Redefining catalyst reconstruction and Cl^--repulsion correlation to delineate a dynamic protective skeleton for seawater splitting

Domain: chemistry
Published: 2026-02-21
Editorial outcome: major_revision

Abstract
From the perspective of sustainable energy and green hydrogen, the critical challenge in seawater electrolysis lies in suppressing competitive chloride oxidation and corrosion at the anode. Despite the proposed protection mechanism based on surface reconstruction-derived anion repellence, a fundamental question of selectivity remains: why does the protective layer exclude Cl- but not the formally equivalent OH-? We now redefine the connection between reconstruction and interfacial protection in NiFeS anodes to elucidate the origin of underlying selectivity. In-situ Raman, in-situ X-ray absorption spectroscopy, and molecular dynamic calculations further confirm that the reconstructed SO 2- establishes an enhanced hydrogen-bond network. This network significantly weakens the hydrogen-bonding strength of Cl- relative to H O, thereby enhancing the discrimination between OH- and Cl-. In practical application, NiFeS demonstrates energy-saving seawater splitting performance (261.8 mV @ 100 mA cm-2), long-term durability (2000 h @ 1.0 A cm-2), and enables stable intermittent electrolysis (1500 start-stop cycles @ 500 mA cm-2). When integrated into an anion exchange membrane cell, it shows promising electrolysis performance for industrial applications (1.970 V @ 1.0 A cm-2).

Introduction
The growing global energy shortage, driven by excessive consumption of traditional fuels like coal, oil, and natural gas, has spurred increasing interest in clean energy alternatives1,2. Water electrolysis for green hydrogen (H ) production has emerged as a promising renewable energy solution to address global energy transition challenges and the climate crisis3,4. However, the conventional alkaline and proton exchange membrane electrolysis systems rely on ultrapure water feedstock, posing limitations due to fresh water scarcity and energy-intensive purification processes. Seawater, constituting 96.5% of global water reserves, presents an environmentally sustainable and cost-effective alternative electrolyte for water electrolysis5,6.

Compared to pure-water electrolysis systems, chloride ions (Cl−) in seawater create critical challenges for the anode. The competing chlorine evolution reaction (CER) not only increases energy consumption but also causes irreversible electrode degradation1,7. Although NiFe-layered double hydroxide (LDH), a benchmark catalyst in alkaline water electrolysis, is considered an ideal candidate, the harsh oxidative conditions at the anode induce corrosion and dissolution of metal active sites, leading to instability8,9. Notably, recent literatures have shown that during the anodic reconstruction of catalysts, non-metallic elements (such as sulfur and phosphorus) undergo oxidation and leaching, which lead to the formation of a protective layer10–15. This layer enhances the corrosion resistance of NiFe LDH by establishing electrostatic repulsion against Cl−, thereby extending the catalyst’s lifespan16. Following this reconstruction process, the OER performance is significantly improved17,18. However, a fundamental question warrants deeper investigation: why, despite sharing the equivalent charge, is only Cl− selectively repelled? This phenomenon cannot be adequately explained by the theory of electrostatic repulsion alone.

This work seeks to correlate catalyst reconstruction with chloride exclusion and unravel the mechanistic origin of the observed corrosion resistance. We synthesized a high-performance NiFeS catalyst through Fe3+ etching of a nickel substrate followed by sulfurization. The resulting catalyst demonstrates notable OER activity and selectivity in natural seawater. During electrolysis, NiFeS preferentially reconstructs, forming a SO 2−-rich interface and highly active NiFe LDH. SO 2− enhances the local hydrogen bond network, creating a region with strengthened hydrogen bonds. Compared to pure water hydrogen bonds, the enhanced network exhibits hydrogen bond strengths in the order: OH− > H O > Cl−. This confirms that SO 2− promotes the mass transport of H O over that of Cl−, while prioritizing OH− access. Consequently, a Cl−-free microenvironment is established at the anode interface. The optimized NiFeS achieves a low cell voltage of 1.885 V @ 500 mA cm−2 for overall seawater splitting, surpassing commercial Pt/C||RuO (2.213 V), while possessing high stability (2000 h @ 1.0 A cm−2 with almost 100% Faradaic efficiency). After 1500 start-shutdown cycles at 500 mA cm−2, the performance of NiFeS is maintained in alkaline seawater. The satisfactory AEM electrolyzer testing (1.970 V @ 1.0 A cm−2 and 200 h @ 1.0 A cm−2) underscore the practical viability of this strategy for seawater electrolysis-driven hydrogen production using intermittent renewable energy sources.

Results
NiFeOOH was first synthesized by etching NF with Fe3+, followed by sulfurization with Na S to produce NiFeS, as shown in Fig. 1a19,20. Optical images (Fig. S1) show that NF is silver-white, NiFeOOH exhibits a brownish-yellow color, and NiFeS appears black. SEM images of NiFeOOH (Figs. S2–S5) reveal that Fe3+ etching induces the growth of a nanosheet array on the initially smooth NF surface. As etching proceeds, the nanosheets become more porous, and their dimensions gradually decrease. NiFeS-5 largely retains the nanosheet morphology, with a nano-lotus structure (Fig. S6). As sulfurization time increases to 10 min (NiFeS-10), the nanosheets persist but begin to adhere and thicken noticeably (Figs. 1b–d and S7). Further extending sulfurization to 15 min (NiFeS-15) causes over-sulfurization, resulting in severe agglomeration into a bulk-like structure (Fig. S8). This aggregated morphology reduces the number of exposed active sites and impedes the mass transport of OH− and O at the electrode interface, ultimately compromising the catalytic kinetics and stability21. TEM images (Fig. S9) further confirmed the nanosheet structure of NiFeS-5. HRTEM and SAED analyses revealed the co-existence of distinct crystalline phases, identified as NiFe LDH and NiS. The measured lattice fringes of 0.212 and 0.265 nm correspond to NiFe LDH and NiS, respectively, matching previously reported values22,23. A slight nanosheet structure is still retained on the NiFeS-10 surface, NiFe LDH, and NiS are still found, which show abundant crystalline-amorphous fringes (Figs. 1e, f and S10). Furthermore, the EDS and color mapping images (Fig. 1g) confirmed the uniform distribution of elements in the product, and demonstrated that the evolution of surface structure was caused by sulfurization. When the sulfurization time reaches 15 min, the surface structure transforms into a block-like aggregated structure, as confirmed by the SEM results, shown in Fig. S11.

XPS analysis was performed to investigate the surface chemical states and the effect of sulfurization (Figs. 1h, i and S12). The survey spectra and elemental composition determined by XPS were consistent with the EDS mapping results. Specifically, in the Fe 2p spectrum, 710.6 and 725.0 eV correspond to Fe2+, while 713.2 and 726.8 eV are assigned to Fe3+24,25. Sulfurization significantly reduced the Fe3+ content from 56.89% to 31.14%. In the Ni 2p spectrum, peaks located around 856.3 and 874.0 eV are assigned to Ni2+4,26. For NiFeS, the Ni 2p and Fe 2p peaks shift to lower binding energies compared to NiFeOOH, indicating an increased electron density around these metal centers due to the lower electronegativity of sulfur. This electron-rich state enhances the anodic reactivity of the precursor, promoting rapid and thorough electrochemical reconstruction. Moreover, the initial presence of low-valence Fe helps suppress dissolution during OER, thereby improving both activity and stability27. The peaks of the S 2p spectrum at 162.2, 163.4, and 168.1 eV were convolved into S2− 2p3/2, S2− 2p1/2, and SO 4 2−, respectively21. The discovery of oxidized S indicates that Fe3+ has oxidized sulfur and converted to Fe2+, which corresponds to the spectral results of Fe 2p. Raman spectroscopy (Fig. S12c) further confirmed the above results, with characteristic bands between 300–800 cm−1 assigned to nickel and iron hydroxides28,29. On the NiFeS surface, except for Fe(OH) located at 731.4 cm−1, the peaks of all other hydroxides are covered by sulfides, and oxidized SO 2− can also be detected21. The FTIR results provide strong evidence for the formation of low-valence Fe and complex sulfide species on the NiFeS surface (as Fig. S12d). Electrode wettability critically influences mass transport. NiFeS exhibits enhanced hydrophilicity, with a contact angle of 74.4° (Fig. S13), which facilitates electrolyte infiltration and promotes efficient bubble detachment from active sites. In contrast, NF and NiFeOOH show poor wettability (contact angles of 122.6° and 121.9°, respectively), resulting in sluggish mass transport. This stark contrast highlights how sulfur-induced hydrophilicity contributes to comparable catalytic performance by optimizing the electrode-electrolyte interface.

Electrocatalytic OER performance was evaluated in alkaline seawater (Seawater + 1.0 M KOH) with the typical three-electrode configuration, and the natural seawater was taken from the Weihai sea area in Shandong Province, China, without any pretreatment (Fig. S14). To determine the optimal catalytic electrode, we first checked the catalysis performance of NiFeOOH in alkaline seawater, as shown in Fig. S15a, b. NiFeOOH-2 exhibited the optimal OER performance, featuring a low Tafel slope of 78.83 mV dec−1, which suggests that an optimal Ni/Fe atomic ratio is crucial for high activity. Therefore, in subsequent work, sulfurization was performed based on NiFeOOH-2 to obtain NiFeS. NiFeS-10 was identified as the optimal catalyst, as shown in Figs. 2a, b and S15c, and it requires an overpotential of only 261.8 mV to achieve 100 mA cm−2, outperforming RuO (290.3 mV), NiFeS-5 (289.9 mV), NiFeS-15 (311.8 mV), and NF (461.1 mV). Furthermore, the catalytic activity of NiFeS-10 is much higher than that of commercial precious metal RuO 2, with a Tafel slope as low as 46.90 mV dec−1 (better than 136.72 mV dec−1). Moreover, NiFeS-10 also provides lower Rct values (0.971 Ω) to achieve lower charge transfer resistance. The electrochemically active surface area is reflected by the double-layer capacitance (C dl) via cyclic voltammetry (CV) curves at different scanning rates. NiFeS-10 has the largest catalytic active area (213.3 cm2), far surpassing NF (161.5 cm2), NiFeS-5 (192.5 cm2), and NiFeS-15 (164.5 cm2), as shown in Fig. S16. After normalizing the active area, NiFeS-10 still exhibits outstanding catalytic activity (Fig. S17). Figure 2e shows the rapid response of NiFeS to fluctuating current in multicurrent step testing from 50 to 400 mA cm−2.

To elucidate the underlying mechanism, the experiments and DFT were carried out to identify the actual OER mechanism. Tetramethylammonium cation (TMA+) was employed to detect O+ 2− generated by LOM based on strong electrostatic interactions, as shown in Figs. 2f and S18. For NiFeS, the performance degradation caused by TMA+ presence far exceeds that of NiFeOOH (42.96 > 31.17 mV dec−1), and the intensified characteristic peaks belonging to TMA+ at 750 and 947 cm−1 on the NiFeS surface are significantly higher than NiFeOOH, supporting the enhancement of LOM (Lattice Oxygen Mechanism)10,30. Figure S19 further shows that NiFeS has a higher proton reaction order than NiFeOOH (ρRHE=∂log(j)/∂pH, 0.665 > 0.445) through performance tests conducted at pH 12.5, 13.0, and 13.5, indicating that it is more pH-dependent and that LOM has a higher priority for NiFeS30,31. Complementary DFT calculations were conducted to further support these findings. To approximate the electrochemical environment, we modeled both a clean NiFe LDH surface and an SO 2−/NiFe LDH surface to evaluate the potential OER mechanisms, as Fig. S20. The reaction step O* + OH− - e− → OOH* emerged as the rate-determining step (RDS). For the LOM pathway, the RDS energy barrier was calculated to be 1.45 eV on the clean surface, while the presence of SO 2− substantially reduced this barrier to 0.67 eV. In contrast, for the AEM (Adsorption Evolution Mechanism), the introduction of SO 2− increased the OER energy barrier from 0.85 to 0.97 eV. These computational results are consistent with the experimental data and indicate that the interfacial SO 2− may function as a modulator that favors the LOM pathway over the AEM, suggesting it could contribute to enhanced OER kinetics within the framework of our model.

Furthermore, we performed in-situ Raman spectroscopy to monitor the evolution and structural transformation of active species, as shown in Fig. 3a. For the NiFeS catalyst, the characteristic sulfide peak disappeared when the potential reached 1.15 V. At 1.20 V, reconstructed-phase peaks emerged at 454 cm−1 (bending) and 530 cm−1 (stretching), assigned to NiFe LDH, with concurrent observation of distinct SO 2− signatures32,33. Notably, sulfur leaching results in NiFe LDH formation with no detectable chloride species. However, in-situ Raman spectroscopy was unable to effectively track the evolution of trace Fe species. We performed the in-situ X-ray absorption spectroscopy (XAS), with results shown in Figs. 3b and S21. Due to the reducing effect of sulfur, the Fe in NiFeS exhibits negative shift than Fe3+ from the XANES results, in line with the XPS results. Meanwhile, we note that the Fe valence state remains essentially unchanged during both reconstruction and OER processes, demonstrating notable stability of Fe species. EXAFS results reveal a decrease in Fe-S/O bond length during reconstruction, which involves sulfur leaching and subsequent OH− incorporation34. The absence of any detected Fe-Cl coordination structures further confirms that the reconstruction of NiFeS following sulfur leaching proceeds preferentially through complexation with OH− to form LDH35. DFT calculations confirm the thermodynamic advantage of OH* in NiFeS vacancies, directing seawater reconstruction toward NiFe LDH formation via OH− oxidation instead of Cl− corrosion (Fig. S22). Only characteristic peaks corresponding to NiFe LDH were identified in the in-situ Raman spectrum of NiFeOOH (Fig. S23). Moreover, we investigated the behavior of interfacial water molecules at different current densities, as shown in Figs. 3c, d and S24. The O-H stretching vibration band in the in-situ Raman spectroscopy was deconvoluted into three peaks: strongly hydrogen-bonded H O (4HB-H O), weakly hydrogen-bonded H O (2HB-H O), and non-hydrogen-bonded H O (Free-H O), enabling calculation of the area percentage for each component36. As the current density increased, the 4HB/Free ratio progressively rose, with NiFeS exhibiting significantly higher ratios than NiFeOOH. This demonstrates enhanced modulation of the interfacial hydrogen-bonding network. Finally, the in-situ Raman spectroscopy of deuterium oxide (D O) on NiFeS surface was conducted (Fig. S25). With increasing current density, the Raman shift of the strong hydrogen-bonding peak shifted toward lower wavenumbers, providing direct evidence for the reinforcing effect of SO 2− on the interfacial hydrogen-bonding network37.

Ab initio molecular dynamics (AIMD) simulations of SO 2−-containing/-deficient electrolytes (Fig. S26 and Supplementary Data 1) elucidated interfacial microenvironment evolution, applying a methodology that has been previously demonstrated as effective for interface analysis38. Trajectory analysis revealed that the SO 2−-containing system exhibited significantly OH− enrichment and effective Cl− exclusion at the anode interface, as shown in Fig. 3e. On the NiFeOOH surface, the closest approach distance of Cl− to the surface is only 2.2 Å. However, in the presence of SO 2−, this distance is significantly extended to 11.6 Å. Meanwhile, the OH− concentration signal near the surface (approximately 1.6–3 Å) shows a notable increase. This interfacial modulation contributes to the strong pH-dependent OER activity of NiFeS (Fig. S19), indicating that SO 2− directly enhances catalytic activity and selectivity. Crucially, SO 2- induces a substantially intensified hydrogen-bond (H-bond) network near the interface, consistent with experimental observations. Cluster model calculation results (Fig. 3f) quantified H-bond strengths: at SO 2−-free interfaces, H O bonds (18.63 kJ/mol) were weaker than OH− (-100.74 kJ/mol) and Cl− (-76.75 kJ/mol), causing interfacial water to preferentially capture both OH− and detrimental Cl−. Within the SO 2−-H O clusters, electrostatic repulsion significantly weakened both H-bonding strengths of OH− and Cl−. Conversely, the H-bond strength of H O substantially increased, exceeding that of Cl−-H O interactions. Meanwhile, the hydrogen-bond network maintains its preferential binding affinity for OH−. Non-covalent interaction (NCI) analysis characterized these weak interactions, as illustrated in Fig. S27. This compelling result demonstrates that SO 2−-hydrate complexes preferentially bind H O even in Cl−-containing environments with limited OH− availability, a preference that is further amplified by the 110-fold excess concentration of H O over Cl−39. To quantitatively reveal how SO 2− affects the distribution of Cl− at the electrode interface from a thermodynamic perspective, we employed the slow-growth thermodynamic integration method to calculate the free energy change of Cl− migrating from the bulk solution to the electrode surface40. As shown in Fig. S28, the free energy profile of Cl− increases monotonically as it approaches the electrode surface, but the energy is significantly higher in the presence of SO 2− (5.314 > 1.530 eV). This result provides a fundamental quantitative thermodynamic explanation for the phenomenon of SO 2− induced interfacial exclusion of Cl−. Essentially, SO 2− enables the interfacial hydrogen-bond network to selectively discriminate between Cl− and OH− while actively depleting interfacial Cl− concentrations. We identify this hydration-shell restructuring as the fundamental mechanism for selective Cl− exclusion despite identical anion charge states.

The ions of the NiFeS surface were detected using time-of-flight secondary ion mass spectrometry (TOF-SIMS) before OER and after 50 h of OER testing at 1.0 A cm−2, as shown in Figs. 3g, S29 and S30, and Table S2. In the TOF-SIMS mapping of NiFeS before OER, strong S2− (S−/HS−) signals were observed. However, only weak SO − was observed in the oxidation state signal, but the intensity was much lower than S2−. The signals of Ni and Fe belonging to the high valence state are also very difficult to determine. But after OER, the peak intensity of S2− significantly decreased, and the distribution density of oxidized S was significantly enhanced, especially SO − and NaSO −, which means that the sulfur is oxidized to form SO 2− adsorbed on the surface. Furthermore, we found that the ratio of Cl−/OH− on the NiFeS surface is about 0.37, much lower than that in seawater, further indicating that NiFeS shows a repulsive Cl− effect41. These results provide compelling evidence for the in-situ formation of SO 2− species and their intrinsic capability for Cl− exclusion at the electrode interface.

To further verify the effectiveness of the hydrogen-bond network selection, we performed electrolysis in seawater using RuIrTi, which favors the CER, at a current density of 250 mA cm−2. The concentration of ClO− in the electrolyte was measured after electrolysis. We propose that this method is highly sensitive since Cl− at the interface can be electrochemically oxidized to ClO− and quantitatively detected. As shown in Fig. S31, the electrolyte containing added SO 2− exhibited only a light pink coloration with a significantly weaker signal intensity. This result clearly demonstrates that the introduction of 0.20 M SO 2− effectively reduces the local Cl concentration at the RuIrTi electrode interface, thereby achieving highly selective OER. We also evaluated the OER performance of NiFeS in chloride-containing and chloride-free KOH solutions to investigate the influence of Cl−, and the results are presented in Fig. S32. In 1.0 M KOH, NiFeS exhibits a 24.63 mV lower overpotential than NiFeOOH at 100 mA cm−2. This enhancement is attributed to the presence of SO 2−, which aligns with our previous analysis in Fig. 2f. The reduced Tafel slope, showing an improvement of 24.14 mV dec−1, provides further evidence for this conclusion. After introducing Cl−, both NiFeS and NiFeOOH experienced OER performance degradation. Notably, the overpotential of NiFeS at 200 mA cm−2 increased by only 10.34 mV compared to its performance in pure KOH, while NiFeOOH showed a much larger increase of 28.49 mV. Furthermore, the smaller increase in Tafel slope observed for NiFeS (6.58 vs. 12.81 mV dec−1 for NiFeOOH) demonstrates that the enhanced hydrogen-bond network at the interface promotes selective hydroxide ion transport while reducing the local chloride concentration.

The potential issue of elemental leaching requires deep investigation due to the inevitable dissolution of elements during the reconstruction process42. We directly employed NiFeS as the anode for stability testing at a current density of 1.0 A cm−2, while monitoring the Fe and S content alongside performance degradation during the OER, as shown in Fig. S33. Throughout the oxygen evolution process, dissolution of both Fe and S elements was observed, with the S leaching being more pronounced than that of Fe. After 1 h of operation, the retention rates of Fe and S decreased to 95.11 and 94.03%, respectively. However, the overpotential at 100 mA cm−2 increased by only 1.95 mV. Following 10 h of OER operation, the Fe and S retention rates remained at 87.98 and 81.50%, respectively, with only a minimal activity loss of 3.78 mV. These results collectively demonstrate that the performance degradation caused by Fe and S leaching during the OER reconstruction process is nearly negligible.

To evaluate the durability of NiFeS in natural seawater, the corrosion test was carried out in seawater + 1.0 M KOH, as shown in Fig. 4a. The corrosion potential increased gradually with the sulfurization time, which reached a comparable 1.273 V after 15 min of sulfurization. And the corrosion potential of NiFeS-10 was also maintained at a satisfactory 1.165 V, with the surface catalytic sites more arduous to leach under Cl− attack. Furthermore, we investigated the Faraday efficiency at different current densities, as shown in Figs. 4b and S34. The almost linear relationship, with values approaching 100%, implies that NiFeS exhibits notable OER specificity. Furthermore, the NiFeS catalyst also demonstrates remarkable hydrogen evolution reaction (HER) performance, shown in Fig. S35. The optimized NiFeS-5 exhibits favorable reaction kinetics (48.22 mV dec−1) and rapid charge transfer capability (Rct = 5.19 Ω), enabling its performance at high current densities to rival that of commercial precious-metal Pt/C. Therefore, we set NiFeS-5 as the cathode and NiFeS-10 as the anode to build a non-membrane electrolysis system (NiFeS-5||NiFeS-10) for overall seawater splitting, and the polarization curves are shown in Fig. 4c. NiFeS-5||NiFeS-10 required only the cell voltage of 1.885 and 2.030 V to achieve 500 and 1000 mA cm−2, respectively. In contrast, the commercial noble metal catalysts (Pt/C||RuO ) require 2.213 and 2.464 V to drive the equivalent current density. Encouraged by the notable corrosion resistance and OER performance, the long-time chronoamperometry test was conducted at 1.0 A cm−2 to determine the durability of the seawater splitting, as Fig. 4d, e. And the catalytic performance of the electrolysis system was checked after the stability test to evaluate its degradation during seawater operation (Fig. S36). The electrode system was operated continuously for 2000 h, and the electrolytic performance degradation was only 25 μV/h. The untreated raw NF and RuO were directly corroded, leading to the collapse of the electrolysis system in less than 1 h, and the cell voltage was also higher than NiFeS-5||NiFeS-10. Moreover, the electrolytic current density and stability of NiFeS were notable to those of the recently reported advanced seawater electrolysis catalysts for anode (Table S1). The active chlorine test result showed that the almost unchanged color demonstrated that hardly any ClO− was generated during electrolysis, supporting energy-efficient electrolysis (Fig. S37)43.

Subsequently, material characterization was employed to further elucidate the transformation of the catalyst, involving SEM, TEM, XPS, and ICP. From the SEM images, it is clearly confirmed that reconstruction of the electrode structure after the stability test (Figs. S38–S43). The nanosheets on the NiFeS-5 surface underwent degradation, leading to a further increase in surface roughness and the formation of a concave-convex morphology. Owing to intense electrochemical oxidation, the surface of NiFeS-10 was densely coated with nanosheets, accompanied by the adhesion of crystalline particles. TEM results further reveal that NiS can still be found on the NiFeS-5 surface after the stability test, and the signal in SAED patterns attributed to NiFe LDH is attenuated to the point of indistinguishability. For the anode, the intense oxidation led to complete leaching of S2−, and SAED patterns indicated the absence of detectable crystalline sulfide phases. After 2000 h of OER, NiFe LDH remained as the sole phase with a contracted lattice spacing. The surface elemental mapping revealed intensified Mg2+ signals on the cathode surface induced by alkaline deposits, while sulfur retained a homogeneous distribution across the catalyst surface to provide effective protection. Notably, this non-catalytic alkaline scale is identified as the notable contributor to the activity degradation, and the XPS results provide compelling evidence supporting this conclusion (Fig. S41). Furthermore, S2− and SO 2− serve as the primary sulfur species in NiFeS after the HER and OER, respectively. We conducted an ICP analysis to quantify the elemental retention of Fe and S in the catalyst, as depicted in Fig. S41. The Fe and S retention rates in NiFeS-5 were measured at notable values of 75.46 and 64.22%, respectively, and 71.96 and 54.46% for the anode catalyst, further corroborating its notable corrosion resistance.

We also explored the performance of NiFeS in low-pH seawater environments by checking the corrosion potential at pH levels of 12.0, 12.9, 13.4, and 13.7. The results demonstrate a clear positive correlation between pH elevation and improvement in corrosion potential, as shown in Fig. S44. Notably, even at pH 12.0, the corrosion potential reached 1.069 V, merely 0.049 V lower than that at pH 13.7 (1.118 V). Encouraged by this corrosion resistance, we conducted a stability test in 0.5 M KOH solution at a current density of 300 mA cm−2. The NiFeS electrode maintained stable operation for 320 h. Post-test SEM characterization revealed that while the Ni foam skeleton remained intact, the nanosheet morphology became less distinct compared to NiFeS tested in 1.0 M KOH. Although partial surface coverage by seawater-derived salts was observed, element mapping confirmed the uniform distribution of Ni, Fe, and S elements, demonstrating the notable durability of NiFeS under low-alkalinity conditions.

Moreover, considering the intermittency of renewable energy sources, we further investigated the evolutionary behavior of this protective mechanism under dynamic power input. The in-situ Raman spectroscopy was performed under periodic fluctuations, as shown in Fig. S45. During the 10-min working/resting cycles, the signals attributed to NiFe LDH of NiFeS showed no significant degradation, and SO 2− species remained detectable at the interface. For the NiFeOOH catalyst, the characteristic peaks of NiFe LDH were observable after 10 min of operation. However, following a 10-min resting period, the peak intensity dramatically decreased, with peaks corresponding to NiO octahedra and FeOOH detected44,45. This indicates that Niδ+ species undergo chloride-induced dissolution followed by complexation with OH−, ultimately triggering phase segregation. Upon resuming operation, the material suffered severe corrosion and dissolution. These findings collectively demonstrate that the SO 2−-enhanced hydrogen bonding network at the resting potential contributes to the resilience of NiFe LDH against chloride-induced corrosion during current fluctuations. Therefore, we investigated the intermittent electrolysis by applying pulsed currents (500 mA cm−2) with varying frequencies: 1/1 (1s work and 1s rest, high-frequency), 10/10 (high-frequency), 60/60 (medium-frequency), and 600/600 (low-frequency), subjecting the system to 1500 continuous start-stop cycles, as illustrated in Figs. 4f and S4646. After the intermittent electrolysis, the performance of the anode and cathode was checked (Fig. S47). Reassuringly, the catalysis activity remained preserved, consistent with prior in-situ Raman spectroscopy results. SEM images in Figs. S48–S55 confirm the structural reconstruction of NiFeS during pulsing, and markedly changed the surface morphology of NiFeS, which appears distinctly different from that observed after steady-state electrolysis. The nano-forest structure vanished after continuous 1/1 work, yet reappeared when the power supply interval was extended to 10 s. As the interval was further prolonged to 60 and 600 s, the resulting nanosheets became progressively shorter and narrower. ICP results confirm that intermittent current operation does not cause significant chloride-induced electrode corrosion, thus maintaining OER activity (Fig. S43). Under intermittent power supply with 600 s intervals, the dissolution of Fe and S was most pronounced, yet remained at acceptable levels of 61.02 and 67.82% retention, respectively. The XPS was employed to evaluate the surface elemental composition, and the inactive components originating from seawater were detected, as shown in Fig. S56 and Table S3. Under low-frequency operation, the sulfur retention on the surface remained relatively high. However, high-frequency operation resulted in considerably higher elemental leaching levels from both the anode and cathode, though within acceptable limits, a trend consistent with the ICP measurements. Furthermore, due to interference from inactive ions such as Mg2+, Ca2+, and Na+ from seawater, the active components of NiFeS became covered, resulting in performance degradation.

Finally, we further integrated an alkaline seawater AEM system (NiFeS-5||NiFeS-10) as shown in Fig. 5a. The device requires only a cell voltage of 1.970 V to achieve a current density of 1.0 A cm−2 in alkaline seawater, as shown in Fig. 5b. The Pt/C||RuO in alkaline AEM achieves a current of only 240 mA cm−2 at an equivalent voltage. Figure 5c shows the AEM electrolysis efficiency and energy consumption for hydrogen production, which still maintains an operating efficiency of 63.69% at 1.0 A cm−2 and consumes only 4.71 kW Nm−3 H . Figure 5d also shows that the price of hydrogen per GGEH ($1.052 at 500 mA cm−2) is well below the U.S. Department of Energy’s target of $2.0 by 202647. Finally, the potential application of AEMWE in industrial applications is further validated. The AEMWE is powered by a solar module, and we monitored cell voltage and current through an electrochemical workstation. The sunlight intensity fluctuates with time (from morning to sunset) and weather (from sunny to cloudy), causing the output voltage and current to fluctuate accordingly (Figs. 5e and S57). Even so, the AEMWE showed a fast response and operated stably for 1 day under outdoor conditions in Harbin, China.

Afterward, the AEMWE was operated stably for another 200 h at 1.0 A cm−2 without any performance degradation (Fig. S58). During the stability assessment, we also investigated the membrane evolution behavior and ion transport in seawater + 1.0 M KOH. Notably, the membrane maintained its geometric structure from the optical images in Fig. S59. The cross-sectional SEM images of the anion exchange membrane were obtained before and after the stability test, as shown in Figs. S60 and S61. Before the stability test, the catalytic elements Pt and Ru were firmly attached to the membrane surface. All elemental signals were concentrated at the membrane surface, with minimal detection within the membrane bulk. After operation, no catalyst detachment was observed, and the distribution of Pt and Ru signals remained consistent with that of the raw membrane. The cathode catalyst layer retained its original thickness, with no signs of agglomeration and full preservation of the micro-nano channel structure. Furthermore, chloride ion signals were scarcely detected within the membrane layer. We also collected electrolytes from the anode and cathode outlets for ion chromatography analysis (IC), with results presented in Fig. S62. The Cl− concentration at the anode outlet was nearly identical to that at the cathode outlet, with this ratio remaining close to 100%. These findings confirm that OH− transport through the anion exchange membrane was not compromised by Cl− crossover. Regarding SO 2−, its concentration decreased in the effluent from the anode side, implying that trace sulfate in seawater was captured at the electrode interface. This observation suggests that performance degradation due to sulfur loss is unlikely. After stability testing, the nanosheet morphology on the NiFeS surface transformed into a porous structure. Elemental mapping further confirmed that Fe and S remained detectable and uniformly distributed post-test (Fig. S63). These results collectively demonstrate the promising application potential of NiFeS in AEM seawater electrolyzer. Moreover, for industrial applications, the large-area electrode preparation technique improves the synthesis efficiency of the catalyst, and we attempted to synthesize the catalyst on a large scale over an area of 1.0 m2 to adapt it to industrial applications, as shown in Figs. 5f and S64. We checked the surface topography and corrosion profiles of the catalysts at different positions, as shown in Figs. 5g and S65, respectively. The surface nanosheets were retained, showing a structure closely resembling that in Fig. 1c, d. The corrosion potential shifts were in the range of 13 mV, proving the uniformity of the expanded synthesis area.

In summary, this work deciphers the dynamic reconstruction of NiFeS catalysts and its critical role in enhancing chloride corrosion resistance in seawater splitting. During operation, the preferential oxidation of S2− and its subsequent interfacial accumulation, coupled with the transformation of NiFe into highly active NiFe LDH, leads to the in-situ formation of a sulfate-rich active interface. In-situ Raman, in-situ XAS, and DFT calculations confirm that SO 2− formed by interface reconstruction establishes a reinforced hydrogen-bonding network. This network enhances the bonding strength of H O beyond that of Cl− while preserving OH− binding priority, thereby creating an anode microenvironment enriched in OH− and depleted of Cl−. This hydrogen-bond network mechanism caused by catalyst reconstruction remains intact at resting potentials, enabling NiFe LDH to withstand 1500 frequent start-stop cycles in natural seawater. NiFeS exhibits notable OER activity (261.8 mV @ 100 mA cm−2) and outstanding stability (2000 h @ 1.0 A cm−2). The seawater AEM electrolyzer assembled PtRu||NiFeS achieves economic efficiency (1.970 V @ 1.0 A cm−2 and $1.052 per GGEH ) and shows quick response performance driven by solar power. Furthermore, large-scale synthesis and tunable interfacial ionic characters highlight the potential application of NiFe LDH in industrial seawater splitting systems. We are confident that our redefined understanding of the correlation between catalyst reconstruction and corrosion resistance establishes a direction for the targeted reconstruction design of catalysts in seawater splitting.

Methods
Reagents and seawater
All reagents involved in this work, including FeCl (99%), Na S·6H O (99%), KOH (95%), NaCl (99%), acetone (99.5%), anhydrous ethanol (99.8%), Tetramethylammonium hydroxide (AR), D O (99%), and N,N-Diethyl-1,4-phenylenediamine (98%) were purchased from Aladdin Company. Hydrochloric acid obtained from Kermel Chemical Reagent Co., Ltd (Tianjin). The half-membrane (PtRu/blank) is obtained from the SCI Materials Hub, where the anion exchange membrane used was Fumasep FAA-3-PK-130. The deionized water (DI water) is based on Hetai purification system (>18.25 MΩ cm), and the natural seawater is taken from Weihai sea area in Shandong Province, China. Seawater was treated with KOH to yield alkaline seawater (1M KOH + seawater, pH=13.72 ± 0.02), which was then sealed to prevent carbonation. The 1 mm thickness nickel foam (NF) comes from Wuhu Eryi New Materials Co., Ltd.

Synthesis method of NiFeS
First, NF was cleaned by acetone, hydrochloric acid, anhydrous ethanol, and DI water, respectively, to remove the oxide layer of the NF surface. FeCl powder was dissolved in 100 mL water to get a 0.1 M FeCl solution that was then heated in an oil bath (95°C). When the transparent solution turned to turbid, a piece of fresh NF was dipped into the solution for certain time to obtain NiFeOOH-X (X is the treatment time in FeCl , minutes). Later, the NiFeOOH was placed in 0.1 M Na S at 60°C, and marked as NiFeS-X (X is treatment time in Na S solution, minutes). The product NiFeS-10 indicates a Fe loading of ~0.52 mg/cm2 and an S loading of ~0.16 mg/cm2.

Preparation of RuO and Pt/C electrodes
To prepare RuO electrodes, 40 mg RuO powder, 60 μL Nafion, and 200 μL ethanol were mixed and sonicated for more than 1 h to obtain uniform RuO ink. Subsequently, the ink was dripped onto a clean NF surface and dried in a vacuum atmosphere at 60°C for 4 h. To obtain Pt/C electrodes, 40 mg Pt/C powder, 60 μL Nafion, and 300 μL ethanol were mixed and sonicated for more than 1 h to obtain uniform Pt/C ink. Subsequently, NF was directly immersed in Pt/C ink and dried for 4 h in a vacuum atmosphere at 60°C.

Materials characterizations
The morphology and microstructure of the material surface were measured by scanning electron microscopy (SEM, ZEISS Sigma 300) and scanning transmission electron microscopy (TEM, JEOL JEM-2100F), combined with energy dispersive X-ray spectroscopy (EDS). The surface species was checked by TOF-SIMS (Time of Flight Secondary Ion Mass Spectrometry, PHI nano TOF II Time-of-Flight SIMS). XRD was carried out using Cu-Ka radiation in Bruker D8 VENTURE diffractometer (Bruker, Germany), and the scanning speed was set at 2°/min, from 10° to 90° (λ=0.15418 nm, 40 kV, 40 mA). X-ray photoelectron spectroscopies (XPS) spectra were recorded on a Thermo Scientific ESCALAB 250Xi spectrometer, and the binding energies were checked with C 1s peak (284.8 eV). Raman measurements were performed at a research Raman microscope system (LabRAM HR Evolution) with a 532 nm laser excitation source. The in-situ X-ray absorption spectroscopy (XAS) tests at the Fe K-edges was performed by Spring 8 in the Japan Synchrotron Radiation Research Institute (JASRI). The Cl− and SO 2− ion concentrations were tested by ion chromatography (ICS-1100 Thermo, American).

Electrochemical measurements
The electrochemical measurements were conducted on the Biologic (VMP3, French) electrochemical workstation, except for the AEMWE test powered by a solar module (DH7001D, Donghua Testing Technology Co., Ltd). The Hg/HgO (1.0 M NaOH) was used as a reference electrode in a typical three-electrode system, and the volume of a single cell is 100 mL. The stirring was carried out at 500 rpm using an IKA C-MAG HP 7 stirrer to promote bubble detachment, enhance OH− diffusion, maintain a constant temperature, and thus obtain reliable performance. The Hg/HgO electrode was calibrated in H using a Pt wire as the working electrode. All measured potentials were calibrated to the reversible hydrogen electrode (RHE) by Eq. (1). The electrochemical impedance spectroscopy (EIS) measurements were recorded in the frequency range from 0.1 Hz to 100 kHz with an amplitude of 5 mV. To demonstrate the true electrode activity, the current scanning speed of LSV testing was set at 1 mV/s, with 100% iR compensation, where the R is 0.82 ± 0.025 Ω of alkaline seawater, as shown in Eq. (2). The electrochemical active surface area (ECSA) were estimated by cyclic voltammetry (CV) at different scanning rates of 20 to 100 mV/s within the non-Faraday current range. C dl and ECSA were estimated by Eq. (3) and Eq. (4), respectively. The Faraday efficiency (FE) of OER was calculated by comparing the amount of O collected from drainage with the theoretical amount, as shown in Eq. (5). Among them, M O is the molar amount of O 2 collected by drainage, N A is the Avogadro 2 number, J is the current density, mA cm−2, S is the effective area of the electrode, cm2, t is time, and e is the electron charge. In AEM testing, the catalytic reaction geometry area is 1×1 cm2. The anion exchange membrane is FAA-3-PK-75, and PtRu is sprayed on the cathode surface. The synthesis process of NiFeS is used for nickel felt, and it is also used as a gas diffusion layer. The half-membrane (1.2×1.2 cm in size) was pre-treated in 1 M KOH for over 12 h before AEM test, and the room temperature is specifically defined as 25 ± 2°C.

E(vs: RHE) = E(vs: Hg=HgO) + 0:0591 × pH + 0:098 (1)

Ecorrected = Emeasured - iR (2)

C dl = Δj = 2v (3)

ECSA = C dl = Cs (Cs = 0:04 mF cm−2 for NF) (4)

FE OER = (4 × M O × NA) = (J × S × t = e) (5)

DFT calculation details
The DFT and Ab initio molecular dynamics (AIMD) simulations were performed via the Vienna Ab initio Simulation Package (VASP). The cutoff energy of the plane wave was set to 420 eV, the exchange-related energy and ion-electron interactions were processed using Perdew Burke Ernzerh’s (PBE) generalized gradient approximation (GGA). For DFT, the structure of the Ni(OH) crystal was obtained from the Materials Project48. The NiFe LDH (001) surfaces were constructed, and the lattice parameter perpendicular to the surface direction was set to 20 Å to eliminate periodic effects along the Z-direction. The k-points mesh used in DFT calculations was set to 3×3×1, and the energy convergence criterion was set to 1×10−4 eV. For surface structures, only the atom positions were relaxed. The internal coordinates of each system were fully optimized until the residual Hellmann-Feynman force was less than 0.02 eV/Å per atom. The Gibbs free energy is obtained by Eq. (6). Among them, G(T) is the Gibbs free energy at temperature T, E DFT is the energy calculated by DFT, and G correct(T) is the Gibbs free energy correction at temperature T, which is obtained from VASPKIT software49. For AIMD calculation, the Monkhorst-Pack mesh of size 1×1×1 was used in k-points. The canonical ensemble NVT conditions were imposed by a Nose-Hoover thermostat with a target temperature of 300 K. The explicit solvent model used in the AIMD simulations consists of 85 H O molecules, 4 Cl−, 1 SO 2−, and 10 OH− ions. The MD time step was 1 fs, and all the systems were run for 5 ps to reach equilibrium. A fixed electric field in the Z direction was applied to each atom with selected voltage values of 0.1 V Å−1, and only the last 50 frames in the calculation results are counted. The results were visualized via Vesta software. The Van der Waals forces were addressed using the DFT-D3 method, while the DFT+U method applied U-J values of 3.29 and 3.40 eV for Fe and Ni, respectively. For slow-growth calculation, the free energy profile along a collective variable (CV, ξ) is scanned by an approximate slow-growth approach, and the INCREM values is setted as = 0.000150. The work required for the transformation from state 1 to state 2 can be calculated through integration using Eq. (7), where the ξ is the velocity of the collective variable and ∂V ∂ξ is the force. As the simulation time is infinitely small, the work corresponds to the free energy change between state 1 and state 2. The hydrogen bond strength calculation was performed using the Gaussian 09 code. Geometry optimization and frequency analysis were conducted at the B3LYP/Def2TZVP level of theory. The SMD implicit solvation model was employed to simulate the solvent environment, combined with D3(BJ) dispersion correction (B3LYP-D3(BJ)). The strength of hydrogen bonds is calculated via Eq. (8), where E H is the energy of hydrogen bond, E slab is the energy of the H O molecule clusters, and E OH−/ Cl− is the energy of OH− or Cl−. The weak intermolecular interaction force is calculated by Multiwfn software and displayed in VMD software.

G(T) = E DFT + G correct(T) (6)

wirrev = ∫ ξ1 ξ2 (∂V(q) = ∂ξ) dξ (7)

E H = E total - E slab - E OH−=Cl− (8)

Peer review rounds
[{"round": 1, "reviewer_comments": {"reviewer_1": "The manuscript describes a NiFeS catalyst for seawater electrolysis and claims that an in-situ-formed SO42- layer kinetically repels Cl⁻ through a hydrogen-bond network. Regrettably, the work cannot be recommended for publication in its present form for the following reasons:\n1. Lack of synthetic novelty. The route—Fe3+ etching of metallic Ni followed by sulfidation—has been extensively reported for both NiFe-LDH and NiFeS systems (e.g., https://doi.org/10.1002/admi.201900788; doi: 10.1039/D0TA02337J). The authors provide no clear advance beyond these precedents.\n2. Performance metrics are not very uncompetitive. At 100 mA cm-2 the overpotential (268.8 mV) is higher than those of recently published non-noble seawater anodes that already achieve < 200 mV at the same current density. Similarly, the AEM cell voltage of 2.134 V at 1 A cm-2 is above the 1.7–12.0 V range now routinely reported for AEM devices.\n3.Critical phase evolution and catalytic behavior against long-term AEM stability are missing. Without evidence of membrane integrity, catalyst delamination, sulfate evolution, or chloride crossover after > 100 h, the industrial relevance of the AEM performance remains speculative.\n4. Additional ambiguities and inconsistencies further undermine the manuscript:\na) The high-resolution TEM images display clear lattice fringes indicative of a crystalline (Ni)FeS phase, yet the powder XRD pattern is essentially featureless, implying an amorphous structure. This stark discrepancy is neither rationalized nor discussed.\nb) The dynamic leaching-redeposition of Fe species during catalysis is not addressed. Without tracking Fe dissolution and its subsequent redeposition (e.g., via ICP-MS) , the impact of Fe specie on activity and stability remains unclear.\nc) The catalytic mechanism relies heavily on oxidation-state changes of Ni and Fe, full operando XANES/EXAFS series under working potentials are required to corroborate the proposed redox pathways.\nd) The DFT section lacks free-energy profiles for key adsorption steps.", "reviewer_3": "Reviewer remarks:\n1. Page 5. Values of binding energies in XPS spectra are given with too high precision (for example710.08 eV). These values should be given only with one decimal point and not with two decimal points, since the shown XPS spectra are pretty broad and noisy.\n2. Page 5. The assignment of the peaks for Fe2+ and Fe3+ in the XPS spectra Fe 2p is wrong. The Fe 2+ is not expressed only at 723.39 eV and 725.94 eV, but mainly around 713 eV.\n3. Page 5. Authors claim that “On the NiFeS surface, the Ni 2p orbitals shift to lower energy levels than NiFeOOH, indicating a higher surface electron cloud density, which is related to weaker sulfur element polarity.” This observation is not correct from the presented spectra. The Ni 2p spectra in Figure S9b do not show this effect. They are rather broad, noisy, fitted in a deficient way (only two peaks for the Ni 2p3/2 spectrum), using the linear background instead of the Shirley type…\n4. Page 5 and page 14/Figure S44. It would be interesting to present the surface composition from XPS spectra of catalysts before and after treatment/exposure to compare the amount of S, Fe, Ni, O, … What is the interpretation for the XPS spectra in Figure S44? It is not visible.\n5. Page 13. Authors described the XPS spectra in Figure S30 as: “Notably, this non-catalytic alkaline scale is identified as the primary contributor to the activity degradation, and the XPS results provide compelling evidence supporting this conclusion (Figure S30).” It is not clear how this is concluded from the XPS spectra in Figure S30?\n6. Page 10, Figure 3i and S20, S21, ToF-SIMS maps of species. Only the qualitative evaluation from SIMS images is given, although some semi-quantitative evaluation may be obtained from the total intensities of the SIMS signal in maps."}, "author_response": "Reviewer #1 (Remarks to the Author):\nThe manuscript describes a NiFeS catalyst for seawater electrolysis and claims that an in-situ-formed SO 2- layer kinetically repels Cl- through a hydrogen-bond network. Regrettably, the work 4 cannot be recommended for publication in its present form for the following reasons:\n1. Lack of synthetic novelty. The route-Fe3+ etching of metallic Ni followed by sulfidation-has been extensively reported for both NiFe-LDH and NiFeS systems (e.g., https://doi.org/10.1002/admi.201900788; doi: 10.1039/D0TA02337J). The authors provide no clear advance beyond these precedents.\nResponse: We sincerely thank the reviewer for raising this critical point and for providing these relevant references. We acknowledge that the foundational chemistry of Fe3+ etching and subsequent sulfidation has been previously reported. However, we respectfully argue that our work represents a substantial advance beyond these precedents, not merely in synthesis, but in practical engineering, fundamental understanding of catalyst evolution, and the discovery of a novel mechanism. Our key advancements are elaborated below.\n1. Beyond initial synthesis: focusing on the dynamic reconstruction and the final active interface. We fully agree with a fundamental principle in the field, which our work strongly supports: most NiFe-based catalysts, including our NiFeS, undergo dynamic surface reconstruction under alkaline OER conditions, ultimately transforming into the active NiFe LDH phase (Adv. Energy Mater. (2025): e04101., Energy Environ. Sci., 2024,17, 6897-6942). Therefore, we contend that the initial synthesis method, while important for creating a specific precursor, is not the ultimate goal. The true scientific importance lies in how the engineered precursor influences the dynamic reconstruction process and dictates the properties of the final active interface (Energy Environ. Sci., 2025,18, 1952-1962, Sci. Adv.11, eads0861(2025).). In this work, we re-examine the relationship between catalyst reconstruction and corrosion resistance mechanisms, and propose an insightful perspective that goes beyond conventional electrostatic repulsion. The unique \" etching-sulfurization\" synthesis creates a precursor that, upon reconstruction, leaves behind a stable, SO 2--rich layer at the catalyst-electrolyte interface. It is precisely this SO 2--rich layer that 4 4 reorganizes the interfacial water structure, thereby enabling selective OH- transport. This tailored interface is the key to the exceptional performance and stability, shifting the focus from \"what we synthesized\" to \"what the material becomes and how it behaves during operation.\" We firmly believe that the insights we have presented reconcile the traditional relationship between catalyst surface reconstruction and the enhancement of corrosion resistance, while also offering novel perspectives on the reconstruction design and underlying mechanisms of catalysts for hydrogen production via seawater electrolysis.\nFurthermore, we have experimentally validated the reliability of this ion- sieving effect using NiFeS as the object. The NiFeS catalyst demonstrated exceptional stability, maintaining continuous operation for over 2000 hours at a high current density of 1.0 A·cm-2, a performance that surpasses most recently reported materials, as shown in Figure R1 and Table R1.\n[... Author response continues with detailed rebuttal and data tables ...]"}, {"round": 2, "reviewer_comments": {"reviewer_1": "The authors have satisfactorily addressed most of my previous comments. I therefore recommend acceptance of the paper in its present form.", "reviewer_2": "The quality of the revised manuscript has improved considerably, and most of my previous concerns have been appropriately addressed. Therefore, the manuscript may be considered for publication after the authors address the following minor revisions:\n1. The authors should clearly and comprehensively describe the AIMD and DFT methodologies, including the construction of the AIMD model and the models used in both AIMD and DFT simulations. This is essential to ensure the reliability and reproducibility of the reported results.\n2. The constructed AIMD results, particularly those obtained using the slow-growth method, are not discussed in the main manuscript. These results should be briefly explained in the manuscript and the full data should be added to the Supporting Information for completeness.\n3. The authors should clearly define the collective variable used in their simulations. For example, if the CV corresponds to the Cl---Fe distance or another structural descriptor, this must be explicitly stated.\n4. Several statements require more detailed explanation for clarity. For instance, on page 10, the authors state: “Trajectory analysis revealed that the SO42-containing system exhibited significantly enriched OH- enrichment and effective Cl-exclusion at the anode interface, as shown in Figure 3e.” Such statements must be thoroughly discussed to ensure the reader can clearly understand the underlying reasoning and evidence.\n5. Several references, such as reference 38, are cited without meaningful context. The authors should explain why each reference is included, e.g., whether it supports the statement made, describes a similar method, or shows that their findings are consistent with previously published results. A careful revision is required to ensure all citations are properly justified and informative.", "reviewer_3": "The authors have successfully considered the main concerns raised by the reviewer regarding the XPS and ToF-SIMS measurements. They have enhanced the quality of data presentation, provided additional data, and expanded on their interpretation of experimental findings.\nHowever, there are still a few minor comments to address:\nIn Table S3, the concentrations derived from XPS analyses should be reported with only one decimal point since the accuracy of XPS data is limited."}, "author_response": "Reviewer #1 (Remarks to the Author):\nThe authors have satisfactorily addressed most of my previous comments. I therefore recommend acceptance of the paper in its present form.\nResponse: Thank you for your positive feedback.\nReviewer #2 (Remarks to the Author):\nThe quality of the revised manuscript has improved considerably, and most of my previous concerns have been appropriately addressed. Therefore, the manuscript may be considered for publication after the authors address the following minor revisions:\n1. The authors should clearly and comprehensively describe the AIMD and DFT methodologies, including the construction of the AIMD model and the models used in both AIMD and DFT simulations. This is essential to ensure the reliability and reproducibility of the reported results.\nResponse: Thank you for your kind suggestion. We acknowledge that a complete computational model and its parameters are essential for computational reproducibility. Accordingly, in response to your comments, we have supplemented the detailed parameters and model descriptions for the DFT, AIMD, and slow-growth calculations, respectively.\n[... Author response continues ...]"}, {"round": 3, "reviewer_comments": {"reviewer_2": "The authors have successfully addressed all my concerns. I recommend acceptance in its current form."}, "author_response": ""}]

Decision letter
Report letter
(Manuscript ID: NCOMMS-25-53658)
In this work, Yang Yu et al. successfully developed a NiFeS catalyst for selective OER, providing an effective strategy to suppress the competing chlorine evolution reaction through a hydrogen bond network that hinders Cl- mass transfer across the interface. This research represents a significant contribution in the development of selective OER catalysts for seawater electrolysis. However, several key concerns must be addressed before reconsideration for publication, as detailed below.
1. Regarding the AIMD methodology, I recommend using DFT+U instead of standard DFT to more accurately optimize the system, as this approach better accounts for the strong electron localization in Fe and Ni.
2. To further clarify the role of the SO₄²⁻ anion in suppressing Cl⁻ migration, I strongly recommend that the authors compute the free energy barrier for Cl⁻ transfer to the interface, both with and without SO₄²⁻ present. The slow-growth method could be employed to quantify these barriers, providing direct mechanistic insight into how sulfate modulates Cl⁻ mobility.
3. I suggest the authors extend this strategy to another catalyst system that favors the CER. By introducing SO₄²⁻ and analyzing its impact, they could further validate its proposed role in suppressing CER. This comparative approach would strengthen the mechanistic conclusions.
4. Could the authors please clarify the pH dependence of this strategy? Understanding how the system's performance varies with pH would provide valuable insights into its practical applicability.
5. While the authors have thoroughly investigated the role of SO₄²⁻ in suppressing CER through hydrogen bond networks, its potential influence on the OER activity and mechanism warrants further discussion. Specifically: (1) How might the SO₄²⁻-induced hydrogen bond network affect the OER kinetics and overpotential? (2) Could this structural modification alter the conventional OER pathway? Clarifying these aspects would provide a more comprehensive understanding of the catalyst's dual functionality.
6. The authors propose the formation of high-valence metal species following OER, but the evidence supporting this claim requires clarification.
7. While the authors attribute anode degradation primarily to Mg²⁺/Ca²⁺ scale deposition and the conversion of NiFeS to NiFe LDH, the role of sulfur vacancy formation remains unclear. Given that sulfur loss was observed (47–67% retention via ICP), it is critical to address whether sulfur vacancies facilitate LDH formation or accelerate corrosion.