Extrinsic Degradation of Ti−V-Based Hydrogen Storage

Sep 27, 2008 - The capacity degradation mechanism of the Ti-V-based hydrogen storage ... The electrochemical reactions occurring in a Ni/MH battery ca...
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16682

J. Phys. Chem. C 2008, 112, 16682–16690

Intrinsic/Extrinsic Degradation of Ti-V-Based Hydrogen Storage Electrode Alloys upon Cycling Yongfeng Liu, Hongge Pan,* Mingxia Gao, Rui Li, and Qidong Wang Department of Materials Science and Engineering, Zhejiang UniVersity, Hangzhou 310027, People’s Republic of China ReceiVed: January 18, 2008; ReVised Manuscript ReceiVed: August 13, 2008

The capacity degradation mechanism of the Ti-V-based hydrogen storage electrode alloys was systematically investigated from the viewpoint of intrinsic/extrinsic degradation for the first time. The oxidation/corrosion of the active components and the appearance of the irreversible hydrogen were found to be the two dominating intrinsic factors for the cycling capacity degradation of the Ti-V-based electrode alloys, rather than the segregation and dissolution of the main hydrogen absorbing element V, the loss of which had some effect on the degradation of its discharge capacity but not dominating. The extrinsic degradation of the electrode alloys was mainly caused by the pulverization of the alloy particles, the increase of the contact resistance between particles, and the reaction resistance on surfaces during charge/discharge cycling. In the meantime the pulverization of the alloy particles accelerates the oxidation/corrosion of the active components, which further promotes the formation of the oxide layer and decreases the reaction speed. The decay of the apparent discharge capacity of the Ti-V-based electrode alloys is jointly affected by these factors. 1. Introduction In recent years, with the development of the small portable electronic devices, the demand for the secondary batteries with a high energy density is growing rapidly. Nickel/metal hydride (Ni/MH) battery is one type of “green” rechargeable batteries with nickel hydroxide electrode as its positive electrode and one hydrogen storage alloy electrode as its negative electrode.1-3 The electrochemical reactions occurring in a Ni/MH battery can be expressed as follows:1 charge

on positive electrode:

Ni(OH)2 + OH- y\z discharge

NiOOH + H2O + echarge

on negative electrode:

M + H2O + e- y\z discharge

MH + OHcharge

overall reaction:

Ni(OH)2 + M y\z NiOOH + MH discharge

Compared with nickel/cadmium (Ni/Cd) battery, Ni/MH battery has many inherent advantages, including a higher energy density, a longer cycle life, a higher rate capacity, better tolerance to overcharge and overdischarge, and higher environmental compatibility.4-8 Additionally, it is very convenient for the user as the voltage of the Ni/MH battery is essentially the same as the conventional Ni/Cd battery. So for a period of time it has been used as the main power source for many electronic devices. Although the Li-ion battery has been widely commercialized nowadays, the Ni/MH battery is still considered a preferred rechargeable battery due to its higher rate capacity, especially in the field of electric powered hand tools.6 * To whom correspondence should be addressed. Phone: +86 571 87952615. Fax: +86 571 87952615. E-mail: [email protected].

The performance of a Ni/MH battery, evaluated with parameters such as capacity, durability (cycle life), and dischargeability (kinetics), depends strongly on the characteristics of the active components or alloys of the electrodes, especially the hydrogen storage alloy used as the negative electrode material.1-3,8-10 Hydrogen storage alloys are a group of new functional intermetallics which can reversibly absorb/desorb a large amount of hydrogen at or around normal temperature and pressure and can be used as anode materials in electrochemical storage batteries. Nowadays, AB5-type rare earth based alloys,1,3 AB3type rare earth magnesium based alloys,8,9 AB2-type multicomponent alloys,2,11 and Ti-V-based multiphase alloys12,13 are extensively being investigated for their commercial usage as negative electrode for Ni/MH batteries. Ti-V-based multiphase alloys with their high discharge capacities are developing as one of the promising candidates for the negative electrode alloy for Ni/MH batteries.12-16 Pan et al. formulated a superstoichiometric (Ti0.8Zr0.2)(V0.533Mn0.107Cr0.16Ni0.2)x (x ) 2, 3, 4, 5, 6) alloy series containing a C14 Laves phase (MgZn2-type structure) and a V-based solid solution phase.12 The V-based solid solution phase is the major hydrogenabsorbing phase, while the C14 Laves phase acts not only as a hydrogen-absorbing phase but also as a catalyst for the electrochemical hydrogenation and dehydrogenation of the V-based phase,15 and their cooperation provides some favorable electrochemical features, especially the higher discharge capacity. The maximum discharge capacity of the (Ti0.8Zr0.2)(V0.533Mn0.107Cr0.16Ni0.2)5 alloy reaches 380 mAh/g,12 which is noticeably higher than that of the commercial AB5-type alloys (∼320 mAh/g). However, its poor cycling stability and poor kinetics prevent it from being adopted in practical applications. Our previous studies have reported that the cycling stability and the electrochemical kinetics of Ti0.8Zr0.2V2.7Mn0.5Cr0.8Nix (x ) 0.0-2.0) alloys were markedly improved by optimizing its Ni content.13,16 But the role played by Ni on the cycling stability has not been fully explored yet. Therefore, it still is of both scientific and application interests to elaborate the capacity

10.1021/jp8005052 CCC: $40.75  2008 American Chemical Society Published on Web 09/27/2008

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degradation mechanism and the function of Ni for the further development of Ti-V-based hydrogen storage alloys. In this work, we first divided the capacity degradation of the hydrogen storage electrode alloys according to the sources of degradation into two categorizations, intrinsic degradation and extrinsic degradation. On the basis of different types of degradation, the capacity degradation mechanism and the function of Ni in the Ti0.8Zr0.2V2.7Mn0.5Cr0.8Nix (x ) 0.75, 1.75) alloys with quite discrepant cycling durability were investigated and formulated. 2. Features of Intrinsic/Extrinsic Degradation The cycling stability of hydrogen storage alloys is affected in a complex way by many different factors. Considerable studies have been carried out on the cycling capacity degradation mechanism of hydrogen storage electrode alloys. The capacity degradation of the rare earth based alloys was mainly caused by the oxidation/corrosion of active elements in alloys, and the selective segregation of manganese and its deposition as the hydroxide as pointed out by Willems et al.1,17,18 and Zu¨ttel et al.19 In addition, the pulverization of alloy particles caused by the lattice expansion/contraction due to hydride formation/ decomposition which results in cracks in alloys was also considered as one of the important reasons for the capacity degradation of the rare earth based alloys.20,21 The fast degradation of Ti-based alloys was attributed mainly to the increase of the polarization resistance caused by the deterioration of surface properties due to the growth of titanium-oxide on the surface by Lee.22 The loss of Mg into the KOH solution and the formation of a gel-type permeable Mg(OH)2 passive film were responsible for capacity degradation of Mg-based alloys.23 Meanwhile, for the V-based solid alloys and Ti-V-based alloys, the appearance of irreversible hydrogen and the changes of the crystallographic structure after hydrogenation were also proved to be the factors for the lowering of the cycling capacity of hydrogen storage alloys.15,24,25 Although many influence factors for the cycling durability of hydrogen storage electrode alloys as listed above have been found out through experimental measurements and theoretical analyses, no systematic analyses and theories have been made on them so far. On the basis of previous studies, we propose that all factors for the cycling capacity degradation of hydrogen storage electrode alloys can be divided into two categories, intrinsic degradation and extrinsic degradation. To clearly elucidate the degradation factors and their influence, we are proposing the following definitions: (1) Intrinsic capacity: a theoretical discharge capacity of an alloy determined by the alloy components and phase structure, which should be the sum of theoretical discharge capacities of the chemical composition in mAh/g. In general, the intrinsic capacity of the metal hydride of MHx can be calculated by the following equation:

C)

xF 3.6M

(1)

where C is the intrinsic capacity (mAh/g), x is the number of hydrogen atoms absorbed, F is the Faraday constant (96500 C/mol), and M is the molar mass of the hydrogen storage alloy (g/mol). (2) Apparent capacity: an actual discharge capacity obtained by an experimental measurement, which is generally lower than the intrinsic capacity. (3) Intrinsic degradation: a decrease in the theoretical or thermodynamic capacity of the electrode alloy due to the irreversible decrease/consumption of active components during

cycling, viz., a real and direct decrease of the capacity of the electrode alloy. (4) Extrinsic degradation: a degradation caused by the decline in the rate of utilization of active components due to the decrease of the kinetic parameters on cycling. In this case, the self of the capacity of the electrode alloy is not decreased, but its utilization efficiency is declined due to some factors, which can decrease the apparent capacity as well. In summary, the intrinsic degradation is closely related to the oxidation/corrosion of the alloy elements, the segregation/ dissolution of the alloy elements, the irreversible hydrogenation, and the change of the phase structure, etc. The extrinsic sources of degradation are mainly composed of the increase of the contact resistance caused by the pulverization, the adverse changes in surface morphology, the increase of the polarization resistance due to the formation of the oxide film, and so on. Moreover, it is interesting to note that for most of hydrogen storage electrode alloys, the intrinsic degradation factors and extrinsic degradation factors are not only varying alone or independently but also affecting each other. They jointly affect the degradation of the apparent capacity. As mentioned above, the oxidation/corrosion of the active components and the pulverization of the alloy particles are the two important factors for the cycling capacity degradation of the hydrogen storage electrode alloys. The oxidation/corrosion of the active components nullifies their ability to store hydrogen reversibly, so it is an intrinsic degradation factor. The pulverization of the alloy particles, which increases the contact resistance between the alloy particles and decreases the charge/discharge efficiency, and consequently decreases the cycling capacity of the electrode alloys, is an extrinsic degradation factor according to our definitions. However, the oxidation/corrosion of the active components deteriorates also the surface morphology of the alloy particles and, consequently, increases the contact resistance and the polarization resistance, hereby decreasing the charge/ discharge efficiency, which belongs to extrinsic degradation. This phenomenon indicates the intrinsic degradation can result in extrinsic degradation. Similarly, the pulverization of the alloy particles intensifies the flushing action of the alkaline electrolyte on the cracks of the alloy particles during cycling, consequently increasing the contact of the fresh alloy surface with the alkaline electrolyte and the oxidation/corrosion rate of the active elements. This shows that the extrinsic capacity degradation accelerates also the intrinsic capacity degradation. Although intrinsic/extrinsic factors can interplay, the capacity degradation caused by their interplay can also be still ascribed to either intrinsic degradation or extrinsic degradation. Since all sources of degradation are classified either into intrinsic degradation or into extrinsic degradation, the apparent capacity degradation can be phenomenologically represented by the addition of the two categories of degradation as shown in Figure 1. Figure 1a schematically illustrates the relationship between the theoretical capacity, the apparent capacity, and the intrinsic/extrinsic degradation of the discharge capacity. The intrinsic and extrinsic factors of degradation jointly depress the theoretical capacity to the apparent discharge capacity during cycling. On cycling, the capacity loss caused by both intrinsic and extrinsic degradation factors increases gradually. Generally, in the initial stage of cycling the intrinsic degradation is the primary factor for the capacity degradation. As cycling goes on, the rate of the intrinsic degradation gradually decreases, and the rate of the extrinsic degradation becomes higher. On further cycling, both the intrinsic degradation rate and the extrinsic degradation rate gradually stop growing and maintain at constant

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Figure 2. Cycling stability curves for Ti0.8Zr0.2V2.7Mn0.5Cr0.8Nix (x ) 0.75, 1.75) electrode alloys.

Figure 1. Schematic graph of the apparent capacity degradation of the hydrogen storage electrode alloys: (a) discharge capacity versus cycle number; (b) capacity degradation rate versus cycle number.

rates. The changes in the intrinsic and extrinsic degradation rates make that the apparent degradation of the discharge capacity is faster in the initial stage of capacity degradation, and gradually slows down as cycling goes on, shown in Figure 1b.1,3,8-29 3. Experimental Section 3.1. Alloy Preparation. Ti0.8Zr0.2V2.7Mn0.5Cr0.8Nix (x ) 0.75, 1.75) alloys were prepared by induction levitation melting under argon atmosphere. The purity of all starting elemental metals is higher than 99%. For homogeneity, the alloys were turned over and remelted twice. The ingots were mechanically crushed and ground into powder for structural and electrochemical experiments. The average particle diameter of the resulting powder was measured by Malvern particle analyzer Mastersizer2000 (Malvern). 3.2. Electrochemical Measurement. The test electrodes for basic electrochemical measurement and scanning electron microscopy (SEM) observation were made by cold pressing a mixture of 100 mg of alloy powder and 400 mg of pure carbonyl nickel powder (INCO) into a pellet electrode under 20 MPa pressure. Here, carbonyl nickel powder works as current collector for decreasing the contact resistance and hence increasing the conductivity of the electrode. For X-ray diffraction (XRD) and Auger electron spectroscopy (AES) investigations, the electrode pellets were prepared by cold pressing 400

mg of pure alloy powder alone to avoid the effect of the additional Ni signal. The pellet electrode was sandwiched within two foamed Ni plates (55 × 20 mm2) with a Ni wire soldered on each of them to form a complete testing negative electrode. Electrochemical measurements were carried out in a standard trielectrode cell, consisting of a working electrode (the MH electrode for studying), a sintered Ni(OH)2/NiOOH counter electrode, and a Hg/HgO reference electrode, immersed in 6 M KOH electrolyte. Each test electrode was charged at 100 mA/g for 5 h followed by a 10 min rest and then discharged at 60 mA/g to the cutoff potential of -0.6 V vs the Hg/HgO reference electrode. After a preset number of charging/discharging cycles, the electrode pellet was taken out, washed with distilled water, and dried in vacuum for structural analysis. Electrochemical impedance spectroscopy (EIS) studies were conducted at 50% depth of discharge (DOD) by using a Solartron SI1287 electrochemical interface together with a 1255B frequency response analyzer (Solartron). EIS curves of the cycled alloy electrodes were obtained in the frequency range from 10 kHz to 5 mHz with an alternating current (ac) amplitude of 5 mV under the open-circuit condition. 3.3. Structural Analysis. X-ray diffraction was carried out on an ARL X-ray diffractometer (ARL) with Cu KR radiation at 45 kV and 40 mA. The surface morphology of the alloy particle was examined by a FEI-SLR10N scanning electron microscope (FEI). To study the element distribution on the cycled alloy surface layer, AES depth profiles were measured by using a PHI-550-type electron spectrometer (Perkin-Elemer) with an electron beam at 3 kV and 10 µA. The alloy surface was sputtered with Ar+ on an area of 1.5 × 1.5 mm2 at 4 kV and 15 mA and a sputtering rate of 3.5 nm/min. 3.4. Analysis for Dissolved Element. To investigate the species and concentrations of alloy components dissolved into the alkaline electrolyte, the KOH solution was analyzed after a preset number of charging/discharging cycles with an inductively coupled plasma spectrometer (ICP; Leeman). 4. Results and Discussion 4.1. Cycling Stability Curves. Figure 2 shows the cycling stability curves of Ti0.8Zr0.2V2.7Mn0.5Cr0.8Nix (x ) 0.75, 1.75) electrode alloys. Obviously, the discharge capacity of the alloy with x ) 0.75 decreases more rapidly with cycling, while the capacity degradation of the alloy with x ) 1.75 is much slower. After 240 charging/discharging cycles, the capacity retentions

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Figure 3. Dissolution amount of active elements from Ti0.8Zr0.2V2.7Mn0.5Cr0.8Ni0.75 (a) and that from Ti0.8Zr0.2V2.7Mn0.5Cr0.8Ni1.75 (b) electrode alloys on cycling.

(C240/Cmax) of the electrode alloys are 5.5 and 89.2% for x ) 0.75 and 1.75, respectively. In other words, the cycling stability of the alloys was noticeably improved by the increase of Ni content. It is very interesting to note that such a large improvement was achieved by the increase of Ni content alone. This phenomenon urged us to investigate the intrinsic and extrinsic factors of the cycling capacity degradation of the Ti-V-based hydrogen storage alloys and the role Ni played in this type of alloy. 4.2. Intrinsic Degradation. We first investigated the segregation/dissolution of the constituent elements of alloys into the electrolyte by analyzing the cycled alkaline solution with ICP. The results are presented in Figure 3. After charge/ discharge cycling, four alloy elements of V, Ti, Zr, and Mn were found in the KOH solution, while the Cr and Ni elements were not detected. This fact indicates the alloy constituents of V, Ti, Zr, and Mn are dissolved into the alkaline solution easily on cycling, while the segregation/dissolution of Cr and Ni from the alloys is much harder. It is important to note that the dissolution amount of V is the highest, much higher compared with other elements. Since V is known as the main hydrogen absorbing element in the Ti-V-based alloys,12,13 the dissolution of V into electrolyte will cause the alloy to lose its intrinsic capacity. It is thought to be a typical intrinsic source of degradation. On analyzing Figure 2 and Figure 3 simultaneously, however, the capacity degradation did not correlate well with the V dissolution amount. The discharge capacity of the alloy with x ) 0.75 decreases from 373.7 to 38.9 mAh/g in 200 cycles (Figure 2), and it loses linearly at the rate of 1.78 mAh/(g · cycle) from the 12th cycle to the 200th cycle, whereas the dissolution amount of V increases quickly within the initial 60 cycles and then slows down gradually on further cycling, as shown in Figure 3a, exhibiting a quite different change in pattern. The total dissolution amount of V was 50.2 µg/mL in 240 cycles. In the initial 60 cycles, the dissolution amount of V has reached 40.2 µg/mL. In the subsequent 180 cycles, only 10 µg/mL of V was dissolved into the alkaline solution. These results indicate that the dissolution of V is not the primary source for the cycling capacity degradation of the Ti-V-based hydrogen storage electrode alloys in the present study, especially after 100 charge/ discharge cycles. The oxidation/corrosion of the active components, which makes the hydrogen storage electrode alloys fail to store

hydrogen reversibly, and inevitably decreases its apparent discharge capacity, should be one of the most important intrinsic factors for capacity degradation. For investigating the rate of the oxidation/corrosion of the different active components of the Ti-V-based alloys, the AES depth profiles of Ti0.8Zr0.2V2.7Mn0.5Cr0.8Nix (x ) 0.75, 1.75) electrode alloys were plotted after cycling 120 times, as shown in Figure 4. The oxygen content in the surface layer of the alloys after cycling 120 times is very high (about 65%), indicating apparently the serious oxidation of the surface elements after charge/discharge cycling in KOH solution. For the alloy with x ) 0.75, the oxygen concentration in the surface layer slowly declines with the distance from the outer surface. Even after being sputtered for 6000 s, the oxygen concentration is still above 35%, revealing that the oxidation/corrosion layer of the alloy with x ) 0.75 is rather thick, which should be one reason for its capacity decay. However, the oxygen concentration in the surface layer of the alloy with x ) 1.75 decreases much more quickly with sputtering, reaching some 20% after being sputtered for 3000 s. These results signify that the antioxidation/anticorrosion ability of the Ti-V-based alloys is markedly improved by the increase of Ni content in the alloy, hence, the marked improvement of the cycling stability. In the meantime it should be noticed that, for these two electrode alloys, the element V almost becomes undetectable in the surface layer and its concentration is close to 0 even after being sputtered for 6000 s due to the segregation and dissolution of V as disclosed previously by the elemental analysis of the cycled alkaline solution (Figure 3). It is well-known that the V-based solid solution alone possesses very little electrochemical discharge capacity in alkaline electrolyte since it does not have any electrocatalytic activity by itself. Only with the presence of a secondary phase acting as a catalyst and a microcurrent collector, such as the TiNi phase or the C14 Laves phase (MgZn2-type structure), the V-based solid solution can reversibly absorb and desorb a large amount of hydrogen in alkaline electrolyte.25 In case of a lack of a catalytic phase in the Ti-V-based electrode alloys, the hydrogen absorbed by the V-base solid solution phase is not able to be released, resulting in a big decrease of the discharge capacity of the alloy. Therefore, the irreversible hydrogenation is also one of the very important intrinsic factors responsible for the cycling capacity degradation. For revealing the structural change of Ti0.8Zr0.2V2.7Mn0.5Cr0.8Nix (x ) 0.75,

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Figure 4. AES depth profiles of Ti0.8Zr0.2V2.7Mn0.5Cr0.8Ni0.75 (a) and Ti0.8Zr0.2V2.7Mn0.5Cr0.8Ni1.75 (b) electrode alloys after 120 cycles.

Figure 5. XRD patterns of Ti0.8Zr0.2V2.7Mn0.5Cr0.8Ni0.75 (a) and Ti0.8Zr0.2V2.7Mn0.5Cr0.8Ni1.75 (b) electrode alloys after different cycles.

1.75) alloys on cycling, the XRD investigations were performed on the electrode alloys before and after cycling. Figure 5 shows the XRD patterns of the electrode alloys before cycling and after 15 and 60 cycles, respectively. It can be seen from the figure that the starting alloys consist mainly of a C14 Laves phase with MgZn2-type hexagonal structure and a V-based solid solution phase with BCC structure as previously reported.13 After being cycled, the main diffraction peaks of the Ti0.8Zr0.2V2.7Mn0.5Cr0.8Ni0.75 alloy obviously shift to lower angles, especially those of V-based solid solution phase, indicating that the cell volume of the V-based solid solution phase expands due to a considerable amount of hydrogen being absorbed by the alloy phases and without being redesorbed, resulting in a decrease in discharge capacity. However, for the Ti0.8Zr0.2V2.7Mn0.5Cr0.8Ni1.75 alloy, the XRD patterns remain almost unchanged and the position of the diffraction peaks shifts very slightly to lower angles after being cycled, implying there is only very little irreversible hydrogen present in the hydrogenated Ti0.8Zr0.2V2.7Mn0.5Cr0.8Ni1.75 alloy due to its high

electrocatalytic activity caused by its high Ni content. This phenomenon is responsible for its superior cycling capacity retention. According to the above analyses, we believe that, for the Ti-V-based electrode alloys, the oxidation/corrosion of the active elements and the presence of the irreversible hydrogen are two dominating factors for the intrinsic degradation during cycling, rather than the segregation and dissolution of the main hydrogen absorption element V, which is certainly another cause for the loss of the discharge capacity but not as important as the two mentioned above. The increase of the Ni content in the Ti-V-based hydrogen storage electrode alloys not only retards the oxidation/corrosion of the active elements in the alloy but also decreases the appearance of the irreversible hydrogen, improving consequently the cycling capacity retention. However, as the capacity retentions of the Ti0.8Zr0.2V2.7Mn0.5Cr0.8Ni0.75 and Ti0.8Zr0.2V2.7Mn0.5Cr0.8Ni1.75 electrode alloys obtained from our experiment are 5.5 and 89.2%, respectively, it is difficult for us to believe that so huge a difference comes all from the

Ti-V-Based Hydrogen Storage Electrode Alloys

Figure 6. Discharge curves of the Ti0.8Zr0.2V2.7Mn0.5Cr0.8Ni0.75 electrode alloy at different cycles.

intrinsic degradation. In addition, further investigations on the discharge curves of the Ti0.8Zr0.2V2.7Mn0.5Cr0.8Ni0.75 alloy electrode (Figure 6) show that, in the initial stage of cycling, the discharge potential plateau is wide and flat, and the discharge capacity is large. But as the cycle number increases, the discharge plateau potential declines gradually, and the discharge capacity decreases continuously. If only the intrinsic capacity degradation factors affect the process, only the discharge capacity would decrease with the discharge plateau potential remaining almost unchanged. There should be other factors which affect capacity degradation here. Besides the factors studied above including the segregation/ dissolution of hydrogen absorption elements, the oxidation/ corrosion of active element, and the presence of irreversible hydrogen, some other factors, such as the pulverization of the alloy particles and the increase of the contact resistance, the reaction resistance, and the polarization resistance which belong to the extrinsic degradation factors, should be investigated also. An analysis of the extrinsic sources of capacity degradation of the Ti-V-based electrode alloys is made in the following section. 4.3. Extrinsic Degradation. In this paper, the extrinsic degradation is defined as the discharge capacity decay caused

J. Phys. Chem. C, Vol. 112, No. 42, 2008 16687 by the decrease in the rate of utilization of active components due to the decline of kinetic parameters. The oxidation/corrosion of the active elements of hydrogen storage electrode alloys generally not only lead to a direct loss of the intrinsic discharge capacity of alloys but also lead to the formation of an oxide layer on the surface of alloy particles which lowers the surface reaction activity and the ionic conductivity and subsequently leads to a decrease in the utilization rate of the active components and decay of the apparent discharge capacity.1,17,18,22,23 From this point of view, the formation of the oxide layer on cycling should also be an important extrinsic factor of capacity degradation. For realizing the surface state of the cycled alloy, the contact resistance and the surface reaction resistance of Ti0.8Zr0.2V2.7Mn0.5Cr0.8Ni0.75 and Ti0.8Zr0.2V2.7Mn0.5Cr0.8Ni1.75 alloy electrodes after 15 and 120 cycles were measured by EIS technique (Figure 7). It can be seen that all EIS spectra consist of a smaller semicircle in the high-frequency region (semicircle A) and a larger semicircle in the low-frequency region (semicircle B) followed by a straight line. The radius of semicircle A represents the particle-particle contact resistance, and that of semicircle B represents the surface reaction resistance as proposed by Kuriyama et al.30 As shown in Figure 7, the radii of both semicircles A and B are increased noticeably after 120 cycles, especially striking for semicircle B, indicating the large increase in the particle-particle contact resistance and the surface reaction resistance. The pulverization of the alloy particles and the formation of the oxide layer worsen the particle-particle contact and increase the contact resistance. In addition, the formation of the oxide layer not only increases the surface reaction resistance but also retards the transfer of hydrogen atoms. These extrinsic sources are responsible more and more for the capacity degradation as cycling goes on. It is worthy noticing that semicircle B of the Ti0.8Zr0.2V2.7Mn0.5Cr0.8Ni0.75 alloy electrode is very noticeably augmented after 120 cycles, and that of the Ti0.8Zr0.2V2.7Mn0.5Cr0.8Ni1.75 alloy electrode is augmented only slightly, which is in good agreement with the discrepancy in their cycling capacity retentions. The pulverization of alloy particles results in many smaller particles with a large amount of cracks in the alloy particles and subsequently an increase in the particle-particle contact resistance and lowering of the conductivity of each particle

Figure 7. EIS of Ti0.8Zr0.2V2.7Mn0.5Cr0.8Ni0.75 and Ti0.8Zr0.2V2.7Mn0.5Cr0.8Ni1.75 alloy electrodes after 15 and 120 cycles.

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Figure 8. Cycling stability curves of Ti0.8Zr0.2V2.7Mn0.5Cr0.8Ni0.75 (a) and Ti0.8Zr0.2V2.7Mn0.5Cr0.8Ni1.75 (b) electrode alloys with different particle sizes.

regardless of the high intrinsic discharge capacity. Pulverization of alloy particles is a typical and important extrinsic factor of degradation. To evaluate the influence of pulverization, the electrochemical performances of Ti0.8Zr0.2V2.7Mn0.5Cr0.8Ni0.75 and Ti0.8Zr0.2V2.7Mn0.5Cr0.8Ni1.75 alloys with different particle sizes were measured. Figure 8 shows the relationship between the discharge capacity and the cycle number of these two alloys. In the particle size range of 60-140 µm, the cycling capacity retention of the Ti0.8Zr0.2V2.7Mn0.5Cr0.8Ni0.75 alloy improves with decreasing particle size, while that of the Ti0.8Zr0.2V2.7Mn0.5Cr0.8Ni1.75 alloy remained almost unchanged. The particle size obviously affects the cycling stability of the Ti0.8Zr0.2V2.7Mn0.5Cr0.8Ni0.75 alloy. Figure 9 shows the SEM micrographs of the Ti0.8Zr0.2V2.7Mn0.5Cr0.8Ni0.75 electrode alloys with different particle sizes after cycling 30 times. It can be seen that the microcracks occurred in all alloy particles after 30 cycles for the particle sizes tested in this work. Pulverization is more serious for the larger alloy particles. The alloy with smaller particles possesses a higher antipulverization ability. This idea agrees well with the finding of Seta and Uchida that the electrodes made of smaller alloy particles relax better after the compressive stress occurred during the hydriding process.31 To further elucidate the role played by Ni in Ti-V-based alloys, the SEM micrographs of the Ti0.8Zr0.2V2.7Mn0.5Cr0.8Ni0.75 and Ti0.8Zr0.2V2.7Mn0.5Cr0.8Ni1.75 alloys with the starting particle size of ∼140 µm after cycling 80 times were observed as shown in Figure 10. Apparently, the particle of the Ti0.8Zr0.2V2.7Mn0.5Cr0.8Ni0.75 alloy breaks into grains even smaller than 20 µm after 80 cycles, while the Ti0.8Zr0.2V2.7Mn0.5Cr0.8Ni1.75 alloy particle is much less pulverized, which is responsible for its better cycling stability. 4.4. Interplay of Intrinsic and Extrinsic Factors. To sum up, the oxidation/corrosion of active components and the pulverization of alloy particles are the two primary factors for the cycling capacity degradation of Ti-V-based electrode alloys. The pulverization of the alloy particles is an important extrinsic factor for capacity degradation, while the oxidation/corrosion of the active elements is a typical and important intrinsic factor of capacity degradation. The pulverization of the alloy particles creates many cracks in the alloy particles and cleaves in the electrode as shown in Figure 10, which intensifies the flushing

action of the alkaline electrolyte through the cleaves of the electrodes and cleaves in the alloy particles during cycling, increasing thereby the contact of the fresh alloy surface with the alkaline electrolyte and, consequently, the oxidation/ corrosion rate of the active elements. Thus, the pulverization of the alloy particles accelerates also the loss of the intrinsic discharge capacity. At the same time, the oxidation/corrosion of the active elements forms oxide layers on the surfaces of the alloy particles, increasing the particle-particle contact resistance and the surface reaction resistance and, accordingly decreasing the utilization of the active components and the apparent capacity besides deteriorating the surface state of the alloy electrode. In this way, the intrinsic degradation promotes the extrinsic degradation. From the above discussions, we certainly believe that the apparent capacity degradation of Ti-V-based electrode alloys comes from the joint action of the intrinsic degradation and the extrinsic degradation. The oxidation/corrosion of the active elements, the appearance of the irreversible hydrogen, and the segregation/dissolution of the main hydrogen absorption element V are responsible for the decay of the intrinsic discharge capacity of Ti-Zr-V-Mn-Cr-Ni electrode alloys. The pulverization of the alloy particles, the formation of the surface oxide layer, the increase in the particle-particle contact resistance, and the increase in surface reaction resistance all decrease the apparent discharge capacity of Ti-Zr-V-Mn-CrNi electrode alloys. In addition, the pulverization of the alloy particles accelerates the oxidation/corrosion of the active components of the electrode alloy, and subsequently the oxidation/corrosion of the active components forms the oxide layers on the surfaces of the alloy particles. Their interactions accelerate further the decay of the apparent discharge capacity of Ti-Zr-V-Mn-Cr-Ni electrode alloys. The increase of Ni content in the Ti-V-based alloy is found to have a good antipulverization and antioxidation/anticorrosion ability, hence, a good component element for improving the cycling stability of the alloys. 5. Conclusion The factors that affect the capacity degradation of the hydrogen storage electrode alloys can be divided into intrinsic

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Figure 10. SEM micrographs of Ti0.8Zr0.2V2.7Mn0.5Cr0.8Ni0.75 (a) and Ti0.8Zr0.2V2.7Mn0.5Cr0.8Ni1.75 (b) electrode alloys with the particle size of ∼140 µm after 80 cycles.

Figure 9. SEM micrographs of the Ti0.8Zr0.2V2.7Mn0.5Cr0.8Ni0.75 electrode alloys with different particle sizes after 30 cycles: (a) 140, (b) 90, and (c) 60 µm.

factors and extrinsic factors. On the basis of these two types of factors, the capacity degradation mechanism of the Ti-V-based

hydrogen storage electrode alloys was elaborated. It was found that the oxidation/corrosion of the active elements and the appearance of the irreversible hydrogen are two dominating factors for the intrinsic degradation of the discharge capacity on cycling, more important than the segregation and dissolution of the main hydrogen absorbing element V. The extrinsic degradation of Ti-V-based electrode alloys was mainly caused by the pulverization of the alloy particles and the formation of the surface oxide layer, which in turn increases the particleparticle contact resistance and the surface reaction resistance and finally decreases the kinetics of the electrochemical reactions of the electrode alloys. In addition, the pulverization of the alloy particles accelerates the oxidation/corrosion of the active components, which in turn promotes the formation of the oxide layer. These factors jointly work on the decay of the apparent discharge capacity of Ti-V-based hydrogen storage electrode alloys. The increase of Ni content in Ti-V-based hydrogen storage alloys effectively decreases the pulverization of the alloy particles and increases the antioxidation/corrosion ability of the active components and consequently improves the cycling stability of Ti-V-based electrode alloys. The approach to understand the intrinsic/extrinsic degradation mechanism de-

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