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Effect of Ni on electrochemical and hydrogen storage properties of V-rich body-centered-cubic solid solution alloys
i n t e r n a t i o n a l j o u r n a l o f h y d r o g e n e n e r g y 4 3 ( 2 0 1 8 ) 8 3 9 5 e8 4 0 3
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Effect of Ni on electrochemical and hydrogen storage properties of V-rich body-centered-cubic solid solution alloys M. Balcerzak University of Technology, Institute of Materials Science and Engineering, Jana Pawła II No 24, 61-138 Poznan , Poland Poznan
article info
abstract
Article history:
V-rich solid solution alloys are potential candidates for Ni-MHx negative electrodes and
Received 8 December 2017
hydrogen sorbing materials. Mechanical alloying (MA) is used in this paper to produce
Received in revised form
Ti0.5V1.5xNix nanocrystalline alloys (x ¼ 0, 0.1, 0.2, 0.3). A SPEX 8000 M mill is used. The aim
12 March 2018
of this work is to study the effect of chemical modification by Ni on hydrogen sorption/
Accepted 16 March 2018
desorption and electrochemical properties of V-rich body-centered-cubic (BCC) alloys.
Available online 5 April 2018
Presented measurements results show formation of BCC phase after 14 h of MA. The nanocrystallinety of obtained materials is confirmed by high resolution transmission
Keywords:
electron microscopy images. MA alloys are tested by a Sievert's device at near room tem-
Energy storage
perature. Partial substitution of V by Ni causes improved hydrogenation kinetics, reduced
Mechanical alloying
hysteresis and increased hydrogenation/dehydrogenation reversibility. Observed proper-
Hydrogen sorption
ties are mainly due to differences in structures of studied materials. Electrochemical
X-ray diffraction
studies on chemically modified V-rich alloys show that capacity retaining rate and
Differential scanning calorimetry
discharge capacity increase with higher Ni content in the material.
Nanomaterials
© 2018 Hydrogen Energy Publications LLC. Published by Elsevier Ltd. All rights reserved.
Introduction V-rich body-centered-cubic (BCC) solid solution alloys are considered for Ni-MHx negative electrode and hydrogen storage materials. BCC-based alloys are mostly synthesized by conventional methods e arc or induction melting. These preparation methods provide production of microcrystalline materials [1e3]. However, due to the fact that elements are characterized by different melting temperatures, some desired chemical compositions are impossible to obtain using conventional methods. Moreover, microcrystalline alloys suffer from slow hydrogen sorption/desorption kinetics and difficult activation process [1]. The kinetics is closely related to
microstructure of material. Reduction of crystallites size and introduction of defects are essential to improve hydrogen storage properties [4]. Recently, V-rich BCC alloys were synthesized directly from mechanical alloying (MA) process [5]. M. Balcerzak has shown that the application of MA technique resulted in formation of nanocrystalline TieV alloys absorbing more than 3.6 wt% of hydrogen without any incubation time. Unfortunately, MA TieV BCC alloys suffer from very high hydride stability and the practical application is limited by high price of V [5]. Hydrogen storage and electrochemical properties can be significantly improved by chemical modification of BCC alloy by different elements like Mn, Fe, Co, Al, Zr, Pd, O, Ce, Dy, Cr [6e14]. Addition of Fe is effective in hydrogen sorption/
E-mail address: [email protected]. https://doi.org/10.1016/j.ijhydene.2018.03.123 0360-3199/© 2018 Hydrogen Energy Publications LLC. Published by Elsevier Ltd. All rights reserved.
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desorption plateau flattening and promotes formation of bphase hydride. The plateau pressure increases with increasing Fe content in TieVeFe alloy. Moreover, activation of iron or zirconium containing alloys is very easy [15e18]. Small amount (1 atomic %) of iron can also improve cyclic durability of BCC alloys [19,20]. Chemical modification by Pd atoms can improve activation properties and cycling stability which is attributed to high electrocatalytic activity and anticorrosion ability of Pd [9]. Cr content in BCC phase is important for improving cycle durability [19]. Zr addition can be used to decrease the hysteresis loss without decreasing hydrogen storage capacity of material [18]. Moreover, V can be also partially replaced by cheaper Zr, VFe, Fe, Mn to reduce the price of hydrogen storage materials [17]. One of elements which can be used for chemical modification of BCC alloys is Ni. Addition of right amount of this element to the alloy composition can improve the maximum discharge capacity, cyclic stability, high-rate dischargeability, kinetic properties of the alloy electrodes and decrease the hydrogen desorption temperature [8,21e23]. However, most of the papers describe the influence of Ni addition/substitution on properties of multi phase V-rich alloys e with BCC as main phase and Laves C14, TiNi, Ti2Ni or other BCC solid solution as secondary phase [21e26]. Presented materials properties resulted from multi phase structure and the exact impact of each phase on hydrogen storage properties is difficult to define. There are no information on the effect of V substitution by Ni atoms in single phase nanocrystalline BCC alloy. Understanding of the effect of chemical composition and structure of single phase BCC alloys on hydrogenation/dehydrogenation properties can provide information on suitability of these materials as hydrogen/energy storage systems. For this reason, in this paper, hydrogen sorption/desorption and electrochemical properties of single phase nanocystalline Ti0.5V1.5xNix (x ¼ 0, 0.1, 0.2, 0.3) alloys were systematically investigated.
Methods Used materials High purity commercial powders were used for the synthesis: Titanium (Alfa Aesar, 325 mesh, 99.5%), Vanadium (Alfa Aesar, 325 mesh, 99.5%), Nickel (Aldrich, 5 mm, 99.99%). For electrochemical measurements a 6 M KOH solution was used [27].
Synthesis of materials MA was used to produce Ti0.5V1.5xNix alloys powders (x ¼ 0, 0.1, 0.2, 0.3). Chemical composition of unmodified TieV alloy was selected based on data reported in previous paper e Ti0.5V1.5 alloy was characterized by the lowest dehydrogenation temperature. Synthesis of all materials lasted 14 h (SPEX 8000 M ball mill). Ball to powder weight ratio was about 6.5:1. Argon was a MA atmosphere. Detailed description of the synthesis can be found in previously published paper [5].
Microstructural analysis Microstructural measurements of synthesized Ti0.5V1.5xNix alloys were performed in order to find correlation between materials structures and observed hydrogen sorption/ desorption and electrochemical properties. X-ray diffraction (XRD) studies were done using Panalytical Empyrean powder diffractometer (Cu Ka1 radiation, l ¼ 1.5405952 A, 45 kV, 40 mA, room temperature). XRD measurements were performed to determine the alloys structure after 14 h of synthesis. Moreover, to understand the hydrogen desorption process, XRD patterns were also obtained for each fully hydrogenated alloy. Rietveld method using the Maud software was used to refine XRD patterns. Average crystallite size and lattice strain were estimated using Williamson-Hall style plot. Scanning electron microscopy (SEM) e Tescan digital microscopy imaging VEGA TS5135, was used to study the morphology and microstructure of TieVeNi alloys particles (secondary electrons images, 20 kV). Moreover, the elements compositions of synthesized materials were obtained from energy dispersive spectroscopy (EDS). SEM pictures were used to create particle size dimension distribution histograms. The detailed microstructure of synthesized Ti0.5V1.5 alloy was studied by high resolution transmission electron microscope e HRTEM (JOEL ARM 200F microscope).
Hydrogen absorption/desorption measurements Ti0.5V1.5xNix alloys were tested by a Sievert's device at near room temperature (Particulate Systems, HPVA-200). The mass of each tested sample was about 1 g. The detailed description of used device was included in previously published paper [5]. There was no activation of studied materials before hydrogen sorption and desorption measurements. Hydrogen absorption kinetics studies (time-depending curves) were performed three times for each of synthesized material (the initial hydrogen pressure was 3 MPa and the measurements temperature was 303 K). Between each measurement, hydrogen was released from the alloys during the degassing process e 673 K, vacuum. After kinetic studies, dehydrogenated (673 K, vacuum) materials were measured to obtain Pressure-Composition-Isotherm (PCI) curves of Ti0.5V1.5xNix alloys e 303 K, up to 7 MPa of hydrogen. More details on hydrogen absorption/desorption studies are included in other work [5].
Differential scanning calorimetry (DSC) studies Fully hydrogenated Ti0.5V1.5xNix alloys were studied by DSC e TA DSC Q20. Each specimen used for DSC measurements is the one used for the PCI measurements. All measurements were performed in argon atmosphere up to 873 K. 10 K/min was the set heating rate.
Electrochemical measurements All TieVeNi alloys with addition of carbonyl nickel powder (10 wt%) were used as negative electrodes in Ni-MHx systems. Discharge capacities are presented as those per unit gram of
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measured material. Electrochemical studies were done at room temperature in a three electrode open cell. Ni(OH)2/ NiOOH was used as counter electrode while Hg/HgO was used as reference electrode. 6 M KOH was used as electrolyte prepared from KOH flakes and 18 MU cm1 water. Electrodes were charged and discharged at a current 40 mAg1. The cutoff voltage was 0.7 V vs. Hg/HgO reference electrode. Detailed description of electrochemical measurements was included in previous work [27]. Cycle stability of materials were evaluated by capacity retaining rate Rh after 50th cycle. Rh ¼ (C50/Cmax) 100%, where C50 and Cmax are discharge capacities at the 50th cycle and maximum discharge capacity, respectively.
Results and discussion Microstructure, structure and phase composition Fig. 1 shows XRD patterns of Ti0.5V1.5xNix alloys powders (x ¼ 0, 0.1, 0.2, 0.3) after 14 h of MA process. All samples are single-phase TieV or TieVeNi BCC solid solution materials BCC structure which crystallizes in a Im-3m space group. The phase composition of TieVeNi alloys strongly depends on chemical composition and preparation method. The MA process was also used in the past to synthesize Ti0.73V1.4Ni0.27 alloy. 30 h of milling resulted in formation of
Fig. 1 e XRD patterns of TieV and TieVeNi alloys after 14 h of mechanical alloying (Cu Ka1, 40 mA, 45 kV).
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single meta-stable BCC phase. However, the heat treatment of MA alloy at 973 K for 1 h resulted in formation of two phase material, composed of BCC solid solution and Ti2Ni intermetallic compound [25]. In other work Dagher et al. milled powders with Ti61V21Ni17 chemical composition for 20 h. The obtained material was composed of amorphous phase, V solid solution and Ti-a solid solution. Hydrogenation of this alloy resulted in formation of Ti and Ti2Ni based hydride phases [28]. The hydrogen storage properties of multi phase materials presented in mentioned work are much worse than properties of nanocrystalline singe phase alloys presented in current paper. The single BCC phase was also obtained for vacuum induction melted TiV2.1Nix (x ¼ 0.1e0.3) alloys. However, alloys with Ni content higher than 0.3 were co-composed with TiNi secondary phase which strongly affected the properties of BCC alloy [22]. An arc melted method was used for production of microcrystalline single BCC phase TiV3.4Ni0.6 alloy [29]. However, Tsukahara et al. reported that arc melted TieVeNi alloys with higher Ni content consist of secondary BCC minor phase [24]. These, authors reported that hydrogen capacity is reduced with higher content of BCC minor phase in the material (its abundance increases with increasing of Ni content). The intensity of BCC peaks decreases with increase of Ni content in alloys. Moreover, peaks positions are shifted to higher angles while the Ni content in Ti0.5V1.5xNix alloys is increasing, indicating the unit cell shrinks (Fig. 1). A lattice parameter values as a function of Ti0.5V1.5xNix alloys chemical compositions are shown in Fig. 2. Rietveld refinement method was used to calculate lattice parameters values (Table 1). Observed reduction of lattice parameter values with higher Ni content in TieVeNi alloys can be determined by a linear function. This is due to the fact that nickel which substitutes V in BCC lattice is characterized by atomic radius smaller than vanadium (0.110 nm and 0.134 nm respectively). Based on XRD data, crystallite size and lattice strain were calculated. The average crystallite size was increased from 5.4 nm for unmodified Ti0.5V1.5 alloy to 31.9 nm for Ti0.5V1.2Ni0.3 alloy. A lattice strain value was increased after partial substitution of V by Ni atoms (Table 2). A chemical composition of materials based on elements weights before MA process and on EDS studies are shown in Table 3. Although, these compositions are not the same, the
Fig. 2 e Lattice parameter as a function of x in Ti0.5V1.5¡xNix. Lattice parameter values are from Rietveld refinement.
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Table 1 e Ti0.5V1.5¡xNix structural parameters (from Rietveld refinement of XRD data). Sample a (A) V ( A3) Rwp Ti0.5V1.5 Ti0.5V1.4Ni0.1 Ti0.5V1.3Ni0.2 Ti0.5V1.2Ni0.3
3.093 3.079 3.068 3.054
29.59 29.19 28.88 28.48
4.3 4.5 6.8 6.6
Rwp-final weighted average Bragg R-factor.
Table 2 e Average particle size, crystallites size and lattice strain of Ti0.5V1.5¡xNix alloys (from SEM micrographs and Wiliamson-Hall style plot). Sample
Average particle size (mm)
Crystallites size (nm)
Lattice strain (%)
14.6 8.4 10.4 10.7
5.4 6.3 21.2 31.9
1.1 1.2 2.1 2.0
Ti0.5V1.5 Ti0.5V1.4Ni0.1 Ti0.5V1.3Ni0.2 Ti0.5V1.2Ni0.3
differences between them are rather small and should be related to accuracy of EDS method. Therefore, it was inferred the chemical compositions of synthesized alloys are close to assumed. Fig. 3 shows SEM micrographs of MA TieVeNi materials. Studied alloys are composed of irregular particles with porous, cleavage and fractured morphology. Moreover, the materials particle size distributions have bimodal character e small grains with size around 1 mm are agglomerated in bigger particles (with size about 50 mm). Insets of Fig. 3 show particle size dimension distribution histograms of Ti0.5V1.5xNix alloys. Table 2 shows the average particle sizes of studied materials. The particles sizes are reduced after substitution of V by Ni atoms - from 14.6 mm for Ti0.5V1.5 alloy to 8.4e10.7 mm for TieVeNi alloys. The average particle size doesn't depend on the amount of Ni in the alloy. Fig. 4 shows a HRTEM image of unmodified Ti0.5V1.5 alloy. This material is characterized by multicrystalline structure. It means that differently oriented nanocrystals are
Table 3 e Chemical composition of Ti0.5V1.5¡xNix alloys determined from materials weighted before synthesis and from EDS studies. Sample
Element
Ti0.5V1.5
Ti V Ti V Ni Ti V Ni Ti V Ni
Ti0.5V1.4Ni0.1
Ti0.5V1.3Ni0.2
Ti0.5V1.2Ni0.3
Weight Weight before based synthesis on EDS (%) (%) 23.88 76.12 23.69 70.48 5.83 23.51 64.94 11.55 23.18 59.12 17.70
24.86 75.14 25.19 70.11 4,70 24.62 66.38 9.00 24.60 60.83 14.57
Formula obtained from EDS Ti0.5V1.42 Ti0.5V1.39Ni0.09
Ti0.5V1.35Ni0.18
Ti0.5V1.24Ni0.3
Fig. 3 e SEM micrograph and corresponding particle size dimension distribution histograms (insets) of TieV and TieVeNi (secondary electrons images, 20 kV).
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Fig. 4 e HRTEM image of Ti0.5V1.5 alloy.
agglomerated in more complex forms. Crystallites sizes observed on HRTEM image and values calculated based on Williamson-Hall style plot are of the same order of magnitude. Furthermore, the d(100) spacing value is nearly the same as that obtained from Rietveld refinement (Fig. 4).
Hydrogenation properties Table 4 shows most significant data from hydrogenation/ dehydrogenation studies of TieVeNi alloys. It should be noticed that these studies were performed without any activation process. Time-capacity curves of Ti0.5V1.5xNix alloys in activation cycles/kinetic studies are shown in Fig. 5. All of studied alloys were unactivated in the first activation/ kinetic measurement. Such effect can be most likely connected with an oxide layer at the surface of the alloys particles which inhibits the hydrogen sorption process [6,30]. S. Challet et al. considered a thermal treatment (heating to 773 K under vacuum) as a method of alloy activation. In this work a similar approach has been applied. First cycle of kinetic measurement with following degassing procedure (heating to 673 K under vacuum) can be considered as activation of studied alloys.
Studied TieVeNi alloys showed much better hydrogenation properties in the second cycle of measurements. The only exception is Ti0.5V1.2Ni0.3 alloy which needed three cycles of hydrogen absorption-desorption to perform the highest hydrogen concentration. However, the differences between the values obtained for all studied materials at second and third cycle are rather small and can be connected to accuracy of used method. The maximum hydrogen concentration value was reduced from 2.75 wt% for Ti0.5V1.5 alloy to 1.63 wt% for Ti0.5V1.2Ni0.3 (Table 4). The reason of this phenomenon is the Ni decreases the lattice parameter of BCC phase reducing the interstitial sites for hydrogen storage and thus the stability of the alloy hydride is decreased. The hydrogen sorption time (95% of maximum hydrogen capacity) depends on the Ni content in the alloy. It was reduced from about 60 min for Ti0.5V1.5 alloy to about 30 min for Ni-rich alloy. TieVeNi alloys absorbed hydrogen without any incubation time. Hydrogenation kinetics of studied materials is much faster than kinetics of presented for amorphous-based TieVeNi alloys [28]. PeC isotherms obtained at 303 K for TieV and TieVeNi alloys are shown on Fig. 6. The shape of curves hinders the determination of plateau pressures. However, based on the data published previously it can be considered that the plateau pressures are lower than the lowest equilibrium pressure obtained during these measurements [5]. Moreover, it was shown in the past that the plateau pressure is far below 0.1 MPa [24,31]. The very low plateau pressures indicate the high stability of these hydrides. On the other hand, as stated K. Nomura et al. low absorption pressure is desired for tritium recovery [15]. An increase of slope with higher Ni content in the materials can be observed. Moreover, the hysteresis between sorption and desorption curves decreases with increasing the content of V substituted by Ni atoms. The shown tendency of hydrogen storage capacity changes on chemical composition corresponds to results of activation/kinetic measurements. A similar dependence was also shown in the past for arc melted TieVeNi alloys, which indicates that the influence of Ni substitution/addition can be similar in wider range of TieVeNi chemical compositions [24,31].
Table 4 e Summarized data from PCI and DSC measurements analyses on Ti0.5V1.5¡xNix samples. Sample
Kinetic measurements H (wt%)
Acycles
RH/DH (%)
TH-des (K)
532.2 631.2 716.4 750.4 750.0 766.1
T95 (min)
1 cycle
2 cycle
3 cycle
1 cycle
2 cycle
3 cycle
Ti0.5V1.5
2
0.02
2.75
2.72
e
61
79
11
Ti0.5V1.4Ni0.1 Ti0.5V1.3Ni0.2 Ti0.5V1.2Ni0.3
3 2 2
0.29 0.07 0.17
2.23 1.79 1.61
2.23 1.79 1.63
e e e
73 45 32
66 48 33
19 18 21
Acycles e number of activation cycles needed to obtain best kinetic properties. H e maximum hydrogen concentration. T95 e time needed to reach 95% of maximum hydrogen capacity. RH/DH e reversibility of hydriding-dehydriding process. TH-des e hydrogen desorption peak temperature.
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Fig. 5 e Time-capacity curves of TieV and TieVeNi alloys in activation/kinetic measurements (3 MPa of initial hydrogen pressure, 303 K). Between each cycle, samples were degassed at 673 K.
A dehydrogenation properties of measured TieVeNi materials are quantitatively presented by reversibility of hydriding-dehydriding process (Table 4). The reversibility of studied Ti0.5V1.5xNix alloys at 303K is very low. A similar situation was observed for other BCC-based alloys [32]. However, it should be noticed, the higher Ni content in studied material is, the greater hydrogenation/dehydrogenation reversibility was obtained. Unfortunately, the desorption plateau pressures are lower than the lowest equilibrium pressure which can be measured on used Sievert's device. This is the reason, why the hydrogen desorption plateau region is not visible on presented curves.
Fig. 6 e PeC Isotherms at 303 K of TieV and TieVeNi alloys (obtained on degassed samples). Arrows indicate direction of hydrogen concentration changes (sorption/desorption).
TieVeNi samples after hydrogenation/dehydrogenation measurements were fully hydrogenated in order to perform DSC and XRD studies. XRD patterns of Ti0.5V1.5xNix alloys after hydrogenation are shown in Fig. 7. All alloys are composed of distorted BCC phase - a-phase. Moreover, unmodified Ti0.5V1.5 and modified Ti0.5V1.4Ni0.1 alloys show a two phases structure e with additional body-centered-tetragonal (BCT) phase. Accordingly, to Rietveld refinement, the cell parameter of BCC phase (for each alloy) was enlarged by about 0.1 A during formation of a-phase hydride solid solution. It is due to the partial occupancy of interstitial sites by H atoms. Maximum hydrogen capacity of this solid solution is 1.5 wt% [6].
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Fig. 7 e XRD patterns of fully hydrogenated TieV and Tie VeNi alloys (Cu Ka1, 40 mA, 45 kV). However, the release of hydrogen from this hydride phase is impossible at 303 K [24]. A BCT phase is a monohydride structure (I4/mmm space group) which corresponds to a BCC phase elongated along the c axis e lattice parameter a remains nearly constant where as c increases with increasing hydrogen concentration in the alloy [33]. BCT phase can store 1.5e1.9 wt% of hydrogen [6]. It can be noticed, the BCT phase abundance decreased with increasing of Ni in Ti0.5V1.5xNix alloys e from 32.3% for Ti0.5V1.5 to 0% for Ti0.5V1.3Ni0.2 alloy. Presented above results of XRD studies fit well to the maximum hydrogen concentrations obtained from activation/kinetic and PCI measurements. However, the hydrogen concentration of Ti0.5V1.5 alloy exceeds the maximum hydrogen capacity of BCT phase. For this reason it can be considered that during hydrogenation this alloy forms also hydride with higher hydrogen concentration e above 1.9 wt%. It is most likely a fully hydrogenated face-centered-cubic (FCC) phase (Fm-3m space group) which can store 1.9e4 wt% of hydrogen [6]. The absence of FCC hydride phase peaks on XRD pattern is probably caused by its small content in Ti0.5V1.5 alloy. The FCC hydride phase was obtained also for fully hydrogenated arc melted Ti2xCrVNix alloys (x ¼ 0, 0.1, 0.2) [8]. As these studies have shown, the contraction of cell volume due
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to Ni substitution, can be compressed by expansion due to high hydrogen sorption. DSC profiles of Ti0.5V1.5xNix alloys hydrides are shown in Fig. 8. The hydrogen desorption peaks temperatures are listed in Table 4. The thermal decomposition of Ti0.5V1.5 alloy hydride consists of three reactions. The low temperature (532.2 K) decomposition process corresponds probably to the FCC full hydride transfer to BCT monohydride phase. In the next decomposition step (631.2 K) the monohydride transferred to BCC hydride phase. Finally, the high temperature peak (716.4 K) corresponds to decomposition of hydrogenated BCC phase to dehydrogenated BCC phase. The three steps decomposition process confirms the presence of additional to BCC and BCT hydride phase e which is most probably FCC hydride phase. However, additional studies must be done to fully prove the phase transformations during hydrogen sorption and desorption. Kumar et al. reported the same hydride decomposition process for TieCreCeCo alloy. It indicates that described dehydrogenation process can be common for all BCC solid solutions characterized by high hydrogen storage capacity [8]. It can be stated that FCC hydride phase is the most desired for hydrogen storage applications due to the lowest thermal stability and high hydrogen content. For TieVeNi alloys only one clearly visible endothermic peak was observed. It corresponds to hydrogen desorption reaction from BCC hydride to dehydrogenated BCC solid solution. It is important to note that the peak temperature increased from 716.4 K for TieV alloy to 766.1 K for Ni-rich alloy. This very high decomposition temperature proves the high hydride stability presented on desorption part of PCI curves. To consider the practical use of studied alloys the dehydrogenation temperature must be reduced. It was shown in the past that dehydrogenation peak temperature can be lowered by cycled hydrogenation/dehydrogenation process which causes reduction of materials particle size [34]. Moreover, the hydrides decomposition temperature can be also lowered by surface modification of TieVeNi alloy by Raney Ni due to reduction of dehydrogenation process energy barrier [35]. The enthalpy of dehydrogenation process reached 240.3, 154.3, 102.3 and 60.2 J/g respectively for x ¼ 0, 0.1, 0.2 and 0.3. It
Fig. 8 e DSC profile of the hydrides of Ti0.5V1.5¡xNix alloys e endothermic downwards, exothermic upwards (argon flow, 10K/min heating rate).
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Fig. 9 e Discharge capacities as a function of cycle number of electrodes prepared with Ti0.5V1.5¡xNix alloys mixed with carbonyl nickel powder (10 wt%). Electrodes were charged and discharged at 40 mAh/g. The cut off potential was ¡0.7 V.
Table 5 e Activation cycles, maximum discharge capacity and capacity retaining rate after 50th cycle of Ti0.5V1.5¡xNix (40mAh/g charging-discharging current, ¡0.7 V cut off potential). Sample
Ti0.5V1.4Ni0.1 Ti0.5V1.3Ni0.2 Ti0.5V1.2Ni0.3
Activation cycles
Maximum discharge capacity [mAh/g]
Capacity retaining rate after 50th cycle [%]
7 5 8
72 97 116
41 69 75
implies that addition of Ni can be effective in reducing the stability of a phase.
Electrochemical properties Fig. 9 shows results of electrochemical measurements of Tie VeNi alloys. The most important data from these studies were summarized in Table 5. Ti0.5V1.5 alloy is not able to desorb electrochemically hydrogen in KOH electrolyte. However, a great improve of electrochemical properties was obtained after partial substitution of V by Ni atoms. Ti0.5V1.5xNix alloys displayed maximum capacity between 5th and 8th cycle. The tendency of activation properties changes depending on Ni amount in BCC alloys was not noticed. These are much better activation properties than previously published for rapidly solidified BCC-based alloy electrodes [36]. It is most likely due to differences in morphology of materials e small grains of rapidly solidified alloys can suppress pulverization of the electrode materials and thus worsen the activation properties. The maximum discharge capacity increased with increasing Ni content in alloys to reach 116 mAh/g for Ti0.5V1.2Ni0.3 alloy. It is due to the well known catalytic properties of Ni [37]. The same tendency was also observed for vacuum induction melted V2.1TiNix (x ¼ 0.1e0.9) alloys [22]. It can be seen that discharge capacities of TieVeNi alloys were degraded during testing (with increasing charge/ discharge cycles). The reason of degradation can be connected with several factors:
dissolution of V and Ti (mainly V) from the electrode material into the KOH electrolyte [24,38,39]. oxidation of alloy particle surface (formation of TiO2 layer) which causes very low hydrogen diffusivity and electrical conductivity [39]. On the other hand, as stated C. Iwakura et al. the formation of oxide layer suppresses the dissolution of V atoms from MgeVeNi alloys into the alkaline solution [40]. pulverization caused by the cell volume expansion and shrinkage during hydrogenation and dehydrogenation. reduction of electrode mass e a part of material can be flaked away from the electrode during charge/discharge cycles. Despite a decrease in discharge capacity value, which was observed for all TieVeNi alloys, chemical modification of Ti0.5V1.5 by Ni atoms caused increase of cycle stability. Ti0.5V1.2Ni0.3 was characterized by the best cycle stability which reached 75%.
Conclusions The Ti0.5V1.5xNix alloys (x ¼ 0, 0.1, 0.2, 0.3) synthesized by mechanical alloying were studied according to their hydrogen storage and electrochemical properties. XRD measurements showed that all samples are single BCC solid solution phase. Nanocrystallinity of obtained MA alloys was proven by HRTEM studies. XRD and hydrogenation/dehydrogenation measurements showed that hydrogen storage capacities derive directly from chemical compositions and lattice parameters. Maximum hydrogen uptake was reduced with higher Ni content in the TieVeNi alloys. It was due to the reduction of cell volume. On the other hand, Ni addition caused improved hydrogenation kinetics, reduced hysteresis and increased hydrogenation/dehydrogenation reversibility. Electrochemical studies showed that Ni can effectively influence the hydrogen absorption in 6 M KOH solution of TieVeNi alloys. The maximum discharge capacity and capacity retaining rate were significantly improved. Furthermore, it is reasonably to consider the substitution of V by Ni atoms in BCC-based NiMHx to reduce the price of negative electrode e nickel is about ten times cheaper than vanadium is.
Acknowledgement Financial assistance from National Science Centre, Poland (no. 2015/17/N/ST8/00271).
references
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Effect of Ni on electrochemical and hydrogen storage properties of V-rich body-centered-cubic solid solution alloys M. Balcerzak University of Technology, Institute of Materials Science and Engineering, Jana Pawła II No 24, 61-138 Poznan , Poland Poznan
article info
abstract
Article history:
V-rich solid solution alloys are potential candidates for Ni-MHx negative electrodes and
Received 8 December 2017
hydrogen sorbing materials. Mechanical alloying (MA) is used in this paper to produce
Received in revised form
Ti0.5V1.5xNix nanocrystalline alloys (x ¼ 0, 0.1, 0.2, 0.3). A SPEX 8000 M mill is used. The aim
12 March 2018
of this work is to study the effect of chemical modification by Ni on hydrogen sorption/
Accepted 16 March 2018
desorption and electrochemical properties of V-rich body-centered-cubic (BCC) alloys.
Available online 5 April 2018
Presented measurements results show formation of BCC phase after 14 h of MA. The nanocrystallinety of obtained materials is confirmed by high resolution transmission
Keywords:
electron microscopy images. MA alloys are tested by a Sievert's device at near room tem-
Energy storage
perature. Partial substitution of V by Ni causes improved hydrogenation kinetics, reduced
Mechanical alloying
hysteresis and increased hydrogenation/dehydrogenation reversibility. Observed proper-
Hydrogen sorption
ties are mainly due to differences in structures of studied materials. Electrochemical
X-ray diffraction
studies on chemically modified V-rich alloys show that capacity retaining rate and
Differential scanning calorimetry
discharge capacity increase with higher Ni content in the material.
Nanomaterials
© 2018 Hydrogen Energy Publications LLC. Published by Elsevier Ltd. All rights reserved.
Introduction V-rich body-centered-cubic (BCC) solid solution alloys are considered for Ni-MHx negative electrode and hydrogen storage materials. BCC-based alloys are mostly synthesized by conventional methods e arc or induction melting. These preparation methods provide production of microcrystalline materials [1e3]. However, due to the fact that elements are characterized by different melting temperatures, some desired chemical compositions are impossible to obtain using conventional methods. Moreover, microcrystalline alloys suffer from slow hydrogen sorption/desorption kinetics and difficult activation process [1]. The kinetics is closely related to
microstructure of material. Reduction of crystallites size and introduction of defects are essential to improve hydrogen storage properties [4]. Recently, V-rich BCC alloys were synthesized directly from mechanical alloying (MA) process [5]. M. Balcerzak has shown that the application of MA technique resulted in formation of nanocrystalline TieV alloys absorbing more than 3.6 wt% of hydrogen without any incubation time. Unfortunately, MA TieV BCC alloys suffer from very high hydride stability and the practical application is limited by high price of V [5]. Hydrogen storage and electrochemical properties can be significantly improved by chemical modification of BCC alloy by different elements like Mn, Fe, Co, Al, Zr, Pd, O, Ce, Dy, Cr [6e14]. Addition of Fe is effective in hydrogen sorption/
E-mail address: [email protected]. https://doi.org/10.1016/j.ijhydene.2018.03.123 0360-3199/© 2018 Hydrogen Energy Publications LLC. Published by Elsevier Ltd. All rights reserved.
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desorption plateau flattening and promotes formation of bphase hydride. The plateau pressure increases with increasing Fe content in TieVeFe alloy. Moreover, activation of iron or zirconium containing alloys is very easy [15e18]. Small amount (1 atomic %) of iron can also improve cyclic durability of BCC alloys [19,20]. Chemical modification by Pd atoms can improve activation properties and cycling stability which is attributed to high electrocatalytic activity and anticorrosion ability of Pd [9]. Cr content in BCC phase is important for improving cycle durability [19]. Zr addition can be used to decrease the hysteresis loss without decreasing hydrogen storage capacity of material [18]. Moreover, V can be also partially replaced by cheaper Zr, VFe, Fe, Mn to reduce the price of hydrogen storage materials [17]. One of elements which can be used for chemical modification of BCC alloys is Ni. Addition of right amount of this element to the alloy composition can improve the maximum discharge capacity, cyclic stability, high-rate dischargeability, kinetic properties of the alloy electrodes and decrease the hydrogen desorption temperature [8,21e23]. However, most of the papers describe the influence of Ni addition/substitution on properties of multi phase V-rich alloys e with BCC as main phase and Laves C14, TiNi, Ti2Ni or other BCC solid solution as secondary phase [21e26]. Presented materials properties resulted from multi phase structure and the exact impact of each phase on hydrogen storage properties is difficult to define. There are no information on the effect of V substitution by Ni atoms in single phase nanocrystalline BCC alloy. Understanding of the effect of chemical composition and structure of single phase BCC alloys on hydrogenation/dehydrogenation properties can provide information on suitability of these materials as hydrogen/energy storage systems. For this reason, in this paper, hydrogen sorption/desorption and electrochemical properties of single phase nanocystalline Ti0.5V1.5xNix (x ¼ 0, 0.1, 0.2, 0.3) alloys were systematically investigated.
Methods Used materials High purity commercial powders were used for the synthesis: Titanium (Alfa Aesar, 325 mesh, 99.5%), Vanadium (Alfa Aesar, 325 mesh, 99.5%), Nickel (Aldrich, 5 mm, 99.99%). For electrochemical measurements a 6 M KOH solution was used [27].
Synthesis of materials MA was used to produce Ti0.5V1.5xNix alloys powders (x ¼ 0, 0.1, 0.2, 0.3). Chemical composition of unmodified TieV alloy was selected based on data reported in previous paper e Ti0.5V1.5 alloy was characterized by the lowest dehydrogenation temperature. Synthesis of all materials lasted 14 h (SPEX 8000 M ball mill). Ball to powder weight ratio was about 6.5:1. Argon was a MA atmosphere. Detailed description of the synthesis can be found in previously published paper [5].
Microstructural analysis Microstructural measurements of synthesized Ti0.5V1.5xNix alloys were performed in order to find correlation between materials structures and observed hydrogen sorption/ desorption and electrochemical properties. X-ray diffraction (XRD) studies were done using Panalytical Empyrean powder diffractometer (Cu Ka1 radiation, l ¼ 1.5405952 A, 45 kV, 40 mA, room temperature). XRD measurements were performed to determine the alloys structure after 14 h of synthesis. Moreover, to understand the hydrogen desorption process, XRD patterns were also obtained for each fully hydrogenated alloy. Rietveld method using the Maud software was used to refine XRD patterns. Average crystallite size and lattice strain were estimated using Williamson-Hall style plot. Scanning electron microscopy (SEM) e Tescan digital microscopy imaging VEGA TS5135, was used to study the morphology and microstructure of TieVeNi alloys particles (secondary electrons images, 20 kV). Moreover, the elements compositions of synthesized materials were obtained from energy dispersive spectroscopy (EDS). SEM pictures were used to create particle size dimension distribution histograms. The detailed microstructure of synthesized Ti0.5V1.5 alloy was studied by high resolution transmission electron microscope e HRTEM (JOEL ARM 200F microscope).
Hydrogen absorption/desorption measurements Ti0.5V1.5xNix alloys were tested by a Sievert's device at near room temperature (Particulate Systems, HPVA-200). The mass of each tested sample was about 1 g. The detailed description of used device was included in previously published paper [5]. There was no activation of studied materials before hydrogen sorption and desorption measurements. Hydrogen absorption kinetics studies (time-depending curves) were performed three times for each of synthesized material (the initial hydrogen pressure was 3 MPa and the measurements temperature was 303 K). Between each measurement, hydrogen was released from the alloys during the degassing process e 673 K, vacuum. After kinetic studies, dehydrogenated (673 K, vacuum) materials were measured to obtain Pressure-Composition-Isotherm (PCI) curves of Ti0.5V1.5xNix alloys e 303 K, up to 7 MPa of hydrogen. More details on hydrogen absorption/desorption studies are included in other work [5].
Differential scanning calorimetry (DSC) studies Fully hydrogenated Ti0.5V1.5xNix alloys were studied by DSC e TA DSC Q20. Each specimen used for DSC measurements is the one used for the PCI measurements. All measurements were performed in argon atmosphere up to 873 K. 10 K/min was the set heating rate.
Electrochemical measurements All TieVeNi alloys with addition of carbonyl nickel powder (10 wt%) were used as negative electrodes in Ni-MHx systems. Discharge capacities are presented as those per unit gram of
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measured material. Electrochemical studies were done at room temperature in a three electrode open cell. Ni(OH)2/ NiOOH was used as counter electrode while Hg/HgO was used as reference electrode. 6 M KOH was used as electrolyte prepared from KOH flakes and 18 MU cm1 water. Electrodes were charged and discharged at a current 40 mAg1. The cutoff voltage was 0.7 V vs. Hg/HgO reference electrode. Detailed description of electrochemical measurements was included in previous work [27]. Cycle stability of materials were evaluated by capacity retaining rate Rh after 50th cycle. Rh ¼ (C50/Cmax) 100%, where C50 and Cmax are discharge capacities at the 50th cycle and maximum discharge capacity, respectively.
Results and discussion Microstructure, structure and phase composition Fig. 1 shows XRD patterns of Ti0.5V1.5xNix alloys powders (x ¼ 0, 0.1, 0.2, 0.3) after 14 h of MA process. All samples are single-phase TieV or TieVeNi BCC solid solution materials BCC structure which crystallizes in a Im-3m space group. The phase composition of TieVeNi alloys strongly depends on chemical composition and preparation method. The MA process was also used in the past to synthesize Ti0.73V1.4Ni0.27 alloy. 30 h of milling resulted in formation of
Fig. 1 e XRD patterns of TieV and TieVeNi alloys after 14 h of mechanical alloying (Cu Ka1, 40 mA, 45 kV).
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single meta-stable BCC phase. However, the heat treatment of MA alloy at 973 K for 1 h resulted in formation of two phase material, composed of BCC solid solution and Ti2Ni intermetallic compound [25]. In other work Dagher et al. milled powders with Ti61V21Ni17 chemical composition for 20 h. The obtained material was composed of amorphous phase, V solid solution and Ti-a solid solution. Hydrogenation of this alloy resulted in formation of Ti and Ti2Ni based hydride phases [28]. The hydrogen storage properties of multi phase materials presented in mentioned work are much worse than properties of nanocrystalline singe phase alloys presented in current paper. The single BCC phase was also obtained for vacuum induction melted TiV2.1Nix (x ¼ 0.1e0.3) alloys. However, alloys with Ni content higher than 0.3 were co-composed with TiNi secondary phase which strongly affected the properties of BCC alloy [22]. An arc melted method was used for production of microcrystalline single BCC phase TiV3.4Ni0.6 alloy [29]. However, Tsukahara et al. reported that arc melted TieVeNi alloys with higher Ni content consist of secondary BCC minor phase [24]. These, authors reported that hydrogen capacity is reduced with higher content of BCC minor phase in the material (its abundance increases with increasing of Ni content). The intensity of BCC peaks decreases with increase of Ni content in alloys. Moreover, peaks positions are shifted to higher angles while the Ni content in Ti0.5V1.5xNix alloys is increasing, indicating the unit cell shrinks (Fig. 1). A lattice parameter values as a function of Ti0.5V1.5xNix alloys chemical compositions are shown in Fig. 2. Rietveld refinement method was used to calculate lattice parameters values (Table 1). Observed reduction of lattice parameter values with higher Ni content in TieVeNi alloys can be determined by a linear function. This is due to the fact that nickel which substitutes V in BCC lattice is characterized by atomic radius smaller than vanadium (0.110 nm and 0.134 nm respectively). Based on XRD data, crystallite size and lattice strain were calculated. The average crystallite size was increased from 5.4 nm for unmodified Ti0.5V1.5 alloy to 31.9 nm for Ti0.5V1.2Ni0.3 alloy. A lattice strain value was increased after partial substitution of V by Ni atoms (Table 2). A chemical composition of materials based on elements weights before MA process and on EDS studies are shown in Table 3. Although, these compositions are not the same, the
Fig. 2 e Lattice parameter as a function of x in Ti0.5V1.5¡xNix. Lattice parameter values are from Rietveld refinement.
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Table 1 e Ti0.5V1.5¡xNix structural parameters (from Rietveld refinement of XRD data). Sample a (A) V ( A3) Rwp Ti0.5V1.5 Ti0.5V1.4Ni0.1 Ti0.5V1.3Ni0.2 Ti0.5V1.2Ni0.3
3.093 3.079 3.068 3.054
29.59 29.19 28.88 28.48
4.3 4.5 6.8 6.6
Rwp-final weighted average Bragg R-factor.
Table 2 e Average particle size, crystallites size and lattice strain of Ti0.5V1.5¡xNix alloys (from SEM micrographs and Wiliamson-Hall style plot). Sample
Average particle size (mm)
Crystallites size (nm)
Lattice strain (%)
14.6 8.4 10.4 10.7
5.4 6.3 21.2 31.9
1.1 1.2 2.1 2.0
Ti0.5V1.5 Ti0.5V1.4Ni0.1 Ti0.5V1.3Ni0.2 Ti0.5V1.2Ni0.3
differences between them are rather small and should be related to accuracy of EDS method. Therefore, it was inferred the chemical compositions of synthesized alloys are close to assumed. Fig. 3 shows SEM micrographs of MA TieVeNi materials. Studied alloys are composed of irregular particles with porous, cleavage and fractured morphology. Moreover, the materials particle size distributions have bimodal character e small grains with size around 1 mm are agglomerated in bigger particles (with size about 50 mm). Insets of Fig. 3 show particle size dimension distribution histograms of Ti0.5V1.5xNix alloys. Table 2 shows the average particle sizes of studied materials. The particles sizes are reduced after substitution of V by Ni atoms - from 14.6 mm for Ti0.5V1.5 alloy to 8.4e10.7 mm for TieVeNi alloys. The average particle size doesn't depend on the amount of Ni in the alloy. Fig. 4 shows a HRTEM image of unmodified Ti0.5V1.5 alloy. This material is characterized by multicrystalline structure. It means that differently oriented nanocrystals are
Table 3 e Chemical composition of Ti0.5V1.5¡xNix alloys determined from materials weighted before synthesis and from EDS studies. Sample
Element
Ti0.5V1.5
Ti V Ti V Ni Ti V Ni Ti V Ni
Ti0.5V1.4Ni0.1
Ti0.5V1.3Ni0.2
Ti0.5V1.2Ni0.3
Weight Weight before based synthesis on EDS (%) (%) 23.88 76.12 23.69 70.48 5.83 23.51 64.94 11.55 23.18 59.12 17.70
24.86 75.14 25.19 70.11 4,70 24.62 66.38 9.00 24.60 60.83 14.57
Formula obtained from EDS Ti0.5V1.42 Ti0.5V1.39Ni0.09
Ti0.5V1.35Ni0.18
Ti0.5V1.24Ni0.3
Fig. 3 e SEM micrograph and corresponding particle size dimension distribution histograms (insets) of TieV and TieVeNi (secondary electrons images, 20 kV).
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Fig. 4 e HRTEM image of Ti0.5V1.5 alloy.
agglomerated in more complex forms. Crystallites sizes observed on HRTEM image and values calculated based on Williamson-Hall style plot are of the same order of magnitude. Furthermore, the d(100) spacing value is nearly the same as that obtained from Rietveld refinement (Fig. 4).
Hydrogenation properties Table 4 shows most significant data from hydrogenation/ dehydrogenation studies of TieVeNi alloys. It should be noticed that these studies were performed without any activation process. Time-capacity curves of Ti0.5V1.5xNix alloys in activation cycles/kinetic studies are shown in Fig. 5. All of studied alloys were unactivated in the first activation/ kinetic measurement. Such effect can be most likely connected with an oxide layer at the surface of the alloys particles which inhibits the hydrogen sorption process [6,30]. S. Challet et al. considered a thermal treatment (heating to 773 K under vacuum) as a method of alloy activation. In this work a similar approach has been applied. First cycle of kinetic measurement with following degassing procedure (heating to 673 K under vacuum) can be considered as activation of studied alloys.
Studied TieVeNi alloys showed much better hydrogenation properties in the second cycle of measurements. The only exception is Ti0.5V1.2Ni0.3 alloy which needed three cycles of hydrogen absorption-desorption to perform the highest hydrogen concentration. However, the differences between the values obtained for all studied materials at second and third cycle are rather small and can be connected to accuracy of used method. The maximum hydrogen concentration value was reduced from 2.75 wt% for Ti0.5V1.5 alloy to 1.63 wt% for Ti0.5V1.2Ni0.3 (Table 4). The reason of this phenomenon is the Ni decreases the lattice parameter of BCC phase reducing the interstitial sites for hydrogen storage and thus the stability of the alloy hydride is decreased. The hydrogen sorption time (95% of maximum hydrogen capacity) depends on the Ni content in the alloy. It was reduced from about 60 min for Ti0.5V1.5 alloy to about 30 min for Ni-rich alloy. TieVeNi alloys absorbed hydrogen without any incubation time. Hydrogenation kinetics of studied materials is much faster than kinetics of presented for amorphous-based TieVeNi alloys [28]. PeC isotherms obtained at 303 K for TieV and TieVeNi alloys are shown on Fig. 6. The shape of curves hinders the determination of plateau pressures. However, based on the data published previously it can be considered that the plateau pressures are lower than the lowest equilibrium pressure obtained during these measurements [5]. Moreover, it was shown in the past that the plateau pressure is far below 0.1 MPa [24,31]. The very low plateau pressures indicate the high stability of these hydrides. On the other hand, as stated K. Nomura et al. low absorption pressure is desired for tritium recovery [15]. An increase of slope with higher Ni content in the materials can be observed. Moreover, the hysteresis between sorption and desorption curves decreases with increasing the content of V substituted by Ni atoms. The shown tendency of hydrogen storage capacity changes on chemical composition corresponds to results of activation/kinetic measurements. A similar dependence was also shown in the past for arc melted TieVeNi alloys, which indicates that the influence of Ni substitution/addition can be similar in wider range of TieVeNi chemical compositions [24,31].
Table 4 e Summarized data from PCI and DSC measurements analyses on Ti0.5V1.5¡xNix samples. Sample
Kinetic measurements H (wt%)
Acycles
RH/DH (%)
TH-des (K)
532.2 631.2 716.4 750.4 750.0 766.1
T95 (min)
1 cycle
2 cycle
3 cycle
1 cycle
2 cycle
3 cycle
Ti0.5V1.5
2
0.02
2.75
2.72
e
61
79
11
Ti0.5V1.4Ni0.1 Ti0.5V1.3Ni0.2 Ti0.5V1.2Ni0.3
3 2 2
0.29 0.07 0.17
2.23 1.79 1.61
2.23 1.79 1.63
e e e
73 45 32
66 48 33
19 18 21
Acycles e number of activation cycles needed to obtain best kinetic properties. H e maximum hydrogen concentration. T95 e time needed to reach 95% of maximum hydrogen capacity. RH/DH e reversibility of hydriding-dehydriding process. TH-des e hydrogen desorption peak temperature.
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Fig. 5 e Time-capacity curves of TieV and TieVeNi alloys in activation/kinetic measurements (3 MPa of initial hydrogen pressure, 303 K). Between each cycle, samples were degassed at 673 K.
A dehydrogenation properties of measured TieVeNi materials are quantitatively presented by reversibility of hydriding-dehydriding process (Table 4). The reversibility of studied Ti0.5V1.5xNix alloys at 303K is very low. A similar situation was observed for other BCC-based alloys [32]. However, it should be noticed, the higher Ni content in studied material is, the greater hydrogenation/dehydrogenation reversibility was obtained. Unfortunately, the desorption plateau pressures are lower than the lowest equilibrium pressure which can be measured on used Sievert's device. This is the reason, why the hydrogen desorption plateau region is not visible on presented curves.
Fig. 6 e PeC Isotherms at 303 K of TieV and TieVeNi alloys (obtained on degassed samples). Arrows indicate direction of hydrogen concentration changes (sorption/desorption).
TieVeNi samples after hydrogenation/dehydrogenation measurements were fully hydrogenated in order to perform DSC and XRD studies. XRD patterns of Ti0.5V1.5xNix alloys after hydrogenation are shown in Fig. 7. All alloys are composed of distorted BCC phase - a-phase. Moreover, unmodified Ti0.5V1.5 and modified Ti0.5V1.4Ni0.1 alloys show a two phases structure e with additional body-centered-tetragonal (BCT) phase. Accordingly, to Rietveld refinement, the cell parameter of BCC phase (for each alloy) was enlarged by about 0.1 A during formation of a-phase hydride solid solution. It is due to the partial occupancy of interstitial sites by H atoms. Maximum hydrogen capacity of this solid solution is 1.5 wt% [6].
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Fig. 7 e XRD patterns of fully hydrogenated TieV and Tie VeNi alloys (Cu Ka1, 40 mA, 45 kV). However, the release of hydrogen from this hydride phase is impossible at 303 K [24]. A BCT phase is a monohydride structure (I4/mmm space group) which corresponds to a BCC phase elongated along the c axis e lattice parameter a remains nearly constant where as c increases with increasing hydrogen concentration in the alloy [33]. BCT phase can store 1.5e1.9 wt% of hydrogen [6]. It can be noticed, the BCT phase abundance decreased with increasing of Ni in Ti0.5V1.5xNix alloys e from 32.3% for Ti0.5V1.5 to 0% for Ti0.5V1.3Ni0.2 alloy. Presented above results of XRD studies fit well to the maximum hydrogen concentrations obtained from activation/kinetic and PCI measurements. However, the hydrogen concentration of Ti0.5V1.5 alloy exceeds the maximum hydrogen capacity of BCT phase. For this reason it can be considered that during hydrogenation this alloy forms also hydride with higher hydrogen concentration e above 1.9 wt%. It is most likely a fully hydrogenated face-centered-cubic (FCC) phase (Fm-3m space group) which can store 1.9e4 wt% of hydrogen [6]. The absence of FCC hydride phase peaks on XRD pattern is probably caused by its small content in Ti0.5V1.5 alloy. The FCC hydride phase was obtained also for fully hydrogenated arc melted Ti2xCrVNix alloys (x ¼ 0, 0.1, 0.2) [8]. As these studies have shown, the contraction of cell volume due
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to Ni substitution, can be compressed by expansion due to high hydrogen sorption. DSC profiles of Ti0.5V1.5xNix alloys hydrides are shown in Fig. 8. The hydrogen desorption peaks temperatures are listed in Table 4. The thermal decomposition of Ti0.5V1.5 alloy hydride consists of three reactions. The low temperature (532.2 K) decomposition process corresponds probably to the FCC full hydride transfer to BCT monohydride phase. In the next decomposition step (631.2 K) the monohydride transferred to BCC hydride phase. Finally, the high temperature peak (716.4 K) corresponds to decomposition of hydrogenated BCC phase to dehydrogenated BCC phase. The three steps decomposition process confirms the presence of additional to BCC and BCT hydride phase e which is most probably FCC hydride phase. However, additional studies must be done to fully prove the phase transformations during hydrogen sorption and desorption. Kumar et al. reported the same hydride decomposition process for TieCreCeCo alloy. It indicates that described dehydrogenation process can be common for all BCC solid solutions characterized by high hydrogen storage capacity [8]. It can be stated that FCC hydride phase is the most desired for hydrogen storage applications due to the lowest thermal stability and high hydrogen content. For TieVeNi alloys only one clearly visible endothermic peak was observed. It corresponds to hydrogen desorption reaction from BCC hydride to dehydrogenated BCC solid solution. It is important to note that the peak temperature increased from 716.4 K for TieV alloy to 766.1 K for Ni-rich alloy. This very high decomposition temperature proves the high hydride stability presented on desorption part of PCI curves. To consider the practical use of studied alloys the dehydrogenation temperature must be reduced. It was shown in the past that dehydrogenation peak temperature can be lowered by cycled hydrogenation/dehydrogenation process which causes reduction of materials particle size [34]. Moreover, the hydrides decomposition temperature can be also lowered by surface modification of TieVeNi alloy by Raney Ni due to reduction of dehydrogenation process energy barrier [35]. The enthalpy of dehydrogenation process reached 240.3, 154.3, 102.3 and 60.2 J/g respectively for x ¼ 0, 0.1, 0.2 and 0.3. It
Fig. 8 e DSC profile of the hydrides of Ti0.5V1.5¡xNix alloys e endothermic downwards, exothermic upwards (argon flow, 10K/min heating rate).
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Fig. 9 e Discharge capacities as a function of cycle number of electrodes prepared with Ti0.5V1.5¡xNix alloys mixed with carbonyl nickel powder (10 wt%). Electrodes were charged and discharged at 40 mAh/g. The cut off potential was ¡0.7 V.
Table 5 e Activation cycles, maximum discharge capacity and capacity retaining rate after 50th cycle of Ti0.5V1.5¡xNix (40mAh/g charging-discharging current, ¡0.7 V cut off potential). Sample
Ti0.5V1.4Ni0.1 Ti0.5V1.3Ni0.2 Ti0.5V1.2Ni0.3
Activation cycles
Maximum discharge capacity [mAh/g]
Capacity retaining rate after 50th cycle [%]
7 5 8
72 97 116
41 69 75
implies that addition of Ni can be effective in reducing the stability of a phase.
Electrochemical properties Fig. 9 shows results of electrochemical measurements of Tie VeNi alloys. The most important data from these studies were summarized in Table 5. Ti0.5V1.5 alloy is not able to desorb electrochemically hydrogen in KOH electrolyte. However, a great improve of electrochemical properties was obtained after partial substitution of V by Ni atoms. Ti0.5V1.5xNix alloys displayed maximum capacity between 5th and 8th cycle. The tendency of activation properties changes depending on Ni amount in BCC alloys was not noticed. These are much better activation properties than previously published for rapidly solidified BCC-based alloy electrodes [36]. It is most likely due to differences in morphology of materials e small grains of rapidly solidified alloys can suppress pulverization of the electrode materials and thus worsen the activation properties. The maximum discharge capacity increased with increasing Ni content in alloys to reach 116 mAh/g for Ti0.5V1.2Ni0.3 alloy. It is due to the well known catalytic properties of Ni [37]. The same tendency was also observed for vacuum induction melted V2.1TiNix (x ¼ 0.1e0.9) alloys [22]. It can be seen that discharge capacities of TieVeNi alloys were degraded during testing (with increasing charge/ discharge cycles). The reason of degradation can be connected with several factors:
dissolution of V and Ti (mainly V) from the electrode material into the KOH electrolyte [24,38,39]. oxidation of alloy particle surface (formation of TiO2 layer) which causes very low hydrogen diffusivity and electrical conductivity [39]. On the other hand, as stated C. Iwakura et al. the formation of oxide layer suppresses the dissolution of V atoms from MgeVeNi alloys into the alkaline solution [40]. pulverization caused by the cell volume expansion and shrinkage during hydrogenation and dehydrogenation. reduction of electrode mass e a part of material can be flaked away from the electrode during charge/discharge cycles. Despite a decrease in discharge capacity value, which was observed for all TieVeNi alloys, chemical modification of Ti0.5V1.5 by Ni atoms caused increase of cycle stability. Ti0.5V1.2Ni0.3 was characterized by the best cycle stability which reached 75%.
Conclusions The Ti0.5V1.5xNix alloys (x ¼ 0, 0.1, 0.2, 0.3) synthesized by mechanical alloying were studied according to their hydrogen storage and electrochemical properties. XRD measurements showed that all samples are single BCC solid solution phase. Nanocrystallinity of obtained MA alloys was proven by HRTEM studies. XRD and hydrogenation/dehydrogenation measurements showed that hydrogen storage capacities derive directly from chemical compositions and lattice parameters. Maximum hydrogen uptake was reduced with higher Ni content in the TieVeNi alloys. It was due to the reduction of cell volume. On the other hand, Ni addition caused improved hydrogenation kinetics, reduced hysteresis and increased hydrogenation/dehydrogenation reversibility. Electrochemical studies showed that Ni can effectively influence the hydrogen absorption in 6 M KOH solution of TieVeNi alloys. The maximum discharge capacity and capacity retaining rate were significantly improved. Furthermore, it is reasonably to consider the substitution of V by Ni atoms in BCC-based NiMHx to reduce the price of negative electrode e nickel is about ten times cheaper than vanadium is.
Acknowledgement Financial assistance from National Science Centre, Poland (no. 2015/17/N/ST8/00271).
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