Solvent-mediated partial ionicity enhances mechanical nanosizing effect of Mg-based hydrogen storage alloys

Domain: chemistry
Published: 2026-01-14
Editorial outcome: minor_revision

Abstract
The high ductility of Mg has posed major challenges for nanofabrication utilizing mechanical ball-milling. While the addition of organic solvents is effective, it has been unclear how they improve the ball-milling effect by modifying the material’s surface properties. Herein, we report that the solvent-mediated partial ionicity plays an important role in enhancing the nanosizing effect of Mg87.5Ni5.5Y7 alloy. This approach enables the Mg87.5Ni5.5Y7 particles to be 88 times smaller than those of the solvent-free procedure. The Mg87.5Ni5.5Y7 nanoparticles underwent complete dehydrogenation in 3 min at 300 °C and in 17 min at 240 °C, which can be stably cycled at least for 500 times. Solvent (THF) adsorption on Mg induces Mgδ+‒Mgδ− dipole structure. This increases the surface hardness of Mg-based alloy and maximizes the ball milling-driven structural deformation, thereby facilitating ion migration. Mg‒Mg bond breaking is caused by the resulting Coulombic repulsion between Mg atoms. These findings provide an affordable approach for nanoparticle fabrication of highly ductile materials.

Introduction
Hydrogen, as a zero-carbon energy carrier, exhibits impressive characteristics of abundance, versatile application, and high energy conversion efficiency. Unfortunately, its widespread application is constrained by critical challenges in hydrogen storage and transportation sectors1. Magnesium (Mg) has emerged as a promising candidate, given its prominent theoretical hydrogen storage capacity (7.6 wt.% H2), environmental friendliness and abundant natural reserves2–10. Despite these advantages, Mg faces limitations due to its sluggish reaction kinetics, which stem from multistep reaction processes, including hydrogen adsorption, dissociation, diffusion and nucleation. These processes necessitate high operation temperature (~350 °C) and substantial energy loss11,12.

To enhance the kinetic processes, a variety of Mg-based alloys, such as Mg-Ni-RE (RE=rare earths), have been developed13–16. They undergo in situ decomposition after the activation process and generate Mg, Mg2Ni, and REHx nanoparticles (NPs). In it, Mg contributes hydrogen storage capacity, while REHx and Mg2Ni serve as catalyst phases for hydrogen adsorption and dissociation, thereby enhancing the rates of hydrogenation and dehydrogenation (de/hydrogenation)17–21. However, the hydrogen diffusion in the Mg phase is hardly accelerated by these catalyst phases, and persistent challenges of coarse particle size distribution and suboptimal morphology control are still unresolved22,23.

Nanostructuring strategies demonstrate transformative potential24,25. Nanoparticles increase their surface energies thanks to the unsaturated coordination number, increased grain boundaries and enriched reaction sites/defects, which accelerate hydrogen adsorption and dissociation26–28. Furthermore, nanoscale dimensions shorten hydrogen diffusion paths between surface and bulk regions, significantly ramping de/hydrogenation rates29. Combining Ni catalysts with nanostructured Mg simultaneously accelerates hydrogen adsorption, dissociation, diffusion and nucleation. Nanofabrication of Mg can drastically enhance de/hydrogenation rates at much milder reaction temperatures4,5,8,29–32. Room-temperature reversible hydrogen storage of Mg can even be achieved at particle sizes below ~5 nm33.

While the current nanofabrication method can improve scalability, it predominantly relies on gas atomization and chemical reduction processes, which requires expensive precursors and complicated synthetic procedures, posing major challenges for industry utilization. In contrast, mechanical ball milling offers a simpler and cost-effective alternative, and a variety of ball-milling studies on both Mg-based compounds, such as MgH2, have been reported34–39. However, unlike the brittle ceramic MgH2, the high ductility of metallic Mg fundamentally hinders its applicability due to severe agglomeration and cold welding40,41. Recent advances in solvent-assisted ball milling show promise. Zou et al.42 successfully synthesized nanoscale Mg via this approach, providing a scalable route for the nanoparticle production of Mg. However, the mechanistic role of solvents in enabling Mg nanosizing is still unclear. In addition, surface instability after ball-milling poses challenges for practical applications. These issues collectively impede the development of milder-temperature operatable Mg-based hydrogen storage materials.

Herein, we report that the partial ionicity induced by solvent adsorption on Mg-based alloy (Mg87.5Ni5.5Y7) can improve nanosizing effects during ball-milling. The solvent-assisted process yielded particles as small as ~511.87 nm, representing an 88-fold size reduction compared to solvent-free milling (~44.89 μm). The nanostructured Mg87.5Ni5.5Y7 exhibited dehydrogenation rates of 1.71 wt.% min‒1 at 300 °C and 0.46 wt.% min‒1 at 240 °C, surpassing the previously reported related Mg-based alloys. Density functional theory (DFT) calculations revealed that solvent (THF) adsorption on Mg surfaces induce electron redistribution. Mg atoms bonded to oxygen lose 0.56 e⁻, while adjacent Mg atoms gain partial electrons. It gives rise to partial ionicity in Mg (Mgδ+‒Mgδ‒ dipole structure), and as a result, the surface hardness (Vickers hardness) is ~22% enhanced by THF-adsorption. The ion migration during balling induces Coulombic repulsion between Mg atoms, maximizing stress-driven structural deformation and subsequent nanoparticle formation. The present data demonstrate the microscopic role of solvents on modifying the material’s surface electronic structure and mechanical properties, offering an affordable strategy for nanoparticle fabrication.

Results
Solvent-Assisted Ball-Milling of Mg87.5Ni5.5Y7

Figure 1a illustrates the powdering process achieved through mechanical impact, such as ball milling, of metallic and ionic compounds. Metallic materials are characterized by delocalized electrons, and structural deformation contributes a little to the metallic bond structure, resulting in more ductile properties. Meanwhile, ionic materials involve both positively and negatively charged ions. The slight modification of crystal deformation immediately gives rise to Coulombic repulsion between identically charged ions and generates cracks. Therefore, it is widely believed that ionic materials are more susceptible to mechanical impact and are easily nanosized by ball-milling. Mg87.5Ni5.5Y7 is one of the Mg-based alloys that is characterized by highly ductile properties. Prior to THF adsorption, ball-milling does not provide a powdering effect to the Mg87.5Ni5.5Y7 alloy, and the mean particle size was estimated to be 44.89 μm (Fig. 1b and Supplementary Figs. S1 and S2). Meanwhile, introducing 1 mL THF significantly enhanced the ball-milling effect. The THF-assisted milling achieved a mean particle size of 511.87 nm, an 88-fold reduction compared to the solvent-free process. Its BET area was measured as 3.72 m2·g‒1 (Fig. 1c). The particle size distribution analysis indicates that 78 % of the particulates fall within the range of 300~900 nm, highlighting the efficacy of THF-assisted milling in achieving uniform dispersion and considerable nanosizing during the ball-milling process.

High-resolution TEM (HRTEM) characterization of the Mg87.5Ni5.5Y7 alloy (Fig. 1d) reveals characteristic lattice fringes with measured d-spacings of 0.19 nm, which corresponds to the (102) plane of hexagonal close-packed Mg. It agrees well with the XRD data, in which Mg and Mg2Ni were observed as major and minor phases, respectively (Supplementary Fig. S3). EDS elemental mapping verifies homogeneous distribution of Mg and Y across the matrix, whereas Ni demonstrates preferential segregation within the Mg2Ni domains. These findings are further supported by XPS analysis of the electronic states of the elements within Mg87.5Ni5.5Y7 NPs. Specifically, Mg displays a Mg0 peak at 1303.5 eV, metallic Ni 2p3/2 signatures at 852.7 eV in Mg2Ni phases, and the Y 3d spectrum reveals spin-orbit resolved doublets where the 3d5/2 component at 156.8 eV is assigned to Yδ+ (Fig. 1e–g). These XPS characteristics are consistent with Mg-Ni-Y alloys reported elsewhere43,44, demonstrating that the Mg87.5Ni5.5Y7 NPs are not oxidized during the nanosizing process.

It is emphasized that not all solvents can enhance the nanosizing effect of ball-milling. As shown in Supplementary Fig. S4, the addition of cyclohexane exhibits little contribution to ball-milling efficiency, and the observed mean particle size is largely consistent with that of the dry-milled samples. The mechanism difference of THF- and cyclohexane-assisted nanosizing effect will be discussed later.

Hydrolysis corrosion properties

While nanosizing enhances the hydrogen storage capability of Mg-based hydrogen storage materials, it also increases their surface areas, thereby worsening the material’s chemical stability. This is because the increased surface area provides more reactive sites for impurities such as water to form Mg(OH)2. The degradation of Mg2Ni further exacerbates the surface corrosion rates45. To tackle the contradiction between hydrogen storage performance and chemical stability of Mg87.5Ni5.5Y7 NPs, we applied PMMA coatings. In this study, we synthesized Mg87.5Ni5.5Y7 NPs with 0, 0.1, 0.5, and 1.0 wt.% PMMA content and successfully suppressed the surface corrosion rates (Fig. 2a). Uncoated nanoparticles exhibited severe hydrogen corrosion with a rapid hydrogen evolution at 51.0 mL·g‒1·min‒1 during the first 3 minutes (Fig. 2b). Meanwhile, the H2 production rate of Mg87.5Ni5.5Y7 NPs with 0.1, 0.5, and 1.0 wt.% PMMA content were estimated to be 14.7 mL·g‒1·min‒1, 7.0 mL·g‒1·min‒1 and 5.6 mL·g‒1·min‒1, respectively, representing a 71.2–89.0% reduction compared to the PMMA-free system (Fig. 2b and Supplementary Figs. S5 and S6). The suppression of H2 evolution is interpreted by the polymer’s (e.g., PMMA) protection effect that suppress direct contact between Mg87.5Ni5.5Y7 NPs and water, which is consistent with previous studies4,22,46–49.

Figure 2c shows the XRD pattern of water-exposed uncoated and 0.1 wt.% PMMA coated Mg87.5Ni5.5Y7 NPs. While the uncoated samples exhibit major Mg(OH)2 peaks, the corresponding peaks are much suppressed for the 0.1 wt.% PMMA-coated samples. To further confirm the surface protection effect of PMMA, we exposed the uncoated and 0.1 wt.% PMMA-coated Mg87.5Ni5.5Y7 NPs under an ambient atmosphere for two weeks (Fig. 2d). Again, a notable Mg(OH)2 peak was confirmed for the uncoated samples, due to the reaction with air moisture. Meanwhile, the XRD is largely unchanged for 0.1 wt.% PMMA-coated Mg87.5Ni5.5Y7 NPs, and both Mg(OH)2 and MgO peaks were not identified. These data demonstrated improved surface stability by the PMMA coating.

Hydrogen storage performances

To investigate the effects of nanosizing and surface modification on the dehydrogenation kinetics of Mg87.5Ni5.5Y7, we conducted dehydrogenation experiments on samples with and without PMMA coating. Prior to the kinetics analysis, all of the samples were activated utilizing de/hydrogenation cycles. While Mg87.5Ni5.5Y7 NPs (THF-assisted ball-milled samples) are easily activated, the dry-milled samples cannot be activated under the same conditions (Supplementary Fig. S7). After the activation, the Mg87.5Ni5.5Y7 NPs transform to MgH2‒Mg2Ni‒YH3 nanocomposite, whose BET surface area is 9.51 m2·g‒1 (Supplementary Fig. S8). In it, both Mg2Ni and YH3 exhibit improved hydrogen adoption and dissociation capability, which serve as catalytically active phases (Supplementary Fig. S9). In YH3, hydrogen deficiencies are required to serve as a catalytically active phase. The de/hydrogenation kinetics of dry-milled sample is fiarly worse than that of Mg87.5Ni5.5Y7 NPs (Supplementary Fig. S7). The dehydrogenation kinetics were analyzed using the JMAK model and the Arrhenius equation to determine the apparent activation energy, and the isothermal kinetics were further explored using the Sharp and Jones method (Supplementary Table S1).

The results revealed that Mg87.5Ni5.5Y7 NPs achieved 5.15 wt.% H2 released within 3 minutes at 300 °C, reaching 95% of its theoretical capacity, and released 4.85 wt.% H2 in 15 min at 240 °C (Fig. 3a). The analysis indicated its apparent activation energy of 69.9 kJ·mol‒1, significantly lower than 158.0 kJ·mol−1 reported for bulk MgH2 (Fig. 3g)50. The isothermal kinetic analysis of Mg87.5Ni5.5Y7 NPs showed that the experimental data were well fitted to the A2 model, suggesting that nucleation-and-growth of Mg/MgH2 phases is the rate-determining step, consistent with previously reported Mg-TM-RE (TM=transition metals, RE=rare earths) hydrogen storage alloys (Supplementary Table S2 and Supplementary Fig. S10)19. These findings demonstrate that nanosizing Mg87.5Ni5.5Y7 notably improves the dehydrogenation kinetics by increasing surface reactivity and reducing diffusion barriers, compared with the reported Mg-TM-RE systems (Fig. 3h and Supplementary Table S3)51–55.

To explore the impacts of PMMA coating on dehydrogenation kinetics, we synthesized Mg87.5Ni5.5Y7 NPs with varying PMMA thicknesses and tested their dehydrogenation capabilities from 240 °C and 300 °C. At 300 °C, 0.1-PMMA@Mg87.5Ni5.5Y7 released 4.85 wt.% H2, while 0.5-PMMA@Mg87.5Ni5.5Y7 showed reduced capacity of 3.65 wt.% H2, with 83.5 kJ·mol‒1 and 81.3 kJ·mol‒1, respectively (Fig. 3b–g). This temperature-dependent trend persisted at 240 °C, with desorption capacities decreasing to 4.46 wt.% and 3.61 wt.% H2, respectively, maintaining the inverse correlation between PMMA loading and hydrogen release efficiency. The kinetic model was fitted to the R2 model (Supplementary Table S4 and Supplementary Fig. S11), indicating that interfacial chemical reactions are the rate-determining step. Accordingly, the PMMA layer on the alloy surface impeded hydrogen molecule access to active sites, thereby suppressing the interfacial reaction kinetics. The suppressed dehydrogenation kinetics corroborate this mechanism observed for 1.0-PMMA@Mg87.5Ni5.5Y7. It only releases 1.94 wt.% H2 at 300 °C and 1.76 wt.% H2 at 240 °C. The excessive coating of 1.0 wt.% PMMA caused a 63.1% hydrogen capacity loss and a 28.0~52.8 % increase in the activation energy (Supplementary Fig. S12). Supplementary Fig. S13 shows the equilibrium pressures of uncoated and 0.1-PMMA@Mg87.5Ni5.5Y7. The reaction enthalpy obtained from van’t Hoff equation shows the nanosizing effect affect little to the thermodynamic properties.

To investigate the effects of different PMMA concentrations on the microstructure and surface chemical states of dehydrogenated Mg87.5Ni5.5Y7 alloys, we systematically analyzed the dehydrogenated samples. All the samples predominantly consist of Mg, Mg2Ni, and YH3 phases after dehydrogenation (Supplementary Fig. S14). The YH3 phase, formed during hydrogen absorption, remains stable without decomposition during the dehydrogenation process. SEM images and EDS elemental mappings of dehydrogenated alloys confirm that the modified samples with varying PMMA concentrations maintain a fine particulate morphology after dehydrogenation (Supplementary Fig. S15). These results indicate that the Mg87.5Ni5.5Y7 alloy maintains a relatively stable phase composition after dehydrogenation, and the different PMMA concentrations have minimal impact on the phase distribution.

HRTEM-EDS mapping (Supplementary Fig. S16) reveals uniform dispersion of Mg and Y elements, whereas Ni predominantly segregates as Mg2Ni clusters within the alloy matrix after dehydrogenation. XPS analysis of uncoated and 0.1 wt.% PMMA-coated samples reveals that the Mg 1s spectra resolves dual peaks, i.e., a metallic magnesium component at 1303.5 eV and an oxidized magnesium species at 1306.3 eV (Supplementary Fig. S17). This indicates that partial surface oxidation exists. Meanwhile, the 0.5 and 1.0 wt.% PMMA samples show a single Mg 1s peak at 1303 eV, verifying complete dehydrogenation without detectable oxidation. The Ni 2p spectra consistently displays spin-orbit doublets (2p3/2 and 2p1/2) accompanied by two characteristic satellite peaks across all compositions. The dominant 2p3/2 component at 852.5 eV corresponds to Mg‒Ni intermetallic bonding, confirming the structural integrity of intermetallic phases during dehydrogenation. The Y 3d spectra reveals spin-orbit split 3d5/2 and 3d3/2 doublets, where the 3d5/2 peak at 157.2 eV is assigned to Y‒H bonding configurations. Optimizing the concentration of PMMA is critical to balance surface protection with hydrogen storage performance. Based on our preliminary analysis of samples with varying PMMA contents, 0.1-PMMA@Mg87.5Ni5.5Y7 exhibited an optimal combination of high hydrogen storage capacity and surface stability. Hence, we selected it for in-depth characterization to further elucidate its cyclic hydrogen storage performance and structural stability.

The cycle stability of uncoated and 0.1-PMMA coated Mg87.5Ni5.5Y7 NPs is presented in Fig. 4a. To each sample, cycle stability was evaluated up to 500 cycles. Both samples exhibit excellent cycling stability, while the uncoated Mg87.5Ni5.5Y7 exhibits persistent degradation in terms of hydrogen storage capacity and reaction rates (Fig. 4a and b). After experiencing 500 cycles, the hydrogen storage capacity changed from 4.86 wt.% to 4.52 wt.%. Notable degradation of hydrogen storage capacity and reaction rates were not identified for the 0.1-PMMA coated Mg87.5Ni5.5Y7 (Fig. 4a and b). The 0.1 wt.% PMMA-coated Mg87.5Ni5.5Y7 material outperforms previously reported Mg-based systems in both de/hydrogenation kinetics and cycling stability (Fig. 4c)56–80. These results collectively confirm that the PMMA-coated nanocrystalline alloy powder possesses exceptional chemical stability and remarkable cyclic durability under repeated hydrogen cycling conditions.

Figure 5a shows the Mg, Ni and Y XPS data of 50 and 500 times cycled uncoated Mg87.5Ni5.5Y7 NPs. The Mg 1s spectrum exhibits a metallic Mg0 peak centered at 1303.5 eV, indicating minimal surface oxidation. The Ni 2p spectrum was deconvoluted into Mg‒Ni alloy components with two satellite peaks, confirming the presence of Mg‒Ni intermetallic bonding. The Y 3d spectrum revealed a Y‒H bonding doublet, suggesting the presence of YH3 species. These are consistent with XRD data of cycled samples (Supplementary Figs. S14 and S18), showing that the Mg, Mg2Ni and YH3 are the major phases for the dehydrogenated samples. It should be highlighted that the XPS data is largely unchanged between the 50th and 500th samples for the uncoated Mg87.5Ni5.5Y7 NPs (Fig. 5a). These suggest that the degradation of hydrogen storage capability is likely associated with the structural or morphological evolution of materials. Figure 5b, c show the TEM images of uncoated and 0.1 wt.% PMMA-coated Mg87.5Ni5.5Y7 NPs after 500 cycles. While the chemical composition keeps relatively uniform, clear particle growth can be identified for the uncoated Mg87.5Ni5.5Y7 NPs relative to the 50th cycled sample (Supplementary Fig. S19), leading to the degraded cycle properties. Meanwhile, the 0.1 wt.% PMMA coated sample kept its finer particle morphology even after the long-cycle experiment (Fig. 5c). These data show that PMMA coating not only improves chemical stability but also helpful to keep particle morphology during long-cycle operation.

Solvent-mediated partial ionicity

To elucidate the formation mechanism of nanostructures, we calculated the adsorption energies of THF at four distinct sites on the Mg (0001) surface, namely FCC-hollow, Bridge, HCP-hollow and on-top sites. Figure 6a shows that the most stable adsorption configuration occurs at the on-top site above Mg atoms, with an adsorption energy of ‒0.48 eV. Bader charge analysis (Fig. 6b and Supplementary Fig. S20) combined with charge distribution difference calculations reveal that upon THF adsorption, the Mg atom coordinating with the oxygen loses 0.56 electrons (manifested as positive charge states in blue), while neighboring Mg atoms acquire 0.03-0.08 electrons, giving rise to the Mgδ+‒Mgδ‒ dipole structure. Additionally, Y-doping induces charge redistribution in proximal Mg atoms, with electron transfer quantities as quantified (Supplementary Fig. S21). Different views of the THF adsorption structure at the on-top site of Mg and the Y-doped structure are shown in Supplementary Fig. S22. Indeed, XPS analysis on the as milled samples show the presence of both positively and negatively charged Mg, consistent with the DFT calculations (Supplementary Fig. S23). The synergistic interplay between THF adsorption and Y-doping induces localized charge redistribution in adjacent Mg atoms, giving rise to the partial ionicity between Mg atoms.

Subsequently, we discuss the impact of cyclohexane adsorption on Mg-based alloy. Unlike THF, cyclohexane is characterized by less polarized molecules and exhibits much weaker interaction (Supplementary Figs. S24 and S25). As a result, adsorption of cyclohexane on Mg-alloy does not cause drastic charge redistribution as was observed for THF, and this is why cyclohexane does not obviously contribute to mechanical nanosizing efficiency (Supplementary Fig. S4).

Such partial ionicity should suppress the ductility of Mg metal and maximize the ball milling-driven structural deformation that facilitates ion migration. The modified surface mechanical properties are indeed demonstrated utilizing Vickers hardness measurements. As shown in Fig. 6c and d, the notable enhancement of Vickers hardness was identified for the THF-adsorbed sample. It achieved an average hardness of 103.1 HV, representing 21.9% enhancement compared to the untreated surface (84.6 HV) (Supplementary Table S5). These results indicate that the Mg surface becomes less ductile and thus more susceptible to mechanical impact, leading to enhanced ball-milling efficiency. Based on the above-mentioned results, we propose the mechanism of solvent-assisted mechanical nanosizing processes in Fig. 6e. First, solvent molecules, such as THF, adsorb on the metal surface. It gives rise to an electron distribution that enhances surface hardness, making the metal powder more susceptible to mechanical impact. Subsequently, the cracks are easier to generate, and the solvent molecules can be adsorbed on the fresh surface, further enabling mechanical nanosizing.

In summary, we successfully prepared Mg87.5Ni5.5Y7 NPs utilizing solvent-assisted mechanical ball-milling. It was found that the adsorption of solvents, such as THF, induces weak positive and negative charge distribution on the Mg surface. This partial ionicity enhances the surface hardness of Mg-based alloys, thereby enhancing mechanical nanosizing effect by the factor of 88. The nanostructured Mg87.5Ni5.5Y7 alloy exhibited exceptional dehydrogenation kinetics, releasing 5.22 wt.% H2 within 5 min at 280 °C and 4.85 wt.% H2 in 15 min even at 240 °C, exceeding previously reported Mg-based hydrogen storage alloys. The Mg87.5Ni5.5Y7 NPs also exhibit improved chemical and cycle stability, which is further improved by PMMA coating. Notable kinetic and capacity degradation were not identified for the PMMA-coated Mg87.5Ni5.5Y7 NPs, at least for 500 cycles. The present data reveals how the charge rearrangement caused by metal-solvent interfaces affect surface mechanical properties and improve nano-fabrication efficiency of ball-milling, which is even applicable to ductile Mg-based alloys, offering an affordable and scalable approach for nanofabrication of inorganic materials.

Methods
Sample preparation
The Mg87.5Ni5.5Y7 alloy was synthesized via vacuum induction melting under a high-purity argon atmosphere. Nanostructured alloy powder was fabricated via solvent-assisted ball milling with the following protocol. In it, 0.25 g of alloy powder was weighed into a 100 ml ball milling jar, followed by uniform dispersion of 1 mL tetrahydrofuran (THF) onto the powder surface using a calibrated micropipette. Ball milling was performed for 20 hours under an argon atmosphere. During the ball-milling, ball-to-powder ratio was set to 40:1 with the rotational speed of 400 rpm. 10 min pause was introduced between every 10 min milling to prevent overheating.

Three polymethylmethacrylate (PMMA, Mw~120,000, Aladdin, Shanghai) solutions with distinct concentrations (0.00025 g/mL, 0.00125 g/mL, 0.0025 g/mL) were prepared by dissolving 0.005 g, 0.025 g, and 0.050 g PMMA in 20 mL anhydrous tetrahydrofuran (99.5%, Macklin). The mixtures were magnetically stirred in argon-protected containers for 12 h to ensure complete dissolution. Three gradient-concentration PMMA solutions were mixed with the Mg87.5Ni5.5Y7 NPs for PMMA coating. The experimental parameters are summarized in Supplementary Table S6. For the test of cyclohexane-assisted ball-milling the solvent was changed to cyclohexane (99.9%, Macklin) while other conditions were retained.

Material characterization
The phase composition of the alloy was characterized by X-ray diffraction (XRD, Rigaku, Cu Kα radiation) at a step size of 0.02° and a scanning rate of 5°·min‒1. The morphological characteristics and elemental distribution of the samples were analyzed before and after hydrogen storage performance tests using environmental scanning electron microscopy (ESEM, Quattro S) and transmission electron microscopy (TEM; Talos F200S). The error bars of particle distribution were estimated utilizing Gaussian fit with statistics from n=147. The elemental composition and stoichiometry of the samples were determined by X-ray photoelectron spectroscopy (XPS, Thermo K-Alpha, Thermo Fisher Scientific). The C 1s peak (284.8 eV) was used as the reference. During the fitting of the Mg 1s, Ni 2p, and Y 3d core-level spectra, the full width at half maximum (FWHM) was constrained to be identical for multiple peaks belonging to the same chemical state, and their area ratios were fixed. For the Ni 2p and Y 3d doublet, the area ratio and splitting set at reference values, thereby ensuring the accuracy of the results.

Vickers hardness testing using a 300 g load and a holding time of 10 s was carried out on a Micro Vickers hardness testing system (UHV-1000F). For each sample, 120 tests were performed and utilized for statics estimation, such as median/mean values, the 25th/75th quartiles and standard deviation. Prior to the Vickers hardness test, the alloys surface was polished. THF was then sprayed to the surface, and the excess THF was dried under an ambient condition.

Hydrolysis corrosion experiments
The hydrolysis corrosion reaction was tested in a hydrolysis H2 generation device based on the drainage method. The equipment consists of a reaction device and a data logger. The amount of gas released was measured by the mass of water discharged from the gas scrubber bottle, which was weighed with a balance throughout. The weight of deionized water was obtained automatically by an electronic balance connected to a computer, which stored data and reaction times. Then, the generation rate-time curve corresponding to the hydrolysis H2 generation process was obtained. For hydrolysis corrosion reaction, in a typical experiment, ~0.05 g of the sample was loaded in a 100 ml flask bottle, and 50 ml deionized water was added to start the reaction at room temperature and under ambient atmosphere. The flask was immersed in a circulating water bath to control reaction temperature and put over a magnetic stirrer to agitate the reactions. The background flow (0.15 ml·min‒1) was subtracted from the data.

Hydrogen storage performance measurements
The isothermal hydrogenation and dehydrogenation characteristics of the samples, including cycling curves, were investigated using a commercial Sieverts-type apparatus (PCT Pro, Setaram and H-Sorb 4600 PCT PRO, Ciqtec). The background hydrogen pressures were set to 30 bar for hydrogenation and 0.01 bar for dehydrogenation processes, respectively. Prior to hydrogen absorption/desorption kinetics testing, the alloy samples were subjected to 5 activation cycles at 300 °C.

To better understand the dehydrogenation kinetic behavior, the equations associated with Sharp’s and Jone’s method model (equation 1 and 2) were used to fit the isothermal dehydrogenation curves at different temperatures81,82

dα/dt = kf(α) (1)

f(α) = A(t/t0.5) (2)

where dα is the differential of conversion extent, dt is the differential of time, k is the reaction rate constant, f(α) is a function that depends on a specific kinetic mechanism, A represents a constant related to the kinetic mechanism, and t0.5 represents the time when α equals to 0.5.

The kinetic model was obtained by plotting the experimental value (t/t0.5)exp against the theoretical value (t/t0.5)theo.

The PCT curves of samples in different cycle stages were tested at temperatures of 280 °C, 300 °C, 320 °C, 340 °C and hydrogen pressures ranging from 0.02 to 3 MPa. The reaction enthalpy (ΔH) was then calculated utilizing van’t Hoff equation.

ln(PH2/P0) = ΔH/RT - ΔS/R (3)

where PH2 denotes the equilibrium pressure of the Mg/MgH2 reaction, ΔH represents the enthalpy change, ΔS is the entropy change, T is the testing temperature, R is the gas constant [8.314 J/(mol·K)], and P0 the standard atmospheric pressure (~1.01×105 Pa).

DFT studies
All structural relaxations and electronic structure calculations were carried out using density functional theory (DFT) as implemented in the Vienna Ab initio Simulation Package (VASP)83. Based on the optimized bulk lattice parameters, a 6×6 Mg(0001) surface model was constructed for the simulations, with a vacuum layer of 15 Å introduced to eliminate interlayer interactions. The supercell consists of five atomic layers, with the bottom three layers fixed to mimic the bulk behavior and improve computational efficiency. During geometry optimization, the plane-wave cutoff energy was set to 450 eV, and a Monkhorst-Pack k-point mesh of 5×5×1 was used. The convergence criteria for electronic self-consistent iteration and ionic relaxation were set to 10−6 eV per atom for the total energy and 0.01 eV·Å−1 for the maximum atomic force, respectively. The exchange-correlation functional was treated within the generalized gradient approximation (GGA) using the Perdew-Burke-Ernzerhof (PBE) formulation84–86. The electron charge density was analyzed using the Bader charge analysis method combined with charge density difference calculations. For the transition state analysis of the hydrogen dissociation reaction, the climbing image nudged elastic band (CI-NEB) was employed to identify saddle points and minimum energy paths (MEPs)87,88. The number of CI-NEB images was determined using the VASP-VTST 978 script, based on the interatomic distances between the initial and final states. All transition-state models were fully optimized until the energy and force convergence thresholds of 1×10−6 eV and 0.01 eV·Å−1, respectively, were reached. The adsorption energy Ead(X) of X (X=tetrahydrofuran or cyclohexane) on the surface is defined as Ead(X) = Etot(X/sub) - E(sub) - E(X), where Etot(X/sub) is the total energy of the optimized configuration of the X molecule adsorbed on the substrate, E(sub) is the total energy of the clean substrate, and E(X) is the total energy of an isolated X molecule.

Peer review rounds
[{"round": 1, "reviewer_comments": {"reviewer_1": "It has long been recognized that adding solvents can enhance ball-milling efficiency, yet the underlying microscopic mechanism has remained elusive. In this study, the authors revealed that surface electronic redistribution as the origin of high mechanical nanosizing effect. Specifically, the formation of a Mgδ+-Mgδ− dipole structure induced by solvent adsorption plays a key role. It is interesting that the mechanism is applicable to Mg-based alloys, a well-known class of ductile materials, leading to a 54-fold improvement in mechanical nanosizing efficiency. The increased surface area of the processed Mg-Ni-Y alloys results in exceptional hydrogen storage kinetics, enabling complete hydrogen release in 3 min at 300 °C. Additionally, the authors employ a polymer coating technique to address surface sensitivity caused by ball-milling, further enhancing material stability. The manuscript is well-organized and this method holds potential application for the hydrogen science community. It also provides valuable microscopic insights into affordable and scalable nanofabrication methods. However, some improvements are needed to justify publication in Nature Communications. My comments are listed below.\n1.The PMMA has been reported used in the Mg-based hydrogen storage materials as a hydrogen-permeable antioxidant organic materials. It effectively improved the ball-milling efficiency, promoted nano-confinement, and prevented the oxidation of the hydrogen storage material. Please clearly state the innovation of the manuscript in the abstract and the introduction part.\n2.One of the most important conclusions of this manuscript is that the Mgδ+‒Mgδ− dipole structure induced by solvent adsorption significantly enhances the mechanical nanosizing efficiency. Yet, the authors solely considered THF and did not evaluate the effects of other solvents. Are there any solvents with weaker interactions with Mg-Ni-Y? If the authors’ claim is valid, such solvents should be less effective in enhancing mechanical ball-milling efficiency.\n3.In addition to the dipole-structure-induced mechanism, the polymer coating is an interesting approach to enhance the chemical stability of ball-milled nanoparticles. However, it seems that the authors failed to clarify its contribution to cycle stability. The cycling properties between polymer-coated and polymer-free systems should be compared. In addition, 50 cycles appears to be too short to reliably assess.\n4.The authors should explain in the introduction that the Mg-Ni-Y alloy decomposes into Mg, Mg2Ni, and YHx during the activation process, so that readers can more easily follow the appearance of the corresponding phases in Figure 4.\n5.This comment is related to comment 3. The authors claimed in the introduction that Mg2Ni contributes to hydrogen adsorption and dissociation. Meanwhile, some literature reports suggest that YHx also contributes to this step. Which is more accurate? This claim should be supported further.\n6.I believe water-resistance capability is also important for practical applications. While I appreciate the authors providing detailed hydrolysis corrosion kinetics, the associated phase modifications are not discussed. XRD data of water-exposed samples should be provided.\n7.The authors did not provide information about the sample activation process. I suggest the authors to compare the activation and hydrogen storage kinetics of Mg-Ni-Y [repared using solvent-mediated and solvent-free processes.\n8.While I appreciate the effect of polymer coating on Mg-Ni-Y nanoparticles, it should be noted that similar approaches have already been reported. Please adequately cite the relevant literature and discuss the similarities and differences.\n9.Please consider optimizing Figures 6b and 6c. While Figure 6c shows opposite charge states between Mg1 and Mg2/3/4/5/6, such electron distributions are difficult to confirm in Figure 6b.\n10.There are some issues with the wording and spelling in the manuscript and supprementary information. For example, the “kinetic” should be “kinetics”. The “temperatur” should be “temperature” in Figure S2.", "reviewer_2": "In this report, the authors present a nanosized Mg-Ni-Y alloy synthesized via a solvent-mediated milling process, demonstrating enhanced hydrogen storage performance. Compared to dry milling, THF-mediated milling yields smaller Mg-based alloy particles with improved hydrogen absorption and desorption properties. However, solvent-mediated milling is a well-established method for particle size reduction (e.g., doi.org/10.1002/er.7046, doi.org/10.1007/s12598-025-03330-x, doi.org/10.1039/C3TA01332D, etc.), and the performance reported here is not particularly exceptional. Although the authors include a comparative chart in Fig. 3h, there are numerous previous studies that report comparable or even superior performance. In addition, the authors claim that THF-adsorption leads to the charge redistribution on Mg surface, promoting the nanosize effect; however, the data provided are insufficient to substantiate this claim.\nWhile it is not necessary for a study to demonstrate record-breaking results, in the absence of a novel synthetic strategy or innovative structural design, I do not believe this work meets the standards for publication in Nature Communications.\n1. Although the authors claim that solvent-free milling results in only a moderate size reduction, other reports have demonstrated a significantly greater decrease in particle size. (ex. doi.org/10.1016/j.ijhydene.2020.09.251, doi.org/10.1016/j.jallcom.2009.11.154, https://doi.org/10.1016/j.jallcom.2004.08.102).\n2. For the particle size characterization, additional high-resolution SEM images are needed, as Figure 1a lacks sufficient clarity to support accurate size estimation. While the mean particle diameter is reported as 813 nm, the resolution of the provided image is too low to justify this calculation. Moreover, the standard deviation should also be reported. Based on Figure 1a, the sample appears to exhibit a broad size distribution, with many larger particles clearly visible—contrary to the histogram presented in the image. In particular, the claim that the particles are \"monodisperse\" (line 95) is not supported by the data and should be reconsidered.\n3. Please provide detailed experimental conditions for the dry ball-milling process. In ball-milling, parameters such as the size and type of milling balls can significantly influence the resulting particle characteristics.\n4. In the Mg 1s XPS results, only the metallic Mg peak was observed, despite the well-known tendency of Mg to form a thin oxide layer—even under ultra-high vacuum conditions—due to its instantaneous reaction with trace amounts of oxygen. It is recommended that the Mg 2p spectrum be provided, as it is more sensitive to detecting surface oxides. Furthermore, the authors claim that THF adsorption induces a change in the charge distribution on the Mg surface. If THF adsorption indeed affects the surface chemistry, the presence of at least a partial oxide layer should be detectable to support this claim.\n5. PCT isotherms, along with the corresponding thermodynamic parameters, should be measured to properly characterize the hydrogen storage properties. This type of analysis is essential for studies involving metal hydride-based hydrogen storage materials."}, "author_response": "Response to the reviewer\nReviewer #1\nIt has long been recognized that adding solvents can enhance ball-milling\nefficiency, yet the underlying microscopic mechanism has remained elusive. In this\nstudy, the authors revealed that surface electronic redistribution as the origin of high\nmechanical nanosizing effect. Specifically, the formation of a Mgδ+-Mgδ− dipole\nstructure induced by solvent adsorption plays a key role. It is interesting that the\nmechanism is applicable to Mg-based alloys, a well-known class of ductile materials,\nleading to a 54-fold improvement in mechanical nanosizing efficiency. The increased\nsurface area of the processed Mg-Ni-Y alloys results in exceptional hydrogen storage\nkinetics, enabling complete hydrogen release in 3 min at 300 °C. Additionally, the\nauthors employ a polymer coating technique to address surface sensitivity caused by\nball-milling, further enhancing material stability. The manuscript is well-organized and\nthis method holds potential application for the hydrogen science community. It also\nprovides valuable microscopic insights into affordable and scalable nanofabrication\nmethods. However, some improvements are needed to justify publication in Nature\nCommunications. My comments are listed below.\nResponse: We appreciate the referee for excellently summarizing our key findings. We\nare also grateful to the referee for the positive evaluations on this work.\n1. The PMMA has been reported used in the Mg-based hydrogen storage materials as a\nhydrogen-permeable antioxidant organic materials. It effectively improved the ball-\nmilling efficiency, promoted nano-confinement, and prevented the oxidation of the\nhydrogen storage material. Please clearly state the innovation of the manuscript in the\nabstract and the introduction part.\nResponse: Thank you for highlighting this point. The application of PMMA in\nMg-based hydrogen storage materials has indeed been extensively reported. As a gas-\nseparation material with high hydrogen selectivity and low oxygen permeability,\nPMMA can form protective coatings via solution blending or mechanical ball milling.\nThe PMMA coating was employed to protect the surface of nano-alloy particles\nsynthesized by THF-assisted ball milling, thereby blocking oxygen/water vapor\npermeation and suppressing surface oxidation. Meannwhile, PMMA is unlikely to\nenhance nano-sizing efficiency. We conducted THF-PMMA-assisted ball-milling\nutilizing identical experimental parameters, but it resulted in worse ball-milling\nefficiency, as observed in Figure R1-1 (Mean particle size: ~2.56 μm). Therefore, it is\nconsidered that PMMA mainly contributes to improve chemical stability of Mg-Ni-Y\nNPs in this study. In the revised manuscript, we clarified these points and included\nappropriate references.\nFigure R1-1 (a) SEM image of the alloy after ball milling assisted by the THF-PMMA\nmixed solution, (b) The corresponding particle size distribution of the alloy.\nThe major finding of this work lies in elucidating the nano-size mechanism of\nsolvent-assisted ball milling. Despite the wide recognition of the metallurgy and nano-\nmaterial community on the effect of some solvents during ball-milling, the microscopic\nmechanism of solvent has remained unsolved. The knowledge gap has forces solvent\nselection to rely on trial-and-error, which is highly time and resource consuming. In this\nsubmission, we discovered that the electron re-distribution caused by the solvent (THF\nin our study) adsorption is the key factor that modifies surface mechanical properties\nof Mg and its alloys. As shown in the Figure R1-2a, the THF adsorption on Mg lattice\ngives rise to the Mgδ+-Mgδ− dipole structure, and such ionic nature should enhance\nsurface mechanical properties. The enhanced surface mechanical properties are indeed\nevinced utilizing Vickers hardness test. As shown in Figure R1-2b, the notable\nenhancement of Vickers hardness was identified by THF adsorption. The THF-sprayed\nsurface achieved an average microhardness of 103.13 HV, representing 22 %\nenhancement compared to the untreated surface (84.58 HV). These results indicate that\nthe Mg surface becomes less ductile and thus more susceptible to mechanical impact,\nleading to enhanced ball-milling efficiency. These findings provide a microscopic\ninsight into optimizing ball-milling parameters.\nFigure R1-2 (a) Bader charge analysis and charge density difference analysis for the\non-top adsorption site and Y-doped configuration. (b) Microhardness distribution map.\nTo further clarify the point, we added more descriptions as follows.\n----------------------------------------------------\nPage 14: Such partial ionicity should suppress the ductility of Mg metal and maximize\nthe ball milling-driven structural deformation that facilitates ion migration. The\nenhanced surface mechanical properties are indeed demonstrated utilizing Vickers\nhardness measurement. As shown in Figure 6c and d, the notable enhancement of\nVickers hardness was identified for the THF-adsorbed sample. It achieved an average\nhardness of 103.13 HV, representing 22 % enhancement compared to the untreated\nsurface (84.58 HV) (Supplementary Table S5). These results indicate that the Mg\nsurface becomes less ductile and thus more susceptible to mechanical impact, leading\nto enhanced ball-milling efficiency. On the basis of the above-mentioned results, we\npropose the mechanism of solvent-assisted mechanical nanosizing processes in Figure\n6e. First, solvent molecules, such as THF, adsorb on metal surface. It gives rise to\nelectron distribution that enhances surface hardness, making the metal powder more\nsusceptible to mechanical impact. Subsequently, the cracks are easier to generate, and\nthe solvent molecules can be adsorbed on the fresh surface, further enabling mechanical\nnanosizing.\n----------------------------------------------------\n2. One of the most important conclusions of this manuscript is that the Mgδ+‒Mgδ−\ndipole structure induced by solvent adsorption significantly enhances the mechanical\nnanosizing efficiency. Yet, the authors solely considered THF and did not evaluate the\neffects of other solvents. Are there any solvents with weaker interactions with Mg-Ni-\nY? If the authors’ claim is valid, such solvents should be less effective in enhancing\nmechanical ball-milling efficiency.\nResponse: Thank you for the comment. Following this suggestion, we prepared\ndry-milled and cyclohexane-assisted ball-milled Mg-Ni-Y alloy samples using identical\nprocessing parameters for comparison. Here, cyclohexane was selected due to its lower\npolarity and its inert characteristics towards Mg and its alloys. Figure R1-3 presents\nthe SEM images and particle size distributions of the dry-milled, cyclohexane-assisted,\nand THF-assisted ball-milled samples, respectively. The mean particle size of\ncyclohexane-assisted ball milling was 40.20 μm, which is largely consistent with dry-\nmilled condition, demonstrating no significant refinement effect compared to dry\nmilling. Notably, the THF-assisted ball milling achieved a substantially reduced\naverage particle size of 511.87 nm, indicating exceptional refinement capability.\nFigure R1-3 (a), (b), and (c) show the SEM images of the dry ball-milled, cyclohexane-\nassisted ball-milled, and THF-assisted ball-milled samples, respectively; while (d), (e),\nand (f) present the corresponding particle size distributions.\nTo unveil the difference of the impact of THF and cyclohexane on ball-milling\nefficiency, we then conducted Bader charge analysis of the cyclohexane-adsorbed Mg.\nFirst, we calculated the adsorption energies of cyclohexane molecules at four\ncharacteristic sites on the Mg (0001) surface, including FCC-hollow, Bridge, HCP-\nhollow, and on-top sites, respectively. Figure R1-4 presents the adsorption energies of\ncyclohexane at various sites on the Mg (0001) surface. The results show that the most\nstable adsorption configuration occurs at the top site of the on-top sites of Mg, with an\nadsorption energy as low as ‒0.16 eV. The adsorption energy shows the weaker\ninteraction between cyclohexane and Mg relative to the case of THF-adsorption.\nBader charge analysis indicates that although adsorbed at the same top site,\ncyclohexane and tetrahydrofuran (THF) exhibit strikingly different electronic behaviors.\nThe interaction between cyclohexane and magnesium is exceptionally weak, resulting\nin a negligible net charge change on its carbon atom (~0.05 e). In stark contrast, the\noxygen atom in THF undergoes significant electron transfer (ΔQ = 1.11 e), with a\ncharge change magnitude 22 times greater than that of cyclohexane. The large charge\ntransfer mechanism resulted in the distinct adsorption energy between the two solvents.\nFigure R1-4 (a) Calculated adsorption energies of cyclohexane at four sites on the\nMg(0001) surface (b) Bader charge analysis for the on-top adsorption site.\nTo further clarify the point, we added more descriptions as follows.\n----------------------------------------------------\nPage 5: Interestingly, not all solvents can enhance the nanosizing effect of ball-milling.\nAs shown in Supplementary Figure S4, the addition of hexane exhibits little\ncontribution to the ball-milling efficiency, and the observed mean particle size is largely\nconsistent with that of the dry-milled samples. The mechanism difference of THF- and\ncyclohexane-assisted nanosizing effect will be discussed later.\nPage 13: Subsequently, we discuss the impact of cyclohexane adsorption on Mg-based\nalloy. Unlike THF, cyclohexane is characterized by the less polarized molecules and\nexhibits much weaker interaction (Supplementary Figure S22 and 23). As a result,\nadsorption of cyclohexane on Mg-alloy does not cause drastic charge redistribution as\nwas observed for THF, and this is why cyclohexane does not contribute to mechanical\nnanosizing efficiency (Supplementary Figure S4).\n----------------------------------------------------\n3. In addition to the dipole-structure-induced mechanism, the polymer coating is an\ninteresting approach to enhance the chemical stability of ball-milled nanoparticles.\nHowever, it seems that the authors failed to clarify its contribution to cycle stability.\nThe cycling properties between polymer-coated and polymer-free systems should be\ncompared. In addition, 50 cycles appears to be too short to reliably assess.\nResponse: Thank you for the comment. Following this suggestion, we have\nsupplemented the data with 500 reversible hydrogen absorption/desorption cycles of\nMg-Ni-Y NPs and 0.1PMMA-Mg-Ni-Y NPs, and compared the cycling stability of the\ntwo materials.\nFigure R1-5 illustrates the isothermal hydrogen absorption/desorption behavior\nof THF-assisted ball-milled Mg Ni Y and 0.1PMMA-Mg Ni Y over 500 long-\n87.5 5.5 7 87.5 5.5 7\nterm cycles at 280 °C. In the first cycle, Mg-Ni-Y and 0.1PMMA-Mg-Ni-Y exhibited\nreversible de/hydrogenation capacities of approximately 4.86 wt.% H and 4.87 wt.%\n2\nwt.% to 4.52 wt.% for Mg-Ni-Y and from 4.87 wt.% to 4.85 wt.% for 0.1PMMA-Mg-\nNi-Y, corresponding to capacity retention rates of 93.0% and 99.6%, respectively. These\nresults demonstrate that a thin PMMA coating effectively suppresses capacity decay\nduring repeated cycling. Furthermore, as isothermal de/hydrogenation profiles of the\n1st and 500th cycles, both materials maintained stable de/hydrogenation kinetics\ndespite minor capacity losses.\nFigure R1-5 (a) Reversible hydrogen adsorption/desorption cycling curves of THF-\nassisted ball-milled Mg Ni Y at 280 °C. (b) Hydrogen adsorption/desorption\n87.5 5.5 7\ncurves of Mg Ni Y at different cycle numbers.\n87.5 5.5 7\nTo further clarify the point, we added more descriptions as follows.\n----------------------------------------------------\nPage 11: Figure 5a shows the Mg, Ni and Y XPS data of 50 and 500 times cycled\nMg Ni Y NPs. The Mg 1s spectrum exhibits a metallic Mg0 peak centered at 1303.5\n87.5 5.5 7\neV, indicating minimal surface oxidation and suggesting that the PMMA coating\neffectively protects the alloy surface. The Ni 2p spectrum was deconvoluted into Mg‒\nNi alloy components with two satellite peaks, confirming the presence of Mg‒Ni\nintermetallic bonding. The Y 3d spectrum revealed a Y‒H bonding doublet, suggesting\nthe presence of YH species. These are consistent with XRD data (Supplementary\n2 3\nthe dehydrogenated samples. It should be highlighted that the XPS data is largely\nunchanged between the 50th and 500th samples. These suggest that the degradation of\nuncoated Mg Ni Y is likely due to the structural evolution. Figure 5b and c show\n87.5 5.5 7\nthe TEM images of uncoated and 0.1 wt.% PMMA coated Mg Ni Y NPs after 500\n87.5 5.5 7\ncycles. While the chemical composition keeps relatively uniform, clear particle growth\ncan be identified for the uncoated Mg Ni Y NPs relative to the 50th cycled sample\n87.5 5.5 7\n(Supplementary Figure S18), leading to the degraded cycle properties. Meanwhile,\nthe 0.1 wt.% PMMA coated sample kept its original morphology even after the long\ncycle experiment (Figure 5c). These data show that PMMA coating improves chemical\nstability but also helpful to keep particle morphology during service.\n----------------------------------------------------\n4. The authors should explain in the introduction that the Mg-Ni-Y alloy decomposes\ninto Mg, Mg Ni, and YH during the activation process, so that readers can more easily\n2 x\nfollow the appearance of the corresponding phases in Figure 4.\nResponse: Thank you for the comment. In Mg-based alloy (e.g., Mg-Ni-Y), most\nof them experiences in-situ decomposition after activation process because of the strong\nMg‒H bond that destroy Mg-Ni-Y alloy structure. After activation, the major phases of\nthe Mg-Ni-Y alloy are the Mg, Mg Ni and YH nanoparticles. In it, Mg contributes\n2 3\nhydrogen storage capacity and Mg Ni serves as catalyst phases, constructing the Mg-\n2 3\nMg-Mg Ni nanocomposite, serving as the pining center that controls particle growth.\nthe improved cycle stability of the Mg-based hydrogen storage materials.\nFigure R1-6 The phase modification of Mg-Ni-Y alloy by the activation process.\nTo further clarify the point, we added more descriptions as follows.\n----------------------------------------------------\nPage 3: They undergo in-situ decomposition after the activation process and generate\nMg, Mg Ni, and REH nanoparticles (NPs). In it, Mg contributes hydrogen storage\n2 x\ncapacity, while REH and Mg Ni serve as catalyst phases for hydrogen adsorption and\nx 2\ndissociation, thereby enhancing the rates of hydrogenation and dehydrogenation\n(de/hydrogenation).17-21. However, the hydrogen diffusion in Mg phase is hardly\naccelerated by these catalyst phases, and persistent challenges of coarse particle size\ndistribution and suboptimal morphology control are still unresolved.22, 23\n----------------------------------------------------\n5. This comment is related to comment 3. The authors claimed in the introduction that\nMg Ni contributes to hydrogen adsorption and dissociation. Meanwhile, some literature\n2\nreports suggest that YH also contributes to this step. Which is more accurate? This\nx\nclaim should be supported further.\nResponse: We appreciate the referee for pointing out the weakness of our presentation.\nTo address the issue, we additionally conducted DFT calculations and simulated\nhydrogen adsorption and dissociation steps on YH and Mg Ni. The YH and Mg Ni\n3 2 3 2\nphases were utilized in this study because these structures were clearly observed in\npowder XRD data.\nAs shown in Figure R1-7a, b, the YH is characterized by the rich hydrogen\n3\noccupation on YH surface, and H molecules are hardly adsorbed on its surface.\n3 2\nMeanwhile, when hydrogen defects were introduced, YH was found effective for\nx\nhydrogen adsorption and dissociation, as well as Mg Ni. Therefore, we consider that\n2\nboth YHx and Mg Ni collectively contribute to hydrogen adsorption and dissociation.\nhydrogen pump effect. We appreciate the referee for pointing out our errors.\nFigure R1-7 (a), (b), and (c) illustrate the dissociation processes of hydrogen molecules\non the YH (110) surface, YH (110) surface, and Mg Ni (100) surface, respectively.\n3 x 2\nTo further clarify the point, we added more descriptions as follows.\n----------------------------------------------------\nPage 3: They undergo in-situ decomposition after the activation process and generate\nMg, Mg Ni, and REH nanoparticles (NPs). In it, Mg contributes hydrogen storage\n2 x\ncapacity, while REH and Mg Ni serve as catalyst phases for hydrogen adsorption and\nx 2\ndissociation, thereby enhancing the rates of hydrogenation and dehydrogenation\n(de/hydrogenation).17-21. However, the hydrogen diffusion in Mg phase is hardly\naccelerated by these catalyst phases, and persistent challenges of coarse particle size\ndistribution and suboptimal morphology control are still unresolved.22, 23\nPage 8: In it, both Mg Ni and YH exhibits improved hydrogen adoption and\n2 x\ndissociation capability, which serve as catalytically active phases (Supplementary\nFigure S8).\n----------------------------------------------------\n6. I believe water-resistance capability is also important for practical applications.\nWhile I appreciate the authors providing detailed hydrolysis corrosion kinetics, the\nassociated phase modifications are not discussed. XRD data of water-exposed samples\nshould be provided.\nResponse: Thank you for pointing this out. To address the issue we conducted\nXRD measurements to the both water- and air-exposed samples.\nAs shown in Figure R1-8, distinct characteristic peaks of Mg(OH) were observed\n2\nin all water-exposed samples. For the unprotected 0-PMMA sample, the hydrolyzed\nMg-Ni-Y alloy predominantly exhibited Mg(OH) phase with the highest peak intensity.\n2\nsystematically decreased relative to the Mg-Ni-Y matrix, eventually transitioning to\nMg-dominated peaks. This result is consistent with the hydrolysis corrosion tests. A\nsimilar trend was observed in air-exposed samples: while prominent Mg(OH) peaks\n2\nwere detected in the 0-PMMA sample, no significant Mg(OH) peaks were observed in\nforms an effective protective layer against both air and water corrosion, significantly\nenhancing the chemical stability of the Mg-Ni-Y nanoparticles.\nFigure R1-8 (a) and (b) present the XRD patterns of the samples exposed to water and\nair, respectively.\nTo further clarify the point, we added more descriptions as follows.\n----------------------------------------------------\nPage 6: Figure 2c shows the XRD pattern of water-exposed uncoated and 0.1 wt.%\nPMMA coated Mg Ni Y NPs. While the uncoated samples exhibit major Mg(OH)\n87.5 5.5 7 2\npeaks, the corresponding peaks are much suppressed for the 0.1 wt.% PMMA-coated\nsamples. To further confirm the surface protection effect of PMMA, we exposed the\nuncoated and 0.1 wt.% PMMA-coated Mg Ni Y NPs under an ambient atmosphere\n87.5 5.5 7\nfor two weeks (Figure 2d). Again, a notable Mg(OH) peak was confirmed for the\n2\nuncoated samples, due to the reaction with air moisture. Meanwhile, the XRD is largely\nunchanged for 0.1 wt.% PMMA-coated Mg Ni Y NPs, and both Mg(OH) and\n87.5 5.5 7 2\nMgO peaks were not identified. These data demonstrated improved surface stability by\nthe PMMA coating.\n----------------------------------------------------\n7. The authors did not provide information about the sample activation process. I\nsuggest the authors to compare the activation and hydrogen storage kinetics of Mg-Ni-\nY repared using solvent-mediated and solvent-free processes.\nResponse: Thank you for the comment. According to this suggestion, we have\nsupplemented the activation data and the corresponding isothermal hydrogen\nde/hydrogenation kinetics data for the Mg-Ni-Y alloy prepared by both the solvent-\nmediated and the solvent-free methods.\nFigure R1-9a, b displays the activation data of Mg-Ni-Y alloy prepared by THF-\nassisted ball milling and dry ball milling over five cycles at 300 °C, respectively. It was\nidentified that the THF-assisted ball milling achieves nanoparticle refinement via\nsolvent mediation. The reduced particle size significantly enhances hydrogenation and\ndehydrogenation kinetics by increasing specific surface area and decreasing hydride\nlayer thickness, facilitating the activation process at lower temperatures. The alloy\nreleases 3.94 wt.% H during the first activation cycle and almost completed activation\n2\nprocess within 5 cycles. In contrast, the dry-ball-milled Mg-Ni-Y alloy fails to complete\nactivation at 300 °C, releasing only 1.05 wt.% H after five cycles.\nball-milled Mg Ni Y alloy after five activation cycles. Owing to its large mean\n87.5 5.5 7\nparticle size (~44.89 μm), the alloy exhibits compromised activation capability,\nresulting in significantly deteriorating kinetics and suppressed hydrogen desorption\ncapacity. At 300 °C, it releases merely 1.53 wt.% H within 15 min. In contrast, the\n2\nsame temperature after activation, demonstrating 3.37 times higher capacity than the\ncounter experiment.\nFigure R1-9 Performance curves of Mg Ni Y alloy: (a, b) Activation curves at\n87.5 5.5 7\n300 °C for THF-assisted and dry ball-milled samples, respectively; (c, d)\nCorresponding isothermal dehydrogenation kinetics curves after activation.\nTo further clarify the point, we added more descriptions as follows.\n----------------------------------------------------\nPage 7: Prior to the de/hydrogenation kinetics analysis, all of samples were activated\nutilizing hydrogenation and dehydrogenation cycles. While Mg Ni Y NPs (THF-\n87.5 5.5 7\nassisted ball-milled samples) are easily activated, the dry-milled samples cannot be\nactivated under the same conditions (Supplementary Figure S6).\n----------------------------------------------------\n8. While I appreciate the effect of polymer coating on Mg-Ni-Y nanoparticles, it should\nbe noted that similar approaches have already been reported. Please adequately cite the\nrelevant literature and discuss the similarities and differences.\nResponse: We appreciate the reviewer's recognition of the polymer coating effects and\nthank you for emphasizing the significance of literature comparison. We have\nconducted a comprehensive review of relevant studies and included these literatures (Li\net al. Int. J. Hydrogen Energy 2023, Yuan et al. Int. J. Hydrogen Energy 2017, Yao et\nal. Prog. Nat. Sci. Mater. Int. 2018, Shen et al. Int. J. Hydrogen Energy 2018, Hashimoto\net al. Mech. Eng. J. 2016) in the main text with associated discussion.\n9. Please consider optimizing Figures 6b and 6c. While Figure 6c shows opposite\ncharge states between Mg1 and Mg2/3/4/5/6, such electron distributions are difficult to\nconfirm in Figure 6b.\nResponse: Thank you for the comment. According to this suggestion, we have\noptimized the color scale range to enhance charge contrast (Figure R1-10). The revised\nfigures now distinctly show that the Mg atom adsorbed at the THF terminal exhibits a\npositive charge state (blue), while adjacent Mg sites display weakly negative charge\nstates (red), consistent with the charge distribution.\nFigure R1-10 Bader charge analysis and charge density difference analysis for the on-\ntop adsorption site and Y-doped configuration.\n10.There are some issues with the wording and spelling in the manuscript and\nsupprementary information. For example, the “kinetic” should be “kinetics”. The\n“temperatur” should be “temperature” in Figure S2.\nResponse: Thank you for your careful reading. We corrected these typos in the\nrevised submission.\nReviewer #2\nIn this report, the authors present a nanosized Mg-Ni-Y alloy synthesized via a\nsolvent-mediated milling process, demonstrating enhanced hydrogen storage\nperformance. Compared to dry milling, THF-mediated milling yields smaller Mg-based\nalloy particles with improved hydrogen absorption and desorption properties. However,\nsolvent-mediated milling is a well-established method for particle size reduction (e.g.,\ndoi.org/10.1002/er.7046, doi.org/10.1007/s12598-025-03330-x,\ndoi.org/10.1039/C3TA01332D, etc.), and the performance reported here is not\nparticularly exceptional. Although the authors include a comparative chart in Fig. 3h,\nthere are numerous previous studies that report comparable or even superior\nperformance. In addition, the authors claim that THF-adsorption leads to the charge\nredistribution on Mg surface, promoting the nanosize effect; however, the data provided\nare insufficient to substantiate this claim.\nWhile it is not necessary for a study to demonstrate record-breaking results, in the\nabsence of a novel synthetic strategy or innovative structural design, I do not believe\nthis work meets the standards for publication in Nature Communications.\nResponse: We appreciate the referee for carefully reviewing the manuscript and\nproviding useful criticisms. We have gone over the report and understand that the major\nconcerns of the referee#2 mainly arises from synthetic methods and hydrogen storage\nproperties.\nFor the first issue, we would like to emphasize that the major finding of this work\nlies in elucidating the nano-size mechanism of solvent-assisted ball milling. Despite\nthe wide recognition of the metallurgy and nano-material community on the effect\nof some solvents during ball-milling, the microscopic mechanism of solvent has\nremained unsolved. Even many of the literatures (Banrejee et al. Int. J. Energy Res.\n2021, Gao et al. Rare Metals 2025, Cui et al. J. Mater. Chem. A 2013), as noted by the\nreferee, reported the effectiveness of adding solution in the ball-milling, to our best of\nknowledge, no literature has reported why such solvent can improve mechanical\nnanosizing effect during ball-milling on a microscopic scale. Is polar solvent\nfavorable, or non-polar solvent provide better results? This knowledge gap has limited\nmaterial development efficiency.\nIn this submission, we discovered that the electron re-distribution caused by the\nsolvent (THF in our study) adsorption is the key factor that modifies surface mechanical\nproperties of Mg and its alloys. As shown in the Figure R2-1a, the THF adsorption on\nMg lattice gives rise to the Mgδ+-Mgδ− dipole structure, and such ionic nature should\nenhance surface mechanical properties. The enhanced surface mechanical properties\nare indeed evinced utilizing Vickers hardness test. As shown in Figure R2-1b, the\nnotable enhances Vickers hardness was identified by THF adsorption. The THF-\nsprayed surface achieved an average microhardness of 103.13 HV, representing 22 %\nenhancement compared to the untreated surface (84.58 HV). These results indicate that\nthe Mg surface becomes less ductile and thus more susceptible to mechanical impact,\nleading to enhanced ball-milling efficiency. These findings provide a microscopic\ninsight into optimizing ball-milling parameters, and explain why polar solvents, such\nas THF, exhibit much higher ball-milling efficiency than that of non-polar solvents such\nas cyclohexane. We believe the discovery provides a contribution to material design\nand would be of potential interest to both metallurgy and nano-material community.\nFigure R2-1 (a) Surface morphology images of the untreated Mg-Ni-Y sample and the\nTHF-sprayed Mg-Ni-Y sample after microhardness testing, respectively. (b)\nMicrohardness distribution map.\nWe are aware that the referee is also concerned by the nano-sizing efficiency by\nciting the literatures like Banrejee et al. Int. J. Energy Res. 2021, Gao et al. Rare Metals\n2025, Cui et al. J. Mater. Chem. A 2013. However, we found that these literatures\nmainly report the ball-milling efficiency of MgH , rather than Mg or Mg-based alloys.\nmechanical property has posed a significant challenge to downsize its particle size\nutilizing ball-milling method. In contrast, MgH is a ceramic material, and thus it is\n2\nmuch easier to prepare nanoparticles utilizing mechanical impact. Therefore, we\nbelieve the reported smaller particle size of MgH does not degrade the impact of our\n2\nfindings.\nFor the second raised concern related to the hydrogen storage properties, we\nbelieve the referee mentions the hydrogen storage property of MgH system. We would\n2\nlike to emphasize that the Mg-Ni-Y system and ball-milling technique were chosen\nbecause of its value for large-scale application. While some MgH , either catalyst added\n2\nor nanosized system, exhibit high hydrogen storage capability, the price of MgH (~800\ndifficulties of MgH production, that need several de/hydrogenation and continuous\n2\nball-milling process, major concerns were posed for large-scale production of MgH .\nMg-based alloy systems so that we can keep the comparison criteria.\nWhen we had gone over the referee report, we noticed our shortage in our\npresentation and understood where we need revision to further clarify our novelty. We\nwould like to thank the referee for providing the useful criticisms. To further clarify the\npoint, we added more descriptions as follows.\n----------------------------------------------------\nPage 3: In contrast, mechanical ball milling offers a simpler and cost-effective\nalternative, and a variety of ball-milling studies on both Mg-based compounds, such\nMgH , have been reported.34-39 However, unlike the brittle ceramic MgH , the high\n2 2\nductility of metallic Mg fundamentally hinders the applicability due to severe\nagglomeration and cold welding.40, 41\nPage 4: Figure 1a illustrates the powdering process achieved through mechanical\nimpact, such as ball milling, of metallic and ionic compounds. Metallic materials are\ncharacterized by delocalized electrons and structural deformation contributes little to\nthe metallic bond, resulting in ductile properties. Meanwhile, ionic materials are\ncharacterized by the presence of positively and negatively charged ions. The slight\nmodification of crystal deformation immediately gives rise to Coulombic repulsion\nbetween identically charged ions and generates cracks. Therefore, it is widely believed\nthat ionic materials are more susceptible to mechanical impact and are easily nanosized\nby ball-milling. Mg Ni Y is one of the Mg-based alloys that is characterized by\n87.5 5.5 7\nhighly ductile properties.\nPage 14: Such partial ionicity should suppress the ductility of Mg metal and maximize\nthe ball milling-driven structural deformation that facilitates ion migration. The\nenhanced surface mechanical properties are indeed demonstrated utilizing Vickers\nhardness measurement. As shown in Figure 6c and d, the notable enhancement of\nVickers hardness was identified for the THF-adsorbed sample. It achieved an average\nhardness of 103.13 HV, representing 22 % enhancement compared to the untreated\nsurface (84.58 HV) (Supplementary Table S5). These results indicate that the Mg\nsurface becomes less ductile and thus more susceptible to mechanical impact, leading\nto enhanced ball-milling efficiency. On the basis of the above-mentioned results, we\npropose the mechanism of solvent-assisted mechanical nanosizing processes in Figure\n6e. First, solvent molecules, such as THF, adsorb on metal surface. It gives rise to\nelectron distribution that enhances surface hardness, making the metal powder more\nsusceptible to mechanical impact. Subsequently, the cracks are easier to generate, and\nthe solvent molecules can be adsorbed on the fresh surface, further enabling mechanical\nnanosizing.\n----------------------------------------------------\n1. Although the authors claim that solvent-free milling results in only a moderate size\nreduction, other reports have demonstrated a significantly greater decrease in particle\nsize. (ex. doi.org/10.1016/j.ijhydene.2020.09.251,\ndoi.org/10.1016/j.jallcom.2009.11.154, https://doi.org/10.1016/j.jallcom.2004.08.102).\nResponse: Thank you for raising this issue. As reported by the referee, numerous\nstudies (Banrejee et al. Int. J. Energy Res. 2021, Gao et al. Rare Metals 2025, Cui et al.\nJ. Mater. Chem. A 2013, Porcheddu et al. Int. J. Hydrogen Energy 2021, Paik et al. J.\nAlloys Compd. 2010, and Schimmel et al. J. Alloys Compd. 2005) reported the\neffectiveness of ball-milling for Mg-based hydrogen storage materials. Meanwhile, we\nare aware that most of the suggested literature is studying the ball-milling effect on\nMgH , rather than Mg or Mg-based alloys. As is widely known, MgH is a brittle\n2 2\nceramic material, and ball-milling is effective to downsize MgH , which is well\n2\ndemonstrated in the suggested literatures. Meanwhile, the Mg and Mg-based alloys,\nwhich have higher large-scale applicability and economic value, suffer from their\nhigh ductile nature during the ball-milling. As shown in Figure R2-2, the\nconventional ball-milling on the Mg-Ni-Y alloy resulted in significant cold-welding dry\nball milling, resulting in adhesion of the alloy powder to the grinding balls or jar walls,\nwhich not only hindered refinement of the final particle size to the nanoscale, but also\ninevitably led to re-agglomeration into coarse particles even when nanoscale\ndimensions were achieved.\nFigure R2-2 (a) and (b) Cold welding phenomenon after dry milling in ball mill jars\nwith stainless steel and zirconia inner linings, respectively. (c) The SEM image of cold-\nwelded particles after dry ball milling and the corresponding particle size distribution.\nMeanwhile, the particle size is significantly suppressed when THF was added\nduring the ball-milling. The nanoparticles prepared by the THF-assisted ball milling\nmethod yielded an average particle size of 511.87 nm, achieving a 88-fold reduction\ncompared to the solvent-free process, while the measured specific surface area reached\n3.72 m2·g-1, representing a 41-fold improvement over the solvent-free method. We note\nthat not all solvents result in the situation, when cyclohexane was chosen, clear\nnanosizing effect was not identified.\nFigure R2-3 (a) Calculated adsorption energies of THF at four different sites on the\nMg(0001) surface. (b) Bader charge analysis and charge density difference analysis (c)\nSurface morphology images of the untreated Mg-Ni-Y sample and the THF-sprayed\nMg-Ni-Y sample after microhardness testing, respectively. (d) Microhardness\ndistribution map.\nTo unveil the microscopic mechanism lying behind the phenomenon, we\nconducted DFT calculations. As shown in Figure R2-3, it was found that upon THF\nadsorption, the magnesium atom coordinating with the oxygen loses 0.56 electrons\n(manifested as positive charge states in blue), while adjacent magnesium atoms gain\n0.03-0.08 electrons, resulting in the formation of Mgδ+-Mgδ- dipole moments.\nFurthermore, charge redistribution in proximal magnesium atoms is induced by Y-\ndoping. (Figure R2-3b). Furthermore, Vicker hardness test demonstrated the\nmodification of surface mechanical properties (Figure R2-3d). The THF-sprayed\nsurface achieved an average microhardness of 103.13 HV, representing 22 %\nenhancement compared to the untreated surface (84.58 HV). It is the consequence of\nthe rise of partial ionicity by the THF adsorption. The enhanced surface hardness leads\nto the Mg-Ni-Y alloy being more susceptible to mechanical impact, and the material\neasily create cracks during the ball-milling. The THF molecules are then adsorbed on\nthe new surface, giving rise to the continuous nanosizing process. As the results the\nTHF-assisted ball-milling realized 88-fold enhancement compared with the dry-milling\ncondition.\nFigure R2-4 (a) Calculated adsorption energies of cyclohexane at four sites on the\nMg(0001) surface (b) Bader charge analysis for the on-top adsorption site.\nWe would like to note that not all solvents resulted in the similar nanosizing effect.\nWe conducted cyclohexane-assisted ball-milled Mg-Ni-Y alloy samples using identical\nprocessing parameters for comparison. Cyclohexane was selected due to its lower\npolarity and its inert characteristics towards Mg and its alloys. Figure R1-3 illustrates\nthe SEM-EDS images and particle size distribution of the cyclohexane-assisted ball-\nmilled sample, where the average particle size is 40.20 μm. This shows no significant\nrefinement compared to the dry-milled alloy (44.89 μm). It is because cyclohexane\nexhibits much weaker interaction with Mg and Mg-based alloy, and it is unlikely to\nrealize ionic bond between Mg (Figure R2-4).\nTo further clarify the point, we added more descriptions as follows.\n----------------------------------------------------\nPage 13: Subsequently, we discuss the impact of cyclohexane adsorption on Mg-based\nalloy. Unlike THF, cyclohexane is characterized by the less polarized molecules and\nexhibits much weaker interaction (Supplementary Figure S22 and 23). As a result,\nadsorption of cyclohexane on Mg-alloy does not cause drastic charge redistribution as\nwas observed for THF, and this is why cyclohexane does not contribute to mechanical\nnanosizing efficiency (Supplementary Figure S4).\n----------------------------------------------------\n2. For the particle size characterization, additional high-resolution SEM images are\nneeded, as Figure 1a lacks sufficient clarity to support accurate size estimation. While\nthe mean particle diameter is reported as 813 nm, the resolution of the provided image\nis too low to justify this calculation. Moreover, the standard deviation should also be\nreported. Based on Figure 1a, the sample appears to exhibit a broad size distribution,\nwith many larger particles clearly visible—contrary to the histogram presented in the\nimage. In particular, the claim that the particles are \"monodisperse\" (line 95) is not\nsupported by the data and should be reconsidered.\nResponse: Thank you for your valuable comments. Your feedback regarding\nFigure 1a and the particle size characterization is highly pertinent and has been crucial\nfor improving the quality of our manuscript. In accordance with your suggestions, we\nhave implemented the following revisions and additions.\nFirst, we have reacquired and supplemented clearer SEM images of the THF-\nassisted ball-milled Mg–Ni–Y alloy, ranging from low to high magnifications (Figure\nR2-5). These images more clearly reveal the morphology and boundaries of individual\nparticles, thereby providing a reliable basis for accurate particle size measurement. This\nfigure has replaced the previous version in the manuscript.\nFigure R2-5 (a), (b), (c), and (d) present SEM images of the THF-assisted ball-milled\nMg–Ni–Y alloy particles at progressively higher magnifications, with scale bars\ncorresponding to 20 μm, 10 μm, 5 μm, and 3 μm, respectively.\nAdditionally, we have included new high-resolution images illustrating the Mg–\nNi–Y alloys processed by dry ball milling, cyclohexane-assisted ball milling, and THF-\nassisted ball milling (Figure R2-6). Statistical measurements were performed on 150\nparticles from the THF-assisted ball-milled alloy using ImageJ software to ensure\naccuracy and to provide standard deviation. The particle size distribution histogram has\nbeen updated accordingly. The results demonstrate that the average particle size of the\nTHF-assisted ball-milled alloy is 511.87 ± 328.52 nm, which is significantly smaller\nthan those obtained by dry ball milling (44.89 ± 11.41 μm) and cyclohexane-assisted\nball milling (40.20 ± 11.23 μm). We also appreciate your comment regarding the term\n“monodisperse”. In the revised submission, we change the term to “the particles exhibit\na uniform distribution”.\nFigure R2-6 (a), (b), and (c) show the SEM images of the dry ball-milled, cyclohexane-\nassisted ball-milled, and THF-assisted ball-milled samples, respectively; while (d), (e),\nand (f) present the corresponding particle size distributions.\n3. Please provide detailed experimental conditions for the dry ball-milling process. In\nball-milling, parameters such as the size and type of milling balls can significantly\ninfluence the resulting particle characteristics.\nResponse: Thank you for raising this issue. All ball milling procedures in this\nstudy were uniformly executed under strictly controlled conditions: 20-hour milling\nunder argon atmosphere (using a Focucy F-P400 mill with 100ml stainless steel-lined\njar), employing 3mm-diameter zirconia balls at a ball-to-powder ratio of 40:1 and\nrotation speed of 400 rpm. To mitigate temperature rise, intermittent 10-minute cooling\nintervals were implemented after every 10 minutes of milling. The equipment\nconfiguration is detailed in Figure R2-7.\nThe detailed parameters were summarized in the Supplementary Table S6.\nFigure R2-7 (a) FOCUCY F-P400 ball mill; (b) Stainless steel milling jar; (c) Zirconia\ngrinding beads.\n4. In the Mg 1s XPS results, only the metallic Mg peak was observed, despite the well-\nknown tendency of Mg to form a thin oxide layer—even under ultra-high vacuum\nconditions—due to its instantaneous reaction with trace amounts of oxygen. It is\nrecommended that the Mg 2p spectrum be provided, as it is more sensitive to detecting\nsurface oxides. Furthermore, the authors claim that THF adsorption induces a change\nin the charge distribution on the Mg surface. If THF adsorption indeed affects the\nsurface chemistry, the presence of at least a partial oxide layer should be detectable to\nsupport this claim.\nResponse: Thank you for your valuable comments on the surface characterization.\nPrior to XPS analysis, in situ cleaning of the Mg surface was performed by Ar+ ion\nsputtering (total depth: 10 nm achieved via 5 sequential etching cycles). This procedure\neffectively eliminates transient oxide layers and ambient adsorbates (e.g.,\nhydrocarbons), exposing the intrinsic metallic surface to ensure the XPS data accurately\nreflect the bulk chemical state of magnesium.\nIn addition, following the reviewer’s suggestion, we have supplemented the Mg\n2p XPS data of the THF-assisted ball-milled Mg–Ni–Y sample without sputtering\ntreatment. Figure R2-8 (a) and (b) show the Mg 2p XPS spectra of the THF-assisted\nball-milled Mg Ni Y alloy before and after fitting, respectively. The binding energy\n87.5 5.5 7\nwas calibrated using the C 1s signal (284.80 eV) to account for inevitable surface\ncarbon contamination.\nA distinct low-binding-energy shoulder can be clearly observed on the right side\nof the main peak in Figure (a), as indicated by an arrow. Fitting analysis in Figure (b)\nreveals that the Mg 2p spectrum consists of three components: the central peak at 49.47\neV is attributed to metallic Mg0, consistent with previous studies (Haider et al. Z.\nNaturforsch. A 1975, Jérôme et al. Appl. Surf. Sci. 1986); the low-binding-energy peak\nat 48.60 eV corresponds to an electron-rich Mgδ– species resulting from THF\nadsorption-induced charge redistribution on the Mg substrate; whereas the high-\nbinding-energy peak at 50.14 eV is significantly higher than that of metallic Mg0 but\nlower than that of MgO (typically around 51 eV, Hoogewijs et al. J. Electron Spectrosc.\nRelat. Phenom. 1977, Inoue et al. Bull. Chem. Soc. Jpn. 1981), which is ascribed to an\nelectron-deficient Mgδ+ species generated by the same THF-induced charge\nredistribution mechanism.\nIt should be emphasized that although THF adsorption on the Mg substrate does\ninduce charge redistribution, the effect is not strong enough to form oxides. According\nto reference (Han et al. ACS Omega 2025), the adsorption energy of molecularly\nadsorbed O on Mg is as high as 8.08 eV, and Bader charge analysis indicates that\n2\ndissociated oxygen atoms (each carrying –1.58 e) bear more negative charge than\nmolecularly adsorbed O₂ (total charge –1.70 e). Both the adsorption energy and the\nextent of charge transfer significantly exceed those of THF adsorption on Mg\n(adsorption energy: 0.48 eV; induced charge transfer at the Mg adsorption site: 0.56 e).\nFigure R2-8 (a) and (b) show the XPS spectra of the Mg Ni Y alloy processed by\n87.5 5.5 7\nTHF-assisted ball milling, before and after spectral fitting, respectively.\nTo further clarify the point, we added more descriptions as follows.\n----------------------------------------------------\nPage 13: Indeed, XPS analysis on the as milled samples show the presence of both\npositively and negatively charged Mg, consistent with the DFT calculations\n(Supplementary Figure S21). The synergistic interplay between THF adsorption and\nY-doping induces localized charge redistribution in adjacent Mg atoms, giving rise to\nthe partial ionicity between Mg atoms.\n----------------------------------------------------\n5. PCT isotherms, along with the corresponding thermodynamic parameters, should be\nmeasured to properly characterize the hydrogen storage properties. This type of analysis\nis essential for studies involving metal hydride-based hydrogen storage materials.\nResponse:Thank you for raising this issue. Generally, reducing the stability of\nmagnesium-based hydrides appropriately can enhance the performance of magnesium-\nbased hydrogen storage alloys, which is crucial for advancing their practical\napplications. The PCT curves of Mg Ni Y and 0.1PMMA-Mg Ni Y alloys\n87.5 5.5 7 87.5 5.5 7\nwere measured at 280, 300, 320, and 340 °C. The enthalpy changes (ΔH) during the\nreactions were calculated using the van't Hoff equation, and the results are presented in\nFigure R2-9.\nFigure R2-9 PCT curves (a, c) of the activated Mg Ni Y and 0.1PMMA-\n87.5 5.5 7\nMg Ni Y samples at different temperatures and the corresponding Van't Hoff plots\n87.5 5.5 7\n(b, d).\nAnalysis of the PCT curves in Figure R2-9(a), (c) reveals two distinct plateaus\nduring both hydrogen absorption and desorption for all samples. The high, short plateau\ncorresponds to the formation/decomposition of Mg Ni/Mg NiH , while the low, long\n2 2 4\nplateau corresponds to the formation/decomposition of Mg/MgH . As revealed in\nabs\n(ΔH ) for the Mg Ni Y and 0.1PMMA-Mg Ni Y alloys are ‒68.5 ± 6.7\ndes 87.5 5.5 7 87.5 5.5 7\nkJ/mol and 83.1 ± 8.1 kJ/mol, ‒70.6 ± 7.1 kJ/mol and 82.8 ± 4.6 kJ/mol, respectively.\nThe two alloys exhibit comparable thermodynamic properties, indicating that the\nreaction enthalpy is not notable in this system. It is because only ~5 nm scale\nnanoparticle can impact thermodynamic property of Mg-based hydrogen storage\nmaterials, consistent with the previous theoretical predictions (Wagemans et al. J. Am.\nChem. Soc. 2005).\nTo further clarify the point, we added more descriptions as follows.\n----------------------------------------------------\nPage 8: Supplementary Figure S12 shows the equilibrium pressures of uncoated and\n0.1-PMMA@Mg Ni Y . The reaction enthalpy obtained from von’t Hoff equation\n87.5 5.5 7\nshows the nanosizing effect affect little to the thermodynamic properties.\n----------------------------------------------------"}, {"round": 2, "reviewer_comments": {"reviewer_1": "This manuscript explores the enhancement of the ball milling effect of magnesium-based hydrogen storage materials by adding PMMA during mechanical ball milling, and presents the mechanism for this improvement. The authors made detailed revisions based on the reviewers' comments, added the necessary experiments, and provided detailed data. I suggest that it be published in the journal of Nature Communications.", "reviewer_2": "The comments are mostly well-addressed. Now I recommend the publication. One minor correction is required: “Page 8: Supplementary Figure S12 shows the equilibrium pressures of uncoated and 0.1-PMMA@Mg87.5Ni5.5Y7. The reaction enthalpy obtained from von’t Hoff equation shows the nanosizing effect affect little to the thermodynamic properties.”. It should be “van’t”, not “von’t”."}, "author_response": "Response to the reviewer\nReviewer #1\nThis manuscript explores the enhancement of the ball milling effect of\nmagnesium-based hydrogen storage materials by adding PMMA during mechanical\nball milling, and presents the mechanism for this improvement. The authors made\ndetailed revisions based on the reviewers' comments, added the necessary experiments,\nand provided detailed data. I suggest that it be published in the journal of Nature\nCommunications.\nResponse: We appreciate the referee for carefully reviewing the manuscript and\nexcellently summarizing our contribution. The referee’s criticisms helped us a lot to\nimprove our revised submission. We highly appreciate the referee for the positive\njudgement and useful comments provided.\nReviewer #2\nThe comments are mostly well-addressed. Now I recommend the publication. One\nminor correction is required: “Page 8: Supplementary Figure S12 shows the equilibrium\npressures of uncoated and 0.1-PMMA@Mg87.5Ni5.5Y7. The reaction enthalpy\nobtained from von’t Hoff equation shows the nanosizing effect affect little to the\nthermodynamic properties.”. It should be “van’t”, not “von’t”.\nResponse: We appreciate the referee for carefully reviewing the manuscript and\nexcellently summarizing our contribution. The referee’s criticisms helped us a lot to\nimprove our revised submission. We highly appreciate the referee for the positive\njudgement and useful comments provided.\nWe appreciate the referee for pointing out the typo. We carefully checked the\ngrammar and any formatting errors during the revision. The “von’t Hoff” was corrected\nto “van’t Hoff”."}]

Decision letter