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International Journal of Hydrogen Energy 29 (2004) 313–318www.elsevier.com/locate/ijhydene

Degradation mechanisms of Ti–V-based multiphase hydrogenstorage alloy electrode

Yunfeng Zhu, Hongge Pan∗, Mingxia Gao, Yongfeng Liu, Rui Li,Yongquan Lei, QidongWang

Department of Materials Science and Engineering, Zhejiang University, Hangzhou 310027, People’s Republic of China

Accepted 2 June 2003

Abstract

The degradation mechanisms of the Ti–V-based multiphase hydrogen storage electrode alloy (Ti0:8Zr0:2)(V0:533Mn0:107Cr0:16Ni0:2)2 during electrochemical cycling in alkaline electrolyte have been studied systematically by using XRD, SEM, XPS,AES and EIS measurements. The results show that the irreversible hydrogen absorbed in the alloy increases with cycling. Thepulverization and oxidation/corrosion of the alloy during cycling are two main factors responsible for the fast degradation ofthe electrode. It was found that a passive Ti oxide Alm formed on the alloy surface, leading to an increase in charge transferresistance on the alloy surface and a decrease in exchange current density I0.? 2003 International Association for Hydrogen Energy. Published by Elsevier Ltd. All rights reserved.

Keywords: Hydrogen absorbing materials; Electrode materials; Electrochemical reactions; Degradation mechanisms; X-ray diCraction

1. Introduction

Recently, hydrogen storage electrode alloys have beenextensively studied, including the rare earth-based AB5=AB3

type alloy, Ti- and Zr-based AB2 type alloy, Mg-basedalloy and V-based solid solution alloy, among which therare earth-based AB5 type alloy was the most successful oneand has been commercialized [1–9]. However, the conven-tional AB5 type alloy has a limited discharge capacity of280–320 mAh=g, and the space for further increasing thedischarge capacity is very small because of its low theo-retical electrochemical capacity. Therefore, the alloys withhigher capacity (400–600 mAh=g), such as the Ti/Zr-, Mg-and V-based alloys were developed. Although these alloyshave a very high discharge capacity, they have not beencommercialized due to several reasons, in which the cyclingstability is considered as an important factor, especially forthe Mg-based alloy and the V-based solid solution alloy.

∗ Corresponding author. Tel.: +86-571-8795-2576;fax: +86-571-8795-2615.

E-mail address: honggepan@zjuem.zju.edu.cn (Hongge Pan).

So it is necessary to study the degradation mechanisms of thehydrogen storage alloy electrodes during cycling in alkalineelectrolyte.

It was reported that the pulverization of alloy particlesduring charging–discharging cycles was responsible for thecapacity degradation of the rare earth-based AB5 type alloys[1,10]. Yu et al. [11] studied the Ti-based hydrogen stor-age electrode alloys and showed that the fast degradationwas mainly caused by the deterioration of surface proper-ties followed by the growth of Ti-oxide on the surface. Thecycle life was improved by partial substitution of Mn by Crin the alloy, which reduced the surface oxidation and pul-verization rate of the alloy. Studies of the Mg-based alloyand V-based solid solution alloy indicated that the drasticdegradation of discharge capacity was ascribed to the seri-ous corrosion of alloy constituents Mg and V in the alkalineelectrolyte during cycling [12,13].

In our previous study [14], we reported a series of multi-phase Ti–V-based hydrogen storage electrode alloys(Ti0:8Zr0:2)(V0:533Mn0:107Cr0:16Ni0:2)x (x=2; 3; 4; 5; 6), whichseems very promising. It was found that these alloys weremainly composed of a C14 Laves phase with hexagonalstructure and a V-based solid solution phase with BCC

0360-3199/$ 30.00 ? 2003 International Association for Hydrogen Energy. Published by Elsevier Ltd. All rights reserved.doi:10.1016/S0360-3199(03)00153-8

314 Y. Zhu et al. / International Journal of Hydrogen Energy 29 (2004) 313–318

structure. In this paper, we will study mainly the degradationmechanisms of the (Ti0:8Zr0:2)(V0:533Mn0:107Cr0:16Ni0:2)2alloy.

2. Experimental

The (Ti0:8Zr0:2)(V0:533Mn0:107Cr0:16Ni0:2)2 alloywas pre-pared by induction levitation melting of the constituentmetals on a water-cooled copper crucible under argon at-mosphere. The ingot was turned over and remelted twiceto ensure a higher homogeneity. Part of the alloy was me-chanically crushed and ground to the powder of 300 meshsize (¡ 47 �m).

The test electrode was prepared by cold pressing 500 mgpure alloy power under a pressure of 800 MPa into a pel-let of 10 mm diameter and about 1:5 mm thickness. Thenthe pellet was sandwiched within two foamed Ni plates(60 × 20 mm2) with a Ni wire soldered on to form anegative electrode. Electrochemical measurements wereperformed in a half-cell consisting of a working electrode(MH electrode), a sintered Ni(OH)2=NiOOH counter elec-trode and a Hg/HgO reference electrode. The electrolytewas a 6 M KOH solution, controlled at 30 ± 1◦C. Inorder to reduce the IR drop during the electrochemicalimpedance spectroscopy measurements, the Hg/HgO ref-erence electrode was equipped with a Luggin tube. Theelectrode was charged at 60 mA=g for 6 h followed by10 min rest and then discharged at 60 mA=g to the cut-oCpotential of −0:6 V vs. the Hg/HgO reference electrode.Electrochemical impedance spectroscopy (EIS) measure-ments were conducted at 50% depth of discharge (DOD)using Solartron SI1287 Electrochemical Interface with a1255B Frequency Response Analyzer. The EIS spectra ofthe electrode after 10 and 30 cycles were obtained in thefrequency range of 100 kHz–5 mHz with an ac amplitudeof 5 mV under open-circuit conditions. The linear polar-ization curves of the electrode after 10 and 30 cycles weremeasured by scanning the electrode potentials (SolartronSI1287 potentionstat) at 0:1 mV=s from −5–5 mV (vs.open-circuit potential) at 50% DOD.

The electrode was taken out after a certain number ofcharging–discharging cycles (10 and 30 cycles) and thenwashed with distilled water followed by drying in vacuum.XRD, SEM, XPS and AES measurements were performedon the electrode before and after cycling. The crystal struc-ture and lattice parameters of the alloy were determined byXRD measurements, which were performed on a PhilipsdiCractometer with Cu K� radiation at 40 kV and 45 mA.For investigation of the surface state of the electrode, XPSmeasurements were carried out on a PHI-550 type spec-trometer at 10−8 Torr using Al K� radiation. Auger electronspectroscopy (AES) depth proAles were measured for inves-tigating the elemental distribution on the electrode surfaceduring cycling by using a PHI-550 type electron spectrome-ter with an electron beam at 3 kV and 10 �A. The electrode

surface was sputtered with Ar+ on an area of 1:5×1:5 mm2

at 4 kV and 15 mA. The KOH solution was analyzed after10 and 30 charging–discharging cycles for investigating thespecies and contents of alloy constituents dissolved into thealkaline electrolyte by using a 721-type spectrophotometer.

3. Results and discussion

For investigating the changes in crystal lattice of the alloyduring cycling, XRD measurements were performed on theelectrode before and after 10 and 30 cycles at the dischargedstate. Fig. 1 shows the XRD patterns of the electrode. It isfound that the as-cast alloy is mainly composed of a C14Laves phase with hexagonal structure and a V-based solidsolution phase with BCC structure. After cycling, the phasestructures are not changed. However, the diCraction peaksof both phases are moved to the lower angle due to the vol-ume expansion of the crystal lattice upon hydrogenation.Besides, it can also be found that the peaks are broadenedafter cycling, especially for the V-based solid solution phase.The lattice parameters and unit cell volumes of both the C14Laves phase and the V-based solid solution phase are calcu-lated and presented in Table 1. It can be seen that the elec-trode expands with cycling due to the hydrogen absorbed inthe alloy. The volume expansion ratio PV=V of both phasesafter 10 and 30 cycles are also calculated. As shown in thetable, the PV=V of both phases increases with cycling. After30 cycles, the PV=V are 12.79% for the C14 Laves phaseand 11.89% for the V-based solid solution phase, respec-tively. The results indicate that a large amount of hydrogenis absorbed in the alloy and cannot be desorbed, which leadsto a large volume expansion of the alloy and also results ina decrease in discharge capacity of the alloy electrode.

Fig. 2 shows the SEM micrographs of the electrodesurfaces before and after cycling. It is found that the alloyparticles are pulverized with cycling due to the volumeexpansion and shrinkage during hydrogenation and dehy-drogenation. The pulverization of the alloy leads to largeamount of cracks in the alloy and interstices in the elec-trode, which increases the contact resistance and decreasesthe conductivity between alloy particles. This may beresponsible for the large amount of irreversible hydrogenabsorbed in the alloy with further cycling.

For investigation of the chemical state of the alloyconstituents on the electrode surface, X-ray photoelectronspectroscopy (XPS) measurements were performed on theelectrode surface before and after cycling. Fig. 3 showsthe core level spectra of Ti2p, Zr3d, V2p, Mn2p, Cr2p andNi2p of (Ti0:8Zr0:2)(V0:533Mn0:107Cr0:16Ni0:2)2 alloy elec-trode. It is found that all six elements can be detected onthe alloy surface. As judged from peak positions, the ele-ments of Ti, Zr, V, Mn, Cr and Ni on the alloy surface areentirely in the oxidized state. During cycling in the alkalineelectrolyte, the binding energy for Ti2p, Zr3d, V2p, Mn2p,Cr2p and Ni2p all increase, especially for the electrode after

Y. Zhu et al. / International Journal of Hydrogen Energy 29 (2004) 313–318 315

Fig. 1. XRD patterns of the (Ti0:8Zr0:2)(V0:533Mn0:107Cr0:16Ni0:2)2 alloy electrode before and after cycling.

Table 1The lattice parameters and unit cell volumes of alloy phases in the (Ti0:8Zr0:2)(V0:533Mn0:107Cr0:16Ni0:2)2 alloy

Samples Phase Lattice parameters ( QA) Cell volumes ( QA3) PV=V (%)

As-cast C14 a = 5:023 c = 8:190 179.0 —BCC a = 3:012 27.33 —

10 Cycles C14 a = 5:168 c = 8:449 195.4 9.16BCC a = 3:095 29.65 8.49

30 Cycles C14 a = 5:228 c = 8:529 201.9 12.79BCC a = 3:127 30.58 11.89

Fig. 2. SEM micrographs of the electrode surface for the (Ti0:8Zr0:2)(V0:533Mn0:107Cr0:16Ni0:2)2 alloy: as-cast (a); 10th cycles (b); 30thcycle (c).

30 charging–discharging cycles, which indicates that thealloy constituents on the electrode surface are oxidizedconstantly during cycling.

The oxidation of alloy constituents generally forms a pas-sive oxide Alm on the electrode surface, such as the TiO2,

ZrO2 and Cr2O3 Alm. However, the oxides cannot absorband desorb hydrogen, so that the alloy constituents on thealloy surface lose the hydrogen storage capability. Besides,the passive Alm has lower conductivity and electro-catalyticactivity, which would decrease the electrochemical

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Fig. 3. XPS spectra of Ti2p, Zr3d, V2p, Mn2p, Cr2p and Ni2p of the (Ti0:8Zr0:2)(V0:533Mn0:107Cr0:16Ni0:2)2 alloy.

hydrogen reaction rate on the electrode surface. Further-more, the passive Alm would also act as a barrier for thehydrogen penetration through the Alm into the inner alloy.All these may lead to a decrease in discharge capacity ofthe alloy electrode with cycling.

Fig. 4 shows the AES depth proAles of the alloy electrodebefore and after cycling. The appearance of K in the surfacelayer of the alloy after cycling is due to the immersion ofelectrode into KOH solution during cycling. It can be seenthat the oxygen content is increased constantly in the alloysurface layer with cycling. The high oxygen content on thesurface of the as-cast alloy may be ascribed to the adsorbedoxygen during exposure to air of the alloy surface. However,it is decreased drastically inside the alloy. After cycling, it isfound that the oxygen content is greatly increased inside thealloy. Combined with the XPS analysis, it is concluded thatthe alloy constituents are oxidized on the surface and theoxidation gradually develops into the inner part of the alloy.Besides oxygen, it can also be found that the Ti contenton the alloy surface is increased with cycling, which makesus believe that a Ti oxide Alm forms on the alloy surface.A comparison of Ni content on the alloy surface showsthat it is decreased with cycling. As we known, Ni has ahigh electro-catalytic activity in the alkaline electrolyte. Sothe decrease in Ni content on the alloy surface may alsodecrease the electrode kinetics of hydriding and dehydriding,which would also be responsible for the degradation of theelectrode.

As the main hydrogen absorbing elements in the al-loys, V, Ti and Zr are analyzed to determine the amount

of each element dissolved into the KOH solution dur-ing cycling. Table 2 shows the results of the (Ti0:8Zr0:2)(V0:533Mn0:107Cr0:16Ni0:2)2 alloy electrode after 10 and 30charging–discharging cycles. It is found that all three ele-ments can be detected in the KOH solution after 10 cycles,which indicates that the alloy constituents of V, Ti and Zrare corroded and dissolved into the KOH solution duringcycling. The dissolved amounts of V, Ti and Zr elementsare increased from 44.93, 17.92 and 9:24 �g=ml to 133.33,25.47 and 14:20 �g=ml after 30 cycles, which indicates thatthe corrosion of alloy constituents proceeds constantly. Thedissolved elements certainly lose the hydrogen storage ca-pability. Therefore, the corrosion of the alloy constituentsduring cycling may also result in a decrease in dischargecapacity of the alloy electrode.

Fig. 5 shows the electrochemical impedance spectra ofthe (Ti0:8Zr0:2)(V0:533Mn0:107Cr0:16Ni0:2)2 alloy electrodeat 50% DOD. According to the analysis model proposedby Kuriyama et al. [15], the larger semicircle in thelow-frequency region is attributed to the charge transferresistance on the alloy surface. As can be seen in the Agure,the radius of the larger semicircle in the low-frequencyregion is greatly increased as the cycle number increasesfrom 10 to 30, which indicates that the charge transferresistance on the alloy surface is greatly increased. The in-crease in charge transfer resistance may be ascribed to theoxidation of surface constituents followed by the formationof a passive oxide Alm on the alloy surface.

Fig. 6 shows the linear polarization curves of the(Ti0:8Zr0:2)(V0:533Mn0:107Cr0:16Ni0:2)2 alloy electrode at

Y. Zhu et al. / International Journal of Hydrogen Energy 29 (2004) 313–318 317

Fig. 4. AES depth proAles of the (Ti0:8Zr0:2)(V0:533Mn0:107Cr0:16Ni0:2)2 alloy: as-cast (a); 10th cycle (b); 30th cycle (c).

Table 2The dissolution amount of V, Ti and Zr elements of the(Ti0:8Zr0:2)(V0:533Mn0:107Cr0:16Ni0:2)2 alloy electrode after 10 and30 charging–discharging cycles

Sample V (�g=ml) Ti (�g=ml) Zr (�g=ml)

10 Cycles 44.93 17.92 9.2430 Cycles 133.33 25.47 14.20

50% DOD. The exchange current density I0 of the alloyelectrode can be calculated by the following formula [3]:

I0 =IRTF�

; (1)

where R; T; F; � denote the gas constant, the absolute tem-perature, the Faraday constant and the overpotential, re-spectively. The results show that the I0 is decreased from46.05 to 14:64 mA=g as the cycle number increases from 10to 30.

Fig. 5. Electrochemical impedance spectra of the (Ti0:8Zr0:2)(V0:533Mn0:107Cr0:16Ni0:2)2 alloy electrode measured at the 50%DOD.

318 Y. Zhu et al. / International Journal of Hydrogen Energy 29 (2004) 313–318

Fig. 6. Linear polarization curves of the (Ti0:8Zr0:2)(V0:533Mn0:107Cr0:16Ni0:2)2 alloy electrode measured at the 50% DOD.

4. Conclusions

In this paper, the degradation mechanisms of the Ti–V-based multiphase hydrogen storage electrode alloy(Ti0:8Zr0:2)(V0:533Mn0:107Cr0:16Ni0:2)2 during electrochemi-cal cycling in alkaline electrolyte have been studied syste-matically. It is concluded that the capacity degradation ofthe alloy electrode is mainly due to the pulverization andoxidation/corrosion of the alloy during cycling, which leadsto an increase in charge transfer resistance on the electrodesurface and a decrease in exchange current density of thealloy electrode.

Acknowledgements

This work was supported by National Natural Foundationof China under Contract No. 50271063.

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