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Hydrogen storage properties of (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1M0.1 (M=Ni, Fe, Cu) alloys easily activated at room temperature
Progress in Natural Science: Materials International 27 (2017) 652–657
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Original Research
Hydrogen storage properties of (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1M0.1 (M=Ni, Fe, Cu) alloys easily activated at room temperature Peng Liua, Xiubo Xiea, Li Xub, Xingguo Lic, Tong Liua,
T
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a
Key Laboratory of Aerospace Materials and Performance (Ministry of Education), School of Materials Science and Engineering, Beihang University, No. 37 Xueyuan Road, Beijing 100191, China b State Grid Smart Grid Research Institute, Future Science and Technology City Changping, Beijing 102211, China c Beijing National Laboratory for Molecular Sciences (BNLMS), The State Key Laboratory of Rare Earth Materials Chemistry and Applications, College of Chemistry and Molecular Engineering, Peking University, Beijing 100871, China
A R T I C L E I N F O
A B S T R A C T
Keywords: Titanium-manganese based alloy Laves phase Unit cell volume Activation properties Hydrogen storage properties
The (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1M0.1 (M=Ni, Fe, Cu) alloys with a single C14-type Laves phase have been fabricated by arc melting. They are able to be easily activated by one hydrogen absorption and desorption cycle under 4 MPa hydrogen pressure and vacuum at room temperature. Partial substitution of M for Mn results in the increase of hydrogenation and dehydrogenation capacities in an order of Ni < Fe < Cu. M elements increase the absorption and desorption plateau pressure in an order of (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1 Cu0.1 < (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Fe0.1 < (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Ni0.1. The (Ti0.85Zr0.15)1.05Mn1.2 Cr0.6V0.1Cu0.1 alloy has reversible hydrogen capacities of 1.81 wt% at 273 K and 1.58 wt% at 318 K with formation enthalpy (ΔHab) of −20.66 kJ mol−1 and decomposition enthalpy (ΔHde) of 27.37 kJ mol−1. The differences in the hydrogen storage properties can be attributed to the increase of the interstitial size for hydrogen accommodation caused by the increase of unit cell volumes in the order of (Ti0.85 Zr0.15)1.05Mn1.2Cr0.6V0.1Ni0.1 < (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Fe0.1 < (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Cu0.1.
1. Introduction Hydrogen alloys have been extensively studied in recent years for their versatile applications [1–7]. LaNi5-based AB5-type alloys are the first kind of alloys that have attracted much attention because of their excellent kinetics and activation properties at room temperature, but they are not so promising due to their limited hydrogen storage capacities and high cost [8]. TiFe based AB-type alloys have some advantages of relatively higher hydrogen storage capacities at room temperature and of low cost of raw materials, nevertheless, they are quite difficult to be activated [9–12]. Ti-Mn based AB2-type Laves phase alloys were reported to be an attractive hydrogen storage materials with better activation property than TiFe based alloys, rapid hydridingdehydriding kinetics, high hydrogen storage capacity and low cost [13,14]. Ti-Mn based AB2-type alloys with Laves phase have been used as hydrogen storage materials in the fuel cells of motor vehicles [15–17]. However, the used TiV0.6Fe0.15Mn1.3 and Ti1.1MnCr alloys were required to work under harsh conditions such as high temperature or high pressure. Gamo and coworkers [18] prepared Ti0.9Zr0.1Mn1.4Cr0.4V0.2 alloy which could uptake 1.96 wt% H2 and release 97% of the hydrogen absorbed at room temperature. Liu et al.
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[19] fabricated Ti0.85Zr0.15MnCr0.8V0.1Cu0.1 alloy which had a 1.90 wt % reversible hydrogen capacity at 303 K. However, these TiMn2-type alloys needed to be activated at the temperature over 673 K. In order to improve comprehensive properties of Ti-Mn based AB2type Laves phase alloys, a lot of researches have been conducted in the way of composition design. Osumi et al. [20] studied the Ti1+xMnCr (0.1⩽x⩽0.3) alloys and found the over-stoichiometry of Ti could improve the activation properties of TiMnCr alloys significantly. Wang et al. [21] also found that Ti1+xMn0.8Cr1.2 (x=0.0, 0.1, 0.2, 0.3) alloys had better hydrogen absorption and desorption kinetics and better activation properties with the increase of x. However, these alloys still needed to be activated under the hydrogen pressure above 7 MPa for several hydrogenation and dehydrogenation cycles. Great deals of transition elements such as V, Fe, Co, Ni, Cu, Al, Nb and Mo [13,21–25] have also been added to improve the comprehensive hydrogen storage properties of TiMn2 type alloys. Among these elements, Ni, Fe and Cu were widely studied. Xu et al. [26] found the partial replacement of Mn by Ni was effective in improving the hydrogen absorption/desorption capacities in Ti-Zr-Mn-(VFe)-M (M=Ni, Cr) alloys. Wang et al. [21] found TiCr1.2Mn0.8 alloy cannot be fully activated after 7 hydrogen absorption/desorption cycles but the alloy
Corresponding author. E-mail address: [email protected] (T. Liu).
https://doi.org/10.1016/j.pnsc.2017.09.007 Received 1 April 2017; Received in revised form 15 September 2017; Accepted 18 September 2017 Available online 08 December 2017 1002-0071/ © 2017 Chinese Materials Research Society. Published by Elsevier B.V. This is an open access article under the CC BY-NC-ND license (http://creativecommons.org/licenses/BY-NC-ND/4.0/).
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could be completely activated experienced 3–6 hydrogen absorption/ desorption cycles when Mn was partially substituted by Fe or Ni. However, the activation hydrogen pressure was still as high as 9 MPa. Wang et al. [21] also found the substitution of Mn by Ni would lead to lower hydrogenation plateau pressure and higher hydrogen capacities. On the contrary, Liu et al. [19] found the replacement of Mn by Ni or Fe would induce an increase of the plateau pressure and a decreased hydrogen storage capacity for Ti-Zr-Mn-Cr-V alloy. Park et al. [27] reported that the replacement of Cr by Cu was effective in decreasing the strain energy and slope for (Ti0.75Zr0.25)1.05Mn0.8Cr1.2 alloy but it also reduced the storage capacity—the reversible capacity was 1.35 wt% for (Ti0.75Zr0.25)1.05Mn0.8Cr1.05V0.05Cu0.1 alloy. In comparison, Hong et al. [24] found the substitution of Mn by Cu contributed to higher hydrogen capacity for Ti0.8Zr0.2Mn2 alloy while the hydrogen capacity was still as low as 1.6 wt%. Up to now, the effects of different transition elements (Fe, Ni and Cu) on the hydrogen storage properties of TiMn2 type alloys have not been studied systematically and the corresponding mechanism needs to be clarified. In order to develop TiMn2 type alloys with good activation properties and high hydrogen storage properties, in this work, we intend to prepare a series of over-stoichiometric TiMn2 type alloys, (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1M0.1 (M=Ni, Fe, Cu), investigate the effects of Ni, Fe and Cu substitution on the hydrogen storage properties in detail, and elucidate the corresponding mechanism.
Fig. 1. XRD patterns of the (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Ni0.1 alloy: (a) as prepared; (b) after the hydrogenation and dehydrogenation.
3. Results and discussion 3.1. Characterization Fig. 1 shows the XRD patterns of the as-prepared (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Ni0.1 alloy and the samples obtained after the hydrogen absorption and desorption. It is observed from Fig. 1(a) that all the diffraction peaks of the as-prepared (Ti0.85Zr0.15) 1.05Mn1.2Cr0.6V0.1Ni0.1 alloy, though shifting to the lower degree less than 0.6°, are in good accordance with that of TiMnCr phase (JCPDS No. 50–1114, space group P63/mmc), which has a typical C14 Laves structure. Elliot [28] found that there were the regular relations between the average number of outer electrons (ANOE) and phase structure for AB2 type alloy when A side atom is Ti or Zr and B side atom is V, Cr, Mn, Fe, Co, Cu or Zn. The Laves phase cannot form while the C15 phase exists for Zr atom when ANOE is less than 5.4 and A site is Ti. The C14 phase structure can form for Ti and Zr when ANOE is in the range of 5.4–7. The C15 Laves phase exists for Ti and Zr atom when ANOE is more than 7. In this work, ANOE of (Ti0.85Zr0.15) 1.05Mn1.2Cr0.6V0.1Ni0.1 alloy is calculated to be 5.811, between 5.7 and 5.9. It can therefore be deduced that the sample has C14 Laves phase structure, which is in agreement with the result of XRD analysis. The strongest diffraction peak of standard TiMnCr C14 Laves phase lies in 43.64°, which is 43.13° for the prepared (Ti0.85Zr0.15) 1.05Mn1.2Cr0.6V0.1Ni0.1 alloy (see Fig. 1(a)), showing 0.51° shift to the lower degree. According to the Bragg equation: 2dsinθ=nλ, the shift of diffraction peaks to lower angle can be contributed to the increase of d, the interplanar spacing, which may be caused by the addition of such elements with larger atomic radius as Zr and V. The grain size of the asprepared (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Ni0.1 alloy is calculated to be 87.9 nm by using three strong peaks and taking the average according to the Scherrer equation, see Table S1 in the Supplementary material. After the hydrogenation and dehydrogenation, the (Ti0.85Zr0.15) 1.05Mn1.2Cr0.6V0.1Ni0.1 alloy keeps its phase unchanged while its diffraction peaks broaden with the evaluated value of crystallite particle size falling to 60.8 nm, see Fig. 1(b). It was reported that the TiMn based alloy will chalk into small powders after absorbing hydrogen [18], which may lead to the decrease of crystallite particle size. Fig. 2 represents the XRD patterns of the as-prepared (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Fe0.1 alloy and the samples obtained after the hydrogen absorption and desorption. It can be seen from Fig. 2(a), (b) that the prepared (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Fe0.1 alloy also consists of TiMnCr C14 Laves phase and the phase keeps stable after the hydrogenation and dehydrogenation. There is 0.55° shift to the lower degree of the strongest diffraction peak for the prepared (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Fe0.1 alloy comparing with the standard
2. Experimental 2.1. Preparation of (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1M0.1 (M=Ni, Fe, Cu) alloys The (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1M0.1 (M=Ni, Fe, Cu) alloys were prepared by arc melting of the constituent metals with the purity more than 99.5% on a water-cooled copper crucible under a high purity argon atmosphere (99.999%). An excess amount of Mn of 5 wt% was added considering the weight loss of Mn during melting due to its low melting point. The button was turned over and re-melted four times to ensure high homogeneity. No annealing treatment was employed after melting. The surface oxide layer of the as-cast ingot was polished before being mechanically crushed into powders with the stainless mortar in the air.
2.2. Characterization The pressure-composition-temperature (PCT) tests of the (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1M0.1 (M=Ni, Fe, Cu) alloys were conducted at different temperatures through the conventional pressurevolume-temperature method using a Sieverts-type apparatus. Prior to the measurements, the samples of (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1M0.1 (M=Ni, Fe, Cu) alloys were subjected to one hydrogen absorption/ desorption cycle to be activated. First, the sample holder was vacuumed for 30 min at room temperature. A hydrogen pressure of 4 MPa was induced in the chamber and the alloy was allowed to react until the hydrogenation reached saturation at room temperature. The chamber was then evacuated and the sample was allowed to release hydrogen at vacuum. The hydrogen absorption or desorption measurement for the PCT curves was considered as reaching equilibrium when the vibration of hydrogen pressure at the certain temperature was less than 20 Pa/s. The X-ray diffraction (XRD) patterns of the samples were recorded with a Rigaku X-ray diffractometer with monochromatic Cu Kα radiation. In order to analyze the crystal structure, high-energy X-ray diffractometer was used to collect the XRD data. Lattice constants were calculated by the MAUD software employing Rietveld model.
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Fig. 2. XRD patterns of the (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Fe0.1 alloy: (a) as prepared; (b) after the hydrogenation and dehydrogenation.
Fig. 4. Results of the Rietveld analysis for (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1M0.1 (M=Ni, Fe, Cu) alloys.
with (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Fe0.1, the lattice parameter a and c rise to 4.8954 Å and 8.0353 Å respectively and the unit cell volume rises to 166.771 Å3 when Fe is further substituted by Cu because of the larger atom radius of Cu, 1.28 Å. Additionally, the value of a/c remains almost constant at about 0.6092 for (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1M0.1 (M=Ni, Fe, Cu) alloys, showing their similar structures.
3.2. Hydrogen storage properties Fig. 3. XRD patterns of the (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Cu0.1 alloy: (a) as prepared; (b) after the hydrogenation and dehydrogenation.
Fig. 5(a) shows the absorption and desorption PCT curves for (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Ni0.1 alloy at different temperatures. The samples were activated after one hydrogen absorption/desorption cycle at room temperature and then the PCT tests were conducted. The (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Ni0.1 alloy can uptake 1.79 wt% H2 and release 1.79 wt% H2 at 273 K. Though the hydrogen capacity is lower than Ti0.9Zr0.1Mn1.4V0.2Cr0.4 [18] and Ti0.85Zr0.15MnCr0.8V0.1Cu0.1 alloy [19] which had about 1.90 wt% reversible hydrogen capacity, both the two alloys needed to be activated at the temperature over 673 K. The better activation properties for AB2 type (Ti0.85Zr0.15)1.05 Mn1.2Cr0.6V0.1Ni0.1 alloy may be contributed to the over-stoichiometry of A side atoms, which was in accordance with the Osumi’s report [20] that the over-stoichiometry of Ti could improve the activation properties of TiMnCr alloys significantly. We propose that the better activation properties for the over-stoichiometry alloys are related to their larger numbers of interstitial positions caused by the excessive Ti. The reversible hydrogen capacity of 1.79 wt% is higher than (Ti0.85Zr0.15)1.1CrMn alloy which had a hydrogen capacity of 1.71 wt% [30]. In addition, the (Ti0.85Zr0.15)1.1CrMn alloy needed to be activated under vacuum at 373 K. The improved hydrogen storage performances for (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Ni0.1 alloy can be accounted for the addition of V and Ni that contributes to the higher hydrogen capacity and better activation properties respectively according to Hong’s report [24] and Wang’s studies [21]. The hydrogen absorption and desorption plateau pressure of (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Ni0.1 alloy is 1.41 and 0.39 MPa respectively at 273 K. With the temperature grows from 273 K to 298 K, the hydrogenation and dehydrogenation capacities are reduced to 1.70 wt% and 1.68 wt% and the hydrogen absorption and
TiMnCr C14 Laves phase, higher than that of the prepared (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Ni0.1 alloy, indicating larger lattice parameters according the Bragg equation. The evaluated value of crystallite particle size also drops after the hydrogen absorption and desorption, decreasing from 87.9 to 47.0 nm, see Table S1. XRD patterns show that the prepared (Ti0.85Zr0.15)1.05 Mn1.2Cr0.6V0.1Cu0.1 alloy contains the same phase as the (Ti0.85 Zr0.15)1.05Mn1.2Cr0.6V0.1Ni0.1 and (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Fe0.1 alloys and the phase maintains stable after absorbing and desorbing hydrogen (see Fig. 3(a), (b)). Like the other two alloys, the evaluated value of crystallite particle size of (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Cu0.1 alloy also decreases after the hydrogen absorption and desorption, dropping from 97.2 to 56.0 nm, see Table S1. The diffraction peak at 43.07° for the (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Cu0.1 alloy shifts 0.57° to the low angle compared with the corresponding peak of the standard TiMnCr C14 Laves phase, larger than the (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Ni0.1 and (Ti0.85 Zr0.15)1.05Mn1.2Cr0.6V0.1Fe0.1 alloy, showing improved lattice parameters. Fig. 4 shows the results of the Rietveld analyses for (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1M0.1 (M=Ni, Fe, Cu) alloys. The calculated lattice constants of these alloys refined by the X-ray Rietveld analysis are listed in Table 1. For (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Ni0.1 alloy, the lattice parameter a is 4.8937 Å, c is 8.0324 Å and the unit cell volume is 166.591 Å3 (see Table 1). When Ni is substituted by Fe, the lattice parameter a increases to 4.8940 Å, c increases to 8.0331 Å and the unit cell volume increases to 166.632 Å3, which can be attributed to the larger atomic radius of Fe (1.26 Å) than Ni (1.25 Å) [29]. Compared 654
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Table 1 Crystallographic data of (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1M0.1 alloys. Alloy
a/Å
c/Å
a/c
Vol/Å3
Radius of atom/Å [29]
(Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Ni0.1 (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Fe0.1 (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Cu0.1
4.8937 4.8940 4.8954
8.0324 8.0331 8.0353
0.6092 0.6092 0.6092
166.591 166.632 166.771
RNi=1.25 RFe=1.26 RCu=1.28
Fig. 5. P-C-isotherm curves at 273, 298, 308, 318 K (a) and van’t Hoff plots (b) for (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Ni0.1 alloy.
capacities decrease but the hydrogen absorption and desorption plateau pressures improve. The hydrogenation and dehydrogenation capacities of (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Fe0.1 alloy are 1.72, 1.68, 1.28 wt% and 1.70, 1.68, 1.28 wt% at 298, 308, 318 K respectively, higher than (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Ni0.1. And the absorption and desorption plateau pressures are 2.38, 3.06, 3.46 MPa and 0.98, 1.47, 1.96 MPa at 298, 308, 318 K respectively, lower than (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Ni0.1. Lundin et al. [31] found that an increase in the unit cell volume would lead to an increase in the interstitial size and a decrease in the plateau pressure. In this work, the unit cell volume of (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Fe0.1 alloy is 166.632 Å3, larger than that of (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Ni0.1 alloy of 166.591 Å3, resulting in increased volume of the interspaces for hydrogen accommodation (see Table 1). Therefore, the (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Fe0.1 alloy has higher hydrogen capacities and lower plateau pressures comparing with (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Ni0.1 alloy. ΔHab and ΔHde for the (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Fe0.1 alloy are determined as −18.94 kJ mol−1 and 27.04 kJ mol−1 respectively according to the van't Hoff plots in Fig. 6(b), higher than that of (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Ni0.1 alloy, indicating an increased
desorption plateau pressures are increased to 2.65 MPa and 1.00 MPa respectively. The (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Ni0.1 alloy can uptake 1.66 wt% H2 and release 1.66 wt% H2 at 308 K with a 3.39 MPa absorption plateau pressure and a 1.49 MPa desorption plateau pressure. The hydrogenation and dehydrogenation capacities decline to 0.94 wt% and 0.92 wt% and the absorption and desorption plateau pressures rise to 3.78 MPa and 1.98 MPa respectively when temperature ascends to 318 K. Fig. 5(b) shows van't Hoff plots of (Ti0.85 Zr0.15)1.05Mn1.2Cr0.6V0.1Ni0.1 alloy based on the PCT curves. ΔHab for the (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Ni0.1 alloy is determined as −16.23 kJ mol−1. ΔHde for the (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Ni0.1 alloy is evaluated to be 26.46 kJ mol−1. Fig. 6(a) shows the absorption and desorption PCT curves for (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Fe0.1 alloy at different temperatures. The PCT tests were conducted after one hydrogen absorption/desorption cycle at room temperature. The (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Fe0.1 alloy can uptake 1.80 wt% H2 and release 1.80 wt% H2 at 273 K, a little higher than (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Ni0.1. The hydrogen absorption and desorption plateau pressure is 1.10 and 0.37 MPa respectively at 273 K, lower than (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Ni0.1. With the temperature growing, the hydrogenation and dehydrogenation
Fig. 6. P-C-isotherm curves at 273, 298, 308, 318 K (a) and van’t Hoff plots (b) for (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Fe0.1 alloy.
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Fig. 7. P-C-isotherm curves at 273, 298, 308, 318 K (a) and van’t Hoff plots (b) for (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Cu0.1 alloy.
desorb 1.81 wt% H2 at 273 K and 1.58 wt% H2 at 318 K respectively. The ΔHab and ΔHde of the (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Cu0.1 alloy is −20.66 kJ mol−1 and 27.37 kJ mol−1 respectively. Larger atom radiuses of Cu and Fe lead to larger cell volumes and lattice interspaces, which results in a higher hydrogen capacity and lower plateau pressures.
stability for these alloy hydrides, which is consistent with the decreased plateau pressures observed in PCT curves. Cao et al. [30] also found that the shrink of the cell volume induced by the substitution of smaller atoms tended to diminish the interspace for hydrogen, resulting in the decreased hydrogen capacity, increased dissociation plateau pressure and decreased ΔHde. Fig. 7(a) exhibits the absorption and desorption PCT curves for (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Cu0.1 alloy at different temperatures. After one hydrogen absorption/desorption cycle at room temperature the PCT tests were conducted. The (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Cu0.1 alloy can uptake 1.81 wt% H2 and release 1.81 wt% H2 at 273 K, higher than (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Ni0.1 and (Ti0.85Zr0.15)1.05 Mn1.2Cr0.6V0.1Fe0.1. The hydrogen absorption and desorption plateau pressure is 0.77 and 0.30 MPa respectively at 273 K, lower than (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Ni0.1 and (Ti0.85Zr0.15)1.05Mn1.2Cr0.6 V0.1Fe0.1. Like the other two alloys, with the temperature rising, the hydrogenation and dehydrogenation capacities fall but the hydrogen absorption and desorption plateau pressures increase. The (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Cu0.1 alloy is able to absorb 1.75, 1.70, 1.58 wt% H2 and desorb 1.75, 1.70, 1.58 wt% H2 at 298, 308 and 318 K respectively, higher than (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Ni0.1 and (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Fe0.1. The absorption and desorption plateau pressures are 1.62, 2.14, 2.61, and 0.85, 1.24, 1.65 MPa at 298, 308, 318 K, respectively, lower than (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Ni0.1 and (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Fe0.1. The higher hydrogen capacities and lower plateau pressures are related to a larger volume of the interspaces resulting from the substitution of Cu with larger atom radius than Fe and Ni (see Table 1). As it shown in Fig. 7(b), the ΔHab and ΔHde for the (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Cu0.1 alloy increase to −20.66 kJ mol−1 and 27.37 kJ mol−1 respectively, higher than that of (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Ni0.1 and (Ti0.85Zr0.15)1.05Mn1.2Cr0.6 V0.1Fe0.1, which is in good accordance with the decreased plateau pressure observed in PCT curves.
Acknowledgements The authors acknowledge the support of this work by the Fund of State Grid Smart Grid Research Institute (No. SGRIDGJSKF[2016]604), the Joint Fund of the National Natural Science Foundation of China and Baosteel Group Corporation (No. U1560106), the Aeronautical Science Foundation of China (No. 2016ZF51050) and the Scientific Research Foundation for the Returned Overseas Chinese Scholars (State Education Ministry). Appendix A. Supplementary material Supplementary data associated with this article can be found in the online version at http://dx.doi.org/10.1016/j.pnsc.2017.09.007. References [1] H.Y. Shao, T. Liu, X.G. Li, Nanotechnology 14 (2003) L1–L3. [2] T. Liu, H.L. Shen, Y. Liu, L. Xie, J.L. Qu, H.Y. Shao, X.G. Li, J. Power Sources 227 (2013) 86–93. [3] X.J. Ma, X.B. Xie, P. Liu, L. Xu, T. Liu, Prog. Nat. Sci. Mater. 27 (2017) 99–104. [4] J.G. Yuan, W. Li, Y. Wu, Prog. Nat. Sci. Mater. 27 (2017) 169–176. [5] V.Y. Zadorozhnyy, G.S. Milovzorov, S.N. Klyamkin, M.Y. Zadorozhnyy, D.V. Strugova, M.V. Gorshenkov, S.D. Kaloshkin, Prog. Nat. Sci. Mater. 27 (2017) 149–155. [6] T. Liu, C.G. Chen, C.G. Qin, X.G. Li, Int. J. Hydrog. Energy 39 (2014) 18273–18279. [7] T. Liu, C.G. Chen, F. Wang, X.G. Li, J. Power Sources 267 (2014) 69–77. [8] K.H. Buschow, H.H.V. Mal, J. Less Common Met. 29 (1972) 203–210. [9] K. Edalati, J. Matsuda, A. Yanagida, E. Akiba, Z. Horita, Int. J. Hydrog. Energy 39 (2014) 15589–15594. [10] J.J. Reilly, R.H. Wiswall, Inorg. Chem. 13 (1974) 77–112. [11] V.Y. Zadorozhnyy, S.N. Klyamkin, M.Y. Zadorozhnyy, O.V. Bermesheva, S.D. Kaloshkin, J. Alloy. Compd. 586 (Supplement 1) (2014) S56–S60. [12] A.V. Bakulin, S.S. Kulkov, S.E. Kulkova, S. Hocker, Int. J. Hydrog. Energy 39 (2014) 12213–12220. [13] T. Huang, Z. Wu, G. Sun, N. Xu, Intermetallics 15 (2007) 593–598. [14] S. Semboshi, N. Masahashi, S. Hanada, J. Alloy. Compd. 352 (2003) 210–217. [15] K. Binder, G. Withalm, in: Proceedings of the 3rd World Hydrogen Energy Conference, Tokyo, Jun, 1980, pp. 1103–1117. [16] J. Töpler, O. Bernauer, H. Buchner, H. Säufferer, J. Less Common Met. 89 (1983) 519–526. [17] Y. Kojima, Y. Kawai, S.I. Towata, T. Matsunaga, T. Shinozawa, M. Kimbara, J. Alloy. Compd. 419 (2006) 256–261. [18] T. Gamo, Y. Moriwaki, N. Yanagihara, T. Yamashita, T. Iwaki, Int. J. Hydrog. Energy 10 (1985) 39–47. [19] B.H. Liu, D.M. Kim, K.Y. Lee, J.Y. Lee, J. Alloy. Compd. 240 (1996) 214–218. [20] Y. Osumi, H. Suzuki, A. Kato, K. Oguro, T. Sugioka, T. Fujita, J. Less Common Met. 89 (1983) 257–262.
4. Conclusions (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1M0.1 (M=Ni, Fe, Cu) alloys composed of a single C14-type Laves phase have been prepared. After hydrogen absorption and desorption, all the alloys keep their phase stable but the grain sizes decrease. All the alloys can be easily activated by one hydrogen absorption and desorption cycle at room temperature under 4 MPa hydrogen pressure and vacuum. The hydrogenation and dehydrogenation capacities and ΔHab and ΔHde increase in an order of (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Ni0.1 < (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1 Fe0.1 < (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Cu0.1. Partial substitution of M for Mn contributes to the increase of plateau pressure of (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1M0.1 (M=Ni, Fe, Cu) alloy in an order of Cu < Fe < Ni. The (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Cu0.1 can absorb and 656
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351–355. [28] R.P. Eliiot, W. Rostoker, Trans. Am. Soc. Met. 50 (1958) 617–632. [29] N.N. Greenwood, A. Earnshaw, Chemistry of the Elements, 2nd ed, ButterworthHeinemann, Oxford, 1997. [30] Z. Cao, L. Ouyang, H. Wang, J. Liu, L. Sun, M. Zhu, J. Alloy. Compd. 639 (2015) 452–457. [31] C.E. Lundin, F.E. Lynch, C.B. Magee, J. Less Common Met. 56 (1977) 19–37.
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Original Research
Hydrogen storage properties of (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1M0.1 (M=Ni, Fe, Cu) alloys easily activated at room temperature Peng Liua, Xiubo Xiea, Li Xub, Xingguo Lic, Tong Liua,
T
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a
Key Laboratory of Aerospace Materials and Performance (Ministry of Education), School of Materials Science and Engineering, Beihang University, No. 37 Xueyuan Road, Beijing 100191, China b State Grid Smart Grid Research Institute, Future Science and Technology City Changping, Beijing 102211, China c Beijing National Laboratory for Molecular Sciences (BNLMS), The State Key Laboratory of Rare Earth Materials Chemistry and Applications, College of Chemistry and Molecular Engineering, Peking University, Beijing 100871, China
A R T I C L E I N F O
A B S T R A C T
Keywords: Titanium-manganese based alloy Laves phase Unit cell volume Activation properties Hydrogen storage properties
The (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1M0.1 (M=Ni, Fe, Cu) alloys with a single C14-type Laves phase have been fabricated by arc melting. They are able to be easily activated by one hydrogen absorption and desorption cycle under 4 MPa hydrogen pressure and vacuum at room temperature. Partial substitution of M for Mn results in the increase of hydrogenation and dehydrogenation capacities in an order of Ni < Fe < Cu. M elements increase the absorption and desorption plateau pressure in an order of (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1 Cu0.1 < (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Fe0.1 < (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Ni0.1. The (Ti0.85Zr0.15)1.05Mn1.2 Cr0.6V0.1Cu0.1 alloy has reversible hydrogen capacities of 1.81 wt% at 273 K and 1.58 wt% at 318 K with formation enthalpy (ΔHab) of −20.66 kJ mol−1 and decomposition enthalpy (ΔHde) of 27.37 kJ mol−1. The differences in the hydrogen storage properties can be attributed to the increase of the interstitial size for hydrogen accommodation caused by the increase of unit cell volumes in the order of (Ti0.85 Zr0.15)1.05Mn1.2Cr0.6V0.1Ni0.1 < (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Fe0.1 < (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Cu0.1.
1. Introduction Hydrogen alloys have been extensively studied in recent years for their versatile applications [1–7]. LaNi5-based AB5-type alloys are the first kind of alloys that have attracted much attention because of their excellent kinetics and activation properties at room temperature, but they are not so promising due to their limited hydrogen storage capacities and high cost [8]. TiFe based AB-type alloys have some advantages of relatively higher hydrogen storage capacities at room temperature and of low cost of raw materials, nevertheless, they are quite difficult to be activated [9–12]. Ti-Mn based AB2-type Laves phase alloys were reported to be an attractive hydrogen storage materials with better activation property than TiFe based alloys, rapid hydridingdehydriding kinetics, high hydrogen storage capacity and low cost [13,14]. Ti-Mn based AB2-type alloys with Laves phase have been used as hydrogen storage materials in the fuel cells of motor vehicles [15–17]. However, the used TiV0.6Fe0.15Mn1.3 and Ti1.1MnCr alloys were required to work under harsh conditions such as high temperature or high pressure. Gamo and coworkers [18] prepared Ti0.9Zr0.1Mn1.4Cr0.4V0.2 alloy which could uptake 1.96 wt% H2 and release 97% of the hydrogen absorbed at room temperature. Liu et al.
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[19] fabricated Ti0.85Zr0.15MnCr0.8V0.1Cu0.1 alloy which had a 1.90 wt % reversible hydrogen capacity at 303 K. However, these TiMn2-type alloys needed to be activated at the temperature over 673 K. In order to improve comprehensive properties of Ti-Mn based AB2type Laves phase alloys, a lot of researches have been conducted in the way of composition design. Osumi et al. [20] studied the Ti1+xMnCr (0.1⩽x⩽0.3) alloys and found the over-stoichiometry of Ti could improve the activation properties of TiMnCr alloys significantly. Wang et al. [21] also found that Ti1+xMn0.8Cr1.2 (x=0.0, 0.1, 0.2, 0.3) alloys had better hydrogen absorption and desorption kinetics and better activation properties with the increase of x. However, these alloys still needed to be activated under the hydrogen pressure above 7 MPa for several hydrogenation and dehydrogenation cycles. Great deals of transition elements such as V, Fe, Co, Ni, Cu, Al, Nb and Mo [13,21–25] have also been added to improve the comprehensive hydrogen storage properties of TiMn2 type alloys. Among these elements, Ni, Fe and Cu were widely studied. Xu et al. [26] found the partial replacement of Mn by Ni was effective in improving the hydrogen absorption/desorption capacities in Ti-Zr-Mn-(VFe)-M (M=Ni, Cr) alloys. Wang et al. [21] found TiCr1.2Mn0.8 alloy cannot be fully activated after 7 hydrogen absorption/desorption cycles but the alloy
Corresponding author. E-mail address: [email protected] (T. Liu).
https://doi.org/10.1016/j.pnsc.2017.09.007 Received 1 April 2017; Received in revised form 15 September 2017; Accepted 18 September 2017 Available online 08 December 2017 1002-0071/ © 2017 Chinese Materials Research Society. Published by Elsevier B.V. This is an open access article under the CC BY-NC-ND license (http://creativecommons.org/licenses/BY-NC-ND/4.0/).
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could be completely activated experienced 3–6 hydrogen absorption/ desorption cycles when Mn was partially substituted by Fe or Ni. However, the activation hydrogen pressure was still as high as 9 MPa. Wang et al. [21] also found the substitution of Mn by Ni would lead to lower hydrogenation plateau pressure and higher hydrogen capacities. On the contrary, Liu et al. [19] found the replacement of Mn by Ni or Fe would induce an increase of the plateau pressure and a decreased hydrogen storage capacity for Ti-Zr-Mn-Cr-V alloy. Park et al. [27] reported that the replacement of Cr by Cu was effective in decreasing the strain energy and slope for (Ti0.75Zr0.25)1.05Mn0.8Cr1.2 alloy but it also reduced the storage capacity—the reversible capacity was 1.35 wt% for (Ti0.75Zr0.25)1.05Mn0.8Cr1.05V0.05Cu0.1 alloy. In comparison, Hong et al. [24] found the substitution of Mn by Cu contributed to higher hydrogen capacity for Ti0.8Zr0.2Mn2 alloy while the hydrogen capacity was still as low as 1.6 wt%. Up to now, the effects of different transition elements (Fe, Ni and Cu) on the hydrogen storage properties of TiMn2 type alloys have not been studied systematically and the corresponding mechanism needs to be clarified. In order to develop TiMn2 type alloys with good activation properties and high hydrogen storage properties, in this work, we intend to prepare a series of over-stoichiometric TiMn2 type alloys, (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1M0.1 (M=Ni, Fe, Cu), investigate the effects of Ni, Fe and Cu substitution on the hydrogen storage properties in detail, and elucidate the corresponding mechanism.
Fig. 1. XRD patterns of the (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Ni0.1 alloy: (a) as prepared; (b) after the hydrogenation and dehydrogenation.
3. Results and discussion 3.1. Characterization Fig. 1 shows the XRD patterns of the as-prepared (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Ni0.1 alloy and the samples obtained after the hydrogen absorption and desorption. It is observed from Fig. 1(a) that all the diffraction peaks of the as-prepared (Ti0.85Zr0.15) 1.05Mn1.2Cr0.6V0.1Ni0.1 alloy, though shifting to the lower degree less than 0.6°, are in good accordance with that of TiMnCr phase (JCPDS No. 50–1114, space group P63/mmc), which has a typical C14 Laves structure. Elliot [28] found that there were the regular relations between the average number of outer electrons (ANOE) and phase structure for AB2 type alloy when A side atom is Ti or Zr and B side atom is V, Cr, Mn, Fe, Co, Cu or Zn. The Laves phase cannot form while the C15 phase exists for Zr atom when ANOE is less than 5.4 and A site is Ti. The C14 phase structure can form for Ti and Zr when ANOE is in the range of 5.4–7. The C15 Laves phase exists for Ti and Zr atom when ANOE is more than 7. In this work, ANOE of (Ti0.85Zr0.15) 1.05Mn1.2Cr0.6V0.1Ni0.1 alloy is calculated to be 5.811, between 5.7 and 5.9. It can therefore be deduced that the sample has C14 Laves phase structure, which is in agreement with the result of XRD analysis. The strongest diffraction peak of standard TiMnCr C14 Laves phase lies in 43.64°, which is 43.13° for the prepared (Ti0.85Zr0.15) 1.05Mn1.2Cr0.6V0.1Ni0.1 alloy (see Fig. 1(a)), showing 0.51° shift to the lower degree. According to the Bragg equation: 2dsinθ=nλ, the shift of diffraction peaks to lower angle can be contributed to the increase of d, the interplanar spacing, which may be caused by the addition of such elements with larger atomic radius as Zr and V. The grain size of the asprepared (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Ni0.1 alloy is calculated to be 87.9 nm by using three strong peaks and taking the average according to the Scherrer equation, see Table S1 in the Supplementary material. After the hydrogenation and dehydrogenation, the (Ti0.85Zr0.15) 1.05Mn1.2Cr0.6V0.1Ni0.1 alloy keeps its phase unchanged while its diffraction peaks broaden with the evaluated value of crystallite particle size falling to 60.8 nm, see Fig. 1(b). It was reported that the TiMn based alloy will chalk into small powders after absorbing hydrogen [18], which may lead to the decrease of crystallite particle size. Fig. 2 represents the XRD patterns of the as-prepared (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Fe0.1 alloy and the samples obtained after the hydrogen absorption and desorption. It can be seen from Fig. 2(a), (b) that the prepared (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Fe0.1 alloy also consists of TiMnCr C14 Laves phase and the phase keeps stable after the hydrogenation and dehydrogenation. There is 0.55° shift to the lower degree of the strongest diffraction peak for the prepared (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Fe0.1 alloy comparing with the standard
2. Experimental 2.1. Preparation of (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1M0.1 (M=Ni, Fe, Cu) alloys The (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1M0.1 (M=Ni, Fe, Cu) alloys were prepared by arc melting of the constituent metals with the purity more than 99.5% on a water-cooled copper crucible under a high purity argon atmosphere (99.999%). An excess amount of Mn of 5 wt% was added considering the weight loss of Mn during melting due to its low melting point. The button was turned over and re-melted four times to ensure high homogeneity. No annealing treatment was employed after melting. The surface oxide layer of the as-cast ingot was polished before being mechanically crushed into powders with the stainless mortar in the air.
2.2. Characterization The pressure-composition-temperature (PCT) tests of the (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1M0.1 (M=Ni, Fe, Cu) alloys were conducted at different temperatures through the conventional pressurevolume-temperature method using a Sieverts-type apparatus. Prior to the measurements, the samples of (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1M0.1 (M=Ni, Fe, Cu) alloys were subjected to one hydrogen absorption/ desorption cycle to be activated. First, the sample holder was vacuumed for 30 min at room temperature. A hydrogen pressure of 4 MPa was induced in the chamber and the alloy was allowed to react until the hydrogenation reached saturation at room temperature. The chamber was then evacuated and the sample was allowed to release hydrogen at vacuum. The hydrogen absorption or desorption measurement for the PCT curves was considered as reaching equilibrium when the vibration of hydrogen pressure at the certain temperature was less than 20 Pa/s. The X-ray diffraction (XRD) patterns of the samples were recorded with a Rigaku X-ray diffractometer with monochromatic Cu Kα radiation. In order to analyze the crystal structure, high-energy X-ray diffractometer was used to collect the XRD data. Lattice constants were calculated by the MAUD software employing Rietveld model.
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Fig. 2. XRD patterns of the (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Fe0.1 alloy: (a) as prepared; (b) after the hydrogenation and dehydrogenation.
Fig. 4. Results of the Rietveld analysis for (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1M0.1 (M=Ni, Fe, Cu) alloys.
with (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Fe0.1, the lattice parameter a and c rise to 4.8954 Å and 8.0353 Å respectively and the unit cell volume rises to 166.771 Å3 when Fe is further substituted by Cu because of the larger atom radius of Cu, 1.28 Å. Additionally, the value of a/c remains almost constant at about 0.6092 for (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1M0.1 (M=Ni, Fe, Cu) alloys, showing their similar structures.
3.2. Hydrogen storage properties Fig. 3. XRD patterns of the (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Cu0.1 alloy: (a) as prepared; (b) after the hydrogenation and dehydrogenation.
Fig. 5(a) shows the absorption and desorption PCT curves for (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Ni0.1 alloy at different temperatures. The samples were activated after one hydrogen absorption/desorption cycle at room temperature and then the PCT tests were conducted. The (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Ni0.1 alloy can uptake 1.79 wt% H2 and release 1.79 wt% H2 at 273 K. Though the hydrogen capacity is lower than Ti0.9Zr0.1Mn1.4V0.2Cr0.4 [18] and Ti0.85Zr0.15MnCr0.8V0.1Cu0.1 alloy [19] which had about 1.90 wt% reversible hydrogen capacity, both the two alloys needed to be activated at the temperature over 673 K. The better activation properties for AB2 type (Ti0.85Zr0.15)1.05 Mn1.2Cr0.6V0.1Ni0.1 alloy may be contributed to the over-stoichiometry of A side atoms, which was in accordance with the Osumi’s report [20] that the over-stoichiometry of Ti could improve the activation properties of TiMnCr alloys significantly. We propose that the better activation properties for the over-stoichiometry alloys are related to their larger numbers of interstitial positions caused by the excessive Ti. The reversible hydrogen capacity of 1.79 wt% is higher than (Ti0.85Zr0.15)1.1CrMn alloy which had a hydrogen capacity of 1.71 wt% [30]. In addition, the (Ti0.85Zr0.15)1.1CrMn alloy needed to be activated under vacuum at 373 K. The improved hydrogen storage performances for (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Ni0.1 alloy can be accounted for the addition of V and Ni that contributes to the higher hydrogen capacity and better activation properties respectively according to Hong’s report [24] and Wang’s studies [21]. The hydrogen absorption and desorption plateau pressure of (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Ni0.1 alloy is 1.41 and 0.39 MPa respectively at 273 K. With the temperature grows from 273 K to 298 K, the hydrogenation and dehydrogenation capacities are reduced to 1.70 wt% and 1.68 wt% and the hydrogen absorption and
TiMnCr C14 Laves phase, higher than that of the prepared (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Ni0.1 alloy, indicating larger lattice parameters according the Bragg equation. The evaluated value of crystallite particle size also drops after the hydrogen absorption and desorption, decreasing from 87.9 to 47.0 nm, see Table S1. XRD patterns show that the prepared (Ti0.85Zr0.15)1.05 Mn1.2Cr0.6V0.1Cu0.1 alloy contains the same phase as the (Ti0.85 Zr0.15)1.05Mn1.2Cr0.6V0.1Ni0.1 and (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Fe0.1 alloys and the phase maintains stable after absorbing and desorbing hydrogen (see Fig. 3(a), (b)). Like the other two alloys, the evaluated value of crystallite particle size of (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Cu0.1 alloy also decreases after the hydrogen absorption and desorption, dropping from 97.2 to 56.0 nm, see Table S1. The diffraction peak at 43.07° for the (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Cu0.1 alloy shifts 0.57° to the low angle compared with the corresponding peak of the standard TiMnCr C14 Laves phase, larger than the (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Ni0.1 and (Ti0.85 Zr0.15)1.05Mn1.2Cr0.6V0.1Fe0.1 alloy, showing improved lattice parameters. Fig. 4 shows the results of the Rietveld analyses for (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1M0.1 (M=Ni, Fe, Cu) alloys. The calculated lattice constants of these alloys refined by the X-ray Rietveld analysis are listed in Table 1. For (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Ni0.1 alloy, the lattice parameter a is 4.8937 Å, c is 8.0324 Å and the unit cell volume is 166.591 Å3 (see Table 1). When Ni is substituted by Fe, the lattice parameter a increases to 4.8940 Å, c increases to 8.0331 Å and the unit cell volume increases to 166.632 Å3, which can be attributed to the larger atomic radius of Fe (1.26 Å) than Ni (1.25 Å) [29]. Compared 654
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Table 1 Crystallographic data of (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1M0.1 alloys. Alloy
a/Å
c/Å
a/c
Vol/Å3
Radius of atom/Å [29]
(Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Ni0.1 (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Fe0.1 (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Cu0.1
4.8937 4.8940 4.8954
8.0324 8.0331 8.0353
0.6092 0.6092 0.6092
166.591 166.632 166.771
RNi=1.25 RFe=1.26 RCu=1.28
Fig. 5. P-C-isotherm curves at 273, 298, 308, 318 K (a) and van’t Hoff plots (b) for (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Ni0.1 alloy.
capacities decrease but the hydrogen absorption and desorption plateau pressures improve. The hydrogenation and dehydrogenation capacities of (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Fe0.1 alloy are 1.72, 1.68, 1.28 wt% and 1.70, 1.68, 1.28 wt% at 298, 308, 318 K respectively, higher than (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Ni0.1. And the absorption and desorption plateau pressures are 2.38, 3.06, 3.46 MPa and 0.98, 1.47, 1.96 MPa at 298, 308, 318 K respectively, lower than (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Ni0.1. Lundin et al. [31] found that an increase in the unit cell volume would lead to an increase in the interstitial size and a decrease in the plateau pressure. In this work, the unit cell volume of (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Fe0.1 alloy is 166.632 Å3, larger than that of (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Ni0.1 alloy of 166.591 Å3, resulting in increased volume of the interspaces for hydrogen accommodation (see Table 1). Therefore, the (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Fe0.1 alloy has higher hydrogen capacities and lower plateau pressures comparing with (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Ni0.1 alloy. ΔHab and ΔHde for the (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Fe0.1 alloy are determined as −18.94 kJ mol−1 and 27.04 kJ mol−1 respectively according to the van't Hoff plots in Fig. 6(b), higher than that of (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Ni0.1 alloy, indicating an increased
desorption plateau pressures are increased to 2.65 MPa and 1.00 MPa respectively. The (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Ni0.1 alloy can uptake 1.66 wt% H2 and release 1.66 wt% H2 at 308 K with a 3.39 MPa absorption plateau pressure and a 1.49 MPa desorption plateau pressure. The hydrogenation and dehydrogenation capacities decline to 0.94 wt% and 0.92 wt% and the absorption and desorption plateau pressures rise to 3.78 MPa and 1.98 MPa respectively when temperature ascends to 318 K. Fig. 5(b) shows van't Hoff plots of (Ti0.85 Zr0.15)1.05Mn1.2Cr0.6V0.1Ni0.1 alloy based on the PCT curves. ΔHab for the (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Ni0.1 alloy is determined as −16.23 kJ mol−1. ΔHde for the (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Ni0.1 alloy is evaluated to be 26.46 kJ mol−1. Fig. 6(a) shows the absorption and desorption PCT curves for (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Fe0.1 alloy at different temperatures. The PCT tests were conducted after one hydrogen absorption/desorption cycle at room temperature. The (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Fe0.1 alloy can uptake 1.80 wt% H2 and release 1.80 wt% H2 at 273 K, a little higher than (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Ni0.1. The hydrogen absorption and desorption plateau pressure is 1.10 and 0.37 MPa respectively at 273 K, lower than (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Ni0.1. With the temperature growing, the hydrogenation and dehydrogenation
Fig. 6. P-C-isotherm curves at 273, 298, 308, 318 K (a) and van’t Hoff plots (b) for (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Fe0.1 alloy.
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Fig. 7. P-C-isotherm curves at 273, 298, 308, 318 K (a) and van’t Hoff plots (b) for (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Cu0.1 alloy.
desorb 1.81 wt% H2 at 273 K and 1.58 wt% H2 at 318 K respectively. The ΔHab and ΔHde of the (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Cu0.1 alloy is −20.66 kJ mol−1 and 27.37 kJ mol−1 respectively. Larger atom radiuses of Cu and Fe lead to larger cell volumes and lattice interspaces, which results in a higher hydrogen capacity and lower plateau pressures.
stability for these alloy hydrides, which is consistent with the decreased plateau pressures observed in PCT curves. Cao et al. [30] also found that the shrink of the cell volume induced by the substitution of smaller atoms tended to diminish the interspace for hydrogen, resulting in the decreased hydrogen capacity, increased dissociation plateau pressure and decreased ΔHde. Fig. 7(a) exhibits the absorption and desorption PCT curves for (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Cu0.1 alloy at different temperatures. After one hydrogen absorption/desorption cycle at room temperature the PCT tests were conducted. The (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Cu0.1 alloy can uptake 1.81 wt% H2 and release 1.81 wt% H2 at 273 K, higher than (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Ni0.1 and (Ti0.85Zr0.15)1.05 Mn1.2Cr0.6V0.1Fe0.1. The hydrogen absorption and desorption plateau pressure is 0.77 and 0.30 MPa respectively at 273 K, lower than (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Ni0.1 and (Ti0.85Zr0.15)1.05Mn1.2Cr0.6 V0.1Fe0.1. Like the other two alloys, with the temperature rising, the hydrogenation and dehydrogenation capacities fall but the hydrogen absorption and desorption plateau pressures increase. The (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Cu0.1 alloy is able to absorb 1.75, 1.70, 1.58 wt% H2 and desorb 1.75, 1.70, 1.58 wt% H2 at 298, 308 and 318 K respectively, higher than (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Ni0.1 and (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Fe0.1. The absorption and desorption plateau pressures are 1.62, 2.14, 2.61, and 0.85, 1.24, 1.65 MPa at 298, 308, 318 K, respectively, lower than (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Ni0.1 and (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Fe0.1. The higher hydrogen capacities and lower plateau pressures are related to a larger volume of the interspaces resulting from the substitution of Cu with larger atom radius than Fe and Ni (see Table 1). As it shown in Fig. 7(b), the ΔHab and ΔHde for the (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Cu0.1 alloy increase to −20.66 kJ mol−1 and 27.37 kJ mol−1 respectively, higher than that of (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Ni0.1 and (Ti0.85Zr0.15)1.05Mn1.2Cr0.6 V0.1Fe0.1, which is in good accordance with the decreased plateau pressure observed in PCT curves.
Acknowledgements The authors acknowledge the support of this work by the Fund of State Grid Smart Grid Research Institute (No. SGRIDGJSKF[2016]604), the Joint Fund of the National Natural Science Foundation of China and Baosteel Group Corporation (No. U1560106), the Aeronautical Science Foundation of China (No. 2016ZF51050) and the Scientific Research Foundation for the Returned Overseas Chinese Scholars (State Education Ministry). Appendix A. Supplementary material Supplementary data associated with this article can be found in the online version at http://dx.doi.org/10.1016/j.pnsc.2017.09.007. References [1] H.Y. Shao, T. Liu, X.G. Li, Nanotechnology 14 (2003) L1–L3. [2] T. Liu, H.L. Shen, Y. Liu, L. Xie, J.L. Qu, H.Y. Shao, X.G. Li, J. Power Sources 227 (2013) 86–93. [3] X.J. Ma, X.B. Xie, P. Liu, L. Xu, T. Liu, Prog. Nat. Sci. Mater. 27 (2017) 99–104. [4] J.G. Yuan, W. Li, Y. Wu, Prog. Nat. Sci. Mater. 27 (2017) 169–176. [5] V.Y. Zadorozhnyy, G.S. Milovzorov, S.N. Klyamkin, M.Y. Zadorozhnyy, D.V. Strugova, M.V. Gorshenkov, S.D. Kaloshkin, Prog. Nat. Sci. Mater. 27 (2017) 149–155. [6] T. Liu, C.G. Chen, C.G. Qin, X.G. Li, Int. J. Hydrog. Energy 39 (2014) 18273–18279. [7] T. Liu, C.G. Chen, F. Wang, X.G. Li, J. Power Sources 267 (2014) 69–77. [8] K.H. Buschow, H.H.V. Mal, J. Less Common Met. 29 (1972) 203–210. [9] K. Edalati, J. Matsuda, A. Yanagida, E. Akiba, Z. Horita, Int. J. Hydrog. Energy 39 (2014) 15589–15594. [10] J.J. Reilly, R.H. Wiswall, Inorg. Chem. 13 (1974) 77–112. [11] V.Y. Zadorozhnyy, S.N. Klyamkin, M.Y. Zadorozhnyy, O.V. Bermesheva, S.D. Kaloshkin, J. Alloy. Compd. 586 (Supplement 1) (2014) S56–S60. [12] A.V. Bakulin, S.S. Kulkov, S.E. Kulkova, S. Hocker, Int. J. Hydrog. Energy 39 (2014) 12213–12220. [13] T. Huang, Z. Wu, G. Sun, N. Xu, Intermetallics 15 (2007) 593–598. [14] S. Semboshi, N. Masahashi, S. Hanada, J. Alloy. Compd. 352 (2003) 210–217. [15] K. Binder, G. Withalm, in: Proceedings of the 3rd World Hydrogen Energy Conference, Tokyo, Jun, 1980, pp. 1103–1117. [16] J. Töpler, O. Bernauer, H. Buchner, H. Säufferer, J. Less Common Met. 89 (1983) 519–526. [17] Y. Kojima, Y. Kawai, S.I. Towata, T. Matsunaga, T. Shinozawa, M. Kimbara, J. Alloy. Compd. 419 (2006) 256–261. [18] T. Gamo, Y. Moriwaki, N. Yanagihara, T. Yamashita, T. Iwaki, Int. J. Hydrog. Energy 10 (1985) 39–47. [19] B.H. Liu, D.M. Kim, K.Y. Lee, J.Y. Lee, J. Alloy. Compd. 240 (1996) 214–218. [20] Y. Osumi, H. Suzuki, A. Kato, K. Oguro, T. Sugioka, T. Fujita, J. Less Common Met. 89 (1983) 257–262.
4. Conclusions (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1M0.1 (M=Ni, Fe, Cu) alloys composed of a single C14-type Laves phase have been prepared. After hydrogen absorption and desorption, all the alloys keep their phase stable but the grain sizes decrease. All the alloys can be easily activated by one hydrogen absorption and desorption cycle at room temperature under 4 MPa hydrogen pressure and vacuum. The hydrogenation and dehydrogenation capacities and ΔHab and ΔHde increase in an order of (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Ni0.1 < (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1 Fe0.1 < (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Cu0.1. Partial substitution of M for Mn contributes to the increase of plateau pressure of (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1M0.1 (M=Ni, Fe, Cu) alloy in an order of Cu < Fe < Ni. The (Ti0.85Zr0.15)1.05Mn1.2Cr0.6V0.1Cu0.1 can absorb and 656
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