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Destabilizing the dehydriding thermodynamics of MgH2 by reversible intermetallics formation in Mg−Ag−Zn ternary alloys
Journal of Power Sources 396 (2018) 796–802
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Destabilizing the dehydriding thermodynamics of MgH2 by reversible intermetallics formation in Mg−Ag−Zn ternary alloys
T
Yanshan Lu, Hui Wang, Jiangwen Liu, Liuzhang Ouyang, Min Zhu∗ School of Materials Science and Engineering and Guangdong Provincial Key Laboratory of Advanced Energy Storage Materials, South China University of Technology, Guangzhou, 510641, China
H I GH L IG H T S
Mg−Ag−Zn ternary alloys show fully reversible de-/hydrogenation performance. • The Mg−Ag−Zn ternary alloys exhibit elevated equilibrium pressure for MgH . • The • Enhanced dehydrogenation kinetics is achieved in the Mg−Ag−Zn ternary alloys. 2
A R T I C LE I N FO
A B S T R A C T
Keywords: Hydrogen storage Thermodynamic destabilization Mg−Ag−Zn Reversibility
In this study, Mg−Ag−Zn ternary alloys are employed as carriers for hydrogen storage, which show improved thermodynamic and kinetic properties compared with pure Mg. Mg90Ag7.5Zn2.5 alloy exhibits a reversible phase transformation during hydriding-dehydriding process with a maximum hydrogen storage capacity of approximately 4.2 wt%. Two-step reactions occur in the dehydriding process of Mg90Ag7.5Zn2.5. In the first step, a portion of MgH2 reacts with MgAg and MgZn1.8Ag0.2 phases to transform to Mg54(Ag, Zn)17 solid solution and release hydrogen. After that, the remaining MgH2 decomposes to Mg. Because of the modified reaction pathway in the first step dehydrogenation, the dehydriding equilibrium pressure at 300 °C increases to 0.28 MPa. Besides, the apparent activation energy (Ea) for dehydrogenation of the Mg90Ag7.5Zn2.5 alloy is lowered to 118.7 kJ mol−1. By fully utilizing the first step reaction, another alloy Mg78Ag16.5Zn5.5 is designed, which presents intact destabilized dehydriding thermodynamics and a lower dehydrogenation Ea of 116.5 kJ mol−1. Our work demonstrates that alloying Mg with elements that can form compounds with low formation enthalpies, and the hydrides of the elements have low stabilities, are the main guidelines for the designing of the Mg-based hydrogen storage alloys with destabilized dehydriding thermodynamics by reversible intermetallics formation.
1. Introduction Mg is a very attractive material for hydrogen storage owing to its reversible reaction with hydrogen to form MgH2 with H capacity of 7.6 wt% [1,2]. However, the strong ionic characteristics of the Mg−H bond, leads to too high dehydriding enthalpy of MgH2, being 74.7 kJ mol−1 H2 [3], which is much larger than that required for dehydriding at room temperature, being 20–40 kJ mol−1 H2 [4]. Thus, the dehydriding temperature of MgH2 is too high to meet the practical application such as hydrogen-fueled vehicles [5,6]. Many attempts have been made to destabilize the thermodynamics of MgH2; however, the destabilization is far from enough in addition to the inducing of other drawbacks. For example, the dehydriding enthalpy can be reduced by alloying [7–9], however, the hydrogen
∗
Corresponding author. E-mail address: [email protected] (M. Zhu).
https://doi.org/10.1016/j.jpowsour.2018.06.060 Received 28 April 2018; Received in revised form 8 June 2018; Accepted 12 June 2018 0378-7753/ © 2018 Elsevier B.V. All rights reserved.
storage capacity also reduces, being 64.5 kJ mol−1 H2 and 3.6 wt%, respectively, in the case of forming Mg2Ni [10]. Although nanostructuring Mg/MgH2 can reduce the dehydriding enthalpy [11–13], the effect is significant only when the size reduces to several nanometer, which causes problems in fabrication of materials, stability of nanostructure and so on. For instance, the decomposition enthalpy of Mg nanowires with diameter of 30–50 nm reduced to 65.3 kJ mol−1 H2, but the Mg nanowires could not keep their structure after 10 hydrogenation and dehydrogenation cycles [14]. Recently, catalyzed MgH2 with a lower dehydriding enthalpy has been reported [15,16]. It is found that the dehydrogenation enthalpy was reduced to 66.34 kJ mol−1 H2 for MgH2 catalyzed by graphene sheet templated Fe3O4 nanoparticles, however, a decrease of the dehydrogenation entropy (ΔS = 125.85 J mol−1 H2) also happens [16], which has a negative
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MgH2 during milling, a hydrogen pressure of 10 MPa was filled in the milling jar. X-ray diffraction (XRD) analysis was conducted on a PANalytical Empyrean X-ray diffractometer equipped with Cu Kα radiation (λ = 1.54060 Å). High purity Si powder was added into XRD samples as an internal reference to correct the instrumental zero-shift. To examine the reaction process, the PCI desorption test was manually terminated at different stages, and the sample holder was quickly cooled while the hydrogen pressure of sample cell was kept. Then the sample was taken out for XRD analysis at room temperature. The lattice constants and the unit cell volume were calculated from the XRD data by Rietveld method using Highscore Plus software. The microstructure of samples at different states was observed by a scanning electron microscope (SEM, Zeiss Merlin) equipped with an energy-dispersive spectrometer (EDS), and a JEOL JEM-2100 transmission electron microscopy (TEM) operating at 200 kV. The TEM samples were prepared by dispersing powder sample in hexane by an ultrasonic cleaner, and then dropped the suspension onto a copper grid. The hydrogen storage properties were measured using a Sieverttype apparatus (SETARAM PCT-Pro2000). Prior to the measurements, the samples were activated through three hydriding-dehydriding cycles at 320 °C under a hydrogen pressure of 4.5 MPa. PCI test was conducted with the hydrogen pressures varied from 4.5 to 0.002 MPa. The dehydrogenation enthalpy ΔH and entropy ΔS were calculated based on the van't Hoff equation by adopting the pressure at the midpoint of the plateau in each PCI curve as the equilibrium pressure. Isothermal hydrogen desorption kinetics was determined at specific temperatures, and the sample cell was pumped to vacuum before testing. The kinetic curves were analyzed using the Johnson−Mehl−Avrami−Kolmogorov η (JMAK) model [37,38], expressed as α= 1 − e−(kt) , where α is the reaction fraction, k is the rate constant, t is the reaction time, and η is the reaction order. Then the apparent activation energy Ea for dehydrogenation was calculated according to the Arrhenius equation: k= Ae−Ea/RT , where A is the pre-exponential factor, R is the gas constant, and T is the Kelvin temperature. All the sample handlings were carried out in a glove box filled with Ar with water and oxygen content less than 3 ppm.
influence on the thermodynamic destabilization of MgH2. Changing the reaction pathway by adding a reactive additive is another attractive route of destabilizing thermodynamics of MgH2 and first proposed by Vajo et al. [17,18]. They found Mg2Si would form upon dehydrogenation when adding Si into MgH2, decreasing reaction enthalpy to 36.4 kJ mol−1, and increasing the dehydriding pressure to 7.5 bar at 300 °C. The thermodynamic destabilization also achieved in MgH2/Ge system with the dehydrogenation temperature decreasing to 130 °C [19]. It is reported that Sn is another additive to destabilize the thermodynamics of MgH2 effectively [20,21]. Unfortunately, the reversibility of those systems is poor. To overcome this problem, a new route, which can be briefly expressed as destabilizing thermodynamics by reversible intermetallics formation (DTRIF), was developed in Mg−In system [22,23]. In this system, dehydriding enthalpy of MgH2 was reduced by reacting with the β intermetallic formed in the hydrogenation of Mg(In) solid solution, and the reaction is highly reversible. Following the DTRIF route, a series of Mg-based alloys have been explored and realized reversible hydrogen storage reactions with destabilized thermodynamics, including Mg–Cd [24,25], Mg–Ga [26], Mg–Ag [27–29], Mg–In–Al [30,31], Mg–In–Cd [32], Mg–In–Ni [33], Mg–Ag–In [34] and Mg–Ag–Al [35] alloy systems. It is obvious that most of those DTRIF systems contain In and Ga, which are expensive rare metals. Besides, the reduction of dehydriding enthalpy of those DTRIF systems is not large enough and most of them suffer from slow hydrogen desorption kinetics, in particular, Mg–In–Al [30,31], Mg–In–Cd [32] and Mg–Ag–Al [35] alloys. Therefore, it is of great significant to explore other DTRIF systems with a reduced amount of expensive elements, and to have larger thermodynamic destabilization and enhanced kinetics, i.e., dual tuning of thermodynamics and kinetics of MgH2. According to the Ag−Zn binary phase diagram, Zn can dissolve in Ag to form a Ag(Zn) solid solution with the maximum solubility of 30 at % at 100 °C. Thus, it is expected that intermetallic may formed in Mg−Ag−Zn ternary alloy, which may lead to a new DTRIF system. In addition, Zn has a low melting point of 420 °C, and thus Zn atoms may diffuse easily during hydrogenation and dehydrogenation at lower temperatures, which would possibly promote the hydrogen sorption kinetics. Indeed, it has been shown that the ultrafine Mg−Zn particles, being 400 nm and consisting of Mg(Zn) solid solution and amorphous, could enhance hydrogen absorption kinetics and decrease the activation energy for hydrogenation to 56.3 kJ mol−1 [36]. In the present work, with the aiming of the dual tuning of thermodynamics and kinetics of MgH2, we explored the DTRIF alloys in Mg−Ag−Zn ternary alloy system. The dehydriding thermodynamic destabilization has been achieved and the hydrogen desorption pressure at 300 °C increased from 0.16 to 0.28 MPa in comparison with pure MgH2. The hydrogen desorption kinetics are also improved with activation energy lowering to 116.5 kJ mol−1. Besides, we proposed the main guidelines for designing DTRIF systems based on the sum up of the systems already known from the literature.
3. Results and discussion 3.1. Achieving a sole reversible thermodynamic destabilized dehydriding reaction in Mg−Ag−Zn ternary alloy Compared with the catalyzed Mg/MgH2 [15,16], the hydrogen storage capacities of the DTRIF systems are usually lower, because a part of Mg would form Mg-based compounds after hydrogen absorption. Therefore, the alloy composition was designed as Mg90Ag7.5Zn2.5 with the consideration that the theoretical hydrogen storage capacity of the alloy can exceed 5 wt% by counting the Mg content. Fig. 1a and b shows the XRD patterns of the powder mixture of Ag and Zn before and after sintering, respectively. As it shown, after sintering at 500 °C for 6 h, the diffraction peaks of Zn disappear, while the Ag peaks are obviously shifted toward higher 2θ angle in comparison with the powder mixture of Ag and Zn. The lattice constant of Ag decreases from 4.0837 Å to 4.0262 Å, indicating the formation of Ag(Zn) solid solution after sintering. Fig. 1c is the XRD pattern of Mg90Ag7.5Zn2.5 ternary alloy prepared by ball milling Ag(Zn) solid solution with MgH2 under a 10 MPa hydrogen pressure. The observed phases are MgH2 and Ag(Zn) solid solution, and no other phase is detected, which means that MgH2 didn't react with Ag(Zn) solid solution during ball milling. The broadening of diffraction peaks suggests the refining of grain size of MgH2 and Ag(Zn) by milling. Fig. S1a shows the back-scattering electron (BSE) image of the ball-milled Mg90Ag7.5Zn2.5, and the gray and bright grains should correspond to MgH2 and Ag(Zn) solid solution, respectively, according to the above XRD result. The EDS result of the ball-milled Mg90Ag7.5Zn2.5 is shown in Fig. S1b, which
2. Experimental section The raw materials MgH2 powder (purity of 96.5%) and Ag powder (purity of 99.9%) were purchased from Sigma-Aldrich, and Zn powder (purity of 99.9%) was bought form Alfa-Aesar. The Mg−Ag−Zn ternary alloys were prepared by a two-step method. First, the Ag(Zn) solid solution was obtained by sintering mixture of Ag and Zn powders at 500 °C for 6 h in a tube furnace under Ar atmosphere. Then, the Ag (Zn) solid solution was milled with MgH2 in an ultrahigh-energy−highpressure planetary mill at room temperature. The ball-to-powder mass ratio was 100:1, and the milling time was 4 h. As Mg is easy to agglomerate and adhere to milling balls and jar in the milling, the brittle MgH2 was chosen as the raw material in this work. Thus, the atom ratio of Mg−Ag−Zn was used to represent the composition of MgH2−Ag (Zn) composite hereafter. Moreover, to avoid the decomposition of 797
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Fig. 1. XRD patterns of Ag and Zn powder mixture before sintering (a) and sintered at 500 °C for 6 h (b), and Mg90Ag7.5Zn2.5 ternary alloy prepared by ball milling MgH2 with Ag(Zn) solid solution under a 10 MPa hydrogen pressure for 4 h at room temperature (c).
proves that the average composition of the as-prepared sample is close to the nominal composition. Fig. 2 shows the PCI curves and van't Hoff plots for hydrogen desorption of Mg90Ag7.5Zn2.5 ternary alloy. As seen from Fig. 2a, the maximum hydrogen storage capacity of Mg90Ag7.5Zn2.5 is approximately 4.2 wt%. In addition to a flat plateau, a sloping plateau is evident in each PCI curve, suggesting two-step dehydrogenation for the Mg90Ag7.5Zn2.5 alloy. The fitted van't Hoff plots are shown in Fig. 2b, by which the hydrogen desorption enthalpy ΔH and entropy ΔS for the two reactions were calculated. The ΔH and ΔS for the sloping plateau region were 76.9 ± 0.3 kJ mol−1 H2 and 142.7 ± 0.4 J K−1 mol−1 H2, respectively. Accordingly, the dehydriding equilibrium pressure in the sloping plateau region at 300 °C was calculated to be 0.28 MPa, higher than those of pure Mg (0.16 MPa) [33], Mg3Ag binary alloy (0.17 MPa) [29], and the first step dehydrogenation of Mg85Ag5Al10 ternary alloy (0.26 MPa) [35]. This indicates more intensive destabilization of dehydriding thermodynamics in the Mg−Ag−Zn alloy comparing with those systems. The ΔH and ΔS for the flat plateau region were determined to be 80.2 ± 0.3 kJ mol−1 H2 and 144.2 ± 0.5 J K−1 mol−1 H2, respectively, and the dehydriding equilibrium pressure in the flat plateau region at 300 °C was 0.16 MPa, which is identical to that of pure Mg [33]. To reveal the two-step dehydriding mechanism of Mg90Ag7.5Zn2.5 alloy, quasi in situ XRD analysis was carried out at different dehydriding stages of hydrogenated sample. Fig. 3 shows the XRD results, in which the patterns a, b, c, d and e correspond to the points a, b, c, d and e, marked on the hydrogen desorption PCI curve measured at 320 °C (Fig. 2a). According to the XRD results, the lattice constants for phases in Mg90Ag7.5Zn2.5 alloy at different states were calculated and summarized in Table 1. Pattern a in Fig. 3 represents the fully hydrogenated Mg90Ag7.5Zn2.5. The hydrogenated products consist of MgH2, MgAg and MgZn1.8Ag0.2 phases. It is noted that the lattice constant of MgAg phase is almost equal to that of the MgAg reference pattern (MgAg: PDF#65–5904, a = b = c = 3.3106 Å). Therefore, the Zn element only exists in MgZn1.8Ag0.2 compound after hydrogenation. When the dehydrogenation progresses to point b (Fig. 2a), the midpoint of the sloping plateau region, both MgAg binary phase and MgZn1.8Ag0.2 ternary phase
Fig. 2. (a) Pressure-composition isotherms and (b) van't Hoff plots for hydrogen desorption of Mg90Ag7.5Zn2.5 ternary alloy.
Fig. 3. XRD patterns of Mg90Ag7.5Zn2.5 ternary alloy at different PCI desorption stages at 320 °C, as indicated in Fig. 2(a). 798
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Table 1 Lattice constants for phases in Mg90Ag7.5Zn2.5 ternary alloy at different states. Sample
Phase
Lattice constants a (Å)
As-milled Fully hydrogenated (pattern a) Partially dehydrogenated (pattern b) Partially dehydrogenated (pattern c) Partially dehydrogenated (pattern d) Fully dehydrogenated (pattern e)
b (Å)
c (Å)
MgH2 Ag(Zn) MgH2 MgAg MgZn1.8Ag0.2 MgH2 Mg54(Ag, Zn)17
4.5140 (2) 4.0265 (2) 4.5150 (3) 3.3110 (1) 5.2407 (1) 4.5144 (6) 14.2613 (4)
3.0190 (2)
14.1765 (2)
89.7011 (5) 3.0195 (4) 14.6402 (3)
MgH2 Mg54(Ag, Zn)17
4.5148 (2) 14.2628 (2)
14.1780 (3)
3.0200 (5) 14.6398 (5)
MgH2 Mg Mg54(Ag, Zn)17 Mg Mg54(Ag, Zn)17
4.5142 (2) 3.2093 (5) 14.2600 (1) 3.2090 (1) 14.2592 (5)
3.0192 (1)
14.1802 (1) 14.1743 (5)
3.0200 (3) 5.2089 (9) 14.6374 (2) 5.2107 (2) 14.6401 (7)
disappear, the MgH2 peaks become less intense, and a new solid solution Mg54(Ag, Zn)17 appear. The Mg54(Ag, Zn)17 solid solution is verified by comparing its unit cell volume (V = 2959.89 Å3) with that of the Mg54Ag17 reference pattern (Mg54Ag17: PDF#65–8314, V = 2966.86 Å3), and the decreasing of cell volume means the dissolving of Zn in Mg54Ag17. At point c, the end point of the sloping plateau region, the MgH2 peaks are further weakened, while the Mg54(Ag, Zn)17 peaks intensify. Therefore, for the dehydrogenation of Mg90Ag7.5Zn2.5 in the high pressure region, a part of MgH2 reacts with MgAg and MgZn1.8Ag0.2, leading to the formation of Mg54(Ag, Zn)17 solid solution and the release of hydrogen. The dehydriding reaction can be described as follows. x MgH2 + (17-2x) MgAg + (37 + x) MgZn1.8Ag0.2 ↔ Mg54(Ag, Zn)17 + x H2 (1) When the dehydrogenation proceeds to point d, the midpoint of the flat plateau region, the Mg54(Ag, Zn)17 peaks remain almost unchanged, but the peak intensities of the MgH2 phase decrease obviously, and meanwhile the Mg phase is detected. After complete dehydrogenation, as shown in pattern e, MgH2 phase disappears, and the peak intensities of the Mg phase reach the maximum values, but the Mg54(Ag, Zn)17 phase is still stable. Thus, it can be confirmed that the decomposition of pure MgH2 occurs in the low pressure region. Although the original structure of the ball-milled Mg90Ag7.5Zn2.5 could not be recovered after the first hydrogen absorption and desorption cycle, the subsequent hydriding-dehydriding process was found to be reversible. Therefore, a new reversible Mg-based hydrogen storage alloy is discovered in Mg−Ag−Zn ternary alloy system.
Fig. 5. (a) Pressure-composition isotherms for hydrogen desorption of Mg78Ag16.5Zn5.5 ternary alloy, and (b) van't Hoff plots for Mg78Ag16.5Zn5.5 and Mg90Ag7.5Zn2.5 (sloping plateau region).
The reversible microstructure evolution of Mg90Ag7.5Zn2.5 was also investigated by SEM observation. Fig. 4a is the BSE image of the fully hydrogenated Mg90Ag7.5Zn2.5, in which different contrasts can be seen. On the basis of the XRD results, the gray matrix in this image corresponds to MgH2, while the bright grains containing a high content of
Fig. 4. Back-scattering electron images of Mg90Ag7.5Zn2.5 ternary alloy: (a) hydrogenated under 5 MPa H2 at 320 °C for 12 h, and (b) dehydrogenated at 320 °C for 12 h. 799
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indexed in the electron diffraction pattern because of the relatively low weight percent, but it was detected in the high-resolution TEM image. After dehydrogenation, as shown in Fig. 7b, Mg54(Ag, Zn)17 solid solution and Mg are formed, which are consistent with the XRD result.
the heavy element Ag are distributed in MgH2, however, they cannot be precisely assigned to MgAg or MgZn1.8Ag0.2 through SEM because of their fine particle sizes. The fully dehydrogenated sample, as shown in Fig. 4b, consists of bright particles Mg54(Ag, Zn)17 solid solution dispersed in gray Mg matrix. Obviously, thermodynamic destabilization is only achieved in the first step dehydriding reaction of Mg90Ag7.5Zn2.5. To obtain the alloy exhibiting an intact thermodynamic destabilized dehydriding reaction, the Mg78Ag16.5Zn5.5 ternary alloy was designed according to the reaction equation (1) and the hydrogen storage capacity of Mg90Ag7.5Zn2.5 in the first step dehydrogenation. The BSE image and EDS result of the ball-milled Mg78Ag16.5Zn5.5 are shown in Fig. S2. Fig. 5a shows the hydrogen desorption PCI curves of Mg78Ag16.5Zn5.5 alloy at different temperatures. Only one sloping plateau can be observed in each PCI curve, and the maximum hydrogen storage capacity is nearly 3.0 wt%, a little higher than that corresponding to the sloping plateau region in the PCI curve of Mg90Ag7.5Zn2.5 (Fig. 2a). The ΔH and ΔS for the dehydrogenation of Mg78Ag16.5Zn5.5 were determined to be 77.2 ± 0.7 kJ mol−1 H2 and 143.4 ± 0.9 J K−1 mol−1 H2, respectively, corresponding to dehydriding equilibrium pressure of 0.28 MPa at 300 °C, the same as that of Mg90Ag7.5Zn2.5 in the first step dehydrogenation. The fitted van't Hoff plot of Mg78Ag16.5Zn5.5, as shown in Fig. 5b, is also coincident to that of the first step dehydrogenation of Mg90Ag7.5Zn2.5. These results demonstrate that the Mg78Ag16.5Zn5.5 alloy realize an intact destabilized dehydriding reaction occurring in the high pressure region of Mg90Ag7.5Zn2.5. Fig. 6 shows the XRD patterns of Mg78Ag16.5Zn5.5 ternary alloy at different states. It can be seen that the phase constituents of the Mg78Ag16.5Zn5.5 alloy at different states were similar to those for Mg90Ag7.5Zn2.5, demonstrating again that the intact reversible thermodynamic destabilized dehydriding reaction occurring in the high pressure region of Mg90Ag7.5Zn2.5 is obtained in the Mg78Ag16.5Zn5.5 alloy. Fig. 7a shows the TEM bright field image, selected area electron diffraction pattern, and high-resolution TEM image of hydrogenated Mg78Ag16.5Zn5.5 alloy. According to the electron diffraction pattern, the dark grains in the bright field image are assigned to MgAg, which are dispersed in the bright MgH2 matrix. The MgZn1.8Ag0.2 phase was not
3.2. Enhanced dehydrogenation kinetics of Mg−Ag−Zn ternary alloys The dehydriding kinetics of those DTRIF systems is generally poor [35]. Thus, the hydrogen desorption kinetics of the DTRIF Mg−Ag−Zn ternary alloys is also a concern. The isothermal dehydrogenation kinetic curves of the Mg90Ag7.5Zn2.5 alloy are shown in Fig. 8, in which the two-step dehydriding process presented in the PCI curves could not be distinguished. The reason may be the two hydrogen desorption reactions occurred simultaneously when the initial pressure for the isothermal kinetics assessment was close to 0. At 320 °C, the alloy exhibits good hydrogen desorption kinetics, releasing hydrogen completely within 10 min. Even at a lower temperature of 240 °C, approximately 1.3 wt% H2 can be desorbed within 120 min. On the basis of the dehydriding kinetic curves, the apparent activation energy Ea for the overall dehydrogenation of Mg90Ag7.5Zn2.5 was determined to be 118.7 ± 2.0 kJ mol−1, which is lower than 1600 kJ mol−1 for pure Mg [39,40] and 124.8 kJ mol−1 for the Mg85Ag5Al10 ternary alloy [35], but higher than 89.8 kJ mol−1 for the Mg6Ag binary alloy [34]. As shown in the SEM image (Fig. 4a), the MgAg and MgZn1.8Ag0.2 are highly dispersed in the MgH2 matrix. The abundant phase boundaries exist in the hydrogenated Mg−Ag−Zn alloy can promote the hydrogen desorption kinetics by providing more hydrogen diffusion channels. Therefore, the decrease of Ea for the dehydrogenation is obtained in the Mg−Ag−Zn alloy. For the Mg78Ag16.5Zn5.5 alloy, it takes 5 min to desorb the entire hydrogen at 320 °C, which is faster than that of Mg90Ag7.5Zn2.5 at the same temperature. The calculated apparent activation energy Ea for the dehydrogenation of Mg78Ag16.5Zn5.5 was 116.5 ± 2.3 kJ mol−1, even lower than that for Mg90Ag7.5Zn2.5. Therefore, the Mg78Ag16.5Zn5.5 exhibits not only intact destabilized dehydriding thermodynamics, but also more excellent dehydriding kinetics. 3.3. Guidelines for designing of Mg-based alloy systems with destabilized dehydriding thermodynamics by reversible intermetallics formation So far, the thermodynamic destabilization of MgH2 has been achieved in some Mg-based alloy systems by reversible intermetallics formation. The reported systems include Mg−In [22,23], Mg−Cd [24,25], Mg − Ga [26], and Mg−Ag [27–29] binary alloys, Mg(In, Al) and Mg(In, Cd) ternary solid solutions [30,32], Mg−In−Ni [33], Mg−Ag−In [34], Mg−Ag−Al [35], and Mg−Ag−Zn ternary alloy systems. However, there is no clear principle for designing of Mg-based alloy systems with destabilized dehydriding thermodynamics by reversible intermetallics formation. Table 2 summarizes the hydrogen storage capacity, dehydriding thermodynamic parameters and apparent activation energy for the dehydrogenation of pure Mg and DTRIF systems. It should be noted that all of the listed alloying elements (In, Ag, Al, Cd, Ni, and Zn) can form intermetallic compounds with Mg upon hydrogenation, such as MgIn, MgAg and Mg4Al6Ag. Furthermore, the formation enthalpies of these compounds are low, ensuring that they can play as the reactive additives for MgH2 to release hydrogen. Although other elements such as Sn, Si, and Pb can form Mg-based compounds Mg2Si, Mg2Sn, and Mg2Pb [41,42], they are too stable to react with MgH2. Besides, alloying elements are readily expelled from the compounds upon MgH2 formation and may react with H to form elemental hydrides. However, most of the hydride mixtures are not reversible. The thermal decomposition temperatures of some hydrides are listed in Table 3. It clearly shows that the thermal decomposition temperatures of the hydrides of In, Ag, Cd, Al and Zn are much lower than that of MgH2, thus inhibiting the disproportionation of Mg-based alloys and promoting the reversibility during hydrogenation and
Fig. 6. XRD patterns of Mg78Ag16.5Zn5.5 ternary alloy at different states: (a) ball milled under 10 MPa H2 for 4 h, (b) hydrogenated under 5 MPa H2 at 320 °C for 12 h, and (c) dehydrogenated at 320 °C for 12 h. 800
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Fig. 7. TEM bright field images, electron diffraction patterns, and high-resolution TEM images of Mg78Ag16.5Zn5.5 ternary alloy at different states: (a) hydrogenated and (b) dehydrogenated. Table 3 The thermal decomposition temperatures (Tdec) of related hydrides (some of them are estimated values) [44].
pure Mg [33] Mg90In10 [30] Mg3Ag [29] Mg90In5Al5 [30] Mg90In5Cd5 [32] Mg14In3Ni3 [33] Mg85Ag5Al10 [35] Mg78Ag16.5Zn5.5
ΔS (J K−1 mol−1 H2)
Peq at 300 °C (MPa)
Ea (kJ mol−1)
7.0 4.2 2.0 4.8 4.3 1.8 3.8 2.5
72.2 65.2 82.3 66.3 86.0 70.1 80.3 77.2
129.8 121.8 147.7 121.2 154.8 132.2 148.2 143.4
0.16 0.27 0.17 0.20 0.18 0.33 0.26 0.28
160 146.0 – 156.0 170.1 78.5 124.8 116.5
InH3
AgH
CdH2
AlH3
ZnH2
CaH2
TiH2
YH2
Tdec (°C)
327
15
> −100
−20
150
90
600
380
510
4. Conclusions
Table 2 Hydrogen storage capacity, dehydrogenation enthalpy (ΔH) and entropy (ΔS), dehydriding equilibrium pressure (Peq) at 300 °C and apparent activation energy for the dehydrogenation (Ea) of pure Mg and DTRIF systems. ΔH (kJ mol−1 H2)
MgH2
dehydrogenation. In this sense, Ca, Ti and Y are not proper, because their hydrides are more stable than MgH2 [43], which need higher thermal decomposition temperatures in comparison with pure MgH2. Based on the sum up of the systems already known, we proposed two guidelines for designing the DTRIF systems. First, the alloying elements can react with Mg to form intermetallic compounds, but the reaction enthalpies should be low. Second, the hydrides of the alloying elements should have low stabilities. It is believed that there is still chance to explore new DTRIF systems according to the above guidelines. However, it should be noted that capacity and kinetics are usually influenced by the alloying addition. Therefore, the selection of alloying elements for Mg-based hydrogen storage alloys needs more consideration and exploration.
Fig. 8. Isothermal dehydrogenation kinetic curves for Mg90Ag7.5Zn2.5 ternary alloys.
Capacity (wt.%)
Hydrides
Mg−Ag−Zn ternary alloys were prepared by ball milling MgH2 with Ag(Zn) solid solution to create a DTRIF system. A two-step hydrogen desorption reactions occurred in Mg90Ag7.5Zn2.5 alloy with the maximum hydrogen storage capacity of 4.2 wt%. In the first step, MgH2 reacts with MgAg and MgZn1.8Ag0.2 to release hydrogen, with an elevated dehydriding equilibrium pressure of 0.28 MPa at 300 °C. This thermodynamic destabilized dehydriding reaction can be intact realized in Mg78Ag16.5Zn5.5 alloy with a maximum hydrogen storage capacity of 3.0 wt%. The improved dehydriding kinetics was also achieved in the Mg−Ag−Zn ternary alloys, and the apparent activation energies for the dehydrogenation of Mg90Ag7.5Zn2.5 and Mg78Ag16.5Zn5.5 were lowered to 118.7 and 116.5 kJ mol−1, respectively. This is due to the abundant phase boundaries exist in the hydrogenated Mg−Ag−Zn alloys, providing more hydrogen diffusion channels. By sum up the already reported DTRIF systems, two guidelines were proposed for designing the DTRIF systems: The alloying elements can react with Mg to form intermetallic compounds with low reaction enthalpies, and the 801
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hydrides of the alloying elements should have low stabilities.
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Acknowledgements This work was financially supported by the Foundation for Innovative Research Groups of the National Natural Science Foundation of China (No. 51621001), by the National Natural Science Foundation of China (Nos. 51471070, 51431001, and 5157010144), by the Guangdong Natural Science Foundation (2014A030313222 and 2016A030312011). The author Yanshan Lu also thanks the Oversea Study Program of Guangzhou Elite Project (No. JY201508). Appendix A. Supplementary data Supplementary data related to this article can be found at http://dx. doi.org/10.1016/j.jpowsour.2018.06.060. References [1] J.C. Crivello, B. Dam, R.V. Denys, M. Dornheim, D.M. Grant, J. Huot, T.R. Jensen, P.D. Jongh, M. Latroche, C. Milanese, Appl. Phys. A 122 (2016) 1–20. [2] H. Wang, H.J. Lin, W.T. Cai, L.Z. Ouyang, M. Zhu, J. Alloy. Comp. 658 (2016) 280–300. [3] K. Bohmhammel, U. Wolf, G. Wolf, E. Königsberger, Thermochim. Acta 337 (1999) 195–199. [4] M. Zhu, Y.S. Lu, L.Z. Ouyang, H. Wang, Materials 6 (2013) 4654–4674. [5] I.P. Jain, C. Lal, A. Jain, Int. J. Hydrogen Energy 35 (2010) 5133–5144. [6] X. Yao, G. Lu, Chin. Sci. Bull. 53 (2008) 2421–2431. [7] X. Tan, L. Wang, C.M.B. Holt, B. Zahiri, M.H. Eikerling, D. Mitlin, Phys. Chem. Chem. Phys. 14 (2012) 10904–10909. [8] A. Anastasopol, T.V. Pfeiffer, J. Middelkoop, U. Lafont, R.J. Canales-Perez, A. Schmidt-Ott, F.M. Mulder, S.W.H. Eijt, J. Am. Chem. Soc. 135 (2013) 7891–7900. [9] T. Takasaki, D. Kyoi, N. Kitamura, S. Tanase, T. Sakai, J. Phys. Chem. B 111 (2007) 14102–14106. [10] J.J. Reilly, R.H. Wiswall, Inorg. Chem. 7 (1968) 2254–2256. [11] R.A. Varin, T. Czujko, C. Chiu, Z. Wronski, J. Alloy. Comp. 424 (2006) 356–364. [12] Z. Zhao-Karger, J. Hu, A. Roth, D. Wang, C. Kubel, W. Lohstroh, M. Fichtner, Chem. Commun. 46 (2010) 8353–8355. [13] L.Z. Ouyang, S.Y. Ye, H.W. Dong, M. Zhu, Appl. Phys. Lett. 90 (2007) 0219171–021917-3. [14] W. Li, C. Li, H. Ma, J. Chen, J. Am. Chem. Soc. 129 (2007) 6710–6711. [15] M.S.L. Hudson, K. Takahashi, A. Ramesh, S. Awasthi, A.K. Ghosh, P. Ravindran, O.N. Srivastava, Catal. Sci. Technol. 6 (2016) 261–268.
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Destabilizing the dehydriding thermodynamics of MgH2 by reversible intermetallics formation in Mg−Ag−Zn ternary alloys
T
Yanshan Lu, Hui Wang, Jiangwen Liu, Liuzhang Ouyang, Min Zhu∗ School of Materials Science and Engineering and Guangdong Provincial Key Laboratory of Advanced Energy Storage Materials, South China University of Technology, Guangzhou, 510641, China
H I GH L IG H T S
Mg−Ag−Zn ternary alloys show fully reversible de-/hydrogenation performance. • The Mg−Ag−Zn ternary alloys exhibit elevated equilibrium pressure for MgH . • The • Enhanced dehydrogenation kinetics is achieved in the Mg−Ag−Zn ternary alloys. 2
A R T I C LE I N FO
A B S T R A C T
Keywords: Hydrogen storage Thermodynamic destabilization Mg−Ag−Zn Reversibility
In this study, Mg−Ag−Zn ternary alloys are employed as carriers for hydrogen storage, which show improved thermodynamic and kinetic properties compared with pure Mg. Mg90Ag7.5Zn2.5 alloy exhibits a reversible phase transformation during hydriding-dehydriding process with a maximum hydrogen storage capacity of approximately 4.2 wt%. Two-step reactions occur in the dehydriding process of Mg90Ag7.5Zn2.5. In the first step, a portion of MgH2 reacts with MgAg and MgZn1.8Ag0.2 phases to transform to Mg54(Ag, Zn)17 solid solution and release hydrogen. After that, the remaining MgH2 decomposes to Mg. Because of the modified reaction pathway in the first step dehydrogenation, the dehydriding equilibrium pressure at 300 °C increases to 0.28 MPa. Besides, the apparent activation energy (Ea) for dehydrogenation of the Mg90Ag7.5Zn2.5 alloy is lowered to 118.7 kJ mol−1. By fully utilizing the first step reaction, another alloy Mg78Ag16.5Zn5.5 is designed, which presents intact destabilized dehydriding thermodynamics and a lower dehydrogenation Ea of 116.5 kJ mol−1. Our work demonstrates that alloying Mg with elements that can form compounds with low formation enthalpies, and the hydrides of the elements have low stabilities, are the main guidelines for the designing of the Mg-based hydrogen storage alloys with destabilized dehydriding thermodynamics by reversible intermetallics formation.
1. Introduction Mg is a very attractive material for hydrogen storage owing to its reversible reaction with hydrogen to form MgH2 with H capacity of 7.6 wt% [1,2]. However, the strong ionic characteristics of the Mg−H bond, leads to too high dehydriding enthalpy of MgH2, being 74.7 kJ mol−1 H2 [3], which is much larger than that required for dehydriding at room temperature, being 20–40 kJ mol−1 H2 [4]. Thus, the dehydriding temperature of MgH2 is too high to meet the practical application such as hydrogen-fueled vehicles [5,6]. Many attempts have been made to destabilize the thermodynamics of MgH2; however, the destabilization is far from enough in addition to the inducing of other drawbacks. For example, the dehydriding enthalpy can be reduced by alloying [7–9], however, the hydrogen
∗
Corresponding author. E-mail address: [email protected] (M. Zhu).
https://doi.org/10.1016/j.jpowsour.2018.06.060 Received 28 April 2018; Received in revised form 8 June 2018; Accepted 12 June 2018 0378-7753/ © 2018 Elsevier B.V. All rights reserved.
storage capacity also reduces, being 64.5 kJ mol−1 H2 and 3.6 wt%, respectively, in the case of forming Mg2Ni [10]. Although nanostructuring Mg/MgH2 can reduce the dehydriding enthalpy [11–13], the effect is significant only when the size reduces to several nanometer, which causes problems in fabrication of materials, stability of nanostructure and so on. For instance, the decomposition enthalpy of Mg nanowires with diameter of 30–50 nm reduced to 65.3 kJ mol−1 H2, but the Mg nanowires could not keep their structure after 10 hydrogenation and dehydrogenation cycles [14]. Recently, catalyzed MgH2 with a lower dehydriding enthalpy has been reported [15,16]. It is found that the dehydrogenation enthalpy was reduced to 66.34 kJ mol−1 H2 for MgH2 catalyzed by graphene sheet templated Fe3O4 nanoparticles, however, a decrease of the dehydrogenation entropy (ΔS = 125.85 J mol−1 H2) also happens [16], which has a negative
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MgH2 during milling, a hydrogen pressure of 10 MPa was filled in the milling jar. X-ray diffraction (XRD) analysis was conducted on a PANalytical Empyrean X-ray diffractometer equipped with Cu Kα radiation (λ = 1.54060 Å). High purity Si powder was added into XRD samples as an internal reference to correct the instrumental zero-shift. To examine the reaction process, the PCI desorption test was manually terminated at different stages, and the sample holder was quickly cooled while the hydrogen pressure of sample cell was kept. Then the sample was taken out for XRD analysis at room temperature. The lattice constants and the unit cell volume were calculated from the XRD data by Rietveld method using Highscore Plus software. The microstructure of samples at different states was observed by a scanning electron microscope (SEM, Zeiss Merlin) equipped with an energy-dispersive spectrometer (EDS), and a JEOL JEM-2100 transmission electron microscopy (TEM) operating at 200 kV. The TEM samples were prepared by dispersing powder sample in hexane by an ultrasonic cleaner, and then dropped the suspension onto a copper grid. The hydrogen storage properties were measured using a Sieverttype apparatus (SETARAM PCT-Pro2000). Prior to the measurements, the samples were activated through three hydriding-dehydriding cycles at 320 °C under a hydrogen pressure of 4.5 MPa. PCI test was conducted with the hydrogen pressures varied from 4.5 to 0.002 MPa. The dehydrogenation enthalpy ΔH and entropy ΔS were calculated based on the van't Hoff equation by adopting the pressure at the midpoint of the plateau in each PCI curve as the equilibrium pressure. Isothermal hydrogen desorption kinetics was determined at specific temperatures, and the sample cell was pumped to vacuum before testing. The kinetic curves were analyzed using the Johnson−Mehl−Avrami−Kolmogorov η (JMAK) model [37,38], expressed as α= 1 − e−(kt) , where α is the reaction fraction, k is the rate constant, t is the reaction time, and η is the reaction order. Then the apparent activation energy Ea for dehydrogenation was calculated according to the Arrhenius equation: k= Ae−Ea/RT , where A is the pre-exponential factor, R is the gas constant, and T is the Kelvin temperature. All the sample handlings were carried out in a glove box filled with Ar with water and oxygen content less than 3 ppm.
influence on the thermodynamic destabilization of MgH2. Changing the reaction pathway by adding a reactive additive is another attractive route of destabilizing thermodynamics of MgH2 and first proposed by Vajo et al. [17,18]. They found Mg2Si would form upon dehydrogenation when adding Si into MgH2, decreasing reaction enthalpy to 36.4 kJ mol−1, and increasing the dehydriding pressure to 7.5 bar at 300 °C. The thermodynamic destabilization also achieved in MgH2/Ge system with the dehydrogenation temperature decreasing to 130 °C [19]. It is reported that Sn is another additive to destabilize the thermodynamics of MgH2 effectively [20,21]. Unfortunately, the reversibility of those systems is poor. To overcome this problem, a new route, which can be briefly expressed as destabilizing thermodynamics by reversible intermetallics formation (DTRIF), was developed in Mg−In system [22,23]. In this system, dehydriding enthalpy of MgH2 was reduced by reacting with the β intermetallic formed in the hydrogenation of Mg(In) solid solution, and the reaction is highly reversible. Following the DTRIF route, a series of Mg-based alloys have been explored and realized reversible hydrogen storage reactions with destabilized thermodynamics, including Mg–Cd [24,25], Mg–Ga [26], Mg–Ag [27–29], Mg–In–Al [30,31], Mg–In–Cd [32], Mg–In–Ni [33], Mg–Ag–In [34] and Mg–Ag–Al [35] alloy systems. It is obvious that most of those DTRIF systems contain In and Ga, which are expensive rare metals. Besides, the reduction of dehydriding enthalpy of those DTRIF systems is not large enough and most of them suffer from slow hydrogen desorption kinetics, in particular, Mg–In–Al [30,31], Mg–In–Cd [32] and Mg–Ag–Al [35] alloys. Therefore, it is of great significant to explore other DTRIF systems with a reduced amount of expensive elements, and to have larger thermodynamic destabilization and enhanced kinetics, i.e., dual tuning of thermodynamics and kinetics of MgH2. According to the Ag−Zn binary phase diagram, Zn can dissolve in Ag to form a Ag(Zn) solid solution with the maximum solubility of 30 at % at 100 °C. Thus, it is expected that intermetallic may formed in Mg−Ag−Zn ternary alloy, which may lead to a new DTRIF system. In addition, Zn has a low melting point of 420 °C, and thus Zn atoms may diffuse easily during hydrogenation and dehydrogenation at lower temperatures, which would possibly promote the hydrogen sorption kinetics. Indeed, it has been shown that the ultrafine Mg−Zn particles, being 400 nm and consisting of Mg(Zn) solid solution and amorphous, could enhance hydrogen absorption kinetics and decrease the activation energy for hydrogenation to 56.3 kJ mol−1 [36]. In the present work, with the aiming of the dual tuning of thermodynamics and kinetics of MgH2, we explored the DTRIF alloys in Mg−Ag−Zn ternary alloy system. The dehydriding thermodynamic destabilization has been achieved and the hydrogen desorption pressure at 300 °C increased from 0.16 to 0.28 MPa in comparison with pure MgH2. The hydrogen desorption kinetics are also improved with activation energy lowering to 116.5 kJ mol−1. Besides, we proposed the main guidelines for designing DTRIF systems based on the sum up of the systems already known from the literature.
3. Results and discussion 3.1. Achieving a sole reversible thermodynamic destabilized dehydriding reaction in Mg−Ag−Zn ternary alloy Compared with the catalyzed Mg/MgH2 [15,16], the hydrogen storage capacities of the DTRIF systems are usually lower, because a part of Mg would form Mg-based compounds after hydrogen absorption. Therefore, the alloy composition was designed as Mg90Ag7.5Zn2.5 with the consideration that the theoretical hydrogen storage capacity of the alloy can exceed 5 wt% by counting the Mg content. Fig. 1a and b shows the XRD patterns of the powder mixture of Ag and Zn before and after sintering, respectively. As it shown, after sintering at 500 °C for 6 h, the diffraction peaks of Zn disappear, while the Ag peaks are obviously shifted toward higher 2θ angle in comparison with the powder mixture of Ag and Zn. The lattice constant of Ag decreases from 4.0837 Å to 4.0262 Å, indicating the formation of Ag(Zn) solid solution after sintering. Fig. 1c is the XRD pattern of Mg90Ag7.5Zn2.5 ternary alloy prepared by ball milling Ag(Zn) solid solution with MgH2 under a 10 MPa hydrogen pressure. The observed phases are MgH2 and Ag(Zn) solid solution, and no other phase is detected, which means that MgH2 didn't react with Ag(Zn) solid solution during ball milling. The broadening of diffraction peaks suggests the refining of grain size of MgH2 and Ag(Zn) by milling. Fig. S1a shows the back-scattering electron (BSE) image of the ball-milled Mg90Ag7.5Zn2.5, and the gray and bright grains should correspond to MgH2 and Ag(Zn) solid solution, respectively, according to the above XRD result. The EDS result of the ball-milled Mg90Ag7.5Zn2.5 is shown in Fig. S1b, which
2. Experimental section The raw materials MgH2 powder (purity of 96.5%) and Ag powder (purity of 99.9%) were purchased from Sigma-Aldrich, and Zn powder (purity of 99.9%) was bought form Alfa-Aesar. The Mg−Ag−Zn ternary alloys were prepared by a two-step method. First, the Ag(Zn) solid solution was obtained by sintering mixture of Ag and Zn powders at 500 °C for 6 h in a tube furnace under Ar atmosphere. Then, the Ag (Zn) solid solution was milled with MgH2 in an ultrahigh-energy−highpressure planetary mill at room temperature. The ball-to-powder mass ratio was 100:1, and the milling time was 4 h. As Mg is easy to agglomerate and adhere to milling balls and jar in the milling, the brittle MgH2 was chosen as the raw material in this work. Thus, the atom ratio of Mg−Ag−Zn was used to represent the composition of MgH2−Ag (Zn) composite hereafter. Moreover, to avoid the decomposition of 797
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Fig. 1. XRD patterns of Ag and Zn powder mixture before sintering (a) and sintered at 500 °C for 6 h (b), and Mg90Ag7.5Zn2.5 ternary alloy prepared by ball milling MgH2 with Ag(Zn) solid solution under a 10 MPa hydrogen pressure for 4 h at room temperature (c).
proves that the average composition of the as-prepared sample is close to the nominal composition. Fig. 2 shows the PCI curves and van't Hoff plots for hydrogen desorption of Mg90Ag7.5Zn2.5 ternary alloy. As seen from Fig. 2a, the maximum hydrogen storage capacity of Mg90Ag7.5Zn2.5 is approximately 4.2 wt%. In addition to a flat plateau, a sloping plateau is evident in each PCI curve, suggesting two-step dehydrogenation for the Mg90Ag7.5Zn2.5 alloy. The fitted van't Hoff plots are shown in Fig. 2b, by which the hydrogen desorption enthalpy ΔH and entropy ΔS for the two reactions were calculated. The ΔH and ΔS for the sloping plateau region were 76.9 ± 0.3 kJ mol−1 H2 and 142.7 ± 0.4 J K−1 mol−1 H2, respectively. Accordingly, the dehydriding equilibrium pressure in the sloping plateau region at 300 °C was calculated to be 0.28 MPa, higher than those of pure Mg (0.16 MPa) [33], Mg3Ag binary alloy (0.17 MPa) [29], and the first step dehydrogenation of Mg85Ag5Al10 ternary alloy (0.26 MPa) [35]. This indicates more intensive destabilization of dehydriding thermodynamics in the Mg−Ag−Zn alloy comparing with those systems. The ΔH and ΔS for the flat plateau region were determined to be 80.2 ± 0.3 kJ mol−1 H2 and 144.2 ± 0.5 J K−1 mol−1 H2, respectively, and the dehydriding equilibrium pressure in the flat plateau region at 300 °C was 0.16 MPa, which is identical to that of pure Mg [33]. To reveal the two-step dehydriding mechanism of Mg90Ag7.5Zn2.5 alloy, quasi in situ XRD analysis was carried out at different dehydriding stages of hydrogenated sample. Fig. 3 shows the XRD results, in which the patterns a, b, c, d and e correspond to the points a, b, c, d and e, marked on the hydrogen desorption PCI curve measured at 320 °C (Fig. 2a). According to the XRD results, the lattice constants for phases in Mg90Ag7.5Zn2.5 alloy at different states were calculated and summarized in Table 1. Pattern a in Fig. 3 represents the fully hydrogenated Mg90Ag7.5Zn2.5. The hydrogenated products consist of MgH2, MgAg and MgZn1.8Ag0.2 phases. It is noted that the lattice constant of MgAg phase is almost equal to that of the MgAg reference pattern (MgAg: PDF#65–5904, a = b = c = 3.3106 Å). Therefore, the Zn element only exists in MgZn1.8Ag0.2 compound after hydrogenation. When the dehydrogenation progresses to point b (Fig. 2a), the midpoint of the sloping plateau region, both MgAg binary phase and MgZn1.8Ag0.2 ternary phase
Fig. 2. (a) Pressure-composition isotherms and (b) van't Hoff plots for hydrogen desorption of Mg90Ag7.5Zn2.5 ternary alloy.
Fig. 3. XRD patterns of Mg90Ag7.5Zn2.5 ternary alloy at different PCI desorption stages at 320 °C, as indicated in Fig. 2(a). 798
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Table 1 Lattice constants for phases in Mg90Ag7.5Zn2.5 ternary alloy at different states. Sample
Phase
Lattice constants a (Å)
As-milled Fully hydrogenated (pattern a) Partially dehydrogenated (pattern b) Partially dehydrogenated (pattern c) Partially dehydrogenated (pattern d) Fully dehydrogenated (pattern e)
b (Å)
c (Å)
MgH2 Ag(Zn) MgH2 MgAg MgZn1.8Ag0.2 MgH2 Mg54(Ag, Zn)17
4.5140 (2) 4.0265 (2) 4.5150 (3) 3.3110 (1) 5.2407 (1) 4.5144 (6) 14.2613 (4)
3.0190 (2)
14.1765 (2)
89.7011 (5) 3.0195 (4) 14.6402 (3)
MgH2 Mg54(Ag, Zn)17
4.5148 (2) 14.2628 (2)
14.1780 (3)
3.0200 (5) 14.6398 (5)
MgH2 Mg Mg54(Ag, Zn)17 Mg Mg54(Ag, Zn)17
4.5142 (2) 3.2093 (5) 14.2600 (1) 3.2090 (1) 14.2592 (5)
3.0192 (1)
14.1802 (1) 14.1743 (5)
3.0200 (3) 5.2089 (9) 14.6374 (2) 5.2107 (2) 14.6401 (7)
disappear, the MgH2 peaks become less intense, and a new solid solution Mg54(Ag, Zn)17 appear. The Mg54(Ag, Zn)17 solid solution is verified by comparing its unit cell volume (V = 2959.89 Å3) with that of the Mg54Ag17 reference pattern (Mg54Ag17: PDF#65–8314, V = 2966.86 Å3), and the decreasing of cell volume means the dissolving of Zn in Mg54Ag17. At point c, the end point of the sloping plateau region, the MgH2 peaks are further weakened, while the Mg54(Ag, Zn)17 peaks intensify. Therefore, for the dehydrogenation of Mg90Ag7.5Zn2.5 in the high pressure region, a part of MgH2 reacts with MgAg and MgZn1.8Ag0.2, leading to the formation of Mg54(Ag, Zn)17 solid solution and the release of hydrogen. The dehydriding reaction can be described as follows. x MgH2 + (17-2x) MgAg + (37 + x) MgZn1.8Ag0.2 ↔ Mg54(Ag, Zn)17 + x H2 (1) When the dehydrogenation proceeds to point d, the midpoint of the flat plateau region, the Mg54(Ag, Zn)17 peaks remain almost unchanged, but the peak intensities of the MgH2 phase decrease obviously, and meanwhile the Mg phase is detected. After complete dehydrogenation, as shown in pattern e, MgH2 phase disappears, and the peak intensities of the Mg phase reach the maximum values, but the Mg54(Ag, Zn)17 phase is still stable. Thus, it can be confirmed that the decomposition of pure MgH2 occurs in the low pressure region. Although the original structure of the ball-milled Mg90Ag7.5Zn2.5 could not be recovered after the first hydrogen absorption and desorption cycle, the subsequent hydriding-dehydriding process was found to be reversible. Therefore, a new reversible Mg-based hydrogen storage alloy is discovered in Mg−Ag−Zn ternary alloy system.
Fig. 5. (a) Pressure-composition isotherms for hydrogen desorption of Mg78Ag16.5Zn5.5 ternary alloy, and (b) van't Hoff plots for Mg78Ag16.5Zn5.5 and Mg90Ag7.5Zn2.5 (sloping plateau region).
The reversible microstructure evolution of Mg90Ag7.5Zn2.5 was also investigated by SEM observation. Fig. 4a is the BSE image of the fully hydrogenated Mg90Ag7.5Zn2.5, in which different contrasts can be seen. On the basis of the XRD results, the gray matrix in this image corresponds to MgH2, while the bright grains containing a high content of
Fig. 4. Back-scattering electron images of Mg90Ag7.5Zn2.5 ternary alloy: (a) hydrogenated under 5 MPa H2 at 320 °C for 12 h, and (b) dehydrogenated at 320 °C for 12 h. 799
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indexed in the electron diffraction pattern because of the relatively low weight percent, but it was detected in the high-resolution TEM image. After dehydrogenation, as shown in Fig. 7b, Mg54(Ag, Zn)17 solid solution and Mg are formed, which are consistent with the XRD result.
the heavy element Ag are distributed in MgH2, however, they cannot be precisely assigned to MgAg or MgZn1.8Ag0.2 through SEM because of their fine particle sizes. The fully dehydrogenated sample, as shown in Fig. 4b, consists of bright particles Mg54(Ag, Zn)17 solid solution dispersed in gray Mg matrix. Obviously, thermodynamic destabilization is only achieved in the first step dehydriding reaction of Mg90Ag7.5Zn2.5. To obtain the alloy exhibiting an intact thermodynamic destabilized dehydriding reaction, the Mg78Ag16.5Zn5.5 ternary alloy was designed according to the reaction equation (1) and the hydrogen storage capacity of Mg90Ag7.5Zn2.5 in the first step dehydrogenation. The BSE image and EDS result of the ball-milled Mg78Ag16.5Zn5.5 are shown in Fig. S2. Fig. 5a shows the hydrogen desorption PCI curves of Mg78Ag16.5Zn5.5 alloy at different temperatures. Only one sloping plateau can be observed in each PCI curve, and the maximum hydrogen storage capacity is nearly 3.0 wt%, a little higher than that corresponding to the sloping plateau region in the PCI curve of Mg90Ag7.5Zn2.5 (Fig. 2a). The ΔH and ΔS for the dehydrogenation of Mg78Ag16.5Zn5.5 were determined to be 77.2 ± 0.7 kJ mol−1 H2 and 143.4 ± 0.9 J K−1 mol−1 H2, respectively, corresponding to dehydriding equilibrium pressure of 0.28 MPa at 300 °C, the same as that of Mg90Ag7.5Zn2.5 in the first step dehydrogenation. The fitted van't Hoff plot of Mg78Ag16.5Zn5.5, as shown in Fig. 5b, is also coincident to that of the first step dehydrogenation of Mg90Ag7.5Zn2.5. These results demonstrate that the Mg78Ag16.5Zn5.5 alloy realize an intact destabilized dehydriding reaction occurring in the high pressure region of Mg90Ag7.5Zn2.5. Fig. 6 shows the XRD patterns of Mg78Ag16.5Zn5.5 ternary alloy at different states. It can be seen that the phase constituents of the Mg78Ag16.5Zn5.5 alloy at different states were similar to those for Mg90Ag7.5Zn2.5, demonstrating again that the intact reversible thermodynamic destabilized dehydriding reaction occurring in the high pressure region of Mg90Ag7.5Zn2.5 is obtained in the Mg78Ag16.5Zn5.5 alloy. Fig. 7a shows the TEM bright field image, selected area electron diffraction pattern, and high-resolution TEM image of hydrogenated Mg78Ag16.5Zn5.5 alloy. According to the electron diffraction pattern, the dark grains in the bright field image are assigned to MgAg, which are dispersed in the bright MgH2 matrix. The MgZn1.8Ag0.2 phase was not
3.2. Enhanced dehydrogenation kinetics of Mg−Ag−Zn ternary alloys The dehydriding kinetics of those DTRIF systems is generally poor [35]. Thus, the hydrogen desorption kinetics of the DTRIF Mg−Ag−Zn ternary alloys is also a concern. The isothermal dehydrogenation kinetic curves of the Mg90Ag7.5Zn2.5 alloy are shown in Fig. 8, in which the two-step dehydriding process presented in the PCI curves could not be distinguished. The reason may be the two hydrogen desorption reactions occurred simultaneously when the initial pressure for the isothermal kinetics assessment was close to 0. At 320 °C, the alloy exhibits good hydrogen desorption kinetics, releasing hydrogen completely within 10 min. Even at a lower temperature of 240 °C, approximately 1.3 wt% H2 can be desorbed within 120 min. On the basis of the dehydriding kinetic curves, the apparent activation energy Ea for the overall dehydrogenation of Mg90Ag7.5Zn2.5 was determined to be 118.7 ± 2.0 kJ mol−1, which is lower than 1600 kJ mol−1 for pure Mg [39,40] and 124.8 kJ mol−1 for the Mg85Ag5Al10 ternary alloy [35], but higher than 89.8 kJ mol−1 for the Mg6Ag binary alloy [34]. As shown in the SEM image (Fig. 4a), the MgAg and MgZn1.8Ag0.2 are highly dispersed in the MgH2 matrix. The abundant phase boundaries exist in the hydrogenated Mg−Ag−Zn alloy can promote the hydrogen desorption kinetics by providing more hydrogen diffusion channels. Therefore, the decrease of Ea for the dehydrogenation is obtained in the Mg−Ag−Zn alloy. For the Mg78Ag16.5Zn5.5 alloy, it takes 5 min to desorb the entire hydrogen at 320 °C, which is faster than that of Mg90Ag7.5Zn2.5 at the same temperature. The calculated apparent activation energy Ea for the dehydrogenation of Mg78Ag16.5Zn5.5 was 116.5 ± 2.3 kJ mol−1, even lower than that for Mg90Ag7.5Zn2.5. Therefore, the Mg78Ag16.5Zn5.5 exhibits not only intact destabilized dehydriding thermodynamics, but also more excellent dehydriding kinetics. 3.3. Guidelines for designing of Mg-based alloy systems with destabilized dehydriding thermodynamics by reversible intermetallics formation So far, the thermodynamic destabilization of MgH2 has been achieved in some Mg-based alloy systems by reversible intermetallics formation. The reported systems include Mg−In [22,23], Mg−Cd [24,25], Mg − Ga [26], and Mg−Ag [27–29] binary alloys, Mg(In, Al) and Mg(In, Cd) ternary solid solutions [30,32], Mg−In−Ni [33], Mg−Ag−In [34], Mg−Ag−Al [35], and Mg−Ag−Zn ternary alloy systems. However, there is no clear principle for designing of Mg-based alloy systems with destabilized dehydriding thermodynamics by reversible intermetallics formation. Table 2 summarizes the hydrogen storage capacity, dehydriding thermodynamic parameters and apparent activation energy for the dehydrogenation of pure Mg and DTRIF systems. It should be noted that all of the listed alloying elements (In, Ag, Al, Cd, Ni, and Zn) can form intermetallic compounds with Mg upon hydrogenation, such as MgIn, MgAg and Mg4Al6Ag. Furthermore, the formation enthalpies of these compounds are low, ensuring that they can play as the reactive additives for MgH2 to release hydrogen. Although other elements such as Sn, Si, and Pb can form Mg-based compounds Mg2Si, Mg2Sn, and Mg2Pb [41,42], they are too stable to react with MgH2. Besides, alloying elements are readily expelled from the compounds upon MgH2 formation and may react with H to form elemental hydrides. However, most of the hydride mixtures are not reversible. The thermal decomposition temperatures of some hydrides are listed in Table 3. It clearly shows that the thermal decomposition temperatures of the hydrides of In, Ag, Cd, Al and Zn are much lower than that of MgH2, thus inhibiting the disproportionation of Mg-based alloys and promoting the reversibility during hydrogenation and
Fig. 6. XRD patterns of Mg78Ag16.5Zn5.5 ternary alloy at different states: (a) ball milled under 10 MPa H2 for 4 h, (b) hydrogenated under 5 MPa H2 at 320 °C for 12 h, and (c) dehydrogenated at 320 °C for 12 h. 800
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Fig. 7. TEM bright field images, electron diffraction patterns, and high-resolution TEM images of Mg78Ag16.5Zn5.5 ternary alloy at different states: (a) hydrogenated and (b) dehydrogenated. Table 3 The thermal decomposition temperatures (Tdec) of related hydrides (some of them are estimated values) [44].
pure Mg [33] Mg90In10 [30] Mg3Ag [29] Mg90In5Al5 [30] Mg90In5Cd5 [32] Mg14In3Ni3 [33] Mg85Ag5Al10 [35] Mg78Ag16.5Zn5.5
ΔS (J K−1 mol−1 H2)
Peq at 300 °C (MPa)
Ea (kJ mol−1)
7.0 4.2 2.0 4.8 4.3 1.8 3.8 2.5
72.2 65.2 82.3 66.3 86.0 70.1 80.3 77.2
129.8 121.8 147.7 121.2 154.8 132.2 148.2 143.4
0.16 0.27 0.17 0.20 0.18 0.33 0.26 0.28
160 146.0 – 156.0 170.1 78.5 124.8 116.5
InH3
AgH
CdH2
AlH3
ZnH2
CaH2
TiH2
YH2
Tdec (°C)
327
15
> −100
−20
150
90
600
380
510
4. Conclusions
Table 2 Hydrogen storage capacity, dehydrogenation enthalpy (ΔH) and entropy (ΔS), dehydriding equilibrium pressure (Peq) at 300 °C and apparent activation energy for the dehydrogenation (Ea) of pure Mg and DTRIF systems. ΔH (kJ mol−1 H2)
MgH2
dehydrogenation. In this sense, Ca, Ti and Y are not proper, because their hydrides are more stable than MgH2 [43], which need higher thermal decomposition temperatures in comparison with pure MgH2. Based on the sum up of the systems already known, we proposed two guidelines for designing the DTRIF systems. First, the alloying elements can react with Mg to form intermetallic compounds, but the reaction enthalpies should be low. Second, the hydrides of the alloying elements should have low stabilities. It is believed that there is still chance to explore new DTRIF systems according to the above guidelines. However, it should be noted that capacity and kinetics are usually influenced by the alloying addition. Therefore, the selection of alloying elements for Mg-based hydrogen storage alloys needs more consideration and exploration.
Fig. 8. Isothermal dehydrogenation kinetic curves for Mg90Ag7.5Zn2.5 ternary alloys.
Capacity (wt.%)
Hydrides
Mg−Ag−Zn ternary alloys were prepared by ball milling MgH2 with Ag(Zn) solid solution to create a DTRIF system. A two-step hydrogen desorption reactions occurred in Mg90Ag7.5Zn2.5 alloy with the maximum hydrogen storage capacity of 4.2 wt%. In the first step, MgH2 reacts with MgAg and MgZn1.8Ag0.2 to release hydrogen, with an elevated dehydriding equilibrium pressure of 0.28 MPa at 300 °C. This thermodynamic destabilized dehydriding reaction can be intact realized in Mg78Ag16.5Zn5.5 alloy with a maximum hydrogen storage capacity of 3.0 wt%. The improved dehydriding kinetics was also achieved in the Mg−Ag−Zn ternary alloys, and the apparent activation energies for the dehydrogenation of Mg90Ag7.5Zn2.5 and Mg78Ag16.5Zn5.5 were lowered to 118.7 and 116.5 kJ mol−1, respectively. This is due to the abundant phase boundaries exist in the hydrogenated Mg−Ag−Zn alloys, providing more hydrogen diffusion channels. By sum up the already reported DTRIF systems, two guidelines were proposed for designing the DTRIF systems: The alloying elements can react with Mg to form intermetallic compounds with low reaction enthalpies, and the 801
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hydrides of the alloying elements should have low stabilities.
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