hydriding of Mg Base Solid Solutions with

Article pubs.acs.org/JPCC

Fully Reversible De/hydriding of Mg Base Solid Solutions with Reduced Reaction Enthalpy and Enhanced Kinetics Hui Wang,† Haichang Zhong,‡ Liuzhang Ouyang,† Jiangwen Liu,§ Dalin Sun,ζ Qingan Zhang,ξ and Min Zhu*,† †

School of Materials Science and Engineering, South China University of Technology, Guangzhou, 510640, China Department of Materials Science and Engineering, Xiamen University of Technology, Xiamen, 361024, China § Key Laboratory of Advanced Energy Storage Materials of Guangdong Province, Guangzhou, 510640, China ζ Department of Materials Science, Fudan University, Shanghai, 200433, China ξ School of Materials Science and Engineering, Anhui University of Technology, Maanshan, 243002, China ‡

S Supporting Information *

ABSTRACT: This paper presents a new approach to tune the de/hydriding thermodynamic properties of Mg via forming reversible Mg base solid solutions in the Mg−In and Mg−In−Al systems by mechanical milling. The effect of solubility of In and Al on the reversible formation of solid solution structure and hydrogen storage properties were investigated. It is found that although the solute atoms unavoidably are rejected upon hydriding, the hydrogenated products of MgH2 and intermediate MgIn compound could fully transform back to solid solution after dehydrogenation. In the hydriding of Mg(In, Al) ternary solid solution, Al would get dissolved into MgIn compound rather than forming free Al like the Mg(Al) binary solid solution. Therefore, the presence of In improves the dehydriding reversibility of Mg(Al) solid solution, and the reversible Al concentration could be increased up to the 8 at. %, which is just the solubility limit of Al in Mg by mechanical milling. The reversible phase transformation is responsible for the reduction in the desorption enthalpy of MgH2, being 12 kJ/(mol·H2) reduction for the alloy Mg0.9In0.1 relative to the desorption enthalpy of pure MgH2. Further, the hydrogen sorption kinetics of Mg(In) solid solutions are enhanced. Comparatively, both the thermodynamic destabilizing effect and the kinetic enhancing effect due to the Al dissolving are inferior to those due to the In dissolving. This work demonstrates a feasible way to improve the thermodynamics and kinetics of Mg base hydrogen storage alloys through traditional metallurgical method.

1. INTRODUCTION The commercial application of hydrogen as clean fuel of fuel cell driven vehicles demands efficient and safe hydrogen storage.1 The solid hydrogen storage via hydrides remains a great challenge because of many critical technical requirements on the storage density, charging/discharging rate, working temperature, and hydrogen pressure.2 Up to date, various metal hydrides (e.g., MgH2) and complex hydrides (e.g., NaAlH4, LiNH2, LiBH4) have been intensively investigated for reversible hydrogen storage.3−5 Unfortunately, these high-capacity hydrides generally suffer from harsh de/rehydrogenation temperature and pressure that are impractical for on-board application. With respect to MgH2 for reversible hydrogen storage, the fundamental issues have been addressed by its overhigh thermodynamic stability and sluggish kinetics.6 A high desorption enthalpy of 74.6 kJ/(mol·H2) means that the decomposition of MgH2 under 1 bar H2 requires a temperature close to 573 K, which would be further elevated because of unfavorable kinetics in association with slow H2 dissociation/ recombination, H diffusion through the Mg/MgH2 lattice, and hydride nucleation. Previous tremendous efforts have made © 2014 American Chemical Society

great progress in enhancing the hydrogen sorption kinetics of Mg by nanostructuring and doping catalysts.7−12 However, the unchanged thermodynamics of MgH2 determines improper pressure−temperature conditions as well as complicated heat management for practical applications. So far, only a few works have succeeded in decreasing the desorption temperature of MgH2 via thermodynamic destabilization approach. For example, Vajo et al. demonstrated a desorption enthalpy of 36.4 kJ/(mol·H2) in the MgH2/Si system and attributed the lowered enthalpy to the formation of Mg2Si that is thermodynamically more stable than pure Mg.13 Similar improvement was achieved in the MgH2−2LiBH4 composite hydride as dehydrogenating into MgB2.14 Unfortunately, these multicomponent composites generally suffer the disadvantages of poor reversibility and kinetics because of severely limited mass transport of elements, and thus, the desorption temperature remains high. Another feasible approach for thermodynamic destabilization is to form a new Received: November 16, 2013 Revised: March 31, 2014 Published: May 20, 2014 12087

dx.doi.org/10.1021/jp411265b | J. Phys. Chem. C 2014, 118, 12087−12096

The Journal of Physical Chemistry C

Article

prepared by high-energy ball-milling. Prior to ball-milling, (Mg + In) and (Mg + In + Al) powder mixtures with designed compositions were pressed into pellet and were sintered for 6 h at 573 K under vacuum. Afterward, the pellets were pulverized and milled on a planetary mill using stainless balls and vials under the protection of pure Ar atmosphere. Mg−Al alloys were directly prepared by milling the (Mg + Al) powder mixture. The ball-to-powder weight ratio was 20:1. To prevent over rising of temperature, milling was interrupted for 30 min after every 30 min of milling until finishing of the program. Pure Mg powder was also milled under the same conditions for comparison. X-ray diffraction (XRD) analysis at ambient temperature and high temperature was, respectively, conducted on PANalytical X’Pert and Rigaku D/Max 2500W diffractometer with Cu−Kα radiation. The lattice constants were calculated by Rietveld refinement using Rietan-2000 program.25 Scanning electron microscope (SEM, Nova NANOSEM 430) and transmission electron microscope (TEM, Jeol JEM-2100) were used to characterize the microstructure. The hydrogen storage properties were evaluated by measuring the pressure−composition isotherms (PCI) and isothermal de/hydriding kinetics on an automatic Sievert-type apparatus (AMC Co., U.S.). The sample chamber was evacuated about 30 min before heating up to target temperature. Activation treatment was conducted by three hydrogenation−dehydrogenation cycles at 633 K prior to PCI and kinetic measurements. The initial pressures of sample chamber for isothermal hydriding and dehydriding kinetic tests were 3 MPa hydrogen and vacuum, respectively.

structure by alloying Mg with other metals; the most successful example is the Mg2Ni alloy with desorption enthalpy of 64.7 kJ/(mol·H2). In addition to Mg2Ni, body-centered cubic (BCC) structures in the Mg−Ti/Mg−Co/Mg−Ni/Mg−Nb systems were synthesized by high-energy ball-milling and physical vapor deposition.15−19 A significant breakthrough was recently reported in the BCC Mg0.75Nb0.25 film that exhibited a reversible 4.0 wt % H capacity at 448 K and a hydride formation enthalpy of −53 kJ/(mol·H2).21 However, these metastable BCC structures tended to disintegrate when de/ hydrogenating at elevated temperature or maximum capacity. For example, the BCC Mg50Co50−H solid solution transformed irreversibly to the tetragonal Mg2CoH5 and a small quantity of additional Co at above 413 K,18 and BCC Mg75Nb25 went through partial decomposition into the equilibrium structure after 40 cycles of dehydrogenation.19 Such hydrogen-induced phase decomposition also occurred in the de/hydriding of many equilibrium or nonequilibrium Mg base alloys.20−23 For instance, the supersaturated Mg12YNi solid solution by rapid solidification irreversibly decomposed into MgH2, YH2(YH3), and Mg2NiH4 upon hydriding,20 and the Mg−rare-earth alloys with long-period stacking ordered structure decomposed into nanocomposites of MgH2 and rare-earth hydrides.22 Therefore, most Mg−base alloys for hydrogen storage are limited because of unavoidable phase separation and structural irreversibility. Recently, our group made a breakthrough in the reversible hydrogen storage of Mg−In solid solution.23 In-situ and ex-situ X-ray diffraction analysis revealed the following hydriding and dehydriding reaction for the solid solution Mg0.95In0.05: Mg 0.95In 0.05 + H 2 ↔ MgH2 + β

(1)

3. RESULTS AND DISCUSSION 3.1. Mg(In) Binary Solid Solutions. 3.1.1. Phase Structure and Microstructure of Mg(In) Solid Solutions. Figure 1 presents the formation of binary Mg(In) solid

where β phase is the disordered MgIn compound existing at elevated temperature, and it would transform to ordered MgIn compound (β″ phase hereafter) with L10 structure at room temperature according to the Mg−In binary phase diagram. The solid solution Mg0.95In0.05 could be completely recovered from the hydrogenated product even after tens of dehydrogenation. More importantly, the solid solution Mg0.95In0.05 presented an elevated desorption equilibrium pressure in comparison with pure Mg, indicating the thermodynamic destabilization effect due to the In dissolving. Following this, Zhou et al. added TiMn2 as catalyst to improve the dehydriding kinetics of Mg(In) solid solutions and achieved a lowtemperature hydrogen release at 423 K.24 Herein, the present work is aimed to investigate the hydrogen storage thermodynamic and kinetic properties of Mg(In) solid solution correlating with the microstructure and In solubility. In addition, the mechanically milled Mg−Al and Mg−In−Al solid solution alloys were explored for reversible hydrogen storage because of relatively high solubility of Al in Mg. The purpose of Al addition was to partially replace expensive and heavy In. We successfully synthesized the Mg(In, Al) ternary solid solution and achieved fully reversible de/ hydriding in the ternary alloy. The reversible solubility limit of In and Al in the binary and ternary alloys in association with the structural change in hydriding and dehydriding was investigated.

Figure 1. XRD patterns of the original Mg powder and the Mg(In) solid solutions with different In content after milling for 50 h; the dashed line indicates the shift of Bragg peaks of Mg because of the dissolving of In.

solutions by milling the sintered powder mixtures with nominal composition of Mg0.9In0.1, Mg0.95In0.05, and Mg0.98In0.02 for 50 h. The absence of Bragg peaks of In indicates the complete dissolution of In in Mg within the resolution of XRD, which is further verified by the slight shift of Mg peaks toward the larger 2θ angle values relative to those of the original Mg powder. As marked by the dashed line, there is more shift for larger In solubility. On the basis of the XRD data, the lattice constants of

2. EXPERIMENTAL SECTION Commercial Mg (99.9% purity), indium (99.9% purity), and aluminum (99.9% purity) powders were used as starting materials. Mg−In, Mg−Al, and Mg−In−Al alloys were 12088



Article

Table 1. Lattice Constants of Mg(In) and Mg(In, Al) Solid Solutions and Their Dehydrogenation Enthalpy (ΔH) and Entropy (ΔS) lattice constants (nm)

a

solid solution

a

c

capacity (wt %)

ΔH (kJ/(mol·H2))

ΔS (J/(K·mol·H2))

Mg0.9In0.1 Mg0.95In0.05 Mg0.98In0.02 pure Mga pure Mgb Mg0.95Al0.05 Mg0.9In0.05Al0.05

0.31929(6) 0.32027(2) 0.32077(1) 0.32113(1)

0.52061(7) 0.52086(3) 0.52106(2) 0.52140(2)

4.2 5.3 6.4

0.3199(5) 0.3198(3)

0.51960(3) 0.51929(1)

65.2 68.1 69.6 77.9 74.6 76.8 66.3

121.8 125.5 126.0 138.9 133.4 138.6 121.2

4.8

Commercial Mg powder measured in this work. bReference 27.

certain aggregated particles of β″ phase corresponding to bright spots, which are dispersed in the gray MgH2 matrix. The deliberate microstructure of hydrogenated Mg0.9In0.1 was further investigated by TEM and high-resolution TEM (HRTEM). The TEM bright field image (Figure 2d) demonstrates that the crystallite of β″ phase is in a size range of 100−300 nm. The HRTEM image of MgH2/β″ phase boundary is shown in Figure 3; one can see that there exists an

Mg(In) solid solutions are deliberately measured by introducing high-purity Si powder as an internal standard, and the results are summarized in Table 1. For the solid solution Mg0.9In0.1, the lattice constant a and c are, respectively, decreased to 0.31929(6) nm and 0.52061(7) nm from 0.32113(1) nm and 0.52140(2) nm for the original Mg powder. Although the atomic radius (rIn = 1.66 Å) of In is larger than that (rMg = 1.60 Å) of Mg, the dissolving of In causes the shrinkage of the Mg unit cell. This unusual variation of lattice constants should arise from a relatively large electronegativity difference between Mg (χMg = 1.23) and In (χIn = 1.49), which enhances the attraction of Mg and neighboring In atoms and which results in the contraction of the unit cell.26 The microstructures of hydride and dehydrided Mg(In) solid solution were revealed by SEM observations on the alloy Mg0.9In0.1. As seen in Figure 2(a, b), backscattering SEM

Figure 3. High-resolution TEM image of hydrogenated Mg0.9In0.1 showing the intermediate layer of ca. 10 nm thickness between the MgH2 and the β″ phases.

intermediate layer of ca. 10 nm thickness, which is deduced to be composed of Mg5In2 and Mg2In compounds according to analysis of HRTEM image. It is thus assumed that there exists a concentration gradient of In at the MgH2/β″ phase boundaries, increasing from MgH2 to β″ phase region. The microstructure of β″ phase is expected to have an important effect on the dehydriding of MgH2, which will be subsequently discussed in combination with the hydrogen storage properties of Mg(In) solid solutions. 3.1.2. Thermodynamic Properties of Mg(In) Solid Solutions. The thermodynamic properties of Mg(In) solid solutions were evaluated by PCI measurement. Figure 4 compares PCI curves of Mg(In) solid solutions containing different In content with that of 80 h milled pure Mg. It is clearly seen that plateau pressures, which reflect the phase equilibrium between alloy and hydride, of three Mg(In) solid solutions at 586 K are elevated compared with that of pure Mg at 588 K and that the Mg(In) solid solution shows positive dependence of plateau pressure on the In solubility. The PCI

Figure 2. Backscattering SEM images of Mg0.9In0.1 alloy at different states: (a) as-milled, (b) as-dehydrogenated, and (c) as-hydrogenated. Both SEM image c and TEM image d of as-hydrogenated Mg0.9In0.1 show that fine grains of β″ phase are embedded in the MgH2 matrix.

images clearly show homogeneous single phase microstructure for Mg 0.9 In 0.1 alloy at both milled (Figure 2a) and dehydrogenated (Figure 2b) states. Comparatively, the particle size of as-milled Mg0.9In0.1 is in tens of micrometers, which is reduced down to finer size after several dehydrogenation cycles. With respect to the hydrogenated product of Mg0.9In0.1, backscattering SEM micrograph (Figure 2c) presents fine and 12089



Article

reaction pathway of MgH2 as shown in eq 1 in which the formation of Mg(In) solid solution instead of Mg would predominately contribute to the enthalpy reduction because the former is more thermodynamically stable than the latter. With larger In solubility, more stable Mg(In) solid solution, and thus lower desorption enthalpy, could be obtained. At the expense of capacity for thermodynamic improvement, the Mg(In) solid solutions show relatively high reversible hydrogen storage capacity, being 6.4, 5.3, and 4.2 wt % for the Mg0.98In0.02, Mg0.95In0.05, and Mg0.9In0.1 alloys, respectively. The measured capacities are slightly less than their theoretical values, which is due to partial oxidation in the alloy preparation or incomplete de/hydrogenation. The alloy Mg0.9In0.1 keeps relatively high hydrogen storage capacity as compared with Mg2Ni alloy having 3.6 wt % H capacity and almost the same ΔH of 64.7 kJ/(mol·H2). As also seen in Figure 4, the PCI curve of pure Mg at 588 K shows evident desorption lag, namely, that the hydrogen release is only initiated under much lower pressure than the Peq, and then the hydrogen pressure gradually increases to the plateau value. This lagging phenomenon is related to the sluggish dehydriding kinetics of MgH2. In contrast, the Mg(In) solid solutions even start desorbing hydrogen above the equilibrium pressure, and the pressure continuously decreases with hydrogen release prior to reaching the plateau region. The desorption slope is particularly obvious for the alloy Mg0.9In0.1, and at least 1.4 wt % H is desorbed at the slope stage. A similar slope characteristic can also be observed in the PCI absorption curves of Mg(In) solid solutions. This phenomenon suggests different de/hydriding behavior at the slope stage from that at the plateau region. In the hydriding of Mg(In) solid solution, the formation of MgH2 is accompanied by the rejection and migration of In atoms, resulting in the increase of In solubility in the unhydrided Mg(In). Along with continuous hydriding, the growth and interconnection of MgH2 grains would isolate several remaining Mg(In) grains that would finally transform into β phase. Therefore, at the later period of hydrogen absorption corresponding to the slope of PCI absorption curve, the diffusion of In into the remaining supersaturated Mg(In) solid solution would become more difficult, and thus, higher hydrogen pressure is required to ensure that H enters the lattice of residual Mg(In) and reacts with Mg to form MgH2. With respect to the dehydriding process, at the initial stage corresponding to the PCI desorption slope, MgH2 would readily react with unstable neighboring nanosized Mg−In compounds (e.g., Mg5In2, or Mg2In) and would decompose into Mg(In). From the kinetic viewpoint, the decomposition reaction at early stage should be very easy to proceed because of high interfacial energy and rather short diffusion range of In atoms. Therefore, we believe that the slope characteristic of the PCI curve is related to the specific thermodynamic and kinetic properties of Mg(In) solid solutions. 3.2. Mg(In, Al) Ternary Solid Solutions. The purpose of Al addition is to partially replace expensive and heavy indium while keeping thermodynamic and kinetic enhancing effects. For comparison with Mg−In−Al ternary alloys, Mg−Al alloys were also prepared by mechanical milling from the starting Mg and Al powders. It was expected that the presence of In would extend the solubility of Al in Mg as well as improve the dehydriding reversibility of Mg−Al alloys. 3.2.1. Preparation of Mg(In, Al) Ternary Solid Solutions. Figure 6 show the XRD patterns of as-milled alloys Mg0.9Al0.1

Figure 4. Pressure−composition isotherms of pure Mg and Mg(In) solid solutions with different In content, showing the elevated plateau pressure due to the In dissolving.

results undoubtedly demonstrate the thermodynamic destabilization of Mg because of the In dissolving. The fitted van’t Hoff plots of pure Mg and Mg(In) solid solutions are given in Figure 5, which shows the linear

Figure 5. Comparison on van’t Hoff plots among pure Mg and Mg(In) solid solutions, showing elevated equilibrium pressures with increasing In content.

relationship between reaction temperature T and equilibrium plateau pressure (Peq hereafter). Peq was adopted as the midpoint value of PCI desorption plateau (see Figure S1 of the Supporting Information) of Mg base solid solution alloys. It is clearly seen that the alloy Mg0.9In0.1 has the highest Peq and that pure Mg has the lowest Peq. On the basis of the fitted van’t Hoff plots and van’t Hoff equation RT ln Peq = ΔH − TΔS, the dehydriding enthalpy (ΔH hereafter) and entropy (ΔS hereafter) are calculated and shown in Table 1. The results indicate that both ΔH (77.9 ± 0.3 kJ/(mol·H2)) and ΔS (138.9 ± 0.3 J/(K·mol·H2)) for pure Mg obtained under present experiment conditions are slightly larger than generally accepted values (ref 27, ΔH = 74.6 ± 0.4 kJ/(mol·H2), ΔS = 133.4 ± 0.7 J/(K·mol·H2)). This difference should originate from systematic instrument errors on the temperature and pressure measurements or from the judgment of equilibrium state. Regardless of this difference, the determined desorption enthalpies are less for Mg(In) solid solutions than that for pure Mg, being 65.2, 68.1, and 69.6 kJ/(mol·H2) for the Mg0.9In0.1, Mg0.95In0.05, and Mg0.98In0.02 alloys, respectively. It is inferred that the lowered dehydriding enthalpy is related to an altered 12090



Article

shows that no Mg−Al intermetallic compounds are formed during milling and that a single solid solution is finally obtained after milling for 80 h. A single solid solution structure is also found in the 80 h milled Mg0.9In0.05Al0.05 alloy, and the lattice constants are shown Table 1. For the alloy Mg0.85In0.05Al0.1 with higher Al content, the milling product consists of Mg17Al12 and Mg (Figure 7d), which is similar to the phase composition of 70 h milled Mg0.9Al0.1 alloy (Figure 6c). The results indicate that the Al solubility could not be extended by alloying with In. 3.2.2. De/hydriding Mechanism of Mg(In, Al) Solid Solutions. Figure 8 shows the XRD patterns of Mg0.95Al0.05

Figure 6. XRD patterns of Mg0.95Al0.05 milled for (a) 15 h and (b) 70 h, and Mg0.9Al0.1 milled for (c) 15 h and (d) 70 h.

and Mg0.95Al0.05. After milling the Mg0.95Al0.05 alloy for 15 h (Figure 6a), the Al peaks become weakened and broadened; there is also a slight shift of Mg peaks toward the larger 2θ value, which indicates the partial dissolving of Al in Mg. Prolonging milling time to 70 h (Figure 6b), a single Mg(Al) solid solution phase is formed. In contrast, the Mg17Al12 phase is formed in the 15 h milled Mg0.9Al0.1. After 70 h of milling, the milling product is composed of Mg and Mg17Al12, and the Al almost completely transforms to Mg17Al12. The difference in the phase structures of Mg0.9Al0.1 and Mg0.95Al0.05 alloys indicates that the solubility limit of Al in Mg at ambient temperature is less than 10 at. %, which is in good accord with previously reported results.28,29 The calculated lattice constants of Mg0.95Al0.05 solid solution are slightly less than those of pure Mg, as shown in Table 1. There exists a similar Al solubility limit in the Mg−In−Al ternary alloy. Figure 7 shows the XRD patterns of Mg0.92In0.05Al0.03 and Mg0.85In0.05Al0.1 alloys with different milling times. For the alloy Mg0.92In0.05Al0.03, Figure 7(a−c)

Figure 8. XRD patterns of (a, b) Mg0.95Al0.05 and (c, d) Mg0.92In0.05Al0.03 at hydrogenated and dehydrogenated states.

and Mg0.92In0.05Al0.03 alloys at hydrogenated and dehydrogenated states. For the alloy Mg0.95Al0.05, it is found that the hydrogenated product consists of main phase MgH2 and a small amount of Al (Figure 8a). The presence of Al indicates the rejection of Al because of hydride formation. After dehydrogenation (Figure 8b), free Al transforms to Mg17Al12 rather than getting dissolved in Mg, which is in agreement with previous literature. 30−33 By contrast, the ternary solid solution Mg0.92In0.05Al0.03 is hydrogenated into MgH2 and β″ (Figure 8c), which transform back to solid solution after dehydrogenation (Figure 8d). No Mg−Al intermetallic compounds or Al phase is found in the hydrogenated or dehydrogenated products. SEM images as shown in Figure 9 further verify the structural reversibility. The alloy Mg0.92In0.05Al0.03 at both milled (Figure 9a) and dehydrogenated (Figure 9b) states presents homogeneous single phase microstructure, while finely distributed β″ phase could be found in the hydrogenated alloy (Figure 9c). Similar microstructure evolution in the de/rehydrogenation cycle is also found in the alloy Mg0.9In0.05Al0.05. Therefore, the ternary solid solutions Mg0.92In0.05Al0.03 and Mg0.9In0.05Al0.05 show the same hydriding and dehydriding behaviors with Mg(In) binary solid solution. To further reveal the reversible dehydriding mechanism of Mg(In, Al) ternary solid solution and the existence of Al, in-situ XRD analysis and SEM observations were performed on the alloy Mg0.9In0.05Al0.05. The hydrogenated sample was heated at a ramping rate of 10 K·min−1 and was radiated at every target temperature for 10 min, and the obtained in-situ XRD result is shown in Figure 10. It is seen that the diffraction intensity of β″

Figure 7. XRD patterns of Mg0.92In0.05Al0.03 milled for (a) 15, (b) 50, and (c) 80 h, and Mg0.85In0.05Al0.1 milled for (d) 80 h. Pattern c shows the formation of Mg(In, Al) ternary solid solution. 12091



Article

Figure 9. Backscattering images of Mg0.92In0.05Al0.03 alloy at (a) milled, (b) dehydrogenated, and (c) hydrogenated states.

Figure 11. Backscattering SEM image of hydrogenated Mg0.9In0.05Al0.05 alloy, and the inset is the Al elemental mapping.

Figure 10. In-situ XRD patterns for hydrogenated Mg0.9In0.05Al0.05 solid solution at different temperatures. The β″ to β phase transition occurs prior to the decomposition of MgH2.

formation of Mg(In, Al) solid solution instead of Al or Mg17Al12. Consequently, although the presence of indium could not extend the solubility limit of Al in Mg, the codissolving of In and Al blocks the precipitation of free Al and the formation of intermetallic compounds in the Mg−Al system; the dehydriding reversibility of Mg(Al) solid solution is thus improved. 3.2.2. Thermodynamic Properties of Mg(In, Al) Solid Solutions. Figure 12 shows the PCI desorption curves of Mg0.9In0.05Al0.05 alloy at different temperatures. Mg0.9In0.05Al0.05 has a reversible capacity of ca. 5.0 wt %, a little higher than that of Mg 0.9 In 0.1 alloy. The desorption enthalpy ΔH of Mg0.9In0.05Al0.05 is determined to be 66.3 kJ/(mol·H2), as also shown in Table 1, which is just slightly lower than that of Mg0.95In0.05 with the same In content. This result indicates that the Al addition further decreases the ΔH of Mg(In) solid solution. Comparatively, the thermodynamic destabilizing effect for the Al dissolving is inferior to that for the In dissolving

phase decreases with increasing temperature. The β″ to β phase transition occurs above 613 K, which is followed by accelerated decomposition of MgH2. At 673 K, a single Mg solid solution phase is fully recovered with the disappearance of MgH2 and β phases. The in-situ XRD result definitely confirms the absence of Al and Mg−Al intermetallic compounds during dehydriding and further indicates same dehydriding mechanism with Mg(In) binary solid solution. Figure 11 shows SEM backscattering image for the hydrogenated Mg0.9In0.05Al0.05 and the elemental mapping of Al. It is evident that the distributions of Al and β″ phases are highly overlapped. It is thus deduced that Al would get dissolved in the β phase upon hydrogenation, that is, would form another Mg−In−Al ternary alloy with higher concentration of Al and In. In the dehydriding process, the reaction of this Mg−In−Al alloy with MgH2 would result in the reversible 12092



Article

seems that the dissolving of In is beneficial to the sorption of H in Mg, but the Al dissolving plays the opposite effect. The hydrogen sorption data shown in Figures 13−15 were fitted using the Johnson−Mehl−Avrami equation, xα = 1 − exp[−(kt)n], where xα is the reaction fraction, k is the reaction rate, t is the reaction time, and n is the reaction exponent. The obtained k value is then used by the Arrhenius equation to calculate the apparent activation energy Ea (see Figures S2−S4 of the Supporting Information). The calculated Ea for hydrogen absorption of Mg0.95In0.05 and Mg0.9In0.1 is 48.1 kJ/(mol·H2) and 56.5 kJ/(mol·H2), respectively, a notable decrease in comparison with the values of 95−120 kJ/(mol·H2) for pure Mg.32,36,37 The Ea for hydrogen desorption is 145.0 kJ/(mol· H2) for Mg0.95In0.05 and 146.0 kJ/(mol·H2) for Mg0.9In0.1, falling within the previously reported range 120−160 kJ/(mol· H2) for pure Mg.32,36,37 Comparatively, the calculated Ea for hydriding and dehydriding of Mg0.9In0.05Al0.05 is 73 kJ/(mol· H2) and 156 kJ/(mol·H2), respectively, both of which are higher than those of Mg0.9In0.1. 3.4. Discussion. As mentioned above, Mg(In) and Mg(In, Al) solid solutions show reversible hydriding and dehydriding properties and lowered desorption enthalpy because of the altered reaction route, and the thermodynamic destabilizing effect is dependent on the solubility of In and Al. The larger the overall solubility of In and Al is, the greater the enthalpy reduction is. To achieve better thermodynamic destabilizing effect in Mg base solid solution, it is feasible to increase the solubility of In and Al. Nevertheless, there exists a respective solubility limit of In and Al in Mg. For the Mg(In) binary solid solution, the maximum solubility of In at room temperature is ca. 11 at. % according to the Mg−In phase diagram. The equilibrium solubility of Al in Mg is ca. 1 at. % at 373 K, which could be further extended to be ca. 8 at. % by mechanical milling. When forming Mg(In, Al) ternary solid solution, the presence of In extends the reversible solubility of Al in Mg, which is still less than 8 at. %. Therefore, the overall reversible solubility of In and Al in Mg and the resultant thermodynamic destabilization effect are limited. As shown in Table 1, the dehydriding reaction entropies of all alloys are also decreased relative to that of pure Mg. With larger solubility of In and Al, there is more entropy reduction for dehydriding. For example, the alloys Mg0.9In0.1 and Mg0.9In0.05Al0.05 have close ΔS of 121.8 and 121.2 J/(K·mol·

Figure 12. Pressure−composition isotherms of Mg0.9In0.05Al0.05.

because the ΔH of Mg0.9In0.05Al0.05 is slightly higher than the desorption enthalpy of 65.2 kJ/(mol·H2) for the alloy Mg0.9In0.1, which has the same overall solubility of In and Al. The reason for which might be because Mg(In) solid solution is thermodynamically more stable than Mg(Al) solid solution.34 3.3. Hydrogen Sorption Kinetic Properties of Mg(In) and Mg(In, Al) Solid Solutions. The sloping PCI curves have implied the enhanced dehydriding kinetics for the Mg(In) and Mg(In, Al) solid solutions. In Figures 13, 14, and 15, the isothermal hydriding and dehydriding curves of the solid solutions Mg0.95In0.05, Mg0.9In0.1, and Mg0.9In0.05Al0.05 are, respectively, presented. In comparison to pure Mg milled under similar conditions,35 both the hydriding and the dehydriding rates of Mg base solid solution alloys are greatly improved. At 609 K, the hydrogen uptake content of Mg0.95In0.05 reaches the maximum in ca. 10 min, while ca. 5.0 wt % hydrogen is released within 30 min at 608 K. Comparatively, the alloy Mg0.9In0.1 releases the maximum hydrogen amount of ca. 2.8 wt % within ca. 20 min; therefore, there is no obvious difference in the dehydrogenation rate between Mg0.9In0.1 and Mg0.95In0.05. Further comparison between the alloys Mg0.9In0.1 and Mg0.9In0.05Al0.05 shows that the binary solid solution has a much faster hydriding and dehydriding rate than the ternary solid solution. Therefore, it

Figure 13. (a) Hydriding and (b) dehydriding curves of Mg0.95In0.05 alloy at different temperatures, showing a faster hydrogen absorption rate than the desorption rate. 12093



Article

Figure 14. (a) Hydriding and (b) dehydriding curves of Mg0.9In0.1 alloy at different temperatures.

Figure 15. (a) Hydriding and (b) dehydriding kinetic curves of Mg0.9In0.05Al0.05 alloy at different temperatures.

entropy−enthalpy compensation effect, which is a well-known rule of thermodynamic behavior in various chemical systems.39 Regardless of the negative effect of entropy reduction on the thermodynamic improvement of MgH2, the dehydriding plateau pressures of all reversible solid solutions are elevated in comparison to that of pure Mg according to their van’t Hoff plots. Unlike other reversible Mg−base composite hydrides (e.g., 2MgH2−Si, MgH2−2LiBH4) with sluggish kinetics, the reversible Mg(In) and Mg(In, Al) solid solutions showed fast hydriding and dehydriding kinetic performances. For the MgH2−Si system, the formation and decomposition of Mg2Si in the hydrogenation and dehydrogenation necessitated harsh temperature and pressure conditions to overcome the problems of long-range diffusion and slow migration rate of Si.13 Actually, the rehydrogenation 2MgH2−Si even required a temperature up to 673 K, a hydrogen pressure up to 10 MPa, and a reaction time longer than 10 h. Similarly for the Mg(In) and Mg(In, Al) solid solutions, the hydrogenation and dehydrogenation involves the diffusion of H as well as the migration of solute atoms in the Mg lattice. Therefore, the kinetic enhancement for Mg(In) solid solution could be attributed to its specific microstructure and fast migration ability of In atoms in Mg. In hydriding of Mg(In) solid solution, the migration of In atoms appears to be very fast and would not increase the activation energy barrier. Conversely, it seems that the diffusion of H in Mg is enhanced because of the dissolution of In because of

H2), respectively, lower than the value 125.5 J/(K·mol·H2) for Mg0.95In0.05 and the value 138.9 J/(K·mol·H2) for pure Mg determined in this work. According to the van’t Hoff equation, the entropy decrease would partially offset the thermodynamic destabilization effect caused by enthalpy reduction. Similar with enthalpy reduction, the reason for entropy decrease is the dissolution of In and Al in Mg and the resultant altered dehydriding route. According to eq 1, the desorption entropy could be described with the formula ΔS = (SH2 + SMg(In) − SMgH2) − Sβ in which SH2 denotes the entropy of gaseous H2 and is generally assigned to be a constant value of ca. 130 J/(K· mol·H2). Therefore, the entropy of β phase (Sβ) dominantly contributes to the entropy decrease of the dehydriding reaction. Simultaneous reduction of reaction entropy and enthalpy is also found in other multicomponent hydrogen storage systems (e.g., Mg2Ni, LaNi5, etc.). For example, the hydride Mg2NiH4 has 64.7 kJ/(mol·H2) and 122.0 J/(K·mol·H2) for desorption ΔH and ΔS, respectively,38 both of which are lower than those of MgH2. In addition, Paskevicius et al. recently reported an entropy decrease of 3.8 J/(K·mol·H2) for MgH2 nanoparticles (down to ∼7 nm in size),27 which was coupled with enthalpy reduction of 2.84 kJ/(mol·H2) relative to that of bulk MgH2. This phenomenon is due to a different increase in the entropy of nanosized Mg and MgH2, namely, a lower reaction ΔS will be generated as the entropy of MgH2 nanoparticle increases by a greater quantity than that of Mg nanoparticle. Such concurrent change in reaction enthalpy and entropy is termed 12094



■

greatly decreased Ea for hydrogen absorption. With respect to the enhanced desorption kinetics, this is because the hydrogenated product contains a large amount of nanosized Mg−In compounds (e.g., Mg5In2, or Mg2In) and abundant phase boundaries, which have been cited as important components of Mg−base materials to improve the kinetics of H2 sorption.40 Namely, those unstable nanograins of Mg−In compounds would readily react with MgH2 at the interfaces and would initialize the decomposition reaction of MgH2. Further, the uniform distribution of fine grains of β phase embedded in the MgH2 matrix avoids the long-range diffusion of In. For the Mg(In, Al) ternary solid solution, the relatively worsening kinetics might be attributed to the slowed diffusion rate of In and H because of the incorporation of Al. Furthermore, the poor Al mobility would further retard the de/hydriding reaction. To further improve the kinetic performance of Mg−base solid solutions, doping catalysts is a feasible way. For instance, the TiMn2 catalyzed Mg−0.1In alloy hydrides even began to dehydrogenate at approximately 373 K because of the dual improvement in the thermodynamics and kinetics,24 and TiMn2 acted as a low-temperature metal hydride to catalyze the physisorption, chemical dissociation, and diffusion of H2. We also found that the addition of TiCl3 precursor greatly accelerated the dehydriding rate of Mg0.9In0.05Al0.05 at 573 K. However, the diffusion rate of solute atoms in Mg could not be accelerated by the catalyst.

AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. Tel: 86-20-87112762. Notes

The authors declare no competing financial interest.

■

ACKNOWLEDGMENTS This work was financially supported by the Ministry of Science and Technology of China under grant no. 2010CB631302; Natural Science Foundation of China under grant no. U1201241, 51071068, 51271078, and U1201241; Key Laboratory of Clean Energy Materials of Guangdong Higher Education Institute under grant no. KLB11003; and the Fundamental Research Funds for the Central Universities of China.

■

REFERENCES

(1) Schlapbach, L.; Zuttel, A. Hydrogen-Storage Materials for Mobile Applications. Nature 2001, 414, 353−358. (2) Yang, J.; Sudik, A.; Wolverton, C.; Siegel, D. J. High Capacity Hydrogen Storage Materials: Attributes for Automotive Applications and Techniques for Materials Discovery. Chem. Soc. Rev. 2010, 39, 656−675. (3) Li, H. W.; Yan, Y. G.; Orimo, S.; Zuttel, A.; Jensen, C. M. Recent Progress in Metal Borohydrides for Hydrogen Storage. Energies 2011, 4, 185−214. (4) Jain, I. P.; Lal, C.; Jain, A. Hydrogen Storage in Mg: A Most Promising Material. Int. J. Hydrogen Energy 2010, 35, 5133−5144. (5) Chen, P.; Zhu, M. Recent Progress in Hydrogen Storage. Mater. Today 2008, 11, 36−43. (6) Aguey-Zinsou, K. F.; Ares-Fernandez, J. R. Hydrogen in Magnesium: New Perspectives toward Functional Stores. Energy Environ. Sci. 2010, 3, 526−543. (7) Paskevicius, M.; Tian, H.-Y.; Sheppard, D. A.; Webb, C. J.; Pitt, M. P.; Gray, E. M.; Kirby, N. M.; Buckley, C. E. Magnesium Hydride Formation within Carbon Aerogel. J. Phys. Chem. C 2011, 115, 1757− 1766. (8) Yao, X. D.; Wu, C. Z.; Du, A. J.; Zou, J.; Zhu, Z. H.; Wang, P.; Cheng, H. M.; Smith, S.; Lu, G. Q. Metallic and Carbon NanotubeCatalyzed Coupling of Hydrogenation in Magnesium. J. Am. Chem. Soc. 2007, 129, 15650−15654. (9) Cheng, F. Y.; Tao, Z. L.; Liang, J.; Chen, J. Efficient Hydrogen Storage with the Combination of Lightweight Mg/MgH2 and Nanostructures. Chem. Commun. 2012, 48, 7334−7343. (10) Yu, X. B.; Guo, Y. H.; Yang, H.; Wu, Z.; Grant, D. M.; Walker, G. S. Improved Hydrogen Storage in Magnesium Hydride Catalyzed by Nanosized Ti0.4Cr0.15Mn0.15V0.3 Alloy. J. Phys. Chem. C 2009, 113, 5324−5328. (11) Xie, L.; Liu, Y.; Wang, Y. T.; Zheng, J.; Li, X. G. Superior Hydrogen Storage Kinetics of MgH2 Nanoparticles Doped with TiF3. Acta Mater. 2007, 55, 4585−4591. (12) Lillo-Ródenas, M. A.; Aguey-Zinsou, K. F.; Cazorla-Amorós, D.; Linares-Solano, A.; Guo, Z. X. Effects of Carbon-Supported Nickel Catalysts on MgH2 Decomposition. J. Phys. Chem. C 2008, 112, 5984− 5992. (13) Vajo, J. J.; Mertens, F.; Ahn, C. C.; Bowman, R. C.; Fultz, B. Altering Hydrogen Storage Properties by Hydride Destabilization through Alloy Formation: LiH and MgH2 Destabilized with Si. J. Phys. Chem. B 2004, 108, 13977−13983. (14) Bösenberg, U.; et al. Hydrogen Sorption Properties of MgH2− LiBH4 Composites. Acta Mater. 2007, 55, 3951−3958. (15) Kalisvaart, W. P.; Harrower, C. T.; Haagsma, J.; Zahiri, B.; Luber, E. J.; Ophus, C.; Poirier, E.; Fritzsche, H.; Mitlin, D. Hydrogen Storage in Binary and Ternary Mg-Based Alloys: A Comprehensive Experimental Study. Int. J. Hydrogen Energy 2010, 35, 2091−2103. (16) Kim, H.; Nakamura, J.; Shao, H.; Nakamura, Y.; Akiba, E.; Chapman, K. W.; Chupas, P. J.; Proffen, T. Insight into the

4. CONCLUSIONS We have demonstrated the synthesis, reversibility, and improved hydrogen storage properties of Mg(In) binary and Mg(In, Al) ternary solid solutions. It is found that Mg(In) binary solid solution and Mg(In, Al) ternary solid solution have the same hydriding products of MgH2 and β phase. In hydriding of Mg(In, Al) ternary solid solution, the Al would get dissolved in β phase, which explains the improved dehydriding reversibility in comparison to binary Mg(Al) solid solution. The reversible solubility limits of In and Al in Mg are approximately 10 and 8 at. %, respectively. The reversible dehydriding reaction of Mg(In) and Mg(In, Al) is featured of lowered desorption enthalpy/entropy and elevated equilibrium pressure in comparison with that of pure Mg. The solid solution Mg0.9In0.1 shows desorption enthalpy of 65.2 kJ/(mol·H2) and relatively high hydrogen storage capacity 4.2 wt %. Further, the alloying with In enhances the de/hydriding kinetics of Mg because of the fast migration ability of In, the uniform distribution of fine β″ particles, and the possible facilitating role of In to H diffusion. Comparatively, the thermodynamic destabilizing effect and the kinetic enhancing effect for the Al dissolving are inferior to that for the In dissolving. Our work opens a new way to tune the thermodynamic and kinetic performances of Mg−base hydrogen storage alloys in the presence of In.

■

Article

ASSOCIATED CONTENT

S Supporting Information *

The PCI curves for Mg0.98In0.02, Mg0.95In0.05, and Mg0.9In0.1 solid solutions at different temperatures. The calculation of the apparent activation energy of hydrogenation and dehydrogenation for Mg0.95In0.05 solid solution and Mg0.9In0.1 solid solution. This material is available free of charge via the Internet at http://pubs.acs.org. 12095



Article

(38) Sandrock, G. A Panoramic Overview of Hydrogen Storage Alloys from a Gas Reaction Point of View. J. Alloys Compd. 1999, 293−295, 877−888. (39) Starikov, E. B.; Nordén, B. Enthalpy−Entropy Compensation: A Phantom or Something Useful? J. Phys.Chem. B 2007, 111, 14431− 14435. (40) Orimo, S.; Fujii, H.; Horie, S. Investigations of Interfacial Materials Designthe Effect of Interface Microstructures in ZrCr1.8Cu0.3/Mg on MgH2 Formation as a Result of Hydrogen Interdiffusion. J. Alloys Compd. 1995, 231, 766−772.

Hydrogenation Properties of Mechanically Alloyed Mg50Co50 from the Local Structure. J. Phys. Chem. C 2011, 115, 20335−20341. (17) Shao, H.; Asano, K.; Enoki, H.; Akiba, E. Preparation and Hydrogen Storage Properties of Nanostructured Mg−Ni Bcc Alloys. J. Alloys Compd. 2009, 477, 301−306. (18) Zhang, Y.; Tsushio, Y.; Enoki, H.; Akiba, E. The Study on Binary Mg−Co Hydrogen Storage Alloys with Bcc Phase. J. Alloys Compd. 2005, 393, 147−153. (19) Tan, X.; Wang, L.; Holt, C. M. B.; Zahiri, B.; Eikerling, M. H.; Mitlin, D. Body Centered Cubic Magnesium Niobium Hydride with Facile Room Temperature Absorption and Four Weight Percent Reversible Capacity. Phys. Chem. Chem. Phys. 2012, 14, 10904−10909. (20) Zhang, Q. A.; Zhang, L. X.; Wang, Q. Q. Crystallization Behavior and Hydrogen Storage Kinetics of Amorphous Mg11Y2Ni2 Alloy. J. Alloys Compd. 2013, 551, 376−381. (21) Si, T. Z.; Liu, Y. F.; Zhang, Q. A. Hydrogen Storage Properties of the Supersaturated Mg12YNi Solid Solution. J. Alloys Compd. 2010, 507, 489−493. (22) Liu, J. W.; Zou, C. C.; Wang, H.; Ouyang, L. Z.; Zhu, M. Facilitating De/Hydrogenation by Long-Period Stacking Ordered Structure in Mg Based Alloys. Int. J. Hydrogen Energy 2013, 38, 10438−10445. (23) Zhong, H. C.; Wang, H.; Liu, J. W.; Sun, D. L.; Zhu, M. Altered Desorption Enthalpy of MgH2 by the Reversible Formation of Mg(in) Solid Solution. Scr. Mater. 2011, 65, 285−287. (24) Zhou, C.; Fang, Z. Z.; Lu, J.; Zhang, X. Thermodynamic and Kinetic Destabilization of Magnesium Hydride Using Mg−in Solid Solution Alloys. J. Am. Chem. Soc. 2013, 135, 10982−10985. (25) Izumi, F.; Ikeda, T. A Rietveld-Analysis Programm Rietan-98 and Its Applications to Zeolites. Mater. Sci. Forum 2000, 321−324, 198−203. (26) Busk, R. S. Lattice Parameters of Magnesium Alloys. Trans. AIME 1950, 188, 1460−1464. (27) Paskevicius, M.; Sheppard, D. A.; Buckley, C. E. Thermodynamic Changes in Mechanochemically Synthesized Magnesium Hydride Nanoparticles. J. Am. Chem. Soc. 2010, 132, 5077−5083. (28) Bououdina, M.; Guo, Z. X. Comparative Study of Mechanical Alloying of (Mg+Al) and (Mg+Al+Ni) Mixtures for Hydrogen Storage. J. Alloys Compd. 2002, 336, 222−231. (29) Tatiparti, S. S. V.; Ebrahimi, F. Electrodeposition of Al−Mg Alloy Powders. J. Electrochem. Soc. 2008, 155, D363−D368. (30) Shang, C. X.; Bououdina, M.; Song, Y.; Guo, Z. X. Mechanical Alloying and Electronic Simulations of (MgH2+M) Systems (M=Al, Ti, Fe, Ni, Cu and Nb) for Hydrogen Storage. Int. J. Hydrogen Energy 2004, 29, 73−80. (31) Bouaricha, S.; Dodelet, J. P.; Guay, D.; Huot, J.; Boily, S.; Schulz, R. Hydriding Behavior of Mg−Al and Leached Mg−Al Compounds Prepared by High-Energy Ball-Milling. J. Alloys Compd. 2000, 297, 282−293. (32) Andreasen, A. Hydrogenation Properties of Mg−Al Alloys. Int. J. Hydrogen Energy 2008, 33, 7489−7497. (33) Crivello, J. C.; Nobuki, T.; Kuji, T. Improvement of Mg−Al Alloys for Hydrogen Storage Applications. Int. J. Hydrogen Energy 2009, 34, 1937−1943. (34) de Boer, F. R.; Boom, R.; Mattens, W. C. M.; Miedema, A. R.; Niessen, A. K. Cohesion in Metals; Elsevier: North-Holland, 1988. (35) Zhong, H. C.; Wang, H.; Ouyang, L. Z.; Zhu, M. Microstructure and Hydrogen Storage Properties of Mg−Sn Nanocomposite by Mechanical Milling. J. Alloys Compd. 2011, 509, 4268−4272. (36) Zlotea, C.; Sahlberg, M.; Ö zbilen, S.; Moretto, P.; Andersson, Y. Hydrogen Desorption Studies of the Mg24Y5−H System: Formation of Mg Tubes, Kinetics and Cycling Effects. Acta Mater. 2008, 56, 2421− 2428. (37) Liang, G.; Huot, J.; Boily, S.; Van Neste, A.; Schulz, R. Catalytic Effect of Transition Metals on Hydrogen Sorption in Nanocrystalline Ball Milled MgH2−Tm (TmTi, V, Mn, Fe and Ni) Systems. J. Alloys Compd. 1999, 292, 247−252. 12096


hydriding of Mg Base Solid Solutions with

Recommend Documents