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Preparation and hydrogenation of multicomponent AB2-type Zr–Mn–V–Co–Ni amorphous alloy
Journal of Alloys and Compounds 307 (2000) 266–271
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Preparation and hydrogenation of multicomponent AB 2 -type Zr–Mn–V–Co–Ni amorphous alloy L.-C. Lai, C.-L. Lee, T.-P. Perng* Department of Materials Science and Engineering, National Tsing Hua University, Hsinchu, Taiwan Received 2 March 2000; accepted 10 March 2000
Abstract Amorphous powder of the multicomponent alloy ZrMn 0.6 V0.1 Co 0.2 Ni 1.2 was prepared by ball milling of the intermetallic compound. The amorphous phase was formed after milling for 80 h. Hydrogenation of the amorphous powder was conducted and a comparison was made with that of the crystalline alloy. The absorption rate increased as the temperature increased, but the rate was slower than that of the crystalline phase. There was no plateau region in the pressure–composition–temperature curve, the maximum hydrogenation capacity was less, and a large amount of hydrogen was trapped at lower pressures. 2000 Elsevier Science S.A. All rights reserved. Keywords: Amorphous materials; Interstitial alloys; Transition metal compounds; hydrogen absorption
1. Introduction The ball-milling technique has been applied extensively to prepare Ti- and Zr-based binary amorphous alloys, e.g. Ni–Zr [1–6], Ti–Mn [7], Ti–Fe [8], Ti–Ni [9–11], Ti–Cu [12], and Fe–Zr [13]. Since both Ti and Zr are hydride formers, they are frequently used as constituent elements to prepare hydrogen storage materials. Conventional crystalline hydrogen storage materials can be classified into several categories according to their constituent elements and the type of structure, e.g. AB 5 , AB 2 , or AB 2 , where A is a hydride former and B is often a transition metal. The hydrogenation characteristics of amorphous hydrogen storage materials are very different from those of traditional crystalline materials. For example, Maeland et al. [14] reported that amorphous TiCu absorbed 35% more hydrogen than crystalline TiCu (TiCu 1.35 compared with TiCu). Recently, the characteristics of hydrogenation of other amorphous or nanocrystalline hydrogen storage alloys prepared by mechanical milling were studied, such as Mg–Ni [15,16], Ti–Fe [17,18], Ti–Mn [19], and Zr–Ni [20] alloys. Chu et al. [18] reported the formation of amorphous TiFe with some nanocrystals embedded in the matrix by mechanical alloying Ti with Fe powder. The hydrogen absorption capacity of this alloy was smaller *Corresponding author. Tel.: 1886-7-3574-2634; fax: 1886-7-35722713. E-mail address: [email protected] (T.-P. Perng).
than that of crystalline TiFe. Zaluski et al. [15] reported the hydrogen absorption properties of nanocrystalline Mg 2 Ni formed by mechanical alloying. They found that the activation for hydrogen absorption was much easier, because of the very active surface of the powders created in the ball-milling process. In addition, Zaluski et al. [17] also found that nanocrystalline TiFe showed a narrower pressure–composition–temperature (P–C–T ) plateau and a higher capacity for hydrogenation at low pressures than its crystalline counterpart. Chin et al. [19] reported that amorphous alloy TiMn could be prepared by mechanical alloying. This alloy could also absorb hydrogen without activation treatment, but no plateau region was observed in the P–C–T curve. In the Zr–Ni system, Spit et al. [20] examined the hydrogen absorption characteristics of amorphous Zr 36 Ni 64 and found that it absorbed only a small amount of hydrogen. Except for the case of TiCu alloy, the hydrogen absorption capacities of amorphous alloys are usually less than those of the corresponding crystalline alloys. Previous reports on the hydrogenation properties of amorphous alloys have mostly been focused on binary alloy systems. In studying the hydrogenation properties of AB 2 -type alloys, we have found that the multicomponent ZrMn 0.6 V0.2 Co 0.1 Ni 1.2 , derived from Zr(MnNi) 2 , could maintain the C15-type Laves phase structure and was an excellent hydrogen storage compound [21]. The hydrogen absorption capacity (H / M) was as high as 1.2, higher than that of a similar binary alloy Zr 36 Ni 64 , the capacity of
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which is only about 0.65. In addition, the hydrogenation rate was very fast, and an electrode made from this composition had a very high discharge capacity (|390 mAh / g) [22]. In this study, an amorphous powder of ZrMn 0.6 V0.2 Co 0.1 Ni 1.2 was prepared from the intermetallic alloy by high-energy ball milling. The amorphization process and the crystalline behavior were studied. The hydrogenation properties of the amorphous phase were examined and compared with those of the crystalline phase.
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conditions, hydrogen was absorbed by the powders rapidly. After repeating several activation treatments until a steadystate absorption rate had been reached, the reaction vessel was cooled to room temperature and evacuated to about 10 24 Pa. The kinetics of hydrogen absorption and P–C–T curves were then measured. For the kinetics measurement, the initial hydrogen pressure was set at 3.3 MPa.
3. Results and discussion
3.1. Mechanical milling 2. Experimental Pure elements Zr, Mn, V, Co, and Ni (purity 99.9– 99.99%) were used as starting materials. They were mixed in the stoichiometry ZrMn 0.6 V0.2 Co 0.1 Ni 1.2 and the intermetallic compound was prepared by arc melting on a water-cooled copper plate under argon atmosphere. The ingots were repeatedly turned and remelted several times to achieve homogeneity. Some ingots were sealed in a quartz tube in vacuum and then annealed at 10008C for 10 h. Both as-melted and annealed ZrMn 0.6 V0.2 Co 0.1 Ni 1.2 ingots were crushed and pulverized mechanically into powders in a glove box filled with argon. Powders between 150 and 2100 mesh were selected for hydrogenation measurements. For ball milling, powders from the asmelted alloy smaller than 100 mesh were used. Approximately 15 g of powder and eight hardened-steel balls were sealed in a hardened-steel vial under a purified argon atmosphere. The volume of the vial was 56 ml. The weight and diameter of each ball were 3.95 g and 10 mm, respectively. The milling procedure was carried out in a Spex-8000 high-energy ball mill at a shaking rate of 1200 rpm. The structures of the powder were characterized by X-ray diffraction (XRD). The morphology was examined by scanning electron microscopy (SEM). Differential scanning calorimetry (DSC) was used to analyze the crystallization behavior of the milled powder in Ar. The heating rate was set at 158C / min. In addition, three temperatures, 490, 560, and 7008C, were selected for annealing the milled powder. The samples were heated in vacuum to these temperatures for 10 min. A Sieverts-type apparatus was employed for the hydrogenation experiment. Approximately 5 g of sample powder was placed in a stainless-steel vessel for the measurement. Since the ball-milled powder was very fine and easily oxidized, it was handled in a glove box under Ar atmosphere. After milling, the vial was kept closed and moved directly to an Ar-filled glove box. The powders were transferred from the vial and sealed in the vessel for hydrogenation test. Activation of the amorphous powders was carried out first in 3.3 MPa hydrogen gas at 3008C. Under such
Ball milling of ZrMn 0.6 V0.2 Co 0.1 Ni 1.2 intermetallic compound led to refinement of the particles. The evolution of morphologies of the ZrMn 0.6 V0.2 Co 0.1 Ni 1.2 particles after various milling times is illustrated by the SEM micrographs in Fig. 1. Before milling (Fig. 1a), the alloy had particle sizes of 50–120 mm and had a typical brittle feature. After 4 h of milling (Fig. 1b), the powders were fractured by the balls to form smaller particles. As the milling proceeded to 36 h (Fig. 1c), the small particles were repeatedly fragmented and cold welded to form larger particles. After 80 h of milling (Fig. 1d), the particles became spherical and the size became more uniform, with an average particle size of about 20 mm diameter. Fig. 2 presents the XRD patterns of the ZrMn 0.6 V0.2 Co 0.1 Ni 1.2 powders after various milling times. The XRD pattern for the starting powder showed a typical C15 intermetallic compound. After 4 h of milling, the peaks broadened and the intensities decreased, which implied a decrease of the grain size and an increase of chemical disorder of the particles. With increased milling time, the intensity of the peaks was further reduced. After 36 h of milling, some peaks became nearly invisible. When milling was extended to 80 h, only one broad peak was observed, indicating that an amorphous structure had been formed. In this multicomponent intermetallic, the atomic sizes of the constituent elements are different from each other. The free energy of the compound after milling was raised due to accumulation of lattice defects or dislocations. The formation sequence of the amorphous phase was similar to that of binary intermetallics by mechanical milling.
3.2. DSC analysis and thermal treatment From the above XRD observation, it was believed that amorphous alloy ZrMn 0.6 V0.2 Co 0.1 Ni 1.2 was formed by ball milling. The crystallization behavior of the amorphous powder was examined by DSC. Fig. 3 shows the DSC curve of the powder that was ball milled for 80 h. A broad exothermic peak ranging from 420 to 5608C with the maximum at 5118C was observed. In order to identify the phase change during heating, heat treatment of the powder at 490, 560, and 7008C for 10 min was performed. The
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Fig. 1. SEM morphologies of ZrMn 0.6 V0.2 Co 0.1 Ni 1.2 powder before and after various ball milling times. (a) Before milling, (b) 4 h, (c) 36 h, and (d) 80 h.
XRD patterns are compared in Fig. 4. Only the crystalline phase was formed, except that the crystallinity improved as the temperature increased. This illustrates that the peak was due to crystallization of the amorphous phase.
3.3. Hydrogenation kinetics The sample that had been ball milled for 80 h was selected for the hydrogenation measurement. The activation was conducted by several hydriding–dehydriding cycles at 3008C in 3.3 MPa of hydrogen gas. If the sample had been exposed to air before the activation treatment, it was found that the activation process became much more difficult, presumably because the particles were too small and easy to oxidize. After the activation treatment, the hydrogen absorption kinetics curves at 40, 80, and 3008C were measured in sequence. Fig. 5 shows the hydrogen absorption rates of amorphous powder at these three temperatures. The absorption rates were all reasonably fast. As the temperature was raised, the reaction rate became even faster. Without exposure to air, the sample could maintain a clean surface for a high hydrogenation rate.
From Fig. 5a it was also found that, as the temperature increased, the steady-state hydrogen pressure became higher, i.e. the hydrogen absorption capacity decreased. Since the enthalpy of formation for the hydride is negative, i.e. hydrogenation is an exothermic reaction, it is reasonable to expect a reduced hydrogen content as the temperature is raised. From Fig. 5b, the times for amorphous ZrMn 0.6 V0.2 Co 0.1 Ni 1.2 to reach 90% of its maximum capacity were 40, 20, and 3 min at 40, 80, and 3008C, respectively. For the crystalline counterpart, however, it has been observed that, in less than 5 min at 258C, the alloy could absorb 90% of its maximum capacity [21]. Therefore, the reaction rates of the amorphous powder were much slower than those of the crystalline sample. There might be several reasons for this phenomenon. In crystalline compounds it has been reported that hydrogen diffusivities in bcc or fcc metals range from 10 26 to 10 211 cm 2 / s at room temperature [23]. In the AB 2 -type Laves phase structure, the interstitial sizes are larger than those of pure bcc or fcc metals [24]. Hydrogen diffusivity in the AB 2 structure would be faster. In amorphous structures,
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Fig. 2. XRD patterns of ZrMn 0.6 V0.2 Co 0.1 Ni 1.2 powder before and after various ball milling times. (a) Before milling, (b) 4 h, (c) 16 h, (d) 36 h, and (e) 80 h.
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Fig. 4. XRD patterns of ZrMn 0.6 V0.2 Co 0.1 Ni 1.2 powder after heating at various temperatures for 10 min. (a) Before heating, (b) 4908C, (c) 5608C, and (d) 7008C.
hydrogen diffusivity is more like that in fcc metals [25]. Therefore, it is expected that hydrogen diffusivity in amorphous ZrMn 0.6 V0.2 Co 0.1 Ni 1.2 would be slower than that in the crystalline phase. In addition, in the crystalline compound, fresh surface is created after activation. The fresh surface would also accelerate the hydrogen absorp-
Fig. 3. DSC curve of amorphous ZrMn 0.6 V0.2 Co 0.1 Ni 1.2 .
Fig. 5. Kinetics curves of hydrogen absorption in amorphous ZrMn 0.6 V0.2 Co 0.1 Ni 1.2 alloy after activation. (a) Pressure change, (b) hydrogen content of the alloy. H S , saturation capacity.
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tion process. For the amorphous powders, they were already too small to crack.
3.4. P–C–T curves Fig. 6 shows the P–C–T curves for the ZrMn 0.6 V0.2 Co 0.1 Ni 1.2 amorphous alloy at 40 and 808C. Three curves for Zr 36 Ni 64 amorphous alloy from Spit et al. are also included for comparison [20]. There is no plateau region in all curves, and the hydrogen absorption capacity decreases as the temperature increases. At low pressures, there is a great deal of hydrogen trapped in both amorphous alloys. This could be due to the larger holes in the amorphous structure which stabilize the dissolved hydrogen atoms. Under the same pressure, ZrMn 0.6 V0.2 Co 0.1 Ni 1.2 absorbs more hydrogen than Zr 36 Ni 64 , but the amount of trapped hydrogen is also higher. The P–C–T curves of the amorphous ZrMn 0.6 V0.2 Co 0.1 Ni 1.2 alloy were also compared with those of the as-melted (AM) and homogenized (H) crystalline alloys at 40 and 808C (Fig. 7). For the crystalline alloys, they both exhibited typical plateau regions in the P–C–T curves. But in the as-melted state, the P–C–T curves had more sloped plateaus, and the maximum hydrogenation capacities were smaller. At low pressures, more hydrogen was trapped in the as-melted alloy than in the annealed alloy. Compared with the crystalline alloys, the amorphous alloy showed very different behavior in the P–C–T curves [17–20,26]. Except for the case of TiCu [14], the maximum hydrogenation capacity in the amor-
Fig. 7. P–C–T curves of ZrMn 0.6 V0.2 Co 0.1 Ni 1.2 . H, homogenized; AM, as-melted; A, amorphous.
phous phase is always not larger, and is usually lower, than that in the corresponding crystalline phase. The P–C–T curves of crystalline hydrogen storage alloys usually have three phase regions, i.e. a, a1b, and b regions. In the a region, hydrogen exists in the metal as a solid solution. The solubility of hydrogen in the metal increases as the hydrogen pressure increases. When it approaches the solubility limit, phase transformation from the solid solution to hydride occurs. The coexistence of two phases (a1b) is indicated by the generation of a plateau pressure because of equal chemical potentials. For an ideal hydride former, a flat plateau slope is observed. When segregation in the alloy occurs, the plateau slope becomes steeper, as seen for the as-melted alloy. Annealing eliminates segregation to yield a more homogeneous phase. Therefore, flatter plateau regions are observed for the annealed alloy. On the contrary, in amorphous alloys, hydrogenation does not result in phase separation. Hydrogen is probably dissolved to form an ‘interstitial’ solid solution, yielding no plateau region. Some hydrogen may even preferentially form a stable hydride with the constituent elements, resulting in more irreversibly trapped hydrogen after desorption.
4. Conclusions
Fig. 6. Comparison of P–C–T curves for amorphous ZrMn 0.6 V0.2 Co 0.1 Ni 1.2 and Zr 36 Ni 64 [20] at different temperatures.
1. The multicomponent amorphous ZrMn 0.6 V0.2 Co 0.1 Ni 1.2 alloy could be prepared by mechanical milling the crystalline intermetallic compound for over 80 h. 2. The crystallization behavior of amorphous
L.-C. Lai et al. / Journal of Alloys and Compounds 307 (2000) 266 – 271
ZrMn 0.6 V0.2 Co 0.1 Ni 1.2 was analyzed by DSC. A broad exothermic peak with a maximum at 5118C was observed. After heating, a C15-type crystalline phase of the same alloy was formed. 3. The hydrogen absorption rate of the amorphous alloy after activation treatment increased as the temperature increased, but the reaction rate was slower than that of the crystalline phase. 4. The hydrogen absorption capacity of amorphous ZrMn 0.6 V0.2 Co 0.1 Ni 1.2 was higher than that of amorphous Zr 36 Ni 64 . 5. Compared with the crystalline phase, amorphous ZrMn 0.6 V0.2 Co 0.1 Ni 1.2 had no plateau region in the P–C–T curves. In addition, the maximum hydrogenation capacity was less, and more irreversible hydrogen was trapped in the amorphous alloy.
Acknowledgements This research was supported by the National Science Council of the Republic of China under contract NSC852216-E-007-019.
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[5] J.H. Harris, W.A. Curtin, L. Schultz, J. Mater. Res. 3 (1988) 872. [6] J. Eckert, L. Schultz, E. Hellstern, K. Urban, J. Appl. Phys. 64 (1988) 3224. [7] B.L. Chu, C.C. Chen, T.P. Perng, Metall. Trans. A 23A (1992) 2107. [8] B.L. Chu, S.M. Lee, T.P. Perng, J. Appl. Phys. 69 (1991) 4211. [9] R.B. Schwarz, R.R. Petrich, C.K. Saw, J. Non-Cryst. Solids 76 (1985) 281. [10] G. Cocco, S. Enzo, L. Schiffini, L. Battezzati, Mater. Sci. Eng. 97 (1988) 43. [11] L. Battezzati, G. Cocco, L. Schiffini, S. Enzo, Mater. Sci. Eng. 97 (1988) 121. [12] C. Politis, W.L. Johnson, J. Appl. Phys. 60 (1986) 1147. [13] E. Hellstern, L. Schultz, Appl. Phys. Lett. 49 (1986) 1163. [14] A.J. Maeland, L.E. Tanner, G.G. Lobowitz, J. Less-Common Met. 74 (1980) 279. [15] L. Zaluski, A. Zaluska, J.O. Strom-Olsen, J. Alloys Comp. 217 (1995) 245. [16] K. Aoki, H. Aoyagi, A. Memezawa, T. Masumoto, J. Alloys Comp. 203 (1994) L7. [17] L. Zaluski, P. Tessier, D.H. Ryan, C.B. Doner, A. Zaluska, J.O. Strom-Olsen, M.L. Trudeau, R. Schulz, J. Mater. Res. 8 (1993) 3059. [18] B.L. Chu, S.M. Lee, T.P. Perng, Int. J. Hydrogen Energy 16 (1991) 413. [19] Z.H. Chin, C.C. Chen, T.P. Perng, Mater. Sci. Eng. A 173 (1993) 383. [20] F.H.M. Spit, J.W. Drijver, S. Radelaar, Scripta Metall. 14 (1980) 1071. [21] Y.S. Hsu, T.P. Perng, J. Alloys Comp. 227 (1995) 180. [22] Y.S. Hsu, Master’s Thesis, National Tsing Hua University, Hsinchu, Taiwan, 1995. [23] T.P. Perng, C.J. Altstetter, Acta Metall. 34 (1986) 1771. [24] K. Girgis, in: R.W. Cahn, P. Haasen (Eds.), Physical Metallurgy, 3rd Edition, Elsevier Science, Amsterdam, 1983, p. 240. [25] J.J. Lin, T.P. Perng, Acta Metall. Mater. 39 (1991) 1101. [26] K. Aoki, M. Kamachi, T. Masumoto, J. Non-Cryst. Solids 61 / 62 (1984) 679.
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Preparation and hydrogenation of multicomponent AB 2 -type Zr–Mn–V–Co–Ni amorphous alloy L.-C. Lai, C.-L. Lee, T.-P. Perng* Department of Materials Science and Engineering, National Tsing Hua University, Hsinchu, Taiwan Received 2 March 2000; accepted 10 March 2000
Abstract Amorphous powder of the multicomponent alloy ZrMn 0.6 V0.1 Co 0.2 Ni 1.2 was prepared by ball milling of the intermetallic compound. The amorphous phase was formed after milling for 80 h. Hydrogenation of the amorphous powder was conducted and a comparison was made with that of the crystalline alloy. The absorption rate increased as the temperature increased, but the rate was slower than that of the crystalline phase. There was no plateau region in the pressure–composition–temperature curve, the maximum hydrogenation capacity was less, and a large amount of hydrogen was trapped at lower pressures. 2000 Elsevier Science S.A. All rights reserved. Keywords: Amorphous materials; Interstitial alloys; Transition metal compounds; hydrogen absorption
1. Introduction The ball-milling technique has been applied extensively to prepare Ti- and Zr-based binary amorphous alloys, e.g. Ni–Zr [1–6], Ti–Mn [7], Ti–Fe [8], Ti–Ni [9–11], Ti–Cu [12], and Fe–Zr [13]. Since both Ti and Zr are hydride formers, they are frequently used as constituent elements to prepare hydrogen storage materials. Conventional crystalline hydrogen storage materials can be classified into several categories according to their constituent elements and the type of structure, e.g. AB 5 , AB 2 , or AB 2 , where A is a hydride former and B is often a transition metal. The hydrogenation characteristics of amorphous hydrogen storage materials are very different from those of traditional crystalline materials. For example, Maeland et al. [14] reported that amorphous TiCu absorbed 35% more hydrogen than crystalline TiCu (TiCu 1.35 compared with TiCu). Recently, the characteristics of hydrogenation of other amorphous or nanocrystalline hydrogen storage alloys prepared by mechanical milling were studied, such as Mg–Ni [15,16], Ti–Fe [17,18], Ti–Mn [19], and Zr–Ni [20] alloys. Chu et al. [18] reported the formation of amorphous TiFe with some nanocrystals embedded in the matrix by mechanical alloying Ti with Fe powder. The hydrogen absorption capacity of this alloy was smaller *Corresponding author. Tel.: 1886-7-3574-2634; fax: 1886-7-35722713. E-mail address: [email protected] (T.-P. Perng).
than that of crystalline TiFe. Zaluski et al. [15] reported the hydrogen absorption properties of nanocrystalline Mg 2 Ni formed by mechanical alloying. They found that the activation for hydrogen absorption was much easier, because of the very active surface of the powders created in the ball-milling process. In addition, Zaluski et al. [17] also found that nanocrystalline TiFe showed a narrower pressure–composition–temperature (P–C–T ) plateau and a higher capacity for hydrogenation at low pressures than its crystalline counterpart. Chin et al. [19] reported that amorphous alloy TiMn could be prepared by mechanical alloying. This alloy could also absorb hydrogen without activation treatment, but no plateau region was observed in the P–C–T curve. In the Zr–Ni system, Spit et al. [20] examined the hydrogen absorption characteristics of amorphous Zr 36 Ni 64 and found that it absorbed only a small amount of hydrogen. Except for the case of TiCu alloy, the hydrogen absorption capacities of amorphous alloys are usually less than those of the corresponding crystalline alloys. Previous reports on the hydrogenation properties of amorphous alloys have mostly been focused on binary alloy systems. In studying the hydrogenation properties of AB 2 -type alloys, we have found that the multicomponent ZrMn 0.6 V0.2 Co 0.1 Ni 1.2 , derived from Zr(MnNi) 2 , could maintain the C15-type Laves phase structure and was an excellent hydrogen storage compound [21]. The hydrogen absorption capacity (H / M) was as high as 1.2, higher than that of a similar binary alloy Zr 36 Ni 64 , the capacity of
0925-8388 / 00 / $ – see front matter 2000 Elsevier Science S.A. All rights reserved. PII: S0925-8388( 00 )00826-4
L.-C. Lai et al. / Journal of Alloys and Compounds 307 (2000) 266 – 271
which is only about 0.65. In addition, the hydrogenation rate was very fast, and an electrode made from this composition had a very high discharge capacity (|390 mAh / g) [22]. In this study, an amorphous powder of ZrMn 0.6 V0.2 Co 0.1 Ni 1.2 was prepared from the intermetallic alloy by high-energy ball milling. The amorphization process and the crystalline behavior were studied. The hydrogenation properties of the amorphous phase were examined and compared with those of the crystalline phase.
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conditions, hydrogen was absorbed by the powders rapidly. After repeating several activation treatments until a steadystate absorption rate had been reached, the reaction vessel was cooled to room temperature and evacuated to about 10 24 Pa. The kinetics of hydrogen absorption and P–C–T curves were then measured. For the kinetics measurement, the initial hydrogen pressure was set at 3.3 MPa.
3. Results and discussion
3.1. Mechanical milling 2. Experimental Pure elements Zr, Mn, V, Co, and Ni (purity 99.9– 99.99%) were used as starting materials. They were mixed in the stoichiometry ZrMn 0.6 V0.2 Co 0.1 Ni 1.2 and the intermetallic compound was prepared by arc melting on a water-cooled copper plate under argon atmosphere. The ingots were repeatedly turned and remelted several times to achieve homogeneity. Some ingots were sealed in a quartz tube in vacuum and then annealed at 10008C for 10 h. Both as-melted and annealed ZrMn 0.6 V0.2 Co 0.1 Ni 1.2 ingots were crushed and pulverized mechanically into powders in a glove box filled with argon. Powders between 150 and 2100 mesh were selected for hydrogenation measurements. For ball milling, powders from the asmelted alloy smaller than 100 mesh were used. Approximately 15 g of powder and eight hardened-steel balls were sealed in a hardened-steel vial under a purified argon atmosphere. The volume of the vial was 56 ml. The weight and diameter of each ball were 3.95 g and 10 mm, respectively. The milling procedure was carried out in a Spex-8000 high-energy ball mill at a shaking rate of 1200 rpm. The structures of the powder were characterized by X-ray diffraction (XRD). The morphology was examined by scanning electron microscopy (SEM). Differential scanning calorimetry (DSC) was used to analyze the crystallization behavior of the milled powder in Ar. The heating rate was set at 158C / min. In addition, three temperatures, 490, 560, and 7008C, were selected for annealing the milled powder. The samples were heated in vacuum to these temperatures for 10 min. A Sieverts-type apparatus was employed for the hydrogenation experiment. Approximately 5 g of sample powder was placed in a stainless-steel vessel for the measurement. Since the ball-milled powder was very fine and easily oxidized, it was handled in a glove box under Ar atmosphere. After milling, the vial was kept closed and moved directly to an Ar-filled glove box. The powders were transferred from the vial and sealed in the vessel for hydrogenation test. Activation of the amorphous powders was carried out first in 3.3 MPa hydrogen gas at 3008C. Under such
Ball milling of ZrMn 0.6 V0.2 Co 0.1 Ni 1.2 intermetallic compound led to refinement of the particles. The evolution of morphologies of the ZrMn 0.6 V0.2 Co 0.1 Ni 1.2 particles after various milling times is illustrated by the SEM micrographs in Fig. 1. Before milling (Fig. 1a), the alloy had particle sizes of 50–120 mm and had a typical brittle feature. After 4 h of milling (Fig. 1b), the powders were fractured by the balls to form smaller particles. As the milling proceeded to 36 h (Fig. 1c), the small particles were repeatedly fragmented and cold welded to form larger particles. After 80 h of milling (Fig. 1d), the particles became spherical and the size became more uniform, with an average particle size of about 20 mm diameter. Fig. 2 presents the XRD patterns of the ZrMn 0.6 V0.2 Co 0.1 Ni 1.2 powders after various milling times. The XRD pattern for the starting powder showed a typical C15 intermetallic compound. After 4 h of milling, the peaks broadened and the intensities decreased, which implied a decrease of the grain size and an increase of chemical disorder of the particles. With increased milling time, the intensity of the peaks was further reduced. After 36 h of milling, some peaks became nearly invisible. When milling was extended to 80 h, only one broad peak was observed, indicating that an amorphous structure had been formed. In this multicomponent intermetallic, the atomic sizes of the constituent elements are different from each other. The free energy of the compound after milling was raised due to accumulation of lattice defects or dislocations. The formation sequence of the amorphous phase was similar to that of binary intermetallics by mechanical milling.
3.2. DSC analysis and thermal treatment From the above XRD observation, it was believed that amorphous alloy ZrMn 0.6 V0.2 Co 0.1 Ni 1.2 was formed by ball milling. The crystallization behavior of the amorphous powder was examined by DSC. Fig. 3 shows the DSC curve of the powder that was ball milled for 80 h. A broad exothermic peak ranging from 420 to 5608C with the maximum at 5118C was observed. In order to identify the phase change during heating, heat treatment of the powder at 490, 560, and 7008C for 10 min was performed. The
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Fig. 1. SEM morphologies of ZrMn 0.6 V0.2 Co 0.1 Ni 1.2 powder before and after various ball milling times. (a) Before milling, (b) 4 h, (c) 36 h, and (d) 80 h.
XRD patterns are compared in Fig. 4. Only the crystalline phase was formed, except that the crystallinity improved as the temperature increased. This illustrates that the peak was due to crystallization of the amorphous phase.
3.3. Hydrogenation kinetics The sample that had been ball milled for 80 h was selected for the hydrogenation measurement. The activation was conducted by several hydriding–dehydriding cycles at 3008C in 3.3 MPa of hydrogen gas. If the sample had been exposed to air before the activation treatment, it was found that the activation process became much more difficult, presumably because the particles were too small and easy to oxidize. After the activation treatment, the hydrogen absorption kinetics curves at 40, 80, and 3008C were measured in sequence. Fig. 5 shows the hydrogen absorption rates of amorphous powder at these three temperatures. The absorption rates were all reasonably fast. As the temperature was raised, the reaction rate became even faster. Without exposure to air, the sample could maintain a clean surface for a high hydrogenation rate.
From Fig. 5a it was also found that, as the temperature increased, the steady-state hydrogen pressure became higher, i.e. the hydrogen absorption capacity decreased. Since the enthalpy of formation for the hydride is negative, i.e. hydrogenation is an exothermic reaction, it is reasonable to expect a reduced hydrogen content as the temperature is raised. From Fig. 5b, the times for amorphous ZrMn 0.6 V0.2 Co 0.1 Ni 1.2 to reach 90% of its maximum capacity were 40, 20, and 3 min at 40, 80, and 3008C, respectively. For the crystalline counterpart, however, it has been observed that, in less than 5 min at 258C, the alloy could absorb 90% of its maximum capacity [21]. Therefore, the reaction rates of the amorphous powder were much slower than those of the crystalline sample. There might be several reasons for this phenomenon. In crystalline compounds it has been reported that hydrogen diffusivities in bcc or fcc metals range from 10 26 to 10 211 cm 2 / s at room temperature [23]. In the AB 2 -type Laves phase structure, the interstitial sizes are larger than those of pure bcc or fcc metals [24]. Hydrogen diffusivity in the AB 2 structure would be faster. In amorphous structures,
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Fig. 2. XRD patterns of ZrMn 0.6 V0.2 Co 0.1 Ni 1.2 powder before and after various ball milling times. (a) Before milling, (b) 4 h, (c) 16 h, (d) 36 h, and (e) 80 h.
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Fig. 4. XRD patterns of ZrMn 0.6 V0.2 Co 0.1 Ni 1.2 powder after heating at various temperatures for 10 min. (a) Before heating, (b) 4908C, (c) 5608C, and (d) 7008C.
hydrogen diffusivity is more like that in fcc metals [25]. Therefore, it is expected that hydrogen diffusivity in amorphous ZrMn 0.6 V0.2 Co 0.1 Ni 1.2 would be slower than that in the crystalline phase. In addition, in the crystalline compound, fresh surface is created after activation. The fresh surface would also accelerate the hydrogen absorp-
Fig. 3. DSC curve of amorphous ZrMn 0.6 V0.2 Co 0.1 Ni 1.2 .
Fig. 5. Kinetics curves of hydrogen absorption in amorphous ZrMn 0.6 V0.2 Co 0.1 Ni 1.2 alloy after activation. (a) Pressure change, (b) hydrogen content of the alloy. H S , saturation capacity.
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tion process. For the amorphous powders, they were already too small to crack.
3.4. P–C–T curves Fig. 6 shows the P–C–T curves for the ZrMn 0.6 V0.2 Co 0.1 Ni 1.2 amorphous alloy at 40 and 808C. Three curves for Zr 36 Ni 64 amorphous alloy from Spit et al. are also included for comparison [20]. There is no plateau region in all curves, and the hydrogen absorption capacity decreases as the temperature increases. At low pressures, there is a great deal of hydrogen trapped in both amorphous alloys. This could be due to the larger holes in the amorphous structure which stabilize the dissolved hydrogen atoms. Under the same pressure, ZrMn 0.6 V0.2 Co 0.1 Ni 1.2 absorbs more hydrogen than Zr 36 Ni 64 , but the amount of trapped hydrogen is also higher. The P–C–T curves of the amorphous ZrMn 0.6 V0.2 Co 0.1 Ni 1.2 alloy were also compared with those of the as-melted (AM) and homogenized (H) crystalline alloys at 40 and 808C (Fig. 7). For the crystalline alloys, they both exhibited typical plateau regions in the P–C–T curves. But in the as-melted state, the P–C–T curves had more sloped plateaus, and the maximum hydrogenation capacities were smaller. At low pressures, more hydrogen was trapped in the as-melted alloy than in the annealed alloy. Compared with the crystalline alloys, the amorphous alloy showed very different behavior in the P–C–T curves [17–20,26]. Except for the case of TiCu [14], the maximum hydrogenation capacity in the amor-
Fig. 7. P–C–T curves of ZrMn 0.6 V0.2 Co 0.1 Ni 1.2 . H, homogenized; AM, as-melted; A, amorphous.
phous phase is always not larger, and is usually lower, than that in the corresponding crystalline phase. The P–C–T curves of crystalline hydrogen storage alloys usually have three phase regions, i.e. a, a1b, and b regions. In the a region, hydrogen exists in the metal as a solid solution. The solubility of hydrogen in the metal increases as the hydrogen pressure increases. When it approaches the solubility limit, phase transformation from the solid solution to hydride occurs. The coexistence of two phases (a1b) is indicated by the generation of a plateau pressure because of equal chemical potentials. For an ideal hydride former, a flat plateau slope is observed. When segregation in the alloy occurs, the plateau slope becomes steeper, as seen for the as-melted alloy. Annealing eliminates segregation to yield a more homogeneous phase. Therefore, flatter plateau regions are observed for the annealed alloy. On the contrary, in amorphous alloys, hydrogenation does not result in phase separation. Hydrogen is probably dissolved to form an ‘interstitial’ solid solution, yielding no plateau region. Some hydrogen may even preferentially form a stable hydride with the constituent elements, resulting in more irreversibly trapped hydrogen after desorption.
4. Conclusions
Fig. 6. Comparison of P–C–T curves for amorphous ZrMn 0.6 V0.2 Co 0.1 Ni 1.2 and Zr 36 Ni 64 [20] at different temperatures.
1. The multicomponent amorphous ZrMn 0.6 V0.2 Co 0.1 Ni 1.2 alloy could be prepared by mechanical milling the crystalline intermetallic compound for over 80 h. 2. The crystallization behavior of amorphous
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ZrMn 0.6 V0.2 Co 0.1 Ni 1.2 was analyzed by DSC. A broad exothermic peak with a maximum at 5118C was observed. After heating, a C15-type crystalline phase of the same alloy was formed. 3. The hydrogen absorption rate of the amorphous alloy after activation treatment increased as the temperature increased, but the reaction rate was slower than that of the crystalline phase. 4. The hydrogen absorption capacity of amorphous ZrMn 0.6 V0.2 Co 0.1 Ni 1.2 was higher than that of amorphous Zr 36 Ni 64 . 5. Compared with the crystalline phase, amorphous ZrMn 0.6 V0.2 Co 0.1 Ni 1.2 had no plateau region in the P–C–T curves. In addition, the maximum hydrogenation capacity was less, and more irreversible hydrogen was trapped in the amorphous alloy.
Acknowledgements This research was supported by the National Science Council of the Republic of China under contract NSC852216-E-007-019.
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