Surface halogenation engineering for reversible silicon-based solid-state batteries

Domain: physics
Published: 2025-12-27
Editorial outcome: major_revision

Abstract
Silicon-based solid-state batteries are promising next-generation high-energy-density technologies. However, poor (electro)chemical compatibility between silicon negative electrodes and solid electrolytes (e.g., Li6PS5Cl) plus sluggish interfacial kinetics severely limit their reversibility and Coulombic efficiency. Here, we propose a surface halogenation strategy that transforms the native amorphous SiO2 passivation layer on silicon particles into a functional Al(Si)OCl composite surface via controlled reaction with AlCl3. This artificial interphase reconciles interfacial incompatibility and enables fast ionic/electronic transport, suppressing irreversible lithium loss. The optimized negative electrode achieves a high initial Coulombic efficiency of 94.3% in half-cells and 85.6% initial Coulombic efficiency (86.6% with pre-lithiation) in full cells paired with LiNi0.88Co0.09Mn0.03O2. Enhanced reversibility further delivers long-term cyclability. The optimized negative electrode delivers 86% capacity retention and 99.998% average Coulombic efficiency over 200 cycles. Even at high-loading (>10 mAh cm-2, and no adhesives/conductive carbon/electrolyte), it retains 72% capacity after 500 cycles. The full cells maintain 80% capacity after 200 cycles at 1C, with an average Coulombic efficiency exceeding 99.95%. The versatility of this halogenation strategy underscores halide chemistry’s broad potential in advancing high-performance, reversible silicon-based solid-state batteries.

Introduction
All solid-state batteries (SSBs) are recognized as a leading next-generation electrochemical energy storage technology due to their superior energy density and inherent safety advantages1,2. Silicon (Si) negative electrodes are particularly promising for SSBs due to their exceptional theoretical specific capacity of 3579 mAh g–¹, which approaches that of metallic lithium (Li) negative electrodes (3860 mAh g–¹)3,4. Moreover, Si negative electrodes demonstrate an intrinsic resistance to Li dendrite formation owing to their moderate operating potential (~0.1–0.5 V vs. Li+/Li), and exhibit higher critical current densities compared to lithium-indium (Li-In) alloys3,5. The inherent safety advantages and favorable electrochemical performance highlight the promising development potential of Si-based SSBs.

However, Si-based SSBs currently face challenges in suboptimal electrochemical reversibility and insufficient Coulombic efficiency (CE), including initial Coulombic efficiency (ICE) and cycling CE, significantly limiting their practical capacity utilization and energy density3,6. The typical ICE ranges between 78 and 90% in half-cell configurations employing Li6PS5Cl (LPSC) solid electrolytes (SEs), while only approximately 80% could be achieved in full cells paired with LiNi1-x-yCoxMnyO2 positive electrodes7–12. Research indicates that irreversible capacity loss during initial charge-discharge cycles includes Li consumption for solid electrolyte interphase (SEI) formation (C-Li) and kinetically trapped Li (K-Li), accounting for approximately 12% and 11% of the total capacity, respectively5. C-Li stems from the poor (electro)chemical compatibility at the Si|LPSC interface3,13. K-Li is primarily caused by inadequate bulk and interfacial ion/electron transport kinetics, resulting in the formation of insulating dead Li-Si alloy14,15 (Fig. 1a). Although it has been suggested that SEI could be “frozen” at the interface once it was formed in SSBs, potentially ensuring cycling stability5,13, subsequent experimental investigations revealed continuous SEI growth and void formation at the 2-dimensional Si|LPSC interface, leading to increased impedance and suboptimal cycling CE (<99% during the first 50 cycles)3,16.

To tackle these issues, pre-lithiation and Li-Si alloys incorporation strategies are the most straightforward approaches for compensating irreversible Li loss4,6,14. For example, a direct Li metal powder pre-lithiation process has been proposed to achieve 85.2% capacity retention after 200 cycles and an average CE of 99.6%17. Similarly, LiSi alloy negative electrode exhibits an enhanced ICE from 75.4% to 80.8% with an average CE of 99.9% when paired with NCM811 positive electrode6. However, excessive Li usage raises safety concerns (such as soft short circuits), while the highly active Li-Si alloys would exacerbate interfacial side reactions4,18,19.

In addition, other major research efforts have focused on stabilizing SEI film. A poly(vinylidene fluoride-co-hexafluoropropylene) (PVDF-HFP)/Li1.3Al0.3Ti1.7(PO4)3 SE is developed to construct a stable LiF-rich SEI layer, achieving 89.6% ICE for Si negative electrodes17. LiBH4-LiI (LBHI) SEs demonstrate electrochemical compatibility with Si negative electrodes, effectively suppressing continuous SEI formation. Although the LBHI-based half-cell delivers a high ICE of 96.2%, it is achieved under a high temperature 60 °C and the cycling performance is limited by compromised ionic transport kinetics due to the substitution of LPSC with LBHI, resulting in an insufficient average CE of 99.34%13.

Despite these advancements, cycling CE remains below the critical 99.98% threshold (calculated for 80% capacity retention over 1000 cycles20). Thus, the ICE of the full cell with commercial NCM-based positive electrodes is still suboptimal, typically below 85%. In this regard, we are inspired by halogen chemistry, as the potential of both halide modifications and halide-based SEs in stabilizing interfaces21–23. Capitalizing on the thermodynamic reactivity between the SiO2 and halide salts such as AlCl3 24,25, a surface halogenation strategy is developed through a simple mixing and heating process of AlCl3 and Si (denoted as Si@AlCl3) (Fig. 1b). It is revealed that the inert amorphous SiO2 (a-SiO2) layer on the Si surface can be converted to Al(Si)OCl. This transformation endows the Si@AlCl3|LPSC interface with enhanced chemical/electrochemical stability, significantly improving ion/electron transport kinetics. This optimization not only effectively suppresses C-Li induced by interfacial side reactions, but also mitigates K-Li, thereby breaking the reversibility barrier. The modified system achieves an ICE of 94.3% in half cell and 85.6% in full cell with commercial LiNi0.88Co0.09Mn0.03O2 (NCM) positive electrodes at 25°C, further boosting a high ICE of 86.6% with pre-lithiation. The improved reversibility improves cycling stability. The Si@AlCl3 negative electrode exhibits a capacity retention increase from 14% (for Si) to 86% after 200 cycles at 3C, with average CE improving from 99.4% to 99.998%. At high loading (>10 mAh cm–2), the Si@AlCl3 negative electrode without any adhesives, conductive carbon or SE maintains 72% capacity retention after 500 cycles with a current density of 5.1 mA cm–2. The assembled NCM-based SSBs show 80% capacity retention after 200 cycles at 200 mA g–1. Furthermore, the universal applicability of halide salts in enhancing the reversibility and capacity of the Si negative electrode is validated, highlighting the significant potential of the surface halogenation engineering for Si-based SSBs.

Results
AlCl3-mediated surface halogenation modification
AlCl3 was selected as the primary research subject owing to its suitably low boiling points among halide salts, as well as its spontaneous reactivity with SiO2 (Fig. S1 and Table S1). As depicted in Fig. 2a, b, pre-experimental validation confirms the thermodynamically driven phase transformation process, where a-SiO2 reacts with AlCl3 to form AlOCl and SiCl4 (SA), as evidenced by X-ray diffraction (XRD), Raman spectroscopy, and X-ray photoelectron spectroscopy (XPS) analyses26,27 (Figs. S2–S4; the detail is discussed in the Supplementary information). The chemical stability of SA to Li has also been experimentally verified (Fig. S5; the detail is provided in the Supplementary information). Consequently, we conducted the halogenation reaction by sealing a mixture of AlCl3 and micron-sized Si particles and heating it at 180°C. Electrochemical evaluation of Si negative electrodes with varying AlCl3 mass fractions reveals a non-monotonic relationship between ICE and AlCl3 content, peaking at 10% incorporation (Fig. S6 and Table S2). Therefore, the optimal 10% AlCl3-incorporated system (designated as Si@AlCl3) is selected for subsequent investigations.

Scanning electron microscopy (SEM) imaging coupled with energy dispersive spectroscopy (EDS) elemental mapping confirms the presence and characteristic distribution of Si, Al, and Cl elements throughout both Si@AlCl3 particles (Fig. S7) and the corresponding electrodes (Figs. S8–S9), showing relatively consistent elemental dispersion. The XRD analysis reveals the main phase of crystalline Si with the emergence of a weak AlCl3 signal at 31° compared to pristine Si (Fig. S10). Raman spectroscopy demonstrates a new high-wavenumber peak alongside the Si-Si vibrational peak, consistent with SiCl(O) observed in pre-experiments (Figs. 2c and S2). The XPS analysis reveals a red shift in the Si 2p and O 1s peaks for Si@AlCl3 particles compared to pure Si, which is also consistent with the shift characteristic of SiCl(O). Specifically, the O 1s peak shifts from 532.4 eV (SiO2)28 to a lower binding energy of 531.7 eV, while the Al 2p peak of Si@AlCl3 displays a slight decrease in binding energy to 74.7 eV relative to pure AlCl3 (75.1 eV). This aligns with reported values for AlOCl29,30. The Cl 2p peaks of Si@AlCl3 at 199.5 and 201 eV exhibit binding energies consistent with pure AlCl3 (Figs. 2d–f and S11; S12a–c). Furthermore, high-resolution transmission electron microscope (HRTEM) imaging reveals significant surface differences between pristine Si and Si@AlCl3. The initial Si particles display an a-SiO2 surface layer of 5-10 nm (Fig. 2g, h). After reacting with AlCl3, a crystalline layer exhibiting the triclinic AlCl3 and AlOCl interplanar spacing characteristic is generated, and a-SiO2 is completely eliminated on Si@AlCl3 surface (Figs. 2i and S13). EDS elemental mapping and line scan profiles further confirm the halogenated modification layer (Figs. 2j and S14). Based on the above experimental observations, we propose a mechanistic pathway wherein AlCl3, functioning as a Lewis acid catalyst, facilitates the chemical modification of the native SiO2 layer31, consequently generating Al(Si)OCl and residual AlCl3 (SA).

Study of irreversible Li loss
As illustrated in Fig. 3a, the reversibility of Si and Si@AlCl3 was explored through electrochemical measurements at 25°C. The Si@AlCl3 half-cell demonstrates an enhanced ICE of 94.3%, compared to 88.4% for the pure Si negative electrode. Additionally, the specific charge capacity of the Si active material improves from 2924 mAh g–1 to 3324.6 mAh g–1 (Fig. 3c). Correspondingly, ICE of the full cell paired with a NCM positive electrode increases from around 79.2% to a maximum of 85.6% (Figs. 3b and S15; Table S3). It is demonstrated that the Si@AlCl3 negative electrode exhibits enhanced reversibility compared with the pure Si negative electrode.

The reversibility in Si-based negative electrodes is closely associated with the mitigation of irreversible Li loss. Neutron depth profiling (NDP), a non-destructive neutron-based characterization technique, is particularly suited for studying Li concentration spatial distribution due to its high sensitivity to Li32. Therefore, NDP (Figs. 3d and S16) coupled with gas chromatography (GC) (Fig. 3e) was employed to analyze Si@AlCl3 and pristine Si negative electrodes during the initial (de)lithiation process. The results reveal a significant reduction in irreversible Li loss in Si@AlCl3 compared to pristine Si negative electrodes. The specific distribution of irreversible Li is characterized as follows: (1) A distinct concentration gradient of residual Li is observed within the Si-based negative electrode layer, exhibiting a gradual decrease from the Si|SE interface towards the bulk region. The interfacial region demonstrates higher Li content (corresponding to C-Li for SEI formation), accounting for the majority of irreversible Li, while the bulk region displays lower Li content (corresponding to K-Li), representing a minor fraction of irreversible Li (Fig. 3f). (2) In Si@AlCl3 negative electrodes, the Li concentration near the interface is substantially reduced compared to pristine Si negative electrodes, indicating suppressed C-Li (decreased from 9.9% to 7.5%). (3) The bulk Li content in Si@AlCl3 negative electrodes exhibits a remarkable decrease (from 1.5% to 0.1%), signifying reduced K-Li. In addition, a constant-current/constant-voltage (CC-CV) charging protocol was implemented for the Si-based negative electrode. Comparative results reveal that the CV process effectively activates residual Li localized near the SEI within the Si-based negative electrode, thereby boosting 96.1% ICE for Si@AlCl3 (Figs. S17–S19).

Furthermore, the irreversible Li loss was investigated following multiple charge-discharge cycles (Fig. 3g–j). After 50 cycles, a prominent signal of kinetically trapped LixSi is observed at 54.5 eV for the Si negative electrode33, while no such signal is detected for the Si@AlCl3 negative electrode, indicating superior cycling reversibility of the latter. Notably, the elemental Al signal appears at 72.8 eV in Al 2p spectra and LiCl signals at 198.4 eV and 200 eV in Cl 2p spectra, attributed to the reduction of lithiated AlCl3. In contrast, the Al 2p (~75 eV) and Cl 2p signatures corresponding to AlOCl remain stable (Fig. S20a–d). Further analysis of the GCD profiles and CV curves confirms negligible capacity contributions from SA and AlCl3 (Fig. S21). These results demonstrate that AlCl3 exhibits weak electrochemical reactivity during cycling, as evidenced by the absence of an Al metal signal in the XPS spectra of Si@AlCl3 after 5 cycles, with the Al 2p and Cl 2p binding energies consistent with reported values for AlOCl-LiCl30 (Fig. S22). The LiCl and Al generated after multiple cycles can serve as components of the SEI, effectively mitigating the reduction of LPSC22. More importantly, AlOCl maintains relatively electrochemically stability throughout cycling, which plays a crucial role in ensuring the reversibility of Si-based negative electrodes in subsequent cycles.

Electrochemical/chemical compatibility of Si-based negative electrode with LPSC
Given that the C-Li caused by interfacial side reactions makes the majority contribution to irreversible Li loss, the chemical compatibility was first investigated to elucidate the underlying mechanisms by which surface halogenation enhances electrochemical reversibility and improves the ICE. A physical mixing approach was employed, where Si-based particles were uniformly mixed with LPSC powder with a 2:1 mass ratio to obtain Si+LPSC and Si@AlCl3+LPSC, respectively (Fig. 4a). XPS spectra reveal that pure Si particles exhibit poor chemical stability with LPSC powder (Figs. 4b and S23a). The Si 2p binding energy for Si+LPSC is observed at 98.7 and 99.4 eV, indicating a red-shift compared to the binding energies of pure Si at 99.0 and 99.6 eV. The S 2p binding energy indicates the presence of minor Li2S signal peaks at 162.1 and 163.5 eV, along with oxidized states at 167.1 and 168.2 eV34. Additionally, the P 2p binding energy indicates the presence of LixP signal peaks at 128.7 and 129.5 eV, along with oxidized states at 133.5 and 134.6 eV35. The substantial formation of oxidized states is hypothesized to result from side reactions between the surface oxide layer of the Si particle and LPSC3. In contrast, the Si 2p, S 2p, P 2p, and Al 2p binding energies for Si@AlCl3+LPSC remain almost unchanged compared with pure LPSC or Si@AlCl3, indicating enhanced stability (Figs. 4c and S23b). The enhancement can be attributed to AlOCl in the SA layer and potentially cycling-derived AlOCl-LiCl, which exhibit higher chemical compatibility with LPSC compared to conventional oxyhalides (e.g., LiTaOCl4) (Figs. S24–S27).

Electrochemical impedance spectroscopy (EIS) analysis further reveals that the impedance of Si+LiPSCl is significantly higher than that of Si@AlCl3+LiPSCl. Even worse, the impedance of Si+LPSC is gradually increasing over time, while that of Si@AlCl3+LPSC slightly decreases and remains stable (Fig. 4d). The distribution of relaxation times analysis indicates that, within the range of 10–5 to 10–4 s, the peaks corresponding to grain boundary and interfacial impedances for Si+LiPSCl are significantly higher36 (Figs. 4e and S28). According to equivalent circuit fitting (Fig. S29 and Table S4), the grain boundary and interfacial impedances in the mid-to-high frequency region are 5.07×104 Ω and 1.23×104 Ω, respectively, with the main contribution coming from the low-frequency region, valued at 5.55×105 Ω. In contrast, the observed impedances in the mid-to-high frequency region for the Si@AlCl3+LPSC system are significantly lower, at only 4938 Ω and 3222 Ω. This validates that the enhanced chemical stability facilitates to form a more stable initial Si@AlCl3|LPSC interface.

To further investigate the electrochemical stability of Si@AlCl3 with LPSC, two half-cell configurations were assembled: Li-In|LPSC|Si@AlCl3 and Li-In|LPSC|Si. EIS measurements were conducted after one charge-discharge cycle, followed by a 12-h rest period (Fig. 4f). The fitted EIS data reveal a significantly higher charge transfer resistance (Rct) of 189 Ω for Li-In|LPSC|Si compared to 27 Ω for Li-In|LPSC|Si@AlCl3 (Fig. 4g, h and Table S5). This substantial difference primarily stems from the formation of a large interfacial impedance (147 Ω) in addition to the bulk impedance (15 Ω) in the Li-In|LPSC|Si, whereas the Li-In|LPSC|Si@AlCl3 exhibits only 5 Ω bulk impedance. SEM and EDS analysis of Si negative electrode after cycling reveals extensive pores generated at the Si|LPSC interface, along with significant accumulation of Cl element, as evidenced in Figs. 4i and S30. This phenomenon is further clearly observed in cross-sectional morphology and elemental line scans, which indicate vertical cracks within the cycled Si negative electrode and notable redistribution of Si, Cl, P, and S (Figs. S31, S32a, b and S34a). XPS depth profiling confirms the formation of LiCl as a side reaction product. These findings suggest poor electrochemical stability between Si and LPSC (Fig. S33). In contrast, the Si@AlCl3 negative electrode maintains a well-defined interface and element distribution with LPSC without detectable side reactions (Figs. 4j and S32c; S34b). It means that the Si negative electrode suffers from both poor chemical and electrochemical stability with LPSC, which in turn detrimentally affects transport dynamics at the interface and bulk of the negative electrode. In contrast, the Si@AlCl3 negative electrode exhibits enhanced chemical/electrochemical compatibility and interfacial stability.

Charge transport kinetics and electrochemical performance
The electrochemical kinetics were systematically investigated. As illustrated in Fig. 5a, b, Li-In|LPSC|Si@AlCl3 exhibits higher rate performance compared to Li-In|LPSC|Si, delivering a specific charge capacity of 3306.3 mAh g–1 at 350 mA g–1 and maintaining 1891.2 mAh g–1 at 7 A g–1, with 57.2% capacity retention. Notably, the system preserves high rate capability even under high mass loading conditions (>9 mAh cm–2) (Fig. 5c). NCM-based full-cell with Si@AlCl3 negative electrode exhibits a stable specific capacity of 124.7 mAh g–1 at 100 mA g–1, surpassing the 56.7 mAh g–1 of the cell with a pristine Si negative electrode (Fig. S35).

The enhanced rate capability can be attributed to improved charge transfer kinetics. As evidenced by direct current (DC) polarization measurements, the electronic conductivity of pristine Si is only 1.98×10–7 S cm–1, whereas Si@AlCl3 exhibits an over 40-fold improvement, reaching 7.93×10–6 S cm–1 (Figs. 5d–f and S36). This enhancement in electronic conductivity can be attributed to the SA layer, which not only disrupts the insulating a-SiO2 layer on the surface of the Si particles but also improves the interface contact, leading to denser particle accumulation (Fig. S37).

Furthermore, distinct differences in Rct can be observed during the charge/discharge processes (Fig. S38 and Fig. 5g). The initial Rct of Li-In|LPSC|Si@AlCl3 is approximately 71 Ω, significantly lower than that of Li-In|LPSC|Si (470 Ω). During discharge, the formation of highly conductive Li-Si alloys results in a significant decrease in Rct for both configurations. The Rct values for Li-In|LPSC|Si@AlCl3 and Li-In|LPSC|Si decrease to 18 Ω and 16.5 Ω, respectively. In the charged state, the Li-In|LPSC|Si@AlCl3 maintains a relatively low Rct of 36 Ω, while the Li-In|LPSC|Si exhibits a higher Rct of 443 Ω, as determined through equivalent circuit fitting. DRT analysis indicates that after lithiation/delithiation, the grain boundary and interphase impedance of Li-In|LPSC|Si in the range of 10–6-10–4 s increase, contributing to the overall higher Rct. Therefore, the stabilized interface ensures the fast interface charge transfer dynamics.

A comparison of the galvanostatic intermittent titration technique (GITT) curves during discharge reveals the differences in the reaction mechanisms (Fig. S39a). For Li-In|LPSC|Si@AlCl3, the discharge process exhibits a direct single-phase transformation from Si to Li3.75Si, bypassing the initial solid-solution reaction (voltage slope) observed in Li-In|LPSC|Si, which first forms LixSi (solid solution) followed by a phase transformation (voltage plateau)37. The calculated diffusion coefficients further support the interpretation (Fig. S39b). In the early discharge stage, the phase transition in Li-In|LPSC|Si is kinetically constrained, thereby activating a solid-solution reaction mechanism (manifested by an elevated diffusion coefficient). Remarkably, during the subsequent phase transition, Li-In|LPSC|Si@AlCl3 exhibits a higher average diffusion coefficient (2.96×10–10 cm2 s–1) compared to Li-In|LPSC|Si (2.15×10–10 cm2 s–1), demonstrating enhanced ion transport and reaction kinetics (Fig. 5h).

Benefiting from the improved electro/chemical compatibility and enhanced ion/electron transport kinetics, the Si@AlCl3 negative electrode shows improved reversibility and cyclic stability. In contrast to the rapid capacity decay observed in the pristine Si negative electrode after tens of cycles, the Si@AlCl3 negative electrode maintains a capacity of approximately 700 mAh g⁻¹ after 400 cycles at 10.5 A g–1, with its average CE increasing from 99.945% to 99.998% and cumulative CE for 400 cycles increasing from 39.2% to 88.5% (Figs. 6a, b and S40). Notably, Li-In|LPSC|Si@AlCl3 shows consistently higher average CE than that of Li-In|LPSC|Si across various current rates (Fig. S41). Furthermore, stable cycling performance is maintained even under low external pressure (20 MPa) (Fig. S42). In high-loading Si-based half-cells with capacities exceeding 10 mAh cm–2 without adhesives, conductive carbon or additional SE additives, the Li-In|LPSC|Si exhibit significant kinetic limitations and Li plating issues, resulting in an ICE as low as 61.5%, and a two-step reaction mechanism during discharge. In contrast, the Li-In|LPSC|Si@AlCl3 achieves a higher ICE of 78.9% and follows a single-step reaction mechanism (Fig. S43). Furthermore, while the high-loading Li-In|LPSC|Si demonstrates rapid capacity degradation, suffering nearly complete capacity loss after only 80 cycles at 25°C and 5.1 mA cm–2, with an average CE exceeding 102% indicative of pronounced lithium plating and side reactions14 (Figs. 6c, d and S44). In contrast, the Li-In|LPSC|Si@AlCl3 system exhibits improved cycling stability, maintaining 72% capacity retention over 500 cycles at 5.1 mA cm–2 with an average CE of 100.06%. Remarkably, at a reduced current density of 1.5 mA cm–2, the system achieved higher capacity utilization (5–6 mAh cm–2) while retaining capacity retention of 88.7% over 50 cycles (Fig. S45).

When paired with a commercial NCM positive electrode, the full cell demonstrates stable cycling performance at 200 mA g-1 for 200 cycles with a high capacity retention of 80% and an average CE exceeding 99.95% (Fig. S46). Furthermore, the pre-lithiated Si@AlCl3||NCM full cell38 achieves an initial charge capacity of 230.8 mAh g–1 and a discharge capacity of 199.9 mAh g–1, with a high ICE of 86.6% (Fig. 6e, f).

Furthermore, comparative analysis of modification strategies reveals the advantages of our halide salt-driven surface halogenation engineering. The results reveal that the incorporation of conductive carbon leads to suboptimal negative electrode capacity and ICE (<85%). Meanwhile, Si-based systems are constrained by the limited active material content, resulting in restricted negative electrode capacity utilization (<2600 mAh g–1). Although the structural engineering of Si negative electrodes ensures capacity retention, it still falls short in achieving satisfactory ICE. In contrast, the halide salt-driven modification strategy demonstrates universally enhanced performance in both ICE and capacity utilization, while maintaining stable cycling in full-cell configurations (Fig. 6g and S47–S48; Table S6).

In summary, we have developed a halide salts-driven surface halogenation engineering, exemplified by AlCl3, which has been demonstrated to form an SA composite layer on Si particles. This composite layer enhances the (electro)chemical compatibility between the Si negative electrode and LPSC electrolyte, ensuring stable interfacial transport kinetics. Simultaneously, it eliminates the intrinsic a-SiO2 layer on the Si surface, thereby improving the ion/electron transport kinetics of the Si-based negative electrode. Notably, the AlOCl component within the composite layer exhibits high electrochemical stability during cycling, ensuring sustained performance enhancement in subsequent cycles. As a result, this strategy effectively reduces C-Li and K-Li during both the initial and subsequent cycles, ultimately promoting capacity utilization and enhanced reversibility of the Si-based negative electrode.

The modified Si@AlCl3 negative electrode demonstrates remarkable performance in half-cells, achieving a high ICE of 94.3% and a specific charge capacity of 3324.6 mAh g–1 at 25°C, with an average CE of 99.998% at 10.5 A g–1. Remarkably, even under challenging high-load conditions (>10 mAh cm–2), the system maintains stable cycling for 500 cycles with 72% capacity retention at 5.1 mA cm–2 under 25°C. When paired with NCM positive electrodes, the full cell delivers stable cycling performance at 200 mA g–1 for 200 cycles with a high capacity retention of 80% and an average CE exceeding 99.95%. Combined with a pre-lithiation process, the full-cell achieves an improved ICE of 86.6% for Si-based SSBs. The universal applicability of the halide salt-driven modification strategy has been further validated, offering a promising pathway for developing high-reversibility, high-energy-density Si-based SSBs that meet practical application requirements.

Methods
Synthesis of materials
μSi (Aladdin, 99.9%, 1-3 μm), aluminum chloride (AlCl3, Macklin, 99%), iron chloride (FeCl3, Aladdin, 99.9%), tantalum chloride (TaCl5, Aladdin, 99.8%), bismuth chloride (BiCl3, Aladdin, 99.99%), niobium chloride (NbCl5, Aladdin, 99%), phosphorus pentachloride (PCl5, Macklin, 98%), magnesium chloride (MgCl2, Aladdin, 99%), antimony trichloride (SbCl3, Aladdin, 99%), zinc chloride (ZnCl2, Macklin, 99%), lithium oxide (Li2O, Aladdin, 99.99%) and Na2SiO3·9H2O (Aladdin, 98%) were purchased as raw materials. Commercial Li6PS5Cl (LPSC, Shenzhen Kejing Star Technology CO., LTD.) and Li3InCl6 (LIC, Shenzhen Kejing Star Technology CO., LTD.) were used as the SE layer. Vapor-grown carbon fiber (VGCF, Hefei Kejing, China) and commercial LiNi0.88Co0.09Mn0.03O2 (NCM) were used for the positive electrode.

All synthesis procedures were performed under an inert argon atmosphere in a glove box (H2O<0.01 ppm, O2<0.01 ppm) using simple thermal treatment.

For μSi@AlCl3, stoichiometric quantities of μSi and AlCl3 were mixed at mass ratios of 90:10, ground manually in an agate mortar for 5 min. The resultant mixture was transferred to a sealed glass bottle and subjected to thermal treatment at 180°C on a hot plate for 2 h, yielding the μSi@AlCl3-10. By adjusting the mass ratios to 95:5 and 85:15 while maintaining identical synthesis conditions, μSi@AlCl3-5 and μSi@AlCl3-15 were prepared, respectively.

For other metal chloride modifications (μSi@MClx (M=P, Ta, Bi, Fe, Nb, Zn, Mg)), the same procedure was followed with modified thermal treatment temperatures: 450°C for ZnCl2 and MgCl2, while temperatures were adjusted according to the respective chloride boiling points for others.

LiTaOCl4 was synthesized via solid-state mechanochemical reaction. A stoichiometric mixture of Li2O and TaCl5 (1 g total) was ball-milled in a ZrO2 pot (80 mL) with ZrO2 balls (30 g) at 500 rpm for 24 h.

Materials characterization
The electrode samples were prepared in a glove box under an argon atmosphere and transferred to the testing equipment via a sealed sample holder.

Phase structures of the as-prepared materials were investigated by XRD on a Bruker D8 diffractometer (equipped with Cu-Kα radiation) with a scan range of 10–90°. (SEM, Regulus 8230) Transmission electron microscopy (TEM, JEM-F200) measurements were employed to characterize the morphology, microstructure of samples and phase characteristics of the samples. EDS mappings were carried out to characterize the element distribution. XPS was recorded on Axis Supra+ to characterize the valence state and composition of each element in a sample. Gas Chromatography (GC, PE Clarus GC 690) was used to analyze the residual Li content of silicon-based negative electrodes. The specific steps include soaking the silica-based negative electrode in absolute ethanol for more than 12 h, during which the residual lithium in the silica-based negative electrode reacts with ethanol to form H2, and finally, the residual lithium content is calculated by detecting the H2 content5. This analytical approach exploits the distinct chemical reactivities of different lithium species: (1) kinetically trapped Li (K-Li) in lithium-silicon alloys reacts vigorously with ethanol to generate H2, while (2) SEI-consumed Li (C-Li) remains chemically inert under these conditions. The quantification protocol involves three key steps: first, the total irreversible Li (xir %) is determined from galvanostatic charge-discharge curves as (100%—ICE); second, the trapped Li fraction (xt %) is calculated from the H2 evolution; finally, the SEI-consumed Li proportion is derived: xs = (xir - xt) %.

The Raman analysis was performed with a Raman spectrometer (LabRAM Odyssey) with an excitation wavelength of 532 nm to detect the structural information of the sample. NDP was acquired using a cold neutron beam of the China Advanced Research Reactor32. The working principle is based on the nuclear reaction between 6Li (natural lithium consists of 6Li with a relative abundance of 7.5 at. %) and a neutron according to equation:

6Li + n ! 4He (2055 keV) + 3H (2727 keV) (1)

Depth calibration. The relationship between the detected energy (E) and the depth at which the charged particles are generated can be established using Stopping and Range of Ions in Matter (SRIM), from which the residual energy of the charged particles generated at different depths of the sample can be calculated. The relationship between the detected energy (E) and its depth (d) in the 1.5 μm LPSC + d μm Si-based negative electrode is determined to be:

d = 4.25719×10−7×E2 − 0.00518×E + 9.17825 (2)

All-solid-state ASSBs assembling
All electrodes were made by a traditional slurry-coating way where μSi/Si@AlCl3, poly-(vinylidene fluoride) (PVDF, 99%, Arkema) binder were ground at a mass ratio of 99.7:0.3, and then homogenized in 1-methyl-2-pyrrolidinone (NMP, 99%, Alfa Aesar). The slurry was then bladed on Cu foil (6 μm) to render a uniform coating, dried for at least 24 h at 60°C in a vacuum oven, and then punched into circular electrodes with a diameter of 10 mm. The electrode mass loading was around 0.8–1.2 mg cm–2. All SSBs were assembled in an argon-filled glove box using a mold with an internal diameter of 10 mm. For half-cell configuration, Li-In|LPSC|Si was made with an In foil (thickness: 50 μm; diameter: 10 mm) and a Li foil (thickness: 50 μm; diameter: 8 mm) as negative electrode (mass ratio of In foil and Li foil is 20), 80 mg LPSC as electrolyte layer, and Si electrode as positive electrode layer. The half-cell was pressed at a pressure of about 180 MPa for 5 min to ensure tight interface contact. Li-In|LPSC|Si@AlCl3 was the same as the above assembly process, only the Si electrode needed to be replaced with Si@AlCl3. The stack pressure applied during testing was about 180 MPa as well. For low external pressure testing, the half-cell assembly followed identical fabrication procedures as described above, with the sole modification of reducing the applied pressure to 20 MPa. For high-loading half cells (>9 mAh cm–2), the silicon electrode layer was directly prepared by cold pressing Si or Si@AlCl3 powder at 50 MPa. For full-cell, Si|LPSC|LIC|NCM was made with NCM-based positive electrode composite powder (NCM:LIC: VGCP=70:25:5) as the positive electrode layer, Si electrode as the negative electrode layer, 30 mg LIC as electrolyte layer on the positive side and 50 mg LPSC as electrolyte layer on the negative side. The full-cell was pressed at a pressure of about 400 MPa for 5 min to ensure tight interface contact. The stack pressure applied during testing was about 180 MPa as well. Si@AlCl3|LPSC|LIC|NCM was the same as the above assembly process, only the Si electrode was replaced with Si@AlCl3. The mass loading of the NCM positive electrode was about 15–20 mg cm–2. n/p is set between 1.1–1.15, where specific capacity of Si-based negative electrode is 3500 mAh g–1 and the specific capacity of NCM is 200 mAh g–1.

Electrochemical tests
LAND CT 2001. A battery testing system was used to conduct galvanostatic measurements of the batteries. Half-cell testing was performed in a voltage window of −0.6–1.4 V (0.02-2.02 V vs. Li+/Li) at 25°C, and full-cell testing was performed in a voltage window of 2.5–4.3 V at 25°C. The EIS was obtained by employing an electrochemical workstation (BioLogic 2) under a potentiostatic mode, with a signal amplitude of 5 mV. The measurements covered a frequency range of 7×106 to 100 Hz, with 10 data points per decade of frequency. GITT measurements to calculate diffusivity were performed at 0.1 C rate with the duration time of each current pulse of 15 min and the resting time of 2 h. DC polarization was performed to measure electronic conductivity by obtaining a stable current at different constant voltage values for 600 s. Cumulative CE is calculated as the product of the CE of each cycle from the first to the nth cycle (i.e., cumulative CE = CE1 × CE2 ×⋯ × CEn), where CEn of each cycle is the ratio of discharge capacity to charge capacity. It quantifies the total cumulative lithium loss and long-term reversibility.

Peer review rounds
[{"round": 1, "reviewer_comments": {"reviewer_1": "This work presents a remarkably innovative and timely contribution to the field of silicon-based all-solid-state batteries. The\nproposed surface halogenation strategy introduces a conceptually innovative paradigm in surface engineering. The authors\nhave identified a critical problem (irreversible lithium loss leading to poor Coulombic efficiency) and developed an ingenious\nsolution that transforms a drawback into an advantage - leveraging the native SiO₂ layer as a halogenation platform to\nconstruct a stable oxyhalide surface layer. Remarkably, this study provides fundamental insights\ninto the mechanism of irreversible lithium mitigation while demonstrating truly\noutstanding and record-setting electrochemical performance. The\nmanuscript is exceptionally well-constructed, the data is robust and the mechanistic investigations is profound.\nUndoubtedly, this work will generate significant interest in the silicon-based solid-state battery community and stimulate\nfurther research into halide-based surface engineering. Considering its scientific rigor, technical innovation, and potential\nimpact, I strongly recommend publication in Nature Communications.\nMinor Consideration (if any):\n1. The silicon-based solid-state half-cells developed in this study demonstrate remarkable electrochemical performance,\nparticularly in terms of cycling Coulombic efficiency. A systematic comparison with state-of-the-art reported values would\nfurther highlight the superior performance achieved through this innovative strategy.\n2. While the capacity contribution evaluation of AlCl3 and SA is generally well-performed, one methodological aspect\ndeserves further consideration. The current verification using pure AlCl3 and SA electrodes may not fully account for\nelectronic conductivity effects. Standard practice typically involves incorporating conductive carbon additives (~20 wt%) in\nsuch control experiments to better simulate actual electrochemical conditions. I recommend supplementing the current data\nwith additional charge-discharge profiles obtained from electrodes containing standard carbon additives.\n3. The cycling stability results under high mass loading conditions (>10 mAh cm-2) are truly impressive. I am particularly\ninterested to see the comparative rate performance.\n4. While not critical to the scientific conclusions, I would recommend using consistent terminology throughout the manuscript\nto enhance academic precision. Specifically:\n1) In Figure 5g, 'charged/discharged' could be unified with 'lithiated/delithiated' used in Figures 3d-g for consistency and\nmore precise description.\n2) The modification process description should consistently use either 'salt-mediated' or 'salt-driven' (currently appearing as\nboth in Fig. 6g and main text respectively).\n5. While the Supplementary Information is generally well-prepared, I noted a few areas where minor revisions would\nimprove consistency. For instance, the Si electrode preparation should be coated on “Cu foil” not “Al foil”. And the verb\ntenses used in the Supplementary Information should be made consistent with those in the main manuscript.\n6. The universal applicability demonstration of various halide salts for surface modification in Figure 6g is particularly\nsignificant and interesting. For a more comprehensive evaluation, the addition of cycling stability data would be highly\nbeneficial. Additionally, preparation methods for these universal modifications should be included in the Supplementary\nInformation for completeness.", "reviewer_2": "This manuscript presents a study on surface halogenation of silicon (Si) anodes for sulfide-based solid-state batteries\n(SSBs). The authors demonstrate that halogenation of the native SiO₂ surface, followed by substitution with AlOCl or AlCl₃\ncoatings, forms a protective interphase that mitigates side reactions with LPSCl and improves initial Coulombic efficiency\n(ICE) and cycling stability. Promising results are reported in both half-cell and full-cell configurations, including ICE values\nabove 90%.\nWhile technically compelling, the novelty of the strategy remains somewhat incremental compared to prior halide- or\noxyhalide-based interface engineering. Additional clarification on the originality, mechanistic insight, and broader\nimplications would strengthen the work. Several points also require clarification or refinement before the manuscript is\nsuitable for publication in Nature Communications.\n1. The manuscript would benefit from a clearer articulation of its novelty. Similar halide-based or oxyhalide strategies have\nbeen previously reported. Please elaborate on what differentiates the Al(Si)OCl surface layer from established approaches.\n2. Prior work (e.g., https://pubs.acs.org/doi/10.1021/acsmaterialslett.1c00055) has shown that LiAlCl₄ significantly enhances\ninterfacial stability in SSBs. How do those findings relate to the Al(Si)OCl phase reported here?\n3. Could the AlOCl layer behave similarly to oxychlorides like LiTaOCl₄ or LiNbOCl₄? Are there fundamental similarities or\ndifferences worth highlighting?\n4. AlCl₃ is selected based on its Lewis acidity, but other salts like TaCl₅ and FeCl₃ have similar or stronger acidity. A more\ndetailed justification for selecting AlCl₃ would improve the scientific rationale.\n5. Since the strategy relies on halogenation, have the authors explored other halogens (F, Br, I)? A broader screening would\nhelp establish generalizability.\n6. The ICE comparison includes both prelithiated and non-prelithiated Si systems. For meaningful benchmarking, please\neither restrict the data to one category or clearly distinguish the two in the figure and discussion.\n7. Figure 6g includes various SSB systems with different solid electrolytes. Restricting the comparison to sulfide-based\nsystems would better contextualize this work.\n8. The ICE difference between modified and unmodified Si anodes is ~6% in half-cells (94.3% vs. 88.4%), but the full-cell\ndifference is ~9%, and capacity loss increases from ~12% to ~25%. Please clarify whether this discrepancy arises from\ncathode limitations or interfacial mismatches.\n9. The loading density of the high-loading Si electrode in Figure 6c should be provided. As a reference, the authors are\nencouraged to include a more general cycling curve—for example, with a silicon loading density of ~4 mAh/cm² and a\ncycling rate of 0.3–0.5 C. How many cycles can such a cell achieve?\"\n10. The manuscript emphasizes achieving over 9 mAh/cm² areal capacity in the first cycle. However, this is not compelling\nwithout demonstrating stable cycling at high loading. Please provide extended cycling data under this condition.\n11. The full-cell shows ~82 mAh/g specific capacity with ~80% retention over 200 cycles. While stable, the capacity itself\nappears modest. Could the authors explain why it is relatively low?\n12. Figure 2j suggests that Cl is not concentrated at the surface despite the presence of a halogenated modification layer.\nPlease clarify the distribution profile and potential diffusion behavior of Cl.\n13. The applied stack pressure during operation is not specified. Please provide this detail. Also, is the cell performance\nsustained under low pressure? Prior studies suggest oxyhalides operate well under low pressure conditions.\n14. There is a typo in the supplementary materials where “M” is written instead of “Mn” for NCM. Please correct.\n15. Figure S30 shows that uSi@AlCl₃ appears denser than pristine uSi under the same pressure. Could this enhanced\ndensification affect ion transport? A comparison of mechanical properties (e.g., nanoindentation) would be helpful.\n16. In Figure 6e, please clarify how lithium fractions (reversible, trapped, SEI-consumed) were distinguished. A schematic or\nexplicit explanation would enhance clarity.\n17. The section titled “Electrochemical/chemical compatibility of Si-based anode with LPSCl” suffers from unclear figure\nreferencing (e.g., 4b, 4c, S20). Please ensure consistent and accessible figure references in both main text and supplement.\n18. In Figure 6a, the y-axis range is 95–100%. To allow fair comparison, the y-axis of Figure 6c should match this scale.\n19. In Figure 3d, “T-Li” appears to be a typo—should this be “K-Li”?\n20. In Figure S21d, what does the letter “g” represent? Please clarify.", "reviewer_3": "This work reports a halide salt-driven modification strategy to attain high Coulombic efficiency and stable cycle life. The\nfunctional Al(Si)OCl composite surface improves interfacial compatibility, enables fast ion/electron transfer, and reconciles\nirreversible Li loss. Some experimental evidence was also provided to support this assumption. Considering the overall\nquality and novelty of this work, the reviewer suggests a major revision before its publication.\n1. The XRD results show signals corresponding to AlCl3. Could the authors quantify its mass fraction, and discuss the\npotential impact of residual AlCl3 on anode performance?\n2. The distribution of Al and Cl seems uneven in the corresponding element mappings. Therefore, the EDS results may not\nsupport the homogeneous distribution of Al and Cl.\n3. The manuscript states that the mixture of AlCl3 and Si was ‘sealed and heated’ at 180 °C. Could the authors specify the\ntype of vessel used (e.g., autoclave, glass bottle) and the reaction atmosphere. In addition, the improvement in ICE of full\ncells relies on prelithiation. Please specify the method.\n4. Since LPSC contains Cl⁻ and the Al(Si)OCl layer also contains Cl⁻, is there any possibility of Cl⁻ migration between them\nduring cycling? The authors may need to consider how such migration could affect electrochemical performance.\n5. While the authors mention that the SA layer can suppress LPSC decomposition, further clarification and supporting\nevidence for the underlying mechanism are required.\n6. The manuscript reports an average Coulombic efficiency exceeding 102% for the\nhigh-loading Si|LPSC|Li-In cell, which requires further clarification and supporting evidence.\n7. The claim of “record-breaking” ICE in NCM full cells is misleading, as some cited comparisons are without prelithiation\nwhile the present work uses prelithiation. Please adjust the claim or benchmark consistently.\n8. The manuscript contains inconsistent terminology, as the compound is first introduced as “Li6PS5Cl (LPSC)” but later\nalso appears as “LiPSCl”; please unify the nomenclature (e.g., consistently use “LPSC”)\n9. Typos such as “Figue” instead of “Figure,” and missing spaces (e.g., “fullcell”) should be corrected. Please double check."}, "author_response": "Response to the Reviewers’ Comments (Manuscript ID: NCOMMS-25-39939)\nDear Reviewers,\nThank you for your comments concerning our manuscript entitled \"Surface\nHalogenation Engineering for Reversible Silicon-Based Solid-State Batteries\"\n(Manuscript ID: NCOMMS-25-39939). Those comments are constructive and very\nhelpful for further improving the quality of the manuscript. We have carefully and\nthoroughly revised the manuscript by following the suggestions. Below are our\nanswers and responses to your comments and questions about our manuscript.\nChanges made in the manuscript and supporting information are highlighted in red.\nThank you again for your time and guidance!\nSincerely,\nWei Xia\nReviewer #1 (Remarks to the Author):\nThis work presents a remarkably innovative and timely contribution to the field of\nsilicon-based all-solid-state batteries. The proposed surface halogenation strategy\nintroduces a conceptually innovative paradigm in surface engineering. The authors\nhave identified a critical problem (irreversible lithium loss leading to poor Coulombic\nefficiency) and developed an ingenious solution that transforms a drawback into an\nadvantage - leveraging the native SiO₂ layer as a halogenation platform to construct a\nstable oxyhalide surface layer. Remarkably, this study provides fundamental insights into the mechanism of irreversible lithium mitigation while demonstrating truly\noutstanding and record-setting electrochemical performance. The manuscript is\nexceptionally well-constructed, the data is robust and the mechanistic investigations is\nprofound.\nUndoubtedly, this work will generate significant interest in the silicon-based\nsolid-state battery community and stimulate further research into halide-based surface\nengineering. Considering its scientific rigor, technical innovation, and potential\nimpact, I strongly recommend publication in Nature Communications.\nMinor Consideration (if any):\nWe sincerely appreciate your constructive feedback and valuable suggestions. We\nhave addressed each of your comments point-by-point and provided detailed revisions\nto further improve the quality of our work.\n1. The silicon-based solid-state half-cells developed in this study demonstrate\nremarkable electrochemical performance, particularly in terms of cycling Coulombic\nefficiency. A systematic comparison with state-of-the-art reported values would\nfurther highlight the superior performance achieved through this innovative strategy.\nWe sincerely appreciate the reviewer's insightful suggestion. As systematically\nsummarized in Figure 6g of our manuscript comparing both areal capacity and initial\nCoulombic efficiency (ICE) between previously reported works and our study in\nhalf-cell configurations, the proposed halide modification demonstrates clear\nadvantages in overall electrochemical performance. Furthermore, Figure 6f presents\nthe ICE superiority in full-cell evaluations.\nTo specifically address the comment regarding cycling Coulombic efficiency (CE),\nwe have supplemented new data (Figure R1) showing cycling CE measurements in\nhalf-cells. Notably, while most reported works exhibit cycling CEs below 99.8%, our\nmodified electrode achieves an unprecedented 99.998% CE - marking the first\nreported system reaching \"four nines\" reliability and approaching \"five nines\"\nperformance class. This exceptionally high cycling CE fundamentally ensures\noutstanding anode cycling stability.\nFigure 6. (e) GCD profiles of Si@AlCl |LPSC|LIC|NCM after pre-lithiation at room temperature.\n3\n(f) ICE statistics of Si-based SSBs with NCM cathode. (g) Comprehensive performance statistics\nof Si@MCl with previously reported results of Si-based anodes in SSBs (M=Al, Ta, Bi, P, Fe, Nb,\nx\nSb; statistical information, including charge specific capacity , ICE and the references).\nanode anode\nFigure R1. Comparison of cycling Coulombic efficiency (CE) between our work and previously\nreported values in half-cell configurations. 1-7 correspond to the following references 1-6.\n2. While the capacity contribution evaluation of AlCl and SA is generally\n3\nwell-performed, one methodological aspect deserves further consideration. The\ncurrent verification using pure AlCl and SA electrodes may not fully account for\n3\nelectronic conductivity effects. Standard practice typically involves incorporating\nconductive carbon additives (~20 wt%) in such control experiments to better simulate\nactual electrochemical conditions. I recommend supplementing the current data with\nadditional charge-discharge profiles obtained from electrodes containing standard\ncarbon additives.\nWe sincerely appreciate the reviewer’s valuable suggestion. Following this\nrecommendation, we have systematically reevaluated the capacity contribution by\nintroducing 20 wt% Super P into both AlCl and SA matrices (Figure R2). The results\n3\ndemonstrate that the capacity contribution from Super P is negligible compared to\nthat of the silicon anode itself.\nFigure R2. Galvanostatic charge/discharge (GCD) profiles of (a) SA+20 wt% Super\nP|LPSC|Li-In and (b) AlCl +20 wt% Super P|LPSC|Li-In. Cyclic Voltammetry (CV) curves of (c)\n3\n3. The cycling stability results under high mass loading conditions (~10 mAh cm-2)\nare truly impressive. I am particularly interested to see the comparative rate\nperformance.\nWe sincerely appreciate the reviewer's valuable suggestion. In response, we have\nconducted additional rate capability tests under extreme high-loading conditions\n(without binder or conductive carbon additives, Figure R3). Under these challenging\nconditions, while the absolute rate performance of μSi@AlCl /Li-In is inevitably\n3\nlimited due to the extremely high mass loading and absence of conductive/binder\ncomponents, it still demonstrates significantly enhanced rate capability compared to\nthe bare μSi/Li-In half-cell system.\nFigure R3. Rate performance of high loading μSi@AlCl /Li-In and high loading μSi/Li-In.\n3\n4. While not critical to the scientific conclusions, I would recommend using consistent\nterminology throughout the manuscript to enhance academic precision. Specifically:\n1) In Figure 5g, 'charged/discharged' could be unified with 'lithiated/delithiated' used\nin Figures 3d-g for consistency and more precise description.\n2) The modification process description should consistently use either 'salt-mediated'\nor 'salt-driven' (currently appearing as both in Fig. 6g and main text respectively).\nWe are grateful for the reviewer's attention to terminology consistency. In our\nrevisions, we have systematically implemented 'lithiated/delithiated' (Figure 5g) and\n'salt-driven' (Figure 6g) throughout the text, ensuring complete standardization of\nthese key terms in all relevant discussions.\nFigure 5. The R statistical results and DRT analysis for (g) Si@AlCl |LPSC|Li-In and\nct 3\nSi|LPSC|Li-In in different electrochemical states.\nFigure 6. (g) Comprehensive performance statistics of Si@MCl with previously reported results\nx\nof Si-based anodes in SSBs (M=Al, Ta, Bi, P, Fe, Nb, Sb; statistical information, including charge\nspecific capacity , ICE and the references).\nanode anode\n5. While the Supplementary Information is generally well-prepared, I noted a few\nareas where minor revisions would improve consistency. For instance, the Si electrode\npreparation should be coated on “Cu foil” not “Al foil”. And the verb tenses used in\nthe Supplementary Information should be made consistent with those in the main\nmanuscript.\nWe sincerely apologize for this oversight. The verb tenses have now been carefully\nrevised and standardized throughout the manuscript to ensure improved consistency\nand accuracy.\n6. The universal applicability demonstration of various halide salts for surface\nmodification in Figure 6g is particularly significant and interesting. For a more\ncomprehensive evaluation, the addition of cycling stability data would be highly\nbeneficial. Additionally, preparation methods for these universal modifications should\nbe included in the Supplementary Information for completeness.\nIn response to the reviewer's constructive comment, we have conducted additional\nfull-cell cycling tests pairing chloride-modified Si-based anodes with NCM cathodes.\nAs shown in Figure R4, the modified anodes consistently demonstrate stable cycling\nperformance, maintaining >80% capacity retention over 200 cycles.\nFigure R4. Cycle performance of μSi@MCl /NCM (M=Zn, Mg, P, Ta, Bi).\nx\nReviewer #2 (Remarks to the Author):\nThis manuscript presents a study on surface halogenation of silicon (Si) anodes for\nsulfide-based solid-state batteries (SSBs). The authors demonstrate that halogenation\nof the native SiO₂ surface, followed by substitution with AlOCl or AlCl₃ coatings,\nforms a protective interphase that mitigates side reactions with LPSCl and improves\ninitial Coulombic efficiency (ICE) and cycling stability. Promising results are reported\nin both half-cell and full-cell configurations, including ICE values above 90%.\nWhile technically compelling, the novelty of the strategy remains somewhat\nincremental compared to prior halide- or oxyhalide-based interface engineering.\nAdditional clarification on the originality, mechanistic insight, and broader\nimplications would strengthen the work. Several points also require clarification or\nrefinement before the manuscript is suitable for publication in Nature\nCommunications.\nWe sincerely appreciate the reviewer's thoughtful evaluation and constructive\nsuggestions regarding our work. In addressing the question of novelty, our\ncomprehensive analysis of high-impact studies on chloride- and oxyhalide-based\nmaterials in solid-state batteries (summarized in Table R1) reveals that research\nefforts predominantly focus on solid electrolyte design for high-voltage cathode\ncompatibility, with only limited exploration of lithium metal anode applications. More\nimportantly, direct utilization of these halides for interface engineering - particularly\nin silicon-based solid-state batteries - remains scarce, thereby establishing the\ngenuine novelty of our approach. While previous studies on halide materials provided\nconceptual inspiration, our interfacial modification strategy represents a significant\ndeparture from conventional applications.\nIt is also worth noting that most reported halide electrolytes rely on electrode\nblending (e.g., 70 wt% active particles + 30 wt% solid electrolyte), making uniform\nelectrode coating challenging. The high proportion of non-redox-active electrolytes\nwithin the electrode results in lower energy density for the overall system. In contrast,\nour proposed method leverages the sublimation of precursor salts to achieve ultrathin\nand uniform electrode coating. Subsequent lithium embedding facilitates potential\nconversion into an ion-conductive layer, significantly reducing interfacial ion transfer\nimpedance. This process reduces halide usage, thereby effectively boosting both\nanode capacity and overall energy density (Fig. 6g).\nMoreover, our halide interfacial engineering strategy breaks new ground by\ntransforming the typically problematic native SiO layer into an asset for in situ\nx\nmodification. This paradigm shift enables remarkable improvements in\nelectrochemical reversibility.\nIn direct response to the reviewer's valuable suggestions, we have conducted\ncomprehensive theoretical calculations, additional characterization and extended\nelectrochemical testing. We believe these supplemental investigations will help better\ndemonstrate the originality of our work, provide deeper mechanistic insights, and\nhighlight its broader implications for the field.\nTable R1. Summary of high-impact studies on halide/oxyhalide materials in\nsolid-state batteries.\nNumber Halide- or oxyhalide Field Reference\n2 x 0.666−x 4\ncathode\n0.388 0.238 0.475 3\nmetal\n3 6\ncathode advances9\n2 6\ncathode Communications10\n5 xLi O-MCly SSE for Nature\n2\ncathode Communications11\n1.3 1.2 4\ncathode\nFigure 6. (g) Comprehensive performance statistics of Si@MCl with previously reported results\nx\nof Si-based anodes in SSBs (M=Al, Ta, Bi, P, Fe, Nb, Sb; statistical information, including charge\nspecific capacity , ICE and the references).\nanode anode\n1. The manuscript would benefit from a clearer articulation of its novelty. Similar\nhalide-based or oxyhalide strategies have been previously reported. Please elaborate\non what differentiates the Al(Si)OCl surface layer from established approaches.\nIn comparison with existing reports on halide/oxyhalide applications in solid-state\nbatteries, our work demonstrates three key distinctive features.\nFirst, the preparation strategy of our modification layer is novel-we ingeniously\nutilize the intrinsic (yet typically detrimental) oxide layer on silicon particles as\nreactive sites, achieving a unique \"waste-to-wealth\" modification effect, which differs\nfundamentally from previous approaches in either solid electrolyte design or\ninterfacial engineering.\nSecond, the modified layer itself is distinct from conventional ionic conductors, as\nit is lithium-free by design; instead, its primary function is to enhance interfacial\ncompatibility, representing a different design direction.\nThird, most reported halide electrolytes rely on electrode blending (e.g., 70 wt%\nactive particles + 30 wt% solid electrolyte), making uniform electrode coating\nchallenging. The high proportion of non-redox-active electrolytes within the electrode\nresults in lower energy density for the overall system. In contrast, our proposed\nmethod leverages the sublimation of precursor salts to achieve ultrathin and uniform\nelectrode coating. Subsequent lithium embedding facilitates potential conversion into\nan ion-conductive layer, significantly reducing interfacial ion transfer impedance. This process reduces halide usage, thereby effectively boosting both\nanode capacity and overall energy density (Fig. 6g).\n2. Prior work (e.g., https://pubs.acs.org/doi/10.1021/acsmaterialslett.1c00055) has\nshown that LiAlCl₄ significantly enhances interfacial stability in SSBs. How do those\nfindings relate to the Al(Si)OCl phase reported here?\nWe sincerely appreciate the provided information. The reported work on LiAlCl₄\nand other halide-based superionic conductors-particularly their role in enhancing\ninterfacial stability-has indeed inspired our selection of AlCl for modifying\n3\nsilicon-based anodes in solid-state batteries. While the initially formed Al(Si)OCl\nlayer is lithium-free, we acknowledge that trace amounts of LiAlCl(O) may form\nduring cycling, whose inherent superionic properties could beneficially facilitate\ninterfacial ion transport. Furthermore, the documented low-temperature (150°C)\nreaction between AlCl and LiCl to form LiAlCl provided valuable reference for our\n4 4\nlow-temperature modification approach at 180°C. In response to this constructive\nsuggestion, we will incorporate citations to this study and related works to further\nstrengthen the scientific rigor of our manuscript.\n3. Could the AlOCl layer behave similarly to oxychlorides like LiTaOCl₄ or LiNbOCl₄?\nAre there fundamental similarities or differences worth highlighting?\nThrough comprehensive analysis, we have identified at least two fundamental\ndifferences between the behavior of AlOCl in silicon-based solid-state battery systems\ncompared to LiTaOCl or LiNbOCl . First, regarding functionality, while\n4 4\nLiTaOCl /LiNbOCl serve as fast-ion conductors in solid electrolytes that enable\n4 4\nrelatively stable high-voltage cathode interfaces, they exhibit poor reduction stability\nat low-voltage anode sides11,13. In contrast, AlOCl, being lithium-free and not a\nfast-ion conductor, functions as an interfacial modification layer to enhance the\n(electro)chemical compatibility between silicon anodes and LPSC. Second,\nconcerning electrochemical compatibility with LPSC, AlOCl demonstrates superior\nperformance compared to LiTaOCl /LiNbOCl . The higher oxidation states of Ta5+\n4 4\nand Nb5+, coupled with their empty E orbitals that tend to accept electrons and\ng\nundergo reduction14, result in poorer compatibility with LPSC. Conversely, Al3+ with\nits lower oxidation state exhibits better stability. To substantiate these observations,\nwe conducted complementary theoretical calculations and validation experiments. We\nsuccessfully synthesized LiTaOCl (LTOC) (Fig. R5 and Fig. R6) and constructed\ntime, confirming the inferior compatibility (Fig. R7). Additional impedance\nmeasurements comparing delithiated-μSi@TaCl /Li-In and\n5\ndelithiated-μSi@NbCl /Li-In with delithiated-μSi@AlCl /Li-In further supported these\n5 3\nfindings, showing higher impedance for the Ta/Nb-based systems (Fig. R8).\nThrough systematic theoretical calculations, we comprehensively evaluated the\nreaction energies between LPSC and various compounds including AlOCl, its\npotential cycling-derived product LiAlOCl, as well as LiTaOCl and LiNbOCl (Fig.\n4 4\nR9). The computational results reveal that LiTaOCl and LiNbOCl exhibit\n4 4\nsignificantly more negative reaction energies with LPSC, indicating both higher\nreactivity and inferior chemical stability at the interface. In marked contrast, AlOCl\nand its possible electrochemical conversion product LiAlOCl demonstrate\nsubstantially more favorable thermodynamic stability, as evidenced by their relatively\npositive reaction energies with LPSC. These computational findings provide\nfundamental insights into the superior interfacial compatibility observed\nexperimentally for AlOCl-modified systems compared to their Ta/Nb-based\ncounterparts.\nFigure R5. XRD spectrum of LiTaOCl .\nevolution of electrochemical impedance.\nFigure R8. EIS results of delithiated-μSi@TaCl /Li-In, delithiated-μSi@NbCl /Li-In and\n5 5\ndelithiated-μSi@AlCl /Li-In.\nand AlOCl, LiAlOCl, LiTaOCl and LiNbOCl 15.\n4 4\n4. AlCl₃ is selected based on its Lewis acidity, but other salts like TaCl₅ and FeCl₃\nhave similar or stronger acidity. A more detailed justification for selecting AlCl₃\nwould improve the scientific rationale.\nWhile it is true that AlCl shares Lewis acid characteristics with other chlorides\n3\nsuch as TaCl and FeCl , our selection of AlCl as the primary focus for this study\n5 3 3\nwas based on two critical considerations. Firstly, from the perspective of reaction\nfeasibility, we specifically leveraged AlCl 's Lewis acidity to react with SiO , a\n3 2\nprocess that proceeds spontaneously at relatively low temperatures. In contrast,\npreliminary thermodynamic calculations of Gibbs free energy revealed that the\nformation of oxychlorides from TaCl or FeCl is not spontaneous at room\n5 3\ntemperature (as detailed in Figures R10, and Table R2). These alternative halides\ngenerally require elevated reaction temperatures while demonstrating inferior\nreaction efficiency. Secondly, comprehensive electrochemical testing of various\nchloride-modified systems consistently showed that AlCl treatment yields\n3\nsilicon-based anodes with superior initial Coulombic efficiency compared to other\nchloride alternatives (Table R3). This combination of favorable reaction\nthermodynamics under mild conditions and demonstrably better electrochemical\nperformance establishes AlCl as the optimal choice for our interfacial modification\n3\nstrategy.\nFigure R10. Theoretical Gibbs free energy analysis (simplified, accounting for unadjusted heat\ncapacity variations and other approximations) for reactions between various chlorides and SiO ,\n2\nshowing critical temperatures for spontaneous reaction. Note: thermodynamic parameters were\nobtained from the NIST database and HSC Chemistry 9.0. A more negative ΔG indicates stronger\nreaction spontaneity (ΔG < 0), while ΔG > 0 denotes non-spontaneous reactions.\nTable R2. Reaction equations between various metal chlorides and SiO .\ntemperature (K) (kJ/mol)\nAlCl AlCl +SiO →AlOCl+SiCl <453 -95.8\n3 3 2 4\nTaCl TaCl +SiO →TaOCl +SiCl 1050 +85.7\n5 5 2 3 4\nFeCl FeCl +SiO →Fe O +SiCl 605 +5.2\n5 3 2 2 3 4\nZnCl ZnCl +SiO →ZnO+SiCl >1200 +142.3\n2 2 2 4\nBiCl BiCl +SiO →BiOCl+SiCl <453 -22.4\n3 3 2 4\nSbCl SbCl +SiO →SbOCl+SiCl 620 +4.1\n3 3 2 4\nNbCl NbCl +SiO →NbOCl+SiCl +Cl <453 -15.3\n5 5 2 4 2\nPCl PCl +SiO →POCl +SiCl <453 -38.5\n5 5 2 3 4\nSnCl SnCl +SiO →SnOCl +SiCl 980 +35.2\n4 4 2 2 4\nCrCl CrCl +SiO →Cr O +SiCl >1200 +115.7\n3 3 2 2 3 4\nMgCl MgCl +SiO →MgO+SiCl >1200 +152.7\n2 2 2 4\nTiCl TiCl +SiO →TiOCl +SiCl 1050 +25.2\n4 4 2 2 4\nTable R3. The performance comparisons of Si@MCl anode (M=Al, Ta, Bi, P, Fe, Nb, Sb, et al.).\nNumber Composite ICE (%)\nanode\n5\n3\n3\n5\n3\n3\n3\n5. Since the strategy relies on halogenation, have the authors explored other halogens\n(F, Br, I)? A broader screening would help establish generalizability.\nWe appreciate the suggestion regarding alternative halide exploration. Our\nselection of AlCl was primarily based on comprehensive considerations of both\n3\nthermodynamic stability and practical feasibility (Table R4). Firstly, from a stability\nperspective, AlCl exhibits optimal physical properties with moderate melting/boiling\n3\npoints compared to other aluminum halides - AlBr and AlI demonstrate severe\n3 3\nhygroscopicity and poor stability unsuitable for practical applications, while AlF\n3\npossesses excessively high melting/boiling points that hinder reaction kinetics.\nSecondly, the theoretical framework supports that AlCl can thermodynamically favor\n3\nthe formation of stable AlOCl through reaction with SiO , whereas hypothetical SiOBr,\nfurther validate this rationale, we experimentally tested AlBr modification (Figure\n3\ninitial Coulombic efficiency at room temperature, showing negligible improvement\nover pristine μSi, which we attribute to the inherent instability of bromide species as\npredicted. These combined theoretical and experimental findings systematically justify\nour focused investigation on AlCl -based modification.\n3 3 3 3\nMelting poin (˚C) boiling point (˚C) reactivity\nAlF 1040 1272 stable\nFigure R12. GCD profiles of μSi@AlBr /Li-In at 0.1C.\n3\n6. The ICE comparison includes both prelithiated and non-prelithiated Si systems. For\nmeaningful benchmarking, please either restrict the data to one category or clearly\ndistinguish the two in the figure and discussion.\nWe appreciate the reviewer's constructive suggestion. In response, we have\nclearly labeled all relevant systems in the figures to explicitly distinguish between\npre-lithiated and non-prelithiated silicon-based systems (Figure R13), ensuring\nunambiguous interpretation of the comparative data.\nFigure R13. (e) GCD profiles of Si@AlCl |LPSC|LIC|NCM after pre-lithiation at room\n3\ntemperature. (f) ICE statistics of Si-based SSBs with NCM cathode (see Figure 6).\n7. Figure 6g includes various SSB systems with different solid electrolytes.\nRestricting the comparison to sulfide-based systems would better contextualize this\nwork.\nWe sincerely appreciate the reviewer's valuable suggestion. In response, we have\nrevised Figure 6 to focus exclusively on comparisons within sulfide-based systems, as\nshown in the updated Figure 6g (Figure R14).\nFigure 6g (Figure R14). Comprehensive performance statistics of Si@MCl with previously\nx\nreported results of Si-based anodes in SSBs (M=Al, Ta, Bi, P, Fe, Nb, Sb; statistical information,\nincluding charge specific capacity , ICE and the references).\nanode anode\n8. The ICE difference between modified and unmodified Si anodes is ~6% in\nhalf-cells (94.3% vs. 88.4%), but the full-cell difference is ~9%, and capacity loss\nincreases from ~12% to ~25%. Please clarify whether this discrepancy arises from\ncathode limitations or interfacial mismatches.\nWe sincerely appreciate the reviewer's insightful comment. The observed\ndiscrepancy between half-cell and full-cell performance can be attributed to several\nfactors, particularly the difference in areal loading (~2 mAh cm-2 for half-cells vs\nsignificantly higher loading in full-cells) and inherent testing variability. As the\nelectrochemical kinetics become increasingly constrained at higher loadings,\nμSi@AlCl demonstrates relatively superior kinetic tolerance compared to pristine μSi,\n3\nexplaining the pronounced performance differences in full-cell configurations. To\nensure statistical reliability and minimize experimental variability, we standardized\nthe full-cell areal loading and conducted triplicate tests under identical conditions\n(Figure R15). The averaged results show consistent improvements: μSi-NCM exhibits\nan average initial Coulombic efficiency of 79.2%, while μSi@AlCl -NCM achieves\n3\n85.2%, representing a 6.0% enhancement that aligns remarkably well with the 5.9%\nimprovement observed in half-cell tests. Furthermore, the revised Figure 3b is shown\nas follows.\nFigure R15. GCD profiles of (a) μSi-NCM, (b) μSi@AlCl -NCM, and (c) statistical data of ICE.\n3\nmeans Li InCl ); (c) Comparison results of the first CE and the charge specific capacity (CSC) of\n3 6\nthe Si active material corresponding to the Si and Si@AlCl anodes.\n3\n9. The loading density of the high-loading Si electrode in Figure 6c should be\nprovided. As a reference, the authors are encouraged to include a more general\ncycling curve—for example, with a silicon loading density of ~4 mAh/cm² and a\ncycling rate of 0.3–0.5 C. How many cycles can such a cell achieve?\"\nWe sincerely appreciate your constructive suggestion regarding the need for\nadditional performance data. In response to your valuable feedback, we have\nsupplemented Figure 6c with detailed silicon loading information and further\nenhanced our electrochemical characterization by including cycling performance\ndata for high-loading silicon-based anodes (5.96 mAh cm-2) (Figure R16). The\ncomprehensive evaluation demonstrates that these high-capacity electrodes can\ndeliver remarkable areal capacities of approximately 2.5 mAh cm-2 when cycled at\n0.3C under ambient temperature conditions. Notably, the system exhibits exceptional\ncycling stability, maintaining 90% of its initial capacity after 200 cycles, which\nsubstantiates the robustness of our electrode design and its potential for practical\napplications.\nFigure R16. GCD profiles of (a) μSi@AlCl /Li-In with a high silicon loading density of\n3\n~6mAh/cm² and (b) cycle performance at 0.3C and room temperature.\n10. The manuscript emphasizes achieving over 9 mAh/cm² areal capacity in the first\ncycle. However, this is not compelling without demonstrating stable cycling at high\nloading. Please provide extended cycling data under this condition.\nWe sincerely appreciate the reviewer's insightful comment regarding the cycling\nstability under high-loading conditions. In response, we would like to clarify that\nFigure 6c in our manuscript already presents the cycling performance under these\nchallenging conditions (>9 mAh/cm2 without binder or conductive additives). While\nthe high current density of 5.1 mA/cm2 (corresponding to ~0.5C rate) does lead to\nrelatively lower capacity utilization during cycling, the system demonstrates excellent\nstability with 72% capacity retention over 500 cycles-a significant achievement for\nsuch demanding conditions\nTo provide more comprehensive evidence, we have supplemented extended cycling\ndata at a reduced current density of 1.5 mA/cm2 (~0.15C). Importantly, at this lower\ncurrent density, the system achieves substantially improved capacity utilization (5-6\nmAh/cm²) while maintaining good cycling stability over 50 cycles.\nFigue 6. Comparison of (a) cycling performance of Si@AlCl |LPSC|Li-In and Si|LPSC|Li-In at\n3\nroom temperature and 3C rate, and (b) the corresponding average CE and Cumulative CE.\nComparison of (c) cycling performance of high-loading Si@AlCl |LPSC|Li-In and high-loading\naverage CE and Cumulative CE.\nFigure R17. GCD profiles of (a) μSi@AlCl /Li-In with a high silicon loading density (>9\n3\nmAh/cm²), (b) GCD profiles at different cycles and (c) cycle performance at 1.5 mA cm-2 (0.15C).\n11. The full-cell shows ~82 mAh/g specific capacity with ~80% retention over 200\ncycles. While stable, the capacity itself appears modest. Could the authors explain\nwhy it is relatively low?\nThe relatively lower capacity utilization observed in full-cell configurations\nprimarily stems from rate capability limitations (Figure S28). As corroborated by\nYing Shirley Meng's team16, silicon-based anodes actually exhibit favorable kinetics,\nevidenced by their higher limiting current density in Li-Si symmetric cells compared\nto Li-In counterparts. Consequently, the rate performance bottleneck in our full-cells\npredominantly originates from the intrinsic kinetic constraints of the NCM cathode\nmaterial. To further elucidate this aspect, we have conducted additional\nelectrochemical evaluations at elevated temperature (50°C), which demonstrate\nsignificantly improved capacity delivery (>100 mAh g-1 at 1C rate) (Figure R18).\nNotably, even at these higher capacity utilizations, the system maintains excellent\ncycling stability, retaining 82.8% of its initial capacity after 200 cycles.\nFigure S28. Rate performance of Si@AlCl |LPSC|LIC|NCM and Si|LPSC|LIC|NCM under room\n3\ntemperature (LIC means Li InCl ).\n3 6\nFigure R18. Rate and cycle performance of μSi@AlCl /NCM.\n3\n12. Figure 2j suggests that Cl is not concentrated at the surface despite the presence of\na halogenated modification layer. Please clarify the distribution profile and potential\ndiffusion behavior of Cl.\nWe have supplemented our analysis with TEM/EDS elemental mapping, including\nline-scan profiles, which reveal that Cl, O, and Al exhibit similar surface-enriched\ndistributions on silicon particles-a trend consistently confirmed by both mapping and\nline-scan results (Figure R19). The primary distinction lies in the relatively lower\nconcentration of Cl compared to other elements, which likely contributes to the less\npronounced surface accumulation observed in Figure 2j. This reduced Cl signal\nintensity may originate from several technical factors: (1) the inherent volatility of Cl\nunder high-energy electron beam irradiation and (2) potential sample preparation\nartifacts.\nWe acknowledge that current understanding of chlorine ion behavior in silicon\ncrystals remains predominantly focused on surface phenomena, particularly\nconcerning the surface SiO layer or adsorbed species on Si(100)17,18, while direct\n2\nevidence for bulk phase diffusion remains limited. Our elemental line-scanning\nanalysis reveals minimal chlorine content within the silicon bulk phase, supporting\nthe conclusion that chlorine primarily interacts with the surface oxide layer and tends\nto accumulate preferentially at surface regions rather than penetrating significantly\ninto the crystalline lattice.\nFigure R19. (a) EDS elemental mapping results and (b) line-scan profiles of μSi@AlCl .\n3\n13. The applied stack pressure during operation is not specified. Please provide this\ndetail. Also, is the cell performance sustained under low pressure? Prior studies\nsuggest oxyhalides operate well under low pressure conditions.\nWe appreciate the reviewer's insightful comments. The detailed parameters of\napplied pressure have been provided in the Supporting Information for clarity. We\nwould like to clarify that our initial tests were conducted under relatively high\nexternal pressures (180 MPa for half-cells and 400 MPa for full-cells), rather than\nthe low-pressure conditions mentioned. In direct response to the reviewer's valuable\nsuggestion, we have conducted additional electrochemical evaluations under reduced\nexternal pressure (20 MPa) to systematically investigate its influence on cycling\nstability (Figure R20). The results demonstrate a capacity retention of 84.3% after 50\ncycles under these low-pressure conditions, which indeed confirms that reduced\nstacking pressure negatively impacts the cycling performance to a certain extent.\nFigure R20. Cycle performance of μSi@AlCl /Li-In at low pressure of 20 MPa.\n3\n14. There is a typo in the supplementary materials where “M” is written instead of\n“Mn” for NCM. Please correct.\nWe sincerely appreciate the reviewer's careful reading and valuable comment. In\nresponse, we have corrected the indicated typographical error in the revised\nmanuscript and have conducted a comprehensive proofreading of the entire document\nto identify and eliminate any additional typographical or grammatical errors, thereby\nensuring the highest quality of presentation in our work.\n15. Figure S30 shows that uSi@AlCl₃ appears denser than pristine uSi under the same\npressure. Could this enhanced densification affect ion transport? A comparison of\nmechanical properties (e.g., nanoindentation) would be helpful.\nIndeed, a more compacted state theoretically facilitates improved ionic transport by\nenhancing interfacial contact between solid phases. Following your recommendation,\nwe have conducted additional qualitative mechanical assessments using\nsingle-particle nanoindentation (Figure R21). The results demonstrate that under\ncomparable particle sizes, the μSi@AlCl exhibits a shallower slope in the\n3\nforce-displacement curve compared to untreated μSi, indicating relatively lower\nmechanical hardness. This reduced hardness characteristic is advantageous for\nachieving superior compaction states, as softer particles can deform more readily\nunder pressure to establish more intimate interfacial contacts.\nFigure R21. Force-displacement curves from single-particle nanoindentation tests comparing\nμSi and μSi@AlCl .\n3\n16. In Figure 6e, please clarify how lithium fractions (reversible, trapped,\nSEI-consumed) were distinguished. A schematic or explicit explanation would\nenhance clarity.\nWe sincerely appreciate the reviewer's valuable suggestion regarding the\nquantification of Li distribution. The classification of lithium into reversible, trapped,\nand SEI-consumed fractions in our study follows the well-established gas-phase\nquantitative methodology16 (Science 373, 1494-1499 (2021)). This analytical\napproach exploits the distinct chemical reactivities of different lithium species: (1)\ntrapped Li in lithium-silicon alloys reacts vigorously with ethanol to generate H₂,\nwhile (2) SEI-consumed Li remains chemically inert under these conditions. The\nquantification protocol involves three key steps: first, the total irreversible Li (x %) is\nir\ndetermined from galvanostatic charge-discharge curves as (100 - initial Coulombic\nefficiency); second, the trapped Li fraction (x%) is calculated from the H₂ evolution\nt\nmeasured by Gas Chromatography (GC); finally, the SEI-consumed Li proportion is\nderived: x =(x -x)%.\ns ir t\nFigure 3. (e) Schematic diagram of GC sample preparation, the test results and the distribution of\nirreversible Li.\n17. The section titled “Electrochemical/chemical compatibility of Si-based anode with\nLPSCl” suffers from unclear figure referencing (e.g., 4b, 4c, S20). Please ensure\nconsistent and accessible figure references in both main text and supplement.\nWe appreciate the reviewer's valuable comment regarding the unclear figure\nreferencing in our manuscript. In response, we have systematically addressed this\nissue by carefully revising all figure citations to ensure they precisely correspond to\nthe intended graphical data.\n18. In Figure 6a, the y-axis range is 95–100%. To allow fair comparison, the y-axis of\nFigure 6c should match this scale.\nWe sincerely appreciate the reviewer's insightful observation regarding the y-axis\nscales in Figures 6a and 6c. The variation in y-axis ranges was intentionally\nimplemented to optimally present the distinct Coulombic efficiency fluctuation\ncharacteristics of each system while maintaining graphical clarity. Importantly, we\nmaintained identical testing conditions and direct side-by-side comparison for both\nμSi and μSi@AlCl within each individual figure panel, ensuring the fairness and\n3\nvalidity of our comparative analysis. This visualization strategy successfully balances\naesthetic considerations with rigorous scientific representation.\nFigue 6. Comparison of (a) cycling performance of Si@AlCl |LPSC|Li-In and Si|LPSC|Li-In at\n3\nroom temperature and 3C rate, and (b) the corresponding average CE and Cumulative CE.\nComparison of (c) cycling performance of high-loading Si@AlCl |LPSC|Li-In and high-loading\naverage CE and Cumulative CE.\n19. In Figure 3d, “T-Li” appears to be a typo—should this be “K-Li”?\nWe sincerely appreciate the reviewer's careful reading and valuable correction\nregarding the terminology. In response to this comment, we have revised the\nmanuscript accordingly, replacing all instances of \"T-Li\" with the correct designation\n\"K-Li\" throughout the text, figures, and supplementary materials to ensure full\nconsistency.\n20. In Figure S21d, what does the letter “g” represent? Please clarify.\nWe sincerely appreciate the reviewer's meticulous attention to detail in identifying\nthis oversight. In response, we have carefully corrected the indicated inaccuracy in\nour revised manuscript.\nReviewer #3 (Remarks to the Author):\nThis work reports a halide salt-driven modification strategy to attain high Coulombic\nefficiency and stable cycle life. The functional Al(Si)OCl composite surface improves\ninterfacial compatibility, enables fast ion/electron transfer, and reconciles\nirreversible Li loss. Some experimental evidence was also provided to support this\nassumption. Considering the overall quality and novelty of this work, the reviewer suggests a\nmajor revision before its publication.\nWe sincerely appreciate your positive evaluation of our work and your valuable\nsuggestions for improvement. We are grateful for your recognition of the novelty and\npotential impact of our halide salt-driven modification strategy. In response to your\nexpert comments, we have systematically addressed each point raised through\ncomprehensive revisions to further enhance the research depth and overall quality of\nthe manuscript.\n1. The XRD results show signals corresponding to AlCl3. Could the authors quantify\nits mass fraction, and discuss the potential impact of residual AlCl3 on anode\nperformance?\nWe sincerely appreciate your insightful suggestion regarding the potential influence\nof residual AlCl in our system. In response to this important consideration, we have\n3\nconducted thorough experimental investigations to quantitatively evaluate this aspect.\nThermogravimetric-mass spectrometric (TG-MS) analysis under N₂ atmosphere\n(accounting for AlCl 's boiling point of 178°C) confirms an extremely low residual\n3 3\nperformance (Figure R22a). To further substantiate this conclusion, we performed\ncontrol experiments by physically mixing μSi with AlCl , which revealed that the\n3\npresence of AlCl actually deteriorates the electrochemical performance -\n3\ndemonstrating decreased initial Coulombic efficiencies from 86.8% to 85.2% with\nLCO cathodes and from 80.4% to 74.5% with NCM (Ni88) cathodes, respectively\n(Figure R22b,c).\nFigure R22. (a) TG-MS analysis of μSi@AlCl under N₂ atmosphere; (b) Comparative GCD\n3\nprofiles of μSi/LCO versus μSi-mixture of AlCl /LCO and (c) μSi/Ni88 versus μSi-mixture of\n3\n2. The distribution of Al and Cl seems uneven in the corresponding element mappings.\nTherefore, the EDS results may not support the homogeneous distribution of Al and\nCl.\nWe sincerely appreciate your insightful suggestion. Following your valuable\nrecommendation, we have conducted additional EDS analyses and carefully revised\nour descriptions to more accurately represent the experimental observations (Figure\nR23). The updated elemental mapping results demonstrate reasonably consistent\ndispersion patterns of Si, Al, and Cl elements throughout both μSi@AlCl particles\n3\nand electrode materials, while appropriately acknowledging the inherent variations\nexpected in such composite systems. We have modified our manuscript text\naccordingly as below, and we believe these revisions have significantly strengthened\nthe scientific rigor of our presentation:\nFigure R23. EDS elemental mapping results of μSi@AlCl .\nSEM imaging coupled with EDS elemental mapping confirms the presence and characteristic\ndistribution of Si, Al, and Cl elements throughout both Si@AlCl₃ particles (Figures S6) and the\ncorresponding electrodes (Figures S7-S8), showing relatively consistent elemental dispersion.\n3. The manuscript states that the mixture of AlCl3 and Si was ‘sealed and heated’ at\n180 °C. Could the authors specify the type of vessel used (e.g., autoclave, glass bottle)\nand the reaction atmosphere. In addition, the improvement in ICE of full cells relies\non prelithiation. Please specify the method.\nWe sincerely appreciate your valuable suggestions regarding our experimental\nmethodology and data interpretation. In response to your comments, we would like to\nclarify that all sample preparations were carefully conducted in sealed glass bottles to\nensure controlled reaction conditions, and we have now included detailed synthesis\nprocedures for the μSi@AlCl in Methods for better reproducibility. Regarding the\n3\nobserved improvement in initial Coulombic efficiency (ICE) of full cells, we attribute\nthis enhancement primarily to the significantly improved first-cycle efficiency and\nreversibility of the μSi@AlCl anode compared to bare μSi, as supported by our\n3\nsystematic electrochemical characterizations. Specifically, as shown in Figure R15,\nour comparative analysis of non-prelithiated full cells demonstrates a notable 6%\nincrease in ICE (from 79.2% to 85.2%), while further optimization through\nprelithiation achieves an even higher ICE of 86.6% (Figure R13).\nMethods\nSynthesis of materials\nAll synthesis procedures were performed under inert argon atmosphere in a glove box\n(H O<0.01 ppm, O <0.01 ppm) using simple thermal treatment.\n2 2\nFor μSi@AlCl , stoichiometric quantities of μSi and AlCl were mixed at mass ratios of\n3 3\n90:10, ground manually in an agate mortar for 5 minutes. The resultant mixture was\ntransferred to a sealed glass bottle and subjected to thermal treatment at 180°C on a\nhotplate for 2 hours, yielding the μSi@AlCl -10. By adjusting the mass ratios to 95:5 and\n3\n85:15 while maintaining identical synthesis conditions, μSi@AlCl -5 and μSi@AlCl -15 were\n3 3\nprepared respectively.\nFor other metal chloride modifications (μSi@MCl (M= P, Ta, Bi, Fe, Nb, Zn, Mg)), the\nx\nsame procedure was followed with modified thermal treatment temperatures: 450°C for ZnCl\n2\nand MgCl , while temperatures were adjusted according to the respective chloride boiling\n2\npoints for others.\nFigure R15. GCD profiles of (a) μSi-NCM, (b) μSi@AlCl -NCM, and (c) statistical data of ICE.\n3\ntemperature. (f) ICE statistics of Si-based SSBs with NCM cathode (see Figure 6).\n4. Since LPSC contains Cl⁻ and the Al(Si)OCl layer also contains Cl⁻, is there any\npossibility of Cl⁻ migration between them during cycling? The authors may need to\nconsider how such migration could affect electrochemical performance.\nWe sincerely appreciate your constructive suggestion. In response, we have\nconducted comprehensive EDS mapping to examine the Cl distribution evolution\nduring cycling, which reveals distinct interfacial characteristics (Figure R24 and\nFigure R25): (1) For the pristine μSi/LPSC interface, clear elemental segregation is\nobserved for Si, P, S and Cl; (2) After cycling, the μSi@AlCl /LPSC interface\n3\nmaintains well-defined elemental distributions, while the μSi/LPSC interface shows\nsignificant Cl accumulation, confirming severe interfacial side reactions and Cl\ndiffusion. Complementary XPS analysis further verifies LiCl formation in the cycled\nμSi/LPSC system (Figure R26). These findings collectively demonstrate that: (i) Cl-\nmigration occurs at the μSi/LPSC interface during cycling, and (ii) the\nμSi@AlCl /LPSC interface exhibits enhanced electrochemical stability that effectively\n3\nsuppresses Cl diffusion and parasitic reactions. The improved interfacial stability is\nfurther corroborated by EIS measurements showing lower interfacial impedance and\nmaintained charge transfer capability, consistent with the stabilized Cl distribution in\nthe modified interface (Figure S31).\nFigure R24. Elemental line-scan profiles across the (a) pristine-μSi/LPSC interface, (b) cycled\nμSi@AlCl /LPSC interface and (c) cycled μSi/LPSC interface, showing the relative intensity\n3\ndistributions of constituent elements (Si, Al, Cl, P, S).\nFigure R25. EDS mapping of the cycled μSi/LPSC interface and cycled μSi@AlCl /LPSC\n3\ninterface.\nFigure R26. XPS depth profiling analysis of (a) μSi/LPSC and (b) μSi@AlCl /LPSC interfaces\n3\nshowing Li 1s and Cl 2p core-level spectra evolution across the interfacial regions.\nFigure S31. EIS results of (d) Si@AlCl |LPSC|Li-In and Si|LPSC|Li-In in different\n3\nelectrochemical states.\n5. While the authors mention that the SA layer can suppress LPSC decomposition,\nfurther clarification and supporting evidence for the underlying mechanism are\nrequired.\nWe sincerely appreciate your valuable suggestion. The underlying mechanism by\nwhich the SA layer suppresses LPSC decomposition is attributed to the relatively\nelectrochemically inert SA component formed via surface modification. SA layer\neffectively isolates Si from direct contact with LPSC, thereby mitigating elemental\ninterdiffusion and side reactions at their interface during electrochemical processes;\nconsequently, it prevents degradation of LPSC. As conclusively validated in our\noriginal manuscript (Figure S5), comprehensive XPS characterization and\nsupporting experiments demonstrate that the SA layer exhibits favorable chemical\nstability against lithium metal and manifests distinctly inert electrochemical\ncharacteristics (Figure R2).\nTo provide further clarification and supporting evidence, we have conducted\nadditional cyclic EDS mapping and XPS depth profiling results (Figure R24-R26), we\ncan clearly identify substantial electrochemical decomposition of LPSC at the pristine\nSi|LPSC interface, as evidenced by both the pronounced interfacial accumulation of\nCl species and the formation of LiCl side products. In striking contrast, the modified\nSi@AlCl |LPSC interface demonstrates significantly suppressed LPSC decomposition,\n3\nas revealed by the well-maintained elemental distribution and minimal LiCl\ngeneration.\nFigure S5. (a) Experimental diagram of chemical stability study of SA electrode and Li metal. The\npristine SA electrode was closely bonded with Li foil and heated at 60℃ for 5h to react (labeled\nSA-Li-H5). XPS analysis of (b) Li 1s, (c) Si 2p, (d) Al 2p, (e) Cl 2p and (f) O 1s spectra for the\nSA-Li-H5.\nFigure R2. Galvanostatic charge/discharge (GCD) profiles of (a) SA+20 wt% Super\nP|LPSC|Li-In and (b) AlCl +20 wt% Super P|LPSC|Li-In. Cyclic Voltammetry (CV) curves of (c)\n3\nand CV curves confirms negligible capacity contributions from SA and AlCl .\nμSi@AlCl /LPSC interface and (c) cycled μSi/LPSC interface, showing the relative intensity\n3\ndistributions of constituent elements (Si, Al, Cl, P, S).\nFigure R25. EDS mapping of the cycled μSi/LPSC interface and cycled μSi@AlCl /LPSC\n3\ninterface.\nFigure R26. XPS depth profiling analysis of (a) μSi/LPSC and (b) μSi@AlCl /LPSC interfaces\n3\nshowing Li 1s and Cl 2p core-level spectra evolution across the interfacial regions.\n6. The manuscript reports an average Coulombic efficiency exceeding 102% for the\nhigh-loading Si|LPSC|Li-In cell, which requires further clarification and supporting\nevidence.\nWe appreciate the opportunity to clarify the observed Coulombic efficiency (CE)\nexceeding 100%, which we have systematically investigated through multiple\nexperimental approaches. The anomalous CE behavior predominantly occurs under\nhigh-rate conditions and can be attributed to two primary mechanisms: First, kinetic\nlimitations at the anode under high current densities induce lithium plating, as\nevidenced by our μSi/Li-In half-cell tests showing CE >100% only at rates exceeding\n1.5C (2C-3C), with normal CE values (<100%) restored upon returning to 1C\n(Figure R27a) - consistent with established literature19 (Nature Communications\nvolume 16, Article number: 1013 (2025); Figure 5d). Similar kinetic limitations occur\nunder high mass loading conditions, where our supplementary full-cell\ncharge/discharge profiles clearly exhibit characteristic lithium plating features\n(Figure R27b). Second, parasitic interfacial reactions between the Si anode and\nLPSC electrolyte during delithiation contribute additional charge capacity, as\ndemonstrated by dQ/dV analysis revealing substantial side reaction contributions in\nthe 0.3-1.2V range during charging (Figure R27c), particularly pronounced at high\ncurrents where these reactions account for a significant portion of the total capacity.\nFigure R27. (a) Rate capability of μSi/Li-In half-cells at varying C-rates, (b) Charge profiles of\nhigh-mass-loading μSi/NCM electrodes, and (c) dQ/dV plots at varying C-rates revealing reaction\nkinetics and side processes.\n[Figure Redacted]\nFigure 5d Nature Communications volume 16, Article number: 1013 (2025).\n7. The claim of “record-breaking” ICE in NCM full cells is misleading, as some cited\ncomparisons are without prelithiation while the present work uses prelithiation. Please\nadjust the claim or benchmark consistently.\nWe sincerely appreciate your constructive suggestions for enhancing the clarity and\nprecision of our manuscript. Regarding the statistical comparison of initial\nCoulombic efficiency (ICE) for NMC-based silicon all-solid-state batteries presented\nin Figure 6, we would like to clarify that this comprehensive analysis encompasses\nboth pre-lithiated and non-pre-lithiated systems from various reported studies.\nFollowing your valuable recommendation, we have carefully revised our own work to\nexplicitly distinguish between pre-lithiated and non-pre-lithiated conditions through\nclear labeling and appropriate terminology (Figure R13, updated data). Furthermore,\nwe have replaced potentially subjective descriptors like \"record-breaking\" with more\nscientifically rigorous terms such as \"excellent\" to more accurately characterize the\nperformance while maintaining appropriate academic restraint. These modifications\nhave been systematically implemented throughout the manuscript to ensure\nconsistency and prevent any potential misinterpretation, while preserving the\nscientific validity and technical accuracy of our findings.\nWe sincerely appreciate your constructive suggestions for improving the clarity and\nprecision of our manuscript. In response to your valuable feedback, we have carefully\nrevised our terminology regarding the pre-lithiation treatments to ensure\nunambiguous distinction between pre-lithiated and non-pre-lithiated systems.\nAdditionally, we have modified the language describing the electrochemical\nperformance from \"record-breaking\" to \"excellent\" where appropriate, adopting more\nmeasured and scientifically precise descriptors that better reflect the data while\nmaintaining academic rigor.\nFigure R13. (e) GCD profiles of Si@AlCl |LPSC|LIC|NCM after pre-lithiation at room\n3\ntemperature. (f) ICE statistics of Si-based SSBs with NCM cathode (see Figure 6).\n8. The manuscript contains inconsistent terminology, as the compound is first\nintroduced as “Li6PS5Cl (LPSC)” but later also appears as “LiPSCl”; please unify the\nnomenclature (e.g., consistently use “LPSC”)\nWe sincerely appreciate your thoughtful suggestion regarding terminology\nconsistency. In response to your valuable recommendation, we have carefully\nreviewed the manuscript and systematically standardized all instances of the solid\nelectrolyte notation to exclusively use the abbreviated form \"LPSC\" throughout the\ntext, figures, and supplementary materials.\n9. Typos such as “Figue” instead of “Figure,” and missing spaces (e.g., “fullcell”)\nshould be corrected. Please double check.\nWe sincerely apologize for the typographical errors in our manuscript and greatly\nappreciate you bringing these issues to our attention. In response to your valuable\nfeedback, we have conducted a thorough proofreading of the entire manuscript and\nsupplementary materials to identify and correct all typographical errors. Your careful\nreview has been instrumental in improving the quality of our work, and we deeply\nappreciate your time and attention to detail."}, {"round": 2, "reviewer_comments": {"reviewer_1": "This manuscript has been well revised for publication in Nature Communications.", "reviewer_2": "The authors substantially improved the manuscript based on the reviewer's comments. Now, it can be published.", "reviewer_3": "This manuscript has been thoroughly revised by providing more experimental details, performance analysis, and theoretical explanation. Therefore, it’s acceptable in its current form."}, "author_response": ""}]

Decision letter