X-ray diffraction (XRD, Mini Flex 600) was performed with Cu-Kα radiation (40 kV, 15 mA) at a scan rate of 5° min⁻¹ in the range of 5-80°. The samples were prepared in an Ar atmosphere glovebox and sealed in a custom-designed holder covered by Scotch tape to avoid possible oxidation during the XRD tests. The structure and morphology were studied by transmission electron microscopy (TEM, FEI Talos F200X G2) and scanning electron microscopy (SEM, MIRA3 LHM) equipped with an energy-dispersive X-ray spectrometer (EDS). For TEM observations, the preparation of samples was in the glovebox. The samples were dispersed in cyclohexane, sonicated, dropped cast on a copper grid and rapidly transferred to the equipment. The preparation conditions of samples for SEM analyses were similar to the TEM observations. The only difference was that the samples were dispersed in cyclohexane and dropped to the silicon wafer. X-ray photoelectron spectroscopy (XPS, Kratos AXIS Ultra DLD) was conducted to analyze the valence state and chemical bonding nature of constituent elements of MgH₂-H-V₂O₅ composites. For XPS analyses, a special air-proof transfer vessel was used to transfer the samples from glovebox to the equipment. Thermal gravity analysis (TGA) (PerkinElmer DSC 8000) was performed in air with a heating rate of 10 °C min⁻¹ from room temperature to 700 °C to quantify the content of oxygen vacancies and the entire testing process takes 67.5 min. Raman spectra of samples were recorded with a confocal Raman microscope (Renishaw inVia Qontor, UK). The Brunner-Emmet-Teller (BET) surface area of samples was measured with N₂ at 77 K using an Autosorb-IQ3 apparatus.

Hydrogen sorption behaviors of composites were carried out using a commercialized sievert’s apparatus (HPSA-auto device, China). The temperature-programmed desorption (TPD) curves were obtained for studying hydrogen evolution under vacuum at a heating rate of 2 °C min⁻¹ up to 350 °C. The isothermal hydrogenation with an initial H₂ pressure of 3 MPa and isothermal dehydrogenation under 0.01 MPa initial hydrogen back-pressure were carried out by rapidly heating the sample to preset temperature and stabilizing at this temperature during the entire test. As for cycling test, the desorption behavior was performed under 0.01 MPa initial hydrogen back-pressure at 275 °C in 10 min, and absorption behavior was performed under 3 MPa initial H₂ pressure at 275 °C in 10 min. The thermodynamic properties of the composites are determined by the pressure-composition-isothermal curve (PCI). To analyze the desorption kinetic properties, the composites were heated from room temperature to 500 °C with the heating rates of 3, 5, 7, and 10 °C min⁻¹ under a flowing Ar atmosphere using differential scanning calorimetry (DSC, NETZSCH STA 449 F3) measurements (at different heating rates, the test duration is 158.3 min (3 °C min⁻¹), 95 min (5 °C min⁻¹), 67.8 min (7 °C min⁻¹), and 47.5 min (10 °C min⁻¹), respectively), and the obtained peak temperatures were fitted using the Kissinger equation to calculate the apparent activation energy of the dehydrogenation reaction.

Density functional theory (DFT) calculations were carried out in Vienna ab initio simulation package (VASP) [36,37]. Perdew-Burke-Ernzerhof (PBE) functionals and generalized gradient approximation (GGA) form were used to describe the exchange-correlation energy [38]. The projected augmented wave (PAW) potentials [39,40] were chosen to describe the ionic cores, and valence electrons were taken into account using a plane wave basis set with a kinetic energy cutoff of 520 eV. The GGA + U method was adopted in our calculations. The value of the effective Hubbard U was set as 4.72 for V. Partial occupancies of the Kohn-Sham orbitals were allowed using the Gaussian smearing method with a width of 0.05 eV. The structures were relaxed until the forces and total energy on all atoms converged to less than 0.05 eV Å⁻¹ and 1 × 10^-5 eV. The Brillouin zone integration is performed using 2 × 2 × 1 Monkhorst-Pack k-point sampling for a structure.

As schematically illustrated in Fig. 1a, H-V₂O₅ nanosheets were synthesized by a solvothermal reaction of NH₄VO₃ in a dilute HCl solution followed by thermal treatment under a hydrogen atmosphere. Figure S1 shows the digital photos of V₂O₅ nanosheets at different states. SEM images (Fig. 1b, c) validate that the as-synthesized V₂O₅·xH₂O nanosheets obtained by the solvothermal reaction exhibited the interpenetrating 2D nanosheets forming a 3D nanostructure. The V₂O₅·xH₂O nanosheets were then annealed in air to remove water and obtain V₂O₅-350air nanosheets. After hydrogenation via thermal treatment under a hydrogen atmosphere at 300 °C (H-V₂O₅), the morphology of the interpenetrating H-V₂O₅ nanosheets was perfectly retained as shown in Fig. 1d, e. Impressively, the H-V₂O₅ nanosheets obtained by hydrogenation of V₂O₅-350air nanosheets at different temperatures exhibited different colors ranging from brown to black, in marked contrast with the orange V₂O₅-350air powders (Fig. S1). The color change of V₂O₅ is ascribed to the presence of oxygen vacancies, which coincides well with the previous work [41].

TEM image of pristine V₂O₅·xH₂O nanosheets reveals an intertwined 2D nanosheet structure with a smooth surface (Fig. 1f). High-resolution TEM (HRTEM) image exhibits the lattice fringes of 0.408 nm (Fig. 1f), corresponding to the (101) plane of V₂O₅, which is in good agreement with XRD and XPS results (Fig. 2). Energy-dispersive spectroscopy (EDS) elemental mapping results demonstrate the uniform distribution of V, O, and N elements in the as-synthesized V₂O₅·xH₂O nanosheets (Figs. 1g and S2). The electron-donating ability of N atoms could enhance the charge transfer, facilitating the decomposition of hydrogen molecule, and destabilization of the Mg-H bond [42,43]. However, due to the small content of N, the effect of N element on the improvement of properties is not emphasized in this work. The TEM image of H-V₂O₅ nanosheets (Fig. 1h) reveals a similar morphology as the pristine V₂O₅·xH₂O nanosheets. The thin 2D nanosheets provide a larger contact area with MgH₂, thereby shortening the path enabling fast H diffusion and electron transport during the de/re-hydrogenation, giving rise to an improved hydrogen storage performance of MgH₂.

XRD patterns of the hydrothermal product are depicted in Fig. S3, revealing peaks corresponding to the (001), (003), (005), and (006) planes of the layered V₂O₅·xH₂O [44]. The V₂O₅-350air was obtained after annealing in air at 350 °C to remove absorbed water, and the diffraction peaks can be indexed to the orthorhombic phase V₂O₅ (JCPDS No. 41-1426) (Fig. S3). Furthermore, Fig. 2a shows the XRD patterns of H-V₂O₅ annealed at different temperatures to investigate the effects of hydrogenation temperature on the crystal structure of V₂O₅. It can be seen that H-V₂O₅-200 shows the same orthorhombic phase as V₂O₅-350air. However, when the hydrogenation temperature is increased to 300 °C, the characteristic diffraction peaks of the monoclinic VO₂ phase (JCPDS No. 09-0142) can be detected, indicating that oxygen is partially removed from V₂O₅ at higher temperatures, which coincides well with the color change of digital photos of V₂O₅ nanosheets at different states (Fig. S1). The enlarged image in Fig. 2a reveals that the (001) diffraction peak of the H-V₂O₅-200 shifts to a higher angle, suggesting the decrease in the d space and the presence of oxygen vacancies in the crystal [45,46,47]. The specific surface areas of V₂O₅ nanosheets were further investigated (Fig. 2b). Compared with commercial V₂O₅, the specific surface area of V₂O₅·xH₂O nanosheets was increased by 13 times, attributed to the special intertwined 2D nanosheet structure.

Figure 2c shows the XPS spectrum of V 2p in V₂O₅ nanosheets, which can be well split into two peaks of V at 517.1 and 524.6 eV, ascribing to the V 2p_3/2 and 2p_1/2 of V⁵⁺, respectively. Compared to the V₂O₅·xH₂O and V₂O₅-350air, the slight shift of the V⁵⁺ peak of H-V₂O₅-300 was observed, which could be attributed to the existence of oxygen vacancy. Moreover, we performed TG tests (Fig. 2d) on the hydrogenated H-V₂O₅ nanosheets and found that there was a significant weight gain when samples were heated in an oxygen atmosphere. Assuming the chemical formula of H-V₂O₅-300 is V₂O_5-x, the value of x was calculated as 1.13 according to the TG results, corroborating the presence of oxygen vacancies in the H-V₂O₅-300 nanosheets. Raman spectra (Fig. 2e) were obtained to investigate the oxygen vacancy sites due to different oxygen sites in the crystal having different bond lengths with V. The peaks at 993, 696, and 526 cm⁻¹ were assigned to the stretching vibration of V = O(I), V-O(III), and V-O(IV) bonds, respectively [48,49]. The typical Raman peaks at 480 and 303 cm⁻¹ were attributed to the bending vibrations of V-O(II) and V-O(IV) bonds, whereas those at 406 and 283 cm⁻¹ correspond to the bending vibration of the V = O(I) bonds. The two Raman bands at 194 and 141 cm⁻¹ were assigned to [VO₅]-[VO₅] vibrations [50,51,52]. After hydrogenation at 300 °C, the peak at 480 cm⁻¹ stemming from the bending vibration of the V-O(II) bridge was almost absent, revealing that the oxygen vacancy site in H-V₂O₅ nanosheets is mainly at the bridging O(II) sites. Notably, hydrogenation treatment of V₂O₅ nanosheets can kill two birds with one stone, which not only can introduce oxygen vacancies into the oxide, but also is equivalent to a pre-treatment of the catalyst, ensuring the stability of V₂O₅ in the subsequent hydrogen ab/de-sorption test.

The catalytic effect of V₂O₅ nanosheets in improving hydrogen storage performances of MgH₂ is evaluated by mixing V₂O₅ nanosheets with MgH₂ through mechanical ball-milling. The temperature-programed desorption of H-V₂O₅-doped samples and pristine MgH₂-BM were first tested. As shown in Fig. 3a, the pristine MgH₂-BM starts releasing H₂ at approximately 256 °C with a high terminal temperature at 350 °C, delivering a hydrogen desorption capacity of about 7.51 wt%, which agrees well with its theoretical hydrogen capacity. The results showed that the hydrogen release rate of the H-V₂O₅-doped samples was significantly faster than that of the pristine MgH₂-BM samples, especially at the beginning of the hydrogen desorption. The MgH₂-10 wt% H-V₂O₅ started to release H₂ at about 185 °C, with a hydrogen capacity of 6.54 wt%. Notably, the obvious decrease in the onset and terminal desorption temperatures of H-V₂O₅-doped MgH₂ compared to those of pristine MgH₂-BM indicates smaller particle size induced by the encapsulation effect from sheet-like structure of H-V₂O₅ and uniform distribution between H-V₂O₅ nanosheets and MgH₂ during ball milling process, which could be verified by the microtopography and elemental mapping results, hence promoting the catalytic effect of H-V₂O₅ nanosheets in improving hydrogen desorption performance of MgH₂. To further illustrate the superior catalytic activity of H-V₂O₅ nanosheets with oxygen vacancies towards MgH₂, the DSC curves of V₂O₅·xH₂O-doped MgH₂, H-V₂O₅-doped MgH₂, and pristine MgH₂ with different heating rates were compared (Figs. 3b, c, S4 and S5). The peak desorption temperatures of the H-V₂O₅-doped MgH₂ decreased by about 36.7 and 101.2 °C compared to that of the V₂O₅·xH₂O-doped MgH₂ and pristine MgH₂, respectively, demonstrating the superior catalytic effect of H-V₂O₅ nanosheets composed of both V⁵⁺, V⁴⁺ and oxygen vacancies than V₂O₅·xH₂O nanosheets in improving the hydrogen desorption properties of MgH₂. According to the Kissinger's method [53,54], the apparent activation energies (E_a) of MgH₂-10 wt% H-V₂O₅ (Fig. S4) and pristine MgH₂-BM (Fig. S5) were obtained. The results showed that the E_a of MgH₂-10 wt% H-V₂O₅ and pristine MgH₂-BM were 113.8 and 137.4 kJ mol⁻¹, respectively. Thus, the addition of H-V₂O₅ nanosheets can significantly reduce the apparent activation energy of dehydrogenation for MgH₂ and improve the kinetics of the composites.

In order to comprehensively evaluate the hydrogen storage performance of MgH₂-H-V₂O₅ composites, isothermal hydrogen ab/de-sorption kinetics of MgH₂ under the catalysis of H-V₂O₅ nanosheets are subsequently investigated (Fig. 3d, e). In contrast, the kinetic properties of pristine MgH₂-BM were also tested (Fig. S6). According to the results of the isothermal hydrogenation, it can be seen that MgH₂-H-V₂O₅ composites exhibit excellent hydrogen absorption capability at near room temperature. Under the conditions of 30 °C and 3 MPa hydrogen pressure, the dehydrogenated MgH₂-H-V₂O₅ composites could absorb 2.38 wt% in 60 min, which meet the requirement of the practical application of hydrogen storage systems. In addition to the significant hydrogen absorption properties at room temperature, the composites show superior hydrogen storage kinetics by increasing the operating temperature. As depicted in Fig. 3d, the MgH₂-H-V₂O₅ composites could absorb 4.1 wt% H₂ within 30 min at 100 °C, while pristine MgH₂-BM showed no obvious hydrogen absorption behavior under the same conditions. Additionally, the hydrogen absorption of 4.16 wt% at 200 °C and 5.08 wt% H₂ at 250 °C only takes 60 s, while it takes 11 min for pristine MgH₂-BM to absorb the same amount of H₂ at 200 °C (Fig. S6).

The isothermal dehydrogenation performances of MgH₂-H-V₂O₅ composites were also measured, and the results are shown in Fig. 3e. It is shown that only 5 min is required for the complete dehydrogenation of MgH₂-H-V₂O₅ with a hydrogen desorption capacity of 6.0 wt% at 275 °C. However, pristine MgH₂-BM can only release 0.2 wt% H₂ in 5 min and 4.2 wt% H₂ in 60 min at 275 °C, which exhibited more sluggish kinetics than MgH₂-H-V₂O₅ composites. According to the slope of the desorption curve at 275 °C, the H₂ desorption rate of MgH₂-H-V₂O₅ was calculated to be 2.3 wt% min⁻¹, which is much faster than that of pristine MgH₂-BM. At the temperature of 225 °C, 4.42 wt% H₂ could be released from MgH₂-H-V₂O₅ within 10 min, which is equivalent to 68% of the theoretical hydrogen storage capacity. More impressively, even at the lower temperature of 200 °C, about 2.37 wt% H₂ could be desorbed from MgH₂ under the catalysis of H-V₂O₅ nanosheets within 60 min.

It is worth noting that different activation energies will be obtained when different kinetic models are used for fitting. Therefore, to obtain more accurate activation energy values, an in-depth analysis of de/re-hydrogenation kinetic mechanism model of composites was carried out according to nine different kinetic models proposed by Sharp and Jones [55], and the activation energy was fitted according to a suitable kinetic model. The specific methods are provided bellow:

The general dynamic equation is shown as follows:

where α, T, k(T), g(α), A, and t_0.5 represent the reaction extent, the reaction temperature, the reaction rate constant, the function depending on the specific kinetic mechanism, the constant related to the kinetic mechanism, and the time when α equals 0.5 respectively. Through plotting the dehydrogenation experimental data of t/t_0.5 against the theoretical ones of composites for nine different kinetic models, respectively, the corresponding reliable kinetic model can be obtained (the line with a slope close to 1 is the reliable kinetic model).

As shown in Fig. 4b, we obtained nine curves based on the isothermal dehydrogenation curves of composites at 225 °C. In our case, the dehydrogenation kinetic model should be R3 model, which has the slope value of 0.97493 close to 1. The R3 model is further verified through plotting the g(α) related to R3 against the reaction time at 250 and 275 °C when α ranges from 0.2 to 0.7, which exhibit strong linear relationship (R² more than 0.99) as shown in Fig. 4c. Moreover, the apparent activation energy (E_a) was calculated to evaluate the energy barriers for the dehydrogenation process of MgH₂-10 wt% H-V₂O₅ according to the following Arrhenius equation:

where R and k represent gas and rate constant, respectively. And k is the slope value that can be extracted from Fig. 4c. Through linear fitting the lnk and 1000/T [53,56], the E_a (84.55 ± 1.37 kJ mol⁻¹ H₂) for dehydrogenation was obtained as shown in Fig. 4d. Similarly, we have re-fitted the kinetic model of the hydrogen absorption process as well as the hydrogen absorption activation energy barrier according to the same method, and the results are shown in Fig. S7. According to the calculation results, we get the precise hydrogen absorption activation energy E_a (42.55 ± 0.41 kJ mol⁻¹ H₂). The E_a values for ab/de-sorption of MgH₂-H-V₂O₅ were significantly lower than those of the commercial MgH₂ (71 kJ mol⁻¹ H₂ for absorption and 160 kJ mol⁻¹ H₂ for desorption) [9,54]. Overall, the introduction of H-V₂O₅ nanosheets greatly improved the kinetic performances of MgH₂.

The thermodynamic properties of the MgH₂-H-V₂O₅ composites and pristine MgH₂-BM were determined by the PCI curve (Fig. S8). According to the van’t Hoff plots, the de/re-hydrogenation enthalpy (ΔH) of the composites were determined to be 73.4 and − 73.6 kJ mol⁻¹ H₂ (Fig. S8a, b), respectively, which were comparable to the ΔH of pristine MgH₂-BM (Fig. S8c, d). It demonstrates that the introduction of H-V₂O₅ nanosheets exhibits no effect in tuning the thermodynamics of MgH₂, confirming that the modification of hydrogen storage properties of MgH₂ by H-V₂O₅ nanosheets is mainly a catalytic effect. Notably, compared to pristine MgH₂-BM, the platform pressure of MgH₂-H-V₂O₅ composites obviously increased, which was closer to the requirements of commercial applications.

The hydrogen ab/de-sorption cycling curves of the MgH₂-H-V₂O₅ sample at 275 °C are shown in Figs. 5 and S9. The composites could achieve the reversible de/re-hydrogenation with the hydrogen capacity of about 5.7 wt% during the first cycle. As the number of cycles increased, the hydrogenation capacity of the composites decreased slightly compared with that of the first cycle, which may be due to the formation of a surface passivation oxide layer on the MgH₂/Mg particles and incomplete activation during the first few cycles, limiting the hydrogen uptake/release capacity. Additionally, the H-V₂O₅ catalyst may undergo a reduction process during the initial cycles, which may contribute to the decrease in capacity. Moreover, the metal V generated during the ball milling process could absorb hydrogen to form VH₂ species, which was proved by the presence of the V²⁺ in the XPS results of dehydrogenated samples (Fig. 6b). Due to the fact that V-H bonds are very stable, VH₂ is difficult to be completely dehydrogenated, resulting in a decrease in the hydrogen storage capacity. However, after 10 cycles, the capacity attenuation tended to be stable. More impressively, as the hydrogen ab/de-sorption cycle continued, the capacity of the composites started to recover and slowly increased, gradually returning to the capacity of the first hydrogen absorption cycle. As the cycling continued, the surface oxide layer can gradually be broken down, creating a highly homogeneous and intimate contact between MgH₂ and H-V₂O₅ and forming a stable structure. Additionally, the presence of oxygen vacancies may inspire the decomposition of VH₂ and promote the transformation between V(H) and VH₂, thus, contributing a portion of the hydrogen storage capacity, which is proved later in the results of the theoretical calculations. The H-V₂O₅ catalyst may also undergo a gradual activation process, restoring its catalytic activity and enhancing the hydrogen storage performance of MgH₂/Mg. Furthermore, the cycling-induced changes in the electronic structure of the H-V₂O₅ catalyst introduced the multi-valance V into composites, which can also contribute to the improved kinetics and stable cyclic behavior. After even 100 cycles, the hydrogen capacity of the composite was still maintained at about 5.7 wt%, which was remarkably better than those of the state-of-the-art V₂O₅-catalyzed MgH₂ systems. Regarding the amount of hydrogen absorption in the first cycle as a benchmark, the composites could maintain a higher capacity retention rate of 99% after 100 cycles. In contrast, the capacity of pristine MgH₂-BM without catalysts decreased significantly during the de/re-hydrogenation cycles [25,57].

More interestingly, a decline of the onset temperature from 185 to 175 °C could be clearly observed for MgH₂ under the catalysis of H-V₂O₅ nanosheets since the 30th cycle of hydrogen desorption process (Fig. 5b), indicating the enhancement of the catalytic effect of H-V₂O₅ nanosheets upon cycling process. Even after 100 cycles, the initial hydrogen desorption temperature and hydrogen capacity of MgH₂-H-V₂O₅ composites showed no obvious attenuation. The re-hydrogenated composites could desorb hydrogen at lower temperature ranges than that of the ball-milled samples, indicating that the chemical states of H-V₂O₅ nanosheets might slightly change and the catalytic activity was significantly improved during the de/re-hydrogenation cycling behavior. By comparison, the onset dehydrogenation temperature of pristine MgH₂-BM increased significantly from 256 to 290 °C after 20 cycles, showing a significant decay of performance (Fig. S10). This indicates that the addition of H-V₂O₅ nanosheets can not only provide catalytic activity for MgH₂, but also inhibit the serve agglomeration and growth of MgH₂ particles from the unique sheets-like structure of H-V₂O₅ nanosheets, thereby improving the cycling stability of the MgH₂-H-V₂O₅ composites. Notably, as the test temperature increases, the large MgH₂ particles continue to release hydrogen at high temperatures, and the maximum hydrogen storage capacity of the MgH₂-BM does not decay (Fig. S10). In contrast, the hydrogen storage capacity of the composite decreases after several hydrogen ab/de-sorption cycles, and this capacity cannot be recovered with the increase of test temperature (Fig. 5b). This further proves that the attenuation of capacity is not caused by the particle growth, but by the formation of the new catalytic phase.

To elucidate the catalytic mechanism of the H-V₂O₅ nanosheets for MgH₂, a series of characterizations were carefully performed to study the corresponding structures and compositions of the MgH₂-H-V₂O₅ at different states. The XRD results of MgH₂-H-V₂O₅ before and after de/re-hydrogenation are shown together with the pristine MgH₂ after ball milling in Figs. 6 and S11. The XRD patterns show that the β-MgH₂ with a little γ-MgH₂ is the main phases in the pristine MgH₂ after ball milling (Fig. S11), indicating the high purity of the MgH₂ powder. In addition to the peak of MgH₂, a small amount of V phase and MgO can be detected in the XRD patterns of the ball-milled MgH₂-H-V₂O₅ composites (Fig. 6a) and these two phases were difficult to distinguish due to their overlapping Bragg reflections, demonstrating the partial reduction of H-V₂O₅ by MgH₂ during the ball milling process (5MgH₂ + V₂O₅ = 5MgO + 2V + 5H₂, ΔG = − 1246 kJ mol⁻¹ V₂O₅). The absence of MgO phase in pristine MgH₂-BM further indicates that H-V₂O₅ and MgH₂ undergo a redox reaction during the ball milling process. This reaction in-situ generates fine, dispersed, and catalytically active metal V along with inert MgO, consuming a small amount of Mg, thus, providing evidence for the decay of the capacity in the initial cycling test in the following section. Additionally, previous studies have shown that metal V could serve as hydrogen spillover active sites to promote the dissociation of hydrogen molecules [24,29,58]. After dehydrogenation, the peaks from MgH₂ completely disappeared, and the XRD characteristic peaks of Mg and V appeared. Notably, Mg phase was converted to MgH₂ accompanied by the appearance of VH₂ in the subsequent hydrogen absorption process. Based on the above results, we can deduce that part of H-V₂O₅ is reduced to the V phase, which can then transform to VH₂ after re-hydrogenation. The reversible chemical reaction between V/V(H) and VH₂ plays a crucial role in the catalytic effect of H-V₂O₅ on MgH₂ [29,59], which is evidenced in the XPS and HRTEM analysis. The XRD patterns of the MgH₂-H-V₂O₅ composites after 100 re/de-hydrogenation cycles were also depicted in Figs. S12 and S13. After 100 cycles, the main characteristic diffraction peaks of MgH₂/Mg with the catalysis of H-V₂O₅ nanosheets could be indexed to MgH₂/Mg, and V/VH₂ phase can also be detected, which provides additional evidence for the high reversibility of MgH₂ and stable catalytic activity of V/VH₂.

XPS analysis and HRTEM measurement were further conducted to analyze the state changes of V in the MgH₂-H-V₂O₅ composites. After the ball milling process, the spin-orbit double peaks at 512.08/517.6 eV, 513.38/518.93 eV, and 514.98/521.28 eV in the V 2p XPS spectra (Fig. 6b) could be indexed to V⁰, V²⁺, and V⁴⁺, respectively, indicating that part of the high-valence V⁵⁺ ions were reduced to metallic V⁰. The existence of V⁴⁺ evidenced the presence of oxygen vacancies. It is worth emphasizing that the absence of the signal of V⁵⁺ may be due to the fact that the detection depth of XPS is only a few nanometers below the surface, while the H-V₂O₅ nanosheets may be covered by MgH₂ or partially reduced metallic V nanoparticles in the ball milling process and undetectable by the XPS measurement. After the subsequent activation process of hydrogen ab/de-sorption, the H-V₂O₅ nanosheets were gradually exposed to the surface, which was further proved by the detection of the signal of V⁵⁺ in the XPS spectra of the de/re-hydrogenated samples. The combined analysis of XPS and XRD results showed that a phase transition between VH₂ and V/V(H) occurred during the hydrogen ab/de-sorption process. Notably, the presence of V²⁺ in the dehydrogenated samples indicated that VH₂ was not completely dehydrogenated, which further provided evidence for the capacity decay at the beginning of the cycle. Moreover, compared to the ball-milled samples, the slightly shift of V⁰ peak of the dehydrogenated samples was observed due to the incomplete dehydrogenation of VH₂ and the existence of V(H). Impressively, after 100 cycles, the absence of V²⁺ in the dehydrogenated sample indicated that VH₂ was activated as an active catalyst and completely dehydrogenated to produce V/V(H), attributed to the pre-activation process and the assistance of oxygen vacancies.

The results further indicate that in-situ formed V could facilitate the spontaneous H₂ dissociation, stemming from the strong hybridization between the molecular orbital of adsorbed hydrogen and the V 3d states [60], which could promote hydrogen spillover from the V/H-V₂O₅ multi-phase catalysts to the Mg surface. Moreover, computational study [61] showed that the formation energy of the VH₂/MgH₂ interface had a low absolute value of the negative heat compared with that of MgH₂, demonstrating that VH₂/V enhanced the desorption properties of MgH₂. Hence, it could be concluded that the hydrogen spillover effect originated from V/H-V₂O₅ catalysts and rapid electron transfer between multi-valent V states facilitate the hydrogen ab/de-sorption process of MgH₂. Furthermore, the high-resolution O 1 s XPS spectra (Fig. S14) demonstrated the presence of oxygen vacancy and it remains stable in the process of re/de-hydrogenation. Notably, it is deemed that the bonding energy between V and H atoms is strong. Hence, we deduced that oxygen vacancy might stimulate the phase transition between VH₂ and V, which was further proved by theoretical calculation in the next section.

Furthermore, the morphologies of MgH₂-H-V₂O₅ in the three different states were presented to clarify the evolution of the microstructure of the composites (Fig. 7). The dark regions with lamellar distribution in the BF-TEM images could be identified as the H-V₂O₅ nanosheets, and the well-distributed dark regions in the form of dots indicate V/V(H)/VH₂ nanoparticles from the reduction of H-V₂O₅ nanosheets during ball milling process. It can be seen that the sheet-like structure of H-V₂O₅ still maintained well after re/de-hydrogenation, encapsulating around the surface of MgH₂ nanoparticles, which inhibited the agglomeration and growth of MgH₂ to a certain extent and ensured its excellent cycling stability. It is worth noting that the morphology of MgH₂ nanoparticles remained stable after the hydrogen ab/de-sorption cycling tests (Fig. S15). However, the pristine MgH₂-BM sample without catalyst doping exhibited obvious agglomeration after re/de-hydrogenation cycles (Fig. S16), which coincided well with the poorly tested cyclic stability of pristine MgH₂. Additionally, the elemental mapping of MgH₂-H-V₂O₅ in different states illustrated that Mg, V, and O elements were homogeneously distributed in the composites.

From the HRTEM characterizations of ball-milled composites, it was observed that the spacing values of 0.2515 and 0.2470 nm matched well with those of MgH₂ (101) and VO₂ (211) (Fig. 7d) in MgH₂-H-V₂O₅, respectively. The lattice spacing of 0.2131 nm could be indexed to the (101) planes of metallic V. In the dehydrogenation states, the spacing value of 0.2455 nm corresponded to the Mg (101) and the metallic V distributing around the Mg phase could still be detected. (Fig. 7e). For the rehydrogenated samples, the spacing values of 0.2502 and 0.2474 nm were indexed to MgH₂ (101) and VH₂ (111), respectively (Fig. 7f), indicating that V nanoparticles formed by the reduction of H-V₂O₅ could transform to VH₂ during the rehydrogenation process.

To further quantitatively understand the synergistic catalytic effect of V₂O₅ and oxygen vacancies, especially in the VH₂/H-V₂O₅ interface in the hydrogen storage properties of MgH₂, DFT calculations were performed within MgH₂ (110), Mg (100), V₂O₅ (001), and VH₂ (111) surface model. According to the calculation results, introducing V₂O₅ into the MgH₂ exhibited a positive effect on decreasing the de/re-hydrogenation energy barriers of MgH₂/Mg and destabilizing the Mg-H bonds. Moreover, the introduction of oxygen vacancies will further enhance this effect. As shown in Fig. S17 and Table S1, the length of the Mg-H bond of pristine MgH₂ was 1.72 Å. After the addition of V₂O₅, the Mg-H bond extended to 1.79 Å. Notably, obvious elongation of the Mg-H bond to 1.88 Å was obtained by further introducing the oxygen vacancies into the system. Furthermore, the bond energy of the Mg-H was also reduced from 92.39 kJ mol⁻¹ for pure MgH₂ to 62.36 and 40.93 kJ mol⁻¹ after the introduction of V₂O₅ nanosheets and oxygen vacancies, respectively. Increasing of the bond length and decreasing of bond energy are deemed to be the most direct evidence for the weakening of Mg-H bond, demonstrating that it is easier to break the Mg-H bond with lower desorption temperature [62].

Additionally, the energy barrier for the de/re-hydrogenation process of MgH₂/Mg on V₂O₅ (001) and V₂O_5-x (001) are subsequently calculated, as schematically illustrated in Fig. 8. Hypothetically, the hydrogen desorption process of MgH₂ includes two steps: MgH₂ adsorbed on V₂O₅ (001)/V₂O_5-x (001) substrates (initial state, denoted as IS) undergoes the breakage of Mg-H bonds (transition state, TS), and then the H atoms escape from the substrate and recombine into H₂ (final state, FS). As shown in Fig. 8c, d, the free energy for de/re-hydrogenation of pristine MgH₂/Mg was calculated to be 0.91/-0.15 eV. In contrast, obvious decreased energy barrier can be noticed under the catalysis of V₂O₅ (0.60/-0.44 eV) and oxygen vacancies (-0.62/-1.12 eV), which directly implies their superior catalytic activities in reducing the hydrogen sorption temperatures of MgH₂. Specifically, the negative energy barrier value of the MgH₂-V₂O_5-x system demonstrates that the decomposition of MgH₂ on oxygen vacancies-rich V₂O₅ nanosheets is spontaneous, implying that the dehydrogenation of MgH₂ is further facilitated by the accelerated electron transfer due to the introduction of oxygen vacancies and this can be further proved from the DOS curves. Figure 9a shows the DOS and d-orbital bias density of MgH₂-V₂O₅ system and MgH₂-V₂O_5-x system, it was demonstrated that the d-band centers of the MgH₂-V₂O_5-x system was closer to the Fermi energy level compared to the MgH₂-V₂O₅ system, indicating the increase in electron distribution.

Moreover, based on the results of hydrogen storage performance test, we have concluded that the presence of oxygen vacancies may inspire the decomposition of VH₂ and promote the transformation between V(H) and VH₂. In order to better illuminate the inspiration function of oxygen vacancies on the decomposition of VH₂, the length of the V-H bond of pristine VH₂ model, VH₂ on V₂O₅ (001) surface model, and VH₂ on V₂O_5-x (001) surface model were further calculated (Fig. S18). As shown in Table S2, under the catalysis of V₂O₅ and V₂O_5-x, the length of the V-H bond was significantly increased from 1.87 Å (pristine VH₂) to 1.93 and 2.04 Å, respectively, and the bond energy was also reduced from 75.663 kJ mol⁻¹ in pure (pristine VH₂) to 70.691 and 52.613 kJ mol⁻¹, which directly demonstrated the inspiration effect of oxygen vacancies on the decomposition of VH₂. Moreover, the DOS and d-orbital bias density of VH₂-V₂O₅ system and VH₂-V₂O_5-x system (Fig. 9b) further demonstrated that the introduction of oxygen vacancies could increase the electron distribution, thus weakening the V-H bond.

The above calculation results provide a rational explanation for lower de/ab-sorption temperatures and energy barrier in MgH₂-H-V₂O₅ composites. The presence of oxygen vacancies not only effectively promotes the dissociation of MgH₂, but also reduces the dehydrogenation energy barrier of the VH₂ catalyst, further improving its catalytic activity. In other words, the introduction of oxygen vacancies can kill two birds with one stone, which can directly enhance the hydrogen storage performance of the matrix material (MgH₂) and indirectly improve the overall hydrogen storage performance of the system by affecting the catalytic activity of the catalyst (VH₂). These results confirm the synergistic catalytic effect of transition metal oxide and oxygen vacancies in improving the hydrogen storage performance of MgH₂, which coincides well with previous researches [31].

Based on the above discussions, the reaction equations for the hydrogen storage process of MgH₂-H-V₂O₅ composites are as follows:

And the he possible catalytic mechanisms (Fig. 10) of H-V₂O₅ for the re/de-hydrogenation of MgH₂ were proposed. (1) The sheet-like structure of the H-V₂O₅ with abundant oxygen vacancies was beneficial for exposing more active sites to contact with MgH₂/Mg. The numerous interfaces between H-V₂O₅ nanosheets and MgH₂/Mg provide more diffusion paths for hydrogen and nucleation sites for Mg/MgH₂, improving its hydrogen storage performances. (2) During the ball milling process, the cross-linked 3D H-V₂O₅ nanosheets were broken into 2D nanosheets, which were uniformly covered on MgH₂ surface. By virtue of the unique 2D sheet-like structure of H-V₂O₅, the agglomeration and growth of MgH₂ could be inhibited through the encapsulation effect of H-V₂O₅, ensuring its excellent cycling stability. (3) Additionally, H-V₂O₅ could react with MgH₂ during ball milling process, forming V/H-V₂O₅ multi-phase catalysts. Profiting from the “hydrogen spillover” effect of V/VH₂ phase, hydrogen molecules exhibit a low dissociation energy barrier on the surface of metal V, which promotes the dissociation of H₂ [62], facilitating the hydrogen absorption properties. (4) The presence of oxygen vacancies and multi-valent V could accelerate the electron transfer between the catalyst and Mg, thus weakening the Mg-H bond and promoting the dehydrogenation of MgH₂ [33]. (5) Moreover, the oxygen vacancies in H-V₂O₅ nanosheets may inspire the "hydrogen pump" effect of VH₂/V, promoting the dehydrogenation of VH₂, thus further strengthen the catalytic activity of the VH₂/V. Taken together, these results corroborate that the incorporation of hydrogenated V₂O₅ nanosheets with abundant oxygen vacancies leads to the improved hydrogen storage performances of MgH₂.