Heat-treatment effects of V-based solid solution alloy with TiNi-based network structure on hydrogen storage and electrode properties

ELSEVlER

Journal

uf AlhJys

antl C”mpo”nds

243 (1996) ,33-,3x

Heat-treatment effects of V-based solid solution alloy with TiNi-based network structure on hydrogen storage and electrode properties M. Tsukahara’. K. Takahashi”. T. Mishima”. A. Isomura”, T. Sakd A. 7.arHorhtlm-rito. X”rrvo-ih A!Ch, 4%lopon ‘I,!,&4iMBlWl”l R&Dco. 1

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Vanadium is known as a high capacity hydrogen storage metal, which reacts reversibly around 0.8 MPa a, 327 K as follows [I]: “fiL=“H+

l/ZH,

(1)

V-based allays such as V,TiNi, (.raOS) 12.31 and VzTiNi,, ,,Hf, (y = 0.046 and 0.24) 14.51 have been found to work as a high capacity electrode material for Ni-MH cells. If the reaction of monohydride around 10-l Pa at 353 K [6] is utilized, the total discharge reach to about capacity of V,TiNi, ih would 8WAhkg-‘. Tbe practical capacity of this alloy, however, was about 4M) A h kg-’ be~ausc of the too low dissociation pressure for monohydride. The V,TiNi‘ (x>O.S) based alloys consisted of a V-based solid solution main phase and a TiNi-based secondary phase network for x > 0.25 131. Tbe TiNibased phase precipitated along the grain boundaries of the V-based main phase and toroved a three-dimensional network. Tbe main phase worked as a main hydrogen storage site. Sx secondary phase worked as a micro-current collector and electro-catalyst for electrochemical reactions. With increasing nickel content x in V,lINi,, the volume ratio of the main phase de-

creased and the network of the secondary phase became line and dense. With this change of the rricm+wure. hydrogen capacity was decreased. while high rate capability was improved. Additionally. the increase in the nickel content x also caused an increase in dissociation prescore for Eq. (1) because the titanium content in the main phase wa decreased and it contracted the wit cell of the main phase. Adding bafnium in l-5 at.% to V,TiNi,, (V;FiNi,. .&HI. : Y = 0.046. 0.24) char& the stroctox bf ;he TyE-b&b phase into the (Ti,cf)(V,Ni), Lavestype (C14) st~octwe. This secomlary phase remarkably improved the high rate capability by microaacking inside and around the szcondafy phase during a couple of charge-discharge cycles. The alloy of VjTiNi,s, was easily deteriorated during charge-discharge cycles. By charge-diwbarge cp:ling. cracking in the alloy across the grains and grair, boundaries exposes fresh surface to the KOH sol~~tion. The vanadium in the secondary phase then dissolves into the KOH electrolyte solution sod the secondary phase network collapses [7]. Therefore, the cycle life of this type of alloy electrode is considered to depend strongly oo the properties of the secoodary phase. The compositions of the two phases would be

chunpcd

diffurmn

by kmnill

men,s hcwccn

men& arc considered clectrochrmical will

the

In Ihe oresent

the’ rkcts of V,TW, 5D on microsIrntiu~e. electrochemical properties.

storage and oaocr.

WC of and

details

metals

on a water-coaled

under argon atmorpherc. under vacuum

ele-

heat-ire&

the heat-fre&& hydrogen storsgc

Ingots ot V,TiNx,, (h were prepared conslitucnt

constiluent

Lo change hydrogen

orooerties.

report

L. Experimental

of

,he two phases. Therefore.

(less than

by arc melting of wooer

hearth

The ingots were LeaI-treated IO-’ Pa, a, 973-1473

K for

24 h and suhseq;enlly

cooled to room temperature in vacuum furnace. After the heat-treatment. the surface layer of all ingots was filed off and samples were hydrogenated under hydro&en pressure of 3.3 MPa by heating up to 673 K: they were then pestled into powder. Micmstructure

3. Results and discussion

in the alloys were examined by scannine. electron microscow (SE-M) and rlrctron~ probe X-ray microanaly& (EPMA). Cross-sections perpendicular lo the bottom of two ingots were observed. The crystal slructures were examined by X-ray powder Iiffraclion (XRD) using Cu Kar radi&on fir [he alloy powder evacuated at 673 K for 4 h. Pressure-composition isotherms (Pm curves) for the samples were measured with a Sieverls-type apparatus. Each rample was put in an SUS 316 reactor tube and activated as follows: the reactor was evacuated and heated “II to 673K. and hvdroeen (purity greater than 99.&9%) was admit&i up”to 3.3MPa and then the reactor was cooled to mom temperature. lust before measuring a PCT curve. the reactor was evacuated at 673 K for 4 h in order to get a hydrogen zero point. An alloy electrode was prepared according to the procedure reported previously [Z]. The alloy powder sifted between 63-1136 pm in diameter was coated with 20 mass% copper by the chemical plating method. The copper-coated alloy was mixed with IO mass% FEP binder (D&ken Co. letrafluoroethylene-hexafluoropropylene copolymer), and then the mixture was hotpressed on a nickel mesh at 573K. The electrode performance was examined in a half-cell at 293K &ing an Hg-HgO reference electrode, d 6M KOH solution, and an Ni(OH), eotmter electrode. Each electrode was charged at 1WA kgfor 4 h and discharged to the cut-off voltage -0.7 V vs. Hg-HgO. Discharge capacity vs. discharge current density relations were measured in the range 25-6WA kg-‘. Charge-discharge cycle tests were conducted at a discharge current density of 50 A kg-‘. and composition

of each phase

’

‘IIe XRD patterns in Fig. 1 show that all samples have two body-centered cubic (h.c.c.) phases. The series of stronger peaks for one b.c.c. phase is ascribed to the main phase of the V-base solid solution. The series of weaker peaks for the other b.c.c. phase is ascribed to the TiNi-based secondary phase with the CsCCtype structure. The Ti,Ni phase appears for the heat-treated samples, and the intensity of these peaks increases with increasing heat-treatment temperature T,. Fig. 2 compares the cross-sections of the samples heat-treated at T, =973K. 1273K and 1473K for 24 h. The as-cast alloy L’,TiNi,,,, consisted of V-based solid solution main phase and TiNi-based secondary phase, where Ihe secondary phase precipitated along the grais boundaries of the main phase and formed a nelwork structure [S]. The network structure of the secondary phase is preserved for the samples treated at T, s 1273 K. The sample treated at T, = 1473 K loses the network sbucture and it consists of at least three phases: a main phase. a rounded-shape phase (secondary phase) and a belt-like shaped phase (1~ tiary phase). The SEM-EPMA observation for the alloy treated at T, = 1473 K showed that the compositions of the three phases were V,,Ti,,Ni, for the main phase, V”Ti,,Ni,, for the secondary phase and V,,Ti,,Ni,,for the tertiary phase. From the compositional data and the results of the XRD analysis. those phases would be the V-based solid solution, the TiNi-based phase and the Ti,Ni-based phase respectively. According to the binary phase diagram of Ti-Ni system [a], the melting point of Ti,Ni (1261 K) is lower than T, = 1473 K, while that of vanadium (2183 K) is much higher than

boundaries in a belt-like shape. In the alloy samples hrnl-trcared at T. < 1273 K. the Ti,Ni-based phase was not observed clearly by SEM, although the very small XRD peaks due to the Ti,Ni-based phase were observed (Fig. 1). In Fig. 3 the average grain width of the main phase increases with increasing heat-treatment temperature T,, except for the sample treated at T, = 1473 K. where the networii s1ructurc is lost. The grain width is doubled by lhe heal-treatmenl at T. = 1173 K and 1273 K. It is considered that the constituent elements were considerably diffused tbmugb the grain and grain boundary and several grains are joined to each other. Tbe PCf curves of dewmion at 353K for the areshownin Fig. 4. as-cast and heat-treated sampies The PCT curve for the alloy treated at 973 K is just the same as that for the assast alloy with a sloped pressure plateau, while the heat-treatment at T, = 1073-1273 K brings abOut a Ratter pressure plaleau. The beat-trearment a, T. a 1073 K lmered the dissocialion pressure compared with that of the as-cast and heat-treated at 973 K xamples

Q,

the treatment temperature. The TiNi-based phase would be decomposed at 1473 K lo form the Ti,Nibased phase, though the treatment temperature is lower than the melting point of TiNi (lB3K). It is considered that during the heal-Imatment at T. = 1473 K the alloy would be separated into solid and liauid phases and the mains of the V-bared solid &I& phase would g&v. During the cooling pmcess. the TiNi-based ohase would be solidified into a rounded shape. Subs&ently, the Ti,Ni-based phase would be solidified and precipitated on grain

Fig. 5 shows ,ha, the change of ,he average dissociation pressure t’,, and the unit cell dimension u of ,hr main pbasc wi,b hca,-treatmem temperalure T,. The F’,, value was obtained from PCT corvrs a, 353 K (Fig. 4). The u value was obtained from XRD analysis (Fig. I). The II value incrcaxs with increasing lrealmcn, temperature until 1173K and then i, gradually decreases until 1473 K. The Pd value decreases rapidly with increasing heal-trealmen, temperature from 973 10 1173K and it gradually increases in fhe range 1173%1473K. The inverse relation ol the Pd and a ~a,ocs shows lha, the increase in the unit cell dimension of the main phase causes the decrease in the disswialion pressure. In Fig. 6 the relation between the hydrogen capacity C,,, and volume fraction of Ihc main phaw D,,,, is shown. where C,,,, was defined as the hydrogen content obtained hclween 3.3-0.01 MPa a1 353 K and D ~l,n was obtained from the SEM micrographs by areal analysis. Both D “,“,” and C ,,,.,I values increase

with increasing trcatmen, lempcrature T,, up lo L173 K and dccrcasc ahovc this temperature because of the prccipitalcs 01 the third phase. Compositional change ol the main and the seeondary phases with the treatment temperafure is shown in Fig, 7. With increasing heal-lrcatment tcmpentore, the lltanium content in the main phase slightly incrcascs. while the vanadium content decreases. The nickel conlen, in the main phase is almob constan, up ,o T,, = II73 K. This compositional change would cause the increase in the unit cell dimension of the main phase. Further increase in the heal-treatment temperalnrc lo 1473K causes an increase in nickel content, which results in conlraclion oC the unit cell oi tbc main phase (Fig. 5). In the secondary phase the vanadium content decreases with increasing treatmen, temperature up lo I173 K and becomes cons,an, al higher temperature. Fig. X shows the discharge capaaty vs. the discharge current density plots for Ihc as-CBS, and heat-treated samples. The discharge Fapacifirs a, 200A kg-’ for Ihe samples trcatcd a, T,, = 1273 and 1473 K are lower than lha, of the samples lrealed a, T,, S II73 K. When the alloys were treated B, T., s 1273 K, the grain width reached more than 3Opm (Fig. 3). and the secondary phase network swocturc was damaged (Fig. 2). The

decrease in the high rate capability would be explained by the degradation of the secondary phase nework as a micro-current collector and ekctro-catalyst. Fig. 9 coo~oares caoacitv decav c”wes durinec charge-discharge cycles for electrodes using as-cast and heat-treated samoles at T. = 1073 and ,173 K. The sample heat-w&d at T,, =iO73 K shows the best durability among these samples. In Fig. 10 the crosssectional views of thcsc electrodes after ihe chargedischarge cycle test are shown. The secondary phase in the as-cast and heat-treated at I, ,173 K samples almost disappeared. as described previously [4]. while the secondary phase was somehow presewud in the alloy heat-treated a‘ T,, = 1073 K. During the chargedischarge cycles, vanadium in the secondary phase could be easily dissolved into the alkaline electrolyte. damaging !he network str”ct”re of the secondary phase. The improvement of durability by the heat-

.,

.

=

lreatmrnt at T, = 1073 K would be ascribed to the decrease in the vanadium content in the secondary phase (Fig. 7(b)). The alloys heal-treated at r, = 1173 K. however. show poorer cycle stability than that of the sample treated at T, = 1073 K, in spite of the lower vanadium c”“tent in the secondary phase. This alloy treated at T. = 1173 K shows a lesser volwne fraction of the secondary phase (Fig. 6) end a larger grain wdth of the main phase (Fig. 3) than those of

the sample treated at IO73 K. Therefore. the network s,ructure of ,he secondary phase for T, = 1173K would he more vulnerable to electrical cycles than that for T, = 1173 K.

4 Conelusionr The heat-treatment changed the micmstxcture ot the alloy V,TiNi,,,.. and it improved the hydrogen storage and electrode properties. The heat-treatment at T,, a 1073 K brought about a higher volume fraction of the main phase. and then a Ratter pressure plateau and a larger hydrogen storage capacity. It also reduced the vanadium content as a corrosive constituent in the secondary phase, improving the durability for chargedischarge cycles. The heat-treatment at 7’,aP 1173 K, however, enlarged the grain size of the main phase and damaged the network structure of the secondary abase, which caused a decrcase in the hiah rate iapability and the cycle bfe. Upon heat-treat&t at 1473 K the network structure was comoletelv lost with the precipitation of the tertiary phas; of ?i,Ni. The

most suitable microstructure for the electrode was abtaincd far the heat-treatment around 1073 K.