Engineering non-interfacial hydrogen spillover in a Ni_17W_3-WO_2 heterostructure

Domain: biology
Published: 2026-03-23
Editorial outcome: major_revision

Abstract
The hydrogen spillover effect offers a promising strategy to overcome the kinetic bottleneck of proton desorption in hydrogen evolution reaction catalysts. However, conventional hydrogen spillover mechanisms rely on interfacial proton transfer between distinct phases and suffer from inherent energy barriers. Here, we show a non-interfacial hydrogen spillover mechanism in a Ni17W3–WO2 heterostructure, engineered through the synergistic creation of a built-in strain gradient and directional electron transfer. This design spatially confines the complete hydrogen evolution process within the Ni17W3 phase, thereby circumventing cross-phase migration and reshaping the hydrogen adsorption energy landscape. Experimental and theoretical analyses confirm the elimination of interfacial barriers and establishment of an optimized proton-migration route. The resulting catalyst achieves a low overpotential of 21 mV at 10 mA cm–2 in 0.5 M H2SO4, along with sustained stability (>1500 hours at 500 mA cm–2) and a Faradaic efficiency of 98.6%. This work demonstrates how tailored heterostructures can bypass interfacial bottlenecks, providing guidance for developing efficient non-precious hydrogen spillover catalysts and advancing sustainable hydrogen production.

Introduction
The urgent transition from fossil fuels to sustainable energy systems demands scalable technologies for green hydrogen production, a cornerstone of the future carbon-neutral economy1-3. Acidic electrochemical hydrogen evolution reaction (HER), with its rapid kinetics and compatibility with renewable energy sources like wind and solar, stands out as a key solution. In practice, proton exchange membrane (PEM) electrolyzers, leveraging acidic HER technology, are attractive for industrial-scale hydrogen production due to their proton concentration being orders of magnitude higher than that in alkaline/neutral media4-6. However, its widespread adoption hinges on replacing expensive Pt-group catalysts with earth-abundant alternatives, which is a challenge exacerbated by the intrinsic trade-off between proton adsorption and desorption in non-noble metal catalysts7-8. Such materials readily adsorb protons (Volmer step), but their high hydrogen-binding energy stifles H2 release (Heyrovsky/Tafel steps), resulting in unsatisfactory activity9,10.

Recent advances propose hydrogen spillover (H-spillover) as a remedy, where protons adsorbed on active sites migrate to secondary phases for facilitated desorption. For instance, Wei et al. constructed a V-rich Pt/TiO2 heterostructure, where protons adsorbed on Pt sites and then spillover to TiO2, enhancing desorption kinetics and achieving an overpotential of approximately 60 mV at a current density of 10 mA cm–2, with 24-hour stability11. Similarly, Li et al. demonstrated that Fe-doped Ru/CoP heterostructures reduce interfacial electronic barriers, promoting proton transfer and improving HER activity12. These cases reveal a fundamental dilemma that conventional H-spillover relies on interfacial proton transfer, which unavoidably introduces electronic barriers that hinder kinetics and structural stress that compromises durability. However, existing demonstrations of non-interfacial H-spillover, such as in multi-element single-phase structures (e.g., La2Sr2PtO7+δ and high-entropy oxysulfides)13-14, face inherent limitations. Their reliance on noble-metal dependency or highly complex compositions, coupled with mechanistic ambiguity due to disordered elemental arrangements, hinders the derivation of a general design principle. The rational design of efficient H-spillover pathways requires precise control over the hydrogen adsorption energy landscape across multiple sites. Strain and electronic structure engineering have emerged as powerful levers for such tuning15. Tungsten oxides (particularly WO2) present advantages as heterostructure components due to their structural adaptability and strong propensity for electronic coupling. By constructing an intimate heterointerface, WO2 can act as a strategic modulator to impose a controllable strain field and drive directional electron transfer to an adjacent active phase16. This synergistic modulation is envisioned to create the internal driving force necessary for steering proton migration along a non-interfacial, intra-phase pathway.

Herein, we report a design strategy to bypass interfacial bottlenecks in H-spillover catalysis (Fig. 1). Through in situ phase reconstruction, we construct a Ni17W3–WO2 heterostructure where the WO2 phase is deliberately employed not as a primary catalytic site, but as an essential engineering component. Its critical function is to induce a built-in strain gradient and facilitate electron transfer within the contiguous Ni17W3 phase. This concerted action collectively reconfigures the local hydrogen adsorption energetics, thereby constructing an efficient, non-interfacial proton spillover channel confined entirely within the Ni17W3 phase. This non-interfacial H-spillover mechanism, validated by a combination of theoretical and experimental analyses, empowers the catalyst with comparable performance, achieving an overpotential of 21 mV at 10 mA cm–2 and long-term stability for 1500 hours at an industrial-grade current density of 500 mA cm–2. This work establishes a paradigm of decoupling multi-step reactions within a single phase via interfacial engineering, providing a generalizable platform for the design of advanced electrocatalysts.

Results
Material characterizations and interfacial charge distribution simulation. The Ni17W3–WO2 heterostructure, comprising WO2 and Ni17W3 alloy, was synthesized via in situ phase reconstruction through annealing of NiWO4 precursor in N2/H2 atmosphere (Fig. 2a). Field emission scanning electron microscopy (FE-SEM) characterization reveals the morphological evolution from precursor to final product. The NiWO4 precursor comprises uniform nanoparticles with an average size of 30.3 nm (Supplementary Fig. 1). Following annealing in a N2/H2 atmosphere, the nanoparticulate morphology is largely retained, although the average particle size increases to 38.9 nm due to partial agglomeration during the thermal process (Fig. 2b-c and Supplementary Fig. 2). For reference, the individual components were also characterized. Pure WO2 exhibits a rod-like morphology with distinct surface striations (Supplementary Fig. 3), whereas Ni17W3 forms as irregular, sintered nanoparticles (Supplementary Fig. 4).

The crystal structure was studied by X-ray diffraction (XRD). The XRD pattern of the precursor can be indexed to NiWO4 (JCPDS#15-0755, Supplementary Fig. 5). Additionally, the Raman spectrum for the precursor is consistent with NiWO4 (Supplementary Fig. 6). The diffraction peaks (purple line) of the final product using (NH4)6H2W12O40·xH2O precursor can be ascribed to monoclinic WO2 (JCPDS#32-1393, Fig. 2d). While using NiSO4·6H2O and (NH4)6H2W12O40·xH2O as precursors, the diffraction peaks (cyan line) of the final product can be indexed to the face-centered cubic Ni17W3 (JCPDS#65-4828). The XRD pattern (pink line) of the product obtained by annealing the NiWO4 precursor indicates the formation of a composite of Ni17W3 phase and WO2 phase.

The high-angle annular dark field scanning transmission electron microscopy (HAADF-STEM) images of Ni17W3–WO2 reveal that it exhibits a nanoparticle morphology consistent with SEM images (Fig. 2e). The elemental maps indicate the uniform distribution of Ni, W, and O within the Ni17W3–WO2 (Fig. 2f). The atomic resolution HAADF-STEM image of Ni17W3–WO2 shows lattice fringes with spacings of d=0.35 nm and d=0.25 nm, which are attributed to the (011) and (020) planes of the WO2 phase (Fig. 2g, Supplementary Fig. 7), respectively. The angle between WO2 (011) and WO2 (020) planes is 45°, matching the standard crystal structure of WO2. While the observed lattice fringe d=0.24 nm corresponds to the (111) planes of the Ni17W3 phase (Supplementary Fig. 8). In addition, a distinct interface is observed between Ni17W3 (111) and WO2 (011) planes. The crystallographic structure of the heterointerface was unequivocally confirmed by Fast Fourier Transform analysis (Fig. 2h). The distinct diffraction spot pattern can be indexed to the superposition of the [011] zone axis of the face-centered cubic Ni17W3 phase and the [001] zone axis of the monoclinic WO2 phase. These results provide definitive evidence for the coexistence of Ni17W3 and WO2, confirming the formation of a well-defined heterointerface between them.

Additionally, geometric phase analysis (GPA) was conducted on the heterointerface area (Fig. 2i). The strain distribution along the xy direction reveals distinct strain regions within the Ni17W3 phase but negligible strain variation in the WO2 phase. The strain distributions along the xx and yy directions exhibit strain regions consistent with those observed along the xy direction (Supplementary Fig. 9). Quantitative strain analysis along the trajectory in Fig. 2i reveals a pronounced gradient within the Ni17W3 phase, with values ranging from –0.2% to 1.1% and tapering to near zero (–0.05% to 0.03%) at the heterointerface, while the WO2 phase remains minimally strained (–0.05% to 0, Fig. 2j). This profile signifies a substantial release of intrinsic Ni17W3 lattice stress at the interface, induced by the lattice misfit with WO2. The well-defined Ni17W3–WO2 heterointerface-induced strain gradient is posited to establish a continuous modulation of hydrogen adsorption energy across the Ni17W3 phase, thereby facilitating directional proton spillover and enhancing the hydrogen evolution kinetics18.

Furthermore, the XRD refinement pattern of Ni17W3–WO2 shows that the molar fractions of Ni17W3 and WO2 are 17.7% and 82.3%, respectively (Fig. 3a, b). The X-ray photoelectron spectroscopy (XPS) survey spectra show that W elements can be observed in WO2, Ni17W3, and Ni17W3–WO2, while Ni is only found in Ni17W3 and Ni17W3–WO2 (Fig. 3c). As shown in Fig. 3d, the high-resolution XPS Ni-2p spectra of Ni17W3 and Ni17W3–WO2 indicate characteristic peaks at 856.1/874.0 eV and 858.6/876.2 eV correspond to divalent Ni2+ and Ni3+, resulting from surface oxidation19-21. The peaks at 862.5/880.6 eV are attributed to the satellite peaks22. In addition, the peaks at 852.9 and 870.3 eV for Ni17W3–WO2 are attributed to metallic state Ni0 23, which has shifted by 0.2 eV to higher binding energy compared to the pure Ni17W3. The high-resolution XPS W-4f spectra of Ni17W3, Ni17W3–WO2, and WO2 were analyzed to investigate the chemical states of W (Fig. 3e). The peaks at 31.3 and 33.5 eV correspond to W0 in the Ni17W3 phase. The peaks at 32.6/34.4 eV are attributed to W4+ in the WO2 phase. Peaks at 35.0/37.3 eV and 35.9/37.9 eV are assigned to W5+ and W6+ originating from surface oxidation24-25. The peak at 40.5 eV corresponds to the W-5p orbital. It is worth noting that the binding energies corresponding to W0 in Ni17W3–WO2 are consistent with those in Ni17W3. However, the peaks corresponding to W4+ in Ni17W3–WO2 shifted by 0.2 eV to lower binding energy compared to those in the pure WO2 phase. The peak shift of Ni0 and W4+ in Ni17W3–WO2 compared to the pure phases indicates that electron redistribution occurs at the Ni17W3–WO2 heterointerface. These results confirm the electronic interaction between the metallic Ni and WO2 at the heterointerface, where electrons transfer from Ni in the Ni17W3 phase to W in the WO2 phase.

Density functional theory (DFT) calculations were used to further study the charge redistribution of the heterointerface. The work functions of Ni17W3 and WO2 were calculated independently to clarify the electron transfer path. The work function of pure Ni17W3 is 3.92 eV, while 4.78 eV for pure WO2 (Fig. 3f, Supplementary Fig. 10). Therefore, electrons will transfer from the Ni17W3 phase to the WO2 phase at the Ni17W3–WO2 heterointerface. The electron density difference of the Ni17W3–WO2 heterointerface reveals that electrons accumulate on the W atoms of the WO2 phase, while Ni17W3 is electron-depleted (Fig. 3g). Additionally, the Mulliken charge value of W atoms in the WO2 side at the heterointerface is significantly lower than that at non-interface regions. The opposite trend is observed for the Ni17W3 phase with a higher Mulliken charge value at the heterointerface compared to the non-interface regions (Fig. 3h). These results suggest the presence of electronic interaction at the heterointerface and electrons moving from the Ni17W3 phase to W atoms in the WO2 phase, which is consistent with the XPS results. The charge redistribution at the heterointerface would alter the hydrogen adsorption properties, thus possibly optimizing the electrocatalytic reaction pathway and enhancing the HER activity26.

Electrocatalytic HER performance evaluation. The linear sweep voltammetry (LSV) curves of Ni17W3, WO2, and Ni17W3–WO2 in 0.5 M H2SO4 obtained on glassy carbon (GC) reveal that the Ni17W3–WO2 exhibits improved HER activity (Fig. 4a), requiring an overpotential of only 74 ± 3 mV to afford a current density of 10 mA cm–2, which is much lower than that of pure Ni17W3 (232 ± 7 mV) and pure WO2 (cannot achieve this current density). The corresponding non-iR corrected polarization curves are shown in Supplementary Fig. 11. The Ni17W3–WO2 also demonstrates notable mass activity compared to pure Ni17W3 and WO2, achieving a mass activity of 0.1 A mg–1 at an overpotential of a mere 200 mV (Supplementary Fig. 12). In addition, the Tafel slope of Ni17W3–WO2 is 48.9 mV dec–1 (Fig. 4b), which is smaller than those of pure Ni17W3 (105.0 mV dec–1) and WO2 (406.6 mV dec–1). This indicates that WO2 is controlled by the Volmer-step (V-step), and both Ni17W3 and Ni17W3–WO2 follow the Volmer-Heyrovsky (V-H) reaction mechanism, where the V-step occurs rapidly, and thus the rate-determining step is H-step27. Furthermore, the NiWO4 precursor exhibits poor activity in HER, achieving a current density of 1 mA cm–2 at an overpotential of 307 ± 9 mV (Supplementary Fig. 13a). It also shows a large Tafel slope of 300.3 mV dec–1 (Supplementary Fig. 13b) and a large charge transfer resistance (Supplementary Fig. 13c). The smaller overpotential and Tafel slope of Ni17W3–WO2 indicate enhanced electrocatalytic activity and accelerated hydrogen evolution kinetics.

HER activity depends on the electrochemically active surface area (ECSA), which can be calculated by the electrical double-layer capacitance (Cdl) in the non-Faradic regions based on the cyclic voltammetry (CV) curves (Supplementary Figs. 14 and 15). Ni17W3–WO2 has an ECSA of 1580 cm2, which is distinctly larger than those of pure phase Ni17W3 (526 cm2) and WO2 (427 cm2). The larger ECSA of Ni17W3–WO2 is attributed to the increased amount of active sites induced by the heterointerface between Ni17W3 and WO2 phases to activate the inert atoms. ECSA-normalized activity measurements reveal that the Ni17W3–WO2 exhibits enhanced intrinsic electrocatalytic activity compared to the pure Ni17W3 and WO2 (Fig. 4c). Furthermore, the Ni17W3–WO2 heterostructure exhibits a markedly higher turnover frequency (TOF) than its individual components at identical overpotentials (Fig. 4d), underscoring its improved intrinsic activity. This enhancement is a direct consequence of the heterointerface, where the synergistic interplay between the induced strain field and interfacial electron redistribution optimizes the hydrogen adsorption energy and collectively lowers the kinetic barrier for the reaction, thereby accelerating the HER.

The Ni17W3–WO2 was supported on carbon cloth (Ni17W3–WO2/CC) to assess its potential for industrial implementation in 0.5 M H2SO4. The electrode achieves low overpotentials of 21 ± 2, 138 ± 4, and 204 ± 5 mV to reach current densities of 10, 500, and 1000 mA cm–2, respectively, coupled with a Tafel slope of 47.8 mV dec−1 (Fig. 4e and Supplementary Fig. 16), indicating both high activity and fast kinetics. The corresponding non-iR corrected polarization curves are shown in Supplementary Fig. 17. A comprehensive benchmark against state-of-the-art catalysts, including commercial Pt/C28, Ru29, Pt/OLC28, NiRuPt/Activated XC-7230, Pt@WS2 31, Pt/NC32, CoNiPt@G33, 2H/1T-WS2-1034, hcp-RuCu HUNSs35, Ru/CNT36, Ru-CoNi37, Ru-Cu-MoO2-48h38, Ru@V-RuO2/C HMS39, Ru-6CDs/G40, Ru/PdCu-7041, PtRu/NMHCS-A42, RuSe2 43, NiSA/rGO-CTAB-MoS2 44, p+-WS2-9.845, Ni/WC46, MoS2-NbS2 47, Pb-CoSe2-DETA48, and NiCoSe2/NCHP49 (Supplementary Table 1), reveals that Ni17W3–WO2/CC is highly competitive with commercial Pt/C and even surpasses many advanced noble-metal-based catalysts. Although its mass activity may trail behind that of ultra-low-loading Pt single-atom catalysts, the entirely earth-abundant composition of Ni17W3–WO2 guarantees a high cost-to-performance ratio, a paramount advantage for scalable hydrogen production.

The hydrogen production performance of Ni17W3–WO2/CC was quantified via the water displacement method (Supplementary Fig. 18). At a constant current of −0.1 A, the system generated H2 at a rate of 1.85 ± 0.02 mmol h–1 with a Faradaic efficiency (FE) of 98.65 ± 0.82%, and gas chromatography confirmed the product was 100% pure H2 (Supplementary Fig. 19a, b). Control experiments on the individual components revealed a critical insight. Both the H2 evolution rate and FE of Ni17W3–WO2/CC were nearly identical to those of pure Ni17W3/CC, and both were higher than the values for WO2/CC (1.82 ± 0.02 mmol h–1, 97.83 ± 0.82%; Supplementary Figs. 20 and 21). This data strongly indicate that the Ni17W3 phase serves as the primary active site for HER, while the slightly lower FE of WO2 is likely due to partial proton intercalation into its structure, a known phenomenon where some protons are stored rather than converted to H2 50. Consequently, the role of the heterointerface is not to generate new active sites on WO2, but to profoundly enhance the intrinsic activity of the Ni17W3 phase through electronic and strain engineering, as evidenced by the drastically improved overpotential.

Furthermore, the long-term durability of Ni17W3–WO2/CC was assessed at an industry-relevant current density of 500 mA cm−2. The catalyst exhibited high stability over 1500 hours of continuous operation, with a negligible average potential decay rate of only 28 µV h−1 (Fig. 4f and Supplementary Fig. 22). Post-test characterization confirmed robust structural integrity. The persistence of distinct diffraction peaks in XRD and the clearly resolved lattice fringes at the interface in TEM (Supplementary Figs. 23 and 24) collectively demonstrate that the core heterostructure remains intact. XPS analysis indicated slight superficial oxidation of the Ni17W3 phase, evidenced by an increase in high-valent W and Ni species and a corresponding attenuation of the metallic Ni0 signal (Supplementary Fig. 25), which likely accounts for the minor activity decay observed during the prolonged test, without compromising the structural integrity of the bulk material. Such performance establishes Ni17W3–WO2/CC as a stable non-noble catalyst compared to the reported literature under demanding acidic conditions (Supplementary Table 2).

To assess its viability in practical devices, we integrated Ni17W3–WO2 as a cathode into a PEM electrolyzer. The catalyst formed a homogeneous ink (Supplementary Fig. 26a), enabling its uniform deposition via brush-coating onto carbon paper to form an integrated electrode (Supplementary Fig. 26b), which is a method amenable to industrial scale-up. In a full-cell configuration with an IrO2 anode (Supplementary Fig. 27), the Ni17W3–WO2||IrO2 electrolyzer delivered a low cell voltage of 1.945 V under room temperature (20 °C) (Fig. 4g, corresponding non-iR corrected polarization curves are shown in Supplementary Fig. 28) and 1.650 V at 80 oC for 1000 mA cm−2 (Supplementary Fig. 29), a performance competitive with the Pt/C||IrO2 benchmark. The device demonstrated stable operation at room temperature (20 °C) and at 80 oC, sustaining performance over 1000 cycles at high current densities (500 and 1000 mA cm−1) with negligible degradation in the polarization curve (Supplementary Fig. 30). These results underscore the direct practical utility of Ni17W3–WO2 as a high-performance, non-precious cathode for industrial hydrogen production.

H-spillover mechanism investigation. CV was conducted to study the proton insertion/extraction process of the catalysts during HER in 0.5 M H2SO4 (Supplementary Fig. 31). Generally, a more positive proton insertion potential signifies stronger catalyst adsorption, while a more negative proton release potential indicates stronger release capability51. The proton insertion potential follows the order: Ni17W3 (0.360 V) > Ni17W3–WO2 (0.319 V) > WO2 (0.305 V) vs. RHE (Fig. 5a), indicating a systematic weakening of hydrogen adsorption strength upon heterostructure formation compared to pristine Ni17W3. More critically, Ni17W3–WO2 exhibits a more negative proton extraction potential (0.363 V) than Ni17W3 (0.377 V), signaling a significantly reduced energy barrier for hydrogen desorption. Due to the weak hydrogen adsorption property of WO2, its proton extraction peak could not be detected 52. This optimized energetics with moderated adsorption coupled with facilitated desorption is the key factor that enhances the HER activity of the Ni17W3–WO2 heterostructure, aligning well with the proposed non-interfacial H-spillover pathway.

Electrochemical impedance spectroscopy (EIS) provides further mechanistic insights into the HER process. The Nyquist plot for WO2 shows a dominant Warburg impedance (Fig. 5b), signifying a rate-limiting proton diffusion process due to its poor adsorption capability, which aligns with the CV data53. In contrast, both Ni17W3 and Ni17W3–WO2 display a capacitive arc, characteristic of a surface-controlled reaction involving H* intermediates54-55. The smaller charge-transfer resistance of the Ni17W3–WO2 heterostructure is consistent with its enhanced HER activity. This is corroborated by the Bode plots, where Ni17W3–WO2 exhibits a lower characteristic frequency (Fig. 5c), indicating a prolonged electron residence time50, 56. This finding is consistent with a multi-step reaction pathway, strongly supporting the involvement of an H-spillover process where proton transfer becomes the kinetic bottleneck57.

To unequivocally establish the H-spillover mechanism, we performed complementary kinetic and chemical probing experiments. The kinetic isotope effect (KIE) was investigated by substituting H2SO4/H2O with D2SO4/D2O. A significant suppression of the current density in the deuterated electrolyte was observed (Fig. 5d), yielding a KIE value > 2.0 across a wide potential range (Fig. 5e), which is a definitive signature that a proton-transfer step is rate-determining58. This kinetic evidence is consistent with a spillover-limited process. In stark contrast, both pure Ni17W3 and WO2 exhibit KIE values close to 1 (Supplementary Fig. 32), indicating the absence of a proton-transfer-limited mechanism. Furthermore, the Bode plot showed a characteristic frequency shift to lower values in D2SO4/D2O (Fig. 5f), confirming the slower kinetics originated from deuterium’s higher mass.

Beyond kinetics, the catalyst’s intrinsic ability to dissociate and transport hydrogen atoms was confirmed by the tungsten blue test. When physically mixed with WO3 and heated in a dilute H2 atmosphere, Ni17W3–WO2 induced a rapid color change of WO3 from yellow to dark blue, signaling the formation of HxWO3 (Supplementary Fig. 33a). A control experiment with pure WO2 showed no change, proving that hydrogen dissociation and spillover originate solely from Ni17W3–WO2 catalyst (Supplementary Fig. 33b). This provides tangible proof of the catalyst’s spillover function, corroborating the electrochemical KIE findings.

In situ Raman spectroscopy was employed to directly visualize the proton transfer pathways, with comparative analysis of WO2, Ni17W3, and the Ni17W3–WO2 heterostructure providing critical insights (Supplementary Fig. 34). For pure WO2, applying a cathodic potential caused the characteristic W-O peak (~860 cm−1) to weaken and shift (Fig. 5g)59, a definitive sign of H+ insertion and reduction50. This confirms WO2’s inherent capability to act as a proton acceptor in a conventional spillover process. Strikingly, the behavior of the Ni17W3–WO2 heterostructure is fundamentally different (Fig. 5h). Under cathodic polarization, the W-O bond signal (780 cm–1) remained stable60, showing no evidence of proton insertion. Instead, we observed the gradual disappearance of interfacial Ni-O bonds (515 cm–1 and 603 cm−1) and the concurrent emergence and intensification of a Ni-H bond (980 cm−1)61. This provides evidence that hydrogen is exclusively accumulating and reacting within the Ni17W3 phase, while the adjacent WO2 phase shows limited participation in proton uptake. This contrast is further contextualized by pure Ni17W3 (Fig. 5i), where the Ni-H bond forms at a less negative potential, consistent with its overly strong, performance-limiting hydrogen adsorption.

To evaluate the generality of the strain-gradient design principle beyond the Ni-W composition, we applied the same in situ phase reconstruction protocol to synthesize a Co0.87W0.13–WO2 composite (Supplementary Fig. S35). Electrochemical testing confirmed its HER activity (Supplementary Fig. S36a-b). Critically, this distinct material replicates the key mechanistic fingerprints of non-interfacial spillover established above. A pronounced KIE≈2 (Supplementary Fig. 36c-d) confirms a proton-transfer-limited process. Operando Raman spectroscopy (Supplementary Fig. 37) reveals hydrogen accumulation exclusively within the Co-W alloy phase (weakening of Co-O and emergence of Co-H bond) with no proton insertion into WO2 (stable W-O bond)60-62. The replication of these signature features in a chemically distinct system confirms that the non-interfacial spillover mechanism is a direct consequence of the engineered heterointerface, rather than a peculiarity of the Ni-W composition.

Atomic models of the catalysts were constructed and analyzed using DFT calculations to investigate their catalytic mechanisms (Supplementary Figs. 38 and 39). The electronic structures of WO2, Ni17W3, and the Ni17W3–WO2 heterostructure were first examined. All models exhibit metallic character with continuous states at the Fermi level (Supplementary Figs. 40 and 41), wherein the heterostructure demonstrates a more uniform charge distribution that facilitates efficient electron transport63. The d-band center profile of the catalysts shows that Ni17W3 is closest to the Fermi level at –2.04 eV (Fig. 6a), followed by Ni17W3–WO2 (–2.17 eV) and WO2 (–4.19 eV). The results indicate that Ni17W3 has a stronger hydrogen adsorption capability, while WO2 shows weaker adsorption, which is consistent with the CV results64-65. Therefore, the difficulty in hydrogen desorption at Ni17W3 and adsorption at WO2 leads to a higher overpotential66. The Ni17W3–WO2 presents a balanced hydrogen adsorption behavior compared to the Ni17W3 and WO2, thereby optimizing the reaction thermodynamics with a lower overpotential.

DFT calculations of the hydrogen adsorption free energy (ΔGH*) provide insight into the limitations of the individual phases (Fig. 6b-c). For WO2, the Volmer step presents a prohibitive energy barrier of 1.041 eV, identifying it as the rate-determining step (RDS). This high barrier confirms the fundamental difficulty in initiating the HER on WO2 due to its weak proton adsorption. Conversely, on Ni17W3, a W site exhibits the most favorable ΔGH*, enabling efficient proton adsorption. However, the subsequent hydrogen desorption step (H-step) becomes the RDS with a significant barrier of 0.633 eV. Thus, WO2 and Ni17W3 are intrinsically limited by the initial adsorption and final desorption steps, respectively, which explains their suboptimal HER activities.

Furthermore, we constructed a realistic Ni17W3–WO2 heterostructure model incorporating the strain gradient mapped by GPA results to precisely evaluate the HER pathway (Fig. 6d-e). The strains at the N1 (N = nickel site) and W1 (W = tungsten site) are 1.2%, the N2 is 0.15%, and the N3, W2, and W3 are −0.05%. The calculations reveal a strategically engineered gradient of hydrogen adsorption free energy (ΔGH*) across the Ni17W3 phase, with values decreasing from –0.553 eV at the bulk N1 site to –0.199 eV at the interfacial N3 site. This gradient creates a thermodynamically driven “proton highway” confined entirely within the Ni17W3 phase (N1→N2→N3), where protons spontaneously migrate from strong-adsorption sites to an optimally weak-adsorption site primed for H release. Auxiliary W sites (W1, W2) within the Ni17W3 phase act as efficient proton reservoirs due to their lower ΔGH*, supplying protons to the adjacent Ni sites via W1→N1 and W2→N3 pathways. Most decisively, the energy barrier for proton transfer from the optimal Ni site (N3) to the adjacent WO2 phase (W3 site, ΔGH* = 0.467 eV) is prohibitively high. This starkly demonstrates that cross-interface proton transfer is both thermodynamically and kinetically unfavorable, ruling out conventional interfacial H-spillover. Instead, the reaction follows a non-interfacial mechanism, wherein the entire process from adsorption to diffusion to desorption is completed within the Ni17W3 phase, decisively bypassing the heterointerface.

Thus, a highly efficient non-interfacial H-spillover channel (N1→N2→N3) is established within the Ni17W3 phase of the Ni17W3–WO2 heterostructure. The synergy of interfacial electron transfer and the strain gradient collectively engineers this pathway, which features a minimized RDS energy barrier of 0.199 eV, which shows a 10% reduction compared to the unstrained model (Supplementary Fig. 42). This work demonstrates a successful design principle: by leveraging synergistic effects to construct an optimal reaction path within one phase and strategically circumvent interfacial bottlenecks, unlocking exceptional catalytic activity.

Methods
Synthesis of Ni17W3–WO2. 4.0 mmol of NiC4H6O4·4H2O (Macklin, 99.9%) and 4.0 mmol of Na2WO4·2H2O (Macklin, 99%) were dissolved in 200 mL of deionized water (DW), which was placed in a water bath and heated to 80 oC for 3 hours. The mixture was filtered and washed several times with DW and ethanol (Sinopharm, ≥99.7), followed by drying and further annealing at 600 oC in the air for 1 hour to obtain the NiWO4 precursor. The Ni17W3–WO2 heterostructure was produced by annealing the NiWO4 precursor at 600 oC in N2/H2 (containing 8% H2) atmosphere for 10 hours with a ramp rate of 5 oC min−1.

Synthesis of WO2. 1.0 g of (NH4)6H2W12O40·xH2O (Macklin, 99.5%) was spread in a ceramic boat, which was placed in a tube furnace and heated to 700 oC in an air atmosphere for 1 hour with a ramp rate of 5 oC min−1. The product was then annealed at 600 oC in an N2/H2 (containing 8% H2) atmosphere for 2 hours with a ramp rate of 5 oC min−1. The final product was designated as WO2.

Synthesis of Ni17W3. 0.4 g NiSO4·6H2O (Macklin, 99.9%) and 0.062 g (NH4)6H2W12O40·xH2O were dissolved in 3 mL DW, followed by stirring until the mixture was completely dissolved. Then, 15 mL of ethylene glycol was added, and the mixture was stirred for another 20 min. Then, 2 mL of NH3·H2O was added dropwise to the above aqueous solution, which was then irradiated in the microwave reactor at 200 oC for 8 min. The precipitate was washed and centrifuged with DW and ethanol before drying. The product was then annealed in an N2/H2 (containing 8% H2) atmosphere at 750 oC for 2 hours with a heating rate of 5 oC min−1 to obtain the Ni17W3 alloy.

Synthesis of Co0.87W0.13–WO2. The Co0.87W0.13–WO2 composite was synthesized via the same in situ phase reconstruction method described for Ni17W3–WO2, with CoC4H6O4·4H2O (Macklin, 99.5%) used as the cobalt precursor in place of NiC4H6O4·4H2O.

Characterizations. Scanning electron microscopy (SEM, Hitachi SU3500) and transmission electron microscopy (TEM, Titan GC260-300) combined with energy-dispersive X-ray spectroscopy (EDS) were employed to investigate the surface morphology and elemental distribution of the samples. X-ray diffraction (XRD, LabX XRD-6100, Shimadzu), high-resolution transmission electron microscopy (HR-TEM, FEI Tecnai G2 F20), aberration-corrected scanning transmission electron microscopy (AC-STEM, JEM-ARM300F), and Raman spectroscopy (HR RamLab) were utilized to determine the crystalline structure and phase characteristics of the samples. The surface chemistry of the samples was recorded by X-ray photoelectron spectroscopy (XPS, Thermo Scientific K-Alpha, Thermo Fisher) using a monochromatic Al Kα X-ray source. The purity of the gas was determined by means of gas chromatographic (GC, GC-7820) analysis.

Electrochemical measurements. Electrochemical tests were conducted in 0.5 M H2SO4 aqueous solution using a CHI 760E workstation and a conventional three-electrode setup. The electrolyte (0.5 M H2SO4) was prepared by diluting concentrated sulfuric acid (XiLONG SCIENTIFIC 98%) with DW, followed by stirring for 30 min. The resulting solution was used immediately after preparation. The Hg/Hg2Cl2 reference electrode (CHI150) was calibrated against a fresh commercial one by measuring their stabilized open-circuit potential (OCP) prior to use. The graphite rod served as the counter electrode. The as-prepared electrocatalysts modified glassy carbon electrode (GCE, CHI104, 3 mm in diameter) served as working electrodes to investigate the HER performance in 100 mL 0.5 M H2SO4. To prepare the modified GCE, 16 mg of electrocatalyst was dispersed in a mixture of 600 μL of DW and 400 μL of N, N-dimethylformamide by ultrasonication to form an ink. Then, 4 μL (0.004 × 16 =0.064 mg catalyst) of the ink was dropped onto a clean GCE and dried (0.914 mg cm−2). Then, the GCE was covered with a Nafion solution (5 wt%, 5 μL). Linear sweep voltammetry (LSV) was conducted at a scanning rate of 5 mV s−1. Unless otherwise noted, all potentials were compensated for iR corrected and converted to the reversible hydrogen electrode (RHE) using the Nernst equation: E = E + 0.242 + 0.059 × pH. The pH was measured to be 0.62 ± 0.03 by a pH meter (FE28, Mettler Toledo). The solution resistance was determined by electrochemical impedance spectroscopy (EIS) performed at an initial potential of −0.22 V vs. RHE in the frequency range of 100 kHz to 0.1 Hz. The Tafel slopes were obtained by plotting the overpotentials (η) vs. log currents (log |j|) derived from the polarization curves. The double-layer capacitance (Cdl) was determined from cyclic voltammetry (CV) curves in the non-Faradic region at scan rates ranging from 1 to 10 mV s−1, which was subsequently used to estimate the electrochemically active surface area (ECSA). By plotting Δj/2 at OCP against scan rates, Cdl is estimated with the slope. Subsequently, ECSA is determined as ECSA = A*Cdl/Cs, where A is the geometric area of the catalyst, and Cs is the specific capacitance (Cs = 0.04 mF cm−2). A Kinetic Isotope Effect (KIE) experiment was performed in 0.5 M D2SO4/D2O electrolyte at a scanning rate of 5 mV s−1. The large current HER properties were obtained using Ni17W3–WO2 modified carbon cloth (Ni17W3–WO2/CC) with a mass loading of 0.914 mg cm−2 as a working electrode. The stability test was conducted using Ni17W3–WO2/CC with a constant current density of 500 mA cm−2. The PEM electrolyzer employed 0.5 M H2SO4 as the electrolyte at room temperature (20 °C) or 80 °C with a scan rate of 5 mV s−1. The catalyst ink for PEM was prepared by dispersing 10 mg of catalyst (IrO₂, Pt/C, or Ni17W3–WO2) in 50 mg of Nafion solution (5 wt%) and 100 mg of ethanol. The mixture was ultrasonicated in an ice bath for 60 min to form a homogeneous suspension. The ink was then brush-coated onto carbon paper and dried under ambient conditions. The catalyst loading was determined by weighing the carbon paper before and after coating and was approximately 10 mg cm−2. A proton exchange membrane (Fumasep FS-990-PK) with a thickness of 90 μm and an area of 2 × 2 cm2 was used as received without pre-treatment. The tungsten blue experiment was performed in a tube furnace at 180 °C for 30 min under a N2/H2 mixed atmosphere (8% H2).

Calculation methods. The DFT calculation was implemented in the Vienna ab initio simulation package (VASP)67 via the first-principles plane-wave pseudopotential formation and with the generalized gradient approximation (GGA) of Perdew-Burke-Ernzerhof (PBE) function. For the optimization of the structure, a 420.0 eV cutoff energy, 10−5 eV energy convergence accuracy, and 0.03 eV Å−1 residual force, a 1 × 1 × 1 k-mesh is utilized. For the calculation of free energy for adsorbed hydrogen (∆GH*), a 450.0 eV cutoff energy, and 0.02 eV Å−1 residual force are utilized. Adsorption energy (Eads) was calculated as: ∆E = E(surface+H) – E(surface) – 1/2EH2. ∆GH* could be calculated as: ∆GH* = ∆E + ∆(ZPE) – T∆S, where the ∆S represents the entropy change. E(surface + H) refers to the total energy of the system after a hydrogen atom is adsorbed on the basal plane, while E(surface) represents the energy of the clean surface, and EH2 corresponds to the total energy of an isolated H2 molecule. The ∆(ZPE) denotes the zero-point energy (ZPE) difference between an adsorbed hydrogen atom and a gaseous H2 molecule, as determined through normal mode analysis68-69. In this experiment, the entropy of gaseous hydrogen under standard conditions (1 bar of H2, pH = 0, and temperature T = 298.15 K) was taken into account. The strain model was constructed by stretching (1.2% and 0.15%) and compressing (−0.05%) the crystal structure of the model along the direction parallel to the surface. The corresponding models are included in Supplementary Data 1. It should be noted that while DFT calculations provide valuable insights into the electronic structure, there are inherent limitations in modeling the exact material under actual catalytic conditions.

Data availability. The source data generated in this study are provided in the Source Data file. Source data are provided with this paper.

Peer review rounds
[{"round": 1, "reviewer_comments": {"reviewer_1": "In this manuscript, A non-interfacial H-spillover mechanism is proposed in the Ni17W3–WO2 heterostructure, wherein proton adsorption, migration, and desorption occur entirely within the Ni17W3 phase, bypassing cross-phase diffusion. While the manuscript presents a potentially interesting heterostructured catalyst for hydrogen evolution, the central claim of a \"non-interfacial hydrogen spillover\" mechanism is insufficiently demonstrated, and several critical mechanistic, theoretical, and benchmarking gaps remain. Therefore, I recommend major revision at this stage, pending substantial improvements in mechanistic clarity.\n1. The central claim of a “non-interfacial hydrogen spillover” mechanism lacks clarity and rigorous mechanistic proof. Although DFT simulations show a progression of ΔGH* values across Ni17W3 lattice sites, this does not convincingly exclude interfacial effects. In fact, the presence of a Ni17W3–WO2 heterointerface inherently introduces interfacial contributions. Operando Raman and electrochemical data are insufficient to rule out conventional interfacial H-spillover. The novelty claim must be significantly better substantiated.\n2. The authors rely heavily on geometric phase analysis (GPA) to argue for a built-in strain gradient driving proton migration. However, the strain gradient reported (–0.2% to +1.1%) is relatively modest and may not be sufficient to dominate hydrogen diffusion energetics. There is a lack of quantitative correlation between the observed strain and calculated hydrogen adsorption/diffusion energy changes. The direct influence of strain on HER performance must be supported by more rigorous thermodynamic and kinetic models or atomic-resolution strain-mapping.\n3. While the authors emphasize that the spillover occurs “entirely within the Ni17W3 phase,” the WO2 phase is still an active part of the heterostructure. The WO2 component contributes >80% by mass (XRD refinement), yet its electrochemical role is limited and primarily described as an electron sink. If WO2 is inactive in H* adsorption and only provides strain/electronic tuning, its inclusion requires more justification, especially as it may reduce the overall electrochemical surface area-to-mass ratio. More rigorous site-specific activity analyses are needed.\n4. Although the manuscript compares the overpotential and Tafel slopes against other Ni/W-based systems (e.g., Supplementary Table S1), it fails to adequately benchmark against Pt/C or commercial HER electrocatalysts in acidic media. A side-by-side comparison, including mass activity and durability at industrially relevant current densities, is essential. Additionally, it is unclear whether the catalyst retains stability and activity in a full cell configuration of PEM under a practical operational conditions.\n5. The reported Faradaic efficiency of 98.6% (Fig. S14) is not supported by gas chromatography or a precise control experiment to rule out side reactions. Moreover, the exceptionally high stability at 500 mA cm⁻² for 1500 h is striking but post-mortem characterization to confirm catalyst preservation is lacking. Providing comparative degradation profiles and microscopy/XPS data after long-term operation would strengthen the stability claim.\n6. The manuscript suffers from several inconsistencies and exaggerated statements. For example, the title and abstract overclaim novelty without clearly defining “non-interfacial spillover” or distinguishing this from conventional intraphase diffusion. Also, some Figures, especially DFT diagrams (e.g., Fig. 5e), are difficult to interpret and should include clearer legends and consistent units.", "reviewer_2": "In this work, the authors report a Ni17W3-WO2 heterostructure designed to enable non-interfacial hydrogen spillover for efficient HER. A strain gradient within the Ni17W3 phase and electron transfer to WO2 are proposed to facilitate hydrogen adsorption, migration, and desorption within the Ni17W3 phase. While the concept of confining hydrogen spillover to a single phase of the heterostructure is intriguing, the manuscript lacks sufficient experimental evidence to substantiate the occurrence of spillover and does not provide a detailed, site-specific mechanistic explanation. Given these unresolved mechanistic ambiguities and the absence of convincing supporting data, the manuscript is not suitable for publication in Nature communications.\n1. The synthesis method employed in this study appears nearly identical to that reported for an alkaline HER catalyst (Small 2020, 16, 2005184). The authors should clarify what distinguishes the present system and how these differences enable improved performance under acidic conditions.\n2. Non-interfacial hydrogen spillover has been reported previously (e.g., Nat. Commun. 2022, 13, 1189; Small 2025, 21, 2411883), with the latter demonstrating higher HER activity than the present study using non-precious metals. The authors should clarify what advantages the current catalyst provides compared to these previously reported systems.\n3. The d-spacing reported in Fig.1h deviates from the known value of the Ni17W3 (111) plane. Moreover, the SAED pattern shown in Fig.1g is insufficiently clear to confirm the crystal planes. More rigorous structural characterization is therefore required to substantiate the phase identification.\n4. In the DFT calculations, the hydrogen adsorption energies were evaluated at arbitrary sites without distinguishing between Ni and W atoms in the Ni17W3 phase. To properly elucidate the hydrogen spillover mechanism, the authors should conduct site-specific adsorption analyses to clarify the distinct roles of Ni and W in hydrogen adsorption and migration within the Ni17W3 phase.\n5. Experimental validation of the proposed hydrogen spillover mechanism is insufficient. The authors should therefore provide direct evidence of spillover, for example through WO3-based HxWO3 coloration and KIE experiments using D2SO4/D2O.\n6. Supplementary Fig.12 reveals evident performance degradation after the 1500-hour durability test, which contradicts the authors' claims of long-term stability. Furthermore, to substantiate the chemical and structural stability of the catalyst, post-durability characterization data (e.g., XPS, XRD, and SEM) should be provided.\n7. In Fig.4h, the disappearance of the Ni-O peak suggests H* adsorption at Ni sites. To clarify the spillover pathway, additional in situ Raman measurements at more negative potentials are needed to detect possible M-H vibration peaks as direct evidence of H* accumulation.\n8. The manuscript contains several inconsistencies, such as inconsistent peak labeling for identical analytical results and spelling errors in axis labels. A thorough revision is necessary to improve the clarity, consistency, and overall quality of the manuscript.", "reviewer_3": "Review of “Boosting hydrogen evolution reaction via non‐interfacial hydrogen spillover” for Nature Communications.\nRecommendation: Major revision\n#Comments to the Authors\nIn this manuscript, the authors report the synthesis of a Ni17W3‐WO2 heterostructure and its application as a non-conventional, non-interfacial hydrogen spillover catalyst for promoting the hydrogen evolution reaction. The study highlights the modulation of hydrogen adsorption and desorption mechanisms by WO2, with Ni17W3 proposed as active sites contributing to enhance HER performance. The conceptual framework is supported by theoretical investigations, and the mechanistic details are elaborated in the manuscript. However, several points require clarification. Therefore, I recommend a major revision of the manuscript to address the following comments.\n1. The authors assert that the morphology and particle size of NiWO4 precursor remain unchanged after annealing. However, as evidenced by the high-resolution images of NiWO4 precursor before (Figure S1b) and after annealing (Ni17W3-WO2) (Figure 1c), the morphology is largely similar, but the particle size has increased. Specifically, particle agglomeration has occurred, leading to larger particle sizes. The authors should recognize this observation and update the manuscript accordingly.\n2. The amount of evolved hydrogen is reported based on both experimental measurements and theoretical calculations. Additional details regarding the methodology used for the theoretical determination of hydrogen evolution should be provided to ensure clarity and reproducibility. Furthermore, it is observed that the experimental and theoretical hydrogen evolution values at each time interval are remarkably consistent, which is somewhat unusual. The authors are encouraged to include specific numerical values (for both experimental and theoretical results) to facilitate verification and to substantiate the accuracy of the data presented.\n3. The hydrogen production rates and Faradaic efficiencies for the individual Ni17W3 and WO2 components should also be calculated. Including these parameters will enable a clearer comparison with the heterostructure and help demonstrate the potential superiority of the interface in enhancing hydrogen evolution activity.\n4. The authors have calculated the mass activity of Ni17W3‐WO2 and compared it with that of Ni17W3 and WO2. However, the data points corresponding to the measurement are missing in Figure 3a. The authors should revise this figure carefully to accurately include all relevant data points for proper comparison.\n5. In addition to LSV curves of Ni17W3-WO2/CC before and after long-term tests, the structural stability of the catalyst be confirmed through structural (XRD)/morphological (TEM)/elemental analysis before and after long term tests. While the authors mention that the LSV curves after 1500 hours are nearly identical, a slight shift is nonetheless observed. The authors should address this discrepancy in the revised manuscript and discuss the possible underlying reasons for the observed shift in the LSV curves.\n6. The proton insertion and extraction processes in the Ni17W3‐WO2 catalyst, investigated via CV analysis, show that Ni17W3‐WO2 has a more negative hydrogen desorption potential compared to Ni17W3, while its hydrogen insertion potential lies between those of Ni17W3 and WO2. These findings appear to contradict the assumption of stronger hydrogen adsorption for enhanced hydrogen evolution performance attributed to Ni17W3‐WO2. Since the core concept of the manuscript revolves around hydrogen spillover to improve hydrogen generation, the authors are recommended to provide a clear discussion addressing this apparent inconsistency in the revised manuscript.\n7. The manuscript does not include a comparison with a standard commercial catalyst, such as Pt/C. Incorporating data for a well-established reference material would provide valuable context and facilitate meaningful comparisons of catalytic performance. The authors could refer Adv. Energy Mater. 2024, 14, 2401426 while revising the manuscript.\n8. The Nyquist plot of Ni17W3‐WO2 is not clearly visible in the Figure 4e. The enlarged figure demonstrating the semicircular Nyquist plot should be moved from supporting information to main text in the inset of Figure 4e.\n9. Color corresponding to each atom in atomic models should be specified in the captions of Figure 5e, Figure S17, and S18."}, "author_response": "Point-by-Point Responses to Reviewers’ Comments\nWe express our sincere gratitude to the editor and all reviewers for their invaluable\nfeedback, which we have utilized to enhance the quality of our manuscript (NCOMMS-\n25-38365). The reviewer comments are presented in italic and bold font, while our\nresponses, including the incorporation of additional figures, tables, descriptions, and\nother elements, are highlighted in blue text.\nPoint-by-point response to the reviewers #1\nReviewer #1 (Remarks to the Author): In this manuscript, A non-interfacial H-\nspillover mechanism is proposed in the Ni W –WO heterostructure, wherein proton\n17 3 2\nadsorption, migration, and desorption occur entirely within the Ni W phase,\n17 3\nbypassing cross-phase diffusion. While the manuscript presents a potentially\ninteresting heterostructured catalyst for hydrogen evolution, the central claim of a\n\"non-interfacial hydrogen spillover\" mechanism is insufficiently demonstrated, and\nseveral critical mechanistic, theoretical, and benchmarking gaps remain. Therefore,\nI recommend major revision at this stage, pending substantial improvements in\nmechanistic clarity.\nGeneral Response: Thanks for your positive assessment of our work’s potential and\nfor the constructive comments. We have thoroughly addressed the concerns regarding\nthe demonstration of the non-interfacial H-spillover mechanism and the identified\ngaps through substantial new experiments, theoretical calculations, and textual\nrevisions, as detailed in our point-by-point responses below. We believe the revised\nmanuscript now provides compelling evidence and significantly improved clarity.\nComment 1: The central claim of a “non-interfacial hydrogen spillover” mechanism\nlacks clarity and rigorous mechanistic proof. Although DFT simulations show a\nprogression of ΔG values across Ni W lattice sites, this does not convincingly\nH* 17 3\nexclude interfacial effects. In fact, the presence of a Ni W –WO heterointerface\n17 3 2\ninherently introduces interfacial contributions. Operando Raman and\nelectrochemical data are insufficient to rule out conventional interfacial H-spillover.\nResponse 1: Thank you for this critical comment, which pushed us to seek more direct\nand conclusive evidence to substantiate our central claim. We agree that the presence\nof a heterointerface does not automatically imply that interfacial spillover is the\ndominant mechanism. The key question is where the protons physically go. To clarify\nthe real H-spillover path in Ni W -WO , we have supplemented in situ Raman\n17 3 2\nscattering of the Ni W -WO heterostructure, as well as pure phases of WO and\n17 3 2 2\nNi W . Furthermore, more rigorous theoretical models were constructed for the\n17 3\ncomputational analysis of the reaction pathways. Our new experimental and\ntheoretical data now provide a definitive answer: the spillover pathway is confined\nwithin the Ni W phase, and the role of the WO phase is to electronically and\n17 3 2\nstructurally modulate Ni W , not to accept protons.\n17 3\n(1) Direct spectroscopic comparison via in situ Raman\nWe conducted in situ Raman spectroscopy on three samples under operating\nconditions. This comparative approach serves as a critical control experiment.\nFor pure WO : At the initial state (0 mV vs. RHE), the Raman peak at\n2\napproximately 860 cm−1 corresponds to the totally symmetric stretching vibration of\nthe W-O bond 1 (Fig. R1a or Fig. 4g). As the potential is scanned cathodically, the\ncharacteristic W–O peak significantly weakens and shifts, which is a classic signature\nof H* insertion into the WO lattice to disrupt the original W-O bond structure2. This\n2\nconfirms that pure WO can accept protons via interfacial spillover if available.\n2\na W-O (symmetric) WO2 b Ni-O W-O (asymmetric) Ni17W3-WO2 c Ni17W3\nNi-H Ni-H\nI\n( a )\nn\n. u\nt e\n.\nn\nsit y - - -\n- - -\n1 2 1\n2 3 4\n-5 5 0 0\n5 0 0\n0 0 0 0 0\n0 0 0\nm m m m m\nm m m\nV V V V V\nV V V I nt e n\nsit y ( a. u.) - - -\n- - -\n1 2 1\n2 3 4\n-5 5 0 0\n5 0 0\n0 0 0 0 0\n0 0 0\nm m m m m\nm m m\nV V V V V\nV V V I nt e n\nsit y ( a. u.) - - -\n- - -\n1 2 1\n2 3 4\n-5 5 0 0\n5 0 0\n0 0 0 0 0\n0 0 0\nm m m m m\nm m m\nV V V V V\nV V V\n-500 mV -500 mV -500 mV\n-600 mV -600 mV -600 mV\n600 800 1000 1200 600 800 1000 1200 600 800 1000 1200\nRaman shift (cm-1) Raman shift (cm-1) Raman shift (cm-1)\nFig. R1 | In situ Raman scattering of (a) WO , (b) Ni W –WO , and (c) Ni W\n2 17 3 2 17 3\nelectrodes recorded at different cathodic potentials.\n17 3 2\nshowed strong characteristic peaks at 515 cm–1 and 603 cm–1, attributed to the\nvibrational modes of Ni-O bonds at the Ni W -WO heterointerface (Fig. R1b or Fig.\n17 3 2\n4h). Additionally, the peak observed at 780 cm–1 is attributed to the asymmetric\nstretching vibration of W-O bonds, resulting from lattice distortions at the\nheterostructure interface 3. As the applied cathodic potential increased, the intensity\nof the Ni-O bonds gradually decreased. Additionally, a Raman peak assigned to the Ni-\nH bond emerged at 980 cm–1 4, which subsequently intensified with the increase of\ncathodic potential. Notably, the W-O bond signal remains stable, showing no\nsignificant weakening or shifting indicative of no H* insertion. This provides direct,\nreal-time evidence that hydrogen is accumulating in the Ni W phase, while the\n17 3\nadjacent WO phase remains essentially \"invisible\" to protons during HER. This\n2\ndecisively rules out WO as a proton acceptor in our system.\n17 3\n4i), consistent with its inherently strong hydrogen adsorption, which leads to sluggish\ndesorption.\n(2) Refined and realistic DFT modeling\nTo move beyond a simplified model, we used Geometric Phase Analysis (GPA)\n(please refer to Figs. 1i,j) from HR-TEM images to build a DFT model that incorporates\nthe actual strain field at the heterointerface to accurately simulate and characterize\nthe HER process (Figs. R2a, b or Fig. R5d, e).\nAcross the Ni W phase of Ni W -WO , the ΔG values of the N1, N2, and N3 (N\n17 3 17 3 2 H*\nrepresents Ni) sites are –0.553 eV, –0.391 eV, and –0.199 eV, respectively. This\nindicates that the hydrogen adsorption capability of Ni sites on the Ni W phase\n17 3\nprogressively weakens towards the interface. Therefore, H+ can adsorb on N1, and\nthen spillover to the interfacial N3 site via the N2 site (N1→N2→N3).\nFurthermore, W1 and W2 (W represents tungsten) sites also adsorb H+ due to their\nlower ΔG relative to adjacent Ni sites, which then spills over, supplying protons to\nH*\nthe adjacent Ni sites (W1→N1, W2→N3).\nH* 2\nenergy barrier for a proton to jump from the optimal Ni site (N3) at the interface to\nthe nearest W site (W3) on WO is calculated to be prohibitively high, making\n2\ninterfacial transfer thermodynamically and kinetically unfavorable compared to direct\nH desorption from the N3 site of the Ni W surface.\n2 17 3\nThe combination of these new results presents a coherent and compelling picture.\nThe in situ Raman data directly observe the absence of proton flow to WO , while the\n2\nrefined DFT theoretically explains why this is the case and identifies the favorable intra-\nphase pathway. Therefore, they provide rigorous and multifaceted proof for our\nproposed non-interfacial H-spillover mechanism. We have integrated these findings\nand clarified the novelty of our work in the revised manuscript.\nFig. R2 | (a) The ΔG for HER on Ni W -WO . (b) Schematic illustration of the H*\nH* 17 3 2\nadsorption models on various active sites and the proposed reaction pathways for\nNi W -WO .\n17 3 2\nComment 2: The authors rely heavily on geometric phase analysis (GPA) to argue for\na built-in strain gradient driving proton migration. However, the strain gradient\nreported (–0.2% to +1.1%) is relatively modest and may not be sufficient to dominate\nhydrogen diffusion energetics. There is a lack of quantitative correlation between\n4\nthe observed strain and calculated hydrogen adsorption/diffusion energy changes.\nThe direct influence of strain on HER performance must be supported by more\nrigorous thermodynamic and kinetic models or atomic-resolution strain-mapping.\nResponse 2: Thank you for your insightful comment. We agree that the strain\nmagnitudes are modest; however, our central argument is not that strain alone\ndominates the process, but that it acts as a fine-tuning mechanism which, when\ncoupled with the pre-existing electronic redistribution effects, collectively optimizes\nthe reaction pathway to achieve a lower overall barrier. To establish a direct\nquantitative correlation between the measured strain and the energetics, as the\nreviewer suggested, we have performed a rigorous comparative DFT analysis.\nWe first quantified the local strain at the atomic level using Geometric Phase\nAnalysis (GPA) (Figs. R3a, b or Figs. 1i, j). The strain profile reveals a clear gradient\n(Table R1), transitioning from ~1.2% tensile strain at sites (N1, W1) away from the\ninterface, to near-zero (0.15%) at intermediate sites (N2), and finally to a slight\ncompression (–0.05%) at the interfacial sites (N3, W2, W3).\nε\na b\nxy\n0.3 e\ng Ni W a c WO 1.0\n17 3\ne\nrf 2\n)\nnt %\ne r\noi\nn\n(\net 0.2a i\nh r\nt\nS\n0.0\n-0.3 -0.2\n0.0 2.5 5.0\nPosition (nm)\nFig. R3 | (a) GPA for the lattice strain maps alone with the xy direction. (b) Strain\nprofiles taken along with the white arrow.\nH*\nActive ΔG (Unstrained ΔG (Strained Change\nH* H*\nStrain state\nsite model) model) (ΔΔG )\nH*\nN1 1.2% tensile –0.509 eV –0.553 eV –0.044 eV\nN2 0.15% tensile –0.386 eV –0.391 eV –0.005 eV\nN3 −0.05% compressive –0.221 eV –0.199 eV 0.022 eV\nW1 1.2% tensile –0.552 eV –0.612 eV –0.060 eV\nW2 −0.05% compressive –0.315 eV –0.309 eV 0.006 eV\nW3 −0.05% compressive 0.441 eV 0.467 eV 0.026 eV\nWe then constructed two distinct DFT models to deconvolve the effects of\nelectronics and strain.\nModel A (electronic effect only without strain):\nAn unstrained Ni W -WO model is used to evaluate the hydrogen adsorption\n17 3 2\ncapability of active sites under electronic interaction only (Fig. R4a or Supplementary\nFig. 34). The ΔG values at various sites are listed in Table R1. Therefore, H+ can adsorb\nH*\non N1, and then spill over to the interfacial N3 site via the N2 site (N1→N2→N3).\nFurthermore, W1 and W2 sites also adsorb H due to the lower ΔG relative to adjacent\nH*\nNi sites, which then spills over, supplying protons to the adjacent Ni sites (W1→N1,\nW2→N3). At the interface, the ΔG of the W3 site (on the WO phase) is 0.441 eV,\nH* 2\nand the energy barrier for proton migration from N3 to W3 across the interface is much\nhigher than that for direct hydrogen release at the N3 site. Consequently, hydrogen is\nultimately released through non-interfacial H-spillover at the N3 site. The energy of\nthe rate-determining step (RDS) in the entire process is 0.221 eV.\nH* 17 3 2\non (b) Ni W -WO with applied strain coefficients. (c) Schematic illustration of the H*\n17 3 2\nadsorption models on various active sites and the proposed reaction pathways for\nNi W -WO .\n17 3 2\nModel B (electronic effect with strain effect):\nIn the realistic model with strain, the tensile strain at N1/N2 strengthens H*\nadsorption (Fig. R4b (or Fig. 5d) and Table R1, more negative ΔG ), facilitating proton\nH*\ninflux, while the slight compressive strain at the interfacial N3 site intentionally\nweakens the H adsorption, bringing its ΔG closer to the ideal value of zero. Its HER\nH*\nmechanism is also non-interfacial H-spillover (N1→N2→N3), identical to that of the\nunstrained model. This collective modulation results in a lower RDS barrier of 0.199\neV for desorption from the N3 site. The overall reduction of 0.022 eV is significant in\ncatalysis, which directly rationalizes the enhanced HER activity.\nIn conclusion, our analysis moves beyond a qualitative claim and provides the\n7\nrigorous quantitative correlation requested. We demonstrate that the GPA-mapped\nstrain gradient, while modest in magnitude, functions precisely to create a more\nfavorable energy profile across the spillover pathway, synergizing with electronic\neffects to achieve superior kinetics. This refined understanding has been incorporated\ninto the revised manuscript.\nComment 3: While the authors emphasize that the spillover occurs “entirely within\nthe Ni W phase,” the WO phase is still an active part of the heterostructure. The\n17 3 2\nWO component contributes >80% by mass (XRD refinement), yet its electrochemical\n2\nrole is limited and primarily described as an electron sink. If WO is inactive in H*\n2\nadsorption and only provides strain/electronic tuning, its inclusion requires more\njustification, especially as it may reduce the overall electrochemical surface area-to-\nmass ratio. More rigorous site-specific activity analyses are needed.\nResponse 3: We thank the reviewer for raising this critical point. We agree that a high\nmass fraction of a component that is not a direct active site requires compelling\njustification. Our data demonstrate that WO is not merely an \"electron sink\" but an\n2\nessential catalytic promoter that dramatically enhances the intrinsic activity of the\nNi W active sites and ensures their stability, with the net benefit far outweighing the\n17 3\n“cost” of its mass. We provide three key justifications:\n(1) Dramatically enhancing the intrinsic activity\nThe reviewer rightly notes that WO might reduce the electrochemical surface area\n2\n(ECSA)-to-mass ratio. However, the critical metric is the intrinsic turnover frequency\n(TOF), which reflects the efficiency of each active site. Ni W -WO exhibited a higher\n17 3 2\nTOF at the same overpotential compared to pure Ni W and WO (Fig. 3d), confirming\n17 3 2\nits improved intrinsic activity in HER, which originates from the formation of a\nheterointerface between the two phases. The heterointerface-induced strain field and\nelectron redistribution synergistically optimize the hydrogen adsorption capability of\nthe atoms at the heterointerface, thereby enhancing the intrinsic HER activity. This\nenhancement in intrinsic activity proves that the role of WO is not passive. By creating\n2\nthe optimal electronic and strain environment, it transforms the Ni W phase from a\n17 3\n8\nmediocre catalyst into an exceptionally active one. The gain in quality (TOF) per active\nsite completely compensates for any loss in the quantity of active sites per mass.\n(2) Creating the novel spillover pathway\nAs detailed in our response to Comments 1 and 2, the non-interfacial H-spillover\npathway is the core innovation of this work. This pathway does not exist in pure Ni W .\n17 3\nIt is exclusively enabled by the synergistic effects arising from the heterointerface. On\nthe one hand, the electron transfer from Ni W to WO fine-tunes the d-band center\n17 3 2\nof Ni sites. On the other hand, the strain gradient imposed by the WO substrate\n2\ncreates the energy landscape for directional proton migration. Without WO , Ni W\n2 17 3\nexhibits strong hydrogen adsorption but poor desorption, following a conventional\nand inefficient HER mechanism. Thus, WO is fundamental to achieving the novel and\n2\nhighly efficient catalytic function.\n(3) Critical enhancing stability\nWO with a high fraction acts as a stable, conductive, and acid-resistant scaffold that\n2\nanchors and protects the metastable Ni W alloy. The 1500-hour stability is a direct\n17 3\nconsequence of this protective role.\nTherefore, WO is not a diluent but a performance-enabling component, making its\n2\nmass fraction a worthwhile investment for the achieved performance. We have added\nthis discussion to the revised manuscript.\nComment 4: Although the manuscript compares the overpotential and Tafel slopes\nagainst other Ni/W-based systems (e.g., Supplementary Table S1), it fails to\nadequately benchmark against Pt/C or commercial HER electrocatalysts in acidic\nmedia. A side-by-side comparison, including mass activity and durability at\nindustrially relevant current densities, is essential. Additionally, it is unclear whether\nthe catalyst retains stability and activity in a full cell configuration of PEM under\na practical operational conditions.\nResponse 4: We sincerely thank the reviewer for this critical suggestion regarding\nbenchmarking and practical validation. We have conducted a comprehensive\nperformance comparison and demonstrated the exceptional performance of our\n9\ncatalyst in a practical PEM water electrolyzer, which strongly underscores its industrial\nrelevance.\nAs suggested, we have performed a detailed side-by-side comparison of our Ni W -\n17 3\nWO catalyst with commercial Pt/C, Ru, and a wide range of recently reported precious\n2\nand non-precious metal catalysts (Fig. R5 (Fig. 3f) and Table R2 (Supplementary Table\n1). Our catalyst achieves an overpotential of 21 mV at 10 mA cm−2 and a Tafel slope of\n47.8 mV dec−1, which is comparable to, and in some cases superior to, commercial Pt/C\nand advanced Pt-based catalysts. We acknowledge the reviewer's point on mass\nactivity. While there is a disparity with ultra-low loading Pt single-atom catalysts, the\nentirely non-precious composition of our material ensures an exceptional cost-to-\nperformance ratio, which is a paramount consideration for industrial scaling.\n1.6\nM\na 20%Pt/C\ns\ns1.2 Pt-based\nlo\nRu\na\nd Ru-based\nin0.8\nNon-noble\ng\n( This work\n.\nm\ng0.4\nc\nm\n)\n-20.0\nO ver 1\np\n5\no t\n0\nen\nt\n1\nia\n0\nl\n0\n(m\nV )\n50\n4\nf\n0\nel\nslop\n6\ne\n0\n(\nmV\n8\nd\n0 ec-1)\nFig. R5 | Comparison of the overpotential at 10 mA cm−2, Tafel slope, and catalyst\nloading for the Ni W -WO /CC against commercial Pt/C, Ru, Pt-based, Ru-based, and\n17 3 2\nother non-precious metal catalysts in 0.5 M H SO .\n2 4\nTable R2. Comparison of the overpotentials (at 10 mA cm−2), Tafel slopes, and catalyst\n10\nloading of different catalysts in 0.5 M H SO .\n2 4\nOverpotential Tafel slope Loading\nCatalyst Ref.\n(mV) (mV dec−1) (mg cm−2)\n17 3 2\n20%Pt/C 36 35 0.51 5\nRu 67 80.2 0.16 6\nPt /OLC 38 36 0.51 5\nPt@WS 31 27.2 0.256 8\nCoNiPt @G 23 55 1.6 10\nSA\n2H/1T-WS -10 113 55 0.52 11\n2\nhcp-RuCu HUNSs 98 75 0.1 12\nRu/CNT 27 59.6 0.34 13\nRu-Co Ni 57 52 1.2 14\nRu/Pd Cu-70 28 46 1 18\n1 1\nRuSe 59 52 0.25 20\n2\np+-WS -9.8 92 33 0.52 22\n2\nMoS -NbS 159 53 1.41 24\n2-x x\nPb-CoSe -DETA 74 45 1.21 25\n0.1 0.9 2\nMore critically, we benchmarked performance at industrially relevant current\ndensities (Table R3 or Supplementary Table 2). At 500 mA cm−2, Ni W -WO requires\n17 3 2\nonly 138 mV overpotential, a performance on par with complex multi-metallic\nprecious metal catalysts. Furthermore, it sustains this high-rate hydrogen production\nfor 1500 hours with a negligible degradation rate of 28 µV h−1, significantly\noutperforming most reported catalysts in both activity and durability at this\nchallenging current density.\nTable R3. Comparison of the overpotential at 500 mA cm−2 and stability of Ni W -\n17 3\nWO /CC against state-of-the-art electrocatalysts.\nCatalyst Electrolyte Ref\n@500 mA cm−2 (h)\nNi W -WO /CC 138 1500 0.5 M H SO This work\n17 3 2 2 4\nPt-Ru/CNT 113 100 0.5 M H SO 27\n2 4\nMoS -P2 302 240 0.5 M H SO 28\n2 2 4\nPt V 240 100 0.5 M H SO 29\n3 2 4\nPtRuFeCo/CFP, 80.1 500 0.5 M H SO 30\n2 4\nCo/NC-HP@Si-NWs 440 24 0.5 M H SO 31\n2 4\nNi P-Ni P @Ni S /NF 410 24 0.5 M H SO 32\n2 12 5 3 2 2 4\n2 2 4\nNi/Cu/MoO @LC 330 120 0.5 M H SO 34\n2 2 4\nPd@CIPS 182 20 0.5 M H SO 35\n2 4\nPt -Ru P@Ru/CNT 90 120 0.5 M H SO 36\n0.095 2 2 4\nIr -NiFe LDH/NiMo 139 500 1.0 M KOH 37\nSA\nMN-NMN09/NF 124 120 1.0 M KOH 38\nFe P/Co N 131 42 1.0 M KOH 39\n2 2\nRu−LC−Ni(OH) 110 300 1.0 M KOH 40\nR\nNMFSOH 210 350 1.0 M KOH 42\nTo directly address the reviewer's question about full-cell performance, we\nfabricated membrane electrode assemblies (MEAs) and tested our catalyst under\npractically relevant PEM electrolyzer conditions (Figs. R6a, b or Supplementary Fig.\n25). The cathode and anode are fabricated by depositing Ni W -WO and IrO onto\n17 3 2 2\ncarbon paper substrates, respectively. The Ni W -WO ||IrO cell demonstrated\n17 3 2 2\noutstanding performance, requiring a cell voltage of only 1.945 V to achieve 1000 mA\ncm−2at room temperature, which is highly competitive with the Pt/C||IrO benchmark\n2\n(Fig. R6c or Fig. 3h). Most importantly, the cell exhibited exceptional operational\nstability over 1000 cycles at current densities of 500 and 1000 mA cm−1 (Fig. R6d or\nSupplementary Fig. 26a), with the polarization curves before and after cycling showing\nnearly perfect overlap (Fig. R6e or Supplementary Fig. 26b).\nelectrolyzer during testing. (c) LSV curves of PEM electrolyzers assembled with Ni W -\n17 3\nWO ||IrO and Pt/C||IrO configurations at room temperature. (d) The voltages\n2 2 2\ncorresponding to 500 and 1000 mA cm−1 during prolonged cycling of the Ni W -\n17 3\nWO ||IrO assembled PEM electrolyzer. (e) LSV curves for the Ni W -WO ||IrO\n2 2 17 3 2 2\nassembled PEM electrolyzer before and after 1000 cycles.\nTherefore, these new results provide compelling evidence that our Ni W -WO\n17 3 2\ncatalyst is not only a high-performance non-precious HER catalyst in fundamental\nthree-electrode tests but also a viable and durable replacement for Pt/C in practical\nPEM electrolyzers. We have integrated these critical benchmarks and the PEM\nelectrolyzer data into the revised manuscript to underscore the practical impact of our\nwork.\nComment 5: The reported Faradaic efficiency of 98.6% (Fig. S14) is not supported by\ngas chromatography or a precise control experiment to rule out side reactions.\nMoreover, the exceptionally high stability at 500 mA cm-2 for 1500 h is striking but\npost-mortem characterization to confirm catalyst preservation is lacking. Providing\ncomparative degradation profiles and microscopy/XPS data after long-term\noperation would strengthen the stability claim.\nvalidation of Faradaic efficiency (FE) and stability. We have now conducted the\nrecommended control experiments and comprehensive post-mortem characterization,\nwhich robustly confirm both the selectivity and the structural integrity of our catalyst.\nTo unequivocally determine the FE and rule out any side reactions, we performed\ngas chromatography (GC) analysis on the gaseous products evolved during HER at the\nNi W -WO /CC electrode. As shown in Fig. R7 (Supplementary Fig. 17a), the GC trace\n17 3 2\nexhibits a single, sharp peak at a retention time of ~1.5 minutes, which is exclusively\nattributed to H . Crucially, no other gaseous products were detected. This result\n2\nprovides direct and unambiguous evidence that the electrochemical current is\ndedicated solely to hydrogen production, thereby validating the high FE (98.65%) and\nconfirming the exceptional selectivity of our catalyst.\nH\n2\n)\n. u\n. a\n( . y H 100%\n2\nt\ni\ns\nn\ne\nt\nn\nI\nNi W -WO /CC\n17 3 2\n0 3 6 9 12 15\nRetention (min)\nFig. R7 | Gas chromatography analysis of the gaseous products from the Ni W -\n17 3\nWO /CC electrode during HER.\n17 3 2\nthe long-term stability test to investigate its structural and chemical preservation. XRD\npatterns (Fig. R8a or Supplementary Fig. 21) confirm that the crystalline phases of\nboth Ni W and WO remain intact after the prolonged test. A slight attenuation in\n17 3 2\nthe Ni W peak intensity suggests minimal, superficial dissolution, which is expected\n17 3\n15\nfor any non-noble metal in acid over such an extended period, yet the core structure\nis preserved.\nTEM and HR-TEM images (Figs. R8b, c or Supplementary Fig. 22) provide direct\nvisual evidence of stability. The granular morphology is maintained, and the clear\nlattice fringes corresponding to Ni W (111) (0.24 nm) and WO (011) (0.35 nm)\n17 3 2\nremain resolvable. Most importantly, the heterointerface between the two phases is\nstill distinctly visible, indicating that the core catalytic architecture survives virtually\nunchanged.\nFig. R8 | (a) XRD patterns of Ni W -WO before and after long-term stability testing.\n17 3 2\n(b) TEM and (c) HR-TEM images of Ni W -WO after long-term stability testing.\n17 3 2\nThe surface chemical states of post-tested Ni W -WO were thoroughly\n17 3 2\ninvestigated using XPS. The W-4f spectrum (Fig. 9a or Supplementary Fig. 23a) shows\nthe persistence of the W4+ state from the WO phase. The increased proportion of\n2\nhigher oxidation states (W5+/ W6+) is attributed to the partial oxidation of the\noutermost surface layer upon exposure to the acidic electrolyte, a common and self-\nlimiting phenomenon that actually protects the bulk material. This led to a decrease\nin the intensity of W4+ in the XPS results and the masking of the low-abundance W0\nsignal from the alloy phase.\nFor the high-resolution Ni-2p XPS spectra (Fig. 9b or Supplementary Fig. 23b), the\nclear presence of the Ni0 state even after long-term operation is remarkable and\nprovides direct evidence that the metallic Ni active sites within the Ni W alloy are\n17 3\neffectively stabilized against complete oxidation or leaching, ascribed to the protective\n16\nrole of the WO matrix and the electronic interaction between the phases.\n2\na b\nW5+\nW5+\nNi2+\nu . ) u . )\na . a .\ns i t y\n(\ns i t y\n(\nsat. Ni3+ Ni2+ sat.\nn n\ne W6+ W6+ e Ni3+\nI n t W 5p W4+ W4+ I n t Ni0 Ni0\nafter test after test\n44 40 36 32 880 870 860 850\nBinding energy (eV) Binding energy (eV)\nFig. R9 | (a) W-4f and (b) Ni-2p high-resolution XPS spectra of Ni W -WO after long-\n17 3 2\nterm stability testing.\nTherefore, this comprehensive dataset strongly reinforces the credibility of our\nreported performance metrics. We have incorporated these new findings into the\nrevised manuscript.\nComment 6: The manuscript suffers from several inconsistencies and exaggerated\nstatements. For example, the title and abstract overclaim novelty without clearly\ndefining “non-interfacial spillover” or distinguishing this from conventional\nintraphase diffusion. Also, some Figures, especially DFT diagrams (e.g., Fig. 5e), are\ndifficult to interpret and should include clearer legends and consistent units.\nResponse 6: We sincerely thank the reviewer for pointing out these critical issues. We\nhave thoroughly revised the manuscript to eliminate exaggerated language, provide\nrigorous definitions, and significantly improve the clarity of our figures. The changes\nare detailed below.\nWe have meticulously revised the manuscript to adopt a more objective and precise\ntone and to unambiguously define our core concept.\nIn the abstract, the novelty of this work is now stated as: “This work demonstrates\nhow tailored heterostructures can effectively bypass interfacial bottlenecks, providing\n17\na new paradigm for developing efficient non-precious H-spillover catalysts and\nadvancing sustainable hydrogen production.”\nIn the introduction, we now clearly define \"non-interfacial H-spillover\" in direct\ncontrast to the conventional mechanism. “Herein, we report a generalizable design\nstrategy to bypass interfacial bottlenecks in H-spillover catalysis (Scheme 1). Through\nin situ phase reconstruction, we construct a Ni W -WO heterostructure where a\n17 3 2\nsynergistic combination of interfacial electron transfer and a built-in strain gradient\ncollectively reconfigures the hydrogen adsorption landscape within the Ni W phase.\n17 3\nThis engineering creates an efficient, single-phase-confined proton migration pathway,\nenabling the complete HER process from adsorption to desorption without cross-\ninterface proton transfer.”\nIn addition, we have completely redesigned the key DFT diagram (Fig. R10 or Fig. 5e)\nto ensure it is immediately interpretable. The new figure includes a comprehensive\nlegend that explicitly defines all colors and symbols. All energy values and spatial\nscales are displayed with consistent, easy-to-read units.\nFig. R10 | Schematic illustration of the H* adsorption models on various active sites\nand the proposed reaction pathways for Ni W -WO .\n17 3 2\nReviewer #2 (Remarks to the Author): In this work, the authors report a Ni W –WO\n17 3 2\nheterostructure designed to enable non-interfacial hydrogen spillover for efficient\nHER. A strain gradient within the Ni W phase and electron transfer to WO are\n17 3 2\nproposed to facilitate hydrogen adsorption, migration, and desorption within the\nNi W phase. While the concept of confining hydrogen spillover to a single phase of\n17 3\nthe heterostructure is intriguing, the manuscript lacks sufficient experimental\nevidence to substantiate the occurrence of spillover and does not provide a detailed,\nsite-specific mechanistic explanation. Given these unresolved mechanistic\nambiguities and the absence of convincing supporting data, the manuscript is not\nsuitable for publication in Nature communications.\nGeneral Response: Thank you for your constructive feedback and for acknowledging\nthe intriguing concept of our work. We have taken the concerns regarding the\nexperimental evidence for spillover and the site-specific mechanistic explanation very\nseriously. In this version, we have conducted a comprehensive set of new experiments\nand theoretical calculations, which we believe decisively address these points. These\ninclude direct spectroscopic, chemical, and electrochemical validation of the hydrogen\nspillover process, as well as a refined, site-specific DFT model that clarifies the distinct\nroles of Ni and W atoms in the mechanism.\nWe hope these substantial additions have fully resolved the previous mechanistic\nambiguities and provided the convincing supporting data required. We appreciate the\nopportunity to have significantly improved our manuscript and believe it now meets\nthe high standard of Nature Communications. Our detailed point-by-point responses\nfollow below.\nComment 1: The synthesis method employed in this study appears nearly identical\nto that reported for an alkaline HER catalyst (Small 2020, 16, 2005184). The authors\nshould clarify what distinguishes the present system and how these differences\nenable improved performance under acidic conditions.\nResponse 1: Thanks for your insightful comment, which allows us to clarify the\n19\nfundamental distinctions between our work and the prior synthesis. While the\nprecursor (NiWO ) is similar, the synthesis conditions, final phase composition, and the\n4\nresulting catalytic mechanism are fundamentally different, leading to the unique\nperformance in acidic HER.\n(1) Distinct material and synthesis goal\nLi et al. (Small 2020, 16, 2005184): Aimed to create a Ni W -W heterostructure via\n17 3\nhydrogen reduction. The presence of the metallic W phase is crucial for H O\n2\ndissociation in alkaline media.\nOur Work: We carefully controlled the reduction kinetics to achieve a Ni W -WO\n17 3 2\nheterostructure. The preservation of the WO phase is intentional, as it acts as an\n2\nelectron acceptor for work function-driven electronic modulation, a role distinct from\nmetallic W.\n(2) Fundamentally different catalytic mechanisms\nLi et al. (Small 2020, 16, 2005184): Operates via a traditional interfacial spillover\nmechanism in alkali: H O dissociates on the W phase, and H atoms spill over to the\n17 3\nbetween two active phases.\nOur Work: We propose a non-interfacial H-spillover mechanism. The directional\nelectron transfer and built-in strain gradient reconfigure the energy landscape within\nthe Ni W phase itself. This creates a proton migration pathway where adsorption,\n17 3\nspillover, and desorption all occur exclusively on the Ni W phase (Fig. R11 or Scheme\n17 3\n1), bypassing any need for cross-interface transfer to WO (as proven by in situ Raman).\n2 17 3\nact as a proton acceptor. Consequently, this new mechanism directly enables the\nexceptional acidic HER performance, achieving a current density of 10 mA cm−2 with\nan ultralow overpotential of 21 mV.\nproduction.\nTherefore, our work is not a simple application of a known synthesis to a different\npH. It represents a deliberate material design to enable a novel catalytic mechanism.\nWe have clarified this in the revised manuscript.\nComment 2: Non-interfacial hydrogen spillover has been reported previously (e.g.,\nNat. Commun. 2022, 13, 1189; Small 2025, 21, 2411883), with the latter\ndemonstrating higher HER activity than the present study using non-precious metals.\nThe authors should clarify what advantages the current catalyst provides compared\nto these previously reported systems.\nResponse 2: Thanks for raising this important point regarding the advantages of our\nsystem. The previously reported non-interfacial H-spillover systems are commendable\nbut face inherent limitations in cost, generalizability, and mechanistic clarity. Our\nNi W -WO catalyst addresses these challenges head-on, offering a distinct and\n17 3 2\nadvantageous pathway.\nThe table below provides a direct comparison, highlighting our unique contributions.\nLa Sr PtO (Dai et al.43, Nat. High-Entropy Oxysulfides (HEOC) (Jo\n2 2 7+δ\nFeature Our Work: Ni W -WO\n17 3 2\nCommun. 2022, 13, 1189) et al.44, Small 2025, 21, 2411883)\nComplex oxide reliant on noble Complex 7-element high-entropy Simple, cost-effective binary (Ni-W)\nMaterial system\nmetal Pt(IV) system (Co, Ni, Mo, W, O, Se, Te) heterostructure\nDefined by a unique crystal Mechanistically ambiguous due to Clear, quantifiable pathway enabled by a\nMechanistic clarity\nstructure, not generalizable random site distribution designed strain/electron gradient\nProves the single-phase spillover Eliminates interfaces via Engineers a single-phase pathway within\nKey innovation\nconcept compositional complexity a heterostructure via strain/electronics\nPerformance 13 mV @ 10 mA cm−2 52 mV @ 10 mA cm−2 21 mV @ 10 mA cm−2\nSuperior activity, cost-effectiveness,\nAdvantage Conceptual proof Interface elimination\ngeneralizability, & mechanistic insight\nlimitations of its predecessors. It achieves higher activity than the HEOC system\nwithout relying on complex and multi-element compositions and demonstrates a clear\nand tunable non-interfacial mechanism using only earth-abundant elements,\novercoming the noble-metal dependency and specificity of La Sr PtO . Most\n2 2 7+δ\nimportantly, we demonstrate that a heterostructure can be cleverly engineered to\ncreate and regulate a non-interfacial pathway. This offers a generalizable principle for\ncatalyst design.\nWe have revised the manuscript to include this comparative perspective and to\nbetter articulate these advantages.\nComment 3: The d-spacing reported in Fig.1h deviates from the known value of the\nNi W (111) plane. Moreover, the SAED pattern shown in Fig.1g is insufficiently clear\n17 3\nto confirm the crystal planes. More rigorous structural characterization is therefore\nrequired to substantiate the phase identification.\nResponse 3: Thank you for your meticulous observation regarding the structural\ncharacterization. We agree that rigorous phase identification is paramount. In this\nwork, we have performed a comprehensive and multi-technique structural analysis,\nwhich collectively and unequivocally confirms the formation of the Ni W -WO\n17 3 2\nheterostructure.\n(1) Phase composition confirmed by XRD\nThe XRD pattern (Fig. R12a) shows two distinct sets of diffraction peaks, which\nexclusively match the standard patterns for face-centered cubic Ni W (JCPDS#65-\n17 3\n4828) and monoclinic WO (JCPDS#32-1393). No impurity phases are detected,\n2\nconfirming that the sample is purely composed of these two phases.\ndiffraction pattern of the Ni W -WO heterostructure.\n17 3 2\n(2) Lattice spacing identified by HAADF-STEM\nThe HAADF-STEM image of Ni W -WO shows lattice fringes with spacings of\n17 3 2\nd =0.35 nm and d =0.25 nm, which are attributed to the (011) and (020) planes of the\n1 2\nWO phase, respectively (Fig. R12b or Supplementary Fig. 1g). The angle between\n2 2 2\nWhile the observed lattice fringe d =0.24 nm, corresponding to the (111) planes of\n3\nthe Ni W phase, is found to be larger than its standard value of 0.21 nm. The reviewer\n17 3\ncorrectly notes the deviation in the d-spacing of the Ni W (111) plane. This is not a\n17 3\nmeasurement error but rather direct evidence of significant tensile strain within the\nNi W phase induced by the heterointerface with WO . This experimental observation\n17 3 2\nis fully consistent with and provides the foundation for our subsequent Geometric\nPhase Analysis (GPA) (Figs. 1i, j) and the proposed strain-gradient effect that is central\nto our mechanism.\n(3) Crystallographic planes validated by FFT Analysis\nTo address the clarity of the SAED pattern, we performed a Fast Fourier Transform\n(FFT) on an aberration-corrected HR-TEM image of the interface, which provides a\nclearer, localized diffraction pattern (Fig. R12c or Fig. 1h). The FFT pattern can be\nindexed as the superposition of the Ni W [011] zone axis and the WO [001] zone axis.\n17 3 2\nThe combination of XRD (bulk phase purity), HR-TEM (direct lattice imaging), and\nFFT (crystallographic indexing) provides a robust and self-consistent body of evidence\nthat conclusively identifies the material as a Ni W -WO heterostructure. The\n17 3 2\n24\n\"deviant\" lattice spacing is, in fact, a key piece of evidence for the interfacial\nengineering we describe. We have clarified this point in the revised manuscript.\nComment 4: In the DFT calculations, the hydrogen adsorption energies were\nevaluated at arbitrary sites without distinguishing between Ni and W atoms in the\nNi W phase. To properly elucidate the hydrogen spillover mechanism, the authors\n17 3\nshould conduct site-specific adsorption analyses to clarify the distinct roles of Ni and\nW in hydrogen adsorption and migration within the Ni W phase.\n17 3\nResponse 4: Thank you for this crucial suggestion regarding the precision of our DFT\nmodel. We fully agree that distinguishing the specific roles of Ni and W atoms is\nessential to elucidating the spillover mechanism. In direct response, we have\ncompletely rebuilt our computational model to enable site-specific analysis.\nWe replaced the previously used virtual crystal approximation with a realistic\nsupercell model that explicitly distinguishes all Ni and W atoms within the Ni W alloy\n17 3\nstructure (Fig. R13a or Supplementary Fig. 31a), and subsequently, the Ni W -WO\n17 3 2\nheterostructure was built based on HAADF-STEM results (Fig. R13b). This allows for\nthe precise calculation of hydrogen adsorption energies (ΔG ) at distinct Ni and W\nH*\nsites.\nFig. R13 | (a) The initial unit cell model of Ni W constructed via the supercell method.\n17 3\n(b) Theoretical model for the calculation of Ni W -WO heterostructure.\n17 3 2\nUsing this refined model, which also incorporates the experimentally derived strain\n25\nfield from GPA, we calculated the ΔG values across the Ni W phase (Figs. R14a, b\nH* 17 3\nor Figs. 5d, e). The results reveal a clear landscape.\nNi sites: The ΔG at Ni sites (labeled N1, N2, N3) shows a progressive weakening\nH*\nfrom the bulk (−0.553 eV) to the interface (−0.199 eV), creating a natural \"proton\nhighway\" (pathway: N1→N2→N3) that drives hydrogen toward the interface for\nefficient desorption.\nW sites: The adjacent W sites (W1, W2) exhibit even stronger H* adsorption (ΔG\nH*\nof −0.612 eV and −0.309 eV, respectively). This enables them to act as efficient proton\nadsorption reservoirs, subsequently supplying protons to the primary Ni spillover\nchannel (pathways: W1→N1 and W2→N3).\nFig. R14 | (a) The ΔG for HER on Ni W -WO . (b) Schematic illustration of the H*\nH* 17 3 2\nadsorption models on various active sites and the proposed reaction pathways for\nNi W -WO .\n17 3 2\nThis site-specific analysis allows us to propose a more precise mechanism: The H-\nspillover within the Ni W phase is a cooperative process. The W atoms primarily\n17 3\nfunction as initial adsorption and reservoir sites, while the gradient of Ni sites forms\nthe main diffusion and desorption pathway. This synergy, fine-tuned by the strain and\nelectronic effects from WO , collectively establishes the highly efficient non-interfacial\n2\nthe resulting mechanistic insight, which significantly strengthens the theoretical\nfoundation of our work.\nComment 5: Experimental validation of the proposed hydrogen spillover mechanism\nis insufficient. The authors should therefore provide direct evidence of spillover,\nfor example through WO -based H WO coloration and KIE experiments using\n3 x 3\nD SO /D O.\n2 4 2\nResponse 5: Thank you for this excellent suggestion to provide direct experimental\nvalidation for the hydrogen spillover effect. Following your guidance, we have\nconducted both chemical and electrochemical probing experiments, which together\noffer compelling and direct evidence supporting our proposed mechanism.\n(1) Chemical probe: Tungsten blue experiment confirms H-spillover\nWe performed the classic tungsten blue test. When a physical mixture of our Ni W -\n17 3\nWO catalyst and WO was heated in a dilute H /N atmosphere (8% H , 180 °C), a\n2 3 2 2 2\nrapid and distinct color change from yellow to dark blue was observed within 30\nminutes (Fig. R15a or Supplementary Fig. 28). This is a definitive signature of the\nformation of H WO . Under identical conditions, pure WO showed no color change,\nx 3 3\nconfirming that it cannot dissociate H on its own. This result directly demonstrates\n2\nthat our Ni W -WO catalyst efficiently dissociates molecular H into hydrogen atoms,\n17 3 2 2\nwhich then spill over from the catalyst surface onto the adjacent WO particles, leading\n3\nto their reduction. This provides direct chemical evidence of the catalyst’s H-spillover\ncapability.\n17 3 2 3\nand pure WO (down), both before and after being treated for 30 minutes at 180 °C in\n3\na mixed N /H atmosphere (8% H ). (b) Polarization curves of Ni W -WO in 0.5 M\n2 2 2 17 3 2\nH SO solution and 0.5 M D SO solution (without iR correction), and (c) the\n2 4 2 4\nrelationship between kinetic isotope effect value and potential. (d) Bode plots of\nNi W -WO in 0.5 M H SO /H O solution and 0.5 M D SO /D O solution.\n17 3 2 2 4 2 2 4 2\n(2) Electrochemical probe: Kinetic Isotope Effect (KIE) confirms proton-coupled RDS\nTo investigate the role of proton transfer in the electrochemical HER mechanism, we\nconducted KIE measurements by replacing H SO /H O with D SO /D O. Polarization\n2 4 2 2 4 2\ncurves indicate the HER current density in the deuterated electrolyte is significantly\nlower than in the protonated one (Fig. R15b or Fig. 4d). The calculated KIE value\nexceeds 2.0 across a wide potential range (Fig. R15c or Fig. 4e), which is a clear\nindicator that a proton-transfer step is the rate-determining step (RDS) of the\nreaction45. Furthermore, the Bode plot showed a characteristic frequency shift to\n28\nlower values in D SO /D O (Fig. R15d or Fig. 4f), confirming the slower kinetics\n2 4 2\noriginated from deuterium’s higher mass.\nIn summary, the combination of these two independent experiments confirms the\ncatalyst’s intrinsic ability to perform H-spillover, and the electrochemical HER on our\ncatalyst is indeed governed by a proton-transfer-limited process, which is the kinetic\nsignature of the H-spillover mechanism we propose. These new findings, now\nincluded in the revised manuscript, offer direct experimental proof that addresses the\nreviewer’s concern and solidly underpins our mechanistic claims.\nComment 6: Supplementary Fig.12 reveals evident performance degradation after\nthe 1500-hour durability test, which contradicts the authors' claims of long-term\nstability. Furthermore, to substantiate the chemical and structural stability of the\ncatalyst, post-durability characterization data (e.g., XPS, XRD, and SEM) should be\nprovided.\nResponse 6: Thank you for this critical observation. The slight performance\ndegradation observed in Supplementary Fig. 12 (the revised Supplementary Fig. 20)\nis indeed real. We quantify this decay rate as an exceptionally low 28 µV h −1. In the\ncontext of acidic HER for non-noble metal catalysts, where dissolution is a fundamental\nchallenge, this level of degradation over such an extended period is highly competitive\nand underscores outstanding operational stability 46-47. Our post-mortem analysis\nreveals that this minor decay results from superficial, self-limiting processes that\nultimately protect the core catalytic structure, which remains remarkably intact.\nTo understand the origin of this minor decay and assess the catalyst’s structural\nintegrity, we performed detailed characterization after the 1500-hour test. XRD in Fig.\nR16a (Supplementary Fig. 21) confirms the preservation of both the Ni W and WO\n17 3 2\ncrystalline phases. The slight attenuation of the Ni W peak intensity points to\n17 3\nminimal, superficial dissolution, which is the likely cause of the observed\nactivity decay and is expected for any non-noble metal in acid over 1500 hours. TEM\nand HR-TEM images (Figs. R16b, c or Supplementary Fig. 22) provide direct visual\nevidence of stability. The material retains its granular morphology, and the clear lattice\n29\nfringes and, most importantly, the distinct heterointerface remain visible, proving the\ncore catalytic architecture survives.\nFig. R16 | (a) XRD patterns of Ni W -WO before and after long-term stability testing.\n17 3 2\n(b) TEM and (c) HR-TEM images of Ni W -WO after long-term stability testing.\n17 3 2\nThe surface chemical states of post-tested Ni W -WO were thoroughly\n17 3 2\ninvestigated using XPS. In the high-resolution W-4f XPS spectrum, peaks at 32.6/34.4\neV are attributed to W4+ in the WO phase, while those at 35.0/37.3 eV and 35.9/37.9\n2\neV are assigned to W5+ and W6+, respectively (Fig. R17a or Supplementary Fig. 23a).\nCompared to the pre-tested sample, the W-4f spectrum shows an increased\nproportion of W5+/W6+, indicating surface oxidation of the WO phase and the\n2\ndissolved W species from Ni W , forming a passivated surface layer. As shown in Fig.\n17 3\nR17b (Supplementary Fig. 23b), the high-resolution Ni-2p XPS spectra of sample\nNi W -WO after long-term testing exhibited characteristic peaks at 852.9/870.3 eV,\n17 3 2\n856.1/874.0 eV, and 858.6/876.2 eV, corresponding to Ni0, Ni2+, and Ni3+, respectively.\nCritically, the Ni-2p spectrum still shows a clear Ni0 signal, demonstrating that metallic\nNi active sites within the bulk of the Ni W phase are effectively preserved.\n17 3\n30\na b\nW5+\nW5+\nNi2+\nu . ) u . )\na . a .\ns i t y\n(\ns i t y\n(\nsat. Ni3+ Ni2+ sat.\nn n\ne W6+ W6+ e Ni3+\nI n t W 5p W4+ W4+ I n t Ni0 Ni0\nafter test after test\n44 40 36 32 880 870 860 850\nBinding energy (eV) Binding energy (eV)\nFig. 17 | High-resolution XPS (a) W-4f and (b) Ni-2p of Ni W -WO after long-term\n17 3 2\nstability testing.\nTherefore, the minor performance degradation is attributed to a self-limiting\nprocess of superficial dissolution and surface oxidation, which leads to a passivation\nlayer. However, the core heterostructure, the metallic Ni active sites, and the interfacial\nregion remain largely preserved. Therefore, we maintain that the catalyst exhibits\n“excellent long-term stability” as its essential structure and function are maintained\nfor 1500 hours under harsh conditions. We have revised the manuscript to include this\ndiscussion and all post-stability characterization data.\nComment 7: In Fig.4h, the disappearance of the Ni-O peak suggests H* adsorption at\nNi sites. To clarify the spillover pathway, additional in situ Raman measurements at\nmore negative potentials are needed to detect possible M-H vibration peaks as direct\nevidence of H* accumulation.\nResponse 7: Thank you for this excellent suggestion, which has allowed us to obtain\ndirect spectroscopic evidence for the H-spillover pathway. Following the\nrecommendation, we extended our in situ Raman measurements to more negative\npotentials, and the results successfully captured the key spectral signatures of the\nproposed mechanism.\nWe extended the in situ Raman spectroscopic analysis to a potential range from 0\nto −600 mV vs. RHE. At the initial state, characteristic peaks at 515 cm–1 and 603 cm–1\n31\nare attributed to the vibrational modes of Ni-O bonds at the Ni W -WO\n17 3 2\nheterointerface (Fig. R18 or Fig. 4h). Additionally, the peak observed at 780 cm–1 is\nattributed to the asymmetric stretching vibration of W-O bonds, resulting from\nsymmetry breaking or lattice distortions at the heterointerface 3. When the potential\nis scanned cathodically from 0 V to −250 mV vs. RHE, we observe the gradual\ndisappearance of the Ni-O vibrational peaks. This indicates the progressive reduction\nof the interface and the freeing of Ni sites. Critically, upon further lowering the\npotential to −300 mV and beyond, a new peak emerges at 980 cm–1, which is\nunequivocally assigned to the Ni-H vibration 4. The intensity of this Ni-H peak grows\nsteadily as the potential becomes more negative. This temporal sequence in the\nRaman spectra provides direct, real-time evidence for the spillover pathway. The\ndisappearance of Ni-O peaks signifies the activation of Ni sites. The subsequent\nappearance and growth of the Ni-H peak directly confirm that atomic hydrogen (H*)\naccumulates on these interfacial Ni sites during HER.\nNi-O W-O (asymmetric) Ni 17 W 3 -WO 2\nNi-H\n0 mV\n)\nu . -50 mV\n.\na -100 mV\n(\ny -150 mV\ni t -200 mV .\ns\nn -250 mV\ne\nt -300 mV\nn\nI -400 mV\n-500 mV\n-600 mV\n600 800 1000 1200\nRaman shift (cm-1)\nFig. R18 | In situ Raman scattering of Ni W -WO electrodes recorded at different\n17 3 2\napplied cathodic potentials.\nComment 8: The manuscript contains several inconsistencies, such as inconsistent\npeak labeling for identical analytical results and spelling errors in axis labels. A\nthorough revision is necessary to improve the clarity, consistency, and overall quality\n32\nof the manuscript.\nResponse 8: We sincerely thank you for your meticulous review of the manuscript.\nWe deeply apologize for the oversights in labeling consistency and spelling that detracted\nfrom the clarity of our presentation. Regarding Fig. 2e, the x-axis label's spelling has\nbeen rectified, and consistent color assignments for identical peaks have been\nimplemented during the revision, as illustrated in Fig. R19 (Fig. 2e). Additionally, in\nthis updated version, we have conducted a thorough, line-by-line proofreading of the\nentire manuscript and figures.\n.\nNi W\n17 3\n)\n. u\n. a\n( .\ny Ni 17 W 3 -WO 2 W5+ W4+\nt\ni\ns\nn W5+ W4+\ne WW66++\nn t W 5p\nW6+\nI W0 W0\n0.2 eV\nWO\n2\n44 40 36 32\nBinding energy (eV)\nFig. R19 | High-resolution XPS W-4f spectra of Ni W , Ni W -WO , and WO .\n17 3 17 3 2 2\nReviewer #3 (Remarks to the Author): In this manuscript, the authors report the\nsynthesis of a Ni W –WO heterostructure and its application as a non-conventional,\n17 3 2\nnon-interfacial hydrogen spillover catalyst for promoting the hydrogen evolution\nreaction. The study highlights the modulation of hydrogen adsorption and\ndesorption mechanisms by WO , with Ni W proposed as active sites contributing\n2 17 3\nto enhance HER performance. The conceptual framework is supported by theoretical\ninvestigations, and the mechanistic details are elaborated in the manuscript.\nHowever, several points require clarification. Therefore, I recommend a major\nrevision of the manuscript to address the following comments.\nGeneral Response: We sincerely thank you for your positive assessment of our work’s\nconceptual framework and mechanistic elaboration, and for the constructive\nrecommendations for major revision. We are pleased to report that we have\ncomprehensively addressed all points raised. Additionally, we have provided a detailed\npoint-by-point response below and believe the work now robustly meets the high\nstandards of Nature Communications.\nComment 1: The authors assert that the morphology and particle size of NiWO\n4\nprecursor remain unchanged after annealing. However, as evidenced by the high-\nresolution images of NiWO precursor before (Figure S1b) and after annealing\n4\n(Ni W –WO ) (Figure 1c), the morphology is largely similar, but the particle size has\n17 3 2\nincreased. Specifically, particle agglomeration has occurred, leading to larger\nparticle sizes. The authors should recognize this observation and update the\nmanuscript accordingly.\nResponse 1: Thank you for this careful observation. We have performed a statistical\nanalysis of particle sizes from the SEM images to address this point quantitatively.\nThe results confirm the reviewer's assessment: the average particle size increased\nfrom 30.3 nm for the NiWO precursor (Fig. R20a or Supplementary Fig. 1c) to 38.9\n4\nnm for the Ni W -WO heterostructure after annealing (Fig. R20b or Supplementary\n17 3 2\nFig. 2). We attribute this increase to partial particle agglomeration during the high-\n34\ntemperature treatment.\nWe have revised the manuscript text accordingly, and the updated description now\nreads: “Field emission scanning electron microscopy (FE-SEM) characterization reveals\nthe morphological evolution from precursor to final product. The NiWO precursor\n4\ncomprises uniform nanoparticles with an average size of 30.3 nm (Supplementary Fig.\n1). Following annealing in a N /H atmosphere, the nanoparticulate morphology is\n2 2\nlargely retained, although the average particle size increases to 38.9 nm due to partial\nagglomeration during the thermal process (Fig. 1b-c and Supplementary Fig. 2).”\n30 30\na b\nNiWO Ni W -WO\nAD = 30.3 nm 4 17 3 2\n) )\n% %\n(20 (20\ne e\ng g\na a15\nt t\nn n\ne e\nc10 c10\nr r\ne e\nP P\n5\n0 0\n20 30 40 50 20 30 40 50 60\nDiameter (nm) Diameter (nm)\nFig. R20 | Particle size distribution of (a) NiWO and (b) Ni W -WO obtained from\n4 17 3 2\nSEM image statistics.\nComment 2: The amount of evolved hydrogen is reported based on both\nexperimental measurements and theoretical calculations. Additional details\nregarding the methodology used for the theoretical determination of hydrogen\nevolution should be provided to ensure clarity and reproducibility. Furthermore, it is\nobserved that the experimental and theoretical hydrogen evolution values at each\ntime interval are remarkably consistent, which is somewhat unusual. The authors\nare encouraged to include specific numerical values (for both experimental and\ntheoretical results) to facilitate verification and to substantiate the accuracy of the\ndata presented.\nResponse 2: Thank you for raising these important points regarding methodological\n35\nclarity and data verification. Herein, we provide a detailed description of the Faradaic\nefficiency (FE) measurement protocol and the complete dataset to ensure full\ntransparency and reproducibility.\n(1) Experimental methodology for hydrogen measurement\nThe experiment was conducted at a constant temperature of 25 °C, where the molar\nvolume of H is 24.45 mL mmol−1. Gaseous products from the cathode were collected\n2\nin a 100 mL graduated cylinder via the water displacement method. Before\nmeasurement, the system was purged by applying a constant current of −0.1 A until\nthe meniscus stabilized at the 10 mL mark. The formal measurement was then\nconducted in galvanostatic mode at −0.1 A for 80 min. Gas volume was recorded at 10-\nminute intervals. The graduated cylinder has a minimum division of 1 mL, and volumes\nwere read with an estimated precision of ±0.5 mL through interpolation.\n(2) Calculation procedures\nThe FE was calculated using the formula:\nFE=\nn(H2 )\nM×100% (Eq. 1)\nn(H2 )\nT\nwhere n(H ) (mmol) is the molar mass of the experimentally measured amount of\nhydrogen, calculated from Eq. 2:\nn(H ) =\nVL\n(Eq. 2)\nVm\nwhere V (mL) is the volume of hydrogen produced, V is the molar volume, with a\nL m\nvalue of 24.45 mL mmol−1 at 298.15 K and 1.013 × 105 Pa.\nThe n(H ) (mmol) is the theoretical amount of hydrogen, calculated as\ni × t\nn(H ) = ×1000 (Eq. 3)\n2 × F\nwhere i (A) is the current, t (s) is the time duration, and F is Faraday’s constant (96485\nC mol−1). FE is calculated at an applied current of −0.1 A for 80 min.\nThe complete experimental readings and calculations are provided in Table R5.\nTime (min) V (mL) n(H ) (mmol) n(H ) (mmol) FE (%)\nL 2 M 2 T\n10 7.5 0.3067 0.3109 98.65\n20 15 0.6135 0.6219 98.66\n30 22 0.8998 0.9327 96.46\n40 30 1.2269 1.2437 98.66\n50 38 1.5542 1.5546 99.97\n60 45 1.8405 1.8656 98.66\n70 53 2.1677 2.1765 99.59\n80 60 2.4540 2.4874 98.65\nComment 3: The hydrogen production rates and Faradaic efficiencies for the\nindividual Ni W and WO components should also be calculated. Including these\n17 3 2\nparameters will enable a clearer comparison with the heterostructure and help\ndemonstrate the potential superiority of the interface in enhancing hydrogen\nevolution activity.\nResponse 3: Thank you for your valuable suggestion. To directly evaluate the role of\neach component and the interface, we measured the hydrogen production rates and\nFaradaic efficiencies (FE) for the individual WO , Ni W , and the Ni W -WO\n2 17 3 17 3 2\nheterostructure.\nGas chromatography analysis confirmed that the gaseous product from all three\ncatalysts was exclusively H (a single peak at ~1.5 min retention time; Figs. R21a-c or\n37\n3\na d\nm\n) \nu . m 80\n. 2\ns\ni t y (\na\nH 2 100% o f H\n2\n1.82 m mol\nh-1\nE %\nn t\nn 40\nt e u 1\nn o\nI m WO\n2\n/CC\nWO /CC Theoretical H yield\n2 0 2 0\n0 3 6 9 12 15 0 20 40 60 80\nRetention (min) Time (min)\n3\nb e\n2 o\n) m\nu . m 80\n. 2\nn\ns\ni t\ny (\na\nH\n2\n100%\nt\no f\nH 2\n1.85\nm mol\nh-1\nE\n%\nt e u n 1 40\nn o\nI m Ni 17 W 3 /CC 20\nA Experimental H yield\n0 3 6 9 12 15 0 20 40 60 80\nRetention (min)\nTime (min)\n3\nc f\n100\nl\nH o\n) 2 m\nu . m 80\n. 2\nn\ns\ni t y ( a H 2 100%\nt\no f H 2 1.85 m mol h-1 60\nF\nE %\nt e u n 1 40\nn o Ni W -WO /CC\nI m 17 3 2\nNi W -WO /CC Theoretical H yield\n17 3 2 2\n0 0\n0 3 6 9 12 15 0 20 40 60 80\nRetention (min) Time (min)\nFig. R21 | Gas chromatography analysis of the gaseous product from (a) WO /CC, (b)\n17 3 17 3 2 2\nproduced by (d) WO /CC, (e) Ni W /CC, and (f) Ni W -WO/CC, and FE at a current of\n2 17 3 17 3\n−0.1 A.\nUnder identical conditions (−0.1 A, 25 °C), the performance metrics were\ndetermined (Figs. R21d-f or Supplementary Figs. 17b, 19) and listed in Table R6.\nHydrogen production\nCatalysts FE (%)\nrate (mmol h−1)\nWO 1.82 97.83\n17 3\nNi W -WO 1.85 98.65\n17 3 2\nThe data reveal that the nearly identical FE and production rate between Ni W and\n17 3\nthe Ni W -WO heterostructure unequivocally demonstrates that the Ni W phase is\n17 3 2 17 3\nthe primary active site for HER. The slightly lower FE of pure WO is attributed to partial\n2\nproton intercalation into its structure, a known phenomenon where some protons are\nstored rather than converted to H 2. The superiority of the heterostructure lies not in\n2\ncreating new active sites on WO or altering the intrinsic FE of Ni W , but in the\n2 17 3\nelectronic and strain engineering that drastically enhances the specific activity of the\nNi W phase.\n17 3\nComment 4: The authors have calculated the mass activity of Ni W –WO and\n17 3 2\ncompared it with that of Ni W and WO . However, the data points corresponding\n17 3 2\nto the measurement are missing in Figure 3a. The authors should revise this figure\ncarefully to accurately include all relevant data points for proper comparison.\nResponse 4: Thanks for this meticulous observation. The reviewer is correct to point\nout that the original presentation in Figure 3a could lead to misinterpretation. We\nsincerely apologize for any confusion caused by our initial design.\nThe intention was to display the intrinsic relationship between geometric current\ndensity and mass activity (Fig. R22a), which is linearly proportional for a fixed catalyst\nloading on a fixed electrode area. As the reviewer suspected, each data point in the\noriginal figure did indeed correspond to both a current density (left y-axis) and a mass\nactivity (right y-axis), with the axes scaled according to the constant conversion factor\nderived from the electrode area and mass loading. However, we fully agree with the\n39\nreviewer that this dual-axis representation is not the clearest way to present the data\nfor direct comparison. The revised Fig. 3a (Fig. R22b) now shows the LSV curves\n(current density) for Ni W -WO , Ni W , and WO . A new supplementary figure\n17 3 2 17 3 2\nclearly displays the corresponding mass activity curves (Fig. R22c or Supplementary\nFig. 11).\n-2\nm\n) \nA\nc m\n-50\n0 a 0.0\n-1\nA\n) \nm\ng\n-2\nm\n) \nA\nc m\n-50\n0 b\n-1\nA\n) \nm\ng\n-0\n0.\n.\n0\n05\n0 c\n(\nurr e nt d\ne n\nsit y\n-\n-\n1\n1\n5\n0\n0\n-\n0\n0\n.\n.\n2\n(\na s s a\ncti\nvit y (\nurr e nt d\ne n\nsit y\n-\n-\n1\n1\n5\n0\n0\n(\na s s a\ncti\nvit y\n-\n-\n-\n0 0\n0\n. . \n. \n2 1\n1\n0 5\nC-200 C-200\n-0.6 -0.4 -0.2 0.0 -0.6 -0.4 -0.2 0.0 -0.6 -0.4 -0.2 0.0\nPotential (V vs. RHE) Potential (V vs. RHE) Potential (V vs. RHE)\nFig. R22 | (a) LSV curves and corresponding mass activity curves of WO , Ni W , and\n2 17 3\nNi W -WO in 0.5 M H SO electrolyte. (b) LSV curves and (c) mass activity curves of\n17 3 2 2 4\nWO , Ni W , and Ni W -WO in 0.5 M H SO electrolyte.\n2 17 3 17 3 2 2 4\nComment 5: In addition to LSV curves of Ni W –WO /CC before and after long-term\n17 3 2\ntests, the structural stability of the catalyst be confirmed through structural\n(XRD)/morphological (TEM)/elemental analysis before and after long term tests.\nWhile the authors mention that the LSV curves after 1500 hours are nearly identical,\na slight shift is nonetheless observed. The authors should address this discrepancy in\nthe revised manuscript and discuss the possible underlying reasons for the observed\nshift in the LSV curves.\nResponse 5: Thank you for this comprehensive suggestion. We have conducted a\ndetailed investigation into the slight performance shift and performed a full suite of\npost-stability characterizations, which together provide a complete picture of the\ncatalyst’s exceptional durability.\nAs noted by the reviewer, a minor performance degradation is observed after the\n1500-hour test at 500 mA cm−2, with the overpotential increasing from –0.138 V to –\n0.180 V (vs. RHE), corresponding to a very low average decay rate of 28 µV h−1 (Fig.\nR23a or Fig. 3g). This is reflected in the slight shift of the LSV curve at high current\n40\ndensities (Fig. R23b or Supplementary Fig. 20). Our post-mortem characterization\nreveals that this minimal decay originates from superficial changes that do not\ncompromise the core structure.\n)\nR\nH E\n0.0\na\n-2)\nA\nc m\n0\nb\nP\n(\nV\nv s\no\n.\nt\ne\nnti\nal\n-\n-\n-\n0\n0\n0\n.\n.\n.\n3\n2\n( m\nurr e\nnt\nd e\nn\nsit\ny\n-\n-\n1\n1\n-5\n5\n0\n0\n0\n00\n0\n0\nb a e fte fo r r t e e s te t st\n0 300 600 900 1200 1500 -0.6 -0.4 -0.2 0.0\nTime (h) Potential (V vs. RHE)\nFig. R23 | (a) stability test of Ni W -WO /CC. (b) LSV curves of Ni W -WO /CC before\n17 3 2 17 3 2\nand after the long-term test.\nTo thoroughly investigate the reasons behind the slight performance degradation of\nNi W -WO /CC, we conducted XRD, TEM, and XPS analyses on samples after 1500\n17 3 2\nhours of long-term stability testing. XRD confirms the preservation of both Ni W and\n17 3\nWO phases (Fig. R24a or Supplementary Fig. 21). The slight attenuation of the Ni W\n2 17 3\npeak intensity indicates minimal dissolution. TEM and HR-TEM images show that the\ngranular morphology (Fig. R24b or Supplementary Fig. 22a) and, most importantly,\nthe heterointerface between Ni W and WO remain intact (Fig. R24c or\n17 3 2\nSupplementary Fig. 22b).\nFig. R24 | (a) XRD patterns of Ni W -WO before and after long-term stability testing.\n17 3 2\n(b) TEM and (c) HR-TEM of Ni W -WO after long-term stability testing.\n17 3 2\n17 3 2\ninvestigated using XPS. In the high-resolution W-4f XPS spectrum, peaks at 32.6/34.4\neV are attributed to W4+ in the WO phase, while those at 35.0/37.3 eV and 35.9/37.9\n2\neV are assigned to W5+ and W6+, respectively (Fig. 25a). Compared to the pre-tested\nsample, the W-4f spectrum shows an increased proportion of W5+/W6+, indicating\nsurface oxidation of the WO phase and the dissolved W species from Ni W , forming\n2 17 3\na passivated surface layer. As shown in Fig. 25b (or Supplementary Fig. 23a), the high-\nresolution Ni-2p XPS spectra of sample Ni W -WO after long-term testing exhibited\n17 3 2\ncharacteristic peaks at 852.9/870.3 eV, 856.1/874.0 eV, and 858.6/876.2 eV,\ncorresponding to Ni0, Ni2+, and Ni3+, respectively. Critically, the Ni-2p spectrum still\nshows a clear Ni0 signal, demonstrating that metallic Ni active sites within the bulk of\nthe Ni W phase are effectively preserved.\n17 3\na b\nW5+\nW5+\nNi2+\nu . ) u . )\na . a .\ns i t y\n(\ns i t y\n(\nsat. Ni3+ Ni2+ sat.\nn n\ne W6+ W6+ e Ni3+\nI n t W 5p W4+ W4+ I n t Ni0 Ni0\nafter test after test\n44 40 36 32 880 870 860 850\nBinding energy (eV) Binding energy (eV)\nFig. R25 | (a) W-4f and (b) Ni-2p high-resolution XPS spectra of Ni W -WO after long-\n17 3 2\nterm stability testing.\nThe collective evidence supports that the catalyst undergoes a self-limiting, surface-\nlevel passivation in the harsh acidic environment. This leads to the minor, quantifiable\nperformance decay. The core heterostructure remains remarkably stable, enabling the\nunprecedented 1500-hour durability.\n17 3 2\ncatalyst, investigated via CV analysis, show that Ni W –WO has a more negative\n17 3 2\nhydrogen desorption potential compared to Ni W , while its hydrogen insertion\n17 3\npotential lies between those of Ni W and WO . These findings appear to contradict\n17 3 2\nthe assumption of stronger hydrogen adsorption for enhanced hydrogen evolution\nperformance attributed to Ni W –WO . Since the core concept of the manuscript\n17 3 2\nrevolves around hydrogen spillover to improve hydrogen generation, the authors are\nrecommended to provide a clear discussion addressing this apparent inconsistency\nin the revised manuscript.\nResponse 6: Thanks for your insightful comment. The apparent contradiction arises\nfrom the difference between absolute adsorption strength and the optimization of the\nadsorption energy landscape. Our strategy is not about creating the strongest\nadsorption, but about creating the most favorable pathway from adsorption to\ndesorption.\nAs the reviewer correctly infers from the CV data, pure Ni W has the most positive\n17 3\nhydrogen adsorption potential, indicating the strongest hydrogen binding. This is\ntypical for many non-noble metals and is precisely why their HER performance is poor-\nthe strong binding stifles the desorption step (the Heyrovsky/Tafel step), creating a\nhigh energy barrier for H release.\n2\nadsorption, but to precisely moderate it and create a spillover pathway. The Ni W -\n17 3\nWO has an adsorption potential that is weaker than pure Ni W but stronger than\n2 17 3\nWO (Fig. R26 or the revised Fig. 4a). This indicates a weakening of the hydrogen\n2\nadsorption strength on the Ni W phase due to the electronic and strain effects,\n17 3\nmoving it closer to the ideal value of zero. More importantly, the more negative\ndesorption potential for Ni W -WO directly demonstrates that the energy barrier for\n17 3 2\nhydrogen release is significantly lower than that on pure Ni W .\n17 3\n43\n0.40 0.40\n)\nE\nH adsorption\nR\n0.35 0.38\ns . desorption\nv .\nV\n(\na\nl0.30\n0.36\ni\nt\nn\ne N/A\nt\no\nP\n0.00 0.00\nWO\n3\n-WO\n2 17 3 17 3\nWO .\n17 3 2\nbut due to a more efficient overall process. We have transformed the catalyst from a\nmaterial with strong adsorption but poor desorption (Ni W ) into one with\n17 3\nexceptionally facile desorption (Ni W -WO ). The non-interfacial H-spillover pathway\n17 3 2\nensures that protons can readily migrate to the optimal sites for this low-barrier\nrelease.\nIn conclusion, the CV data does not contradict our claim. It empirically shows that\nour heterostructure engineering successfully weakens the overly strong H-binding on\nthe Ni W phase and simultaneously drastically lowers the desorption barrier, which\n17 3\nis the fundamental reason for the enhanced HER activity. We have revised the\nmanuscript to include this clear discussion.\nComment 7: The manuscript does not include a comparison with a standard\ncommercial catalyst, such as Pt/C. Incorporating data for a well-established\nreference material would provide valuable context and facilitate meaningful\ncomparisons of catalytic performance. The authors could refer Adv. Energy Mater.\n2024, 14, 2401426 while revising the manuscript\nResponse 7: Thank you for this essential suggestion. Following the recommendation\nand citing the provided reference (Adv. Energy Mater. 2024, 14, 2401426), we have\n44\nnow included a comprehensive benchmarking of our Ni W -WO /CC catalyst against\n17 3 2\ncommercial Pt/C and a wide range of state-of-the-art catalysts.\nAs detailed in the Fig. R27 (the revised Fig. 3f) and Table R7 (or Supplementary\nTable 1), the comparison spans catalyst loading, overpotential at 10 mA cm−2, and Tafel\nslope. The data demonstrates that our catalyst’s performance is highly competitive.\nThe overpotential and Tafel slope are comparable to those of commercial Pt/C and\nadvanced Pt-based catalysts. It significantly outperforms Ru-based catalysts and other\nnon-precious metal systems.\nWe acknowledge the reviewer’s point regarding mass activity, where a gap with\nultra-low loading Pt-based catalysts remains. However, the entirely non-precious\ncomposition of our material results in an exceptional cost-to-performance ratio, which\nis a paramount advantage for large-scale industrial implementation. We have\nintegrated this comparative analysis into the revised manuscript to provide the\nnecessary context and underscore the practical potential of our work.\nFig. R27 | Comparison of the overpotential at 10 mA cm−2, Tafel slope, and catalyst\nloading for the Ni W -WO /CC against Pt-based, Ru-based, and other non-precious\n17 3 2\nmetal catalysts in 0.5 M H SO .\n2 4\nTable R7. Comparison of the overpotentials (at 10 mA cm−2), Tafel slopes, and catalyst\n45\nloading of different catalysts in 0.5 M H SO .\n2 4\nOverpotential Tafel slope Loading\nCatalysts Ref.\n(mV) (mV dec−1) (mg cm−2)\nNi W -WO /CC 21 47.8 0.914 This work\n17 3 2\n20% Pt/C 36 35 0.51 5\nRu 67 80.2 0.16 6\nPt /OLC 38 36 0.51 5\nPt@WS 31 27.2 0.256 8\nCoNiPt @G 23 55 1.6 10\nSA\n2H/1T-WS -10 113 55 0.52 11\n2\nhcp-RuCu HUNSs 98 75 0.1 12\nRu/CNT 27 59.6 0.34 13\nRu-Co Ni 57 52 1.2 14\nRu/Pd Cu-70 28 46 1 18\n1 1\nRuSe 59 52 0.25 20\n2\n2\np+-WS -9.8 92 33 0.52 22\nMoS -NbS 159 53 1.41 24\n2-x x\nPb-CoSe -DETA 74 45 1.21 25\n0.1 0.9 2\nComment 8: The Nyquist plot of Ni W –WO is not clearly visible in the Figure 4e.\n17 3 2\nThe enlarged figure demonstrating the semicircular Nyquist plot should be moved\nfrom supporting information to main text in the inset of Figure 4e.\nResponse 8: Thank you for this suggestion to improve the clarity of the presentation.\nWe have moved the enlarged view of the Nyquist plot for Ni W -WO from the\n17 3 2\nSupporting Information and inserted it as an inset in Fig. 4e (the revised Fig. 4b) of the\nmain text, as shown below (Fig. R28). This adjustment ensures the data is clearly visible\nto the reader within the primary context of our discussion on charge transfer\nresistance.\n2000\n80\n60\n1500\n40\n)\nm\n20\nh\no1000 0\n( 0 20 40 60 80 .\nZ''\n-\n17 3\nNi W -WO\n0 17 3 2\n0 500 1000 1500 2000\nZ' (ohm)\nFig. R28 | Nyquist plots of WO , Ni W , and Ni W -WO .\n2 17 3 17 3 2\nin the captions of Figure 5e, Figure S17, and S18.\nResponse 9: Thank you for pointing out this oversight. We have revised the captions\nfor Figure 5e (Fig. 5e), S17 (the revised Supplementary Fig. 30), and S18 (the revised\nSupplementary Fig. 31) to explicitly specify the color scheme used in the atomic\nmodels, as shown in Fig. R29. This clarification ensures that the structural models are\nunambiguous and easily interpretable.\nFig. R29 | Schematic illustrations depicting the structures of (a) WO and (b) Ni W .\n2 17 3\n(c) Schematic illustration of the H* adsorption models on various active sites and the\nproposed reaction pathways for Ni W -WO .\n17 3 2\n1. Bastians, S., Crump, G., Griffith, W. P. & Withnall, R. Raspite and studtite: Raman spectra of two unique\nminerals. J. Raman Spectrosc. 35, 726-731 (2004).\n2. Chen, Z. et al. Metallic W/WO solid-acid catalyst boosts hydrogen evolution reaction in alkaline\n2\nelectrolyte. Nat. Commun. 14, 5363 (2023).\n3. Ross-Medgaarden, E. I. & Wachs, I. E. Structural determination of bulk and surface tungsten oxides\nwith UV-vis diffuse reflectance spectroscopy and Raman spectroscopy. J. Phys. Chem. C 111, 15089-\n15099 (2007).\n4. Guo, P. et al. 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Pd@CuinP S core-shell nanospheres with exceptional hydrogen evolution capability\n2 6\nand stability in both alkaline and acidic media under large current density exceeding 1000 mA cm−2.\nSmall 20, 2403005 (2024).\n36. Miao, H. et al. Ultrasmall noble metal doped Ru P@Ru/CNT as high-performance hydrogen\n2\nevolution catalysts. ACS Sustainable Chem. Eng. 9, 15063-15071 (2021).\n37. Wu, Y. et al. Coupling Ir single atom with NiFe LDH/NiMo heterointerface toward efficient and\ndurable water splitting at large current density. Appl. Catal. B-Environ. Energy 360, 124548 (2025).\n38. Zhu, Z. et al. High-performance alkaline freshwater and seawater hydrogen catalysis by sword-head\nstructured Mo N-Ni Mo N tunable interstitial compound electrocatalysts. Adv. Funct. Mater. 34,\n2 3 3\n2306061 (2024).\n39. Zhou, X. et al. Engineering active iron sites on nanoporous bimetal phosphide/nitride\nheterostructure array enabling robust overall water splitting. Adv. Funct. 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High-entropy oxychalcogenide for hydrogen spillover enhanced\nhydrogen evolution reaction in proton and anion exchange membrane water electrolyzers. Small\n21, 2411883 (2025).\n45. Wang, Q. et al. 2D metastable-phase hafnium oxide triggers hydrogen spillover for boosting\nhydrogen production. Adv. Mater. 37, 2415978 (2025).\n46. Zeng, B. et al. Grafting ultra-fine nanoalloys with amorphous skin enables highly active and long-\nlived acidic hydrogen production. Angew. Chem. Int. Ed. 63, e202400582 (2024).\n47. Yang, H. et al. Integrated high-entropy alloy nanowire/carbon nanotube membrane electrode for\nefficient hydrogen evolution in acid solution. Adv. Funct. Mater. 35, 2425156 (2025).\nReviewer #4 (Remarks to the Author): After carefully reading the manuscript and the\nauthors’ response to reviewers’ comments, I believe the authors have in large part\nproperly addressed reviewer #2’s concerns by performing additional experiments\nand calculations. However, the response to specific comment #2 and #5 needs more\nattention.\nGeneral Response: We sincerely thank you for your rigorous and constructive\nevaluation. We are pleased that our previous revisions have largely addressed the\nconcerns raised. In direct response to the specific points highlighted (#2 and #5), we\nhave conducted substantial new experimental and analytical work to further\nstrengthen the mechanistic foundation and practical relevance of our study. The key\nenhancements in this revised manuscript are:\n(1) We have performed comparative KIE studies on the Ni W -WO catalyst as\n17 3 2\nwell as the control samples to better clarify the structure-to-mechanism correlation.\n(2) The Co-W-based alloy-oxide heterostructures have been synthesized via the\nsame in situ phase reconstruction protocol to demonstrate the generalizability of our\nstrategy.\n(3) To unequivocally validate the industrial potential, we have conducted\nadditional PEM electrolyzer test at an elevated temperature of 80 °C, a condition\ncloser to real-world operation.\n(4) Furthermore, we have thoroughly revised the title, abstract, and introduction\nto more precisely articulate our core innovation and to clearly define the indispensable\nrole of WO .\nand elevated the manuscript’s impact and clarity. We detail our point-by-point\nresponses below.\nComment 1: For comment #2, the authors claim that the current work is\nadvantageous over previous works about non-interfacial hydrogen spillover for its\n1\ngeneralizability. However, this is not obvious as the non-interfacial hydrogen\nspillover is only demonstrated on one single material of NiW-WO2. The generality of\nsuch design strategy in other types of materials remain questionable.\nResponse 1: Thank you for this critical question, which allows us to clarify and\nstrengthen the core finding of our work regarding generality. We agree with your\ncomment that a single material demonstration is insufficient to claim a general design\nstrategy. Therefore, we have undertaken additional experimental work to explicitly\ntest the generalizability of our proposed principle.\nOur central hypothesis is that the non-interfacial hydrogen spillover mechanism\nis enabled by heterointerface-induced strain gradients, not by a specific elemental\ncomposition or a well-defined crystalline structure. To validate this, we designed and\nsynthesized a distinct Co W -WO heterostructure using the same in situ phase\n0.87 0.13 2\nreconstruction protocol. The XRD diffraction peaks of the precursor can be fully\nindexed to the CoWO phase (JCPDS#72-0479) (Supplementary Fig. 32a). After\n4\nannealing under a reductive atmosphere (Supplementary Fig. 32b), the diffraction\npeaks can be assigned to a face-centered cubic Co W alloy phase (JCPDS#71-7508)\n0.87 0.13\nand a monoclinic WO phase (JCPDS#32-1393). This new system shares the same\n2\ndesign principle but has a different primary alloy composition (Co-W vs. Ni-W).\nRemarkably, this Co-W-based heterostructure exhibits the same key mechanistic\nsignatures as the Ni-W system. To investigate the hydrogen transfer behavior during\nthe reaction, comparative electrochemical measurements of the Co-W material were\nconducted in H SO /H O and D SO /D O electrolytes (Supplementary Fig. 33). A\n2 4 2 2 4 2\npronounced kinetic isotope effect (KIE≈2) confirms that the reaction is proton-transfer-\nlimited, indicative of a spillover process1.\nIn situ Raman spectroscopy (Supplementary Fig. 34) reveals hydrogen\naccumulation exclusively within the Co W phase (weakening of Co-O bond and\n0.87 0.13\nemergence of Co-H bond) with no proton insertion into WO (stable W-O bond)2-4,\n2\ndirectly visualizing the non-interfacial pathway.\nWhile the activity of the Co-W-based system differs, the fundamental mechanism\nis conserved. This extension demonstrates that our strategy is not material-specific.\nalloy-oxide heterointerface to generate a strain gradient that reconfigures intra-phase\nadsorption energetics and enables a non-interfacial spillover pathway. We have\ndiscussed this in the revised version as follows.\nRevision in the Results section and Supplementary Figures:\nTo evaluate the generality of the strain-gradient design principle beyond the Ni-\nW composition, we applied the same in situ phase reconstruction protocol to\nsynthesize a Co W -WO composite (Supplementary Fig. S32). Electrochemical\n0.87 0.13 2\ntesting confirmed its HER activity (Supplementary Fig. S33a-b). Critically, this distinct\nmaterial replicates the key mechanistic fingerprints of non-interfacial spillover\nestablished above. A pronounced KIE≈2 (Supplementary Fig. 33c-d) confirms a proton-\ntransfer-limited process. Operando Raman spectroscopy (Supplementary Fig. 34)\nreveals hydrogen accumulation exclusively within the Co-W alloy phase (weakening of\nCo-O and emergence of Co-H bond) with no proton insertion into WO (stable W-O\n2\nbond)2-4. The replication of these signature features in a chemically distinct system\nconfirms that the non-interfacial spillover mechanism is a direct consequence of the\nengineered heterointerface, rather than a peculiarity of the Ni-W composition.\nSupplementary Fig. 32. XRD patterns of (a) the CoWO precursor and (b) the\n0.87 0.13 2\nCo W -WO measured in 0.5 M H SO . (c) Polarization curves of Co W -WO\n0.87 0.13 2 2 4 0.87 0.13 2\nin 0.5 M H SO /H O solution and 0.5 M D SO /D O solution (without iR correction). (d)\n2 4 2 2 4 2\nThe relationship between KIE value and potential of Co W -WO .\n0.87 0.13 2\n0.87 0.13 2\nH SO from 0 to −600 mV vs. RHE.\n2 4\nComment 2: For comment #5, the authors have provided KIE experiments to validate\nthe proposed hydrogen spillover mechanism, but only on the NiW-WO2 catalyst. The\nKIE experiments should be performed on control samples and comparatively\nanalyzed to better clarify the structure-to-mechanism correlation.\nResponse 2: Thanks for your insightful suggestion. The comparative KIE analysis on the\ncontrol samples (Ni W and WO ) provides crucial, direct evidence for the structure-\n17 3 2\nmechanism correlation, strongly supporting our central claim that the non-interfacial\nH-spillover is a direct consequence of the engineered heterointerface.\nA pronounced KIE≈2 is observed exclusively for the Ni W -WO heterostructure\n17 3 2\n(Fig. 4d-e), unambiguously identifying proton transfer as the rate-determining step—\nthe kinetic signature of a H-spillover process. In stark contrast, both pure Ni W and\n17 3\nWO exhibit KIE values close to 1 (Supplementary Fig. 29), indicating the absence of a\n2\nproton-transfer-limited mechanism.\nThis dramatic distinction demonstrates that the H-spillover mechanism is not an\nintrinsic property of the individual Ni W or WO phases, but is uniquely enabled by\n17 3 2\ntheir specific integration into a heterostructure. The strain gradient and electron\ntransfer generated at this interface reconfigure the energy landscape to create the\nintra-phase proton migration pathway, which is now kinetically validated by this KIE\nexperiment. We have discussed this in the revised version as follows.\nRevision in the Results section and Supplementary Figures:\nIn stark contrast, both pure Ni W and WO exhibit KIE values close to 1\n17 3 2\n(Supplementary Fig. 29), indicating the absence of a proton-transfer-limited\nmechanism.\n17 3\n(b) WO in 0.5 M H SO /H O and 0.5 M D SO /D O solutions. The relationship between\n2 2 4 2 2 4 2\nKIE value and potential of (c) Ni W and (d) WO .\n17 3 2\nIn addition to these issues, this reviewer have the following issues for the authors to\nconsider.\nComment 3: As discussed in the introduction, the non-interfacial H spillover\nphenomenon has already been reported previously. The significance of this work lies\nin the observation of such phenomenon in a new non-noble metal-based material\nsystem, i.e. NiW-WO2. Accordingly, the title and abstract should be modified by\nincluding the material studied to clarify this. The current version is misleading as it\nimplies that this work is the first to discover the non-interfacial H spillover\nphenomenon.\nResponse 3: We sincerely appreciate your guidance to enhance the precision of the\n6\ntitle and abstract. We have revised both accordingly to explicitly highlight the Ni W -\n17 3\nWO material system and to more accurately frame our contribution.\nhas been reported, prior demonstrations have been confined to systems with limited\ngeneralizability, such as noble-metal-containing perovskites or highly complex high-\nentropy materials. The core advance of this work is not merely the observation of this\nphenomenon in a new Ni-W-based material, but the demonstration of a generalizable\ndesign principle to actively create it.\nSpecifically, we show that through deliberate heterointerface engineering of\ncombining gradient strain and directed electron transfer, we can strategically construct\nan efficient, non-interfacial spillover pathway within a conventional, earth-abundant\nheterostructure. This represents a shift from discovering the phenomenon in\nspecialized systems to programming it into scalable catalyst architectures via a clear\nphysical mechanism.\nThe Ni W -WO system serves as a proof-of-concept platform to establish this\n17 3 2\nprinciple. Our revised manuscript, particularly the title and abstract, now more clearly\ncommunicates that we report a non-interfacial H-spillover mechanism in a Ni W -\n17 3\nWO heterostructure enabled by synergistic strain and electron engineering,\n2\nemphasizing the enabling method over the specific material.\nThe updated title:\nEngineering non-interfacial hydrogen spillover in a Ni W -WO heterostructure\n17 3 2\nThe updated abstract:\nThe hydrogen spillover (H-spillover) effect offers a promising strategy to overcome the\nkinetic bottleneck of proton desorption in hydrogen evolution reaction (HER) catalysts.\nHowever, conventional H-spillover mechanisms rely on interfacial proton transfer\nbetween distinct phases and suffer from inherent energy barriers. Here, we\ndemonstrate a non-interfacial H-spillover mechanism in a Ni W -WO heterostructure,\n17 3 2\nengineered through the synergistic creation of a built-in strain gradient and directional\n7\nelectron transfer. This design spatially confines the complete hydrogen evolution\nprocess within the Ni W phase, thereby circumventing cross-phase migration and\n17 3\nreshaping the hydrogen adsorption energy landscape. Experimental and theoretical\nanalyses confirm the elimination of interfacial barriers and establishment of an\noptimized proton-migration route. The resulting catalyst achieves a low overpotential\nof 21 mV at 10 mA cm–2 in 0.5 M H SO , along with notable stability (>1500 hours at\n2 4\n500 mA cm–2) and a near-unity Faradaic efficiency (98.6%). This work demonstrates\nhow tailored heterostructures can effectively bypass interfacial bottlenecks, providing\na new paradigm for developing efficient non-precious H-spillover catalysts and\nadvancing sustainable hydrogen production.\nComment 4: The function and the necessity of the WO2 phase in the designed\ncomposite should be explained in the introduction part.\nResponse 4: Thank you for highlighting the need to clarify the role of WO . We fully\n2\nagree and have revised the Introduction to explicitly define its function.\nAs now articulated in the revised Introduction, the WO phase is not the primary\n2\nhydrogen-evolution site. Instead, its strategic necessity lies in its dual function as a\nstructural and electronic modulator:\n(1) The lattice mismatch at the Ni W /WO interface induces a built-in strain\n17 3 2\ngradient within the Ni W phase.\n17 3\n(2) WO acts as an electron acceptor, driving directional electron transfer from\n17 3\nThis synergistic “strain + electron” engineering collectively reconfigures the\nhydrogen adsorption energy landscape across the Ni W phase. It is this reconfigured\n17 3\nlandscape that creates the continuous, intra-phase proton migration pathway, which\nis the core of our non-interfacial spillover mechanism. Without WO , this tailored\n2\nenergy gradient and the resulting efficient pathway cannot be established. These\npoints have been clarified in the introduction as follows.\nThe rational design of efficient H-spillover pathways requires precise control over\nthe hydrogen adsorption energy landscape across multiple sites. Strain and electronic\nstructure engineering have emerged as powerful levers for such tuning5. Tungsten\noxides (particularly WO ) present unique advantages as heterostructure components\n2\ndue to their structural adaptability and strong propensity for electronic coupling. By\nconstructing an intimate heterointerface, WO can act as a strategic modulator to\n2\nimpose a controllable strain field and drive directional electron transfer to an adjacent\nactive phase6. This synergistic modulation is envisioned to create the internal driving\nforce necessary for steering proton migration along a non-interfacial, intra-phase\npathway.\nHerein, we report a design strategy to bypass interfacial bottlenecks in H-spillover\ncatalysis (Scheme 1). Through in situ phase reconstruction, we construct a Ni W -WO\n17 3 2\nheterostructure where the WO phase is deliberately employed not as a primary\n2\ncatalytic site, but as an essential engineering component. Its critical function is to\ninduce a built-in strain gradient and facilitate electron transfer within the contiguous\nNi W phase. This concerted action collectively reconfigures the local hydrogen\n17 3\nadsorption energetics, thereby constructing an efficient, non-interfacial proton\nspillover channel confined entirely within the Ni W phase.\n17 3\nComment 5: The activity of the NiW-WO2 catalyst should be dependent on the\ndensity of the grain boundaries. Can the ratio between the NiW and WO2\ncomponents be varied to optimize the performance?\nResponse 5: Thank you for raising this important point regarding the potential\ninfluence of grain boundaries and phase composition. We agree that these factors can\nbe crucial in many heterostructured systems. In our specific case, however, the design\nand synthesis pathway lead to a distinct mechanistic picture. The Ni W -WO\n17 3 2\nheterostructure is formed via a controlled in situ phase reconstruction from a\nstoichiometric NiWO precursor. This one-pot synthesis ensures atomic-scale intimacy\n4\n9\nat the interface and yields a thermodynamically stable phase composition, as dictated\nby elemental conservation and reduction kinetics. Consequently, the Ni W -WO ratio\n17 3 2\nis inherently fixed in our optimized system, and the resulting interface/grain boundary\ndensity is an intrinsic outcome of this synthesis, not an independently tunable\nvariable7-8.\nOur combined experimental and theoretical evidence indicates that the\nexceptional HER activity stems not from maximizing generic grain boundary density,\nbut from the specific electronic and strain effects generated at the well-defined\nNi W -WO heterointerface. As detailed throughout the manuscript, it is the\n17 3 2\nsynergistic gradient strain field and directional electron transfer arising from this\ncoherent interface that reconfigure the hydrogen adsorption landscape within the\nNi W phase, creating the efficient, non-interfacial spillover pathway.\n17 3\nTherefore, while your suggestion is highly relevant for composite catalysts made\nby physical mixing, our work demonstrates an alternative and powerful design\nprinciple: creating a single, optimized heterointerface with tailored electronic/strain\nproperties can be more effective for directing catalytic function than simply increasing\nthe number of grain boundaries. The fixed phase ratio in our system is not a limitation\nbut a feature of this targeted design, which yields a well-defined and highly active\ncatalytic system.\nWe appreciate your insightful comment, which has allowed us to better articulate\nthe unique design logic and mechanistic origin of activity in our catalyst system. We\nhave discussed this in the revised manuscript as follows.\nRevision in the Results section:\nThe well-defined Ni W -WO heterointerface-induced strain gradient is posited\n17 3 2\nto establish a continuous modulation of hydrogen adsorption energy across the Ni W\n17 3\nphase, thereby facilitating directional proton spillover and enhancing the hydrogen\nevolution kinetics\nComment 6: The activity and stability evaluation in the PEM configuration were\n10\nperformed at room temperature. Considering that the stability is highly relevant to\nthe testing conditions and that most studied report the PEM performance at 80\ndegree, the performance evaluation at higher temperature should be provided for\nbenchmarking purpose.\nResponse 6: Thank you for raising this critical point regarding benchmarking under\nrealistic conditions. In direct response, we have conducted new PEM electrolyzer\nevaluations at the industrially relevant temperature of 80 oC.\nAs shown in Supplementary Fig. 26, the Ni W -WO ‖IrO cell demonstrates\n17 3 2 2\noutstanding performance, requiring cell voltages of only 1.60 V and 1.65 V to achieve\n500 and 1000 mA cm-2, respectively. This performance is highly competitive with the\nPt/C‖IrO benchmark under identical conditions.\ntemperature. Over 1000 cycles, the cell voltages at both high current densities\nremained stable with minimal fluctuation (Supplementary Fig. 27). The nearly perfect\noverlap of the polarization curves before and after this rigorous testing confirms\nnegligible performance degradation.\nThese new results provide robust evidence that the Ni W -WO catalyst not only\n17 3 2\nmatches noble-metal benchmarks in activity but also retains its functional integrity\nunder harsh, practical operating conditions. This comprehensive high-temperature\nPEM evaluation solidly underscores the catalyst’s strong potential for industrial\napplication. We have discussed this in the revised manuscript as follows.\nRevision in the Results section and Supplementary Figures:\nIn a full-cell configuration with an IrO anode (Supplementary Fig. 25), the Ni W -\n2 17 3\nWO ||IrO electrolyzer delivered a low cell voltage of 1.945 V under room\n2 2\ntemperature and 1.650 V at 80 oC for 1000 mA cm−2 (Supplementary Fig. 26), a\nperformance competitive with the Pt/C||IrO benchmark (Fig. 3h). Critically, the\n2\ndevice demonstrated exceptional durability under both room temperature and 80 oC,\nsustaining its performance over 1000 cycles at high current densities (500 and 1000\n11\nmA cm−1) with negligible degradation in the polarization curve (Supplementary Fig. 27).\nSupplementary Fig. 26 | The LSV curves of PEM electrolyzers assembled with Ni W -\n17 3\nWO ||IrO and Pt/C||IrO measured at 80 oC.\n2 2 2\ndensities of 500 and 1000 mA cm−1 of Ni W -WO ||IrO assembled PEM electrolyzer\n17 3 2 2\nmeasured at (a) room temperature and (b) 80 oC. LSV curves of the Ni W -WO ||IrO\n17 3 2 2\nassembled PEM electrolyzer measured at (c) room temperature and (d) 80 oC before\nand after 1000 cycles.\n1. Wang, Q. et al. 2D metastable-phase hafnium oxide triggers hydrogen spillover for\nboosting hydrogen production. Adv. Mater. 37, 2415978 (2025).\n2. Ross-Medgaarden, E. I. & Wachs, I. E. Structural determination of bulk and surface\ntungsten oxides with UV-vis diffuse reflectance spectroscopy and Raman\nspectroscopy. J. Phys. Chem. C 111, 15089-15099 (2007).\n3. Guo, P. et al. Heterojunction-induced rapid transformation of Ni3+/Ni2+ sites which\nmediates urea oxidation for energy-efficient hydrogen production. Adv. Mater. 36,\n2311766 (2024).\n4. Wang, Y. et al. Defect-induced Zn-Co pair active site for high-efficiency\nelectrosynthesis of H O . Matter 8, 102479 (2025).\n2 2\n5. Li, J., Ma, Y., Ho, J. C. & Qu, Y. Hydrogen spillover phenomenon at the interface of\nmetal-supported electrocatalysts for hydrogen evolution. Acc. Chem. Res. 57, 895-\n904 (2024).\n6. Peng, Q. et al. Lattice strain induced Ni Se/WO heterostructures accelerate\n0.85 2.90\ndynamic reconstruction for efficient water oxidation. Appl. Catal. B-Environ. Energy\n384, 126158 (2026).\n7. Pei, Z. et al. Theory-guided design of high-entropy alloys with enhanced strength-\nductility synergy. Nat. Commun. 14, 2519 (2023).\n8. Hansen, P. L. et al. Atom-resolved imaging of dynamic shape changes in supported\ncopper nanocrystals. Science 295, 2053-2055 (2002).\n14"}, {"round": 2, "reviewer_comments": {"reviewer_1": "The manuscript has been adequately updated and is now suitable for publication in Nature Communications.", "reviewer_3": "The work is now suitable for publication.", "reviewer_4": "After carefully reading the manuscript and the authors’ response to reviewers’ comments, I believe the authors have in large part properly addressed reviewer #2’s concerns by performing additional experiments and calculations. However, the response to specific comment #2 and #5 needs more attention.\nFor comment #2, the authors claim that the current work is advantageous over previous works about non-interfacial hydrogen spillover for its generalizability. However, this is not obvious as the non-interfacial hydrogen spillover is only demonstrated on one single material of NiW-WO2. The generality of such design strategy in other types of materials remain questionable.\nFor comment #5, the authors have provided KIE experiments to validate the proposed hydrogen spillover mechanism, but only on the NiW-WO2 catalyst. The KIE experiments should be performed on control samples and comparatively analyzed to better clarify the structure-to-mechanism correlation.\nIn addition to these issues, this reviewer have the following issues for the authors to consider.\n1. As discussed in the introduction, the non-interfacial H spillover phenomenon has already been reported previously. The significance of this work lies in the observation of such phenomenon in a new non-noble metal-based material system, i.e. NiW-WO2. Accordingly, the title and abstract should be modified by including the material studied to clarify this. The current version is misleading as it implies that this work is the first to discover the non-interfacial H spillover phenomenon.\n2. The function and the necessity of the WO2 phase in the designed composite should be explained in the introduction part.\n3. The activity of the NiW-WO2 catalyst should be dependent on the density of the grain boundaries. Can the ratio between the NiW and WO2 components be varied to optimize the performance?\n4. The activity and stability evaluation in the PEM configuration were performed at room temperature. Considering that the stability is highly relevant to the testing conditions and that most studied report the PEM performance at 80 degree, the performance evaluation at higher temperature should be provided for benchmarking purpose."}, "author_response": ""}, {"round": 3, "reviewer_comments": {"reviewer_4": "The authors have adequently addressed the reviewer's comment and the revised manuscript is recommended for publication."}, "author_response": ""}]

Decision letter