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The cycle life and surface properties of Ti-based AB2 metal hydride electrodes
Journal of Alloys and Compounds 337 (2002) 214–220
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The cycle life and surface properties of Ti-based AB 2 metal hydride electrodes a, b a b b Yan-hui Xu *, Chang-pin Chen , Xiao-lin Wang , Yong-quan Lei , Qi-dong Wang b
a Department of Chemical Engineering, Tsing Hua University, Beijing 100084, PR China Department of Materials Science and Engineering, Zhejiang University, Hangzhou 310027, PR China
Received 21 August 2001; accepted 10 October 2001
Abstract In this paper, Ti-based AB 2 Laves phase hydrogen storage electrode alloys are prepared and their electrochemical cycle stability is studied, including Ti 0.7 Zr 0.5 V0.2 Mn 1.82x Ni x (x50.4, 0.8, 1.1, 1.4 and 1.7) and Ti 1.0 Zr 0.2 Cr 0.4 Ni 0.8 V0.82x Al x (x50.0, 0.2, 0.4 and 0.6) alloys. Energy disperse spectroscopy (EDS) and X-ray photoelectron energy spectroscopy (XPS) were used to analyze the changes of the elements on the alloy grain surface before and after charging–discharging cycles in a KOH solution in order to understand the nature of the decrease of discharge capacity with cycling. The experimental results show that the linear-polarization resistance increases when cycling. The segregation, oxidization and dissolution of some elements in the alloy are key factors influencing the cycle stability and are also a main problem that needs to be overcome in order to use these materials as negative electrode in a commercial nickel–metal hydride rechargeable battery. 2002 Elsevier Science B.V. All rights reserved. Keywords: Transition metal alloys; Hydrogen-absorbing materials; Electrode materials; Gas–solid reaction; X-Ray diffraction
1. Introduction In recent years, extensive investigations on metal hydride alloys have been carried out. The latter are used as negative electrodes of nickel–metal hydride batteries because of their low environmental impact, compared with Cd negative electrode [1–5]. AB 5 type rare-earth-based hydrogen storage alloys are the main active materials in commercial Ni–MH batteries and their discharge capacity can reach about 330 mAh / g. It is difficulty to further increase their discharge capacity, which is restricted by their theoretical hydrogen storage capacity. Cobalt is the most expensive element and takes up about 45% of the total cost of the Mm(NiCoMnAl) 5 alloy raw materials. Therefore, research and development of new Mm-based alloys with low cobalt or without cobalt is a new direction for AB 5 -type metal hydride alloys in order to lower the Ni–MH battery costs [1]. Zr-based AB 2 Laves phase alloys have high discharge capacity, long cycle life but are difficult to activate. For the near future it will be difficult to use Zr-based AB 2 alloys because of their high cost and low high-current discharge ability. In the past 10 years, *Corresponding author. E-mail address: [email protected] (Y.-h. Xu).
Mg–Ni-based and V–Ti-based metal hydride alloys have been given much attention because of their very high discharge capacity [2–5]. Now the discharge capacity of Mg-based alloy electrode increased from about 170 mAh / g to about 1000 mAh / g (Mg 2 Ni alloy) for Mg 2 Ni1 Ni(1:1.28). But the cycle life of Mg-based alloy electrode is very poor, as reported by Ye et al. [5]. The electrochemical behavior of V–Ti-based solid solution alloy electrode is similar to that of the Mg-based alloy: easy activation, high capacity and poor cycle life [2]. Although adding Cr, Co, Nb, Ta, C, Zr and Hf could improve the electrode properties to some degree, it still cannot satisfy commercial requirements. Vanadium is also an expensive material. According to the corresponding phase diagram or standard electrode potential, the most stable form of V and Mg in alkaline solution is their oxides. Upon cycling, oxidation or dissolution of some elements, especially V and Mg, caused a decrease in discharge capacity. We consider that Mg-based and V-based alloys cannot be used as negative electrode active materials for commercial Ni–MH battery. Several authors [6–10] reported on the electrochemical behavior of Ti–Ni-based alloys. The discharge capacity of TiNi alloy electrodes reached 210–250 mAh / g at 5 mA / g current density and reached 180 mAh / g at 50 mA / g. Gutjahr et al. [6] found that a sintered mixture of
0925-8388 / 02 / $ – see front matter 2002 Elsevier Science B.V. All rights reserved. PII: S0925-8388( 01 )01917-X
Y.-h. Xu et al. / Journal of Alloys and Compounds 337 (2002) 214 – 220
Ti 2 Ni and TiNi phases possessed reversible capacities above 300 mAh / g. The discharge capacity of Ti 3 Ni 2 alloy, which is a composite of the TiNi phase and the Ti 2 Ni phase, increases with increasing temperature [10]. The cycle life of this type of alloy needs further improvement when it is to be used as negative electrode material. Ti-based AB 2 metal hydride alloys are another type of potential candidate because of their easy activation, large hydrogen storage capacity, and very low cost. Many researches [11,12] investigated their electrochemical behavior and found that the key problem which needs to be overcome is the increase of the cycle life when used as negative electrode materials. Many researches investigated the reason for the poor cycle life of some metal hydride alloys using electrochemical techniques or physical analysis methods. According to several reports [13–16], the pulverization, the corrosion of the alloys and the segregation and dissolution of some of the elements (especially V, Mn, Al and rare earth elements) are the main factors influencing the cycle life of metal hydride electrodes. Previous investigations [11] have shown that the discharge capacity of the Ti 0.7 Zr 0.5 V0.2 Mn 1.82x Ni x is 399.1 mAh / g at 323 K at first cycle and decreases fast with cycling. For example, when x is 0.4 the discharge capacity decreases to about 20 mAh / g after three cycles. In this paper, the cycle life of the Ti 0.7 Zr 0.5 V0.2 Mn 1.82x Ni x (x50.4, 0.8, 1.1, 1.4 and 1.7) and Ti 1.0 Zr 0.2 Cr 0.4 Ni 1.8 V0.82x Al x (x50.0, 0.2, 0.4 and 0.6) metal hydride electrodes was investigated. Physical analysis techniques were used in order to find the reason why the discharge capacity of Ti-based hydrogen storage alloy electrodes decreases fast with cycling. We used X-ray diffraction (XRD), energy disperse spectroscopy (EDS) and X-ray photoelectron energy spectroscopy (XPS).
2. Experiment The purity of all used original materials is higher than 99.5%. The ingots were prepared by high-frequency induction melting in a water-cooled boat under argon atmosphere and repeated four times in order to ensure homogeneity. Then the ingots were decrepitated into powders for electrochemical measurement and physical analysis. The electrode preparation method is the following: a mixture of the alloy powders and nickel powder (mass ratio is 1:2) was pressed at 1.5 M Pa into a round disk electrode with a diameter of 10 mm and thickness of about 1 mm. The content of the alloy powder is about 200 mg. The reference electrode is Hg / HgO / 6 N KOH and the counter electrode is Ni(OH) 2 / NiOOH. The discharge capacity of which is more than that of the positive electrode. The electrolyte is a 6 N KOH water solution. Electrochemical measurement was carried out in a V-type open glass three-electrode system. The end point of the
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discharge was set at 20.6 V (versus Hg / HgO). The charging–discharging cycle was measured using a DC-5 Battery Test Instrument. A Philips X’pert diffractometer was used to analyze the phase structure of all used alloys. The power is 40 kV320 mA and the scanning rate is 0.048 s 21 . The energy dispersive analysis was performed with a PV9900-Philips Energy Analyzer in order to determine the change in concentration of the various elements before and after electrochemical cycling. The duration time was set at 100 s. The XPS analysis technique, which was conducted using Al Ka radiation and an Escalab MKII photoelectron energy spectrum instrument, was used to determine the chemical state and content of the constituent elements of the alloys. The vacuum in the analysis room was 10 29 Torr. The ray energy was 50 eV.
3. Results
3.1. The cycle life of T-based AB2 metal hydride electrodes The cycle life of Ti 0.7 Zr 0.5 V0.2 Mn 1.82x Ni x (x50.4, 0.8, 1.1, 1.4 and 1.7) hydrogen storage alloys at 283 K is shown in Fig. 1, from which it can be seen that the cycle stability of this alloy electrode increases as the content of nickel increases. The maximum discharge capacity increases from about 110 mAh / g for x50.4 to about 286 mAh / g for x51.4 except for x51.7. The phase composition for the x51.7 alloy is different from the other alloys. Nickel partly substituting for manganese is very effective for improving the discharge capacity as shown in Fig. 1. It was found from our previous investigations that the environmental temperature influences strongly the discharge capacity and cycle life. For example, the discharge capacity of the Ti 0.7 Zr 0.5 V0.2 Mn 1.4 Ni 0.4 alloy increases from 111.7 mAh / g at 283 K to 399.1 mAh / g at 323 K and then decreases to about 260 mAh / g at 353 K. After 10 charging–discharging cycles the discharge capacity for the x50.4 alloy decreases to about 50 mAh / g, but at 323 and 353 K after only three cycles it decreases to about 15 mAh / g [11]. This confirmed that an increasing environment temperature is not beneficial to cycle life. The cycle life curves of Ti 1.0 Zr 0.2 Cr 0.4 Ni 0.8 V0.82x Al x (x50.0, 0.2, 0.4 and 0.6) series alloy electrodes are shown in Fig. 2. For all used alloy electrodes, increasing temperature promotes the increase of the discharge capacity. When the temperature is 298 K, the discharge capacity of Ti 1.0 Zr 0.2 Cr 0.4 Ni 0.8 V0.8 is 199 mAh / g and it increases from 273 mAh / g at 323 K to 335 mAh / g at 353 K. Many researches found that adding Al to the alloy is effective for improving the cycle life of Ti–Ni-based [17] and AB 5 -type [18] metal hydride electrodes because of the formation of compact oxides of aluminum, which can protect the bulk alloys from oxidization. Our experimental results show
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Fig. 1. Cycle life of the Ti 0.7 Zr 0.5 V0.2 Mn 1.82x Ni x (x50.4, 0.8, 1.1, 1.4 and 1.7) metal hydride electrodes at 283 K.
that although the discharge capacity decreases only little, except for the Ti 1.0 Zr 0.2 Cr 0.4 Ni 0.8 V0.6 Al 0.2 alloy at 298 K, the cycle stability was improved by substituting Al for part of the V element. At 298 K the capacity retention of the Ti 1.0 Zr 0.2 Cr 0.4 Ni 0.8 V0.6 Al 0.2 alloy electrode is 75% after 140 cycles; at 323 K it is 52% after 140 cycles. The increase of the temperature is bad for an improvement of the cycle stability. According to the standard electrode potential or pHelectrode potential, it is well known that Mn and V are easily corroded. Ni and Al are materials having a high resistance to corrosion. The above results show that substituting Al for V in Ti 1.0 Zr 0.2 Cr 0.4 Ni 0.8 V0.82x Al x and substituting Ni for Mn in Ti 0.7 Zr 0.5 V0.2 Mn 1.82x Ni x can lead to an improvement of the cycle stability. We considered that under some conditions, substituting elements of
Fig. 2. Cycle life of the Ti 1.0 Zr 0.2 Cr 0.4 Ni 0.8 V0.82x Al x metal hydride electrodes at 298 and 323 K.
high corrosion resistance for elements that are easily corroded is an effective method to improve cycle stability, as confirmed by above results. The investigation of the mechanism of the decay of the discharge capacity and the determination of the factors which cause this decrease of the discharge capacity, are very important for realizing commercialization of the Tibased AB 2 alloys as negative electrode for Ni–MH battery.
3.2. The X-ray analysis The geometrical size of the atoms is one of the main factors influencing the phase structures of Laves phase alloys. For AB 2 Laves phase alloy, the A atomic diameter is larger than that of B. The theoretical atomic diameter ratio rA /r B is 1.225. The phase composition of the AB 2 alloy is related to the average number of outer electrons (ANOE). For A5Ti or Zr, and B5V, Cr, Mn, Fe, Co, Cu or Zn, when ANOE is less than 5.4 no Laves phase is formed for Ti and the C15 Laves phase is formed for Zr. If ANOE is larger than 5.4 and lower than 7.0, the C14 Laves phase forms regardless of A5Ti or Zr. For ANOE .7.0 the C15 Laves phase forms. Some researchers studied the relationship between phase composition and valence electron concentration for AB 2 type alloys. Elliot and Rostoker [19] considered that, when the valence electron concentration (VEC) is 1.50–1.80, the C15 Laves phase forms. When it is 1.80–2.32, the C14 Laves phase is the main phase and when it is larger than 2.32, the C15 Laves phase still forms. All the above discussions are semi-empirical. In fact, it is very difficult to obtain single phase alloys. The results of the X-ray analysis of the Ti 0.7 Zr 0.5 V0.2 Mn 1.82x Ni x (x50.4, 0.8, 1.1, 1.4 and 1.7) and Ti 1.0 Zr 0.2 Cr 0.4 Ni 0.8 V0.82x Al x (x50.0, 0.2, 0.4 and 0.6)
Y.-h. Xu et al. / Journal of Alloys and Compounds 337 (2002) 214 – 220
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Table 1 Lattice parameters of the Ti 0.7 Zr 0.5 V0.2 Mn 1.82x Ni x (x50.4, 0.8, 1.1, 1.4 and 1.7) and Ti 1.0 Zr 0.2 Cr 0.4 Ni 0.8 V0.82x Al x (x50.0, 0.2, 0.4 and 0.6) alloys desired by XRD Alloy
˚ a (A)
˚ c (A)
˚ 3) V (A
Main phase
Second phase
ANOE
VEC
Ti 1.0 Zr 0.2 Cr 0.4 Ni 0.8 V0.8 Ti 1.0 Zr 0.2 Cr 0.4 Ni 0.8 V0.6 Al 0.2 Ti 1.0 Zr 0.2 Cr 0.4 Ni 0.8 V0.4 Al 0.4 Ti 1.0 Zr 0.2 Cr 0.4 Ni 0.8 V0.2 Al 0.6 Ti 0.7 Zr 0.5 V0.2 Mn 1.4 Ni 0.4 Ti 0.7 Zr 0.5 V0.2 Mn 1.0 Ni 0.8 Ti 0.7 Zr 0.5 V0.2 Mn 0.7 Ni 1.1 Ti 0.7 Zr 0.5 V0.2 Mn 0.4 Ni 1.4 Ti 0.7 Zr 0.5 V0.2 Mn 1.82x Ni x
4.3202 4.3157 4.1909 4.3366 4.9181 4.9322 4.9247 4.9195
8.1152 8.1107 8.0755 8.1024 8.0597 8.0348 8.0172 8.0102 No C14 Laves
151.463 151.067 141.837 152.372 168.828 169.273 168.389 167.887 phase
C14 C14 C14 C14 C14 C14 C14 C14
TiNi TiNi TiNi TiNi TiNi TiNi TiNi TiNi TiNi
6.075 5.875 5.75 5.625 6.13 6.50 6.78 7.06 7.34
1.875 1.9375 2 2.0625 2 2 2 2 2
alloys are listed in Table 1. The main phase is the C14 Laves phase for all studied alloys except for the Ti 0.7 Zr 0.5 V0.2 Mn 0.1 Ni 1.7 alloy. The second phase is the cubic TiNi phase. The ANOE of Ti 0.7 Zr 0.5 V0.2 Mn 0.1 Ni 1.7 alloy is 7.34 and its main phase is the C15 Laves phase although its VEC is 2. For the other alloys it can be seen from Table 1 that the ANOE and valence electron concentration are in the range of 5.625–7.06 and 1.875– 2.0625, respectively, and that the main phase is the C14 Laves phase. The above measurement results are in agreement with what reported in the literature [19].
3.3. Energy disperse analysis EDS is used to analyze the change in element content at the alloy surface before electrochemical cycling and after 20 cycles. The energy disperse spectra of the Ti 0.7 Zr 0.5 V0.2 Mn 1.82x Ni x (x50.4, 0.8, 1.1, 1.4) alloys are plotted in Figs. 3 and 4. The analysis results are listed in Table 2.
Fig. 3. Energy disperse spectra of the cast Ti 0.7 Zr 0.5 V0.2 Mn 1.82x Ni x (x5 0.4, 0.8, 1.1, 1.4 and 1.7) alloys.
The results show that after 20 cycles, for the x50.4 alloy, the Ti and Zr concentrations at the surface decreases. When x is larger than 0.4, the Ti and Zr concentrations increase with charging–discharging cycling. Nickel substitution for part of the Mn is effective in suppressing the dissolution of V. The V content at the alloy surface becomes lower with cycling and its decrease becomes small when the content of nickel increases, as found by EDS. The concentration of the surface elements is different from the average concentration: the contents of Mn and V are larger than the average content but the contents of Ti and Zr are lower than that. The difference between surface and bulk concentration is relatively large for Zr and Mn. This may be caused by differences in surface energy of the different elements.
3.4. The XPS analysis X-ray photoelectron spectroscopy is a useful and powerful technique to investigate the surface state of electrons. The XPS spectra of the Ti 0.7 Zr 0.5 V0.2 Mn 1.82x Ni x (x50.4, 0.8, 1.1, 1.4) alloys before electrochemical cycling and after 100 cycles are shown in Figs. 5 and 6, respectively. The standard binding energy of Ti 2p 3 / 2 , Zr 3d 5 / 2 , Cr 2p 3 / 2 , V 2p 3 / 2 , Mn 2p 3 / 2 , Ni 2p 3 / 2 and Al 2p 3 / 2 is 453.8, 178.7, 574.1, 511.95, 638.8, 852.3 and 72.65 eV, respectively. The standard binding energy of Ti 41 2p 3 / 2 , Cr 31 51 41 21 31 2p 3 / 2 , V 2p 3 / 2 , Mn 2p 3 / 2 , Ni 2p 3 / 2 and Al 2p 3 / 2 is 458.5, 576.6, 517.45, 642.2, 853.3 and 74.7 eV, respectively, for their oxides and that of K 11 2p 3 / 2 for the KBr compound is 292.9 eV. For the Ti 0.7 Zr 0.5 V0.2 Mn 1.82x Ni x (x50.4, 0.8, 1.1, 1.4) series alloys, V and Ni are not detected. Only Ti, Zr and Mn are found by XPS, which is different from the results obtained by EDS technique. The XPS analysis has shown that there is a serious segregation of Mn and its content at the alloy surface is larger than the average content. When the nickel content increases its segregation becomes more serious, i.e. the difference between surface content and average content becomes larger. The binding energy of Ti 2p 3 / 2 , Zr 3d 5 / 2 and Mn 2p 3 / 2 is 458.0, 182 and 642 eV, respectively, and it is
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Fig. 4. Energy disperse spectra for the cast Ti 0.7 Zr 0.5 V0.2 Mn 1.82x Ni x (x50.4, 0.8, 1.1, 1.4 and 1.7) alloys after 100 cycles.
larger than the standard binding energy of Ti, Zr and Mn but slightly lower than that of their oxides. Therefore we considered that the elements Ti, Zr and Mn at the alloy surface exist in the form ZrO x (0,x#2), TiO x (0,x#2) and MnO x (0,x#2) type, respectively. After 100 electrochemical cycles, V is not detected and Ni is found by XPS, except for Ti 0.7 Zr 0.5 V0.2 Mn 1.4 Ni 0.4 . When x51.4, Mn is not detected after charging–discharging cycles. Based on this phenomenon, we guessed that the segregation of Mn can be suppressed when the content of Ni increases. However, the XPS experiments show the fact that the content of some elements at the alloy surface is different Table 2 Elemental concentration (atom%) of the Ti 0.7 Zr 0.5 V0.2 Mn 1.82x Ni x (x50.4, 0.8, 1.1, 1.4 and 1.7) alloys desired by the EDS technique Element x50.4 Average x50.8 Average x51.1 Average x51.4 Average x51.7 Average
Cast * content Cast * content Cast * content Cast * content Cast * content
Ti
Zr
V
Mn
Ni
19.90 18.09 21.875 19.26 20.35 21.875 20.40 25.03 21.875 20.49 21.06 21.875 20.59 19.89 21.875
7.18 5.80 15.625 6.45 6.70 15.625 7.66 8.11 15.625 8.1 8.85 15.625 8.10 8.04 15.625
7.96 5.07 6.25 8.19 7.09 6.25 7.80 6.85 6.25 7.99 7.73 6.25 7.64 7.47 6.25
52.44 60.34 43.75 39.13 42.60 31.25 28.76 25.95 21.875 15.93 14.89 12.5 6.28 5.36 3.125
12.52 10.69 12.5 26.97 23.28 25.00 35.38 34.06 34.375 47.49 47.46 43.75 57.39 59.23 53.125
*After electrochemical cycles.
from that in the bulk and that some elements exist in their oxides. When charging–discharging, new fresh surfaces appear due to pulverization, and some elements such as V, Mn and Ti dissolve into the KOH solution. EDS analysis still confirmed this phenomenon. Continuous dissolution and segregation of some elements from the alloy surface causes a deviation of the local composition from the AB 2 composition. The AB 2 Laves phase slowly decomposes upon electrochemical cycling. As results, the relative content of the Laves phase, which plays an important part in the hydrogen storage capacity, decreases slowly with cycling. The existence of the oxides layer at the surface is bad for the conductivity and hydrogen diffusion, which causes an increase of the linear-polarization resistance. The XPS spectra of the Ti 1.0 Zr 0.2 Cr 0.4 Ni 0.8 V0.82x Al x
Fig. 5. XPS spectra of the cast Ti 0.7 Zr 0.5 V0.2 Mn 1.82x Ni x alloys.
Y.-h. Xu et al. / Journal of Alloys and Compounds 337 (2002) 214 – 220
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Fig. 6. XPS spectra for the Ti 0.7 Zr 0.5 V0.2 Mn 1.82x Ni x alloys after electrochemical cycles.
Fig. 8. XPS spectra of the Ti 0.7 Zr 0.5 Cr 0.4 Ni 0.8 V0.82x Al x (x50.0, 0.2, 0.4 and 0.6) after 50 cycles.
(x50.0, 0.2, 0.4 and 0.6) hydrogen storage alloys are given in Figs. 7 and 8. For the cast alloys, Al is detected only when x is larger than 0.4, and Ni and V are not found by XPS. For alloys cycled in KOH solution, Cr, V and Al are not detected as shown in Fig. 8. For the Ti 1.0 Zr 0.2 Cr 0.4 Ni 0.8 V0.82x Al x (x50.4 and 0.6) alloys, the Al content at the alloy surface is about 40 at. and larger than the average content. The Zr content at the surface is also larger than the average composition. Because Al is apt to dissolve in alkaline solution and exists as AlO 2 2 , Al at the alloy grain surface dissolves uninterruptedly, which is possibly the reason why the Al was not detected after electrochemical cycling. It was confirmed by the ICP [20] technique that some elements are easily dissolved in alkaline solution. In short, the results of the XPS experiments for Ti 1.0 Zr 0.2 Cr 0.4 Ni 0.8 V0.82x Al x alloys are similar to those of the Ti 0.7 Zr 0.5 V0.2 Mn 1.82x Ni x (x50.4, 0.8, 1.1,
1.4) series alloys. The XPS analysis of this series alloys also confirmed that the segregation (Ti, Zr, Mn, etc.) and dissolution (V, Al, etc.) of some elements are key factors influencing the cycle stability.
3.5. The linear-polarization resistance The formation of oxides at the alloy surface, which is confirmed by the XPS tests, could cause an increase of the linear-polarization resistance. The corresponding electrochemical tests demonstrate this phenomenon, as shown in Table 3, which represents the change of the linear-polarization resistance (R V ) with cycling for all studied alloys. For example, the R V increases from 1.012 V?g at the first cycle to 24.8 V?g at the 15th cycle at 323 K, while its discharge capacity decreases from 399.1 to about 15 mAh / g [11]. Similar results are found for the Ti 1.0 Zr 0.2 Cr 0.4 Ni 0.8 V0.82x Al x series alloys: the R V inTable 3 Change of the linear-polarization resistance (R V ) with cycling for the Ti 0.7 Zr 0.5 V0.2 Mn 1.82x Ni x (x50.4, 0.8, 1.1 and 1.4) and Ti 1.0 Zr 0.2 Cr 0.4 Ni 0.8 V0.82x Al x (x50.0, 0.2, 0.4 and 0.6) alloys
Fig. 7. XPS spectra for the cast Ti 0.7 Zr 0.5 Cr 0.4 Ni 0.8 V0.82x Al x (x50.0, 0.2, 0.4 and 0.6).
Alloy
Linear-polarization resistance (V?g, 323 K)
Cycle Ti 1.0 Zr 0.2 Cr 0.4 Ni 0.8 V0.8 Ti 1.0 Zr 0.2 Cr 0.4 Ni 0.8 V0.6 Al 0.2 Ti 1.0 Zr 0.2 Cr 0.4 Ni 0.8 V0.4 Al 0.4 Ti 1.0 Zr 0.2 Cr 0.4 Ni 0.8 V0.2 Al 0.6
* 0.303 0.286 0.254 0.233
60th 2.04 1.92 1.87 1.89
Cycle Ti 0.7 Zr 0.5 V0.2 Mn 1.4 Ni 0.4 Ti 0.7 Zr 0.5 V0.2 Mn 1.0 Ni 0.8 Ti 0.7 Zr 0.5 V0.2 Mn 0.7 Ni 1.1 Ti 0.7 Zr 0.5 V0.2 Mn 0.4 Ni 1.4
1st cycle 1.012 0.470 0.181 0.042
15th 24.8 14.0 9.02 3.71
120th 2.54 2.38 2.38 2.26
* For the cycle at which the alloy electrode reached the max discharge capacity.
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creases with cycling while the discharge capacity decreases. The change tendency of the linear-polarization change with cycling is in agreement with the XPS analysis.
4. Conclusions For Ti-based AB 2 hydrogen storage electrode alloys, the segregation, dissolution and oxidization of some elements are key factors causing the decrease of the discharge capacity with cycling. This is especially the case for V, Mn. We believe that the nature of the decrease of the discharge capacity is the decrease of the relative content of the Laves phase, which is effective for the storage of hydrogen. Some problems need to be overcome in order to use the Ti-based Laves phase alloys as the negative electrode materials for commercial Ni–MH batteries. Firstly, it is necessary to increase the resistance to corrosion. Secondly, it is necessary to suppress the segregation or dissolution of some elements, especially V, Mn.
Acknowledgements The financial support of the National Advanced Materials Committee of China (Number 863-715-004-0060) is acknowledged.
References [1] W.K. Hu, H. Lee, D.M. Kim, J.Y. Lee, S.W. Jeon, J. Alloys Comp. 268 (1998) 261–265.
[2] W.K. Chio, T. Tanaka, R. Miyauchi, T. Morikawa, H. Inoue, C. Iwakura, J. Alloys Comp. 299 (2000) 141. [3] J. Chen, D.H. Bradhurst, S.X. Dou, H.K. Liu, J. Alloys Comp. 280 (1998) 290. [4] N. Cui, P. He, J.L. Luo, Acta Mater. 47 (14) (1999) 3737. [5] H. Ye, Y.Q. Lei, L.X. Chen, H. Zhang, J. Alloys Comp. 311 (2000) 194–199. [6] M.A. Gutjahr, H. Buchner, K.D. Beccu, Power Sources Symp., Brighton, 1972, Power Sources 4 (1973) 79. [7] S. Wakao, H. Sawa, H. Nakano, S. Chubachi, M. Abe, J. LessComm. Met. 131 (1987) 311–319. [8] B. Luan, N. Cui, H.J. Zhao, H.K. Liu, S.X. Dou, Int. J. Hydrogen Energy 21 (1996) 373. [9] Y.H. Xu, C.P. Chen, Q.D. Wang, Mater. Chem. Phys. 71 (2001) 190–194. [10] Y.-H. Xu, C.-P. Chen, Q.-D. Wang, L.X. Chen, Int. J. Hydrogen Energy 26 (2001) 1177–1181. [11] Y.-H. Xu, C.-P. Chen, S.-Q. Li, T. Ying, Q.D. Wang, Trans. Nonferrous Met. Soc. China 3 (2001) 350. [12] J.H. Jung, H.H. Lee, D.M. Kim, K.J. Jang, J.Y. Lee, J. Alloys Comp. 266 (1998) 266–270. [13] H.Y. Park, W.I. Cho, B.W. Cho, S.R. Lee, K.S. Yun, J. Power Sources 92 (2001) 149. [14] S.N. Jeng, H.W. Yang, C.C. Wan, Y.Y. Wang, Mater. Chem. Phys. 48 (1997) 10. [15] G.D. Adzic, J.R. Johnson, J.J. Reilly, J. McBreen, S. Mukerjee, M.P.S. Kumar, W. Zhang, S. Srinivasan, J. Electrochem. Soc. 10 (1995) 3424. [16] W.K. Chio, T. Tanaka, R. Miyauchi, T. Morikawa, H. Inoue, C. Iwahura, J. Alloys Comp. 299 (2000) 141. [17] B. Luan, N. Cui, H. Zhao, S. Zhong, H.K. Liu, S.X. Dou, J. Alloys Comp. 233 (1996) 225. [18] T. Sakai, H. Miyamura, H. Kuriyama, A. Kato, K. Oguro, H. Ishikawa, J. Less-Comm. Met. 159 (1990) 127. [19] R.P. Elliot, W. Rostoker, Trans. Am. Soc. Met. 50 (1958) 617. [20] M.S. Wu, H.R. Wu, Y.Y. Wang, C.C. Wan, J. Alloys Comp. 302 (2000) 248.
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The cycle life and surface properties of Ti-based AB 2 metal hydride electrodes a, b a b b Yan-hui Xu *, Chang-pin Chen , Xiao-lin Wang , Yong-quan Lei , Qi-dong Wang b
a Department of Chemical Engineering, Tsing Hua University, Beijing 100084, PR China Department of Materials Science and Engineering, Zhejiang University, Hangzhou 310027, PR China
Received 21 August 2001; accepted 10 October 2001
Abstract In this paper, Ti-based AB 2 Laves phase hydrogen storage electrode alloys are prepared and their electrochemical cycle stability is studied, including Ti 0.7 Zr 0.5 V0.2 Mn 1.82x Ni x (x50.4, 0.8, 1.1, 1.4 and 1.7) and Ti 1.0 Zr 0.2 Cr 0.4 Ni 0.8 V0.82x Al x (x50.0, 0.2, 0.4 and 0.6) alloys. Energy disperse spectroscopy (EDS) and X-ray photoelectron energy spectroscopy (XPS) were used to analyze the changes of the elements on the alloy grain surface before and after charging–discharging cycles in a KOH solution in order to understand the nature of the decrease of discharge capacity with cycling. The experimental results show that the linear-polarization resistance increases when cycling. The segregation, oxidization and dissolution of some elements in the alloy are key factors influencing the cycle stability and are also a main problem that needs to be overcome in order to use these materials as negative electrode in a commercial nickel–metal hydride rechargeable battery. 2002 Elsevier Science B.V. All rights reserved. Keywords: Transition metal alloys; Hydrogen-absorbing materials; Electrode materials; Gas–solid reaction; X-Ray diffraction
1. Introduction In recent years, extensive investigations on metal hydride alloys have been carried out. The latter are used as negative electrodes of nickel–metal hydride batteries because of their low environmental impact, compared with Cd negative electrode [1–5]. AB 5 type rare-earth-based hydrogen storage alloys are the main active materials in commercial Ni–MH batteries and their discharge capacity can reach about 330 mAh / g. It is difficulty to further increase their discharge capacity, which is restricted by their theoretical hydrogen storage capacity. Cobalt is the most expensive element and takes up about 45% of the total cost of the Mm(NiCoMnAl) 5 alloy raw materials. Therefore, research and development of new Mm-based alloys with low cobalt or without cobalt is a new direction for AB 5 -type metal hydride alloys in order to lower the Ni–MH battery costs [1]. Zr-based AB 2 Laves phase alloys have high discharge capacity, long cycle life but are difficult to activate. For the near future it will be difficult to use Zr-based AB 2 alloys because of their high cost and low high-current discharge ability. In the past 10 years, *Corresponding author. E-mail address: [email protected] (Y.-h. Xu).
Mg–Ni-based and V–Ti-based metal hydride alloys have been given much attention because of their very high discharge capacity [2–5]. Now the discharge capacity of Mg-based alloy electrode increased from about 170 mAh / g to about 1000 mAh / g (Mg 2 Ni alloy) for Mg 2 Ni1 Ni(1:1.28). But the cycle life of Mg-based alloy electrode is very poor, as reported by Ye et al. [5]. The electrochemical behavior of V–Ti-based solid solution alloy electrode is similar to that of the Mg-based alloy: easy activation, high capacity and poor cycle life [2]. Although adding Cr, Co, Nb, Ta, C, Zr and Hf could improve the electrode properties to some degree, it still cannot satisfy commercial requirements. Vanadium is also an expensive material. According to the corresponding phase diagram or standard electrode potential, the most stable form of V and Mg in alkaline solution is their oxides. Upon cycling, oxidation or dissolution of some elements, especially V and Mg, caused a decrease in discharge capacity. We consider that Mg-based and V-based alloys cannot be used as negative electrode active materials for commercial Ni–MH battery. Several authors [6–10] reported on the electrochemical behavior of Ti–Ni-based alloys. The discharge capacity of TiNi alloy electrodes reached 210–250 mAh / g at 5 mA / g current density and reached 180 mAh / g at 50 mA / g. Gutjahr et al. [6] found that a sintered mixture of
0925-8388 / 02 / $ – see front matter 2002 Elsevier Science B.V. All rights reserved. PII: S0925-8388( 01 )01917-X
Y.-h. Xu et al. / Journal of Alloys and Compounds 337 (2002) 214 – 220
Ti 2 Ni and TiNi phases possessed reversible capacities above 300 mAh / g. The discharge capacity of Ti 3 Ni 2 alloy, which is a composite of the TiNi phase and the Ti 2 Ni phase, increases with increasing temperature [10]. The cycle life of this type of alloy needs further improvement when it is to be used as negative electrode material. Ti-based AB 2 metal hydride alloys are another type of potential candidate because of their easy activation, large hydrogen storage capacity, and very low cost. Many researches [11,12] investigated their electrochemical behavior and found that the key problem which needs to be overcome is the increase of the cycle life when used as negative electrode materials. Many researches investigated the reason for the poor cycle life of some metal hydride alloys using electrochemical techniques or physical analysis methods. According to several reports [13–16], the pulverization, the corrosion of the alloys and the segregation and dissolution of some of the elements (especially V, Mn, Al and rare earth elements) are the main factors influencing the cycle life of metal hydride electrodes. Previous investigations [11] have shown that the discharge capacity of the Ti 0.7 Zr 0.5 V0.2 Mn 1.82x Ni x is 399.1 mAh / g at 323 K at first cycle and decreases fast with cycling. For example, when x is 0.4 the discharge capacity decreases to about 20 mAh / g after three cycles. In this paper, the cycle life of the Ti 0.7 Zr 0.5 V0.2 Mn 1.82x Ni x (x50.4, 0.8, 1.1, 1.4 and 1.7) and Ti 1.0 Zr 0.2 Cr 0.4 Ni 1.8 V0.82x Al x (x50.0, 0.2, 0.4 and 0.6) metal hydride electrodes was investigated. Physical analysis techniques were used in order to find the reason why the discharge capacity of Ti-based hydrogen storage alloy electrodes decreases fast with cycling. We used X-ray diffraction (XRD), energy disperse spectroscopy (EDS) and X-ray photoelectron energy spectroscopy (XPS).
2. Experiment The purity of all used original materials is higher than 99.5%. The ingots were prepared by high-frequency induction melting in a water-cooled boat under argon atmosphere and repeated four times in order to ensure homogeneity. Then the ingots were decrepitated into powders for electrochemical measurement and physical analysis. The electrode preparation method is the following: a mixture of the alloy powders and nickel powder (mass ratio is 1:2) was pressed at 1.5 M Pa into a round disk electrode with a diameter of 10 mm and thickness of about 1 mm. The content of the alloy powder is about 200 mg. The reference electrode is Hg / HgO / 6 N KOH and the counter electrode is Ni(OH) 2 / NiOOH. The discharge capacity of which is more than that of the positive electrode. The electrolyte is a 6 N KOH water solution. Electrochemical measurement was carried out in a V-type open glass three-electrode system. The end point of the
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discharge was set at 20.6 V (versus Hg / HgO). The charging–discharging cycle was measured using a DC-5 Battery Test Instrument. A Philips X’pert diffractometer was used to analyze the phase structure of all used alloys. The power is 40 kV320 mA and the scanning rate is 0.048 s 21 . The energy dispersive analysis was performed with a PV9900-Philips Energy Analyzer in order to determine the change in concentration of the various elements before and after electrochemical cycling. The duration time was set at 100 s. The XPS analysis technique, which was conducted using Al Ka radiation and an Escalab MKII photoelectron energy spectrum instrument, was used to determine the chemical state and content of the constituent elements of the alloys. The vacuum in the analysis room was 10 29 Torr. The ray energy was 50 eV.
3. Results
3.1. The cycle life of T-based AB2 metal hydride electrodes The cycle life of Ti 0.7 Zr 0.5 V0.2 Mn 1.82x Ni x (x50.4, 0.8, 1.1, 1.4 and 1.7) hydrogen storage alloys at 283 K is shown in Fig. 1, from which it can be seen that the cycle stability of this alloy electrode increases as the content of nickel increases. The maximum discharge capacity increases from about 110 mAh / g for x50.4 to about 286 mAh / g for x51.4 except for x51.7. The phase composition for the x51.7 alloy is different from the other alloys. Nickel partly substituting for manganese is very effective for improving the discharge capacity as shown in Fig. 1. It was found from our previous investigations that the environmental temperature influences strongly the discharge capacity and cycle life. For example, the discharge capacity of the Ti 0.7 Zr 0.5 V0.2 Mn 1.4 Ni 0.4 alloy increases from 111.7 mAh / g at 283 K to 399.1 mAh / g at 323 K and then decreases to about 260 mAh / g at 353 K. After 10 charging–discharging cycles the discharge capacity for the x50.4 alloy decreases to about 50 mAh / g, but at 323 and 353 K after only three cycles it decreases to about 15 mAh / g [11]. This confirmed that an increasing environment temperature is not beneficial to cycle life. The cycle life curves of Ti 1.0 Zr 0.2 Cr 0.4 Ni 0.8 V0.82x Al x (x50.0, 0.2, 0.4 and 0.6) series alloy electrodes are shown in Fig. 2. For all used alloy electrodes, increasing temperature promotes the increase of the discharge capacity. When the temperature is 298 K, the discharge capacity of Ti 1.0 Zr 0.2 Cr 0.4 Ni 0.8 V0.8 is 199 mAh / g and it increases from 273 mAh / g at 323 K to 335 mAh / g at 353 K. Many researches found that adding Al to the alloy is effective for improving the cycle life of Ti–Ni-based [17] and AB 5 -type [18] metal hydride electrodes because of the formation of compact oxides of aluminum, which can protect the bulk alloys from oxidization. Our experimental results show
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Fig. 1. Cycle life of the Ti 0.7 Zr 0.5 V0.2 Mn 1.82x Ni x (x50.4, 0.8, 1.1, 1.4 and 1.7) metal hydride electrodes at 283 K.
that although the discharge capacity decreases only little, except for the Ti 1.0 Zr 0.2 Cr 0.4 Ni 0.8 V0.6 Al 0.2 alloy at 298 K, the cycle stability was improved by substituting Al for part of the V element. At 298 K the capacity retention of the Ti 1.0 Zr 0.2 Cr 0.4 Ni 0.8 V0.6 Al 0.2 alloy electrode is 75% after 140 cycles; at 323 K it is 52% after 140 cycles. The increase of the temperature is bad for an improvement of the cycle stability. According to the standard electrode potential or pHelectrode potential, it is well known that Mn and V are easily corroded. Ni and Al are materials having a high resistance to corrosion. The above results show that substituting Al for V in Ti 1.0 Zr 0.2 Cr 0.4 Ni 0.8 V0.82x Al x and substituting Ni for Mn in Ti 0.7 Zr 0.5 V0.2 Mn 1.82x Ni x can lead to an improvement of the cycle stability. We considered that under some conditions, substituting elements of
Fig. 2. Cycle life of the Ti 1.0 Zr 0.2 Cr 0.4 Ni 0.8 V0.82x Al x metal hydride electrodes at 298 and 323 K.
high corrosion resistance for elements that are easily corroded is an effective method to improve cycle stability, as confirmed by above results. The investigation of the mechanism of the decay of the discharge capacity and the determination of the factors which cause this decrease of the discharge capacity, are very important for realizing commercialization of the Tibased AB 2 alloys as negative electrode for Ni–MH battery.
3.2. The X-ray analysis The geometrical size of the atoms is one of the main factors influencing the phase structures of Laves phase alloys. For AB 2 Laves phase alloy, the A atomic diameter is larger than that of B. The theoretical atomic diameter ratio rA /r B is 1.225. The phase composition of the AB 2 alloy is related to the average number of outer electrons (ANOE). For A5Ti or Zr, and B5V, Cr, Mn, Fe, Co, Cu or Zn, when ANOE is less than 5.4 no Laves phase is formed for Ti and the C15 Laves phase is formed for Zr. If ANOE is larger than 5.4 and lower than 7.0, the C14 Laves phase forms regardless of A5Ti or Zr. For ANOE .7.0 the C15 Laves phase forms. Some researchers studied the relationship between phase composition and valence electron concentration for AB 2 type alloys. Elliot and Rostoker [19] considered that, when the valence electron concentration (VEC) is 1.50–1.80, the C15 Laves phase forms. When it is 1.80–2.32, the C14 Laves phase is the main phase and when it is larger than 2.32, the C15 Laves phase still forms. All the above discussions are semi-empirical. In fact, it is very difficult to obtain single phase alloys. The results of the X-ray analysis of the Ti 0.7 Zr 0.5 V0.2 Mn 1.82x Ni x (x50.4, 0.8, 1.1, 1.4 and 1.7) and Ti 1.0 Zr 0.2 Cr 0.4 Ni 0.8 V0.82x Al x (x50.0, 0.2, 0.4 and 0.6)
Y.-h. Xu et al. / Journal of Alloys and Compounds 337 (2002) 214 – 220
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Table 1 Lattice parameters of the Ti 0.7 Zr 0.5 V0.2 Mn 1.82x Ni x (x50.4, 0.8, 1.1, 1.4 and 1.7) and Ti 1.0 Zr 0.2 Cr 0.4 Ni 0.8 V0.82x Al x (x50.0, 0.2, 0.4 and 0.6) alloys desired by XRD Alloy
˚ a (A)
˚ c (A)
˚ 3) V (A
Main phase
Second phase
ANOE
VEC
Ti 1.0 Zr 0.2 Cr 0.4 Ni 0.8 V0.8 Ti 1.0 Zr 0.2 Cr 0.4 Ni 0.8 V0.6 Al 0.2 Ti 1.0 Zr 0.2 Cr 0.4 Ni 0.8 V0.4 Al 0.4 Ti 1.0 Zr 0.2 Cr 0.4 Ni 0.8 V0.2 Al 0.6 Ti 0.7 Zr 0.5 V0.2 Mn 1.4 Ni 0.4 Ti 0.7 Zr 0.5 V0.2 Mn 1.0 Ni 0.8 Ti 0.7 Zr 0.5 V0.2 Mn 0.7 Ni 1.1 Ti 0.7 Zr 0.5 V0.2 Mn 0.4 Ni 1.4 Ti 0.7 Zr 0.5 V0.2 Mn 1.82x Ni x
4.3202 4.3157 4.1909 4.3366 4.9181 4.9322 4.9247 4.9195
8.1152 8.1107 8.0755 8.1024 8.0597 8.0348 8.0172 8.0102 No C14 Laves
151.463 151.067 141.837 152.372 168.828 169.273 168.389 167.887 phase
C14 C14 C14 C14 C14 C14 C14 C14
TiNi TiNi TiNi TiNi TiNi TiNi TiNi TiNi TiNi
6.075 5.875 5.75 5.625 6.13 6.50 6.78 7.06 7.34
1.875 1.9375 2 2.0625 2 2 2 2 2
alloys are listed in Table 1. The main phase is the C14 Laves phase for all studied alloys except for the Ti 0.7 Zr 0.5 V0.2 Mn 0.1 Ni 1.7 alloy. The second phase is the cubic TiNi phase. The ANOE of Ti 0.7 Zr 0.5 V0.2 Mn 0.1 Ni 1.7 alloy is 7.34 and its main phase is the C15 Laves phase although its VEC is 2. For the other alloys it can be seen from Table 1 that the ANOE and valence electron concentration are in the range of 5.625–7.06 and 1.875– 2.0625, respectively, and that the main phase is the C14 Laves phase. The above measurement results are in agreement with what reported in the literature [19].
3.3. Energy disperse analysis EDS is used to analyze the change in element content at the alloy surface before electrochemical cycling and after 20 cycles. The energy disperse spectra of the Ti 0.7 Zr 0.5 V0.2 Mn 1.82x Ni x (x50.4, 0.8, 1.1, 1.4) alloys are plotted in Figs. 3 and 4. The analysis results are listed in Table 2.
Fig. 3. Energy disperse spectra of the cast Ti 0.7 Zr 0.5 V0.2 Mn 1.82x Ni x (x5 0.4, 0.8, 1.1, 1.4 and 1.7) alloys.
The results show that after 20 cycles, for the x50.4 alloy, the Ti and Zr concentrations at the surface decreases. When x is larger than 0.4, the Ti and Zr concentrations increase with charging–discharging cycling. Nickel substitution for part of the Mn is effective in suppressing the dissolution of V. The V content at the alloy surface becomes lower with cycling and its decrease becomes small when the content of nickel increases, as found by EDS. The concentration of the surface elements is different from the average concentration: the contents of Mn and V are larger than the average content but the contents of Ti and Zr are lower than that. The difference between surface and bulk concentration is relatively large for Zr and Mn. This may be caused by differences in surface energy of the different elements.
3.4. The XPS analysis X-ray photoelectron spectroscopy is a useful and powerful technique to investigate the surface state of electrons. The XPS spectra of the Ti 0.7 Zr 0.5 V0.2 Mn 1.82x Ni x (x50.4, 0.8, 1.1, 1.4) alloys before electrochemical cycling and after 100 cycles are shown in Figs. 5 and 6, respectively. The standard binding energy of Ti 2p 3 / 2 , Zr 3d 5 / 2 , Cr 2p 3 / 2 , V 2p 3 / 2 , Mn 2p 3 / 2 , Ni 2p 3 / 2 and Al 2p 3 / 2 is 453.8, 178.7, 574.1, 511.95, 638.8, 852.3 and 72.65 eV, respectively. The standard binding energy of Ti 41 2p 3 / 2 , Cr 31 51 41 21 31 2p 3 / 2 , V 2p 3 / 2 , Mn 2p 3 / 2 , Ni 2p 3 / 2 and Al 2p 3 / 2 is 458.5, 576.6, 517.45, 642.2, 853.3 and 74.7 eV, respectively, for their oxides and that of K 11 2p 3 / 2 for the KBr compound is 292.9 eV. For the Ti 0.7 Zr 0.5 V0.2 Mn 1.82x Ni x (x50.4, 0.8, 1.1, 1.4) series alloys, V and Ni are not detected. Only Ti, Zr and Mn are found by XPS, which is different from the results obtained by EDS technique. The XPS analysis has shown that there is a serious segregation of Mn and its content at the alloy surface is larger than the average content. When the nickel content increases its segregation becomes more serious, i.e. the difference between surface content and average content becomes larger. The binding energy of Ti 2p 3 / 2 , Zr 3d 5 / 2 and Mn 2p 3 / 2 is 458.0, 182 and 642 eV, respectively, and it is
Y.-h. Xu et al. / Journal of Alloys and Compounds 337 (2002) 214 – 220
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Fig. 4. Energy disperse spectra for the cast Ti 0.7 Zr 0.5 V0.2 Mn 1.82x Ni x (x50.4, 0.8, 1.1, 1.4 and 1.7) alloys after 100 cycles.
larger than the standard binding energy of Ti, Zr and Mn but slightly lower than that of their oxides. Therefore we considered that the elements Ti, Zr and Mn at the alloy surface exist in the form ZrO x (0,x#2), TiO x (0,x#2) and MnO x (0,x#2) type, respectively. After 100 electrochemical cycles, V is not detected and Ni is found by XPS, except for Ti 0.7 Zr 0.5 V0.2 Mn 1.4 Ni 0.4 . When x51.4, Mn is not detected after charging–discharging cycles. Based on this phenomenon, we guessed that the segregation of Mn can be suppressed when the content of Ni increases. However, the XPS experiments show the fact that the content of some elements at the alloy surface is different Table 2 Elemental concentration (atom%) of the Ti 0.7 Zr 0.5 V0.2 Mn 1.82x Ni x (x50.4, 0.8, 1.1, 1.4 and 1.7) alloys desired by the EDS technique Element x50.4 Average x50.8 Average x51.1 Average x51.4 Average x51.7 Average
Cast * content Cast * content Cast * content Cast * content Cast * content
Ti
Zr
V
Mn
Ni
19.90 18.09 21.875 19.26 20.35 21.875 20.40 25.03 21.875 20.49 21.06 21.875 20.59 19.89 21.875
7.18 5.80 15.625 6.45 6.70 15.625 7.66 8.11 15.625 8.1 8.85 15.625 8.10 8.04 15.625
7.96 5.07 6.25 8.19 7.09 6.25 7.80 6.85 6.25 7.99 7.73 6.25 7.64 7.47 6.25
52.44 60.34 43.75 39.13 42.60 31.25 28.76 25.95 21.875 15.93 14.89 12.5 6.28 5.36 3.125
12.52 10.69 12.5 26.97 23.28 25.00 35.38 34.06 34.375 47.49 47.46 43.75 57.39 59.23 53.125
*After electrochemical cycles.
from that in the bulk and that some elements exist in their oxides. When charging–discharging, new fresh surfaces appear due to pulverization, and some elements such as V, Mn and Ti dissolve into the KOH solution. EDS analysis still confirmed this phenomenon. Continuous dissolution and segregation of some elements from the alloy surface causes a deviation of the local composition from the AB 2 composition. The AB 2 Laves phase slowly decomposes upon electrochemical cycling. As results, the relative content of the Laves phase, which plays an important part in the hydrogen storage capacity, decreases slowly with cycling. The existence of the oxides layer at the surface is bad for the conductivity and hydrogen diffusion, which causes an increase of the linear-polarization resistance. The XPS spectra of the Ti 1.0 Zr 0.2 Cr 0.4 Ni 0.8 V0.82x Al x
Fig. 5. XPS spectra of the cast Ti 0.7 Zr 0.5 V0.2 Mn 1.82x Ni x alloys.
Y.-h. Xu et al. / Journal of Alloys and Compounds 337 (2002) 214 – 220
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Fig. 6. XPS spectra for the Ti 0.7 Zr 0.5 V0.2 Mn 1.82x Ni x alloys after electrochemical cycles.
Fig. 8. XPS spectra of the Ti 0.7 Zr 0.5 Cr 0.4 Ni 0.8 V0.82x Al x (x50.0, 0.2, 0.4 and 0.6) after 50 cycles.
(x50.0, 0.2, 0.4 and 0.6) hydrogen storage alloys are given in Figs. 7 and 8. For the cast alloys, Al is detected only when x is larger than 0.4, and Ni and V are not found by XPS. For alloys cycled in KOH solution, Cr, V and Al are not detected as shown in Fig. 8. For the Ti 1.0 Zr 0.2 Cr 0.4 Ni 0.8 V0.82x Al x (x50.4 and 0.6) alloys, the Al content at the alloy surface is about 40 at. and larger than the average content. The Zr content at the surface is also larger than the average composition. Because Al is apt to dissolve in alkaline solution and exists as AlO 2 2 , Al at the alloy grain surface dissolves uninterruptedly, which is possibly the reason why the Al was not detected after electrochemical cycling. It was confirmed by the ICP [20] technique that some elements are easily dissolved in alkaline solution. In short, the results of the XPS experiments for Ti 1.0 Zr 0.2 Cr 0.4 Ni 0.8 V0.82x Al x alloys are similar to those of the Ti 0.7 Zr 0.5 V0.2 Mn 1.82x Ni x (x50.4, 0.8, 1.1,
1.4) series alloys. The XPS analysis of this series alloys also confirmed that the segregation (Ti, Zr, Mn, etc.) and dissolution (V, Al, etc.) of some elements are key factors influencing the cycle stability.
3.5. The linear-polarization resistance The formation of oxides at the alloy surface, which is confirmed by the XPS tests, could cause an increase of the linear-polarization resistance. The corresponding electrochemical tests demonstrate this phenomenon, as shown in Table 3, which represents the change of the linear-polarization resistance (R V ) with cycling for all studied alloys. For example, the R V increases from 1.012 V?g at the first cycle to 24.8 V?g at the 15th cycle at 323 K, while its discharge capacity decreases from 399.1 to about 15 mAh / g [11]. Similar results are found for the Ti 1.0 Zr 0.2 Cr 0.4 Ni 0.8 V0.82x Al x series alloys: the R V inTable 3 Change of the linear-polarization resistance (R V ) with cycling for the Ti 0.7 Zr 0.5 V0.2 Mn 1.82x Ni x (x50.4, 0.8, 1.1 and 1.4) and Ti 1.0 Zr 0.2 Cr 0.4 Ni 0.8 V0.82x Al x (x50.0, 0.2, 0.4 and 0.6) alloys
Fig. 7. XPS spectra for the cast Ti 0.7 Zr 0.5 Cr 0.4 Ni 0.8 V0.82x Al x (x50.0, 0.2, 0.4 and 0.6).
Alloy
Linear-polarization resistance (V?g, 323 K)
Cycle Ti 1.0 Zr 0.2 Cr 0.4 Ni 0.8 V0.8 Ti 1.0 Zr 0.2 Cr 0.4 Ni 0.8 V0.6 Al 0.2 Ti 1.0 Zr 0.2 Cr 0.4 Ni 0.8 V0.4 Al 0.4 Ti 1.0 Zr 0.2 Cr 0.4 Ni 0.8 V0.2 Al 0.6
* 0.303 0.286 0.254 0.233
60th 2.04 1.92 1.87 1.89
Cycle Ti 0.7 Zr 0.5 V0.2 Mn 1.4 Ni 0.4 Ti 0.7 Zr 0.5 V0.2 Mn 1.0 Ni 0.8 Ti 0.7 Zr 0.5 V0.2 Mn 0.7 Ni 1.1 Ti 0.7 Zr 0.5 V0.2 Mn 0.4 Ni 1.4
1st cycle 1.012 0.470 0.181 0.042
15th 24.8 14.0 9.02 3.71
120th 2.54 2.38 2.38 2.26
* For the cycle at which the alloy electrode reached the max discharge capacity.
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creases with cycling while the discharge capacity decreases. The change tendency of the linear-polarization change with cycling is in agreement with the XPS analysis.
4. Conclusions For Ti-based AB 2 hydrogen storage electrode alloys, the segregation, dissolution and oxidization of some elements are key factors causing the decrease of the discharge capacity with cycling. This is especially the case for V, Mn. We believe that the nature of the decrease of the discharge capacity is the decrease of the relative content of the Laves phase, which is effective for the storage of hydrogen. Some problems need to be overcome in order to use the Ti-based Laves phase alloys as the negative electrode materials for commercial Ni–MH batteries. Firstly, it is necessary to increase the resistance to corrosion. Secondly, it is necessary to suppress the segregation or dissolution of some elements, especially V, Mn.
Acknowledgements The financial support of the National Advanced Materials Committee of China (Number 863-715-004-0060) is acknowledged.
References [1] W.K. Hu, H. Lee, D.M. Kim, J.Y. Lee, S.W. Jeon, J. Alloys Comp. 268 (1998) 261–265.
[2] W.K. Chio, T. Tanaka, R. Miyauchi, T. Morikawa, H. Inoue, C. Iwakura, J. Alloys Comp. 299 (2000) 141. [3] J. Chen, D.H. Bradhurst, S.X. Dou, H.K. Liu, J. Alloys Comp. 280 (1998) 290. [4] N. Cui, P. He, J.L. Luo, Acta Mater. 47 (14) (1999) 3737. [5] H. Ye, Y.Q. Lei, L.X. Chen, H. Zhang, J. Alloys Comp. 311 (2000) 194–199. [6] M.A. Gutjahr, H. Buchner, K.D. Beccu, Power Sources Symp., Brighton, 1972, Power Sources 4 (1973) 79. [7] S. Wakao, H. Sawa, H. Nakano, S. Chubachi, M. Abe, J. LessComm. Met. 131 (1987) 311–319. [8] B. Luan, N. Cui, H.J. Zhao, H.K. Liu, S.X. Dou, Int. J. Hydrogen Energy 21 (1996) 373. [9] Y.H. Xu, C.P. Chen, Q.D. Wang, Mater. Chem. Phys. 71 (2001) 190–194. [10] Y.-H. Xu, C.-P. Chen, Q.-D. Wang, L.X. Chen, Int. J. Hydrogen Energy 26 (2001) 1177–1181. [11] Y.-H. Xu, C.-P. Chen, S.-Q. Li, T. Ying, Q.D. Wang, Trans. Nonferrous Met. Soc. China 3 (2001) 350. [12] J.H. Jung, H.H. Lee, D.M. Kim, K.J. Jang, J.Y. Lee, J. Alloys Comp. 266 (1998) 266–270. [13] H.Y. Park, W.I. Cho, B.W. Cho, S.R. Lee, K.S. Yun, J. Power Sources 92 (2001) 149. [14] S.N. Jeng, H.W. Yang, C.C. Wan, Y.Y. Wang, Mater. Chem. Phys. 48 (1997) 10. [15] G.D. Adzic, J.R. Johnson, J.J. Reilly, J. McBreen, S. Mukerjee, M.P.S. Kumar, W. Zhang, S. Srinivasan, J. Electrochem. Soc. 10 (1995) 3424. [16] W.K. Chio, T. Tanaka, R. Miyauchi, T. Morikawa, H. Inoue, C. Iwahura, J. Alloys Comp. 299 (2000) 141. [17] B. Luan, N. Cui, H. Zhao, S. Zhong, H.K. Liu, S.X. Dou, J. Alloys Comp. 233 (1996) 225. [18] T. Sakai, H. Miyamura, H. Kuriyama, A. Kato, K. Oguro, H. Ishikawa, J. Less-Comm. Met. 159 (1990) 127. [19] R.P. Elliot, W. Rostoker, Trans. Am. Soc. Met. 50 (1958) 617. [20] M.S. Wu, H.R. Wu, Y.Y. Wang, C.C. Wan, J. Alloys Comp. 302 (2000) 248.