Trans. Nonferrous Met. Soc. China 28(2018) 1828-1837

Hydrogen storage performances of as-milled REMg₁₁Ni (RE=Y, Sm) alloys catalyzed by MoS₂

Yang-huan ZHANG^1,2, Wei ZHANG^1,2, Ze-ming YUAN^1,2, Wen-gang BU², Yan QI², Xiao-ping DONG³, Shi-hai GUO²

1. Key Laboratory of Integrated Exploitation of Baiyun Obo Multi-Metal Resources, Inner Mongolia University of Science and Technology, Baotou 014010, China;

2. Department of Functional Material Research, Central Iron and Steel Research Institute, Beijing 100081, China;

3. Department of Mechanical Engineering, Hebei University, Baoding 071002, China

Received 1 June 2017; accepted 30 September 2017

Abstract: To compare the hydrogen storage performances of as-milled REMg₁₁Ni-5MoS₂ (mass fraction) (RE=Y, Sm) alloys, which were catalyzed by MoS₂, the corresponding alloys were prepared. The hydrogen storage performaces of these alloys were measured by various methods, such as XRD, TEM, automatic Sievert apparatus, TG and DSC. The results reveal that both of the as-milled alloys exhibit a nanocrystalline and amorphous structure. The RE=Y alloy shows a larger hydrogen absorption capacity, faster hydriding rate, lower initial hydrogen desorption temperature, superior hydrogen desorption property, and lower hydrogen desorption activation energy, which is thought to be the reason of its better hydrogen storage kinetics, as compared with RE=Sm alloy.

Key words: Mg-based alloy; ball milling; catalyst; rare earth element; hydrogen storage performance

1 Introduction

As is well known, the end of the Stone Age is not because of the depletion of stone. Likewise, the termination of the fossil fuel era does not need to wait until it dries up. In other words, human beings cannot be forced to abandon the car due to the exhaustion of fossil fuels; instead, mankind is very likely to initiatively give up fossil fuels in the case of the great development of automobile. Hydrogen is considered to be the best fuel for fuel cells owing to its outstanding versatility, utilization efficiency, safety, environment compatibility and inexhaustible reserves [1-3]. However, hydrogen storage is the major technology barrier to the introduction of hydrogen economy [4,5]. That is to say, the schedule of acclaiming hydrogen fuel cell vehicle as a wide commercial application of the 21st century will principally depend on hydrogen storage technology [6]. Metal hydride hydrogen storage, one of hydrogen storage methods, is regarded to be the most promising alternative to satisfy the requirements for mobile application [4]. Mg and Mg-based alloys, as promising hydrogen storage materials, have been extensively investigated [7-9] due to some of their unique advantages, such as high gravimetric (7.6%, mass fraction) and volumetric (110 kg/m³) hydrogen storage densities, good reversibility in de-/hydrogenation processes and low cost [10]. However, some of their disadvantages, such as the relatively high thermal stability and slow de-/hydrogenation kinetics, impede their commercial application in hybrid electric vehicle [11]. Therefore, the hydrogen storage thermodynamics and kinetics of Mg and Mg-based alloys need to be further ameliorated.

It has been ascertained that the change of structures can dramatically affect the de-/hydrogenation thermo- dynamics and kinetics of Mg and Mg-based alloys [12,13]. Especially, reducing the grain size of Mg- based alloys far below micron size can significantly ameliorate their de-/hydrogenation characteristics [14]. Mechanical milling [15,16] and melt spinning [17] are generally accepted as two commonly used strategies for obtaining nanostructure and ameliorating hydrogen storage performance. However, whether ball mill can improve hydrogen storage thermodynamics is still a questionable issue. LASS [18] considered that ball milling, in theory, can improve the thermodynamics of de-/hydrogenation reactions by introducing capillarity effects when particle size is reduced to a small enough size. While AGARWAL et al [19] stated that ameliorating thermodynamics by ball milling is possible only when the particle size is less than 5 nm, which has not yet been achieved by using mechanical milling for now.

Furthermore, it was confirmed that the dissociation of hydrogen molecules on Mg surface requires a fairly large energy (1.15 eV), which is the rate controlling factor for the hydrogenation of Mg. Catalysts may reduce this dissociation energy for hydrogen storage in Mg. For example, the addition of transition metals Pd, Cu, Ni and Co can reduce it to 0.39, 0.56, 0.06, and 0.03 eV, respectively [20]. It was proven that the partial substitution of rare earth elements (La, Ce, Pr, Nd, Y, Sm) for Mg [21-23] and the addition of transition metals (Ni, Ti, V, Fe, Cu) considerably ameliorate the thermo- dynamics of de-/hydrogenation [20,24,25]. Theoretically, the substitutional atoms in MgH₂ can weaken Mg—H bond through interactions between the valance electron of H and the unsaturated d/f electron shells of transition metals as catalyst, and thus the desorption behavior of MgH₂ can be ameliorated greatly [26]. Generally, rare earth elements [23], transition metallic elements [27] as well as their oxides [28], fluorides [29], halides [7], hydrides [21] and intermetallic compounds [19] are verified to be good catalysts which can destabilize MgH₂ and enhance the de-/hydriding rate of alloy hydrides. SADHASIVAM et al [14] investigated the effect of adding Mm (Ce and La as the dominant components) oxide on hydrogen desorption kinetics of MgH₂ and the results revealed that Mm-oxide is an effective catalyst for improving the hydrogen sorption behavior of MgH₂. The initial desorption temperature for 5% (mass fraction) Mm-oxide catalyzed MgH₂ was decreased from 654 (ball-milled) to 578 K, viz. a decrease by 76 K. DARYANI et al [30] reported that adding catalyst TiO₂ can significantly enhance the milling efficiency and accelerate the size reduction of MgH₂ during ball milling. Moreover, the addition of 6% (molar fraction) TiO₂ markedly improved the hydrogen absorption kinetics and made the decomposition temperature of as-milled MgH₂ decreased by 100 K. MoS₂ is a naturally occurring layered solid and has been widely used as dry lubricant and catalyst [31-33]. WANG et al [34] have added MoS₂ into 2LiBH₄-MgH₂ system as catalyst, which showed a good catalysis in both de-/hydrogenation kinetics and thermodynamics. HUANG et al [35] studied the catalysis of Mo, MoO₂, MoO₃ and MoS₂ in catalyzing the hydrolysis of Mg composites in seawater, one of the promising hydrogen generation technologies, and got a good outcome.

In the present work, a trace of MoS₂ was added into the alloys. According to the recent result made by EL-ESKANDARANY et al [36], 5% (mass fraction) is thought to be the optimum amount of added MoS₂. Ball milling technology has been utilized to synthesize REMg₁₁Ni-5MoS₂ (RE=Y, Sm) alloys. A comparison of structure and hydrogen storage performance of the as-milled alloys containing different rare earth elements was performed.

2 Experimental

The REMg₁₁Ni (RE=Y, Sm) alloy ingots were synthesized in a vacuum induction furnace under the protective atmosphere of helium to prevent the evaporation of Mg element. Prior to testing de-/hydriding performances, experimental samples were mechanically pulverized into powders of about 50 μm in diameter, and then 5% (mass fraction) MoS₂ was added into the alloy powder. The hardened steel balls (210 g) and mixed powder (6 g) were homogeneously mixed under argon protection in Cr-Ni stainless steel vials. The mechanical mill is a planetary-type mill with 135 r/min of disc revolution speed. In each cycle, 50 min milling and 10 min resting were programmed with a total milling time of 20 h. The experimental samples were treated in an Ar-filled glove box to avoid contaminating.

The X-ray diffraction patterns after milling were obtained by an X-ray diffractometer (XRD) (D/max-2400) with Cu K_α1 radiation. The scan rate is 10 (°)/min. The microstructure and crystalline states of alloys dealt by milling were characterized by a high resolution transmission electron microscope (HRTEM) (JEM-2100F, operated at 200 kV) and electron diffraction (ED).

The isothermal de-/hydriding kinetics was performed on an automatically Sieverts apparatus. The hydrogenation performance was measured at 593, 613, 633 and 653 K under 3 MPa H₂, while the dehydrogenation performance was conducted at the same temperature under 1×10^-4 MPa. In each experiment, the stainless steel reactor needs to be sealed with 300 mg sample powder in it. Thermogravimetry (TG) and differential scanning calorimetry (DSC) (SDT-Q600) were adopted to investigate the non-isothermal dehydrogenation performance. The heating runs were performed at different rates (5, 10, 15 and 20 K/min). The test sample mass was typically 20 mg.

3 Results and discussion

3.1 Microstructure characteristics

The XRD diffraction patterns of as-milled REMg₁₁Ni-5MoS₂ (RE=Y, Sm) alloys before and after de-/hydrogenation are presented in Fig. 1. The hydrogenation reaction was performed at 593 K and 3 MPa, and dehydrogenation reaction at 593 K and 1×10^-4 MPa. It is found that adding catalyst MoS₂ does not create any new phase in the patterns, implying that there is no reaction between MoS₂ and any element in the alloy. The ICDD identification of XRD patterns reveals that the as-cast RE=Y alloy (not present here) includes major phase Mg₂₄Y₅ and secondary phase Mg₂Ni, while RE=Sm alloy contains major phase Sm₅Mg₄₁ and secondary phases SmMg₃ and Mg₂Ni. Mechanical milling makes diffraction peaks merge and broadened dramatically, showing a typical nanocrystalline and amorphous structure. After being hydrogenated, the alloy exhibits an obvious crystalline reaction, and three hydrides can be found in RE=Y alloy, viz. MgH₂, Mg₂NiH₄ and YH₃. Based on XRD analysis, possible formation pathways of the hydrides can be inferred as follows [37]:

                   (1)

                            (2)

                      (3)

But for RE=Sm alloy, the hydride phases are MgH₂, Mg₂NiH₄ and Sm₃H₇. The possible formation pathways of these hydrides can be inferred as follows [38]:

               (4)

                 (5)

                      (6)

It can be seen from Fig. 1(a) that three phases appear in dehydrogenated RE=Y alloy, including YH₂, Mg and Mg₂Ni. The path of the dehydrogenation reactions can be inferred as follows:

                           (7)

                      (8)

                            (9)

It is evident that the YH₂ phase keeps undecomposed, which is assigned to its high thermal stability. Meanwhile, three phases appear in dehydrogenated RE=Sm alloy, including Sm₃H₇, Mg and Mg₂Ni, as seen from Fig. 1(b). The path of these dehydrogenation reactions can be drawn as follows:

                           (10)

                     (11)

Apparently, the Sm₂H₇ phase keeps unchanged, which is assigned to its high thermal stability. It can be found that the reversible hydriding and dehydriding cycles of the activated RE=Y alloy include the following reactions:

                           (12)

                     (13)

                           (14)

And for RE=Sm alloy, the reversible cycles include the following reactions:

                           (15)

                     (16)

Fig. 1 XRD patterns of as-milled REMg₁₁Ni-5MoS₂ (RE=Y, Sm) alloys before and after hydrogen absorption and desorption

Figure 2 shows the microstructure and crystalline states of as-milled REMg₁₁Ni-5MoS₂ (RE=Y, Sm) alloys before and after hydrogen absorption and desorption. According to the HRTEM and ED analyses, a nanocrystalline and amorphous structure is observed. Furthermore, measured by linear intercept method, the average grain size is gauged to be 20 nm. The MoS₂ additive does not change the phase composition of REMg₁₁Ni (RE=Y, Sm) alloys, containing Mg₂₄Y₅ and Mg₂Ni phases in RE=Y alloy and Sm₅Mg₄₁, Mg₂Ni and SmMg₃ phases in RE=Sm alloy, which is also evidenced by the analysis of ED patterns. After being hydrogenated, the amount of the amorphous phase of as-milled REMg₁₁Ni (RE=Y, Sm) alloys obviously decreases, although the nanocrystalline and amorphous structure (Figs. 2(b) and (d)) still exists. This indicates that hydrogen absorption facilitates crystalline reaction. Three hydrides, MgH₂, Mg₂NiH₄ and YH₃, appear in hydrogenated RE=Y alloy, and MgH₂, Mg₂NiH₄ and Sm₃H₇ appear in hydrogenated RE=Sm alloy, which are supported by ED patterns. It is found from Figs. 2(c) and 2(f) that dehydrogenated REMg₁₁Ni (RE=Y, Sm) alloys both have a completely crystalline structure, and the grain sizes grow visibly. A similar result was reported by PUKAZHSELVAN et al [15]. The structural analysis and the index of ED rings of alloys reveal that there are three phases in dehydrogenated RE=Y alloy, viz. Mg, Mg₂Ni and YH₂ and three phases Mg, Mg₂Ni and Sm₃H₇ in dehydrogenated RE=Sm alloy. Apparently, YH₂ and Sm₃H₇ maintain undecomposed in the process of hydrogen desorption, which is consistent with XRD detections.

Fig. 2 HRTEM micrographs and SAD patterns of as-milled alloys in different states

3.2 Hydrogen absorption and desorption kinetics

The non-isothermal dehydrogenation performance of as-milled REMg₁₁Ni-5MoS₂ (RE=Y, Sm) alloys was investigated by thermogravimetry (TG). The temperature programmed desorption curve of the alloys by saturated hydriding at 593 K and 3 MPa, are illustrated in Fig. 3 at a heating rate of 5 K/min. In order to avoid the effect of desorption temperature due to incremental pressure, the mass of sample is set equal in each closed chamber. It is found that the as-milled REMg₁₁Ni-5MoS₂ (RE=Y, Sm) alloys start to release hydrogen at about 525.8 and 545.7 K, respectively. This suggests that the change of rare earth elements affects the initial desorption temperature of alloy, which might be ascribed to the weakened Mg—H bond caused by the electronic exchange reaction between rare earth elements and MgH₂. Lower initial desorption temperature of RE=Y alloy means that its Mg—H bond is weakened more seriously.

Fig. 3 Temperature programmed desorption curves of as-milled REMg₁₁Ni-5MoS₂ (RE=Y, Sm) alloys after hydrogen absorption at heating rate of 5 K/min

To inspect the different influences of elements Y and Sm on hydriding kinetics, the variations of hydriding capacity of as-milled REMg₁₁Ni-5MoS₂ (RE=Y, Sm) alloys with hydriding reaction time were operated at 593, 613, 633 and 653 K and 3 MPa. A rapid hydriding rate exhibits in the first few minutes, after which the value of hydrogen absorption capacity increases tardily at a long hydriding time, and the maximal hydrogenation capacity is nearly saturated, as shown in Fig. 4. These hydrogenation curve characteristics are most likely determined by the fact that within the first few minutes a surface-near hydride layer is formed, so that the diffusion of hydrogen through this layer becomes a rate-limiting factor. For facilitating comparison, 60 s is taken as a reference of hydrogen absorption time. It can be derived from Fig. 4 that the hydrogen absorption capacities within 60 s at 593, 613, 633 and 653 K are 4.363%, 4.577%, 4.708% and 4.827% (mass fraction), respectively, for RE=Y alloy, and 4.233%, 4.481%, 4.520% and 4.674% (mass fraction) for RE=Sm alloy. Apparently, the hydrogen absorption capacity for the fixed time is in order of RE=Y>RE=Sm, meaning that the element Y can facilitate the hydriding kinetics more than element Sm. It has been well known that hydrogenation process of MgH₂ is dominated by three critical steps, namely, 1) H₂ molecules decomposing into H atoms on the surface of alloy, 2) H atoms diffusing along grain boundaries, and 3) hydrogenation of Mg atoms and transforming into MgH₂ molecules. It is convinced that the decomposing energy of H₂ molecules is fairly high, which considerably dominates the hydriding rate of Mg [39]. POURABDOLI et al [40] considered that the addition of rare earth elements or their oxides may reduce dissociation energy and thus can be used as catalysts for hydrogen storage in Mg. Thereby, it is well-founded to believe that the addition of Y or Sm particles facilitates step 1), i.e., the dissociation of hydrogen molecules into atoms, and the addition of Y has a better improvement.

Fig. 4 Hydrogen absorption kinetic curves of as-milled REMg₁₁Ni-5MoS₂ (RE=Y, Sm) alloys at different temperatures

To investigate the effect of different rare earth elements on hydrogen desorption kinetics, isothermal dehydrogenation measurements of as-milled REMg₁₁Ni- 5MoS₂ (RE=Y, Sm) alloys were conducted at different temperatures, as illustrated in Fig. 5. The isothermal dehydrogenation curves of as-milled alloys were studied at 593, 613, 633, and 653 K. Distinctly, the temperature change causes a significant influence on dehydro- genation kinetics. The rise of temperature markedly enhances the dehydriding rate of alloys. It can be derived from Fig. 5 that the time needed by desorbing 3% (mass fraction) H₂ at 593, 613, 633 and 653 K is 507, 208, 125 and 86 s, respectively, for RE=Y alloy, and 938, 586, 296 and 140 s for RE=Sm alloy. This indicates that the variation of rare earth elements has a significant effect on improving the dehydrogenation kinetics of REMg₁₁Ni- 5MoS₂ (RE=Y, Sm) alloys. Obviously, the hydrogen desorption capacity for the fixed time is in order of RE=Y>RE=Sm, indicating that the hydrogen desorption kinetics can be improved more seriously by the addition of rare earth element Y.

Fig. 5 Hydrogen desorption kinetic curves of as-milled REMg₁₁Ni-5MoS₂ (RE=Y, Sm) alloys at different temperatures

3.3 Dehydrogenation activation energy

To understand the mechanism of RE=Y alloy possessing a faster dehydriding rate as compared with RE=Sm alloy, the dehydrogenation activation energy is estimated by Arrhenius and Kissinger methods for REMg₁₁Ni-5MoS₂ (RE=Y, Sm) alloys due to the fact that the energy barrier for MgH₂ releasing H₂ is deemed to be a dominated factor for dehydrogenation kinetics. Generally, the total energy barrier, which must be overcome in a gas-solid reaction, is signified by dehydrogenation activation energy. It is well known that the hydrogen desorption reaction is completed through a nucleation and growth process. Usually, the nucleation and growth process taking place during hydrogen desorption can be simulated by Johnsone-Mehle-Avramie (JMA) model, which can be expressed by the following equation [40]:

                   (17)

where α is the reaction rate of MgH₂ transforming into Mg at time t, k is an effective kinetic parameter and η is the reaction order or Avrami exponent. The JMA graphs of ln[-ln(1-α)] vs ln t at 593, 613, 633 and 653 K can be constructed by using logarithmic transform of Eq. (17), as presented in Fig. 6. The JMA plots are found to be nearly linear, implying that the dehydriding reaction of alloys starts from an instantaneous nucleation followed by interface controlled three-dimensional growth  process [41]. Thus, the η and ηln k values at different temperatures can be derived from the slope and intercept of JMA plots and the rate constant k can be easily calculated. The dehydrogenation activation energy () can be evaluated from Arrhenius equation [42]:

                           (18)

where A and R are the pre-exponential factor and gas constant, respectively.

Fig. 6 JMA graphs and Arrhenius plots of as-milled REMg₁₁Ni-5MoS₂ (RE=Y, Sm) alloys

The Arrhenius plots of ln k vs 1/T for dehydrogenation process are given in Fig. 6. Therefore, the activation energy can be derived from the  slopes. The  values are 85.32 and 87.89 kJ/mol for RE=Y and RE=Sm alloys, viz. a reduction by 2.57 kJ/mol in activation energy.

Here, Kissinger method is also used to evaluate the dehydrogenation activation energy  for comparison. Kissinger equation is as follows [43]:

                         (19)

where T_P is thermodynamic temperature, R is mole gas constant and β is heating rate.

The DSC curves of hydrogenated REMg₁₁Ni- 5MoS₂ (RE=Y, Sm) alloys under different heating rates were measured to determine the dehydrogenation activation energy . As displayed in Fig. 7, a dehydrogenated endothermic peak is obviously seen, and the peak shapes are similar for all alloys, indicating that the reaction processes are same. Additionally, for RE=Y alloy, the endothermic peak shows a drift tendency to low temperature, suggesting that element Y can also improve the dehydrogenation reaction rate. Based on the data in Fig. 7, the graphs of  vs 1/T_P can be built by using logarithmic transform of Eq. (19), which is termed as Kissinger plots, as inserted in Fig. 7. The Kissinger plots are found to be nearly linear, and thus the activation energy  can be easily calculated from the slopes of the fitting line. The  values of REMg₁₁Ni-5MoS₂ (RE=Y, Sm) alloys are 77.58 and 82.62 kJ/mol, respectively. Obviously, the activation energy evaluated from Arrhenius equation is larger compared with that from Kissinger equation. BARICCO et al [44] reported a similar result. Furthermore, it has a same result that the activation energy of RE=Y alloy is always lower than that of RE=Sm alloy, whichever method is used. Thereby, a conclusion can be drawn that the decrease of dehydrogenation activation energy is the real driving force for ameliorating dehydrogenation kinetics. FAN et al [45] considered that the activation energy (E_a) for hydrogen desorption is an important indicator to evaluate the dehydrogenation performance. The reduction of activation energy E_a means that the energy barriers of hydrogen releasing from the system are reduced.

Fig. 7 DSC curves and Kissinger plots of as-milled REMg₁₁Ni-5MoS₂ (RE=Y, Sm) alloys at various heating rates

3.4 P-C-T curves and hydrogen storage thermodynamics

To inspect the different effects of elements Sm and Y on de-/hydrogenation thermodynamics, P-C-T curves are measured at 593, 613, 633 and 653 K for the as-milled REMg₁₁Ni-5MoS₂ (RE=Y, Sm) alloys. It is presented in Fig. 8 that the absorption and desorption pressure plateaus are fairly flat and the hysteresis (H_f=ln(P_a/P_d)) is quite small. The variations of rare earth elements have an insignificant effect on the plateau features of P-C-T curves. It is clear that each P-C-T curve has two pressure plateaus. Among them, the lower pressure corresponds to the formation and dissociation of MgH₂, while the higher one corresponds to Mg₂NiH₄. Based on the plateau pressures (P_a and P_d) in Fig. 8, thermodynamics parameters, viz. enthalpy ΔH and entropy ΔS, can be derived from Van’t Hoff equation [46]:

                         (20)

where P₀ is 1.01325×10⁵ Pa, R is the mole gas constant and is the equilibrium plateau pressure corresponding to MgH₂. The van’t Hoff graphs of  vs 1/T for as-milled REMg₁₁Ni-5MoS₂ (RE=Y, Sm) alloys can be plotted by using logarithmic transform of Eq. (20), as inserted in Fig. 8. According to the slopes and intercepts of the fitting lines, the enthalpy ΔH and entropy ΔS can be easily calculated, as summarized in Table 1. This indicates that the substitution of Y for Sm renders a slight improvement on hydrogen storage thermodynamics of the alloy as the absolute values of both ΔH and ΔS of RE=Y are a little smaller than those of RE=Sm alloy for both hydrogen absorption and desorption processes.

Fig. 8 P-C-T curves and van’t Hoff plots of as-milled REMg₁₁Ni-5MoS₂ (RE=Y, Sm) alloys in temperature range of 593-653 K

Table 1 Enthalpy ΔH and entropy ΔS of as-milled REMg₁₁Ni-5MoS₂ (RE=Y, Sm) alloys

4 Conclusions

1) The substitution of Y for Sm significantly improves the de-/hydriding rates of REMg₁₁Ni-5MoS₂ (RE=Y, Sm) alloys, for which the reduction of dehydrogenation activation energy produced by substituting Sm with Y is responsible.

2) The dehydrogenation activation energy is evaluated from Arrhenius and Kissinger equations. The results reveal that the dehydrogenation activation energy is visibly reduced by substituting Sm with Y, which is considered to be the real driving force of the dehydrogenation kinetics improved by substituting Sm with Y.

3) The substitution of Y for Sm incurs a slight reduction in enthalpy ΔH and entropy ΔS for the alloys. Furthermore, it facilitates to reduce the stability of alloy hydride. The initial hydrogen desorption temperature of hydrogenated REMg₁₁Ni-5MoS₂ (RE=Y, Sm) alloys is reduced from 545.7 to 525.8 K by substituting Sm with Y.

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MoS₂催化的球磨态REMg₁₁Ni (RE=Y, Sm)合金的储氢性能

张羊换^1,2，张 巍^1,2，袁泽明^1,2，卜文刚²，祁 焱²，董小平³，郭世海²

1. 内蒙古科技大学 内蒙古自治区白云鄂博矿多金属资源综合利用重点实验室，包头 014010；

2. 钢铁研究总院 功能材料研究所，北京 100081；

3. 河北大学 机械系，保定 071002

摘  要：制备MoS₂催化作用下的球磨态REMg₁₁Ni-5MoS₂ (质量分数)(RE=Y, Sm)合金，用以比较其储氢性能。储氢性能通过多种方法测定，包括XRD、TEM、自动Sievert设备、TG和DSC。结果显示，两种球磨态的合金都具有纳米晶和非晶结构。与RE=Sm合金相比，RE=Y合金具有较大的吸氢量、较快的吸氢速率、较低的初始放氢温度、较好的放氢性能和较低的放氢活化能，其中较低的放氢活化能被视为其具备较好储氢动力学的原因。

关键词：镁基合金；球磨；催化剂；稀土元素；储氢性能

(Edited by Bing YANG)

Foundation item: Projects (51761032, 51471054, 51871125) supported by the National Natural Science Foundation of China

Corresponding author: Yang-huan ZHANG; Tel: +86-10-62183115; Fax: +86-10-62187102; E-mail: zhangyh59@sina.com

DOI: 10.1016/S1003-6326(18)64827-7