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Article

Direct Synthesis of LiAlH4 from Ti-Doped Active LiAl Alloy

by
Yan Chu
1,†,
Shiwei Fang
1,†,
Yingjue Chen
1,
Xiaoqi Zhang
1,
Jie Zheng
1,
Zhenglong Li
1,
Wubin Du
1,
Wengang Cui
1,
Jian Miao
1,
Yaxiong Yang
1,*,
Yongfeng Liu
2,
Mingxia Gao
2 and
Hongge Pan
1,2,*
1
Institute of Science and Technology for New Energy, Xi’an Technological University, Xi’an 710021, China
2
State Key Laboratory of Silicon Materials and School of Materials Science and Engineering, Zhejiang University, Hangzhou 310027, China
*
Authors to whom correspondence should be addressed.
†
These authors contribute equally to this work.
Inorganics 2025, 13(3), 74; https://doi.org/10.3390/inorganics13030074
Submission received: 19 January 2025 / Revised: 20 February 2025 / Accepted: 24 February 2025 / Published: 1 March 2025
(This article belongs to the Special Issue Advanced Inorganic Nanomaterials for Energy Conversion and Catalysis Applications)
Browse Figures
Graphical abstract
"> Figure 1
<p>(<b>a</b>) Schematic diagram of the preparation process of active LAT alloys; optical photographs of (<b>b</b>) as-cast alloy by induction smelting, (<b>c</b>) alloy flakes by melt spinning, and (<b>d</b>) alloy powder after cryomilling.</p> "> Figure 2
<p>XRD patterns of LAT–b, LAT–2000, and LAT–4000 alloy flakes.</p> "> Figure 3
<p>SEM of (<b>a</b>) LAT-b, (<b>b</b>) LAT-2000, and (<b>c</b>) LAT-4000; corresponding (<b>d</b>–<b>f</b>) EDX mapping (<b>g</b>–<b>i</b>) corresponding BSE images.</p> "> Figure 4
<p>(<b>a</b>) XRD patterns and (<b>b</b>) corresponding crystallite size and strain of LAT–b–CM, LAT–2000–CM, and LAT–4000–CM; SEM images of (<b>c</b>) LAT–b–CM, (<b>d</b>) LAT–2000–CM, and (<b>e</b>) LAT–4000–CM, respectively; (<b>f</b>) TEM and corresponding EDX mapping of LAT–4000–CM.</p> "> Figure 5
<p>(<b>a</b>) XRD patterns and (<b>b</b>) Non-isothermal dehydrogen curves of LAT–b–CM–H, LAT–2000–CM–H, LAT–4000–CM–H samples; (<b>c</b>) hydrogen mass spectrum of LAT–4000–CM–H sample; (<b>d</b>) survey XPS spectra, and (<b>e</b>) Al 2p XPS spectra of LAT–4000–CM–H and as-received LiAlH<sub>4</sub>.</p> "> Figure 6
<p>(<b>a</b>) XRD patterns of LAT–4000–CM and LA–b alloy, (<b>b</b>) XRD patterns, (<b>c</b>) non-isothermal dehydrogen curves, (<b>d</b>) survey XPS spectra, and (<b>e</b>) Al 2p XPS spectra of LAT–4000–CM–H and LA–b–H samples.</p> "> Figure 7
<p>(<b>a</b>) Non-isothermal dehydrogen curves, (<b>b</b>) survey XPS spectra, (<b>c</b>) Al 2p XPS spectrum, and (<b>d</b>) Li 1s XPS spectrum of rehydrogenated LAT–4000–CM–H.</p> ">
Versions Notes

Abstract

:
LiAlH4, characterized by high hydrogen capacity and metastable properties, is regarded as a promising hydrogen source under mild conditions. However, its reversible regeneration from dehydrogenated production is hindered thermodynamically and kinetically. Herein, we demonstrate an active Li–Al–Ti nanocrystalline alloy prepared by melt spinning and cryomilling to enable directly synthesizing nano-LiAlH4. Due to the non-equilibrium preparation methods, the grain/particle size of the alloy was reduced, stress defects were introduced, and the dispersion of the Ti catalyst was promoted. The refined Li–Al–Ti nanocrystalline alloy with abundant defects and uniform catalytic sites demonstrated a high reactivity of the particle surface, thereby enhancing hydrogen absorption and desorption kinetics. Nano-LiAlH4 was directly obtained by ball milling a 5% Ti containing Li–Al–Ti nanocrystalline alloy with a grain size of 17.4 nm and Al3Ti catalytic phase distributed under 20 bar hydrogen pressure for 16 h. The obtained LiAlH4 exhibited room temperature dehydrogenation performance and good reversibility. This finding provides a potential strategy for the non-solvent synthesis and direct hydrogenation of metastable LiAlH4 hydrogen storage materials.

Graphical Abstract">
Graphical Abstract

1. Introduction

Hydrogen is regarded as a critical carrier for achieving low-carbon energy transition due to its zero-carbon emission properties. However, because of hydrogen’s ultra-low density and high diffusivity, the development of safe and efficient hydrogen storage technologies remains a major bottleneck for its large-scale application [1,2,3,4]. Solid-state hydrogen storage materials with high safety, high hydrogen storage density, and convenient transportability have emerged as a highly promising solution [5,6]. Among these materials, LiAlH4 with a high hydrogen capacity (gravimetric: 10.5 wt%, volumetric: 96.6 g·L−1) and relatively moderate dehydrogenation temperature (~180 °C) has been extensively studied [7,8,9,10,11,12,13,14]. Nevertheless, the dehydrogenation process is accompanied by poor reversibility, severely limiting the potential of LiAlH4 for cyclic applications and direct synthesis [15].
The decomposition of LiAlH4 is generally accepted to proceed through three distinct steps (LiAlH4 → Li3AlH6 → LiH/Al → LiAl), with the initial step being an exothermic reaction (−10 kJ·mol−1·H2), rendering the regeneration of LiAlH4 thermodynamically unfavorable [16,17]. At ambient temperature, the direct hydrogenation of Li3AlH6 to LiAlH4 requires hydrogen pressures exceeding 103 bar [18,19]. Only part of Li3AlH6 can be reversibly formed without LiAlH4 when LiH/Al hydrogenation is performed with a titanium-based catalyst [20,21,22,23]. Furthermore, sluggish hydrogenation kinetics impose stringent synthesis conditions [20,21,22,23]. One strategy for regulating reaction thermodynamics is stabilizing LiAlH4 by forming a complex with the ether (THF, Et2O, Me2O) [24,25,26,27]. In 1947, LiAlH4 was first successfully achieved via the reaction of LiH with AlCl3 in an ether solution [28]. This kind of liquid-phase synthesis method to stabilize LiAlH4 by forming complexes with ethers is a feasible synthesis strategy. In addition, additional active sites, heterogeneous interfacial reactions, and [AlH4]– nucleation can be promoted by the introduction of a catalyst, thereby improving the hydrogen absorption and desorption kinetics of metal hydrides [29,30,31,32,33,34,35,36]. Nevertheless, liquid-phase synthesis suffers complicated subsequent purification because of extensive organic solvent use. Therefore, a possible greener solid-phase synthesis method should urgently be explored and developed.
LiAl alloy, the fully dehydrogenated phase of LiAlH4, crystallizes in a NaTl-type structure (B32), where Li and Al are mixed at the atomic level in a 1:1 ratio, which shows theoretical advantages in hydrogenation to form LiAlH4. Moreover, metallic Ti, as an effective metal catalyst for aluminum hydride, can be flexibly introduced into the alloy preparation. However, given the low solubility of Ti in Al (only 0.2 at%) and its immiscibility with Li [37], ensuring the uniform dispersion of Ti within the LiAl alloy matrix presents a critical technical challenge [37,38]. Herein, two non-equilibrium preparation methods, melt spinning and cryomilling, were combined to effectively reduce the alloy grain/particle size and promote the uniform dispersion of the Ti catalyst. A Li–Al–Ti nanocrystalline alloy with a grain size of 17.4 nm and Al3Ti catalytic phase dispersive distribution was successfully prepared. Nano-LiAlH4 could be directly synthesized by the Li–Al–Ti alloy with the addition of 5% Ti through ball milling for 16 h under a hydrogen pressure of 20 bar. The obtained LiAlH4 exhibited room temperature dehydrogenation performance and good reversibility.

2. Results and Discussion

The lithium–aluminum–titanium (denoted as LAT) alloys were prepared by suspension smelting, melt spinning, and cryomilling techniques, as schematically illustrated in Figure 1a. Firstly, lithium foils, aluminum granules, and titanium granules were induced to melt into an alloy (Figure 1b) with an atomic ratio of 1:1:0.037 (the mass ratio of titanium in the alloy was 5 wt%). The alloy flakes (Figure 1c) with different cooling rates were obtained by controlling the rotation speed of the copper roller in a vacuum chamber. In this experiment, the copper roller speed varied between 0, 2000, and 4000 rpm, and the samples obtained were denoted as LAT–b (the as–cast alloy), LAT–2000, and LAT–4000. Then, the obtained alloy flakes were cryomilled at liquid nitrogen temperature to obtain alloy powder (Figure 1d) as active LAT alloys, which were denoted as LAT–b–CM, LAT–2000–CM, and LAT–4000–CM, respectively.
The X-ray diffraction (XRD) patterns of LAT alloys prepared with different cooling rates are displayed in Figure 2. It was found that after adding 5 wt% of metallic Ti to the LiAl alloy, the ternary alloy contains two phases, the characteristic reflections of LiAl dominate the XRD profile, along with a small number of diffraction peaks of Al3Ti. It is worth noting that the intensities of the LiAl alloy diffraction peaks decreased while the diffraction peaks of the Al3Ti phases became stronger as the rotational speed of the copper roll increased. Only the (103) crystal plane diffraction peak of the Al3Ti phase was observed in the LAT–b sample. With the increase in cooling rate, other diffraction peaks of the Al3Ti phase could be further detected in the XRD spectra of the LAT–2000 and LAT–4000 samples.
The morphology of the LAT alloy prepared at different rotational speeds is demonstrated as shown in Figure 3a–c. A relatively smooth and homogeneous surface morphology was observed in the as-cast alloy LAT–b. For the LAT–2000 alloy flakes, two phases with different contrasts were observed: the substrate was 1–3 μm dark contrast spherical particles with some ~10 μm bright contrast rectangular lamellas embedded (Figure 3b). When the rotational speed of the copper roller was increased to 4000 rpm, the high contrast phases in the obtained LAT–4000 sample were transformed into particles with a size of 5~8 μm (Figure 3a). The elemental distribution of Ti in the three alloys under different cooling rates was further characterized through energy-dispersive X-ray spectroscopy (EDX) mapping analysis (Figure 3d–f) and backscattering electron (BSE) imaging (Figure 3g–i). According to the EDX spectra, the alloy was mainly composed of Al and Ti. Owing to the limited solubility of Ti in Al alloys (0.2 at.%), the as-cast LAT-b alloy exhibited significant Ti segregation, with distinct Ti-depleted and Ti-enriched regions (Figure 3d,g). In Ti-depleted regions, finer Al3Ti grains were observed (Figure 3d), whereas Ti-rich regions contained coarser particles (Figure 3g, displaying brighter contrast due to the higher atomic number of Ti). In the melt-spun LAT–2000 alloy flake, the distribution of Ti elements coincided with the rectangular lamellae structure with high contrast (Figure 3e,h). Similarly, in the LAT–4000 alloy, Ti was predominantly concentrated in high-contrast particulate regions (Figure 3f,i). Combined with XRD analysis, the bright rectangular lamellae in LAT–2000 and the bright particles in LAT–4000 were identified as the Al3Ti phase, while the dark, fine-grained matrix corresponded to the LiAl alloy phase. Notably, as the cooling rate increased, the Al3Ti phase distribution became more homogeneous, accompanied by a reduction in grain size.
In the as-cast alloy, large metal grains and inhomogeneous phase distribution existed due to the slow cooling of the high-temperature metal liquid in the crucible. During the melt spinning process, the cooling rate increased as the copper roll speed increased to 2000 rpm. The Al3Ti with a higher melting point (1200 °C) [39] in LAT–2000 nucleated preferentially, which provided the nucleation site for the grain refinement of the LiAl alloy. At this time, Al3Ti phases presented rectangular lamellar structure due to a moderate cooling rate [40] and were embedded in the LiAl alloy matrix. When the copper roller speed increased to 4000 rpm, the nucleation rate of Al3Ti in the LAT–4000 sample was further accelerated due to the higher cooling rate, and the grain refinement of the LiAl alloy was also promoted. The LAT–4000 alloy with irregular Al3Ti particles embedded in a small-particle LiAl alloy matrix was obtained. By the non-equilibrium preparation method, the grain size of the LiAl alloy was refined and the catalyst phase Al3Ti was evenly distributed in the alloy phase [41], which is of critical importance for improving the hydrogen storage kinetics of Ti-doped LiAl alloy.
In order to further refine the alloy particle and grain size, the LAT–b, LAT–2000, and LAT–4000 samples were cryomilled for 1 h, and the as prepared samples were noted as LAT–b–CM, LAT–2000–CM, and LAT–4000–CM, respectively. XRD profiles of the LAT alloys after cryomilling show the obvious broadening of the diffraction peak and decreasing of the diffraction intensity (Figure 4a), implying that the grains of the alloy were obviously refined after freezing ball milling, and a large number of stress–strain and other defects were introduced [42,43]. The grain size and strain of the alloy were calculated via the Williamson–Hall method [44]. As shown in Figure 4b, the grain sizes of LAT–b–CM, LAT–2000–CM, and LAT–4000–CM are 34.4, 20.5, and 17.4 nm, respectively. In addition, the microstrain of all samples reached ~0.4% owing to the severe plastic deformation induced by cryomilling. It is worth pointing out that the Al3Ti phase in the alloys changed from the tetragonal structure (PDF#65–7847) to the metastable simple cubic structure (PDF#49–1446) after cryomilling [45]. Further scanning electron microscope (SEM) observations revealed that after 1 h of cryomilling, the alloy flakes were transformed into small near-spherical alloy particles, and the particle size of the alloys was reduced to less than 5 μm (Figure 4c–e). Moreover, the particle size distribution of LAT–4000–CM is more uniform (Figure 4e). The transmission electron microscope (TEM) results further confirm that the element Ti is uniformly dispersed in the alloy particles of the LAT–4000–CM sample after the cryomilling (Figure 4f). These results demonstrate the significant grain/particle refinement, Al3Ti phase transformation, and Ti element uniform distribution in the alloy particles after cryomilling treatment, leading to the formation of highly active Li–Al–Ti alloy powder.
The active alloy powders, after being cryomilled, were hydrogenated via ball milling for 16 h under 20 bar hydrogen pressure. The hydrogenated samples were labeled as LAT–b–CM–H, LAT–2000–CM–H, and LAT–4000–CM–H. The XRD patterns (Figure 5a) show that crystalline phases of LiH and Al appeared in all hydrogenated samples, indicating that the same reaction, 2LiAl + H2 → 2LiH + 2Al, occurred during ball milling due to the favorable hydrogen affinity of Li [46]. The non-isothermal hydrogen desorption results are shown in Figure 5b, LAT–4000–CM–H began to desorb at room temperature ~25 °C, while LAT–b–CM–H and LAT–2000–CM–H started at approximately 40 °C. Upon heating to 180 °C, about 0.25, 0.44, and 0.58 wt% H2 were desorbed by LAT–b–CM–H, LAT–2000–CM–H, and LAT–4000–CM–H, respectively, and after being further heated to 300 °C, 0.32, 0.60, and 0.79 wt% dehydrogenation capacities were obtained. The LAT–4000–CM–H sample exhibited a low initial desorption temperature and better hydrogen storage capability within 300 °C. The mass spectrum of LAT–4000–CM–H (Figure 5c) showed three desorption peaks at 124, 404, and about 500 °C, respectively, which further confirmed that the gas released in the low temperature (<180 °C) was hydrogen. Combined with the XRD results, the produced LiH with high thermal stability is responsible for the dehydrogenation behavior above 380 °C (Figure 5a,b). X-ray photoelectron spectroscopy (XPS) characterization was further used to analyze the possible hydrogenation products of the alloy powder after hydrogenation. As shown in Figure 5d, the survey spectra of LAT–4000–CM–H indicated the presence of Li, Al, and oxide species. Although the vacuum level was strictly controlled throughout the experiment, and all samples were stored and transferred without exposure to water or oxygen, substantial amounts of lithium oxides and aluminum oxides were inevitably generated on the surfaces of the hydrogenated products due to the high chemical activity of the Li–Al–Ti alloy powders [47,48,49,50]. The presence of characteristic diffraction peaks for oxidizing species of aluminum and lithium (primarily composed of Al2O3 and Li2O) between 30 and 35 degrees in the XRD pattern provides additional evidence to support this observation (Figure 5a). In the Al 2p spectra of the LAT–4000–CM–H sample (Figure 5e), four distinct peaks were observed at 74.9 eV, 74.1 eV, 72.7 eV, and 72.4 eV, corresponding to the binding energies of Al2O3, LiAlH4 [8], Al, and Al3Ti [51], respectively. It indicates that LiAlH4 formed on the particle surface under mechanical chemical forces with Al3Ti being catalytic, which contributes to the unique low-temperature hydrogen desorption behavior beginning at room temperature (Figure 5b).
According to the above hydrogenation properties of LAT alloys with different cooling rates, it can be seen that the LAT–4000–CM samples with small grain/particle sizes and uniform distribution of the Ti catalyst demonstrated better hydrogenation reaction ability. To compare the key roles of cooling rate, Ti addition, and defect-rich metastable structure on the hydrogenation kinetics of LiAl alloy, the as-cast LiAl alloy (LA–b) without a catalyst was prepared by direct melting. The structure and hydrogen ab/desorption properties of LA–b and LAT–4000–CM samples were compared. The XRD pattern confirmed that LA–b had a pure NaTl–type crystal structure (PDF#71–0362) and a sharp diffraction peak (Figure 6a). The calculated grain size of LA–b was about 139.3 nm, which is 10 times that of the Ti-doped LAT–4000–CM treated by melt spinning and cryomilling (Figure 4b, 17.4 nm).
After the hydrogenation reaction ball milling at 20 bar hydrogen pressure for 16 h, only diffraction peaks of LiH and Al were observed in both obtained products LA–b–H and LAT–4000–CM–H (Figure 6b). However, as shown in the non-isothermal hydrogen desorption results in Figure 6c, there were significant differences in their hydrogen dehydrogenation behaviors. LA–b–H began hydrogen desorption at 98 °C, and only ~0.10 wt% was released when heating up to 300 °C, while LAT–4000–CM–H released hydrogen at room temperature (25 °C) and the dehydrogenation capacity at 300 °C was determined to be 0.79 wt%, nearly 8 times that of LA–b–H. These results suggested that more aluminum hydrides were generated during the hydrogenation of LAT–4000–CM. The XPS survey spectra (Figure 6d) confirmed the presence of Li, Al, and oxide species in both the LA–b–H and LAT–4000–CM–H samples. Owing to the large grain/particle size, limited specific surface area, and relatively low chemical activity of the as-cast LA–b alloy, the extent of surface oxidation of Al in the samples (Figure 6e) was comparatively mild. Al 2p spectra of LAT–4000–CM–H and LA–b–H samples revealed that the surface of LA–b–H was mainly dominated by Al, and only trace amounts of thermodynamically stable Li3AlH6 were found, which is due to the poor interfacial reaction kinetics of as-cast LiAl alloy. However, for the LAT–4000–CM sample, which consists of Ti-doped LiAl nanocrystalline alloys, the hydrogenation reaction tends to directly produce nano-LiAlH4. In this nanocrystalline alloy, Al3Ti nanocatalysts are uniformly dispersed within the LiAl alloy matrix. These Al3Ti particles serve as catalysts for hydrogen dissociation and recombination [39], while simultaneously generating numerous heterogeneous nucleation sites. During hydrogenation, hydrogen molecules dissociate into individual H atoms on the surface of the Al3Ti nanocatalysts and initially gather around the heterogeneous nucleation sites formed by the Al3Ti distribution [52,53,54]. The H atoms then diffuse and dissolve into the surrounding high-defect-state LiAl alloy nanocrystalline, leading to the formation of nano-LiAlH4. The synergistic effects of the enhanced kinetics driven by these catalysts and heterogeneous nucleation sites, combined with the modulation of hydrogenation thermodynamics by the nanocrystalline and metastable defect-rich structures, collectively enable the direct hydrogenation formation of nano-LiAlH4 from the Li–Al–Ti nanocrystalline alloy and its room-temperature hydrogen desorption behavior [6,55,56].
The reversibility of the aluminum hydride obtained from the hydrogenation of Ti-doped LiAl nanocrystalline alloys was further studied. The product of LAT–4000–CM–H dehydrogenated at 300 °C was subjected to 100 bar hydrogen pressure and held at 100 °C for 24 h. The non-isothermal hydrogen desorption results of the rehydrogenated product are shown in Figure 7a, the sample still began hydrogen desorption at room temperature with 0.49 wt% hydrogen released when heated to 300 °C, maintaining 62% of the initial dehydrogenation capacity. We observed that oxides remain present both before and after hydrogen release (Figure 7b–d). In contrast to the dense Al2O3 layer in pure Al, which completely blocks hydrogen diffusion [57], the surface oxide layer of the Li–Al–Ti alloy consists of a mixture of lithium and aluminum oxides (Figure 7c,d). The multicomponent oxide layer, primarily composed of Al2O3 and Li2O grains, serves a dual purpose: it protects the alloy and hydride from further oxidation, while its non-dense structure, combined with the presence of numerous oxide grain boundaries, facilitates potential pathways for hydrogen adsorption and diffusion during the processes of hydrogen absorption and desorption [58]. The high-resolution Al 2p XPS spectrum also confirmed that the surface of the product was rehydrogenated back to LiAlH4 (Figure 7c), implying that the locally generated nano-LiAlH4 on the alloy particle surface showed certain reversibility. The reported catalysts, including TiC [59], ZrC [59], TiN [59], MnFe2O4 [60], and K2NiF6 [61], doped with LiAlH4 dehydrogenation products, had shown no reversibility for hydrogenation. The after-dehydrogenation absorption capacities of LiAlH4 doped with 5 wt% nano Fe [62] and 7 wt% NiTiO3@h-BN [23] were 0.5 wt% and 1 wt%, respectively, with only Li3AlH6 identified as the hydrogen absorption product. However, no reversible formation of LiAlH4 was achieved. In contrast, nano-LiAlH4 was directly obtained by ball milling a 5% Ti containing Li–Al–Ti nanocrystalline alloy under 20 bar hydrogen pressure for 16 h.

3. Materials and Methods

3.1. Preparation of the Li–Al–Ti Alloy

Lithium tablets (Φ13 mm, purity 99.99%, Xinjiang Nonferrous Metals Industry (Group) Co., Ltd. (Urumqi, China), aluminum granules (Φ6 × 3 mm, purity 99.999%, Zhongnuo Advanced Materials (Beijing) Technology Co., Ltd., Beijing, China), and titanium powder (Φ6 × 3 mm, purity 99.999%, Aladdin, Shanghai, China) were purchased and used as received without further purification. The titanium powder was pressed with a flake grinding tool and then put into a boron nitride crucible with lithium flakes and aluminum particles and transferred to a high-vacuum melting and quenching equipment. The chamber vacuum reached 10−3 bar, then the chamber was filled with argon gas to reach atmospheric pressure. The pressure of the injected molten cast was set to 2 bar, and the distance between the copper rod and the crucible was adjusted to 1 mm. The rotational speeds of the copper rollers were adjusted to be 0, 2000, and 4000 rpm, and then the heating power was elevated until the granules in the crucible were in a molten state. Next, the cast melt was sprayed with argon to obtain Li–aluminum–titanium alloy flakes. A total of 1 g of the alloy flakes and 100 g of hardened steel balls was sealed in a 50 mL hardened steel vial; then, the vial was transferred to a Retsch CryoMill (Haan, Germany) and cryomilled at liquid nitrogen temperature (77 K) with 30 Hz impact frequency for 60 min. In a typical process, 20 min of pre-cooling was performed first to bring the temperature of the tank to the temperature of the liquid nitrogen, and then, ball milling was performed for 5 min to make the alloy with higher energy, with a break of 2 min to eliminate the heat generated in the process of the ball milling.

3.2. Synthesis of LiAlH4 Through Hydrogen-Reactive Ball Milling

LiAlH4 was prepared through hydrogen-reactive ball milling (HRBM) employing a planetary ball milling apparatus (QM–3SP4, Nanjing China). In an argon-atmosphere protected glove box, approximately 1 g of Li–Al–Ti flakes and 100 g of stainless-steel spherical grinding beads were filled into a custom-engineered milling vessel. Then, the jar underwent an evacuation process and was pressurized with hydrogen at 20 bar (purity 99.999%; Xi’an Tianhai Gas Co., Ltd., Xi’an, China). The ball milling speed was set to 400 rpm, the hydrogenated samples were obtained after 16 h of HRBM. For comparison, the as-cast LiAl alloy was hydrogenated by HRBM following an identical grinding process. All samples were handled within an argon-purged glove box (Lab2000, Etelux, Beijing, China, H2O < 0.1 ppm; O2 < 0.1 ppm).

3.3. Structure and Morphology Characterizations

X-ray diffraction, via a PANalytical X’Pert X-ray diffractometer sourced from the Netherlands and Cu Kα (λ = 0.15406 nm, 40 kV, 75 mA) radiation, was applied to characterize the crystal textures within a nitrogen-filled glove box (MBRAUN 200B, Bayern, Germany). Data acquisition was performed within a 2θ angular span ranging from 10° to 90°; a step interval of 0.02° was employed. A transmission electron microscope (TEM, Tecnai G2 F20, Waltham, MA, USA) and a scanning electron microscope (SEM, Hitachi SU–70, Ibaraki Prefecture, Japan) integrated with an energy-dispersive X-ray spectroscopy (EDX) and backscattered electron (BSE) system were employed to examine the microstructural features and elemental distributions of the samples. X-ray photoelectron spectroscopy (XPS, ThermoFisher ESCALAB Xi, Waltham, MA, USA) employing an Al Kα radiation source was utilized to characterize the elemental chemical states. The pass energy for the full spectra and Al–2p high–resolution spectra in the XPS measurements was set to 100 eV and 30 eV, respectively. The studied samples were subjected to XPS testing directly without any prior surface ion-etching.

3.4. Hydrogen Storage Property Measurements

The hydrogen desorption characteristics of the samples were evaluated with a meticulously calibrated Siever-type instrument. Experiments were conducted in a non-isothermal mode, wherein approximately 50–60 mg of the sample was loaded into a stainless-steel reactor and heated to a specified temperature at a rate of 2 °C/min under a vacuum of 10−2 Pa. The hydrogen absorption properties were measured isothermally. The dehydrogenated samples were heated to 100 °C under vacuum, and then 10 MPa H2 was rapidly charged to start the hydrogenation. Hydrogen levels at the analyzer outlet were continuously monitored using a quadrupole mass spectrometer (Hiden HPR–20EGA, Cheshire, UK).

4. Conclusions

Through the high cooling rate of melting spinning, the grain size of 5% Ti containing Li–Al–Ti alloy was refined, and the LAT–4000 alloy flakes with a dispersed Al3Ti catalyst phase were obtained. Further, by cryomilling for 1 h, the grain/particle size of LAT–4000–CM was reduced to ~17.4 nm/2.5 μm, ~0.4% microstrain was introduced, and the Ti element uniform distribution was further promoted. Owing to the enhancement of hydrogenation kinetics by the uniformly dispersed Al3Ti catalyst phase and the regulation of hydrogenation reaction thermodynamics and kinetics by nanocrystalline and metastable defect-rich structure, LiAlH4 could be directly formed by the high-activity LAT–4000–CM alloy powder via reactive milling under 20 bar hydrogen. The composite hydride containing nano-LiAlH4 began hydrogen desorption at room temperature (25 °C) and the dehydrogenation capacity at 300 °C was determined to be 0.79 wt%. Furthermore, the locally generated nano-LiAlH4 exhibited a certain reversibility due to the LiAlH4 that was also obtained after rehydrogenation of the dehydrogenation product. Although the synthesized LiAlH4 content was limited and the hydrogen storage capacity remained low, this finding provides a novel approach for the non-solvent synthesis of metastable LiAlH4 at room temperature. Building on this, strategies such as multi-component alloying and interface modification could potentially further enhance the synthetic yield of LiAlH4. Furthermore, by combining this method with liquid-phase synthesis assisted by organic solvents and replacing LiH/AlCl3 with Li–Al–Ti alloy, the challenges associated with the removal of inert byproducts such as LiCl and NaCl might be mitigated, thereby optimizing the existing synthesis strategy for LiAlH4.

Author Contributions

Conceptualization, Y.C. (Yan Chu) and Y.Y.; methodology, S.F.; validation, Y.C. (Yingjue Chen) and X.Z.; investigation, Y.C. (Yingjue Chen) and J.Z.; formal analysis, Z.L. and W.D.; writing—original draft preparation, Y.C. (Yan Chu) and S.F.; writing—review and editing, Y.Y., W.C., and Y.L.; visualization, W.D.; resources, J.M. and H.P.; supervision, M.G.; project administration, J.M.; funding acquisition, Y.Y., Z.L., and H.P. All authors have read and agreed to the published version of the manuscript.

Funding

This research was funded by the National Key Research and Development Program of China (2021YFB4000602), the National Natural Science Foundation of China (52271227, 52307250), the Natural Science Foundation of Zhejiang Province, PR China (LZ23E010002), the Two-chain integration key project of Shaanxi Province (2021LLRH-09), and the Key Research and Development Program of Shaanxi (S2024-YF-ZDXM-CY-0359, 2024GX-ZDCYL-04-06).

Data Availability Statement

The original contributions presented in this study are included in the article. Further inquiries can be directed to the corresponding authors.

Conflicts of Interest

The authors declare no conflicts of interest.

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Figure 1. (a) Schematic diagram of the preparation process of active LAT alloys; optical photographs of (b) as-cast alloy by induction smelting, (c) alloy flakes by melt spinning, and (d) alloy powder after cryomilling.
Figure 1. (a) Schematic diagram of the preparation process of active LAT alloys; optical photographs of (b) as-cast alloy by induction smelting, (c) alloy flakes by melt spinning, and (d) alloy powder after cryomilling.
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Figure 2. XRD patterns of LAT–b, LAT–2000, and LAT–4000 alloy flakes.
Figure 2. XRD patterns of LAT–b, LAT–2000, and LAT–4000 alloy flakes.
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Figure 3. SEM of (a) LAT-b, (b) LAT-2000, and (c) LAT-4000; corresponding (d–f) EDX mapping (g–i) corresponding BSE images.
Figure 3. SEM of (a) LAT-b, (b) LAT-2000, and (c) LAT-4000; corresponding (d–f) EDX mapping (g–i) corresponding BSE images.
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Figure 4. (a) XRD patterns and (b) corresponding crystallite size and strain of LAT–b–CM, LAT–2000–CM, and LAT–4000–CM; SEM images of (c) LAT–b–CM, (d) LAT–2000–CM, and (e) LAT–4000–CM, respectively; (f) TEM and corresponding EDX mapping of LAT–4000–CM.
Figure 4. (a) XRD patterns and (b) corresponding crystallite size and strain of LAT–b–CM, LAT–2000–CM, and LAT–4000–CM; SEM images of (c) LAT–b–CM, (d) LAT–2000–CM, and (e) LAT–4000–CM, respectively; (f) TEM and corresponding EDX mapping of LAT–4000–CM.
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Figure 5. (a) XRD patterns and (b) Non-isothermal dehydrogen curves of LAT–b–CM–H, LAT–2000–CM–H, LAT–4000–CM–H samples; (c) hydrogen mass spectrum of LAT–4000–CM–H sample; (d) survey XPS spectra, and (e) Al 2p XPS spectra of LAT–4000–CM–H and as-received LiAlH4.
Figure 5. (a) XRD patterns and (b) Non-isothermal dehydrogen curves of LAT–b–CM–H, LAT–2000–CM–H, LAT–4000–CM–H samples; (c) hydrogen mass spectrum of LAT–4000–CM–H sample; (d) survey XPS spectra, and (e) Al 2p XPS spectra of LAT–4000–CM–H and as-received LiAlH4.
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Figure 6. (a) XRD patterns of LAT–4000–CM and LA–b alloy, (b) XRD patterns, (c) non-isothermal dehydrogen curves, (d) survey XPS spectra, and (e) Al 2p XPS spectra of LAT–4000–CM–H and LA–b–H samples.
Figure 6. (a) XRD patterns of LAT–4000–CM and LA–b alloy, (b) XRD patterns, (c) non-isothermal dehydrogen curves, (d) survey XPS spectra, and (e) Al 2p XPS spectra of LAT–4000–CM–H and LA–b–H samples.
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Figure 7. (a) Non-isothermal dehydrogen curves, (b) survey XPS spectra, (c) Al 2p XPS spectrum, and (d) Li 1s XPS spectrum of rehydrogenated LAT–4000–CM–H.
Figure 7. (a) Non-isothermal dehydrogen curves, (b) survey XPS spectra, (c) Al 2p XPS spectrum, and (d) Li 1s XPS spectrum of rehydrogenated LAT–4000–CM–H.
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MDPI and ACS Style

Chu, Y.; Fang, S.; Chen, Y.; Zhang, X.; Zheng, J.; Li, Z.; Du, W.; Cui, W.; Miao, J.; Yang, Y.; et al. Direct Synthesis of LiAlH4 from Ti-Doped Active LiAl Alloy. Inorganics 2025, 13, 74. https://doi.org/10.3390/inorganics13030074

AMA Style

Chu Y, Fang S, Chen Y, Zhang X, Zheng J, Li Z, Du W, Cui W, Miao J, Yang Y, et al. Direct Synthesis of LiAlH4 from Ti-Doped Active LiAl Alloy. Inorganics. 2025; 13(3):74. https://doi.org/10.3390/inorganics13030074

Chicago/Turabian Style

Chu, Yan, Shiwei Fang, Yingjue Chen, Xiaoqi Zhang, Jie Zheng, Zhenglong Li, Wubin Du, Wengang Cui, Jian Miao, Yaxiong Yang, and et al. 2025. "Direct Synthesis of LiAlH4 from Ti-Doped Active LiAl Alloy" Inorganics 13, no. 3: 74. https://doi.org/10.3390/inorganics13030074

APA Style

Chu, Y., Fang, S., Chen, Y., Zhang, X., Zheng, J., Li, Z., Du, W., Cui, W., Miao, J., Yang, Y., Liu, Y., Gao, M., & Pan, H. (2025). Direct Synthesis of LiAlH4 from Ti-Doped Active LiAl Alloy. Inorganics, 13(3), 74. https://doi.org/10.3390/inorganics13030074

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