermic peak appears in the specimen aged below 1023 K. Although the A1* and A2* in the specimens aged above 1073 K for prolonged period are slightly higher than those for 3.6 ks, the same successive transformation takes place essentially. As apparently the compositional and/or structural changes may occur during aging in the specimens aged below 1023 K. On the other hand, above 1073 K the phase equilibrium is completed within aging for 3.6 ks and/or no diffusional process, such as decomposition occurs during aging. The slight increase of A1* and A2* in the specimen aged above 1073 K for prolonged period is probably due to the decrease of Ti concentration with evaporation and/or oxidation of Ti during the aging. Comparing Fig. 3(a) and (b), we can estimate roughly the homogeneity range of TiPd compound. The A2* of Ti–47 at.% Pd alloy aged at 873 K for prolonged period in Fig. 3(b) is about 853 K which can be extrapolated to the A2* of Ti–49.4 at.% Pd in Fig. 3(a). This indicates that the matrix composition of the specimen in equilibrium with Ti2 Pd phase is about Ti–49.4 at.% Pd at 873 K. In the same way, the equilibrium compositions of matrix are estimated to be Ti–49.0, 48.7 and 48.2 at.% Pd at 973, 1023 and 1073 K, respectively. These compositions are plotted with the closed circles in Fig. 1. It is also noted that the phase boundary of TiPd single and TiPd + Ti2 Pd duplex phases in the Ti–47 at.% Pd alloy is considered to be temperature range between 1073 and 1123 K. These results indicate that the phase boundary TiPd compound in Ti side extends from near-equiatomic composition about 50.6 at.% Ti at around 873 K to about 51.8 at.% Ti at 1073 K, which is confirmed by following diffusion couple experiments. The concentration–penetration curves for Pd in the vicinity of the diffusion zone in diffusion couples examined are divided into three types as shown in Fig. 4. Fig. 4(a) shows the curve in Ti/Ti–50 at.% Pd couple annealed at 1273 K for 36 ks. Instead of concentration gap corresponding to boundary between -Ti and TiPd compound, there is a clear singularity in the profile of Pd near 30 at.% Ti, which suggests that the critical A2/B2 ordering boundary is located at this concentration [5]. The profile of Pd in Fig. 4(b) is obtained from
Ti/Ti–46 at.% Pd couple annealed at 998 K for 1440 ks. There are three concentration gaps corresponding to ␣-Ti/-Ti, -Ti/Ti2 Pd and Ti2 Pd/TiPd boundaries. Fig. 4(c) shows the profile of Pd in Ti/Ti–46 at.% Pd couple annealed at 948 K for 3600 ks. There are three concentration gaps corresponding to ␣-Ti/-Ti, -Ti/Ti3 Pd and Ti3 Pd/Ti2 Pd boundaries. Although the Ti2 Pd/TiPd boundary was also observed in the profile, we omit it because of space limitation. In order to confirm the formation of Ti3 Pd, TEM observations are carried out in Ti/Ti–46 at.% Pd couple annealed at 948 K for 1800 ks. Fig. 5(a) shows a bright field image of the diffusion product between -Ti and Ti2 Pd in the same diffusion couple in Fig. 4(c). Electron diffraction patterns in Fig. 5(b)–(d) are taken from areas B, C and D in (a).
Fig. 4. Three typical concentration profile in diffusion couple experiments. (a) Ti/Ti–50 at.% Pd couple annealed at 1273 K for 36 ks, (b) Ti/Ti–46 at.% Pd couple annealed at 998 K for 1440 ks and (c) Ti/Ti–46 at.% Pd couple annealed at 948 K for 3600 ks.
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Fig. 5. (a) Bright field image, (b–d) electron diffraction patterns taken from areas B, C and D in (a), showing Ti3 Pd compound formed in pure Ti/Ti–46 at.% Pd diffusion couple annealed at 948 K for 1800 ks.
These patterns can be indexed consistently with -Ti of bcc, Ti3 Pd of A15 and Ti2 Pd of C11b structures, respectively. It is inevitably that there is Ti3 Pd phase between -Ti and Ti2 Pd in the diffusion couple of Ti/Ti–46 at.% Pd annealed at 948 K. However, the formation temperature range of Ti3 Pd is higher than that in the currently available phase diagram in Fig. 1 [2,6], which will be discussed later. All the results obtained by
the diffusion couple experiments are plotted by open circle in Fig. 1. In order to confirm the formation temperature and process of the Ti3 Pd compound, Ti–25 at.% Pd alloy was heat-treated from 948 to 998 K with 5 K interval for 360 ks. The Ti3 Pd could be detected with SEM scale in the specimen aged below 963 K. There is no Ti3 Pd in the specimens aged above 968 K. The for-
Fig. 6. SEM micrographs showing the formation of Ti3 Pd compound with peritectoid reaction between -Ti and Ti2 Pd compound in Ti–25 at.% Pd alloy aged at (a) 973 K for 360 ks, (b–d) 948 K for 230.4, 360 and 518.4 ks, respectively.
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mation process of Ti3 Pd is demonstrated by SEM micrographs in Fig. 6. There is duplex structure consisting of -Ti and Ti2 Pd in the specimen aged at 973 K for 360 ks in Fig. 6(a), which is the initial microstructure of following experiments. Fig. 6(b)–(d) show the microstructure changes in the specimens aged 948 K for 230.4, 360 and 518.4 ks after those aged at 973 K for 360 ks, respectively. The interface reaction is clearly seen at the boundary between -Ti and Ti2 Pd. The amount of the reaction product increases with increasing the aging period. This is a typical microstructure change upon peritectoid reaction. We can conclude that the Ti3 Pd compound is formed by the peritectoid reaction between -Ti and Ti2 Pd at 963 ± 5 K.
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firmed that Ti3 Pd compound is formed with peritectoid reaction between -Ti and Ti2 Pd at 963 ± 5 K. From these results, we have proposed the new phase diagram of Ti Pd binary system. Acknowledgements This work was supported by Grant-in-Aid for Exploratory Research from the Ministry of Education, Culture, Sports, Science and Technology and Grant-in-Aid for Scientific Research (B) from Japan Science Promotion Society, COE-directing Research Program B at Kumamoto University on Nano-space Electrochemistry and Special Program at Kumamoto University for Promoting Research on Nano-space Electrochemistry.
4. Conclusions The phase diagram of near-equiatomic and Ti-rich side of Ti Pd binary system has been reinvestigated. The phase boundary of TiPd compound in Pd side is almost constant about 50.5 at.% Pd. On the other hand, that of in Ti side extends from 50.6 at.% Ti at around 873 K to about 51.8 at.% Ti at 1073 K. From the diffusion couple experiments, it is recognized that there is A2(bcc) to B2 order–disorder transformation around 30 at.% Pd, instead of -Ti and TiPd duplex phase region. It is also con-
References [1] H.C. Donkersloot, J.H.N. Van Vucht, J. Less-Common Met. 20 (1970) 83–91. [2] H. Okamoto, J. Phase Equilib. 14 (1993) 1–9. [3] V.C. Solomon, M. Nishida, Mater. Trans. 43 (2002) 897–901. [4] V.C. Solomon, M. Nishida, Mater. Trans. 43 (2002) 908–915. [5] R. Kainuma, I. Ohnuma, K. Ishikawa, K. Ishida, Intermetallics 8 (2000) 869–875. [6] A.F. Jankowski, J. Alloys Compd. 182 (1992) 35–42.
Phase equilibrium and martensitic transformation in near equiatomic Ti Pd alloys T. Yamamuro ∗ , Y. Morizono, J. Honjyo, M. Nishida Department of Materials Science and Engineering, Kumamoto University, 2-39-1 Kurokami, Kumamoto 860-8555, Japan Received 27 May 2005; received in revised form 11 November 2005; accepted 10 January 2006
Abstract The homogeneity range of the TiPd compound is experimentally determined from martensitic transformation behaviors of solution treated and aged specimens with various compositions and diffusion couple method, and then the phase diagram of Ti Pd binary system in Ti side is reinvestigated. The phase boundary of TiPd compound on the Pd side is almost constant at about 50.5 at.% Pd. On the other hand, that on the Ti side extends from 50.6 at.% Ti at around 873 K to about 51.8 at.% Ti at 1073 K. From the diffusion couple experiments, it is recognized that there is an A2(bcc) to B2 order–disorder transformation around 30 at.% Pd instead of a -Ti and TiPd duplex phase region. It is also confirmed that the Ti3 Pd compound is formed by a peritectoid reaction between -Ti and Ti2 Pd. From these results, we propose a new phase diagram of the Ti Pd binary system. © 2006 Elsevier B.V. All rights reserved. Keywords: High temperature shape memory alloy; Homogeneity range of TiPd compound; A2/B2 transformation; Peritecroid reaction; Ti3 Pd compound
1. Introduction Ti Pd shape memory alloys of near equiatomic compositions are highlighted as potential materials for high temperature applications, since the intermetallic phase of TiPd undergoes a thermoelastic martensitic transformation around 800 K [1]. However, there is no sufficient information about the phase equilibrium around the TiPd compound, i.e., near-equiatomic composition. In the present study, the homogeneity range of the TiPd compound is experimentally determined from martensitic transformation behaviors of solution treated and aged specimens with various compositions. Subsequently, the diffusion couple method is applied to examine the phase equilibrium. Based on these results, the phase diagram of the Ti Pd binary system on the Ti side is reconstructed. The formation of Ti3 Pd compound by a peritectoid reaction of -Ti and Ti2 Pd is also described. 2. Experimental procedure Ti–15 to 60 at.% Pd alloys were prepared from 99.7 mass% Ti and 99.8 mass% Pd by arc melting in an argon atmosphere. The ∗
Corresponding author. Tel.: +81 96 342 3707; fax: +81 96 342 3710. E-mail address: [email protected] (T. Yamamuro).
0921-5093/$ – see front matter © 2006 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2006.01.120
ingots were homogenized at 1273 K for 36 ks in vacuum. Rods of 3 mm in diameter were spark cut from the ingots. Disks of 1.0 mm in thickness for differential scanning calorimetry (DSC) were cut from the rods. They were solution-treated in vacuum at 1273 K for 3.6 ks, and then quenched into ice water. Some of the disks were aged from 873 to 1273 K with 100 K intervals for 36–360 ks. Diffusion couples with combinations of appropriate alloys were prepared by partial arc melting technique. They were annealed from 948 to 1273 K for 36–1800 ks, and then quenched into ice water. Concentration–penetration profiles in the diffusion zone were obtained by electron probe microanalysis. Transmission electron microscopy (TEM) was also applied to the identification of diffusion products in the couple. Microstructure examination of some alloys after heat treatment was carried out by scanning electron microscope (SEM). 3. Results and discussion Fig. 1 shows the summary of experimental results plotted at the Ti-rich portion in Ti–Pd binary phase diagrams assessed by Okamoto [2]. Major modifications are summarized as follows. The phase boundary of TiPd compound in Pd side is almost constant about 50.5 at.% Pd. On the other hand, that of in Ti side extends from 50.6 at.% Ti at around 873 K to about 51.8 at.% Ti at 1073 K. There is A2/B2 order–disorder transformation as
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Fig. 1. Summary of experimental results plotted at the Ti-rich portion in Ti–Pd binary phase diagrams assessed by Okamoto [2]. Open circles are obtained from diffusion couple experiments and closed ones are deduced from transformation temperature of quenched and aged specimens measured by DSC.
plotted by the broken line around 30 at.% Pd above 1180 K, instead of the two phase region consisting of -Ti and TiPd. The other new finding is that Ti3 Pd compound is formed below 963 ± 10 K with peritectoid reaction between -Ti and Ti2 Pd. The evidence of these modifications is discussed below. The homogeneity range of TiPd compound is estimated from transformation behaviors of quenched and aged near-equiatomic alloys. Fig. 2 shows two typical DSC curves of quenched alloys after solution treatment. In the Ti–45 at.% Pd alloy, two endothermic and exothermic peaks are clearly observed upon heating and cooling. In other words, there is successive transformation in the Ti-rich alloys. We define the first and second peaks on heating as A2* and A1* , respectively. In the same way, the first and second peaks on cooling are denoted as M1* and M2* . The inversion of A2* and M2* is clearly seen. However, it must be also noted that A1* lies very close to the reverse transformation peak temperature of the Ti–50 at.% Pd alloy. In the Ti–50 at.% Pd alloy, only a single sharp peak is observed upon heating and cooling and there is no inversion between reverse and forward transformations. The origin of these phenomena has been described in our previous reports [3,4] and is briefly summarized as follows. There is no decomposition in the Ti–50 at.% Pd alloy, during the DSC measurement. On the other hand, the precipitation of fine Ti2 Pd particles takes place in the Ti-rich alloys during the DSC measurement. Therefore, the peak of A1* in the DSC heating curve represents the reverse martensitic transformation of the TiPd matrix. The peak temperature increases with increasing Pd content and reaches a maximum at 50 at.% Pd as shown later in Fig. 3. In other word, the first peak overlaps the second peak at 50 at.% Pd. The peak of A2* is due to the reverse martensitic transformation in the local neighboring areas around the precipitates in which the Pd concentration is considered to be equiatomic, since the peak temperature is almost constant with Pd content and is identical to that of the Ti–50 at.% Pd alloy. Fig. 3(a) shows the compositional dependence of transformation peak temperatures during heating process in the quenched and furnace cooled alloys. There is no difference of transformation behavior and temperature in the quenched and the furnace
Fig. 2. Typical DSC curves of near-equiatomic Ti Pd alloys quenched from 1273 K.
cooled Ti–50 at.% Pd alloy. In the furnace cooled Ti-rich alloys, only one peak was observed in heating curve. The temperature of this peak for all the compositions lies very close to the peak temperature of Ti–50 at.% Pd alloy. On the other hand, two peaks appear on the heating curves of quenched Ti-rich alloys. The temperature of A2* increases with increasing Pd content and reaches the maximum at 50 at.% Pd. The temperature of A1* is almost constant with Pd content and is identical to the transformation peak temperatures A* of the furnace cooled specimens. The change of transformation behavior with Pd content and thermal history suggests that the phase boundary of TiPd compound in Ti-rich side has significant variation with temperature and concentration. Fig. 3(b) shows the aging time and temperature dependence of transformation peak temperatures in the Ti–47 at.% Pd alloys. Here, we plot A1* and A2* in the specimens aged at various temperature for 3.6 ks and relatively prolonged period. The relatively prolonged period is defined as the aging time by which the transformation behavior has become stable. All the specimens aged for 3.6 ks show two endothermic peaks independently of aging temperature as described above. In this case, the A1* presents no significant variation with the aging temperature, while A2* firstly increases and then decreases with decreasing aging temperature below 1073 K. In the specimens aged above 1123 K for 3.6 ks, the A2* shows almost constant temperature. On the other hand, after relatively long time aging
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Fig. 3. (a) Compositional dependence of transformation peak temperatures during heating in near-equiatomic Ti Pd alloys quenched from 1273 K. (b) Effect of aging temperature on the transformation peaks in Ti–47 at.% Pd.
only one endothermic peak appears in the specimen aged below 1023 K. Although the A1* and A2* in the specimens aged above 1073 K for prolonged period are slightly higher than those for 3.6 ks, the same successive transformation takes place essentially. As apparently the compositional and/or structural changes may occur during aging in the specimens aged below 1023 K. On the other hand, above 1073 K the phase equilibrium is completed within aging for 3.6 ks and/or no diffusional process, such as decomposition occurs during aging. The slight increase of A1* and A2* in the specimen aged above 1073 K for prolonged period is probably due to the decrease of Ti concentration with evaporation and/or oxidation of Ti during the aging. Comparing Fig. 3(a) and (b), we can estimate roughly the homogeneity range of TiPd compound. The A2* of Ti–47 at.% Pd alloy aged at 873 K for prolonged period in Fig. 3(b) is about 853 K which can be extrapolated to the A2* of Ti–49.4 at.% Pd in Fig. 3(a). This indicates that the matrix composition of the specimen in equilibrium with Ti2 Pd phase is about Ti–49.4 at.% Pd at 873 K. In the same way, the equilibrium compositions of matrix are estimated to be Ti–49.0, 48.7 and 48.2 at.% Pd at 973, 1023 and 1073 K, respectively. These compositions are plotted with the closed circles in Fig. 1. It is also noted that the phase boundary of TiPd single and TiPd + Ti2 Pd duplex phases in the Ti–47 at.% Pd alloy is considered to be temperature range between 1073 and 1123 K. These results indicate that the phase boundary TiPd compound in Ti side extends from near-equiatomic composition about 50.6 at.% Ti at around 873 K to about 51.8 at.% Ti at 1073 K, which is confirmed by following diffusion couple experiments. The concentration–penetration curves for Pd in the vicinity of the diffusion zone in diffusion couples examined are divided into three types as shown in Fig. 4. Fig. 4(a) shows the curve in Ti/Ti–50 at.% Pd couple annealed at 1273 K for 36 ks. Instead of concentration gap corresponding to boundary between -Ti and TiPd compound, there is a clear singularity in the profile of Pd near 30 at.% Ti, which suggests that the critical A2/B2 ordering boundary is located at this concentration [5]. The profile of Pd in Fig. 4(b) is obtained from
Ti/Ti–46 at.% Pd couple annealed at 998 K for 1440 ks. There are three concentration gaps corresponding to ␣-Ti/-Ti, -Ti/Ti2 Pd and Ti2 Pd/TiPd boundaries. Fig. 4(c) shows the profile of Pd in Ti/Ti–46 at.% Pd couple annealed at 948 K for 3600 ks. There are three concentration gaps corresponding to ␣-Ti/-Ti, -Ti/Ti3 Pd and Ti3 Pd/Ti2 Pd boundaries. Although the Ti2 Pd/TiPd boundary was also observed in the profile, we omit it because of space limitation. In order to confirm the formation of Ti3 Pd, TEM observations are carried out in Ti/Ti–46 at.% Pd couple annealed at 948 K for 1800 ks. Fig. 5(a) shows a bright field image of the diffusion product between -Ti and Ti2 Pd in the same diffusion couple in Fig. 4(c). Electron diffraction patterns in Fig. 5(b)–(d) are taken from areas B, C and D in (a).
Fig. 4. Three typical concentration profile in diffusion couple experiments. (a) Ti/Ti–50 at.% Pd couple annealed at 1273 K for 36 ks, (b) Ti/Ti–46 at.% Pd couple annealed at 998 K for 1440 ks and (c) Ti/Ti–46 at.% Pd couple annealed at 948 K for 3600 ks.
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Fig. 5. (a) Bright field image, (b–d) electron diffraction patterns taken from areas B, C and D in (a), showing Ti3 Pd compound formed in pure Ti/Ti–46 at.% Pd diffusion couple annealed at 948 K for 1800 ks.
These patterns can be indexed consistently with -Ti of bcc, Ti3 Pd of A15 and Ti2 Pd of C11b structures, respectively. It is inevitably that there is Ti3 Pd phase between -Ti and Ti2 Pd in the diffusion couple of Ti/Ti–46 at.% Pd annealed at 948 K. However, the formation temperature range of Ti3 Pd is higher than that in the currently available phase diagram in Fig. 1 [2,6], which will be discussed later. All the results obtained by
the diffusion couple experiments are plotted by open circle in Fig. 1. In order to confirm the formation temperature and process of the Ti3 Pd compound, Ti–25 at.% Pd alloy was heat-treated from 948 to 998 K with 5 K interval for 360 ks. The Ti3 Pd could be detected with SEM scale in the specimen aged below 963 K. There is no Ti3 Pd in the specimens aged above 968 K. The for-
Fig. 6. SEM micrographs showing the formation of Ti3 Pd compound with peritectoid reaction between -Ti and Ti2 Pd compound in Ti–25 at.% Pd alloy aged at (a) 973 K for 360 ks, (b–d) 948 K for 230.4, 360 and 518.4 ks, respectively.
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mation process of Ti3 Pd is demonstrated by SEM micrographs in Fig. 6. There is duplex structure consisting of -Ti and Ti2 Pd in the specimen aged at 973 K for 360 ks in Fig. 6(a), which is the initial microstructure of following experiments. Fig. 6(b)–(d) show the microstructure changes in the specimens aged 948 K for 230.4, 360 and 518.4 ks after those aged at 973 K for 360 ks, respectively. The interface reaction is clearly seen at the boundary between -Ti and Ti2 Pd. The amount of the reaction product increases with increasing the aging period. This is a typical microstructure change upon peritectoid reaction. We can conclude that the Ti3 Pd compound is formed by the peritectoid reaction between -Ti and Ti2 Pd at 963 ± 5 K.
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firmed that Ti3 Pd compound is formed with peritectoid reaction between -Ti and Ti2 Pd at 963 ± 5 K. From these results, we have proposed the new phase diagram of Ti Pd binary system. Acknowledgements This work was supported by Grant-in-Aid for Exploratory Research from the Ministry of Education, Culture, Sports, Science and Technology and Grant-in-Aid for Scientific Research (B) from Japan Science Promotion Society, COE-directing Research Program B at Kumamoto University on Nano-space Electrochemistry and Special Program at Kumamoto University for Promoting Research on Nano-space Electrochemistry.
4. Conclusions The phase diagram of near-equiatomic and Ti-rich side of Ti Pd binary system has been reinvestigated. The phase boundary of TiPd compound in Pd side is almost constant about 50.5 at.% Pd. On the other hand, that of in Ti side extends from 50.6 at.% Ti at around 873 K to about 51.8 at.% Ti at 1073 K. From the diffusion couple experiments, it is recognized that there is A2(bcc) to B2 order–disorder transformation around 30 at.% Pd, instead of -Ti and TiPd duplex phase region. It is also con-
References [1] H.C. Donkersloot, J.H.N. Van Vucht, J. Less-Common Met. 20 (1970) 83–91. [2] H. Okamoto, J. Phase Equilib. 14 (1993) 1–9. [3] V.C. Solomon, M. Nishida, Mater. Trans. 43 (2002) 897–901. [4] V.C. Solomon, M. Nishida, Mater. Trans. 43 (2002) 908–915. [5] R. Kainuma, I. Ohnuma, K. Ishikawa, K. Ishida, Intermetallics 8 (2000) 869–875. [6] A.F. Jankowski, J. Alloys Compd. 182 (1992) 35–42.