1 Introduction

Hydrogen is world widely regarded as a sustainable energy carrier for solving the fossil energy shortage and environmental pollution issues [1]. Hydrogen can be produced directly from water via electrochemical or photochemical water splitting [2], and it can be converted into water to produce heat or electricity for industrial or domestic use. Compared to the carbon-based energy society, hydrogen-based energy cycle can well solve the energy sustainability and pollution problems because the hydrogen energy technologies are in a small burden to our environment.

The utilizations of the hydrogen consist of four main parts, including hydrogen production, hydrogen storage and transportation, hydrogen conversion and hydrogen utilization, among which hydrogen storage is one of the most difficult and challenging issues for scientists or engineers. Until now, there is no solution that can perfectly solve the hydrogen storage issue [3]. In 2015, Toyota announced the success of the first commercial hydrogen fuel cell vehicle, MIRAI. The vehicle stores energy as compressed hydrogen gas with pressure as high as 70 MPa, which is combined in a fuel cell stack with oxygen taken from the air to generate electricity. Although with a great success selling in Europe and USA, it is regarded that storage of hydrogen in a high pressure compressed gas way is only a short-term solution, and to store hydrogen in solid-state materials is considered as a more promising and safer way [4,5,6,7].

Hydrogen can be stored in metals or compounds by chemical bonding inside the materials [8]. Typically, LaNi5 can readily absorb hydrogen to form LaNi5H6 and reversibly recover upon heating or pumping at ambient conditions [9]. The study on hydrogen storage materials is one of the hottest and most important research topics in the fields of materials science and energy technology. Here, some classical hydrogen storage alloys are summarized, and their main hydrogen storage properties are listed in Table 1 [10,11,12,13,14,15,16,17,18,19,20,21].

Table 1 Classical hydrogen storage alloys and metastable hydrogen storage materials
Full size table

By simply comparing the classic hydrogen storage alloys in Table 1, it can be easily seen that there is a trade-off relation between the hydrogen storage capacity and the hydrogen storage conditions. The materials with higher capacity will only absorb and desorb hydrogen at higher temperatures and pressures, and vice versa. For MgH2, catalysts doping, nanostructuring and combined strategies were adopted and compared to develop promising hydrogen storage systems; however, the practical application was still limited owing to its high thermal stability [1, 22,23,24]. More hydrogen due to that properties data can also be found in other published literatures [4, 25,26,27].

In order to deal with the trade-off issue, many efforts have been devoted to developing metastable alloys for hydrogen storage. Indeed, metastable alloys show very much different hydrogen storage kinetic and thermodynamic properties compared with their thermodynamically stable crystalline counterparts. The metastable alloys may absorb and desorb hydrogen at much lower temperatures with much faster kinetics [28], and in some cases, the hydrogen storage capacities may be higher [16]. Therefore, the new metastable structures may open new pathways for developing hydrogen storage alloys with greatly improved properties.

In this review, recent advances in metastable alloys for hydrogen storage applications have been comprehensively reviewed. Firstly, materials preparation methods to synthesize these metastable hydrogen storage alloys were summarized. Afterwards, hydrogen storage properties of the metastable alloys are summarized, focusing on the unique kinetics and thermodynamics properties resulting from the unique metastable structures. In addition to highlighting the recent achievements of metastable alloys in the field of hydrogen storage, the remaining challenges and trends of the emerging research are also discussed.

2 Techniques to prepare metastable hydrogen storage materials

Metastable materials are thermodynamically metastable, and the metastable structures are usually formed in non-equilibrium conditions. As proposed by Prof. Turnbull [29], these structures are configurationally frozen, such as glassy metals, highly supersaturated crystalline alloys and new alloys with exceptionally high interfacial densities. In order to obtain metastable structures, preparation methods under non-equilibrium conditions are usually applied. For example, high-energy ball milling, rapid quenching of the melt, high-pressure synthesis methods, severe plastic deformation and physical vapor deposition are some of the commonly used techniques to fabricate metastable alloys. Figure 1 shows the schematic illustrations of several typical preparation methods for metastable hydrogen storage materials.

Fig. 1
figure 1

Schematic diagrams of typical preparation methods for metastable hydrogen storage materials

Full size image

2.1 High-energy ball milling

High-energy ball milling is one of the most efficient and commonly used techniques to prepare metastable hydrogen storage alloys [30], such as nanocrystalline alloys, amorphous alloys and high-entropy alloys. Particularly, the powder materials can be easily prepared by high-energy ball milling with very well controlled chemical compositions and particle/crystallite sizes. A ball milling machine is easily accessible in a laboratory because its size is generally small, and not in a high price, and the operation procedure is quite simple.

In the past decades, a large number of metastable hydrogen storage alloys have been explored by high-energy ball milling, such as Mg–Ni nano-/amorphous alloys [31], Mg–Ag, Mg–Cd and Mg–In supersaturated solid solution alloys [32,33,34,35]. Zaluski et al. [36,37,38] synthesized additives/catalysts-doped hydrogen storage materials via high-energy ball milling [39,40,41,42,43], in which process the materials are usually homogenous and thus quite beneficial for the improved hydrogen storage kinetics.

2.2 Rapid quenching

Metastable alloys, including monolithic glasses, amorphous nanocomposites and quasi-crystals, can be directly produced from copper-mold casting or quenching of the melt. Under very rapid cooling conditions, metastable phases can be obtained from the liquid state [44]. Amorphous alloys, quasicrystalline alloys and saturated solid solution alloys have been abundantly prepared and studied as materials for hydrogen storage applications. Single Cu-roller melt-spinning have been abundantly adapted in laboratory to prepare amorphous and quasicrystalline alloy ribbons. Owing to the very high quenching rate of 1 × 105 K·s−1, metastable phases are quite easily formed in the melt-spun alloys. In the past decades, melt-spun Mg–Ni–RE (RE = rare earth metals) alloys [45,46,47,48,49,50,51,52,53,54,55,56], melt-spun Ti-based quasicrystals [57,58,59] have been widely studied for hydrogen storage applications. Particularly, improvements on hydrogen storage kinetics and cycling performances could be easily achieved. Hu et al. [59] studied the electrochemical hydrogen storage performances of (Ti1−xVx)2Ni (x = 0.05–0.3) alloys containing icosahedral quasicrystalline phase and found that the capacity first improved due to the activation process followed by decays due to the corrosion of V. The existence of Ti could restrain the capacity deterioration. The cycling capacity retention rates were about 80% after 30 cycles. Since the melt-spun alloys are usually in the form of ribbons, they should be crushed into powders either by hand milling or by ball milling prior to the hydrogen storage measurements.

2.3 High-pressure synthesis

Materials in metastable structures can be easily formed under high-pressure conditions. High pressure has long been used as an efficient tool to fabricate metastable materials. Many types of materials, including carbon-based materials, oxides and intermetallic compounds, can be transformed into metastable amorphous materials under high pressure [60]. High pressure up to tens of gigapascal can be achieved in a diamond anvil cell. It was reported by Vajeeston et al. [61] that the typical hydrogen storage material MgH2 would undergo several metastable phase transformations under pressure up to 20 GPa. Kyoi et al. [62,63,64,65,66] prepared a number of metastable Mg-rich hydrides Mg6–7MH12–16 (M = Ti, V, Nb, Ta, Zr and Hf) by using a gigapascal high-pressures thermal technique, which could not be successful under conventional preparation conditions.

On the other hand, high hydrogen pressure synthesis can be used to synthesize metastable materials for hydrogen storage. For example, Kamegawa et al. [67] prepared a number of Mg–RE–H hydrides (RE = Y, La, Ce, Pr, Sm, Gd, Tb, Dy) by high-pressure synthesis. Selvam et al. [68] synthesized a group of hydrides, including Mg2FeH6, Mg2CoH5 and Mg2NiH4, by high-pressure sintering of the elements, under temperature of 450–500 °C and hydrogen pressure up to 9 MPa.

2.4 Severe plastic deformation

Severe plastic deformation (SPD) has been developed to fabricate bulk nanostructured metals with advanced properties during the past decades [69, 70]. Particularly, equal-channel angular pressing (ECAP) [71] and high-pressure torsion (HPT) [72, 73] are two of the most efficient methods to achieve severe plastic deformation for the materials. Edalati et al. reported that immiscible Mg–Zr [72] and Mg–Ti [73] alloys system could be produced by HPT-treatment under high pressure, leading to greatly improved hydrogen storage kinetics at low temperatures. The HPT treated immiscible Mg–Zr alloys can even reversibly absorb and desorb 1 wt% of hydrogen due to the Mg nanoclusters generated by the HPT treatment. It was also shown by Révész et al. [74] that HPT can provoke a drastic decrease in the hydrogen absorption temperature in an amorphous Mg–Ni–Cu–Y alloy.

2.5 Physical vapor deposition

Physical vapor deposition is a process during which the ingredient elements evaporate into atoms, molecules or ions in vacuum, and then condense on the substrate surface to form a desired thin film [75]. The deposition is a non-equilibrium process which is suitable for preparing metastable films for hydrogen storage or hydrogen sensing [28, 76,77,78]. It has been well established that nanostructuring of the materials can lead to greatly improved hydrogen storage kinetics, and thin films in nanoscale can be easily and efficiently prepared by physical vapor deposition through magnetron sputtering or pulsed laser deposition technologies.

In the past decades, tremendous efforts have been devoted to study the hydrogen storage property of thin films. Besides the improvements in hydrogen storage kinetics, the tunable thermodynamics of these materials is even more attractive to researchers. For example, Baldi et al. [79] reported the destabilization of the Mg–H system through elastic constraints in multi-layer films with a sandwich structure. Moreover, because of the optical transformation of thin films upon hydrogenation/dehydrogenation, they have also been studied as hydrogen sensors in the past years [28, 80].

3 Metastable hydrogen storage alloys

3.1 Nanocrystalline hydrogen storage alloys

Mg is regarded as a promising hydrogen storage material because of its high hydrogen storage capacity and low cost. However, the dehydrogenation of Mg hydride (MgH2) requires a temperature of 300 °C or higher. This is mainly because MgH2 is thermodynamically too stable. The hydrogen desorption enthalpy (ΔH) of MgH2 is as high as 75 kJ·mol−1 H2. In recent years, metastable nanocrystalline Mg and Mg-based hydrogen storage alloys were investigated widely by scholars around the world to improve the performance of the materials. Zhang et al. [81, 82] adopted wet-chemical milling to synthesize nanocrystalline Mg and Mg–NiTiO3 and found hydrogenation was started at room temperature for both samples (Fig. 2). Due to the synergistic effect of nanosizing and catalyst doping, the Mg–NiTiO3 composite presented lower absorption apparent activation energy and desorption enthalpy compared to those of nanocrystalline Mg. Shao et al. [83] prepared Mg2Ni under 3 MPa hydrogen pressure at 280 °C from magnesium and nickel ultrafine particles by ball milling. The prepared Mg2Ni readily absorbs 1.74 wt%, 2.07 wt%, 2.31 wt% and 2.82 wt% hydrogen at 20, 75, 153 and 220 °C, respectively. Electron-beam deposition was adopted by Vermeulen et al. [84] to synthesize metastable MgyTi(1−y) thin films with y ranging from 0.50 to 0.95 at room temperature. Hydrogen storage properties characterization revealed that metastable Mg0.80Ti0.20 showed a superior reversible hydrogen storage capacity with an excellent rate-capability compared with that of pure single-phase Mg thin films. High-pressure torsion (HPT) was also considered as a processing route to synthesize bulk hydrogen storage materials with nano-grains. Edalati et al. [73] produced four metastable phases with grain sizes of 5–10 nm by HPT processing on the Mg–Ti immiscible alloys. The metastable alloys decompose to Ti and Mg at 370 °C and theoretical calculations showed that both the binding energy of hydrogen and the thermodynamic hydride stability increase with the addition of Ti to Mg. Li et al. [15] prepared a Mg–Co-based metastable nanoalloys with a body-centered cubic (BCC) lattice structure by mechanical alloying method. Subsequent investigation exhibited that Mg50Co50 and Mg55Co45 metastable BCC alloys can take up 2.67 wt%–3.24 wt% hydrogen at − 15 °C, which is almost the lowest temperature reported in the literature for Mg-based materials to absorb hydrogen. Compared with other structures, the BCC structure has three times higher interstitial sites (octahedral and tetrahedral ones) for possible hydrogen occupation, and thus it is more promising to be used for future hydrogen storage area.

Fig. 2
figure 2

Reproduced with permission from Ref. [81]. Copyright 2020, Elsevier. b Hydrogen desorption, c hydrogen absorption properties, and d modification mechanism for nanocrystalline Mg and Mg−NiTiO3. Reproduced with permission from Ref. [82]. Copyright 2021, Elsevier

a Synthesis process of nanocrystalline Mg.

Full size image

As a common metastable phase of MgH2, γ-MgH2 as a hydrogen storage material also attracts the researchers worldwide. Shen and Aguey-Zinsou [85] reported that a mixed γ/β hydride phase with high γ-MgH2 content (29.6%) was formed during the electrochemical synthesis of nanosized Mg upon hydrogenation. The existence of γ-MgH2 was found to bring a rapid reduction of the apparent activation energy (Ea) from (106.2 ± 4.0) to (69.1 ± 2.9) kJ·mol−1 as well as a reduction of the Mg/H2 reaction enthalpy (ΔH) from (74.8 ± 1.0) to (57.7 ± 5.3) kJ·mol−1 H2. Mechanism analysis testified that the improved hydrogen sorption properties are due to the formation of the metastable γ-MgH2 phase. In the meanwhile, Xiao et al. [86] also synthesized a β-/γ-MgH2 nanocomposite composed of tetragonal β-MgH2 and 18% orthorhombic γ-MgH2. The β-/γ-MgH2 nanocomposite starts to desorb H2 at ~ 260 °C, 140 °C lower than commercial MgH2, with a final hydrogen capacity of 6.6 wt%. The Ea for hydrogen desorption of the β-/γ-MgH2 nanocomposite was calculated to be (115 ± 3) kJ·mol−1, 46% lower than that of commercial MgH2. The significant improvement in desorption performance of the β-/γ-MgH2 nanocomposite over the β-MgH2 itself could be attributed to two main causes: a reactive, defective nanostructure which could result in enhanced hydrogen diffusion and an exothermic process associated with the γ- to β-MgH2 transformation, leading to the improved hydrogen desorption kinetics.

3.2 Amorphous hydrogen storage alloys

Amorphous alloys were occasionally developed by rapid quenching of the Au–Si alloy in 1960s [87], almost at the same time when hydrogen storage alloys were also accidentally invented. During the past decades, extensive efforts have been devoted to preparing Mg-based Mg–RE–Ni (RE = Rare earth, such as La and Ce) alloys by rapid solidification, resulting in amorphous phases. Amorphous Mg–La–Ni alloys were firstly reported by Inoue et al. [88, 89] in early 1990s. Afterwards, some groups started to study the hydrogen storage properties of melt-spun Mg-based alloys, including Köster's group [56, 90], Tanaka's group [91], Yartys's group [53, 92, 93], Schultz's group [94], Rontzsch's group [95,96,97], Zhang's group [98], Zhu's group [16, 99], etc.

Spassov et al. [54, 90] reported that some nanocrystalline and/or amorphous Mg–RE–Ni alloys show much superior hydrogen absorption/desorption properties than the corresponding crystalline counterparts. They found that the melt-spun Mg75Ni20Mm5 (Mm = Ce or La-rich mischmetal) alloy shows quick hydrogen absorption kinetics. Kalinichenka et al. [100] compared hydrogen absorption and desorption properties of several melt-spun Mg-based alloys, including Mg90Cu2.5Ni2.5Y5, Mg85Cu5Ni5Y5 and Mg80Cu5Ni5Y10 alloys. The activation procedure and the hydrogen sorption kinetics of these alloys were studied by thermogravimetry at different temperatures in the range of 100–380 °C. It has been found that these alloys can reach reversible gravimetric hydrogen storage densities of up to 4.8 wt%. Lass [52] studied nanostructured Mg85Ni15−xMx (M = Y or La, x = 0 or 5) alloys by devitrification of amorphous and amorphous-nanocrystalline precursors which were produced via melt-spinning, and he found that all three alloys exhibited hydrogen storage capacity up to 5 wt% H at 300 °C. Zhang et al. [51] prepared an amorphous Mg–Y–Ni alloy together with a long-period stacking ordered (LPSO) phase. It was demonstrated that the nanostructured YH2 and Mg2Ni significantly improved the hydrogen absorption and desorption kinetics of MgH2. Zhang et al. [101] also comparatively studied the hydrogen storage properties of melt-spun Mg10NiRE alloys, and it was found that the hydrogen absorption/desorption kinetics were mainly influenced by the particle size but not by the species of the REHx, and hydrides with finer particles display faster kinetics than those with coarser particles. Wu et al. [93, 102] studied the effects of solidification rate on the microstructures and hydrogen storage kinetics of melt-spun Mg–Ni–Mm alloys. The hydrogenation properties showed closed relation with their microstructures; the refined microstructure is very beneficial for the hydrogen storage properties of Mg-based materials.

It should be pointed out that because the crystallization temperature of Mg-based alloys is generally below 200 °C, the Mg-based amorphous alloys can be easily crystallized during hydrogenation/dehydrogenation processes. In order to avoid the crystallization issue, Lin et al. [16] studied the hydrogen storage properties of a series of Mg–Ce–Ni amorphous alloys at a temperature as low as 120 °C. As shown in Fig. 3a, the amorphous Mg80Ce10Ni10 alloy can readily absorb hydrogen at 120 °C without pre-activation treatment. X-ray diffraction (XRD) patterns show the evolution of amorphous alloys to amorphous hydride upon hydrogenation (Fig. 3b), and high-angle annular dark field scanning transmission electron microscopy (HADDF-STEM) images clearly show the differences of atomic structures of the amorphous alloy and hydride (Fig. 3c, d). Because this glassy system can be amorphized in a wide chemical compositional range and it may be tuned by alloying, results show that by alloying of only 5 at% of Cu, the dehydrogenation temperature of the Mg–Ce–Ni alloys can be reduced by about 150 °C (Fig. 3e). The dehydrogenation temperature of the alloys shows a close relation with the enthalpy of mixing for the X–H pairs, which shows a reduction tendency of dehydrogenation temperature with the increase of mixing enthalpy (Fig. 3f). Moreover, hydrogen storage capacity of the amorphous Mg–Ce–Ni alloys may be higher than that of their crystalline counterparts. The results suggest that the amorphous structure may provide more hydrogen occupation sites.

Fig. 3
figure 3

Reproduced with permission from Ref. [16]. Copyright 2016, Elsevier

a Hydrogenation kinetics of Mg80Ce10Ni10 amorphous alloy under an initial pressure of 4.5 MPa at 120 °C; b XRD patterns with increase of hydrogen concentration; HAADF STEM images of Mg80Ce10Ni10 MG c before and d after full hydrogenation; e TG profiles of a series of fully hydrogenated quaternary Mg−Ce−Ni−X (X = Ni, Ti, Co, Cu, Zn and Ag) amorphous hydrides; f enthalpy of mixing for X−H pairs, which shows a reduction tendency of dehydrogenation temperature with increase of mixing enthalpy.

Full size image

Alloying is an efficient strategy to tune the hydrogen storage properties of materials. It was reported by Lin et al. [103] that alloying of only 1 at% Ag greatly improves the hydrogen storage capacity and reduces the dehydrogenation temperature of the amorphous Mg65Cu25Y10 alloy. Moreover, it was reported by Lin et al. [104] that alloying with Zn shows negligible impact on the hydrogenation kinetics and storage capacity of the Mg–Ce–Ni–Cu amorphous alloys; however, alloying with Co remarkably improves the hydrogenation kinetics and storage capacities. Combined with theoretical calculations and hydrogen storage kinetics, it seems that geometry issue is not a key factor influencing the hydrogenation properties of the studied Mg-based amorphous alloys. It is proposed that chemical composition is the key to improve the hydrogenation properties of Mg-based amorphous alloys.

Besides Mg-based amorphous alloys, Zr-based glassy materials have also been studied as potential hydrogen storage materials. Ismail et al. [105] studied the hydrogen absorption properties of a Zr55Cu30Al10Ni5 amorphous alloy, and found that the material could absorb hydrogen up to 95 wt%. Hydrogenation would cause the formation of Zr-hydride. This will cause the depletion in the number of free Zr atoms, leading to different crystalline phases upon crystallization when the temperature is above 350 °C. As a result, these Zr-based amorphous alloys are lack of hydrogen storage reversibility, thus reducing the interest in their study. Hydrogen-sorption experiments on a Mg54Cu28Ag7Y11 bulk metallic glass indicate that the as-cast alloy exhibits significantly larger enthalpy of desorption, compared to the fully crystallized state; therefore, it was assumed that the local atomic structure of the glass can be accountant for the observes of hydrogen release [106].

3.3 Nanoglass hydrogen storage alloys

Nanoglass is a new concept of materials consisting of amorphous nanoscale grains connected by amorphous interfaces. The concept of nanoglass alloys was firstly reported by Gleiter et al. in late 1980s by consolidating glassy droplets with diameters of less than 10 nm [107]. The Pd70Fe3Si27 droplets were produced by evaporating the initial substances in a cold He atmosphere. In the latter two decades, metallic nanoglass alloys were reported only by few groups [108,109,110,111,112], and only very limited chemical compositions were successfully prepared due to the harsh preparation conditions and atmosphere as well as the limitations of the chemical composition suitable for the preparation.

Lin et al. recently reported a Mg-based Mg–Ce–Ni–Cu nanoglass alloy which was prepared by HPT upon melt-spun amorphous alloy [113]. Figure 4a shows optical image of the bulk nanoglass alloy. Although some crystallites were introduced by severe plastic deformation (Fig. 4b), the main matrix of the amorphous alloy remains amorphous after HPT treatment. However, abundant boundary-like defects were generated in the amorphous alloy matrix to generate a new nanoglass structure (Fig. 4c). As a result, the hydrogen absorption temperature of the nanoglass was greatly reduced with significant improvements on the kinetics upon HPT treatments (Fig. 4d).

Fig. 4
figure 4

Reproduced with permission from Ref. [17]. Copyright 2018, Elsevier

a Optical image of a nanoglass disk; b XRD patterns of melt-spun Mg65Ce10Ni20Cu5 alloy before and after HPT process for 1, 5 and 10 turns; c high resolution transmission electron microscope (HRTEM) image of sample after 1 turn; d hydrogenation kinetics of melt-spun and HPT-treated alloys under initial pressure of 3.0 MPa and temperature of 120 °C.

Full size image

3.4 Quasicrystal hydrogen storage alloys

Titanium-based quasicrystals have been studied as hydrogen storage materials since the 1990s. Viano et al. [18] reported that the icosahedral phase (i phase) in a Ti–Zr–Ni alloy demonstrates hydrogen absorption from the gas phase at a temperature of 260 °C and a pressure of 4 MPa. Hydrogenation causes the quasilattice to expand with an increase of about 7%. The i phase absorbs the most hydrogen among the phases, giving a hydrogen atom to metal atom ratio (H/M) of 1.6. Stroud et al. [57] reported a stable Ti45Zr38Ni17–H quasicrystal, which shows reversible hydrogen storage in a quasicrystalline phase.

Wang et al. studied the electrochemical hydrogen storage properties of a series of melt-spun quasicrystalline (Ti1−xVx)2Ni (x = 0.05, 0.1, 0.15, 0.2 and 0.3) ribbons [58]. They found that the discharge capacity of the alloy electrodes reached a maximum value of 271.3 mAh·g−1, and the best high-rate discharge ability of 82.7% at the discharge current density of 240 mAh·g−1 as x is 0.3. In a later study, they studied the electrochemical hydrogen storage properties of quasicrystalline Ti1.6V0.4Ni1−xCox (x = 0.02–0.1) alloys. The addition of appropriate amount of Co could enhance the cycling stability yet at the expense of high-rate discharging ability [59].

Lee et al. [114] prepared a metastable TiZrNi quasicrystal via a rapid quenching method and a series of pressure–composites–temperature (P–C–T) curves for the sample were measured. Studies showed that the coherence length was increased from 18 to 26 nm after introducing hydrogen into the metastable quasicrystal ribbons, suggesting an improved atomic order in quasicrystal. In the meanwhile, no hydride phase was formed after hydrogen absorption and a quasicrystal phase was still preserved, demonstrating a superior thermal stability, which would possibly be applied as hydrogen storage materials. Later, they found the equilibrium vapor pressure of TiZrNi quasicrystal was significantly increased by adding Pb and V, while the hydrogen storage content was decreased [115, 116]. The quasi-lattice constants of hydrogenated Ti53Zr27−xNi20Pbx (where 0 ≤ x ≤ 8) sample were changed from 5.13 to 5.38 nm. In addition, it was obviously seen that the formation of a Laves phase preliminary manifested at x = 13 and completely transformed at x = 15 for TiZrNi quasicrystals. Moreover, Balcerzak [117] demonstrated that larger content of Zr could improve the hydrogen storage capacity and equilibrium pressure of TiZrNi, but corresponding low reversibility of TiZrNi was not satisfied for commercial hydrogen storage, as demonstrated in Fig. 5. Huang et al. [118] confirmed the possibility that two stable structural models of W–TiZrNi could be obtained according to the characteristics and preference site of W–TiZrNi, and the values of specific heat and vibrational entropy were further decreased with Zr content decreasing.

Fig. 5
figure 5

Reproduced with permission from Ref. [117]. Copyright 2016, Elsevier

PCI a absorption and b desorption curves obtained at different temperatures for Ti1.5Zr0.5Ni and Ti1.75Zr0.25Ni.

Full size image

Luo et al. [119] studied the gaseous hydrogen storage properties of quasicrystalline Mg–Zn–Y alloys. The pressure–composition–isotherm (PCI) curve of the quasicrystalline Mg–Zn–Y alloy shows a flat plateau with a pressure of 0.35 MPa at 300 °C, indicating that the dehydrogenation enthalpy is lower than that of MgH2.

3.5 High-entropy hydrogen storage alloys

High-entropy alloys (HEAs) are a new class of metallic materials which have received increasing attention due to their excellent mechanical and functional properties [120, 121]. As a type of metastable alloys, HEAs have also been tremendously studied as hydrogen storage materials in recent years [122,123,124,125,126,127].

Kunce et al. [19] synthesized a TiZrNbMoV HEA by laser engineered net shaping (LENS) from a blend of elemental powders. The TiZrNbMoV HEA can absorb hydrogen at 50 °C under a pressure of 8.5 MPa without any prior activation. The maximum hydrogen storage capacity of the alloy was 2.30 wt% after synthesis and 1.78 wt% after additional heat treatment. Zhang et al. [126] reported the hydrogen absorption properties, including hydrogen storage capacity, thermodynamics, kinetics and cyclic properties, of a sophisticatedly designed TiZrNbTa HEA. It was showed that multiple hydrides including ε-ZrH2, ε-TiH2 and β-(Nb,Ta)H were found in the TiZrNbTa HEA upon hydrogen absorption. The maximum hydrogen storage capacity decreased from 1.67 wt% to 1.25 wt% as temperature increased from 20 to 220 °C. It was also showed that the TiZrNbTa HEA exhibited a rapid hydrogen absorption kinetic even at room temperature with a short incubation time, indicating a nucleation and growth process for hydrogenation.

Edalati et al. [125] reported the reversible hydrogen storage performance of a TiZrCrMnFeNi HEA. The sample can absorb 1.7 wt% hydrogen without thermal hydrogenation/dehydrogenation activation process with very fast kinetics and fully hydrogenation within 1 min. The cell volume of hydride phase is 40% larger than that of the HEA, which is due to the occupation of hydrogen atoms that expand the lattice. The cycling PCT isotherms show that the material remains almost the same capacity after three cycles, with a capacity of 1.7 wt% in the third cycle. Zlotea et al. [123] studied the hydrogen absorption and desorption as well as the cycling property of a TiZrNbHfTa HEA by in situ synchrotron XRD, pressure-composition-isotherm, thermal desorption spectroscopy and differential scanning calorimetry (DSC). The alloy undergoes a two-stage hydrogen absorption reaction to an FCC dihydride phase with an intermediate tetragonal monohydride, which is quite similar to the V–H system. The hydrogen absorption/desorption in TiZrNbHfTa is completely reversible with a hydrogen storage capacity H/M of 2. Karlsson et al. [128] studied the hydrogenation properties of a HfNbTiVZr HEA, which was prepared by arc melting followed by ball milling. They found that the large lattice strain in the HEA seems favorable for absorption in both octahedral and tetrahedral sites; therefore, HEAs may show great potential for hydrogen storage applications.

3.6 Immiscible hydrogen storage alloys

Under equilibrium state, Mg is immiscible with some transition metals such as Ti, V, Cr, Mn, Fe, Co, Zr, Nb, Mo, Hf, Ta. Ultra-high-pressure technique is one of the attractive methods to synthesize novel metal hydrides. Kyoi et al. synthesized a series of Mg-rich hydrides Mg6–7MH12–16 (M = Ti [62], V [63], Nb [64], Ta [65] and Zr [66]) by using a gigapascal high-pressure thermal technique. Compared to MgH2, the hydrogen desorption temperature of Mg7MHx hydrides is lowered by 130–190 °C. For M = Ti, V, Nb, Hf and Ta, the metal atoms have the Ca7Ge-type super-lattice structure, leading to a prototype unit formula of Mg7MHx. During hydrogen desorption, Ca7Ge-type lattices decompose into Mg and M-hydride phases. With respect to M = Zr, an FCC hydride phase with no super-lattice structure was formed. Even after hydrogen release, the Mg–Zr FCC lattice is maintained. Dong et al. [129] studied the structures of high pressure synthesized MgZr2H6 and MgNb2H6. Both MgZr2H6 and MgNb2H6 are described in the R-3m space group with the hexagonal unit cells.

Compared with FCC and HCP lattices, BCC lattice has lower packing efficiency and consequently provides more interstitial sites for hydrogen occupancy. Akiba et al. prepared Mg–Co metastable alloys with BCC structure by mechanical alloying [130]. The BCC structure formation range was from 37 at% to 80 at% Co [131]. Most of the Mg–Co BCC alloys absorbed hydrogen at 100 °C under a hydrogen pressure of 6 MPa. However, all the Mg–Co alloys could not release hydrogen at 100 °C. After hydrogenation at a temperature below 100 °C, the BCC structure in the Mg50Co50 [130] and Mg55Co45 [15] alloys is maintained. Even at a low temperature of − 15 °C, the Mg50Co50 and Mg55Co45 alloys absorbed 2.67 wt% and 3.24 wt% hydrogen, respectively, under a hydrogen pressure up to 8 MPa. The excellent hydrogen absorption performance was due to the nanometer-scale crystallites, the large surface area, the numerous defects, and the BCC structure.

HPT is a technique to create ultrafine-grained structures with promising mechanical and functional properties by severe plastic deformation. Edalati et al. found that several new metastable phases with the HCP or/and BCC or/and FCC structures were formed in the immiscible systems of Mg–Ti [73] and Mg–Zr [132] through the process of HPT. At room temperature, the HPT-processed Mg–Zr phases can absorb ~ 1 wt% hydrogen under 9 MPa within 20 s and completely release the hydrogen in the air atmosphere [132]. The studies showed that the HPT method was effective not only for increasing the hydrogenation kinetics but also for improving the hydrogenation activity. In the immiscible Mg–Hf system, a metastable FCC structure was also generated by HPT [133]; however, the preparation of Mg–Hf hydrides has been unsuccessful so far.

Anastasopol et al. [134] synthesized Mg–Ti nanocomposites containing a metastable BCC Mg–Ti alloy phase via spark discharge generation. Upon hydrogenation, the BCC Mg–Ti alloy phase transformed into the FCC fluorite-type Mg1−yTiyHx phase, which exhibits promising hydrogen storage properties. A much less negative enthalpy of formation of about − 45 kJ·mol−1 H2 was observed for the Mg–Ti–H nanocomposites than that for bulk MgH2 (ΔH =  − 75 kJ·mol−1 H2). At the same time, the reaction entropy was reduced to 84 J·K−1·mol−1 H2. Calizzi et al. [135] prepared Mg–Ti nanostructured samples by inert gas condensation. When Ti content was 22 at%, a metastable Mg–Ti–H FCC phase was observed after in situ hydrogenation. However, the enthalpy of hydride formation was almost the same as that of bulk MgH2.

Lu et al. [136, 137] have synthesized MgH2–Mn and MgH2–Cr composites through high-energy ball milling, in which nanometer-sized MgH2 embedded in an immiscible Mn or Cr matrix is formed. The samples can reversibly absorb and desorb hydrogen at a low temperature of 200 °C. The large lattice distortion in the nanometer-sized MgH2 should be contributed to the destabilized thermodynamics and improved kinetics. Ding et al. [138] prepared micrometer-sized Mg–Zr–H composites by the reactive ball milling and isothermal treatment of MgH2 and Zr powders under hydrogen gas. The Mg–Zr–H composites absorbed 4.0 wt% and 2.5 wt% H2 in 1 min at 200 and 150 °C, which showed accelerated hydrogen sorption at low temperatures. This improvement is attributed to the synergistic effect of chemical catalysis, grain refinement, dislocations and interfacial coupling induced by the presence of immiscible Zr. The mechanism for the enhanced hydrogen absorption and desorption performance of the Mg−Zr−H composites is proposed, as shown in Fig. 6.

Fig. 6
figure 6

Reproduced with permission from Ref. [137]. Copyrights 2017, The Royal Society of Chemistry

Schematic diagram of proposed mechanism for fast hydrogen desorption/absorption of MgH2 through Mg–Zr–H interfaces, which provides a mystical channel for shuttling back and forth (note that interface is marked by a solid yellow line, green particle is ZrH2 phase, dark yellow bulk is MgH2 matrix, blue bulk is Mg matrix and symbol ⊥ represents dislocation)

Full size image

3.7 2D hydrogen storage films

Thin films often exhibit different properties from those of bulk and power materials. Mg/Ti thin film multilayers were deposited on rotating substrates under an argon pressure using an ultra-high vacuum (UHV) multi-target sputter system. A multilayer stack of Mg/Ti with the thickness of Mg layers below 10 nm showed thermodynamic destabilization of MgH2 due to the energy difference between the MgH2(110)|TiH2(111) and Mg(001)|TiH2(111) interface [76]. Asano et al. [139] reported that the nanometer-sized MgH2 clusters coherently embedded in a TiH2 matrix were formed after hydrogenation of Ti–rich Mg–Ti alloy thin films. After that, they obtained even a more destabilized Mg hydride with an increased interface energy difference by the addition of Cr, which is immiscible with Mg as Ti. Moreover, Cr could catalyze the hydrogen absorption [140]. Vermeulen et al. [141] studied the electrochemical properties of MgyTi1−y thin films (0.50 ≤ y ≤ 0.95) during de-/hydrogenation. These metastable alloys were prepared by electron-beam deposition at room temperature. Mg0.80Ti0.20 alloy showed a superior reversible hydrogen storage capacity as well as an excellent rate-capability. Upon hydrogenation, an FCC-structured hydride was formed.

Recently, Lu et al. [142] found large thermodynamic destabilization of MgH2 in the Mn-rich MgxMn1−x thin films. For x = 0.30, the increase in hydrogen desorption pressure was ~ 2.5 orders in magnitude, which allowed the MgH2 to reversibly absorb and desorb hydrogen at room temperature. Zheng et al. [143] prepared Mg1−xFex thin films (x = 0–0.30) capped with Pd by electron beam co-deposition and found a substantial improvement in hydrogen absorption and desorption kinetics. For x = 0.05–0.15, over 3.5 wt% hydrogen can be absorbed under hydrogen pressure of 0.1 MPa within 2 min, and more than 3.0 wt% hydrogen can be released in 15 min. The presence of Fe layers percolating throughout the Mg matrix was responsible for the improvement. A metastable hydride (Mg0.75Nb0.25)H2 with BCC structure was discovered in the hydrogenated Mg−Nb alloy thin film [130]. The reversible hydrogen capacity is as high as 4 wt%. A significant thermodynamic destabilization is achieved in this BCC Mg–Nb hydride. Besides, the hydrogen sorption kinetics is superior to that of pure MgH2.

As Y and Zr are immiscible, Ngene et al. [144] studied the hydrogen storage properties of Y–Zr thin films prepared at room temperature in a multi-target ultrahigh-vacuum (UHV) direct current/radio frequency (DC/RF) magnetron sputtering system. The equilibrium hydrogen pressure of YH3 was tuned up to five orders of magnitude at room temperature, which indicated large destabilization. By adding Zr into Y, a compression of the Y lattice is generated, which can be maintained during de-/hydrogenation cycles.

3.8 Metastable metal hydrides for potential hydrogen storage applications

Metastable metal hydrides can offer high hydrogen capacity and release hydrogen at a relatively low temperature. Particularly, aluminum-based metal hydrides have received considerable interests. These hydrides include aluminum hydride (AlH3), lithium alanate (LiAlH4), magnesium alanate (Mg(AlH4)2) and calcium alanate (Ca(AlH4)2) [8, 21, 145, 146]. Owing to their low decomposition enthalpy, these materials require small activation to release hydrogen at practical temperatures and pressures. In addition, these materials can release hydrogen with a rapid rate at a temperature below 100 °C. The aluminum-based metastable metal hydrides cannot be recovered by direct hydrogenation of the elements under moderate temperatures and pressures. Therefore, they were generally used on-board and then regenerated off-board by chemical regeneration method.

3.8.1 Aluminum hydride

AlH3 has a high hydrogen density of 10.01 wt%, which is twice of the density of liquid hydrogen. Its volumetric hydrogen density is 1.48 kg·m−3 under normal temperature and pressure [147]. The discovery of AlH3 can be tracked back to 1947 when Finholt et al. [148] first prepared AlH3 ether complex by a reaction of LiH and AlCl3 in ether solution. In the past half century, AlH3 has been used as rocket propellant, strong reducing agent and polymerization catalyst. In 2005, Sandrock et al. [149] reported that AlH3 could be a potential high-performance hydrogen storage material, which attracted considerable attention of the researchers in the field of hydrogen storage.

There are seven known polymorphs of AlH3, that is α, α′, β, γ, δ, ε and ζ, among which, α-AlH3 is the most stable one [150]. Different polymorphs exhibit different morphologies. Figure 7a displays SEM images of α-AlH3, α′-AlH3, β-AlH3 and γ-AlH3 [151]. The structures of AlH3 polymorphs are all composed of A1–H–A1 bonds with different arrangement of free atoms. In 1969, Turley and Rinn [152] determined that α-AlH3 is a trigonal structure with lattice parameters of a = 0.4449 nm and c = 1.1804 nm. Its structure is composed of octahedral AlH6, Al atoms and H atoms alternately form a plane structure, with a complete bridge, as shown in Fig. 7b. In 2006, Brinks et al. [153] determined the structure of α′-AlH3 to be orthorhombic with lattice parameters of a = 0.6470 nm, b = 1.1117 nm and c = 0.6562 nm. In 2007, Brinks et al. [154] carried out a similar study and determined the structure of β-AlD3 to be cubic with lattice parameter of a = 0.9004 nm. In the same year, Yartys et al. [155] characterized the structure of γ-AlH3 by synchrotron X-ray powder diffraction and determined that the structure of γ-AlH3 is orthorhombic, and the lattice parameters are a = 0.53806 nm, b = 0.73555 nm and c = 0.577509 nm. The crystal structure of γ-AlH3 consists of two types of AlH6 octahedron. Besides the bridge bond Al–H–Al, there are bifurcated double bridge bonds Al–2H–Al.

Fig. 7
figure 7

Reproduced with permission from Ref. [146]. Copyright 2011, Elsevier

a Scanning electron microscopy (SEM) images and b crystal structures of α-AlH3, α′-AlH3, β-AlH3 and γ-AlH3.

Full size image

The decomposition of AlH3 is a one-step process to form the elements. The thermal decomposition of the AlH3 polymorphs was characterized by DSC, as shown in Fig. 8a. α-AlH3 undergoes direct decomposition to form Al and H2, while the other polymorphs (α′-AlH3, β-AlH3, γ-AlH3) will first transform to the more stable α-AlH3 and then decompose. The decomposition of the as-prepared AlH3 generally occurs in the temperature range of 150−200 °C, which is still too high for practical application. Therefore, some efforts have been made to reduce the decomposition temperature and improve the kinetics. The strategies include nanoscaling [149, 156, 157] and addition of catalyst [158, 159]. Nanoconfinement and nanocrystallization are two methods to achieve nanoscaling [160]. Wang et al. [161] used high surface area graphite (HSAG) as support carrier and loaded α-AlH3 into the pore structure of HSAG by solution immersion method. The initial dehydrogenation temperature of the composite decreased significantly from 155 to 60 °C. Moreover, the composite achieves slight reversibility under the conditions of 150 °C and 6 MPa H2. It was also found that the content of AlH3 in HSAG was only 14.4 wt% ± 0.2 wt%, which results in the total hydrogen content of only 1.4 wt%. Nanocrystallization by ball milling or other techniques is another way to achieve nanoscaling. In 2005, Sandrock et al. [149] used ball milling to reduce the particle size of α-AlH3 from 100 to 1 μm, and the initial dehydrogenation temperature of the as-milled α-AlH3 was reduced from 175 to 125 °C. Liu et al. [157] prepared γ-AlH3 by organometallic synthesis and used ball milling to improve the dehydrogenation performance of γ-AlH3. At 97 °C, if 90% hydrogen is released, 280 min is needed for as-prepared γ-AlH3, but only 82 min is needed after 10-h ball milling. Utilization of additives can adjust the activation energy in the reaction process and improve the dehydrogenation performance of AlH3. Sandrock et al. [158] utilized alkaline metal hydrides (LiH, NaH, KH) to tailor the dehydrogenation performance of AlH3 and proved that the modification effect of LiH is the most significant. The initial dehydrogenation temperature of AlH3 system was lower than 100 °C when the addition amount is 20 mol%. In addition, they also found that Ti doping could further improve the low-temperature decomposition performance of AlH3 [158]. Chen et al. [162] introduced Nb2O5 and NbF5 to reduce the initial dehydrogenation temperature of AlH3 to lower than 70 and 60 °C, respectively. Nakagawa et al. [159] suggested that the reaction between NbF5 and AlH3 (surface Al2O3) results in the formation of fine and uniform Nb or AlF3, which enhance the dehydrogenation of AlH3.

Fig. 8
figure 8

Differential scanning calorimeter (DSC) curves of a AlH3 polymorphs and b some metal alanates [144]

Full size image

AlH3 cannot be recovered by direct hydrogenation of the elements under moderate conditions. Wet method refers to reaction occurring in solution. AlH3 can be synthesized by the reaction of LiAlH4 and AlCl3 in the ether solution, followed by desolvation [146]. The chemical reaction formula is as Reaction (1):

$$3{\text{LiAlH}}_{4} + {\text{ AlCl}}_{3} + n{\text{Et}}_{2} {\text{O }} \to \, 4{\text{AlH}}_{3} \cdot n{\text{Et}}_{2} \rm {O} \, + \, 3{\text{LiCl}}$$
(1)

The key to the preparation of nonsolvated AlH3 was to add excessive LiAlH4/LiBH4, and carefully control the desolvation conditions (e.g., time, temperature, atmosphere). Owing to the use of a large quantity of organic solvent, this method has the following disadvantages: (1) high preparation cost, (2) a large amount of waste liquid is produced, and (3) the experiment is dangerous and toxic. To overcome the production of waste liquid, dry synthesis method was developed. Paskevicius et al. [163] used mechanical alloying to synthesize AlH3 by ball milling LiAlH4 and AlCl3 raw materials and a small amount of LiCl buffer at low temperature (− 196 °C), followed by washing with CH3NO2/AlCl3 mixture. Hlova et al. [164] synthesized AlH3 by ball milling LiH and AlCl3 at room temperature and suggested that excessive LiH in the initial mixture is essential for the formation and stability of Al–H bond. The gradual addition of AlCl3 during ball milling can prevent the decomposition of AlH3 and inhibit the production of unstable intermediates. It is found that the competitive reaction of AlH3 decomposition to metal Al can be effectively inhibited at 30 MPa. Mechanical alloying has the advantages of environmental protection and energy saving, but the surface structure of AlH3 will be destroyed in the process of ball milling, resulting in the increase in material defects and the decrease in stability.

At the same time, a certain amount of hydrogen will be lost during the preparation of AlH3 by long-time mechanical ball milling, thus reducing the hydrogen content. Electrochemical regeneration is another method to synthesize AlH3. Zidan et al. [165] reported the synthesis of AlH3 using a reversible circulating electrolyte. Using MAlH4 (M = Li, Na) dissolved in anhydrous (tetrahydrofuran) THF or ether as electrolyte and aluminum sheet as anode, they produced tetrahydrofuran complex of AlH3 as anode product and MH or M3AlH6 (M = Li, Na) as cathode product by consuming electrolyte MAlH4. It was found that Al reacts with NaH and H2 to form electrolyte MAlH4. The key of the preparation method is how to inhibit the reaction of H+ adsorbed on the surface of Al electrode to generate gas. Crouch-Baker [166] used a fluidized bed reactor to prepare AlH3 ether complex. The anode and cathode are aluminum electrodes, and the electrolyte is NaAlH4. NaAlH4 electrolyte was formed in situ by blowing hydrogen into cathode, and then tetrahydrofuran complex solution of AlH3 was formed by in situ electrolysis. This method simplifies the processing steps of cathode products, which is simple and low costly.

Graetz et al. [167] found that the enthalpy of formation and standard Gibbs free energy of α-AlH3 at 25 °C respectively are − 9.9 and 48.5 kJ·mol−1. The equilibrium hydrogen fugacity of α-AlH3 is equivalent to 700 MPa hydrogen pressure. Since is not feasible to regenerate waste aluminum by direct hydrogenation, researchers have been looking for other feasible methods. They found that the amine complex AlH3–TEDA can be obtained by using Ti doped active aluminum and triethylenediamine (TEDA) at low temperature and low pressure. However, AlH3–TEDA complex is too stable to recover AlH3 without decomposition. Lacina et al. [168] obtained a relatively unstable AlH3–TEA complex by substituting triethylenediamine (TEA) for triethylenediamine (TEDA). AlH3 can be obtained by heating the AlH3–TEA complex at 70 °C. This research has a certain theoretical significance for the regeneration of AlH3.

In the early stage, AlH3 was mainly used as rocket propellant and explosive additive, and then AlH3 was gradually used as strong reducing agent, polymerization catalyst, etc. In recent decades, AlH3 is used as a high-performance hydrogen storage material due to its high hydrogen storage capacity. It can be used as an excellent carrier material for hydrogen energy and as a hydrogen supply system for hydrogen fuel cells. Ahluwalia et al. [169] developed a kind of AlH3 slurry, which can easily extract aluminum powder from fuel tank for external regeneration. In addition, they developed a model of on-board hydrogen storage system and studied in detail the performance of on-board hydrogen storage system using aluminum slurry as hydrogen carrier. In 2012, Grew et al. [170] proposed a concept of portable hydrogen fuel cell using AlH3 as the hydrogen supply and calculated and analyzed the feasibility of AlH3 as hydrogen supply for portable hydrogen fuel cell system. The hydrogen fuel cell system contains 0.4 kg AlH3 material. The volume and mass of the system are about 1 L and 1 kg, respectively. The system can operate continuously for about 25 h and provide 600 Wh·kg−1 or 600 Wh·L−1 power supply. In recent years, Thampan et al. [171, 172] developed a new type of wearable portable power supply system, which uses AlH3 material as hydrogen source, and uses a 20-W proton exchange membrane hydrogen fuel cell system to meet the long-term power demand of individual portable power supply. The system has the advantages of small size, light weight and easy to carry (0.7 kg, 6.22 L). It can provide 300-Wh·kg−1 power density output for 24-h missions.

3.8.2 Metal alanates

Metal alanates metastable metal hydrides include lithium alanate (LiAlH4), magnesium alanate (Mg(AlH4)2), calcium alanate (Ca(AlH4)2), etc. The metal alanates are more stable than AlH3 because the complex anion AlH4− is stabilized by the cation such as Li+, Na+, Mg2+, Ca2+ [173,174,175,176,177]. LiAlH4 has a P21/c crystal structure with the isolated AlH4− tetrahedra surrounded by Li atoms [178]. The space group of Mg(AlH4)2 is P3m1 with AlH4− tetrahedra surrounded by six Mg atoms [179]. Ca(AlH4)2 has a crystal structure of Pbca with four-coordinated H in tetrahedra around Al and H eight-coordinated around Ca [180].

Metastable metal alanates are not easily synthesized by direct hydrogenation of the elements under reasonable temperatures and pressures. They are generally synthesized by solution methods. The most common method is a reaction of a binary hydride with a metal halide to form the metal alanates. For example, LiAlH4 can be formed by a reaction of LiH and AlCl3 (Reaction (2)) [139]. Mg(AlH4)2 can be synthesized by a metathesis reaction of MgCl2 and NaAlH4 (Reaction (3)) [174, 181]. However, such methods are generally costly because the expensive alkali metals (Li, Na, etc.) change to stable salts (LiCl, NaCl, etc.), which are difficult to separate. Metastable metal alanates can also be synthesized by the reaction of aluminum hydride with a binary metal hydride following a mechanochemical reaction. For example, Ca(AlH4)2 can be synthesized by a mechanochemical reaction of AlH3 and CaH2 (Reaction (4)) [182].

$$4{\text{LiH }} + {\text{ AlCl}}_{3} + n{\text{Et}}_{2} {\text{O }} \to {\text{ LiAlH}}_{4} + \, 3{\text{LiCl }} + n{\text{Et}}_{2} {\text{O}}$$
(2)
$$2{\text{NaAlH}}_{4} + {\text{ MgCl}}_{2} + n{\text{DEE }} \to {\text{ Mg}}\left( {{\text{AlH}}_{4} } \right)_{2} + \, 2{\text{NaCl }} + n{\text{DEE}}$$
(3)
$${\text{CaH}}_{2} + \, 2{\text{AlH}}_{3} \to {\text{ Ca}}\left( {{\text{AlH}}_{4} } \right)_{2}$$
(4)

The decompositions of metal alanates are not as simple as that of AlH3. LiAlH4 generally first decomposes to Li3AlH6, Al, and H2 at 150 °C and then to LiH, Al, H2 at 200 °C (Fig. 8b) [183]. The first decomposition step is slightly exothermic, while the second step is slightly endothermic. Mg(AlH4)2 undergoes a slightly endothermic decomposition to MgH2 and Al at 110 °C and then the further decomposition of MgH2 (Fig. 8b) [184, 185]. Ca(AlH4)2 will decompose to form CaAlH5 at 150 °C slightly exothermically (Fig. 8b). It should be noted that the first decomposition steps of those metastable metal alanates are generally irreversible under moderate temperatures and pressures. For the decomposition of metastable metal hydrides, additives can significantly reduce the dehydrogenation temperature and improve the decomposition kinetics. These additives include halides, oxides, carbides, etc. [186,187,188,189,190,191,192,193,194].

Like AlH3, metal alanates can be regenerated by chemical methods, which involves a multistep chemical process. For example, LiAlH4 cannot be formed from the direct hydrogenation of LiH + Al, but in the presence of THF, LiAlH4–THF adduct can be formed at 35 MPa H2 [195]. Addition of Ti catalyst will lower the hydrogenation pressure at room temperature [196]. The regeneration process of LiAlH4 can be given as Reaction (5). Liu et al. [197] used dimethyl ether (DME) rather than THF as the stabilizing ligand to regenerate LiAlH4. In this process, DME is a gas at room temperature and is easily removed after hydrogenation. The as-synthesized LiAlH4 possessed 6 wt% capacity after several cycles.

$${\text{LiH }} + {\text{ Al}}* \, + n{\text{THF }} + \, 3/2{\text{H}}_{2} \to {\text{ LiAlH}}_{4} {-}n{\text{THF }} \to {\text{ LiAlH}}_{4} + n{\text{THF}}$$
(5)

The kinetically stabilized aluminum hydride and metal alanates offer new materials for hydrogen storage, which are promising for portable application. These materials are not reversible at practical temperatures and pressures. Therefore, a cost-effective regeneration method is required. In addition, strategies should be carried out to achieve the controllable hydrogen releasing from the metastable metal hydrides.

4 Summary and outlook

In summary, this work on metastable alloys for hydrogen storage applications comprehensively reviewed the recent trends in this field. As metastable structures can be formed under non-equilibrium conditions, several techniques have been developed to efficiently prepare metastable hydrogen storage alloys, including high-energy ball milling, rapid quenching, high-pressure synthesis methods, severe plastic deformation and physical vapor deposition.

Owing to the metastable characteristics of the alloys, their hydrogen storage properties have many excellent features compared with conventional crystalline hydrogen storage alloys. For example, these metastable alloys usually show faster hydrogen storage kinetics and altered thermodynamics compared with the conventional crystalline alloys. In some cases, the metastable alloys may possess higher hydrogen storage capacity than their crystalline counterparts, such as amorphous alloys.

However, it should be pointed out that the metastable phases also show some disadvantages for hydrogenation and dehydrogenation. Particularly, the metastable alloys may decompose into crystalline alloys during the hydrogen storage process because of heating. For example, the crystallization temperatures of Mg-based amorphous alloys are generally below 200 °C; however, they should generally be activated at temperatures as high as 300 °C. Therefore, the hydrogen storage reversibility of metastable alloys in some cases has some obvious limitations.

It is believed that study on metastable hydrogen storage alloys will continue to be a research hotspot in the field of hydrogen storage alloys in the future. We think that the following aspects are worthy to be studied: (1) further improving the preparation efficiency of metastable alloys and developing new preparation techniques is of great significance. (2) Tuning the metastable characteristics of the alloy by altering chemical composition, crystal structure, lattice distortion, etc., has very important influence on the hydrogen storage kinetics and thermodynamic properties of the metastable alloys. (3) Using multi-strategy techniques to further improve the hydrogen storage reversibility of metastable alloys remains a major challenge in this field.