Cyclic hydrogenation stability of γ-hydrides for Ti25V35Cr40 alloys doped with carbon

Journal of Alloys and Compounds 648 (2015) 534e539

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Cyclic hydrogenation stability of g-hydrides for Ti25V35Cr40 alloys doped with carbon Chia-Chieh Shen a, b, c, *, Hsueh-Chih Li b a

Department of Mechanical Engineering, Yuan Ze University, Chungli 32003, Taiwan Graduate School of Renewable Energy and Engineering, Yuan Ze University, Chungli 32003, Taiwan c Fuel Cell Center, Yuan Ze University, Chungli 32003, Taiwan b

a r t i c l e i n f o

a b s t r a c t

Article history: Received 20 March 2015 Received in revised form 1 July 2015 Accepted 2 July 2015 Available online 7 July 2015

An automatic Sievert's apparatus equipped with a temperature-programmed desorption spectrometer was constructed to study the stability of annealed Ti25V35Cr40Cx (x ¼ 0 and 0.1) alloy under cyclic hydrogenation at 6 N H2. The specimens were tested at 30  C with a hydrogen loading of around 1.00 H/M, which enabled the phase transformation from b-to g-hydrides. After 500 cycles, 83% and 90% of the initial hydrogen capacities were preserved for Ti25V35Cr40 and Ti25V35Cr40C0.1, respectively. Therefore, a small amount of C doping was effective in reducing the hydrogenation degradation of Ti25V35Cr40. The hydrogenation degradation of Ti25V35Cr40 was examined by measuring the PeC isotherms, temperatureprogrammed desorption spectra, and X-ray diffraction patterns. The degradation was ascribed to intrinsic disproportionation, i.e., Ti0.25V0.35Cr0.40 þ 0.88H2 / yTiH2 þ Ti0.25yV0.35Cr0.40H1.76e2y, where the coefficient y indicates the amount of Ti-rich precipitate. The better cyclic hydrogenation stability of Ti25V35Cr40C0.1 was related to the suppression of intrinsic disproportionation by the presence of carbon atoms in the body-centered-cubic lattice. © 2015 Elsevier B.V. All rights reserved.

Keywords: TiVCr Hydrogenation PeC isotherm Cycling Disproportionation

1. Introduction TiV-based body-centered-cubic (BCC)-type alloys have been explored as advanced hydrogen storage materials because of their excellent maximum hydrogenation capacities of 3.5 wt%. It is important to investigate the stability of the TiV-based alloys under repeated hydrogenationedehydrogenation applications, such as hydrogen storage tanks, NieMH batteries, metal hydride heat pumps, separation of hydrogen isotopes, and H2 purification membranes [1e9]. To date, few studies have described the cyclic hydrogenation properties under the gas-phase for TiV-based alloys [10e16]. Kamegawa et al. reported that the annealed Ti40Cr57.5Mo2.5 alloy maintained only 43% of its initial maximum protium capacity after 300 cycles, which was ascribed to the formation of protium trapping sites caused by an increase in lattice defects that formed during cycling [10]. By alloying with a high vanadium content, the annealed Ti8Cr12V80 alloy lost only 0.8% of the initial hydrogenation

* Corresponding author. Department of Mechanical Engineering, Yuan Ze University, Chungli 32003, Taiwan. E-mail address: [email protected] (C.-C. Shen). http://dx.doi.org/10.1016/j.jallcom.2015.07.021 0925-8388/© 2015 Elsevier B.V. All rights reserved.

capacity after 500 cycles [11]. On the other hand, the cast Ti35V31Cr34 alloys were susceptible to hydrogenation degradation within five cycles via the formation of irreversible TiH2, probably due to microsegregation in the specimen [12]. Similar phenomena have been observed in other cast TiVCr alloys [13]. Only about 23% of the initial hydrogenation capacity of the cast V55Ti20.5Cr18.1Fe6.4 alloy remained after 150 cycles. Impurities in the H2 gas were responsible for some of the degradation because of the presence of metal oxides and adsorption of hydrocarbons onto the surfaces of the specimens [14]. Therefore, the hydrogenation degradation mechanism of TiV-based alloys by either intrinsic and/or extrinsic factors requires closer examination. Approaches that replace the use of the expensive vanadium element in alloys but maintain their cyclic hydrogenation properties are also required [11,16]. Recently, Chandra et al. investigated the effect of interstitial carbon on the improvement in the stability of vanadium hydrides during thermal cycling. After 1000 cycles, the degradation of effective hydrogen capacity was reduced from 20% for VHx to 5% for V0.995C0.005Hx. In addition, a significant reduction in both the plastic deformation of g-VH2 and premature fracture of vessels can be obtained by adding a small amount (0.5 at%) of interstitial carbon into the vanadium [17,18]. Kim et al. also suggested that using an

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(1)

At room temperature, only the hydrogen storage capacity discharged from g to b of the Ti25V35Cr40 alloy can be utilized. Thereby, this study focused on the hydrogenation stability of Ti25V35Cr40 alloys subjected to hydrogen loading under b 4 g cyclic transitions. An automatic Sievert's apparatus equipped with a temperatureprogrammed desorption (TPD) spectrometer was constructed to study the stability of annealed Ti25V35Cr40Cx (x ¼ 0 and 0.1) alloys under cyclic hydrogenation. A TPD technique was employed to examine whether the cycled specimen was degraded extrinsically via surface contamination by impurities in the H2 gas [14,15,22]. Additional measurements, including hydrogenation kinetics, PeC isotherm curves, maximum hydrogenation capacities, and the changes in phase structure for the specimens, were also included to determine the hydrogenation degradation behaviors of the specimens. 2. Material and methods Ti25V35Cr40 and Ti25V35Cr40C0.1 alloys were prepared. Elemental powders of Ti (purity 99.7%), V (purity 99.9%), Cr (purity 99.5%), and C (purity 99.5%) were mixed at the desired stoichiometries, and arcmelted on a water-cooled copper hearth in an argon atmosphere with a flow rate of 15 L/min. The ingots were turned and re-melted five times for homogenization. They were further sealed in a quartz tube under vacuum (102 Torr), followed by annealing at 1200  C for 2 h to eliminate segregation [7]. The carbon contents in the specimens were measured by a combustion method using an infrared absorption carbonesulfur analyzer. Previous report has shown that the carbon concentrations in Ti25V35Cr40 and Ti25V35Cr40C0.1 were 0.063(4) at% and 0.156(4) at%, respectively [23]. A modified automatic Sievert's apparatus equipped with a TPD spectrometer was constructed for the cyclic hydrogenation test [22,24]. High-purity (99.9999%) hydrogen gas was used. All specimens were activated at 400  C under a hydrogen pressure of 0.50 MPa. To completely remove the residual hydrogen, the hydrided specimens were heated under vacuum at 600  C for 1 h. After two cycles of hydriding and dehydriding (i.e., pulverization), PeC isotherm measurements at a maximum hydrogen pressure of 4 MPa were conducted. Cyclic hydrogenation of the alloy was conducted with as many as 500 cycles at 30  C, and the hydrogen loading in equilibrium was controlled at 1.15 H/M by the volumetric method. The H concentration was conducive to the b to g phase transformation. Dehydrogenation via the transition from g to b was accomplished by outgassing the hydrided specimens under vacuum using a rotary pump. Hydrogen absorption for 300 s and

3. Results and discussion 3.1. Hydrogenation kinetics and capacity The kinetics of hydrogen absorption for Ti25V35Cr40 and Ti25V35Cr40C0.1 at 30  C was recorded after each cycle, as shown in Fig. 1(a) and (b), respectively. For each cycle, hydrogen loading of around 1.00 H/M into the specimen was achieved within 100 s. However, the absorption pressures that were required to reach a saturated equilibrium increased with the cycle number. Fig. 2 shows the variations in the cyclic hydrogenation capacities at 30  C for Ti25V35Cr40 and Ti25V35Cr40C0.1. Both specimens displayed a similar tendency toward degradation of the hydrogenation capacity. Rapid degradation was observed within the first 100 cycles, followed by a gradual leveling-off. This phenomenon was correlated with the increase in the final pressures, as observed in Fig. 1. After 500 cycles, 83% and 90% of the initial hydrogen capacities were preserved in Ti25V35Cr40 and Ti25V35Cr40C0.1, respectively.

3.5

Cycle 1 Cycle 100 Cycle 300 Cycle 500

Ti25V35Cr40

H2 Pressure [MPa]

a4a þ b4b4b þ g4g

hydrogen desorption for 900 s were performed each cycle. The TPD test was conducted under a vacuum in which the pressure of H2 desorbed from the hydrided specimen was monitored as the temperature increased from 30  C to 600  C at a heating rate of 7  C/ min. The phase structures of the annealed and cycled specimens, which had been dehydrided under vacuum at 600  C for 2 h, were examined by X-ray diffraction (XRD) at two scanning rates (4 /min and 1 /min). Si was added as an internal standard for the XRD measurements.

o

30 C

3.4 3.3 3.2 3.1

0

100

200

300

time [s] (a)

3.5

H2 Pressure [MPa]

interstitial nitrogen element could slow down the hydrogenation degradation of annealed Ti33V37Mn30 alloys [19]. These results encouraged us to investigate whether a carbon element has the same positive effect for TiV-based alloys. Our previous study showed that the effective hydrogen desorption capacities (desorption pressure S0.1 MPa) at T ¼ 30  C increased from 0.80 H/ M for the annealed Ti25V35Cr40 to 0.87 H/M for the annealed Ti25V35Cr40C0.1 due to the destabilization of the g-hydride [7]. Therefore, the aim of this study was to investigate the cyclic hydrogenation behavior of the annealed Ti25V35Cr40 and Ti25V35Cr40C0.1 alloys. Based on the strong similarities in composition between Ti25V35Cr40 and other TiVCr-based specimens, reversible phase transformations of Ti25V35Cr40 occurred during hydrogenation and dehydrogenation, as described in equation (1). The phases a, b, and g represent the solid solution (BCC), the mono-hydride (BCT), and the dihydride (FCC), respectively [20e22].

535

Cycle 1 Cycle 100 Cycle 300 Cycle 500

Ti25V35Cr40C0.1

3.4

o

30 C

3.3 3.2 3.1

0

100

200 time [s]

300

(b) Fig. 1. Comparison of hydrogen absorption kinetic curves at 30  C for various cycles. (a) Ti25V35Cr40, and (b) Ti25V35Cr40C0.1 alloys.

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1.2

Ti25V35Cr40 Ti25V35Cr40C0.1

H Content [H/M]

1.1

1.0

0.9

o

30 C 0.8

0

100

200

300

400

500

Cycle Number Fig. 2. Variation of hydrogenation capacities for Ti25V35Cr40 and Ti25V35Cr40C0.1 at 30  C during 500 hydrogenation cycles.

Clearly, Ti25V35Cr40C0.1 exhibited a better resistance to hydrogenation degradation than Ti25V35Cr40. 3.2. Temperature-programmed desorption analysis It is important to verify whether the cycled specimens were degraded by extrinsic or intrinsic factors [25]. TPD spectra were used to examine the contributions, if any, of extrinsic factors [22]. The TPD spectra for the activated and cycled Ti25V35Cr40, which were charged to the b-hydride phase with a hydrogen content of 0.75 H/M and a hydrogen pressure of 0.02 MPa, are presented in Fig. 3(a). The starting desorption temperature from b to a phases (denoted as Ts) for the activated Ti25V35Cr40 was immediately observed at room temperature, as shown in Table 1, and then the desorption reaction was maximal at 261  C. The maximum desorption transition for the cycled Ti25V35Cr40 was observed at 257  C, although a portion of the TPD spectra above 257  C was not recorded due to data acquisition problems in the Sievert system. The surface of the cycled specimen did not present any indication of poisoning sealing due to impurities from the hydrogen source [26], and so the recombination of H atoms that had desorbed from the b hydride to form H2 molecules was not retarded. Some reaction conditions resulted in a significant shift in Ts from room temperature for an activated specimen to 170  C for a poisoned specimen that was sealed with impurities at the surface. These conditions have been discussed in a previous publication [22]. The TPD spectra of the activated and cycled Ti25V35Cr40C0.1 in the b-hydride form are presented in Fig. 3(b). Again, the Ts for Ti25V35Cr40C0.1 before and after 500 cycles starts at room temperature. This result indicates that the cycled specimen was not degraded by the contamination process, and that the extrinsic factors could be excluded. 3.3. PeC isotherms and crystal structure changes The features of the PeC isotherms, such as the plateau pressure and maximum hydrogenation capacity, can be employed to study the occurrence of phase transformations during hydrogenation [24,27,28]. Prior to measuring the PeC isotherms, specimens were vacuum-treated at 600  C for 2 h to remove residual hydrogen in the activated and cycled specimens. Fig. 4(a) and (b) present the PeC isotherms at 30  C before and after 500 cycles of hydrogenation for Ti25V35Cr40 and Ti25V35Cr40C0.1, respectively. The PeC isotherm shape of the cycled Ti25V35Cr40 was similar to that of the activated Ti25V35Cr40, indicating that the crystal structure before and after 500 cycles of hydrogenation was almost unchanged. This

Fig. 3. TPD curves of the activated and cycled specimens charged to the phase region of b-hydride for (a) Ti25V35Cr40 and (b) Ti25V35Cr40C0.1.

conclusion is supported by the XRD measurement shown in Fig. 5(a). The main diffraction peaks of the annealed and cycled Ti25V35Cr40 correspond to a BCC structure. Fig. 5(a) shows that the (110) diffraction peak was shifted to higher angles after cycling, indicating that the cell volume had decreased. The lattice constants for the annealed and cycled Ti25V35Cr40 were 0.30277(3) and 0.30007(27) nm, respectively (Table 1). The decrease in cell volume was accompanied by an increase in the absorption plateau pressure at a hydrogen content of 1.25 H/M in the PeC isotherm, i.e., from 0.71 MPa for the activated specimen to 1.25 MPa for the cycled specimen. Moreover, two extra diffraction peaks at 2q ¼ 38e40 for the cycled Ti25V35Cr40 were clearly observed at a reduced scan rate, as shown in Fig. 5(b). These peaks were assigned to a (002) and a (101) of the Ti-rich (HCP) precipitate. Formation of this phase led to a drop in the reversible capacity of the PeC isotherm curve, i.e., from 0.97 H/M in the activated specimen to 0.87 H/M in the cycled specimen. These results indicate that hydrogenation degradation of Ti25V35Cr40 can be ascribed to intrinsic microstructural phase separation. Precipitation may occur via disproportionation of Ti25V35Cr40 during cyclic hydrogenation, as shown in equation (2):

Ti0:25 V0:35 Cr0:40 þ 0:88H2 /yTiH2 þ Ti0:25y V0:35 Cr0:40 H1:762y (2) where the coefficient y indicates the amount of Ti-rich precipitate.

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Table 1 Parameters derived from the TPD spectra, PeC isotherms, and lattice constants of specimens before and after 500 cycles of hydrogenation. Sample

Test condition

Ts [ C]

Plateau pressurea [MPa]

Reversible capacity [H/M]

Lattice constant [nm]

Ti25V35Cr40

Activated Cycled Poisoned Activated Cycled

R.T. R.T. 170d R.T. R.T.

0.71 1.25 e 0.91 0.96

0.97 0.87, (10%)c e 1.00 0.96, (4%)c

0.30277(3)b 0.30007(27), (0.9%)c e 0.30287(7)b 0.30210(14), (0.3%)c

Ti25V35Cr40C0.1 a b c d

The plateau pressures were derived from the H2 absorption curves at H/M ¼ 1.25. The lattice constants of the annealed specimens are cited from Ref. [23]. The figures in parentheses indicate the percentages of hydrogenation capacity loss or changes in the lattice constant of the specimens after 500 hydrogenation cycles. From Ref. [22].

Disproportionation resulted from the intrinsic composition of Ti25V35Cr40, which was composed of elements that form stable hydrides, i.e., Ti and V (DHTieH ¼ 136 kJ/mol H2 and DHVeH ¼ 70 kJ/mol H2), and an element that forms less stable hydrides, Cr (DHCreH ¼ 16 kJ/mol H2) [29]. During extended cycling, a Ti-rich hydride precipitate may form in the matrix phase of Ti25V35Cr40 so that Ti25V35Cr40 gradually loses its reversible hydrogenation capacity. This titanium hydride precipitate may be dehydrided by heating to 600  C under vacuum during the TPD measurement, as reported previously [28]. Meanwhile, an interdiffusion reaction of the elements between the matrix and the precipitate will take place. Considering that the atomic radius decreases over the series Ti, V, and Cr (0.145, 0.132, and 0.125 nm,

Fig. 4. PeC isotherm curves for (a) Ti25V35Cr40 and (b) Ti25V35Cr40C0.1 alloys before and after 500 hydrogenation cycles at 30  C.

respectively), it is reasonable to expect that the cell volume of the cycled Ti25V35Cr40, with less Ti content in the matrix phase, would decrease, as evidenced by the reduction in the lattice constant. Similar disproportionation reactions in metalehydrogen systems such as LaNi5 and Sm2Fe17 have been known for a long time [24,25,30]. For a solid-solution TiV-based alloy, the H-induced TiH2 precipitate was first reported in the cast Ti35V31Cr34 alloy after only five cycles of hydrogen absorption (25  C) and desorption (80  C) [12], which led to a 48% reduction in the reversible hydrogenation capacity. However, a smaller reduction (10%) was found in the

Fig. 5. XRD patterns of the specimens, recorded at a scan rate of (a) 4 /min and (b) 1 / min. The patterns of a, b, c, d, e, and f correspond to the annealed Ti25V35Cr40, cycled Ti25V35Cr40, annealed Ti25V35Cr40C0.1, cycled Ti25V35Cr40C0.1, cycled Ti25V35Cr40, and cycled Ti25V35Cr40C0.1 alloys, respectively. The precipitate was identified as an a Ti-rich phase with an HCP structure based on JCPDS 44-1294. Si was added as an internal standard.

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annealed Ti25V35Cr40 due to the formation of a small amount of Tirich precipitate after 500 cycles in this study. Because the composition between the Ti35V31Cr34 and Ti25V35Cr40 was similar, it suggested that the homogeneous treatment was beneficial for elimination of the microstructural segregation. Therefore, the intrinsic microstructural phase separation could be partially suppressed to maintain the hydrogenation stability of the annealed Ti25V35Cr40. For the Ti25V35Cr40C0.1 alloy after 500 cycles of hydrogenation, the absorption plateau pressure at a hydrogen content of 1.25 H/M in the PeC isotherm increased slightly from 0.91 to 0.96 MPa, and the reversible capacity decreased from 1.00 to 0.96 H/M, as shown in Fig. 4(b) and Table 1. The loss of reversible hydrogenation capacity and the reduction in the lattice constant were 4% and 0.3%, respectively, compared to 10% and 0.9% for Ti25V35Cr40. Moreover, the amount of Ti-rich precipitate induced by disproportionation in Ti25V35Cr40C0.1 was less, as demonstrated by the XRD measurements shown in Fig. 5(b). A small amount of C doping was very effective in suppressing the disproportionation of the Ti25V35Cr40 alloy. Furthermore, anisotropic lattice was found in both specimens after the cyclic hydrogenation test, as seen by the fact that the intensities of the (211) peak become higher than those of the (200) peak in Fig. 5(a). This phenomenon may be related to the formation of Ti-rich precipitate that induces a tiny preferred orientation of the (211) plane. 3.4. Stabilization of g-hydride by C inclusion The cyclic stability of the g-hydride of Ti25V35Cr40C0.1 was influenced by the presence of carbon atoms in the BCC lattice. In Fig. 5(a), XRD peaks of the annealed Ti25V35Cr40C0.1 were identified as being a single BCC phase. Note that the maximum solubilities of C in b-Ti and V are 1.8 at% and 4.3 at%, respectively [31,32], both of which are significantly higher than the amount of carbon doping (0.1 at%) in Ti25V35Cr40. Therefore, there is a high possibility that this amount of carbon would be completely dissolved into the interstitial sites in Ti25V35Cr40 after the annealing, even though an extremely small amount of Ti- or V-carbide may be present in the amorphous state, which has not been detected by X-ray diffraction. Previously, a very small expansion in the lattice constant was found after the carbon doping, from 0.30277(3) nm for the annealed Ti25V35Cr40 to 0.30287(7) nm for the annealed Ti25V35Cr40C0.1, using the Pawley refinement from the X-ray diffraction patterns, as seen in Table 1. This indicates that the carbon has entered the interstitial sites [23]. The presence of carbon in the interstitial sites of the BCC lattice could also be checked by one of Hume-Rothery's rules to determine the range of solid solutions, i.e., the solute solubility increases when the ratio of electrons to atoms (e/a) for the elements in the solid solution is increased [33, 34]. The dissolution of the carbon into the lattice of the BCC structure can be referred to the fact that the e/a ratio of 3.2508 for Ti25V35Cr40C0.1 is slightly larger than that of 3.2500 for Ti25V35Cr40. The presence of the interstitial elements, such as O, N, B, and C elements, in TiV-based hydrogen storage alloys has also been found by other groups [35e40]. It is well known that the lattice will be distorted by the inclusion of the interstitial elements, and then the atomic diffusion will be blocked by the strained lattice [41]. Thus, inclusion of interstitial carbon atoms could act as an obstacle that blocks the interdiffusion of Ti, V, and Cr elements driven by Hinduced disproportionation during cycling in this study. Moreover, the heat to overcome the activation energy of disproportionation will be internally supplied from the heat of formation (DHf ) released from the transition of b-to g-hydrides [25]. Because the DHf was reduced from 43.0 kJ/mol for Ti25V35Cr40 to 39.5 kJ/ mol for Ti25V35Cr40C0.1 [23], the disproportionation rate was

accordingly and partially retarded. Such hydride destabilization by the inclusion of carbon has also been seen in the VeH system, as for the reduction of the DHf of b2 4 g from 83 kJ/mol for pure V to 60 kJ/mol for V containing 0.5 at% C [9]. As a result, better stability of the annealed Ti25V35Cr40C0.1 during hydrogenation cycling could be obtained. 4. Conclusions An automatic Sievert's apparatus equipped with a TPD spectrometer was constructed to study the cyclic hydrogenation stability of annealed Ti25V35Cr40Cx (x ¼ 0 and 0.1). The conclusions to be drawn from this study include: (1) After 500 cycles of hydrogenation, 83% and 90% of the initial hydrogen capacities were preserved for Ti25V35Cr40 and Ti25V35Cr40C0.1, respectively. (2) The TPD results show that the starting desorption from b to a phases for Ti25V35Cr40 and Ti25V35Cr40C0.1 remained at around room temperature before and after cycling. Therefore, cyclic degradation was caused by intrinsic factors. (3) The PeC isotherms and XRD results indicate that hydrogenation degradation of Ti25V35Cr40 was due to intrinsic disproportionation, as described by Ti0.25V0.35Cr0.40 þ 0.88H2 / yTiH2 þ Ti0.25yV0.35Cr0.40H1.762y. Formation of a Ti-rich precipitate led to a decrease in the maximum hydrogenation capacity. (4) The improvement in the cyclic hydrogenation stability of Ti25V35Cr40C0.1 was related to the suppression of intrinsic disproportionation by the presence of carbon atoms in the BCC lattice. Acknowledgments This work was partially supported the National Chung-Shan Institute of Science and Technology of Taiwan under Contract CSIST-769-V412 and the Ministry of Science and Technology of Taiwan under Contract MOST 103-2221-E-155-017. References [1] E. Akiba, H. Iba, Hydrogen absorption by Laves phase related BCC solid solution, Intermetallics 6 (1998) 461e470. [2] M. Aoki, T. Noritake, A. Ito, M. Ishikiriyama, S.I. Towata, Improvement of cyclic durability of Ti-Cr-V alloy by Fe substitution, Int. J. Hydrog. Energy 36 (2011) 12329e12332. [3] D. Mori, K. Hirose, Recent challenges of hydrogen storage technologies for fuel cell vehicles, Int. J. Hydrog. Energy 34 (2009) 4569e4574. [4] H. Miao, W.G. Wang, Mechanisms of improving the cyclic stability of V-Tibased hydrogen storage electrode alloys, J. Alloys Compd. 508 (2010) 592e598. [5] J.H. Yoo, G. Shim, J.S. Yoon, S.W. Cho, Effects of substituting Al for Cr in the Ti0.32Cr0.43V0.25 alloy on its microstructure and hydrogen storage properties, Int. J. Hydrog. Energy 34 (2009) 1463e1467. [6] A. Gueguen, J.M. Joubert, M. Latroche, Influence of the C14 Ti35.4V32.3Fe32.3 Laves phase on the hydrogenation properties of the body-centered cubic compound Ti24.5V59.3Fe16.2, J. Alloys Compd. 509 (2011) 3013e3018. [7] C.C. Shen, J.C.P. Chou, H.C. Li, Y.P. Wu, T.P. Perng, Effect of interstitial boron and carbon on the hydrogenation properties of Ti25V35Cr40 alloy, Int. J. Hydrog. Energy 35 (2010) 11975e11980. [8] M.D. Dolan, Non-Pd BCC alloy membranes for industrial hydrogen separation, J. Membr. Sci. 362 (2010) 12e28. [9] J. Lamb, D. Chandra, M. Coleman, A. Sharma, W.N. Cathey, S.N. Paglieri, J.R. Wermer, R.C. Bowman Jr., F.E. Lynch, Low and high-pressure hydriding of Ve0.5at.%C, J. Nucl. Mater. 399 (2010) 55e61. [10] A. Kamegawa, T. Tamura, H. Takamura, M. Okada, Protium absorptiondesorption properties of Ti-Cr-Mo bcc solid solution alloys, J. Alloys Compd. 356e357 (2003) 447e451. [11] H. Itoh, H. Arashima, K. Kubo, T. Kabutomori, K. Ohnishi, Improvement of cyclic durability of BCC structured Ti-Cr-V alloys, J. Alloys Compd. 404e406 (2005) 417e420. [12] J.Y. Wang, R.R. Jeng, J.K. Nieh, S. Lee, S.L. Lee, H.Y. Bor, Comparing the

C.-C. Shen, H.-C. Li / Journal of Alloys and Compounds 648 (2015) 534e539

[13]

[14]

[15]

[16]

[17]

[18]

[19] [20]

[21]

[22] [23]

[24]

[25]

hydrogen storage alloys-TiCrV and vanadium-rich TiCrMnV, Int. J. Hydrog. Energy 32 (2007) 3959e3964. H.C. Lin, K.M. Lin, K.C. Wu, H.H. Hsiung, H.K. Tsai, Cyclic hydrogen absorptiondesorption characteristics of TiCrV and Ti0.8Cr1.2V alloys, Int. J. Hydrog. Energy 32 (2007) 4966e4972. L. Hao, Y. Chen, Y. Yan, C. Wu, M. Tao, Cyclic properties of hydrogen absorption and desorption in V-Ti-Cr-Fe (Al, Si) alloy, Mater. Sci. Eng. A 448 (2007) 128e134. H. Tanaka, N. Kuriyama, S. Ichikawa, H. Senoh, N. Naka, K. Aihara, H. Itoh, M. Tsukahara, Degrading mechanism on hydrogen absorbing-desorbing cycle durability of V- and Ti-Cr-based bcc-type solid solutions, Mater. Sci. Forum 475e479 (2005) 2481e2484. T. Kuriiwa, T. Maruyama, A. Kamegawa, M. Okada, Effects of V content on hydrogen storage properties of V-Ti-Cr alloys with high desorption pressure, Int. J. Hydrog. Energy 35 (2010) 9082e9087. D. Chandra, A. Sharma, R. Chellappa, W.N. Cathey, F.E. Lynch, R.C. Bowman Jr., J.R. Wermer, S.N. Paglieri, Hydriding and structural characteristics of thermally cycled and cold-worked V-0.5 at.% C alloy, J. Alloys Compd. 452 (2008) 312e324. J.S. Cantrell, R.C. Bowman Jr., Phase composition and the effect of thermal cycling for VHx, V0.995C0.005Hx, and V0.975Zr0.020C0.005Hx, J. Alloys Compd. 293e295 (1999) 156e160. H. Kim, K. Sakaki, Y. Nakamura, Improving the cyclic stability of V-Ti-Mn alloys using interstitial elements, Mater. Trans. 55 (2014) 1144e1148. T. Kazumi, T. Tamura, A. Kamegawa, H. Takamura, M. Okada, Effect of absorption-desorption cycles on structure and stability of protides in Ti-Cr-V alloys, Mater. Trans. 43 (2002) 2748e2752. G. Mazzolai, B. Coluzzi, A. Biscarini, F.M. Mazzolai, A. Tuissi, F. Agresti, S.L. Russo, A. Maddalena, P. Palade, G. Principi, J. Alloys Compd. 466 (2008) 133e139. C.C. Shen, H.C. Li, Passivation and reactivation of Ti25V35Cr40 hydrides by cycling with impure hydrogen, Int. J. Hydrog. Energy 40 (2015) 3277e3282. C.C. Shen, K.C. Wu, H.C. Li, Y.P. Wu, Influence of interstitial carbon on the formation of monohydride and dihydride of Ti25V35Cr40 alloys, Mater. Chem. Phys. 151 (2015) 87e92. C.C. Shen, T.P. Perng, On the cyclic hydrogenation stability of an Lm(NiAl)5based alloy with different hydrogen loadings, J. Alloys Compd. 392 (2005) 187e191. P.D. Goodell, Stability of rechargeable hydriding alloys during extended cycling, J. Less-Common Met. 99 (1984) 1e14.

539

[26] H. Sakaguchi, T. Tsujimoto, G.Y. Adachi, The confinement of hydrogen in LaNi5 by poisoning of the hydride surface, J. Alloys Compd. 223 (1995) 122e126. [27] D. Chandra, Intermetallics for hydrogen storage, in: G. Walker (Ed.), Solidstate Hydrogen Storage, Woodhead Publishing Limited, Cambridge, 2012, pp. 317e318. [28] C.C. Shen, T.P. Perng, Pressure-composition isotherms and reversible hydrogen-induced phase transformations in Ti-6Al-4V, Acta Mater. 55 (2007) 1053e1058. [29] R. Griessen, T. Riesterer, Heat of formation models, in: L. Schlapbach (Ed.), Hydrogen in Intermetallic Compounds I, Springer-Verlag, Berlin, 1988, p. 266. [30] S. Sugimoto, T. Maeda, D. book, T. Kagotani, K. Inomata, M. Homma, H. Ota, Y. Houjou, R. Sato, GHz microwave absorption of a fine a-Fe structure produced by the disproportionation of Sm2Fe17 in hydrogen, J. Alloys Compd. 330e332 (2002) 301e306. [31] J.L. Murray, Phase Diagrams of Binary Titanium Alloys, ASM International, Metals Park, OH, 1987. [32] J.F. Smith, Phase Diagrams of Binary Vanadium Alloys, ASM International, Metals Park, OH, 1989. [33] G.P. Tiwari, R.V. Ramanujan, The relation between the electron to atom ratio and some properties of metallic systems, J. Mater. Sci. 36 (2001) 271e283. [34] R.E. Smallman, A.H.W. Ngan, Phase diagrams and alloy theory, in: Modern Physical Metallurgy, Elsevier Ltd., 2014, pp. 79e84. [35] H.Y. Chang, C.A. Wert, The solubility and trapping of hydrogen in vanadium, Acta Metall. 21 (1973) 1233e1242. [36] Y. Nakamura, J. Nakamura, K. Sakaki, K. Asano, E. Akiba, Hydrogenation properties of Ti-V-Mn alloys with a BCC structure containing high and low oxygen concentrations, J. Alloys Compd. 509 (2011) 1841e1847. [37] M. Uno, K. Takahashi, T. Maruyamy, H. Muta, S. Yamanaka, Hydrogen solubility of BCC titanium alloys, J. Alloys Compd. 366 (2004) 213e216. [38] J. Shi, T. Sakai, H.T. Takeshita, N. Kuriyama, M. Tsukahara, Influence of carbon impurity on microstructures and electrode properties for V-based battery alloys, J. Alloys Compd. 290 (1999) 267e272. [39] S.W. Cho, J.H. Yoo, G. Shim, C.N. Park, J. Choi, Effect of B addition on the hydrogen absorption-desorption property of Ti0.32Cr0.43V0.25 alloy, Int. J. Hydrog. Energy 33 (2008) 1700e1705. [40] S. Yamanaka, Y. Kashiwara, H. Sugiyama, M. Katsura, Influence of interstitial oxygen and nitrogen on hydrogen solubility in vanadium, J. Nucl. Mater. 247 (1997) 244e248. [41] D.R. Askeland, P.P. Fulay, W.J. Wright, The Science and Engineering of Materials, Sixth ed., Cengage Learning, CT, 2011.