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Effects of rapid solidification on the phase structures and electrochemical properties of a V3TiNi0.56Co0.14Nb0.047Ta0.047 alloy
Journal of Alloys and Compounds 370 (2004) 321–325
Effects of rapid solidification on the phase structures and electrochemical properties of a V3TiNi0.56Co0.14 Nb0.047Ta0.047 alloy Q.A. Zhang a,∗ , Y.Q. Lei b a
School of Materials Science and Engineering, Anhui University of Technology, Maanshan, Anhui 243002, PR China b Department of Materials Science and Engineering, Zhejiang University, Hangzhou 310027, PR China Received 29 August 2003; received in revised form 22 September 2003; accepted 22 September 2003
Abstract The crystal structures and electrochemical properties of a V3 TiNi0.56 Co0.14 Nb0.047 Ta0.047 alloy prepared by the melt-spinning method were studied. The rapidly solidified alloy consists of a major Vanadium-based solid solution phase and a secondary Ti2 Ni-based phase. The secondary phase exists between the dendritic arms of the major phase, showing a typical eutectic characteristic. With increasing cooling rate, the amount of the secondary phase decreases but the lattice parameters of both the major phase and the secondary phase show no obvious change. Because the disappearance of the catalytic effect of the secondary phase, the rapidly solidified alloy shows bad electrochemical properties except cycle stability. © 2003 Elsevier B.V. All rights reserved. Keywords: Hydrogen absorbing materials; Rapid solidification; Crystal structure; Electrochemical properties
1. Introduction Vanadium-based solid solution alloys with a secondary phase network have been developed as negative electrode materials in nickel/metal hydride (Ni/MH) batteries [1–6]. In these alloys, the major phase working as a hydrogen storage phase is a (Ti, V) solid solution with body centered cubic (bcc) structure, while the secondary phase improving the reaction rate as a catalyst is a TiNi-based phase with bcc structure or a C14-type Laves phase with hexagonal structure [4–6]. These electrodes exhibit a large discharge capacity, but their cycle life should be further improved. Rapid solidification was believed as an effective method to improve the cycle stability without the decrease of discharge capacity in AB5 - and AB2 -type hydrogen storage alloys [7–9]. The improvement of the cycle life is attributed to the special microstructures in the rapidly solidified alloys. In our previous work, we studied a rapidly solidified vanadium-based solid solution alloy V3 TiNi0.56 Hf0.24 Mn0.15 Cr0.1 with a secondary C14 phase [10]. It was found that the cycle life of the rapidly solidified ∗ Corresponding author. Tel.: +86-555-2400615; fax: +86-555-2471263. E-mail address: [email protected] (Q.A. Zhang).
0925-8388/$ – see front matter © 2003 Elsevier B.V. All rights reserved. doi:10.1016/j.jallcom.2003.09.115
alloy was improved but its discharge capacity decreased remarkably. The decrease of discharge capacity was caused by the decrease of the amount of the secondary C14 phase and the degradation of the secondary phase network as an electro-catalyst. The aim of the present study is to understand the effects of rapid solidification on the phase structures and electrode properties of the V-based solid solution alloy with a secondary TiNi-based phase. As reported by Tsukahara [11], the V3 TiNi0.56 Co0.14 Nb0.047 Ta0.047 alloy, consisting of the major V-based solid solution phase and the secondary TiNi-based phase, showed good electrode properties. Therefore, we choose this alloy to study the effects of rapid solidification on the crystal structures and electrochemical properties.
2. Experimental details The V3 TiNi0.56 Co0.14 Nb0.047 Ta0.047 alloy was prepared by vacuum magnetic induction melting of appropriate amounts of pure metals. The sample was remelted and then rapidly solidified at different cooling rates using a melt-spinning machine with a single Mo wheel. The ribbon thicknesses were measured and thus the cooling rates were
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Table 1 Measured thicknesses of melt spun ribbons and estimated cooling rates Sample
Thickness (m)
Cooling rate (106 K s−1 )
Sample 1 Sample 2 Sample 3
41 32 26
1.49 2.44 3.70
estimated by the following formula [12]: R = Ad−2 where R is the cooling rate (K s−1 ), d the ribbon thickness (m) and A a constant which was determined to be 2.5 × 10−3 m2 K s−1 . The cooling rates of the ribbons are shown in Table 1. The microstructures were examined using transmission electron microscopy (TEM). The foils for TEM were prepared by twin-jet electropolishing the discs of Ø 3 mm in a solution containing 15 ml perchloric acid, 90 ml n-butanol and 150 ml methanol. The ribbon samples were also pulverized by hydrogenation and then crushed mechanically into powders of 300 mesh. To evaluate the phase structures of the samples, XRD measurements were performed using a Rigaku D/MAX-3B diffractometer with Cu K␣ radiation at 40 kV and 30 mA. The XRD profiles were analysed with the Rietveld refinement program RIETAN-2000 [13]. All test electrodes were prepared by mixing alloy powders with copper powder in a weight rate of 1:2 and then cold-pressing the mixture into pellets. The characteristics of the test electrodes were measured at 298 K by using a sintered Ni(OH)2 /NiOOH positive counter electrode and a Hg/HgO (6 N KOH) reference electrode. The test electrodes were charged at 100 mA g−1 for 5 h and discharged at different currents to the cut-off potential of −0.6 V versus Hg/HgO reference electrode.
3. Results and discussion Fig. 1a shows the result of the Rietveld analysis for sample 1. It can be seen that this sample consists of a major V-based solid solution phase with the bcc structure and a secondary Ti2 Ni-based phase with the fcc structure. This result is quite different from the fact that a secondary
TiNi-based phase can be formed in the conventional as-cast V3 TiNi0.56 Co0.14 Nb0.047 Ta0.047 alloy [11,14]. This indicates that the secondary phase changed from TiNi-based phase to Ti2 Ni-based phase due to rapid solidification. For the Rietveld refinement, the structure models for the major phase and the secondary phase were taken from the reported structures for V [15] and Ti2 Ni [16], respectively. As shown in the figure, the diffraction pattern calculated from the structure models is in good agreement with that measured. The abundance of the secondary phase was calculated to be 7 wt.%. Fig. 2a presents a transmission electron micrograph of sample 1. We can see that the secondary phase exists along the grain boundaries of the major phase. Fig. 2b is the magnified image showing the microstructure of the boundaries. A fine dendritic morphology can be seen clearly and the secondary Ti2 Ni-based phase exists between the dendritic arms of the major phase. This microstructure shows a typical eutectic characteristic. As reported by Boettinger, eutectic microstructures often occur in rapidly solidified samples at compositions other than the thermodynamic eutectic because of competitive growth of two phases [17]. From the Fig. 2a, we believe that the V-based solid solution phase is first formed at the initial solidification stage. The nickel content of the primary solid solution phase must be over-saturated due to rapid solidification, leading to a decrease of the nickel content in the remaining liquid. Thus, the solid solution phase and the Ti2 Ni-based phase cooperatively grow at the final solidification stage. Fig. 1b shows the result of the Rietveld analysis for sample 2. It can be seen that this sample also consists of the major V-based solid solution phase and the secondary Ti2 Ni-based phase. A similar result is further found in sample 3, as shown in Fig. 1c. This means that rapid solidification leads to the formation of the secondary Ti2 Ni-based phase instead of TiNi phase, even though the cooling rates of these rapidly solidified samples are different. Table 2 summarizes the lattice parameters and phase abundance of the rapidly solidified samples. It can be seen that the lattice parameters of both the major phase and the secondary phase have no obvious change at different rapid solidification processes. However, the amount of the secondary Ti2 Ni-based phase decreases with the increase of the cooling rates, which is related to an extension of the solid solubility limits of these elements in the V-based solid solution phase.
Table 2 Structural parameters and phase abundance of the rapidly solidified samples refined by the X-ray Rietveld analysis Alloy
Phase
Space group
RI (%)
Lattice parameters a = b = c (Å)
Abundance (wt.%)
Sample 1 Rwp = 8.85%, S = 1.78
Major phase Secondary phase
¯ Im3m ¯ Fd3m
1.62 2.75
3.042(1) 11.260(2)
93 7
Sample 2 Rwp = 9.23%, S = 1.97
Major phase Secondary phase
¯ Im3m ¯ Fd3m
1.14 3.32
3.044(1) 11.268(2)
95 5
Sample 3 Rwp = 9.07%, S = 1.92
Major phase Secondary phase
¯ Im3m ¯ Fd3m
1.21 2.87
3.044(1) 11.266(2)
97 3
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323
Fig. 1. Calculated (line) and observed (+) X-ray diffraction patterns for (a) sample 1 prepared by the melt-spinning method at a cooling rate of 1.49 × 106 K s−1 ; (b) sample 2 prepared by the melt-spinning method at a cooling rate of 2.44 × 106 K s−1 ; and (c) sample 3 prepared by the melt-spinning method at a cooling rate of 3.70 × 106 K s−1 . Reflection markers are for the V-based solid solution phase and the Ti2 Ni-based phase, respectively.
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Fig. 2. (a) Transmission electron micrograph of sample 1 prepared by the melt-spinning method at a cooling rate of 1.49 × 106 K s−1 ; and (b) magnified image showing the microstructure of the boundaries.
Fig. 3 shows the discharge capacity versus charging– discharging cycle number for the alloy electrodes at a discharge current of 50 mA g−1 . It can be seen that the rapidly solidified alloy electrodes have a poorer activation property than the conventional as-cast one [11,14]. This result is in good agreement with the fact that a rapidly solidified vanadium-based solid solution alloy V3 TiNi0.56 Hf0.24 Mn0.15 Cr0.1 with a secondary C14 phase shows poor activation behavior [10]. This is related to the small grains in the rapidly solidified alloy because the small grains can suppress pulverization of the alloy powder during charge-discharge cycling [18]. However, the lower pulverization rate of the rapidly solidified alloy powder contributes to the better cycle stability. Compared with the conventional as-cast V3 TiNi0.56 Co0.14 Nb0.047 Ta0.047 alloy [11,14], we found that the cycle stability of the rapidly solidified alloy was improved further. It should be noted that the discharge capacity of the rapidly solidified V3 TiNi0.56 Co0.14 Nb0.047 Ta0.047 alloy decreases remarkably. Fig. 4 shows the discharge capacity of
Fig. 3. Discharge capacity vs. cycle number for the rapidly solidified alloy electrodes (charged at 100 mA g−1 for 5 h, discharged to −0.6 V at 50 mA g−1 ).
the rapidly solidified alloy electrodes at different discharge currents. We can see that the discharge capacity at a discharge current of 10 mA g−1 is about 270 mAh g−1 , which is much lower than that of the conventional as-cast alloy electrode (360 mAh g−1 [14]). Moreover, the high-rate dischargeability (C100 /C10 × 100%) of the rapidly solidified alloy electrodes also becomes worse. Actually, the low discharge capacity and the bad high-rate dischargeability are predominantly ascribed to the secondary Ti2 Ni-based phase. In the conventional as-cast V3 TiNi0.56 Co0.14 Nb0.047 Ta0.047 alloy, the major solid solution phase works as a hydrogen storage phase while the secondary TiNi-based phase improves the reaction rate as a catalyst [11]. In the present case, however, the secondary phase is Ti2 Ni-based phase and the phase abundance is too low. Thus, the catalytic effect of the secondary phase almost disappears in the rapidly solidified V3 TiNi0.56 Co0.14 Nb0.047 Ta0.047 alloy. From Figs. 3 and 4, we can also see that the cooling rate of rapid solidification under the present experimental condition has no large effect on the electrochemical properties. All these samples show bad electrochemical properties except
Fig. 4. Discharge capacities for the rapidly solidified alloy electrodes at different discharge currents.
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cycle stability. This implies that rapid solidification with a higher or lower cooling rate leads to the disappearance of the catalytic effect of the secondary phase, which is similar to that occurring in the V-based solid solution alloys with a secondary C14 phase [10]. Therefore, we believe that rapid solidification is not suitable for preparing the V-based solid solution alloys with a secondary catalytic phase network as negative electrode materials in Ni/MH batteries. 4. Conclusions The rapidly solidified V3 TiNi0.56 Co0.14 Nb0.047 Ta0.047 alloy consists of a major V-based solid solution phase and a secondary Ti2 Ni-based phase. A catalytic TiNi-based phase does not exist due to rapid solidification. The secondary phase exists between the dendritic arms of the major phase, showing a typical eutectic characteristic. With increasing cooling rate, the amount of the secondary phase decreases but the lattice parameters of both the major phase and the secondary phase show no obvious change. Because the disappearance of the catalytic effect of the secondary phase, the rapidly solidified V3 TiNi0.56 Co0.14 Nb0.047 Ta0.047 alloy shows bad electrochemical properties except cycle stability. From the present results and our previous work [10], we believe that rapid solidification is not suitable for preparing the V-based solid solution alloys with a secondary catalytic phase network as negative electrode materials in Ni/MH batteries.
325
References [1] M. Tsukahara, K. Takahashi, T. Mishima, T. Sakai, H. Miyamura, N. Kuriyama, I. Uehara, J. Alloys Compd. 224 (1995) 162. [2] M. Tsukahara, K. Takahashi, T. Mishima, T. Sakai, H. Miyamura, N. Kuriyama, I. Uehara, J. Alloys Compd. 226 (1995) 203. [3] M. Tsukahara, K. Takahashi, T. Mishima, A. Isomura, T. Sakai, J. Alloys Compd. 243 (1996) 133. [4] M. Tsukahara, K. Takahashi, T. Mishima, A. Isomura, T. Sakai, J. Alloys Compd. 245 (1996) 59. [5] M. Tsukahara, K. Takahashi, T. Mishima, A. Isomura, T. Sakai, J. Alloys Compd. 236 (1996) 151. [6] M. Tsukahara, K. Takahashi, T. Mishima, A. Isomura, T. Sakai, J. Alloys Compd. 253 (1997) 583. [7] R. Mishima, H. Miyamura, T. Sakai, N. Kuriyama, H. Ishikawa, I. Uehara, J. Alloys Compd. 192 (1993) 176. [8] N. Higashiyama, Y. Matsura, H. Nakamura, M. Kimoto, M. Nogami, I. Yonezu, K. Nishio, J. Alloys Compd. 253–254 (1997) 648. [9] K. Shu, Y. Lei, X. Yang, G. Lin, Q. Wang, J. Alloys Compd. 293–295 (1999) 756. [10] Q.A. Zhang, Y.Q. Lei, L.X. Chen, Q.D. Wang, J. Alloys Compd. 308 (2000) 104. [11] M. Tsukahara, K. Takahashi, A. Isomura, T. Sakai, J. Alloys Compd. 287 (1999) 215. [12] H. Jones, Rep. Prog. Phys. 36 (1973) 1425. [13] F. Izumi, T. Ikeda, Mater. Sci. Forum 321–324 (2000) 198. [14] Q.A. Zhang, Y.Q. Lei, L.X. Chen, Q.D. Wang, Mater. Chem. Phys. 69 (2001) 241. [15] V.A. Finkel, V.I. Glamazda, G.P. Kovtun, Soviet Phys. 30 (1970) 581. [16] G.A. Yurko, J.W. Barton, J.G. Parr, Acta Crystallogr. 12 (1959) 909. [17] W.J. Boettinger, Mater. Sci. Eng. 98 (1988) 123. [18] C. Li, X. Wang, C. Wang, J. Alloys Compd. 266 (1998) 300.
Effects of rapid solidification on the phase structures and electrochemical properties of a V3TiNi0.56Co0.14 Nb0.047Ta0.047 alloy Q.A. Zhang a,∗ , Y.Q. Lei b a
School of Materials Science and Engineering, Anhui University of Technology, Maanshan, Anhui 243002, PR China b Department of Materials Science and Engineering, Zhejiang University, Hangzhou 310027, PR China Received 29 August 2003; received in revised form 22 September 2003; accepted 22 September 2003
Abstract The crystal structures and electrochemical properties of a V3 TiNi0.56 Co0.14 Nb0.047 Ta0.047 alloy prepared by the melt-spinning method were studied. The rapidly solidified alloy consists of a major Vanadium-based solid solution phase and a secondary Ti2 Ni-based phase. The secondary phase exists between the dendritic arms of the major phase, showing a typical eutectic characteristic. With increasing cooling rate, the amount of the secondary phase decreases but the lattice parameters of both the major phase and the secondary phase show no obvious change. Because the disappearance of the catalytic effect of the secondary phase, the rapidly solidified alloy shows bad electrochemical properties except cycle stability. © 2003 Elsevier B.V. All rights reserved. Keywords: Hydrogen absorbing materials; Rapid solidification; Crystal structure; Electrochemical properties
1. Introduction Vanadium-based solid solution alloys with a secondary phase network have been developed as negative electrode materials in nickel/metal hydride (Ni/MH) batteries [1–6]. In these alloys, the major phase working as a hydrogen storage phase is a (Ti, V) solid solution with body centered cubic (bcc) structure, while the secondary phase improving the reaction rate as a catalyst is a TiNi-based phase with bcc structure or a C14-type Laves phase with hexagonal structure [4–6]. These electrodes exhibit a large discharge capacity, but their cycle life should be further improved. Rapid solidification was believed as an effective method to improve the cycle stability without the decrease of discharge capacity in AB5 - and AB2 -type hydrogen storage alloys [7–9]. The improvement of the cycle life is attributed to the special microstructures in the rapidly solidified alloys. In our previous work, we studied a rapidly solidified vanadium-based solid solution alloy V3 TiNi0.56 Hf0.24 Mn0.15 Cr0.1 with a secondary C14 phase [10]. It was found that the cycle life of the rapidly solidified ∗ Corresponding author. Tel.: +86-555-2400615; fax: +86-555-2471263. E-mail address: [email protected] (Q.A. Zhang).
0925-8388/$ – see front matter © 2003 Elsevier B.V. All rights reserved. doi:10.1016/j.jallcom.2003.09.115
alloy was improved but its discharge capacity decreased remarkably. The decrease of discharge capacity was caused by the decrease of the amount of the secondary C14 phase and the degradation of the secondary phase network as an electro-catalyst. The aim of the present study is to understand the effects of rapid solidification on the phase structures and electrode properties of the V-based solid solution alloy with a secondary TiNi-based phase. As reported by Tsukahara [11], the V3 TiNi0.56 Co0.14 Nb0.047 Ta0.047 alloy, consisting of the major V-based solid solution phase and the secondary TiNi-based phase, showed good electrode properties. Therefore, we choose this alloy to study the effects of rapid solidification on the crystal structures and electrochemical properties.
2. Experimental details The V3 TiNi0.56 Co0.14 Nb0.047 Ta0.047 alloy was prepared by vacuum magnetic induction melting of appropriate amounts of pure metals. The sample was remelted and then rapidly solidified at different cooling rates using a melt-spinning machine with a single Mo wheel. The ribbon thicknesses were measured and thus the cooling rates were
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Table 1 Measured thicknesses of melt spun ribbons and estimated cooling rates Sample
Thickness (m)
Cooling rate (106 K s−1 )
Sample 1 Sample 2 Sample 3
41 32 26
1.49 2.44 3.70
estimated by the following formula [12]: R = Ad−2 where R is the cooling rate (K s−1 ), d the ribbon thickness (m) and A a constant which was determined to be 2.5 × 10−3 m2 K s−1 . The cooling rates of the ribbons are shown in Table 1. The microstructures were examined using transmission electron microscopy (TEM). The foils for TEM were prepared by twin-jet electropolishing the discs of Ø 3 mm in a solution containing 15 ml perchloric acid, 90 ml n-butanol and 150 ml methanol. The ribbon samples were also pulverized by hydrogenation and then crushed mechanically into powders of 300 mesh. To evaluate the phase structures of the samples, XRD measurements were performed using a Rigaku D/MAX-3B diffractometer with Cu K␣ radiation at 40 kV and 30 mA. The XRD profiles were analysed with the Rietveld refinement program RIETAN-2000 [13]. All test electrodes were prepared by mixing alloy powders with copper powder in a weight rate of 1:2 and then cold-pressing the mixture into pellets. The characteristics of the test electrodes were measured at 298 K by using a sintered Ni(OH)2 /NiOOH positive counter electrode and a Hg/HgO (6 N KOH) reference electrode. The test electrodes were charged at 100 mA g−1 for 5 h and discharged at different currents to the cut-off potential of −0.6 V versus Hg/HgO reference electrode.
3. Results and discussion Fig. 1a shows the result of the Rietveld analysis for sample 1. It can be seen that this sample consists of a major V-based solid solution phase with the bcc structure and a secondary Ti2 Ni-based phase with the fcc structure. This result is quite different from the fact that a secondary
TiNi-based phase can be formed in the conventional as-cast V3 TiNi0.56 Co0.14 Nb0.047 Ta0.047 alloy [11,14]. This indicates that the secondary phase changed from TiNi-based phase to Ti2 Ni-based phase due to rapid solidification. For the Rietveld refinement, the structure models for the major phase and the secondary phase were taken from the reported structures for V [15] and Ti2 Ni [16], respectively. As shown in the figure, the diffraction pattern calculated from the structure models is in good agreement with that measured. The abundance of the secondary phase was calculated to be 7 wt.%. Fig. 2a presents a transmission electron micrograph of sample 1. We can see that the secondary phase exists along the grain boundaries of the major phase. Fig. 2b is the magnified image showing the microstructure of the boundaries. A fine dendritic morphology can be seen clearly and the secondary Ti2 Ni-based phase exists between the dendritic arms of the major phase. This microstructure shows a typical eutectic characteristic. As reported by Boettinger, eutectic microstructures often occur in rapidly solidified samples at compositions other than the thermodynamic eutectic because of competitive growth of two phases [17]. From the Fig. 2a, we believe that the V-based solid solution phase is first formed at the initial solidification stage. The nickel content of the primary solid solution phase must be over-saturated due to rapid solidification, leading to a decrease of the nickel content in the remaining liquid. Thus, the solid solution phase and the Ti2 Ni-based phase cooperatively grow at the final solidification stage. Fig. 1b shows the result of the Rietveld analysis for sample 2. It can be seen that this sample also consists of the major V-based solid solution phase and the secondary Ti2 Ni-based phase. A similar result is further found in sample 3, as shown in Fig. 1c. This means that rapid solidification leads to the formation of the secondary Ti2 Ni-based phase instead of TiNi phase, even though the cooling rates of these rapidly solidified samples are different. Table 2 summarizes the lattice parameters and phase abundance of the rapidly solidified samples. It can be seen that the lattice parameters of both the major phase and the secondary phase have no obvious change at different rapid solidification processes. However, the amount of the secondary Ti2 Ni-based phase decreases with the increase of the cooling rates, which is related to an extension of the solid solubility limits of these elements in the V-based solid solution phase.
Table 2 Structural parameters and phase abundance of the rapidly solidified samples refined by the X-ray Rietveld analysis Alloy
Phase
Space group
RI (%)
Lattice parameters a = b = c (Å)
Abundance (wt.%)
Sample 1 Rwp = 8.85%, S = 1.78
Major phase Secondary phase
¯ Im3m ¯ Fd3m
1.62 2.75
3.042(1) 11.260(2)
93 7
Sample 2 Rwp = 9.23%, S = 1.97
Major phase Secondary phase
¯ Im3m ¯ Fd3m
1.14 3.32
3.044(1) 11.268(2)
95 5
Sample 3 Rwp = 9.07%, S = 1.92
Major phase Secondary phase
¯ Im3m ¯ Fd3m
1.21 2.87
3.044(1) 11.266(2)
97 3
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323
Fig. 1. Calculated (line) and observed (+) X-ray diffraction patterns for (a) sample 1 prepared by the melt-spinning method at a cooling rate of 1.49 × 106 K s−1 ; (b) sample 2 prepared by the melt-spinning method at a cooling rate of 2.44 × 106 K s−1 ; and (c) sample 3 prepared by the melt-spinning method at a cooling rate of 3.70 × 106 K s−1 . Reflection markers are for the V-based solid solution phase and the Ti2 Ni-based phase, respectively.
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Fig. 2. (a) Transmission electron micrograph of sample 1 prepared by the melt-spinning method at a cooling rate of 1.49 × 106 K s−1 ; and (b) magnified image showing the microstructure of the boundaries.
Fig. 3 shows the discharge capacity versus charging– discharging cycle number for the alloy electrodes at a discharge current of 50 mA g−1 . It can be seen that the rapidly solidified alloy electrodes have a poorer activation property than the conventional as-cast one [11,14]. This result is in good agreement with the fact that a rapidly solidified vanadium-based solid solution alloy V3 TiNi0.56 Hf0.24 Mn0.15 Cr0.1 with a secondary C14 phase shows poor activation behavior [10]. This is related to the small grains in the rapidly solidified alloy because the small grains can suppress pulverization of the alloy powder during charge-discharge cycling [18]. However, the lower pulverization rate of the rapidly solidified alloy powder contributes to the better cycle stability. Compared with the conventional as-cast V3 TiNi0.56 Co0.14 Nb0.047 Ta0.047 alloy [11,14], we found that the cycle stability of the rapidly solidified alloy was improved further. It should be noted that the discharge capacity of the rapidly solidified V3 TiNi0.56 Co0.14 Nb0.047 Ta0.047 alloy decreases remarkably. Fig. 4 shows the discharge capacity of
Fig. 3. Discharge capacity vs. cycle number for the rapidly solidified alloy electrodes (charged at 100 mA g−1 for 5 h, discharged to −0.6 V at 50 mA g−1 ).
the rapidly solidified alloy electrodes at different discharge currents. We can see that the discharge capacity at a discharge current of 10 mA g−1 is about 270 mAh g−1 , which is much lower than that of the conventional as-cast alloy electrode (360 mAh g−1 [14]). Moreover, the high-rate dischargeability (C100 /C10 × 100%) of the rapidly solidified alloy electrodes also becomes worse. Actually, the low discharge capacity and the bad high-rate dischargeability are predominantly ascribed to the secondary Ti2 Ni-based phase. In the conventional as-cast V3 TiNi0.56 Co0.14 Nb0.047 Ta0.047 alloy, the major solid solution phase works as a hydrogen storage phase while the secondary TiNi-based phase improves the reaction rate as a catalyst [11]. In the present case, however, the secondary phase is Ti2 Ni-based phase and the phase abundance is too low. Thus, the catalytic effect of the secondary phase almost disappears in the rapidly solidified V3 TiNi0.56 Co0.14 Nb0.047 Ta0.047 alloy. From Figs. 3 and 4, we can also see that the cooling rate of rapid solidification under the present experimental condition has no large effect on the electrochemical properties. All these samples show bad electrochemical properties except
Fig. 4. Discharge capacities for the rapidly solidified alloy electrodes at different discharge currents.
Q.A. Zhang, Y.Q. Lei / Journal of Alloys and Compounds 370 (2004) 321–325
cycle stability. This implies that rapid solidification with a higher or lower cooling rate leads to the disappearance of the catalytic effect of the secondary phase, which is similar to that occurring in the V-based solid solution alloys with a secondary C14 phase [10]. Therefore, we believe that rapid solidification is not suitable for preparing the V-based solid solution alloys with a secondary catalytic phase network as negative electrode materials in Ni/MH batteries. 4. Conclusions The rapidly solidified V3 TiNi0.56 Co0.14 Nb0.047 Ta0.047 alloy consists of a major V-based solid solution phase and a secondary Ti2 Ni-based phase. A catalytic TiNi-based phase does not exist due to rapid solidification. The secondary phase exists between the dendritic arms of the major phase, showing a typical eutectic characteristic. With increasing cooling rate, the amount of the secondary phase decreases but the lattice parameters of both the major phase and the secondary phase show no obvious change. Because the disappearance of the catalytic effect of the secondary phase, the rapidly solidified V3 TiNi0.56 Co0.14 Nb0.047 Ta0.047 alloy shows bad electrochemical properties except cycle stability. From the present results and our previous work [10], we believe that rapid solidification is not suitable for preparing the V-based solid solution alloys with a secondary catalytic phase network as negative electrode materials in Ni/MH batteries.
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