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Hydrogen-Induced Reversible Phase Transformations and Hydrogen Storage Properties of Mg−Ag−Al Ternary Alloys Yanshan Lu,† Hui Wang,†,‡ Jiangwen Liu,†,‡ Ziming Li,† Liuzhang Ouyang,†,‡ and 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 ‡ China-Australia Joint Laboratory for Energy & Environmental Materials, South China University of Technology, Guangzhou, 510641, China ABSTRACT: Overly stable thermodynamics and sluggish kinetics hinder the practical applications of Mg-based hydrogen storage alloys. Compositional and structural modifications are important strategies in tuning these hydrogen storage properties. In this study, Mg-based Mg−Ag−Al ternary alloys were investigated to explore their performance as hydrogen storage alloys. Mg80Ag15Al5 exhibits a reaction pathway that differs from that in pure Mg, in which the intermediate phase, consisting of a new ternary solid solution MgAg(Al), reacts with MgH2 during dehydriding and contributes to an increase in the dehydriding equilibrium pressure (0.22 MPa at 300 °C) and to a reversible hydrogen storage capacity of 1.7 wt %. Adjusting the composition to Mg85Ag5Al10 results in a reversible hydrogen storage capacity of approximately 3.8 wt % and an elevated equilibrium pressure (0.26 MPa at 300 °C). These Mg−Ag−Al ternary alloys also show enhanced hydrogen sorption kinetics relative to that of Mg, and the apparent activation energies for hydrogenation and dehydrogenation of the Mg85Ag5Al10 sample are lowered to 74.5 kJ mol−1 and 124.8 kJ mol−1, respectively. This work demonstrates the possibility of exploring new hydrogen storage alloys via creating different reaction pathways for hydrogen-induced reversible phase transformations.

1. INTRODUCTION Mg is a promising metallic hydrogen storage material owing to its high hydrogen storage capacity and its abundance in the Earth’s crust. However, the sluggish hydrogen sorption kinetics of this metal and the high thermodynamic stability of MgH2, which results in the need for elevated dehydriding temperature, hinder its practical applications.1−4 Several methods, such as nanostructuring,5−8 catalyzing,9−14 and forming Mg-based composites15−18 have been researched as means of addressing these issues, and substantial progress has been made in solving the kinetics problem. The challenge of the excessively high dehydrogenation temperature (up to 300 °C) of MgH2 is attributed to the strong ionic characteristics of the Mg−H bond. Compositional and structural modifications have been proven to be effective in destabilizing the thermodynamics of MgH2.19,20 One approach is to form ternary or multinary hydrides with reduced stabilities by alloying Mg with other elements.21−25 In 1968, Reilly et al.19 reported that a Mg2Ni alloy is transformed to Mg2NiH4 upon hydrogenation and that the hydrogen desorption enthalpy of Mg2NiH4 (ΔH = 64.5 kJ mol−1 H219) is lower than that of MgH2 (ΔH = 74.7 kJ mol−1 H226). However, it is difficult to search these types of hydrides because the majority of Mgbased alloys show no tendency to form ternary or multinary hydrides in the presence of hydrogen and instead decompose when reacting with hydrogen.27−32 © XXXX American Chemical Society

Unlike Ni, other elements can form compounds with Mg that destabilize the thermodynamics of MgH2 by altering the dehydriding reaction pathway. These elements include Al,33−35 Cu,36−38 Ag,39−41 Cd,42,43 and In.44−47 In the Mg−Cu binary system, two intermetallic compounds exist: Mg2Cu and MgCu2. Reilly et al.36 found that the former would be disproportionate to MgCu2 and MgH2 phases during hydrogenation and could be recovered upon removing H2, with a lower dehydrogenation enthalpy (ΔH = 70.0 kJ mol−1 H2) compared to that of MgH2. This reaction pathway, which is different from that of pure MgH2, explains the thermodynamic destabilization of MgH2. Liang48 was the first to report that alloying In with Mg to form a Mg(In) solid solution was effective at reducing the stability of MgH2. The hydriding reaction of Mg(In) leads to the formation of MgH2 and MgIn, with the original structure being recovered after dehydrogenation.44 It has been reported that the hydrogen desorption enthalpy is reduced to 65.2 kJ mol−1 H2 when the concentration of In dissolved in the Mg lattice is 10 at. %.45 In other works, partly placing In with Al or Cd produced the ternary solid solutions Mg(In, Al)45 and Mg(In, Cd),49 which exhibited reversible structural transformations and altered thermodynamics. Recently, two new Received: August 16, 2016 Revised: November 6, 2016 Published: November 8, 2016 A

DOI: 10.1021/acs.jpcc.6b08291 J. Phys. Chem. C XXXX, XXX, XXX−XXX

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The Journal of Physical Chemistry C

Figure 1. Observed and calculated XRD patterns, their difference plots, and all possible Bragg positions of all refined phases for a Mg80Ag15Al5 sample at different states: (a) As-prepared following 20 h milling, (b) hydrogenated under 5 MPa H2 at 320 °C for 12 h, and (c) dehydrogenated at 320 °C for 12 h.

reversible transitions during de/hydrogenation cycles.50,51 Compared with MgH2, both the dehydriding thermodynamics

ternary intermetallic compounds, Mg14In3Ni3 and Mg2InNi, were found in the Mg−In−Ni system, showing completely B


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Table 1. Lattice Constants and Abundance Obtained from Rietveld Refinement for Phases in a Mg80Ag15Al5 Sample at Different States lattice constants Mg80Ag15Al5 before milling (Rwp = 11.70%, S = 1.65)

as-milled Mg80Ag15Al5 (Rwp = 10.10%, S = 1.40) hydrogenated Mg80Ag15Al5 (Rwp = 9.82%, S = 1.69) dehydrogenated Mg80Ag15Al5 (Rwp = 11.90%, S = 1.68)

phase

a (Å)

Mg Mg54Ag17 Al Mg(Al) Mg54(Ag, Al)17 MgH2 MgAg(Al) Mg(Al) Mg54Ag17

3.2231(2) 14.2032(2) 4.0496(6) 3.2082(3) 14.2892(7) 4.5341(1) 3.3129(8) 3.1995(2) 14.1996(7)

c (Å)

abundance (wt %)

14.2309(3)

5.2020(2) 14.6098(1)

22 74 4 25 75 23 77 27 73

14.2349(9)

14.2311(1)

5.1986(2) 14.6585(6) 3.0317(6) 5.1949(8) 14.5980(6)

comparison to MgH2 as a result of the altered reversible de/ hydriding reactions. The associated mechanism was revealed by investigating the phase transformation process.

and kinetics are improved by the novel phase transition between these two Mg−In−Ni alloys, leading to a minimum hydrogen desorption temperature of 230 °C. Similar to the Mg−Cu and Mg−In systems, the Mg−Ag binary alloys Mg3Ag39 and Mg4Ag40 can be fully recovered from the hydrogenated products MgH2 and MgAg after dehydrogenation, and the dehydrogenation enthalpies are decreased to 69.8 and 69.1 kJ mol−1 H2 for Mg3Ag and Mg4Ag, respectively. To further improve the thermodynamics of the Mg−Ag system, Zhang et al.52 added In to synthesize a (Mg, In)3Ag ternary solid solution alloy and successfully lowered the reaction enthalpy to 62.6 kJ mol−1 H2. The reversible de/hydriding reaction of the (Mg, In)3Ag solid solution can be expressed as in eq 1, and the disproportionation of such alloy is similar to the hydrogenation of Mg2Cu.36 In this system, it is evident that In would be dissolved in MgAg in the hydrogenated state. (Mg, In)3 Ag + H 2 ↔ MgH2 + (Mg, In)Ag

b (Å)

2. EXPERIMENTAL SECTION Mg−Ag−Al ternary alloys were prepared using a two-step approach. Mg and Ag powders were first mixed and sintered at 600 °C for 20 h in a tube furnace under pure argon. Subsequently, the sintered Mg−Ag powder was milled with Al powder at 300 rpm with a ball-to-powder mass ratio of 20:1 for 20 h (milling for 30 min followed by a 30 min pause, repeated 40 times) under an argon atmosphere in a QM-3SP2 planetary mill. X-ray diffraction (XRD) patterns were obtained with a Rigaku MiniFlex 600 X-ray diffractometer, using Cu Kα radiation. Si powder (99.9999% purity) was added into XRD samples as an internal reference to correct the instrumental zero-shift. These XRD patterns were refined by the Rietveld method using the program FULLPROF.53 Scanning electron microscopy (SEM) observations were performed with a Zeiss Supra 40/VP. Transmission electron microscopy (TEM) and scanning transmission electron microscopy (STEM) images were obtained using a JEOL JEM-2100 transmission electron microscope operating at 200 kV. In situ TEM observations were carried out during the heating process to assess the phase transformations. The hydrogen storage properties of the Mg−Ag−Al alloys were measured using an automatic Sievert-type apparatus (GRC, Advanced Materials Corp.). Pressure−composition isotherms (PCI) were acquired at specific temperatures and at hydrogen pressures ranging from 4 to 0.002 MPa. Isothermal hydrogen sorption kinetics was determined at various temperatures, using an initial pressure of approximately 4 MPa for hydrogenation and applying a vacuum during dehydrogenation. All sample handlings were performed in a glovebox filled with argon gas.

(1)

Despite the above success, In is a rare and expensive metal, and the thermodynamic destabilizing effect is still limited in these In alloys. Thus, it is necessary to explore other alloying elements capable of generating larger destabilization effects and with the advantage of lower cost. In this regard, Al may be a possible alloying element. Although the equilibrium solubility of Al in Mg is negligible at room temperature, it can be extended up to 8 at.% using ball milling method to form a Mg(Al) solid solution, which is able to destabilize the de/ hydriding thermodynamics.35 The Ag−Al binary phase diagram also indicates that Al can dissolve in the Ag lattice with a maximum solubility of 10 at.%. In addition, it has been found that alloying Mg with Ag to form Mg−Ag binary alloys can decrease the dehydrogenation enthalpies compared to pure Mg.39,40 Therefore, it is expected that Al may be able to dissolve in Mg−Ag alloys to generate new reversible de/ hydriding routes with destabilized thermodynamics. As an added benefit, Al is the most abundant metal in the Earth’s crust and is inexpensive. It should be noted that the de/hydriding kinetic mechanism also changes with variations in the de/hydriding reaction pathway, because additional intermetallic phases can form or disappear via solid state phase transformations during the de/ hydriding process.45,51 In fact, such phase transformations may be the key factor determining the de/hydriding kinetics. Therefore, it is also necessary to assess the kinetics and phase transformation mechanism. In this work, we explored the hydrogen storage performance of Mg−Ag−Al ternary alloys. These alloys exhibited destabilized dehydriding thermodynamics and improved kinetics in

3. RESULTS AND DISCUSSION 3.1. Phase Transformations of Mg−Ag−Al Ternary Alloys during De/Hydriding. Figure 1 shows the XRD patterns and Rietveld refinement results for a Mg80Ag15Al5 sample at different states. Prior to ball milling, the powder mixture consisted of Mg, Mg54Ag17, and Al phases. As can be seen from Figure 1a, the as-prepared sample generated broad diffraction peaks attributed to Mg and Mg54Ag17 phases, while there is no evidence for Mg−Al or Ag−Al intermetallic compounds or Al phase. These results suggest that Al dissolved in Mg and/or Mg54Ag17 phases. The lattice constants obtained C


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Figure 2. Backscattering SEM images of a Mg80Ag15Al5 sample: (a) hydrogenated under 5 MPa H2 at 320 °C for 12 h and (b) dehydrogenated at 320 °C for 12 h.

reactions of the Mg80Ag15Al5 sample can be expressed as follows.

from the Rietveld refinement are summarized in Table 1, and these values verify that Al dissolved in both Mg and Mg54Ag17 to form Mg(Al) and Mg54(Ag, Al)17 solid solutions during the milling process. Therefore, a new ternary solid solution alloy Mg54(Ag, Al)17 was successfully prepared by our two-step method consisting of sintering followed by ball milling. The quantities of the Mg(Al) and Mg54(Ag, Al)17 phases were 25 and 75 wt %, respectively. It should be pointed out that the broad peak in the range of 18−28° in the XRD pattern of the as-milled sample is due to the glass holder used for the XRD test. As reported in previous work,39,41 Mg3Ag when hydrogenated forms MgH2 and MgAg phases. In the case of the present Mg80Ag15Al5 sample, as shown in Figure 1b, the hydrogenated products consist of MgH2 and MgAg(Al) solid solution. The presence of MgAg(Al) can be identified because the positions of the MgAg diffraction peaks relative to the MgH2 peaks are slightly shifted toward higher angles compared with those in the reference pattern (MgAg: no. 65-5904). In addition, the lattice parameters are reduced compared to those of MgAg phase in the hydrogenated Mg6Ag (a = 3.3232(5) Å),52 further verifying the formation of MgAg(Al) solid solution. Our prior research has also demonstrated that Al is able to dissolve in the β phase (MgIn compound) within the hydrogenated Mg90In5Al5 solid solution alloy.45 Therefore, another new ternary solid solution, MgAg(Al), is formed as a result of the hydrogen-induced reaction. Figure 2a presents a backscattering SEM image of a hydrogenated Mg80Ag15Al5 sample, in which two obvious contrasts can be seen. Based on the XRD results, the fine, bright particles in this image correspond to MgAg(Al), containing a high level of elemental Ag, that are distributed throughout the gray MgH2 matrix. Figure 1c demonstrates that both Mg(Al) solid solution and Mg54Ag17 result from complete dehydrogenation. As shown in Table 1, the lattice parameters of the Mg54Ag17 phase in the dehydrogenated Mg80Ag15Al5 sample are close to those of the sintered Mg−Ag and Al powder mixture. However, the lattice parameters of the Mg(Al) phase in the dehydrogenated Mg80Ag15Al5 sample are less than those of the ball-milled sample, indicating that Al is preferentially dissolved in Mg phase rather than in Mg54Ag17 during dehydrogenation. Figure 2b provides an SEM image of the microstructure of the dehydrogenated Mg80Ag15Al5. In this image, the gray matrix is Mg(Al), while the bright spots are attributed to Mg54Ag17 phase. Based on the above results, the initial de/hydriding

Mg(Al) + Mg54(Ag, Al)17 + H 2 → MgH2 + MgAg(Al) (2)

MgH2 + MgAg(Al) ↔ Mg(Al) + Mg54Ag17 + H 2

(3)

Although the initial phase constituents of the ball-milled Mg80Ag15Al5 were not regenerated following the first de/ hydrogenation cycle, our results demonstrate that the subsequent de/hydriding process is reversible and can be expressed as in eq 3. Varying the relative proportions of the elements, a different Mg−Ag−Al sample, Mg85Ag5Al10, was synthesized, and the XRD patterns and Rietveld refinement results for Mg85Ag5Al10 at different states are shown in Figure 3 and Table 2. The phase constituents of the as-prepared material were evidently similar to those described above for Mg80Ag15Al5 and consist of solid solutions Mg(Al) (67 wt %) and Mg54(Ag, Al)17 (33 wt %) (Figure 3a). In the case of the hydrogenated Mg85Ag5Al10, Figure 3b demonstrates that MgH2 and MgAg(Al) solid solution were present in addition to a ternary phase Mg4Al6Ag. Figure 3c shows that the hydrogenated products also transform to Mg(Al) and Mg54Ag17 after the first de/hydrogenation cycle. The subsequent, reversible hydrogen ab/desorption can be described as follows. MgH2 + MgAg(Al) + Mg4Al 6Ag ↔ Mg(Al) + Mg54Ag17 + H 2

(4)

This reversible microstructural evolution of the Mg85Ag5Al10 sample was also investigated on the basis of SEM observations. Two evident contrasts appear in the backscattering SEM micrograph of the Mg85Ag5Al10 sample in the fully hydrogenated state (Figure 4a). Here, the gray matrix is MgH2, while the bright spots embedding in the MgH2 particles are believed to consist of MgAg(Al) or Mg4Al6Ag phases, although these two phases cannot be precisely distinguished from one another in this image. Figure 4b displays the SEM image of the dehydrogenated Mg85Ag5Al10, which is almost identical to that of the dehydrogenated Mg80Ag15Al5, as the result of the equivalent phase compositions of these two samples. To further demonstrate that Al dissolves in the Mg rather than the Mg54Ag17 during dehydrogenation, a STEM-EDS mapping analysis was conducted. It can be seen from Figure 5 that the D


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Figure 3. Observed and calculated XRD patterns, their difference plots, and all possible Bragg positions of all refined phases for a Mg85Ag5Al10 sample at different states: (a) As-prepared following 20 h milling, (b) hydrogenated under 5 MPa H2 at 320 °C for 12 h, and (c) dehydrogenated at 320 °C for 12 h.

3.2. Hydrogen Storage Properties of Mg−Ag−Al Ternary Alloys. Figure 6a shows the hydrogen desorption

distribution of Al is different from that of Ag, but close to that of Mg, verifying that Mg(Al) solid solution was formed. E


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Table 2. Lattice Constants and Abundance Obtained from Rietveld Refinement for Phases in a Mg85Ag5Al10 Sample at Different States lattice constants Mg85Ag5Al10 before milling (Rwp = 12.50%, S = 1.92)

as-milled Mg85Ag5Al10 (Rwp = 9.93%, S = 1.55) hydrogenated Mg85Ag5Al10 (Rwp = 14.90%, S = 2.25)

dehydrogenated Mg85Ag5Al10 (Rwp = 16.80%, S = 2.60)

phase

a (Å)

Mg Mg54Ag17 Al Mg(Al) Mg54(Ag, Al)17 MgH2 MgAg(Al) Mg4Al6Ag Mg(Al) Mg54Ag17

3.2103(2) 14.1923(4) 4.0452(8) 3.2037(1) 14.2986(7) 4.5183(8) 3.3048(2) 14.3695(1) 3.1887(9) 14.1917(5)

b (Å)

c (Å)

abundance (wt %)

14.2154(4)

5.2118(5) 14.6057(8)

62 30 8 67 33 52 25 23 69 31

14.2253(1)

14.2016(1)

5.1971(3) 14.6841(1) 3.0223(1)

5.1820(1) 14.5566(3)

Figure 4. Backscattering SEM images of a Mg85Ag5Al10 sample: (a) hydrogenated under 5 MPa H2 at 320 °C for 12 h and (b) dehydrogenated at 320 °C for 12 h.

Figure 5. STEM-EDS mapping of a dehydrogenated Mg85Ag5Al10 sample.

determined to be 84.3 kJ mol−1 H2 and 153.5 J K−1 mol−1 H2, respectively, based on the van’t Hoff plot in Figure 6b. Using the values of ΔH and ΔS, the equilibrium pressure of Mg80Ag15Al5 at 300 °C was calculated to be 0.22 MPa, a value that exceeds the pressures of 0.16 MPa for pure Mg51 and 0.17 MPa for Mg3Ag alloy.41 This result indicates that the dehydriding thermodynamics of the Mg−Ag−Al ternary system is destabilized compared to that in pure MgH2. As shown in Figure 7a, the hydrogen storage capacity of the Mg85Ag5Al10 sample was increased to 3.8 wt %, a value more than twice that of the Mg80Ag15Al5, even though there was only

PCI curves obtained from a Mg80Ag15Al5 sample at various temperatures, from which the reversible hydrogen storage capacity of this sample was determined to be approximately 1.7 wt %. The intensities of the MgAg(Al) peaks in Figure 1b are much greater than those of the MgH2 peaks, which explains the relatively low hydrogen storage capacity. Only one slightly sloping plateau is evident in each PCI curve, suggesting onestep dehydrogenation in the case of the Mg80Ag15Al5 via the reaction in eq 3. Adopting the pressure at the midpoint of the plateau as the equilibrium pressure, the hydrogen desorption enthalpy, ΔH, and entropy, ΔS, of Mg80Ag15Al5 were F


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Figure 6. (a) Pressure−composition isotherms (PCI) and (b) van’t Hoff plot for hydrogen desorption of a Mg80Ag15Al5 sample.

Figure 7. (a) Pressure−composition isotherms (PCI) and (b) van’t Hoff plots for hydrogen desorption of a Mg85Ag5Al10 sample.

a 5 at.% increase in the Mg content. Two plateaus are evident in each PCI curve, suggesting two-step dehydrogenation. One of these plateaus is flat, while the other one is sloping. In the sloping region, the dehydriding reaction ΔH and ΔS values were determined to be 80.3 kJ mol−1 H2 and 148.2 J K−1 mol−1 H2, respectively (Figure 7b). From these, the equilibrium pressure in the sloping region at 300 °C was calculated to be 0.26 MPa. This pressure is even higher than that of the Mg80Ag15Al5, which indicates that the dehydriding pressure associated with the Mg−Ag−Al ternary alloys is elevated to a greater extent when MgH2 reacts with MgAg(Al) and Mg4Al6Ag phases compared to a scenario in which MgH2 reacts only with MgAg(Al) phase. In the flat region, the equilibrium pressure at 300 °C was found to be 0.16 MPa, a value the same as that for pure Mg. The ΔS for these two Mg−Ag−Al alloys are much higher than the standard dehydriding reaction entropy change of ∼130 J K−1 mol−1 H2. The higher ΔS may be caused by the dehydrogenation involving more phases. The dehydriding reaction of the Mg−Ag−Al ternary alloy is more complicated than that of Mg (eq 4). In the dehydriding reaction, two additional compounds are involved into the decomposition of

MgH2, and the Mg54Ag17 phase and Mg(Al) solid solution instead of Mg are formed in the dehydrogenation, which would cause the different enthalpy and entropy change from those of pure MgH2. This topic needs to be further investigated. Considering that the hydrogen storage capacity of the Mg80Ag15Al5 sample was less than 2.0 wt %, the kinetics investigation was performed only for the Mg85Ag5Al10 sample, and Figure 8 presents the isothermal hydrogenation and dehydrogenation kinetics curves for this material. At 260 °C, the material exhibits good hydrogen absorption kinetics, reaching its maximum hydrogen content within 20 min (Figure 8a). Even at a lower temperature of 200 °C, approximately 2.5 wt % hydrogen is absorbed within 60 min. During the dehydrogenation kinetics measurements, the two-step dehydriding process could not be distinguished, possible due to the different test conditions employed during PCI and isothermal kinetics assessments. The initial pressure during measurements of the isothermal dehydriding kinetics was almost nil, which may have promoted the simultaneous occurrence of the two hydrogen desorption reactions. It is, however, evident from Figure 8b that the hydrogen is completely released over the G


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3.3. Dehydriding Mechanism of Mg−Ag−Al Ternary Alloys. The reversible de/hydriding process of the Mg80Ag15Al5 sample can be expressed as in eq 3, and this modified dehydriding reaction pathway explains the higher hydrogen desorption pressure for this material compared to pure Mg. The reversible reaction pathway of Mg80Ag15Al5 is quite similar to that of (Mg, In)3Ag ternary alloy and so is not discussed here in detail. In contrast, the reversible reaction path of Mg85Ag5Al10 is more complicated. To further investigate the dehydriding mechanism, quasi in situ XRD analyses of the Mg85Ag5Al10 sample were performed at various dehydriding stages during the PCI test process. The details of the experimental process have been described elsewhere.50 The results are shown in Figure 9, in which the profiles labeled a, b, c, d, and e correspond to the dehydriding points a, b, c, d, and e, respectively, indicated on the PCI curve measured at 290 °C (Figure 7a).

Figure 8. (a) Isothermal hydrogenation kinetic curve and (b) isothermal dehydrogenation kinetics curves for a Mg85Ag5Al10 sample. Figure 9. XRD patterns of a Mg85Ag5Al10 sample at different PCI desorption stages at 290 °C, as indicated in Figure 7a.

span of 10 min at 380 °C. In addition, at 320 °C, approximately 3.3 wt % hydrogen can be desorbed in 2 h. The kinetic curves of the Mg85Ag5Al10 sample were analyzed using the Johnson−Mehl−Avrami−Kolmogorov (JMAK) model54,55 and the Arrhenius equation, and the apparent activation energies for hydrogen absorption/desorption were calculated. The apparent activation energy for the hydrogenation was determined to be 74.5 kJ mol−1, a value below the energies of 80−120 kJ mol−1 for pure Mg56 and close to the value of 73 kJ mol−1 for the Mg90In5Al5 alloy.45 The apparent activation energy for the overall dehydrogenation of Mg85Ag5Al10 was found to be 124.8 kJ mol−1, a value that is much lower than the 160 kJ mol−1 for pure MgH2,56,57 but higher than the 89.8 kJ mol−1 for the Mg6Ag binary alloy.52 Although the kinetics of the Mg85Ag5Al10 sample is superior to that of pure Mg, it remains slow. This can possibly be attributed to the reduced rate of formation of an intermetallic phase and a solid solution as the result of the slow diffusion of Al. Our previous work demonstrated that the sluggish dehydriding kinetics of the Mg(In, Al) solid solution was due to the slow diffusion of In and H resulting from the poor mobility of Al.45 Therefore, in the case of the Mg−Ag−Al system, alloying elements or additives that can enhance the diffusion of Al should be considered in future.

As noted in the previous section, MgH2, MgAg(Al) solid solution and Mg4Al6Ag are formed at the fully hydrogenated state corresponding to point a in Figure 7a. When the dehydrogenation proceeds to point b (that is, the midpoint of the sloping region), the Mg4Al6Ag phase is absent; the MgH2 and MgAg(Al) solid solution peaks are less intense, and the presence of Mg54Ag17 is observed. At point c (the end point of the sloping plateau), the MgAg(Al) solid solution peaks completely disappear, while the Mg(Al) solid solution is present. Therefore, during the dehydriding process occurring in the high pressure sloping region, a small portion of MgH2 reacts with MgAg(Al) solid solution and Mg4Al6Ag, which leads to the formation of Mg(Al) solid solution and Mg54Ag17 with a minor release of hydrogen. The dehydriding reaction can be described as in eq 4. As the dehydrogenation progresses to the midpoint of the flat region (point d), the Mg54Ag17 phase remains almost unchanged, while the MgH2 peaks are further weakened, accompanied by an intensification of the Mg(Al) peaks. At point e, MgH2 disappears completely, and the Mg(Al) peak intensities are maximized. It is evident that the majority of the MgH2 decomposition occurs in this flat region, explaining the relatively prolonged plateau. Accordingly, the dehydriding H


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Figure 10. TEM bright field images, electron diffraction patterns, and high-resolution TEM image of a Mg85Ag5Al10 sample: (a) hydrogenated and (b) dehydrogenated.

These Mg−Ag−Al alloys generate higher dehydriding pressures compared with that for pure Mg, with values as high as 0.22 MPa for the dehydrogenation of Mg80Ag15Al5 and 0.26 MPa for the first step dehydrogenation of Mg85Ag5Al10 at 300 °C. These ternary alloys also show improved hydrogen absorption kinetics, and the apparent activation energy for the hydrogenation of the Mg85Ag5Al10 sample is lowered to 74.5 kJ mol−1. Although the dehydriding kinetics of these Mg−Ag−Al ternary alloys are superior to that of pure Mg, the kinetics remain slow because poor Al mobility retards the diffusion of Mg, Ag, and H. Therefore, future research should involve alloying elements or additives capable of enhancing the diffusion of Al.

reaction corresponding to the low pressure plateau can be expressed as follows. MgH2 + Mg(Al) ↔ Mg(Al) + H 2

(5)

It should be noted that the Mg(Al) solid solution on the right side of eq 5 will have a lower Al concentration than that on the left side. In situ TEM was also carried out to further study the dehydriding mechanism of the Mg85Ag5Al10 sample. A bright field image of hydrogenated Mg85Ag5Al10 is shown in Figure 10a, in which the three regions of different contrasts are assigned to MgH2, MgAg(Al), and Mg4Al6Ag, respectively, based on the electron diffraction patterns. This same sample was subsequently heated in the TEM holder to release hydrogen, and the image contrast was found to change as the temperature increased. The MgH2 and MgAg(Al) initially reacted with Mg4Al6Ag to form Mg54Ag17 and Mg(Al), after which the remaining MgH2 decomposed to form Mg, which reacted with Mg(Al) to generate Mg(Al) solid solution with a higher weight percent of Mg. It should be noted that the Mg54Ag17 phase was not indexed in the selected area electron diffraction pattern of the dehydrogenated Mg85Ag5Al10 (Figure 10b) due to the lower diffraction intensity compared to that of the Mg(Al) solid solution phase, although this material can be detected in the high-resolution TEM image.

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. Tel.: +86-20-8711 3924. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work was financially supported by the National Natural Science Foundation of China (Nos. 51471070, 51431001, and 5157010144), by the Foundation for Innovative Research Groups of the National Natural Science Foundation of China (No. 51621001), by the Guangdong Natural Science Foundation (2014A030313222 and 2016A030312011), and by the International Science & Technology Cooperation Program of China (2015DFA51750).

4. CONCLUSIONS Mg−Ag−Al ternary alloys were successfully prepared by sintering Mg and Ag powders followed by ball milling with Al powder. In the case of a Mg80Ag15Al5 sample, the product was found to consist of Mg(Al) and Mg54(Ag, Al)17 solid solutions after ball milling, which transformed to MgH2 and MgAg(Al) during hydriding. Following an initial dehydrogenation, Mg(Al) and Mg54Ag17 were obtained, while subsequent de/hydriding resulted in a reversible phase transition with a hydrogen storage capacity of 1.7 wt %. Another sample, Mg85Ag5Al10, exhibited two-step dehydrogenation and a reversible hydrogen storage capacity of approximately 3.8 wt %. In addition to MgH2 and MgAg(Al), the ternary phase Mg4Al6Ag was observed in the hydrogenated Mg85Ag5Al10.

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REFERENCES

(1) Crivello, J. C.; et al. Review of Magnesium Hydride-Based Materials: Development and Optimisation. Applied Physics A-Materials Science & Processing 2016, 122, 97. (2) Jia, Y.; Sun, C.; Shen, S.; Zou, J.; Mao, S. S.; Yao, X. Combination of Nanosizing and Interfacial Effect: Future Perspective for Designing Mg-Based Nanomaterials for Hydrogen Storage. Renewable Sustainable Energy Rev. 2015, 44, 289−303. (3) Zhu, M.; Lu, Y. S.; Ouyang, L. Z.; Wang, H. Thermodynamic Tuning of Mg-Based Hydrogen Storage Alloys: A Review. Materials 2013, 6, 4654−4674.

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