Improved Absorption and Desorption Kinetics of Mg–Ni–Ce Alloy Activated under Elevated Hydrogen Pressure

Abstract

The large-scale application of Mg as a hydrogen storage material has been limited by its slow absorption and desorption kinetics at moderate temperatures. Refining the microstructures is an effective way to improve the hydrogen storage performance. Aiming at improving the de-/absorption kinetics of Mg-based alloys by in situ formed superfine catalysts, Mg–7Ni4Ce (mass%) alloy has been prepared and activated by controlling the activation hydrogen pressure. The phase components, microstructure and hydrogen storage properties have been systematically investigated. It is found that an 18R-type long-period stacking ordered (LPSO) phase is formed in the eutectic region of as-cast Mg–Ce–Ni ternary alloy. The observed LPSO structure is a variant of Mg₁₂Ce rather than Mg. Abundant secondary phase particles are obtained in the as-activated alloy. For the sample which is activated at 300°C under 7.5 MPa hydrogen pressure, the particle size of secondary phase is much finer than that of the sample activated under 3 MPa hydrogen pressure. The sample activated under higher hydrogen pressure shows superior absorption and desorption kinetics.

1. Introduction

With the increasing concern on environmental deterioration and energy shortage, hydrogen has been universally regarded as an optimal energy carrier in the future due to its considerable merits of high energy efficiency, zero emission of greenhouse gases and inexhaustible supply from water.¹⁾ Hydrogen storage technology in solids is considered one of the most promising choice for applications in fuel cell and hybrid vehicles.²⁾ Unfortunately, the widespread use of hydrogen is still impeded by its storage technologies, especially the storage in solid hydrides.

Among all kinds of hydrogen storage materials, Mg and Mg based alloys are considered to be the most promising candidates owing to the high gravimetric hydrogen storage capacity (7.6 mass% in MgH₂), environmental friendliness and abundance in the earth.^3–6) The major drawbacks with pure Mg as a reversible hydrogen storage material are its high thermodynamic stability and poor kinetics during process of hydrogenation and dehydrogenation.^7,8) Encouragingly, Mg based composites in nanoparticle, nanowire and thin film forms, which are synthesized by alloying or adding catalysts, are currently subject to intense scientific inquiry and extraordinary achievements have been made.^9,10)

Combining Mg with a small amount of other materials acting as catalysts or reaction path modifiers has shown significant improvements in absorption and desorption kinetics and ameliorates the thermodynamics to some extent.^11,12) The additives are reported to accelerate the dissociation and recombination of hydrogen molecules, shorten the diffusion distance of hydrogen atoms, inhibit the grain growth and particle aggregation of Mg matrix.^13,14) Moreover, several catalysts are reported to be able to weaken the Mg–H bond, thus decreasing the thermodynamic stability and lowering the dehydrogenation temperature.^15,16)

Recently it is shown that in situ formed catalysts show higher catalytic activity and superior stability than those of the externally added catalysts due to the better homogeneity and finer particle size.^13,17,18) Of these additives with positive effect, rare earth (RE) metals have been proved to be able to significantly lower the dissociation barrier of hydrogen molecule and achieve homogeneous distribution.^19–21) Due to a much stronger affinity between RE and hydrogen, RE hydrides usually keep stable in the temperature range of MgH₂ dehydrogenation and act as catalysts accelerating absorption and desorption kinetics of the matrix.²²⁾ RE hydrides can chemisorb hydrogen atoms and transfer them to the matrix-RE hydrides interfaces, which serve as active nucleation sites for magnesium hydride.²³⁾ Correspondingly, the kinetics properties are dramatically improved. Mg₃RE alloys can absorb hydrogen to form MgH₂ and RE hydride even at room temperature via a synergetic effect of in situ formed hydrides on improving the kinetic properties.²⁴⁾ Considering the requirement of hydrogen storage capacity, however, the RE and/or other additive contents should be kept in a reasonable range.

Understanding from the de-/hydrogenation process, the absorption and desorption kinetics of Mg depend on both surface and internal features. For hydrogenation, the surface characteristics determine the adsorption and dissociation of hydrogen molecules. The subsequent hydrogen diffusion and occupation process is strongly associated with the internal microstructure.²⁵⁾ Ultrafine nanostructures of Mg-based materials with substantial surface and short hydrogen diffusion distance are always favorable to absorption and desorption kinetics. For pure Mg particles with particle size in the range of 3∼8 nm, the hydrogen absorption capacity can even reach 7.6 mass% at near room temperature.²⁶⁾ Improved absorption and desorption kinetics have also been observed through increasing the solidification rate, thus refining the nanocrystalline microstructures.²⁷⁾ It is found that Mg nanowires with smallest diameter (30–50 nm) can absorb 7.6 mass% and desorb 6.8 mass% hydrogen within 30 min at 300°C, while the largest diameter nanowires (150–170 nm) absorb 6.5 mass% and desorb 5 mass% hydrogen at the same condition.²⁸⁾ High-energy ball milling (HEBM) has been generally regarded as an effective method to prepare Mg-based materials with nanocrystalline. HEBM can refine both grain and particle size, introduce dislocations and distribute catalysts simultaneously.²⁹⁾

Inspired by the mentioned above and considering the requirement of hydrogen storage capacity, Mg-rich Mg–Ni–Ce alloy is fabricated in the present work. Moreover, a simple but valid strategy via controlling the hydrogen pressure during activation process is proposed. Alloying Mg with Ce and Ni elements is based on the synergistic catalytic effect between transition metals and RE elements. The phase component, microstructure and hydrogen storage properties of prepared samples are investigated systematically. It is found that increasing hydrogen pressure is beneficial for obtaining fine hydrides, thus leading to improved hydrogen absorption and desorption kinetics.

2. Experimental

2.1 Materials preparation

The nominal composition of the alloy prepared in the present work is Mg–6Ni–6Ce (mass%). The detailed preparation process has been described elsewhere.^30,31) In short, it can be summarized as melting Mg ingot, Ce ingot and Mg–Ni intermediate alloy in graphite crucible under the protection of covering agent (RJ-6) in the resistance furnace. To obtain near equilibrium solidification, furnace cooling was performed. Prior to activation and de-/hydrogenation tests, the as-cast alloy was mechanically cut and pulverized to ∼300 µm, followed by HEBM (Spex 8000D) for 2 h in argon. The ball to powder weight ratio was 20:1 and the speed of the mill was 875 rpm.

2.2 Characterization and measurement

After solidification, the specific composition of the as-cast alloy was determined by inductive coupled plasma atomic emission spectrometer (ICP-AES), which was measured to be Mg–7.34Ni–4.17Ce (mass%). Hereafter, the alloy was conveniently denoted as Mg–7Ni4Ce. The microstructure of the as-cast alloy and activated sample was characterized by scanning electron microscopy (SEM, equipped with EDS) in a backscattered mode and transmission electron microscopy (TEM) (FEI Tecnai G2 F30). Phase identification was also carried out via X-ray diffraction (XRD, DX-2700) with Cu Kα radiation. To compare the effects of hydrogen pressure during activation process on the microstructure and hydrogen storage properties of Mg-based alloys, the ball milled particles were activated at 300°C under an initial hydrogen pressure of 3 MPa and 7.5 MPa, respectively, using a Sieverts-type (PCT Pro2000) apparatus. The isothermal absorption and desorption kinetics of as-activated samples were measured by PCT. The desorption properties were also measured by differential scanning calorimetry (DSC) on a simultaneous TG-DTA/DSC apparatus (STA449C) at the heating rate of 5, 10 and 15°C/min, respectively, up to a maximum temperature of 450°C under 50 ml/min flowing argon gas.

3. Results and Discussions

3.1 Microstructure of as-cast alloy

The XRD pattern of as-cast Mg–7Ni4Ce alloy is illustrated in Fig. 1. Bragg peaks from Mg, Mg₂Ni and Mg₁₂Ce are clearly observed. Unexpectedly, Bragg peaks around 36.1° and 36.8° are also obviously observed in the XRD pattern. According to the Mg–Ni–Ce ternary phase diagram,³²⁾ the phases in the Mg-rich corner are only Mg, Mg₂Ni, Mg₁₂Ce or Mg₁₇Ce₂. Moreover, the observed Bragg peaks around 36.1° and 36.8° can not correspond to any phase in Mg–Ni, Mg–Ce, Ni–Ce or Mg–Ni–Ce phase diagrams available.^32–35) However, the first-principle calculation results indicate that the extraordinary observed Bragg peaks can be ascribed to the existence of a long period stacking ordered (LPSO) phase in Mg–TM–RE system.²⁰⁾ It is temporarily marked as LPSO phase for the observed Bragg peaks at 36.1° and 36.8° in Fig. 1 and further investigation by TEM technique will be given below.

Fig. 1

XRD pattern of as-cast Mg–7Ni4Ce alloy.

Figure 2(a) presents the back-scattered SEM (SEM/BSE) image of the as-cast Mg–7Ni4Ce alloy, showing typical Mg primary dendrites (marked with A) together with extensive eutectic structures and a few gray blocks (marked with B). The compositions of consistent phases marked with letter A–D measured by EDS are given in Table 1. Ten different points are selected for each consistent phase during EDS measurement and the data in Table 1 are the average. The gray block marked with B is identified as Mg₁₂Ce phase with minor Ni solution by EDS. The microstructure details of as-cast Mg–7Ni4Ce alloy in eutectic region are displayed in Fig. 2(b), showing dark microstructures (Mg) and gray microstructures with some bright strips embedded in. The bright strips marked with C are measured to be Mg₂Ni phase and further confirmation by TEM technique are given below. Unexpectedly, the composition of gray microstructures marked with D in Fig. 2(b) can not be accurately determined. In other words, the composition of gray microstructures fluctuates greatly with the change of selected points during EDS measurement.

Fig. 2

(a) SEM/BSE image of as-cast Mg–7Ni4Ce alloy. (b) is the magnified image in (a), showing the microstructure details in the eutectic region.

Table 1 EDS results of the as-cast alloy indicating the phase composition.

Figure 3 gives the bright-field (BF) image of the as-cast Mg–7Ni4Ce alloy in the eutectic region and corresponding selected area electron diffraction (SAED) patterns. The phase distribution characteristics in BF image coincide well with the microstructure observed in Fig. 2(b). The diffraction spots in SAED1 confirm Mg₂Ni phase with a hexagonal crystal structure, which is consistent with the EDS results in Table 1. The strip-like Mg₂Ni phase has also been observed in Mg–Ni–Y ternary hydrogen storage alloy.²⁰⁾ The Mg₁₂Ce phase is identified by SAED2 pattern, which corresponds to the gray microstructures shown in Fig. 2(b). However, a new phase with fairly new diffracting pattern is observed in SAED3 which has not been reported before in Mg–Ni–Ce system.^32,33) As shown in SAED3 pattern, five extra and uniformly distributed diffraction spots are obviously observed between two adjacent diffracting spots of Mg₁₂Ce. The features are commonly adopted evidences proving the existence of the 18R type LPSO phase in Mg–TM (transition metals)–RE systems.^36–38) LPSO structures reported in Mg-based alloys are usually stacking variants of hcp-Mg crystal³⁹⁾ and the existence of LPSO phase as a variant of hcp-Mg in Mg–Ni–Ce alloy has been excluded by DFT.⁴⁰⁾ However, distinct from the reported results, the observed 18R-type LPSO structure in the present work is a variant of Mg₁₂Ce with tetragonal structure (I4/mmm, a = 1.033 nm, c = 0.596 nm, JCPDS 19-0289). The orientation relation between the 18R phase and the Mg₁₂Ce matrix is (002)_Mg12Ce//(0018)_18R. The observed LPSO phase in the eutectic region accounts for the Bragg peaks at 36.1° and 36.8° in the XRD pattern in Fig. 1 and the composition fluctuation during EDS measurement. Therefore the eutectic structure of Mg–7Ni4Ce alloy is composed of Mg, Mg₂Ni, Mg₁₂Ce and 18R-type LPSO phase. The specific composition and structure of the LPSO phase need to be further investigated.

Fig. 3

TEM image in the eutectic region of as-cast Mg–7Ni4Ce alloy and corresponding SAED patterns confirming the corresponding consistent phases.

It has been reported that the LPSO phase is unstable during high energy ball milling and subsequent hydrogenation processes.^20,30,41) It has been found in Mg–Ni–Y system that the 18R-type LPSO phase decomposes into ultrafine YH₂, MgH₂ and Mg₂Ni phases during the first hydrogenation process by in-situ synchrotron X-ray diffraction and TEM.²⁰⁾ Although LPSO phase is unstable during hydrogen absorption process and the decomposition is irreversible, it can make Mg, Ni and Ce elements distribute uniformly at the atomic scale during solidification process, and can decompose into ultrafine catalytic phases after hydrogenation, promoting hydrogen absorption and desorption. As a precursor, it is easy to achieve superior de-/hydrogenation kinetics by reasonable alloying and processing.

3.2 Microstructure of as-activated samples

During the process of first hydrogenation, Mg₁₂Ce undertakes disproportion reaction and decomposes into MgH₂ and CeH_2.73. CeH_2.73 remains unchanged in subsequent hydrogenation and dehydrogenation processes, while Mg₂Ni can release and absorb hydrogen again.³⁰⁾ Figure 4(a) shows the SEM/BSE image of as-activated particle which is activated at 300°C under an initial hydrogen pressure of 3 MPa. Abundant secondary phase particles with particle size in the range of 1∼3 µm are clearly observed on the surface of MgH₂ matrix. With respect to the particles activated under an initial hydrogen pressure of 7.5 MPa as displayed in Fig. 4(b), the secondary phase particles are much finer. According to the scale bar, the particle size of secondary phase is estimated to be ∼0.1 µm, indicating a significant effect of activation pressure on the microstructure characteristics.

Fig. 4

SEM/BSE images of the hydrogenated samples activated under hydrogen pressure of (a) 3 MPa and (b) 7.5 MPa.

TEM technique is applied to further reveal the microstructure details of samples activated under different hydrogen pressures. The BF image of dehydrogenated alloy activated under 3 MPa hydrogen pressure is shown in Fig. 5(a). Seen from the BF image, secondary phases with irregular shape are clearly observed. According to the scale bar, the particle size of secondary phase is in the range of 100–200 nm. Figure 5(b) gives the dark-field (DF) image of dehydrogenated alloy which is activated under 7.5 MPa hydrogen pressure. Prevalent Mg₂Ni particles with particle size in the range of several nm to ∼20 nm are clearly observed, displaying much finer secondary particles in the Mg matrix.

Fig. 5

(a) BF image of as-activated alloy under 3 MPa hydrogen pressure. (b) is the DF image of as-activated alloy under 7.5 MPa hydrogen pressure and corresponding SAED pattern.

3.3 Hydrogen storage performance

The isothermal hydrogenation and dehydrogenation kinetic curves at 300°C of as-activated alloys are illustrated in Fig. 6. For absorption at 300°C as displayed in Fig. 6(a), the alloy which is activated under 3 MPa hydrogen pressure absorbs 5.0 mass% hydrogen within 15 min, while it is 3 min for the alloy activated under 7.5 MPa hydrogen. Although the theoretical hydrogen storage capacities of the two samples are the same, the sample activated at 7.5 MPa hydrogen can absorb 6.3 mass% hydrogen, which is ∼0.5 mass% higher than that of the sample activated at 3 MPa hydrogen. With respect to dehydrogenation, similar phenomenon is observed in Fig. 6(b). The desorption rate of the sample activated under 7.5 MPa hydrogen is much faster than that of the sample activated under 3 MPa hydrogen at 300°C. Another point worth mentioning is that the hydrogen absorption capacity of the sample activated under 7.5 MPa hydrogen pressure is still lower than the theoretical capacity (∼6.8 mass%). During hydrogenation of Mg–Ni–RE alloys, the transition from Mg₂NiH_0.3 to Mg₂NiH₄ is the last and the slowest hydrogen absorption process.⁴²⁾ Only 6% of Mg₂NiH_0.3 is transformed into Mg₂NiH₄ after 450 min of hydrogenation at 350°C under an initial hydrogen pressure of 2.2 MPa. Correspondingly, the absorbed hydrogen is usually lower than the theoretical value for Mg–Ni–RE alloys.

Fig. 6

Hydrogen (a) absorption and (b) desorption kinetics of samples activated under different hydrogen pressures at 300°C.

To evaluate the effects of hydrogen pressure during activation process on hydrogen storage properties, the desorption performance of as-activated samples under different pressures has been measured by DSC. As illustrated in Fig. 7(a), both the onset temperature and peak temperature of the sample activated under 3 MPa hydrogen pressure are much higher than those of the sample activated under 7.5 MPa hydrogen pressure at each heating rate, suggesting superior desorption kinetics of sample activated under a higher hydrogen pressure. The desorption activation energy has been also calculated according to the DSC curves using Kissinger equation:⁴³⁾   

\begin{equation} \frac{d(\ln \beta /T_{p}^{2})}{d(1/{T_{p}})} = -\frac{E_{A}}{R} \end{equation} (1)

where T_p is the peak temperature, β is the heating rate, E_A is the activation energy and R is the gas constant. The plot of ln β/T_p² versus 1000/T_p is shown in Fig. 7(b). The desorption activation energies of the two as-activated samples are obtained from the slope of the fitted line. The desorption activation energy of as-activated sample under 7.5 MPa hydrogen pressure is ∼99 kJ/mol, much lower than that of the as-activated sample under 3 MPa hydrogen pressure (∼126 kJ/mol). It is the finer microstructure induced by higher hydrogen pressure during activation that accounts for the improved desorption performance. In other words, increasing hydrogen pressure is beneficial for obtaining fine microstructures, thus bringing improved hydrogen absorption and desorption kinetics.

Fig. 7

(a) DSC curves at various heating rates of as-activated samples under different hydrogen pressures and (b) corresponding plots of ln (β/T_p²) versus 1000/T_p of the samples in (a).

4. Conclusions

In summary, the Mg–7Ni4Ce alloy has been successfully prepared in the present work and an effective strategy to obtain superfine microstructures of Mg-based hydrogen storage alloy via activation under elevated hydrogen pressure is proposed. A fairly new LPSO phase with an 18R-type structure, which is a variant of Mg₁₂Ce, is observed in the eutectic region in as-cast Mg–7Ni4Ce ternary alloy together with Mg, Mg₂Ni and Mg₁₂Ce phases. The microstructure of the sample activated under 7.5 MPa hydrogen is much finer than that of the sample activated under 3 MPa hydrogen. The particle size of secondary phase can be decreased to even several nm. Correspondingly, improved absorption and desorption kinetics are obtained for the sample activated under higher hydrogen pressure. The desorption activation energy of the sample activated under 7.5 MPa hydrogen is calculated to be 99 kJ/mol, much lower than that of as-activated sample under 3 MPa hydrogen pressure.

Acknowledgments

This work is financially supported by the Natural Science Foundation of Jiangsu Province (Grant No. BK20191020) and the Scientific Research Foundation of Nanjing Institute of Technology (No. YKJ201804).

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