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A new reversible Mg3Ag–H2 system for hydrogen storage
Journal of Alloys and Compounds 581 (2013) 246–249
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Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jalcom
A new reversible Mg3Ag–H2 system for hydrogen storage T.Z. Si, J.B. Zhang, D.M. Liu, Q.A. Zhang ⇑ School of Materials Science and Engineering, Anhui University of Technology, Maanshan, Anhui 243002, PR China
a r t i c l e
i n f o
Article history: Received 30 April 2013 Received in revised form 5 July 2013 Accepted 8 July 2013 Available online 16 July 2013 Keywords: Hydrogen storage Mg-based compound Thermodynamics Cyclic stability
a b s t r a c t For the first time, the compound Mg3Ag was employed as a medium for hydrogen storage. It has been demonstrated that the hydriding/dehydriding process of Mg3Ag is reversible through the reaction Mg3Ag + 2H2 M 2MgH2 + MgAg with obtaining altered thermodynamics. An enhanced cycling stability is also achieved by the capacity retention of 95% after 30 cycles, much higher than 70% for the pure Mg sample, which can be explained that the agglomeration and sintering of the resulting MgH2 are efficiently prevented by the formation of hard and brittle MgAg phase upon multi-cycling. Ó 2013 Elsevier B.V. All rights reserved.
1. Introduction Magnesium as a lightweight, abundant and low-cost metal with a high hydrogen storage capacity (7.6 wt.%) is regarded as one of the most promising material for solid-state hydrogen storage [1,2]. For pure Mg, the sluggish hydrogen ab-/desorption kinetics were significantly improved by alloying [3–6] and/or catalytic doping over the past several decades [2,7–9], but no substantive advance was obtained in the large enthalpy for desorption (DHd 75 kJ mol1 H2) [10], which retards its progress in practical application. Several novel strategies thus have been proposed to overcome the drawback of thermodynamics. One is to reduce the Mg or MgH2 particles size to nanoscale [11–15]. For example, nano-MgH2 with particle size of 3 nm exhibited a reduction in DHd by 11 kJ mol1 H2 as compared with the bulk MgH2 [13]. However, the desorption temperature of this nanosized MgH2 is still too high (at 290 °C) to the practical application. In addition, the susceptibility to oxidation makes it difficult to scale up the obtaining of Mg or MgH2 nanoparticles [14]. Another approach for altering the desorption thermodynamics is to destabilize MgH2 by adding reactive elements such as Al [16], Si [17,18] and Ge [10] to form intermetallic compounds. Among these destabilization systems, the MgH2–Si system is the most interesting one that shows a significantly reduced DHd value of 36.4 kJ mol1 H2 and a hydrogen desorption temperature as low as 23 °C under 0.1 MPa H2 based on the theoretical prediction [18]. However, the rehydrogenation of Mg2Si was unsuccessful [19,20], even though a slight hydrogenation of Mg2Si can be realized by ⇑ Corresponding author. Tel.: +86 555 2311871; fax: +86 555 2311570. E-mail address: [email protected] (Q.A. Zhang). 0925-8388/$ - see front matter Ó 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.jallcom.2013.07.054
reactive milling under H2 atmosphere [21]. Hence, it is necessary to explore new Mg-based compound – H2 systems for reversible hydrogen storage. Recently, several Mg3RE (RE = Mm [22], La [23,24] and Pr [25] etc.) compounds have attracted much attention because of their relatively moderate operating conditions for hydrogen storage. However, the higher hydride formation enthalpies of Mg3RE alloys than pure Mg make it not to form a single hydride (e.g., Mg3REHx) but to decompose into MgH2 and REH3, and these hydrogenation products cannot be regenerated back to Mg3RE after dehydrogenation. Moving down the Mg–Ag system, a similar compound Mg3Ag can also be obtained, but a complete understanding of the Mg3Ag– H2 system is still lacking up to now. Thus, in the present work, the Mg3Ag compound was employed as hydrogen storage medium and its hydriding/dehydriding behaviors were investigated in detail. The results obtained demonstrate that the Mg3Ag–H2 system not only has an excellent reversible manner, but also exhibits an altered thermodynamics properties. 2. Experimental details In view of the advantage of hydrogen metallurgy in synthesizing high-purity compound [26,27], the Mg3Ag sample was prepared by reactive sintering of compressed pellets of MgH2 (98%, 70 lm, Alfa Aesar) and Ag (99.9%, 1.0 lm, Alfa Aesar) powders in a molar ratio of 3:1 at 400 °C for 2 h under Ar atmosphere. Subsequently, the sample was crushed mechanically into powders with a particle size of 300 mesh. To evaluate the phase structures of the samples, XRD measurement was carried out using a Rigaku D/Max 2500VL/PC diffractometer with Cu Ka radiation at 50 kV and 200 mA. Based on the structural models, the XRD profiles were finally refined by the Rietveld program RIETAN-2000 [28]. The morphologies of the samples were observed using a Nova NanoSEM 430 scanning electron microscope (SEM). Differential scanning calorimetry (DSC) analysis was performed on a Netzsch STA 409 PC/PG unit under argon flow (30 ml min1) with a heating rate of 3 °C min1.
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Fig. 1. Rietveld refinement of the observed XRD pattern for the as-prepared Mg3Ag sample.
The kinetics properties of the Mg3Ag sample were measured using an automated Sieverts-type apparatus under the initial hydrogen pressures of 3.0 MPa for absorption and 0.001 MPa for desorption. To examine the cyclic stability, the samples were repeatedly hydrogenated for 1 h and dehydrogenated for 2 h at 300 °C.
3. Results and discussion 3.1. Hydrogen ab-/desorption mechanisms The observed XRD pattern and the Rietveld analysis result for the as-prepared Mg3Ag sample are shown in Fig. 1. It can be seen that the initial MgH2 and Ag disappear completely while the Mg3Ag and a small amount of Mg54Ag17 are obtained in the as-prepared sample. For the Rietveld refinement, the structure models of Mg3Ag and Mg54Ag17 were taken from the reported models [29], and the diffraction pattern calculated from the structure models is in good agreement with that measured. The Rietveld analysis results as listed in Table 1 further reveal that the phase abundances of Mg3Ag and Mg54Ag17 are 91 and 9 wt.%, respectively. Fig. 2a shows the first hydrogen absorption and desorption kinetic curves for the Mg3Ag sample at 300 °C. It can be seen that the Mg3Ag sample without activation can absorb 2.0 wt.% of hydrogen within 10 min and desorb almost the same content of absorption within 20 min. This indicates that the Mg3Ag–H2 system not only has a rapid hydrogen ab-/desorption process but also exhibits a good reversibility under the present conditions. In order to reveal the hydrogen storage mechanism, the XRD patterns of the Mg3Ag sample after first hydrogen absorption and desorption were further examined and shown in Fig. 2b. It is found that the hydrogenated sample consisting of MgH2 and MgAg shows the occurrence of hydrogen-induced decomposition during hydrogenation process of the Mg3Ag sample. The further hydrogenation of the MgAg compound is desired, however, the hydrogenation reaction of MgAg was not observed even under a hydrogen pressure of 10.0 MPa. According to the results of the Rietveld analysis (see Table 1), the amount of resulting MgH2 is 28 wt.%, suggesting that the theoretical hydrogen capacity of this system is about 2.1 wt.%
Fig. 2. (a) Hydriding and dehydriding kinetic curves for the Mg3Ag sample at 300 °C and (b) the corresponding XRD patterns after first hydrogen absorption and desorption.
(28% 7.6% = 2.1%), which is quite close to the measured value as shown in Fig. 2a. Significantly, the Mg3Ag and a small amount of Mg54Ag17 phases (see Fig. 1b) are regenerated after hydrogen desorption. Note that, Mg54Ag17, having a composition close to Mg3Ag, usually exists in an unequilibrium state because that it could be fully transformed into the equilibrium Mg3Ag by prolonging hydrogen desorption time [29–31]. Therefore, the results of phase analysis further indicate that the Mg3Ag–H2 system is reversible during the hydrogen ab-/desorption cycling, and the hydrogen ab-/desorption reactions can be described as follows:
Mg3 Ag þ 2H2 $ 2MgH2 þ MgAg
ð1Þ
3.2. Thermodynamics of the Mg3Ag–H2 system To understand the thermodynamic characteristics of the Mg3Ag–H2 system, the P–C isotherms of the Mg3Ag sample were measured and shown in Fig. 3a, in which the hydrogen storage properties of pure Mg obtained by dehydrogenating MgH2 is also compared. It can be seen that the Mg3Ag sample has an evidently elevated plateau pressure at 300 °C (0.26 MPa for absorption and 0.21 MPa for desorption) compared to pure Mg, implying the reduced enthalpy changes of hydrogen ab-/desorption. The midpoint of each plateau is taken as the equilibrium pressure, and then the van’t Hoff plots for the Mg3Ag–H2 and Mg–H2 systems can be drawn as in Fig. 3b. By the van’t Hoff equation, the reaction enthalpy changes of hydrogen absorption and desorption for the Mg3Ag– H2 system were calculated to be 68.2 and 69.8 kJ mol1 H2, respectively; which are lower than the values for pure Mg–H2 system, i.e., 74.3 kJ mol1 H2 for absorption and 74.5 kJ mol1 H2 for desorption. The decreased enthalpy changes suggest that Eq. (1) is a thermodynamically preferred way, which would lead to lower hydrogen desorption temperature. This deduction is further
Table 1 Structural parameters and phase abundances of the as-prepared, hydrogenated and dehydrogenated Mg3Ag samples. Sample
Phase
Space group
RI (%)
Lattice parameters
As-prepared Mg3Ag Rwp = 10.65% S = 2.91
Mg3Ag Mg54Ag17 MgH2 MgAg Mg3Ag Mg54Ag17
Fm-3 Immm P42/mnm Pm-3m Fm-3 Immm
3.75 2.88 2.55 1.89 2.89 3.48
17.5913(4) 14.2226(4) 4.4984(5) 3.3117(4) 17.6156(5) 14.2609(1)
a (Å)
Hydrogenated Mg3Ag Rwp = 8.97% S = 2.74 Dehydrogenated Mg3Ag Rwp = 10.84% S = 2.82
Abundance (wt.%)
b (Å)
c (Å)
14.1784(3)
14.6971(2) 3.0178(2)
14.2068(6)
14.6364(2)
91 9 28 72 93 7
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Fig. 3. (a) Pressure–composition isotherms and (b) van’t Hoff plots for the Mg3Ag–H2 and Mg–H2 systems, and (c) DSC curves for the hydrogenated Mg3Ag and Mg samples.
confirmed by DSC measurement (see Fig. 3c). The DSC curves show that the hydrogenated Mg3Ag sample (i.e., MgH2 + MgAg) has an endothermic peak centered at about 317 °C, corresponding to a decrease in dehydrogenation temperature by 33 °C, as compared with the pure MgH2. In turn, this suggests that the phase MgAg acts as a destabilization agent for MgH2 through the inverse reaction of Eq. (1). Different from the previous MgH2–M (M = reactive element) systems that suffer from poor reversibility [10,16,19], the present result opens an interesting topic to reversibly destabilize MgH2 by the solid-state reaction with an intermetallic compound instead of a pure element. 3.3. Cyclic stability of the Mg3Ag–H2 system
Fig. 4. (a) Capacity retention of hydrogen absorption versus cycle number for the Mg3Ag and pure Mg at 300 °C, as well as hydrogen (b) absorption and (c) desorption kinetic curves for the Mg3Ag sample after multiple cycling.
Fig. 4a presents the hydrogen capacity as a function of the cycling number for the Mg3Ag and pure Mg samples where the capacities are normalized to unity for the first value. For the pure Mg–H2 system, the capacity retention is about 75% after 30 cycles, which should be ascribed to the inevitable sintering of Mg upon multi-cycling [32]. In contrast, for the Mg3Ag–H2 system, the capacity retention is larger than 95% after 30 cycles, showing a noticeable improvement in cyclic stability, even as compared to the Nb hydride-doped [33] or TiF3-doped [34] Mg–H2 system. Note that, Fig. 4b and c clearly show that the degradation of hydrogen ab-/desorption kinetics is not obvious for the Mg3Ag sample with
Fig. 5. (a) XRD patterns and (b) SEM images of the Mg3Ag sample after first and 30th hydrogen absorption.
T.Z. Si et al. / Journal of Alloys and Compounds 581 (2013) 246–249
cycling. The good cyclic stability may be associated with the improvement in particles agglomeration during cycling. To verify this point, XRD measurements and SEM observations for the Mg3Ag sample after multiple cycling were carried out. Fig. 5a compares the XRD patterns of the Mg3Ag sample after first and 30th hydrogen absorption, respectively; showing that both samples are composed of MgH2 and MgAg phases. This means that the reversible phase transformation as shown in Eq. (1) keeps a high stability during cycling. The SEM observation shows that the particle size of the 30th hydrogenated sample obviously decreases, as compared with the first hydrogenated sample (see Fig. 5b). Given the fact that MgAg is a hard and brittle compound [35], it can be deduced that the phase MgAg acts as a buffer to prevent MgH2 particles from agglomerating and sintering during cycling, thus resulting in an excellent cyclic stability of the Mg3Ag–H2 system. 4. Conclusions In this work, for the first time, we have found a reversible Mg3Ag–H2 system for hydrogen storage with altered hydrogen ab-/ desorption thermodynamics. During dehydrogenation process, the phase MgAg acts as a thermodynamic destabilization agent for MgH2, resulting in decrease in enthalpy change and temperature of hydrogen desorption. Moreover, the hard and brittle MgAg phase acts as a barrier to prevent MgH2 particles from agglomerating and sintering during hydrogen ab-/desorption cycling, thus leading to a good cyclic stability. These results demonstrate a new pathway to reversibly destabilize MgH2 by the solid-state reaction with an intermetallic compound instead of a pure element. Acknowledgements This work was financially supported by the National Natural Science Foundation of China (No. 51271002) and Natural Science Foundation of Anhui Province (No. 1208085ME83). References [1] L. Schlapbach, A. Zuttel, Nature 414 (2001) 353. [2] J. Cermak, L. Karl, J. Alloys Comp. 546 (2013) 129.
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[3] H.X. Chen, Z.M. Wang, H.Y. Zhou, C.Y. Ni, J.Q. Deng, Q.R. Yao, J. Alloys Comp. 563 (2013) 1. [4] M. Ponthieu, J.F. Fernandez, F. Cuevas, J.R. Ares, F. Leardini, J. Bodega, C. Sanchez, J. Alloys Comp. 548 (2013) 96. [5] Q.A. Zhang, L.X. Zhang, Q.Q. Wang, J. Alloys Comp. 551 (2013) 376. [6] Y.L. Du, L. Xu, Y. Shen, W. Zhuang, S.H. Zhang, G. Chen, Int. J. Hydrogen Energy 38 (2013) 4670. [7] B.H. Chen, C.H. Kuo, J.R. Ku, P.S. Yan, C.J. Huang, M.S. Jeng, F.H. Tsau, J. Alloys Comp. 568 (2013) 78. [8] M.O.T. da Conceicao, M.C. Brum, D.S. dos Santos, M.L. Dias, J. Alloys Comp. 550 (2013) 179. [9] X.L. Zhu, L.C. Pei, Z.Y. Zhao, B.Z. Liu, S.M. Han, R.B. Wang, J. Alloys Comp. 577 (2013) 64. [10] G.S. Walker, M. Abbas, D.M. Grant, C. Udeh, Chem. Commun. 47 (2011) 8001. [11] T. Liu, H.L. Shen, Y. Liu, L. Xie, J.L. Qu, H.Y. Shao, X.G. Li, J. Power Sources 227 (2013) 86. [12] M. Paskevicius, D.A. Sheppard, C.E. Buckley, J. Am. Chem. Soc. 132 (2010) 5077. [13] Z. Zhao-Karger, J.J. Hu, A. Roth, D. Wang, C. Kubel, W. Lohstroh, M. Fichtner, Chem. Commun. 46 (2010) 8353. [14] B. Peng, J. Liang, Z.L. Tao, J. Chen, J. Mater. Chem. 19 (2009) 2877. [15] K.J. Jeon, H.R. Moon, A.M. Ruminski, B. Jiang, C. Kisielowski, R. Bardhan, J.J. Urban, Nat. Mater. 10 (2011) 286. [16] A. Zaluska, L. Zaluski, J.O. Strom-Olsen, Appl. Phys. A: Mater. Sci. Process 72 (2001) 157. [17] J.J. Vajo, F. Mertens, C.C. Ahn, R.C. Bowman, B. Fultz, J. Phys. Chem. B 108 (2004) 13977. [18] Y.W. Cho, J.H. Shim, B.J. Lee, Calphad 30 (2006) 65. [19] S.W.H. Eijt, R. Kind, S. Singh, H. Schut, W.J. Legerstee, R.W.A. Hendrikx, V.L. Svetchnikov, R.J. Westerwaal, B. Dam, J. Appl. Phys. 105 (2009) 043514. [20] S.T. Kelly, S.L. Van Atta, J.J. Vajo, G.L. Olson, B.M. Clemens, Nanotechnology 20 (2009) 204017. [21] R. Janot, F. Cuevas, M. Latroche, A. Percheron-Guégan, Intermetallics 14 (2006) 163. [22] L.Z. Ouyang, L. Yao, X.S. Yang, L.Q. Li, M. Zhu, Int. J. Hydrogen Energy 35 (2010) 8275. [23] H.J. Lin, H. Wang, J.W. Liu, M. Zhu, Int. J. Hydrogen Energy 37 (2012) 1145. [24] L.Z. Ouyang, J.M. Huang, C.J. Fang, H. Wang, J.W. Liu, Q.A. Zhang, D.L. Sun, M. Zhu, J. Alloys Comp. (2013), http://dx.doi.org/10.1016/j.jallcom.2013.03.153. [25] L.Z. Ouyang, X.S. Yang, H.W. Dong, M. Zhu, Scr. Mater. 61 (2009) 339. [26] R.V. Denys, A.A. Poletaev, J.K. Solberg, B.P. Tarasov, V.A. Yartys, Acta Mater. 58 (2010) 2510. [27] B. Zhao, F. Fang, D.L. Sun, Q.A. Zhang, S.Q. Wei, F.L. Cao, H. Sun, L.Z. Ouyang, M. Zhu, J. Solid State Chem. 192 (2012) 210. [28] F. Izumi, T. Ikeda, Mater. Sci. Forum 321–323 (2000) 198. [29] A.V. Arakcheeva, O.G. Karpinskii, V.E. Kolesnichenko, Sov. Phys. Crystallogr. 33 (1988) 907. [30] M.V. Prokofev, V.E. Kolesnichenko, V.V. Karonik, Inorg. Mater. 21 (1985) 1168. [31] A. Rakowska, M. Podosek, R. Ciach, Mater. Des. 18 (1997) 279. [32] W.P. Kalisvaart, A. Kubis, M. Danaie, B.S. Amirkhiz, D. Mitlin, Acta Mater. 59 (2011) 2083. [33] S.A. Jin, J.H. Shim, J.P. Ahn, Y.W. Cho, K.W. Yi, Acta Mater. 55 (2007) 5073. [34] L.P. Ma, P. Wang, H.M. Cheng, J. Alloys Comp. 432 (2007) L1. [35] Z.W. Lu, D.W. Zhou, J.P. Bai, C. Lu, Z.G. Zhong, G.Q. Li, J. Alloys Comp. 550 (2013) 406.
Contents lists available at ScienceDirect
Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jalcom
A new reversible Mg3Ag–H2 system for hydrogen storage T.Z. Si, J.B. Zhang, D.M. Liu, Q.A. Zhang ⇑ School of Materials Science and Engineering, Anhui University of Technology, Maanshan, Anhui 243002, PR China
a r t i c l e
i n f o
Article history: Received 30 April 2013 Received in revised form 5 July 2013 Accepted 8 July 2013 Available online 16 July 2013 Keywords: Hydrogen storage Mg-based compound Thermodynamics Cyclic stability
a b s t r a c t For the first time, the compound Mg3Ag was employed as a medium for hydrogen storage. It has been demonstrated that the hydriding/dehydriding process of Mg3Ag is reversible through the reaction Mg3Ag + 2H2 M 2MgH2 + MgAg with obtaining altered thermodynamics. An enhanced cycling stability is also achieved by the capacity retention of 95% after 30 cycles, much higher than 70% for the pure Mg sample, which can be explained that the agglomeration and sintering of the resulting MgH2 are efficiently prevented by the formation of hard and brittle MgAg phase upon multi-cycling. Ó 2013 Elsevier B.V. All rights reserved.
1. Introduction Magnesium as a lightweight, abundant and low-cost metal with a high hydrogen storage capacity (7.6 wt.%) is regarded as one of the most promising material for solid-state hydrogen storage [1,2]. For pure Mg, the sluggish hydrogen ab-/desorption kinetics were significantly improved by alloying [3–6] and/or catalytic doping over the past several decades [2,7–9], but no substantive advance was obtained in the large enthalpy for desorption (DHd 75 kJ mol1 H2) [10], which retards its progress in practical application. Several novel strategies thus have been proposed to overcome the drawback of thermodynamics. One is to reduce the Mg or MgH2 particles size to nanoscale [11–15]. For example, nano-MgH2 with particle size of 3 nm exhibited a reduction in DHd by 11 kJ mol1 H2 as compared with the bulk MgH2 [13]. However, the desorption temperature of this nanosized MgH2 is still too high (at 290 °C) to the practical application. In addition, the susceptibility to oxidation makes it difficult to scale up the obtaining of Mg or MgH2 nanoparticles [14]. Another approach for altering the desorption thermodynamics is to destabilize MgH2 by adding reactive elements such as Al [16], Si [17,18] and Ge [10] to form intermetallic compounds. Among these destabilization systems, the MgH2–Si system is the most interesting one that shows a significantly reduced DHd value of 36.4 kJ mol1 H2 and a hydrogen desorption temperature as low as 23 °C under 0.1 MPa H2 based on the theoretical prediction [18]. However, the rehydrogenation of Mg2Si was unsuccessful [19,20], even though a slight hydrogenation of Mg2Si can be realized by ⇑ Corresponding author. Tel.: +86 555 2311871; fax: +86 555 2311570. E-mail address: [email protected] (Q.A. Zhang). 0925-8388/$ - see front matter Ó 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.jallcom.2013.07.054
reactive milling under H2 atmosphere [21]. Hence, it is necessary to explore new Mg-based compound – H2 systems for reversible hydrogen storage. Recently, several Mg3RE (RE = Mm [22], La [23,24] and Pr [25] etc.) compounds have attracted much attention because of their relatively moderate operating conditions for hydrogen storage. However, the higher hydride formation enthalpies of Mg3RE alloys than pure Mg make it not to form a single hydride (e.g., Mg3REHx) but to decompose into MgH2 and REH3, and these hydrogenation products cannot be regenerated back to Mg3RE after dehydrogenation. Moving down the Mg–Ag system, a similar compound Mg3Ag can also be obtained, but a complete understanding of the Mg3Ag– H2 system is still lacking up to now. Thus, in the present work, the Mg3Ag compound was employed as hydrogen storage medium and its hydriding/dehydriding behaviors were investigated in detail. The results obtained demonstrate that the Mg3Ag–H2 system not only has an excellent reversible manner, but also exhibits an altered thermodynamics properties. 2. Experimental details In view of the advantage of hydrogen metallurgy in synthesizing high-purity compound [26,27], the Mg3Ag sample was prepared by reactive sintering of compressed pellets of MgH2 (98%, 70 lm, Alfa Aesar) and Ag (99.9%, 1.0 lm, Alfa Aesar) powders in a molar ratio of 3:1 at 400 °C for 2 h under Ar atmosphere. Subsequently, the sample was crushed mechanically into powders with a particle size of 300 mesh. To evaluate the phase structures of the samples, XRD measurement was carried out using a Rigaku D/Max 2500VL/PC diffractometer with Cu Ka radiation at 50 kV and 200 mA. Based on the structural models, the XRD profiles were finally refined by the Rietveld program RIETAN-2000 [28]. The morphologies of the samples were observed using a Nova NanoSEM 430 scanning electron microscope (SEM). Differential scanning calorimetry (DSC) analysis was performed on a Netzsch STA 409 PC/PG unit under argon flow (30 ml min1) with a heating rate of 3 °C min1.
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Fig. 1. Rietveld refinement of the observed XRD pattern for the as-prepared Mg3Ag sample.
The kinetics properties of the Mg3Ag sample were measured using an automated Sieverts-type apparatus under the initial hydrogen pressures of 3.0 MPa for absorption and 0.001 MPa for desorption. To examine the cyclic stability, the samples were repeatedly hydrogenated for 1 h and dehydrogenated for 2 h at 300 °C.
3. Results and discussion 3.1. Hydrogen ab-/desorption mechanisms The observed XRD pattern and the Rietveld analysis result for the as-prepared Mg3Ag sample are shown in Fig. 1. It can be seen that the initial MgH2 and Ag disappear completely while the Mg3Ag and a small amount of Mg54Ag17 are obtained in the as-prepared sample. For the Rietveld refinement, the structure models of Mg3Ag and Mg54Ag17 were taken from the reported models [29], and the diffraction pattern calculated from the structure models is in good agreement with that measured. The Rietveld analysis results as listed in Table 1 further reveal that the phase abundances of Mg3Ag and Mg54Ag17 are 91 and 9 wt.%, respectively. Fig. 2a shows the first hydrogen absorption and desorption kinetic curves for the Mg3Ag sample at 300 °C. It can be seen that the Mg3Ag sample without activation can absorb 2.0 wt.% of hydrogen within 10 min and desorb almost the same content of absorption within 20 min. This indicates that the Mg3Ag–H2 system not only has a rapid hydrogen ab-/desorption process but also exhibits a good reversibility under the present conditions. In order to reveal the hydrogen storage mechanism, the XRD patterns of the Mg3Ag sample after first hydrogen absorption and desorption were further examined and shown in Fig. 2b. It is found that the hydrogenated sample consisting of MgH2 and MgAg shows the occurrence of hydrogen-induced decomposition during hydrogenation process of the Mg3Ag sample. The further hydrogenation of the MgAg compound is desired, however, the hydrogenation reaction of MgAg was not observed even under a hydrogen pressure of 10.0 MPa. According to the results of the Rietveld analysis (see Table 1), the amount of resulting MgH2 is 28 wt.%, suggesting that the theoretical hydrogen capacity of this system is about 2.1 wt.%
Fig. 2. (a) Hydriding and dehydriding kinetic curves for the Mg3Ag sample at 300 °C and (b) the corresponding XRD patterns after first hydrogen absorption and desorption.
(28% 7.6% = 2.1%), which is quite close to the measured value as shown in Fig. 2a. Significantly, the Mg3Ag and a small amount of Mg54Ag17 phases (see Fig. 1b) are regenerated after hydrogen desorption. Note that, Mg54Ag17, having a composition close to Mg3Ag, usually exists in an unequilibrium state because that it could be fully transformed into the equilibrium Mg3Ag by prolonging hydrogen desorption time [29–31]. Therefore, the results of phase analysis further indicate that the Mg3Ag–H2 system is reversible during the hydrogen ab-/desorption cycling, and the hydrogen ab-/desorption reactions can be described as follows:
Mg3 Ag þ 2H2 $ 2MgH2 þ MgAg
ð1Þ
3.2. Thermodynamics of the Mg3Ag–H2 system To understand the thermodynamic characteristics of the Mg3Ag–H2 system, the P–C isotherms of the Mg3Ag sample were measured and shown in Fig. 3a, in which the hydrogen storage properties of pure Mg obtained by dehydrogenating MgH2 is also compared. It can be seen that the Mg3Ag sample has an evidently elevated plateau pressure at 300 °C (0.26 MPa for absorption and 0.21 MPa for desorption) compared to pure Mg, implying the reduced enthalpy changes of hydrogen ab-/desorption. The midpoint of each plateau is taken as the equilibrium pressure, and then the van’t Hoff plots for the Mg3Ag–H2 and Mg–H2 systems can be drawn as in Fig. 3b. By the van’t Hoff equation, the reaction enthalpy changes of hydrogen absorption and desorption for the Mg3Ag– H2 system were calculated to be 68.2 and 69.8 kJ mol1 H2, respectively; which are lower than the values for pure Mg–H2 system, i.e., 74.3 kJ mol1 H2 for absorption and 74.5 kJ mol1 H2 for desorption. The decreased enthalpy changes suggest that Eq. (1) is a thermodynamically preferred way, which would lead to lower hydrogen desorption temperature. This deduction is further
Table 1 Structural parameters and phase abundances of the as-prepared, hydrogenated and dehydrogenated Mg3Ag samples. Sample
Phase
Space group
RI (%)
Lattice parameters
As-prepared Mg3Ag Rwp = 10.65% S = 2.91
Mg3Ag Mg54Ag17 MgH2 MgAg Mg3Ag Mg54Ag17
Fm-3 Immm P42/mnm Pm-3m Fm-3 Immm
3.75 2.88 2.55 1.89 2.89 3.48
17.5913(4) 14.2226(4) 4.4984(5) 3.3117(4) 17.6156(5) 14.2609(1)
a (Å)
Hydrogenated Mg3Ag Rwp = 8.97% S = 2.74 Dehydrogenated Mg3Ag Rwp = 10.84% S = 2.82
Abundance (wt.%)
b (Å)
c (Å)
14.1784(3)
14.6971(2) 3.0178(2)
14.2068(6)
14.6364(2)
91 9 28 72 93 7
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Fig. 3. (a) Pressure–composition isotherms and (b) van’t Hoff plots for the Mg3Ag–H2 and Mg–H2 systems, and (c) DSC curves for the hydrogenated Mg3Ag and Mg samples.
confirmed by DSC measurement (see Fig. 3c). The DSC curves show that the hydrogenated Mg3Ag sample (i.e., MgH2 + MgAg) has an endothermic peak centered at about 317 °C, corresponding to a decrease in dehydrogenation temperature by 33 °C, as compared with the pure MgH2. In turn, this suggests that the phase MgAg acts as a destabilization agent for MgH2 through the inverse reaction of Eq. (1). Different from the previous MgH2–M (M = reactive element) systems that suffer from poor reversibility [10,16,19], the present result opens an interesting topic to reversibly destabilize MgH2 by the solid-state reaction with an intermetallic compound instead of a pure element. 3.3. Cyclic stability of the Mg3Ag–H2 system
Fig. 4. (a) Capacity retention of hydrogen absorption versus cycle number for the Mg3Ag and pure Mg at 300 °C, as well as hydrogen (b) absorption and (c) desorption kinetic curves for the Mg3Ag sample after multiple cycling.
Fig. 4a presents the hydrogen capacity as a function of the cycling number for the Mg3Ag and pure Mg samples where the capacities are normalized to unity for the first value. For the pure Mg–H2 system, the capacity retention is about 75% after 30 cycles, which should be ascribed to the inevitable sintering of Mg upon multi-cycling [32]. In contrast, for the Mg3Ag–H2 system, the capacity retention is larger than 95% after 30 cycles, showing a noticeable improvement in cyclic stability, even as compared to the Nb hydride-doped [33] or TiF3-doped [34] Mg–H2 system. Note that, Fig. 4b and c clearly show that the degradation of hydrogen ab-/desorption kinetics is not obvious for the Mg3Ag sample with
Fig. 5. (a) XRD patterns and (b) SEM images of the Mg3Ag sample after first and 30th hydrogen absorption.
T.Z. Si et al. / Journal of Alloys and Compounds 581 (2013) 246–249
cycling. The good cyclic stability may be associated with the improvement in particles agglomeration during cycling. To verify this point, XRD measurements and SEM observations for the Mg3Ag sample after multiple cycling were carried out. Fig. 5a compares the XRD patterns of the Mg3Ag sample after first and 30th hydrogen absorption, respectively; showing that both samples are composed of MgH2 and MgAg phases. This means that the reversible phase transformation as shown in Eq. (1) keeps a high stability during cycling. The SEM observation shows that the particle size of the 30th hydrogenated sample obviously decreases, as compared with the first hydrogenated sample (see Fig. 5b). Given the fact that MgAg is a hard and brittle compound [35], it can be deduced that the phase MgAg acts as a buffer to prevent MgH2 particles from agglomerating and sintering during cycling, thus resulting in an excellent cyclic stability of the Mg3Ag–H2 system. 4. Conclusions In this work, for the first time, we have found a reversible Mg3Ag–H2 system for hydrogen storage with altered hydrogen ab-/ desorption thermodynamics. During dehydrogenation process, the phase MgAg acts as a thermodynamic destabilization agent for MgH2, resulting in decrease in enthalpy change and temperature of hydrogen desorption. Moreover, the hard and brittle MgAg phase acts as a barrier to prevent MgH2 particles from agglomerating and sintering during hydrogen ab-/desorption cycling, thus leading to a good cyclic stability. These results demonstrate a new pathway to reversibly destabilize MgH2 by the solid-state reaction with an intermetallic compound instead of a pure element. Acknowledgements This work was financially supported by the National Natural Science Foundation of China (No. 51271002) and Natural Science Foundation of Anhui Province (No. 1208085ME83). References [1] L. Schlapbach, A. Zuttel, Nature 414 (2001) 353. [2] J. Cermak, L. Karl, J. Alloys Comp. 546 (2013) 129.
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[3] H.X. Chen, Z.M. Wang, H.Y. Zhou, C.Y. Ni, J.Q. Deng, Q.R. Yao, J. Alloys Comp. 563 (2013) 1. [4] M. Ponthieu, J.F. Fernandez, F. Cuevas, J.R. Ares, F. Leardini, J. Bodega, C. Sanchez, J. Alloys Comp. 548 (2013) 96. [5] Q.A. Zhang, L.X. Zhang, Q.Q. Wang, J. Alloys Comp. 551 (2013) 376. [6] Y.L. Du, L. Xu, Y. Shen, W. Zhuang, S.H. Zhang, G. Chen, Int. J. Hydrogen Energy 38 (2013) 4670. [7] B.H. Chen, C.H. Kuo, J.R. Ku, P.S. Yan, C.J. Huang, M.S. Jeng, F.H. Tsau, J. Alloys Comp. 568 (2013) 78. [8] M.O.T. da Conceicao, M.C. Brum, D.S. dos Santos, M.L. Dias, J. Alloys Comp. 550 (2013) 179. [9] X.L. Zhu, L.C. Pei, Z.Y. Zhao, B.Z. Liu, S.M. Han, R.B. Wang, J. Alloys Comp. 577 (2013) 64. [10] G.S. Walker, M. Abbas, D.M. Grant, C. Udeh, Chem. Commun. 47 (2011) 8001. [11] T. Liu, H.L. Shen, Y. Liu, L. Xie, J.L. Qu, H.Y. Shao, X.G. Li, J. Power Sources 227 (2013) 86. [12] M. Paskevicius, D.A. Sheppard, C.E. Buckley, J. Am. Chem. Soc. 132 (2010) 5077. [13] Z. Zhao-Karger, J.J. Hu, A. Roth, D. Wang, C. Kubel, W. Lohstroh, M. Fichtner, Chem. Commun. 46 (2010) 8353. [14] B. Peng, J. Liang, Z.L. Tao, J. Chen, J. Mater. Chem. 19 (2009) 2877. [15] K.J. Jeon, H.R. Moon, A.M. Ruminski, B. Jiang, C. Kisielowski, R. Bardhan, J.J. Urban, Nat. Mater. 10 (2011) 286. [16] A. Zaluska, L. Zaluski, J.O. Strom-Olsen, Appl. Phys. A: Mater. Sci. Process 72 (2001) 157. [17] J.J. Vajo, F. Mertens, C.C. Ahn, R.C. Bowman, B. Fultz, J. Phys. Chem. B 108 (2004) 13977. [18] Y.W. Cho, J.H. Shim, B.J. Lee, Calphad 30 (2006) 65. [19] S.W.H. Eijt, R. Kind, S. Singh, H. Schut, W.J. Legerstee, R.W.A. Hendrikx, V.L. Svetchnikov, R.J. Westerwaal, B. Dam, J. Appl. Phys. 105 (2009) 043514. [20] S.T. Kelly, S.L. Van Atta, J.J. Vajo, G.L. Olson, B.M. Clemens, Nanotechnology 20 (2009) 204017. [21] R. Janot, F. Cuevas, M. Latroche, A. Percheron-Guégan, Intermetallics 14 (2006) 163. [22] L.Z. Ouyang, L. Yao, X.S. Yang, L.Q. Li, M. Zhu, Int. J. Hydrogen Energy 35 (2010) 8275. [23] H.J. Lin, H. Wang, J.W. Liu, M. Zhu, Int. J. Hydrogen Energy 37 (2012) 1145. [24] L.Z. Ouyang, J.M. Huang, C.J. Fang, H. Wang, J.W. Liu, Q.A. Zhang, D.L. Sun, M. Zhu, J. Alloys Comp. (2013), http://dx.doi.org/10.1016/j.jallcom.2013.03.153. [25] L.Z. Ouyang, X.S. Yang, H.W. Dong, M. Zhu, Scr. Mater. 61 (2009) 339. [26] R.V. Denys, A.A. Poletaev, J.K. Solberg, B.P. Tarasov, V.A. Yartys, Acta Mater. 58 (2010) 2510. [27] B. Zhao, F. Fang, D.L. Sun, Q.A. Zhang, S.Q. Wei, F.L. Cao, H. Sun, L.Z. Ouyang, M. Zhu, J. Solid State Chem. 192 (2012) 210. [28] F. Izumi, T. Ikeda, Mater. Sci. Forum 321–323 (2000) 198. [29] A.V. Arakcheeva, O.G. Karpinskii, V.E. Kolesnichenko, Sov. Phys. Crystallogr. 33 (1988) 907. [30] M.V. Prokofev, V.E. Kolesnichenko, V.V. Karonik, Inorg. Mater. 21 (1985) 1168. [31] A. Rakowska, M. Podosek, R. Ciach, Mater. Des. 18 (1997) 279. [32] W.P. Kalisvaart, A. Kubis, M. Danaie, B.S. Amirkhiz, D. Mitlin, Acta Mater. 59 (2011) 2083. [33] S.A. Jin, J.H. Shim, J.P. Ahn, Y.W. Cho, K.W. Yi, Acta Mater. 55 (2007) 5073. [34] L.P. Ma, P. Wang, H.M. Cheng, J. Alloys Comp. 432 (2007) L1. [35] Z.W. Lu, D.W. Zhou, J.P. Bai, C. Lu, Z.G. Zhong, G.Q. Li, J. Alloys Comp. 550 (2013) 406.