Interfacial Redox Buffering Stabilizes Halide Solid Electrolytes Against Low-Potential Anodes

Interfacial Redox Buffering Stabilizes Halide Solid Electrolytes Against Low-Potential Anodes | Research Square window.SnipcartSettings = { analytics: { enabled: false } }; (function() { var accessVector = localStorage.getItem('access_vector') || ''; window.dataLayer = window.dataLayer || []; if (accessVector) { window.dataLayer.push({ user: { profile: { profileInfo: { snid: accessVector } } } }); } })(); (function(w,d,s,l,i){w[l]=w[l]||[];w[l].push({'gtm.start':new Date().getTime(),event:'gtm.js'});var f=d.getElementsByTagName(s)[0],j=d.createElement(s),dl=l!='dataLayer'?'&l='+l:'';j.async=true;j.src='https://www.googletagmanager.com/gtm.js?id='+i+dl;f.parentNode.insertBefore(j,f);})(window,document,'script','dataLayer','GTM-K279D39R'); Browse Preprints In Review Journals COVID-19 Preprints AJE Video Bytes Research Tools Research Promotion AJE Professional Editing AJE Rubriq About Preprint Platform In Review Editorial Policies Our Team Advisory Board Help Center Sign In Submit a Preprint Cite Share Download PDF Article Interfacial Redox Buffering Stabilizes Halide Solid Electrolytes Against Low-Potential Anodes Guoying Chen, Hyunwon Chu, Shuhao Yang, Sanggyu Chong, Pravin Didwal, and 1 more This is a preprint; it has not been peer reviewed by a journal. https://doi.org/ 10.21203/rs.3.rs-8928802/v1 This work is licensed under a CC BY 4.0 License Status: Under Review Version 1 posted You are reading this latest preprint version Abstract The practical adoption of solid-state batteries requires solid electrolytes that sustain stable interfaces with electrodes. Conventionally, electrolyte stability has been defined by its thermodynamic limits, assuming that operating beyond these boundaries causes irreversible decomposition. Such perspectives have deterred the integration of many solid electrolytes with low-potential anodes, leaving their behavior beyond the limit largely unexplored. Here, we demonstrate that the trivalent metal halide Li 3 YCl 6 (LYC) can be electrochemically driven below its calculated reduction limit while exhibiting a predominantly reversible redox response. Lowering the potential below the thermodynamic window activates a reversible lithiation process consistent with Y-centered redox, while the halide framework is retained. At the anode interface, this reversibility enables LYC to accommodate low-potential anodes by forming a self-limiting lithiation layer that kinetically suppresses continued reduction. Comparative anode screening with Li-In, Li-Si, and Li metal identifies practical operating window below the nominal reduction threshold, within which LYC maintains both phase retention and ion transport with manageable impedance growth. Utilizing the reversible regime, we realize coating-free LYC full cells with an NMC811 cathode and a Li-Si anode that cycle stably for over 500 cycles at high capacity (~ 140 mAh g -1 ) and extended cell voltages (> 4 V), expanding the practical anode selection beyond conventional thermodynamic-window constraints and enabling high-performance all-solid-state cells. Physical sciences/Materials science/Materials for energy and catalysis/Batteries Physical sciences/Chemistry/Electrochemistry/Batteries Figures Figure 1 Figure 2 Figure 3 Figure 4 Figure 5 1. Introduction The search for solid electrolytes (SEs) that combine high ionic conductivity and stable electrode interfaces remains a central challenge in solid-state battery (SSB) research 1-2 . Recent advances in halide electrolytes have begun to address both requirements simultaneously, as a new generation of chloride or bromide-based materials has been shown to exhibit room-temperature ionic conductivities around 1 mS cm -1 3 and, in some cases, exceeding 10 mS cm -1 4 , while maintaining strong compatibility with high-voltage (HV) lithium-ion battery (LIB) cathodes 5 . For many years, halide compounds were explored as Li-ion conductors 6-7 but attracted limited attention because their ionic conductivities lagged behind those of sulfide-based electrolytes 8-9 . This perception has shifted only recently with the discovery of highly conductive halide electrolytes, repositioning halides from a niche materials class to prime candidates for practical all-solid-state batteries 10 . Beyond their high ionic conductivity, halide electrolytes offer a distinct advantage in oxidative stability around 4 V vs. Li/Li + 11,12 , based on the intrinsic robustness of the halide anion framework. Compared to oxide or sulfide counterparts, the chloride sublattice lacks easily oxidizable species, minimizing parasitic interfacial reactions at high-voltage oxide cathodes 13 and enabling simplified composite cathode designs without protective layers 5,14-15 . Within this family, Li 3 YCl 6 (LYC) has been widely investigated as a promising electrolyte for all-solid-state battery designs targeting high energy density. LYC offers high ionic conductivity (0.51 mS cm -1 at room temperature 5 ), good processability (Poisson’s ratio of 0.274 16 ), and excellent oxidation stability (up to 4.21 V vs. Li/Li + 17 ). These attributes have driven extensive studies of LYC-based full cells in combination with conventional LIB cathodes and even high voltage electrodes 13,18-19 , where the electrolyte remains largely chemically inert. However, LYC, like other metal halide electrolytes, has often been regarded as intrinsically unstable when paired with low-voltage anodes. Thermodynamic calculations predict that LYC enters a thermodynamically reducing regime below 0.62 V vs Li/Li + 17 , a trend that is consistent with electrochemical observations 20-21 . Such behavior has been taken as indicative of phase decomposition driven by reduction of the cationic sublattice. Specifically, reduction of Y 3+ is predicted to promote the formation of reduced yttrium chlorides (Y 2 Cl 3 ) and, under sufficiently reducing conditions, metallic Y (Y 0 ) accompanied by LiCl-rich phases 17 . Importantly, the decomposition-based interpretation is not unique to LYC that most transition-metal halides are assumed to always decompose into electronically conductive and ionically blocking products beyond its thermodynamic limit 12 . Because of this prevailing assumption, the low-potential behavior of halide electrolytes has rarely been examined systematically, except a few studies focused on direct Li-metal contact 21-22 . Instead, most LYC-based solid-state batteries have been designed to avoid any reductive conditions below its intrinsic threshold, mostly by pairing the electrolyte with elevated-potential anodes such as Li-In alloys 5,19,23 . While this strategy ensures stable operation within the thermodynamic window, the high anode potential substantially narrows the accessible cell voltage (< 4 V for Li-In full cells) 5 and limits the achievable energy density. While adding a sulfide-based interlayer has been demonstrated with some success, the complexity of adding another layer limit its practical use. As such, the perceived reduction instability of halide electrolytes has effectively acted as an inherent design constraint. Against this long-standing view, recent works have started to re-examine whether reduction of halide electrolytes must necessarily result in catastrophic breakdown. A growing body of literature reported that redox activity within halide frameworks can proceed in a reversible and non-destructive manner when mediated by multivalent cations such as Fe (Li 2 FeCl 4 24 and Li 3 FeCl 6 25 ), V (Li 3 VCl 6 26 ), and Ti (Li 3 TiCl 6 27 ). In these systems, reduction does not immediately disrupt the host structure but instead occurs through well-defined redox processes that preserve the integrity of the halide lattice 28 . Such reversible redox behavior has been demonstrated most clearly on the cathode side, where transition-metal redox in the halide framework falls within the typical operating potential window of oxide cathodes (i.e., Li 2 FeCl 4 at ~3.6 V vs. Li/Li + 24 ). The dynamic redox activity of the halide electrolyte contributes an additional 20-50% reversible capacity to cathodes while maintaining structural integrity under extended cycling conditions 28 . At a broader level, similar behavior has also been reported in Na + -conducting halide electrolytes 29 , further underscoring that reduction does not universally induce irreversible decomposition of halide-based electrolytes. While redox behavior has primarily been explored to extend cathode capacity, a few studies have newly extended this redox-based view to anode-relevant voltages using lower-redox-potential cations 30 . For instance, in mixed-halide Li 3 YCl 3 Br 3 (LYCB), stable lithiation/de-lithiation has been observed below the nominal stability window (< 0.6 V) driven by reversible Y redox 31 . When incorporated into anodic composites, the low-voltage redox activity enables extended cycling of red phosphorus (RP) anodes 31 . Nevertheless, in these anode “composite” configurations, the redox activity remains confined within the electrode region and is leveraged mainly for anode material activation and stabilization. A key remaining question is how the same redox behavior manifests when the halide electrolyte itself serves as a stand-alone separator and directly interfaces with a low-potential anode in a practical solid-state battery configuration. In this work, we show that LYC is not limited by its thermodynamic stability window, but instead supports reversible Li + insertion/extraction that expands its operating range toward low-potential anodes. When the potential is pushed below the nominal reduction limit, we observe that a reversible reduction-oxidation response emerges in LYC framework with the main lithiation feature centered at 0.25 V vs. Li/Li + . After the initial activation, the process becomes highly reversible and delivers a stable cycling capacity of ~ 150 mAh g -1 with a high Coulombic efficiency (CE) over 99 %. Synchrotron-based characterization across selected discharge/charge states verifies that the Li uptake and release proceed without loss of framework integrity. Computational modeling further confirms the structural stability of the accessible lithiation states and elucidates the local coordination changes during lithiation. Benefiting from LYC’s intrinsic redox, we next evaluate whether LYC can function as a stand-alone separator against low-potential anodes. In-situ tracking of charge-transport kinetics indicates that lithiation of LYC forms a self-limiting interphase that buffers further reduction. Separator-level application, however, requires careful control over an operating lithiation range, since deeper lithiation introduces excessive ionic polarization. Based on this mechanistic understanding, anode-side behavior of LYC was compared across three representative electrode materials spanning progressively lower potentials: Li-In alloy (within the thermodynamic window), Li-Si alloy (partial lithiation), and Li metal (deep lithiation). Access to the reversible Li-exchange regime enables stable operation of halide-based Si full cells (LiNi 0.8 Mn 0.1 Co 0.1 O 2 (NMC) | LYC | Li-Si) for over 500 cycles without any protective layers. This protection-free compatibility directly allows the lower Si anode potential to be utilized so that the cell operating voltage can be extended up to 4.1 V with NMC electrodes (compared to 3.7 V for Li-In cells). In contrast, Li metal drives excessive lithiation at the interface, creating a highly resistive contact that promotes current localization and rapid failure, defining the practical limit of redox buffering that can be utilized in LYC-based cells. Overall, the results establish reversible Li exchange as a practical lever to extend halide separators beyond their nominal reduction limits, opening a viable route to pair LYC with low-potential, high-capacity anodes in coating-free full cells without sacrificing long-term stability. 2. Results & Discussion 2.1 Reversible Li Exchange Beyond Thermodynamic Limit Before examining the low-voltage electrochemical behavior of the LYC electrolyte, we first confirmed the structural and morphological uniformity of as-synthesized powder to ensure that the observed redox features originate inherently. The LYC electrolyte was prepared by high-energy ball milling, and the resulting particle morphology consists of sub-micron primary particles that aggregate into secondary clusters extending to several tens of microns ( Supplementary Fig. 1 ), which is typical of halide electrolytes prepared by mechanochemical synthesis routes 20 . The X-ray Diffraction (XRD) pattern displays reflections consistent with the targeted LYC phase and does not show detectable impurity peaks. The samples’ homogeneity is essential because any redox activity discussed later should be attributed to the intrinsic LYC rather than to secondary phases or processing byproducts. After confirming the phase purity, we performed cyclic voltammetry (CV) at progressively lower cutoff voltages to determine whether LYC engages in a reversible electrochemical process ( Fig. 1a ). A Stainless Steel (SS) | LYC | Lithium-Indium (Li-In) cell was used, where the Li-In alloy prevents direct chemical reactivity at the counter electrode and the SS plate minimizes capacitive artifacts in the working-electrode response. LYC is predicted to have a nominal reduction limit at 0.62 V vs. Li/Li + , associated with the Y 3+ /Y 0 redox potential 17 . Potentials below this range are shown by theory to drive thermodynamically favorable reduction reactions 17,21 , although the extent to which these reactions proceed under operating conditions remains an experimental question. When the lower cutoff is held at 0.62 V, the current response remains essentially featureless, as expected within its thermodynamic stability window. As the cutoff is gradually lowered, however, a new reduction feature starts to develop. At 0.2 V cutoff, a distinct cathodic peak (R1) becomes clearly visible, and extending the cutoff to 0 V sharpens this feature into a well-defined peak centered at 0.25 V. The systematic emergence of the peak only under deeper reducing conditions suggests activation of an additional electrochemical process beyond the thermodynamic window. Reversing the potential scan provides further insight into whether LYC participates in a reversible electrochemical process. Interestingly, we observed an oxidation peak (O1) emerges near 0.7 V on the return sweep, forming a redox pair with the low-voltage reduction (R1) feature. In the 1st cycle, the CE ranges from 30 to 45 % depending upon the cutoff voltages, implying that a portion of the initial reduction introduces irreversible processes. However, the efficiency increases steeply with continued cycling, surpassing 90 % by the 2nd cycle and approaching 99 % by the 20th CV cycle ( Supplementary Fig. 2 ). The rapid improvement supports that, after an initial conditioning step, LYC begins to follow a predominantly reversible redox pathway rather than sustained, cumulative side reactions dominating the response. To additionally verify whether the observed redox activity comes from LYC itself rather than other electrode components, we also varied cell configurations to include or exclude conductive carbon (C) additives ( Supplementary Fig. 3 ). Regardless of configuration, the same reduction feature near 0.3 V and oxidation feature near 0.7 V were observed. The reproducibility across different electrode setups supports that the dominant redox features originate from LYC rather than the current collector or carbon additive. Such redox behavior becomes more evident under galvanostatic cycling, which allows a more quantitative assessment of its magnitude and reversibility. For these measurements, the cell configuration was slightly adjusted to utilize an LYC/C composite electrode so that the active mass could be precisely defined for capacity determination. Under a constant-current discharge to 0 V, the LYC/C composite delivers an initial capacity of 162 mAh g -1 ( Fig. 1b ). The subsequent first-charge capacity reaches 113 mAh g -1 , corresponding to an initial CE of about 70 %. The lower efficiency reflects the initial irreversible portion of the reaction that was also observed in the CV sweeps. With continued cycling, however, the efficiency rapidly increases and exceeds 97 % by the 10th cycle ( Supplementary Fig. 4 ). The main contributors to the high capacity can be more clearly identified when the voltage profiles are converted to differential capacity (dQ/dE) plots ( Supplementary Fig. 4 ). During reduction, two distinct processes appear: a prominent peak near 0.4 V and a second rise below 0.1 V. The higher-voltage peak closely aligns with the R1 reduction feature observed in the CV scans, whereas the low-voltage contribution near 0 V likely reflects a decomposition-related process that lowers CE during the early galvanostatic cycles. This secondary pathway is far less apparent in the CV measurements, likely because the SS electrode offers limited interfacial contact area. In contrast, the composite electrode used for galvanostatic operation presents a much larger active surface, which accentuates the surface-related decomposition contribution. Even with this additional feature, CV and galvanostatic data together show that once LYC passes through its initial chemical adjustment, the predominant electrochemical response settles into a highly reversible regime when the lower cutoff is limited to avoid the most reducing potentials. The voltage profile of the LYC/C composite under galvanostatic discharge shows a broad sloping region often linked to solid-solution-type insertion in Li-ion electrode materials 32 . The similar response motivated us to examine whether LYC undergoes measurable changes in Li stoichiometry through reversible lithiation and de-lithiation. To evaluate dynamic Li-ion exchange in LYC, we first quantified the number of electrons transferred from the measured capacity and converted the value into the equivalent amount of inserted Li + . The electrochemically estimated Li content was then compared with the Li/Y ratio obtained from Inductively Coupled Plasma-Optical Emission Spectroscopy (ICP-OES) analysis of cycled samples ( Fig. 1c ). The initial discharge capacity of 162 mAh g -1 corresponds to an estimated 1.95 mol of inserted Li + per mol of LYC. When the discharged sample was analyzed, the ICP measurement gives the Li/Y ratio of 5.09. Assuming no Y loss during cycling, the ratio corresponds to a LYC stoichiometry of Li 5.09 YCl 6 . The measured composition is in close agreement with the expected value of Li 4.95 YCl 6 derived from electrochemical insertion of ~1.95 mol of Li. Furthermore, the Li content decreases upon the 1st charge, and the subsequent discharge step re-introduces additional Li⁺ in a manner that closely follows the capacity-based Li estimate. Overall, the electrochemical capacity aligns with reversible changes in Li stoichiometry, indicating that the redox response originates from Li-exchange within the LYC framework. The reversible Li exchange within the LYC framework remains stable in multiple insertion and extraction cycles. Since deep discharge near 0 V led to a current rise that likely initiated minor decomposition and a small CE loss, we narrowed the operation potential window and set the lower cutoff at 0.1 V. As the majority of the capacity arises from the plateau region between 0.2 and 0.3 V, the discharge capacity of the LYC/C composite stayed around 150 mAh g -1 while maintaining high reversibility over 50 cycles with an average CE of 99.63 % ( Fig. 1d ). Cycling at higher current resulted in some capacity reduction due to increased polarization, yet the cell still delivered about 110 mAh g -1 with an average CE near 99 %. The sustained efficiency shows that the activated electrochemical process continues to operate in a reversible manner well beyond LYC’s thermodynamic stability window. The extended reversibility of lithiation and de-lithiation demonstrates that LYC repeatedly accommodates Li + without loss of function, prompting a detailed examination of how its lattice and phase stability evolve as Li content varies. 2.2 Structural Integrity and Redox-Driven Chemical Changes The crystal structure of Li 3 YCl 6 (trigonal, space group P3̅m1 ) shows that Li⁺ ions intrinsically occupy the Wyckoff 6g and 6h positions with occupancies of 100 and 50 %, respectively 33 . A simple stoichiometric consideration indicates that these two intrinsic Li sites can accommodate an additional 1.5 mol of Li + , allowing lithiation up to approximately Li 4.5 YCl 6 . Prior studies have also reported that LYC contains tetrahedral voids that can serve as additional insertion sites 5,33 ( Fig. 2a ). Structurally, therefore, LYC possesses sufficient Li-accessible sites to support dynamic Li exchange without requiring a phase transformation. To examine how the LYC lattice responds during reversible Li exchange, we recovered LYC/C powders cycled to different voltages and performed synchrotron XRD to examine structural evolution. Across all states of charge, the major reflections match exclusively to the LYC phase, and no crystalline decomposition products such as Y metal, LiCl, or YCl 3 are detected within the XRD detection limits ( Supplementary Fig. 5 ). More importantly, the diffraction patterns exhibit clear and systematic peak shifts as a function of the electrochemical state ( Fig. 2b ). In pristine LYC, three representative reflections appear in the 2θ range of 12-17°, corresponding to the (002), (301), and (112) planes. During the 1st discharge, no noticeable peak shift was observed up to 0.5 V, which aligns with the electrochemical response where lithiation begins near 0.3 V. Upon subsequent discharge to 0 V, however, the (112) reflection shifts to higher 2θ from 15.72° to 16.22°, and comparable shifts occur for the (002) and (301) peaks. The shift to higher 2θ is consistent with a decrease in d-spacing (lattice contraction). Using the Bragg’s equation, we find that the c-axis lattice parameter decreases from 6.07 Å to 5.93 Å during the 1st discharge, bringing roughly 2 % shrinkage in unit-cell volume ( Supplementary Fig. 6 ). The trend largely mirrors the early-stage lithiation behavior of layered oxide cathode materials such as LiNi x Mn y Co 1-x-y O 2 (NMC), where lithiation induces c-axis contraction from reduced repulsion between neighboring anion planes 34 . Notably, the contraction is fully reversible upon de-lithiation. When charged back to 1 V, the diffraction peaks of LYC/C return to lower 2θ and restore the lattice parameters close to their pristine values. Such reversible lattice breathing persists over the initial five cycles we examined. The structural stability is further corroborated by atomistic simulation with Machine-Learning Interatomic Potential (MLIP, PET-OAM-XL 35 ) and Density Functional Theory (DFT) calculations, where structural relaxations confirm that the LYC networks are reasonably retained across the lithiation range of Li 3 to Li 5 ( Supplementary Fig. 7 ). Throughout all conditions examined, both spectroscopic and computational analyses show no evidence of significant phase decomposition. reassuring that the LYC framework remains structurally coherent and resilient during repeated lithiation and de-lithiation. Building on this bulk-scale structural analysis, we then investigated whether the local bonding environment exhibits the same level of stability using X-ray Absorption Spectroscopy (XAS). XAS data are commonly interpreted by separating the near-edge (XANES) region, which reflects oxidation-state and electronic changes, from the extended fine structure (EXAFS) region, which captures the local coordination and bond environment 21,36 . When comparing the EXAFS region of the Y K-edge across different states of charge, the spectra overlap almost completely as the local coordination around Y remains largely unchanged during discharge and charge ( Fig. 2c ). In contrast, the XANES region captures clear signatures of Y-centered redox activity. The half-maximum edge position shifts by 1.1 eV toward lower energy upon 1st discharge (1D, from 17043.6 to 17042.5 eV), showing that Y 3+ in pristine LYC is partially reduced during lithiation. In the subsequent charging step (1C), the edge returns to its initial position, a clear indication that the Y valence state is restored based on the highly reversible redox process. The similar shift pattern recurs over multiple cycles and continues through at least the 10th discharge. Independent analysis of the white-line peak maximum shows a comparable energy evolution as the half-maximum edge positions ( Supplementary Fig. 8 ): the peak moves to lower energy during discharge and shifts back to higher energy upon charge. Comparison with reference compounds YCl 3 (Y 3+ ) and Y 2 Cl 3 (Y 1.5+ ) suggests that the discharged state approaches an average Y valence near +1, consistent with the ICP-derived stoichiometry of Li 5.09 YCl 6 (Y 0.91+ ). We note that a quantitative oxidation-state estimate depends on reference selection and fitting model. These XAS results, from both EXAFS and XANES regions, confirm that LYC undergoes a controlled Y-centered redox process rather than a decomposition-driven transition. The Cl K-edge spectra also support this understanding. Across all states, the Cl K-edge data retain its overall line shape and edge position, resembling a stable mixture of Li-Cl and Y-Cl coordination motifs. Complementing the XAS measurements, radial distribution function (RDF) analysis of Molecular Dynamics (MD) simulations with the MLIP model confirms that the local coordination environments of Li-Cl and Y-Cl experience no severe changes across varying lithiation levels ( Supplementary Fig. 9 ). The absence of coordination changes indicates that chloride ions do not participate in the redox process and that the halide sublattice remains intact even under deep lithiation. X-ray Photoelectron Spectroscopy (XPS) measurements further monitor the surface evolution of elemental valence states during Li exchange ( Fig. 2d ). In the Y 3d core-level spectra, pristine LYC/C shows only Y 3+ presence with the Y 3d 5/2 peak centered around 159.2 eV. This value is consistent with reported binding energies for trivalent Y in halide environments 37 . After the 1st discharge, an additional component emerges at lower binding energy near 157.7 eV, which indicates the formation of partially reduced Y states (Y n+ , n < 3, exact peak position provided in Supplementary Fig. 10 ). Reference binding energies place metallic Y 0 near 155.8 eV and intermediate-valence Y 1.5+ species (such as Y 2 S 3 ) around 157.4 eV 38 . The new feature at the discharged states remains well above the energy range characteristic of Y 0 , suggesting that reduction in LYC does not progress all the way toward metal formation. Instead, its position aligns with a partially reduced Y state expected within the lithiated framework. Quantitative fitting shows that the Y 3+ to Y n+ ratio shifts to roughly 30:70 after the 1st discharge. Upon charging, the ratio partially recovers to 61:39, showing substantial reversible reoxidation. During the 2nd discharge, the reduced component increases again in a reproducible manner, matching the cycle-to-cycle evolution and reversible Y-centered redox behavior observed in XAS. Throughout all states of charge, the Cl 2p spectra remain essentially unchanged in both binding energy and line shape, reaffirming that the chlorine sublattice maintains chemical stability during cycling. Together, these structural and spectroscopic results cohesively support that LYC accommodates Li + insertion through reversible changes in its Y-centered electronic states while preserving the integrity of the halide framework. 2.3 Ion Transport Kinetics Across Lithiation States Although the structural framework remains intact, lithiation inevitably reduces the population of vacant Li sites in the LYC lattice, which can be expected to lower ionic conductivity as Li content increases. Computational Li-Y RDF profiles reveal that lithium atoms increasingly occupy shorter-distance sites starting around the Li 4 level ( Supplementary Fig. 9 ), indicating the progressive filling of tetrahedral interstitials that normally function as conduction pathways 17 . Understanding how ionic transport evolves under progressive lithiation is therefore important for describing the electrolyte’s behavior when in contact with low-voltage anodes. To trace how transport kinetics change, we collected the in-situ Electrochemical Impedance Spectroscopy (EIS) response of an LYC/C | LYC | Li-In cell across different lithiation states ( Fig. 3a ). During discharge, we observe that the impedance shows a systematic rise in the low-frequency region. The response is dominated by a semicircle centered around 0.64 Hz, which reflects ion-transport processes in the LYC/C composite 31 . Once the potential enters the main lithiation regime below 0.3 V, the impedance begins to increase and continues to grow toward 0 V. The same evolution appears in the Bode plot and in the single-frequency transients collected at 0.64 Hz ( Supplementary Fig. 11 ). The resistance increase during lithiation can be interpreted through progressive filling of Li vacancies. As fewer vacancies become available, the number of mobile carriers declines and Li + transport through the composite becomes increasingly hindered at deep lithiation states. It is also worth noting that the impedance of LYC/C increases rather than decreases. In systems where an ionic conductor develops substantial electronic conductivity, the added electronic pathway often lowers the overall impedance 39 . Here, however, the increase in resistance implies that any electronic conduction that may develop does not compensate for the loss of Li + carriers. Therefore, the charge transport in lithiated LYC is still primarily dominated by ionic motion and does not show evidence of an electronically percolating pathway that dominates the impedance response over the examined window. On the other hand, as the cell enters the charge step, the low-frequency impedance begins to decrease once the potential rises above roughly 0.5 V. The resistance drop reflects the reopening of vacancy sites that were filled at high Li content. By the time the cell approaches 1 V, the semicircle associated with the LYC/C contribution returns to values similar to the pristine state, showing a near-complete recovery of optimal ionic transport pathways. Distribution of Relaxation Times (DRT) analysis provides a clearer quantitative picture of how the impedance changes arise 40 . In the pristine LYC/C cell, the DRT spectrum contains three distinguishable peaks ( Supplementary Fig. 12 ). As discharge proceeds, these features reorganize into two dominant peaks: a short time-constant (τ) component near 10 -6 s associated with bulk electrolyte conduction (P1), and a long time-constant contribution near 10 0 s that reflects ionic transport in the LYC/C electrode (P2). To visualize how the kinetic signatures evolve across the full voltage range, the extracted DRT intensities (γ) were mapped as a function of potential ( Fig. 3b ). With increasing lithiation, we find that the amplitude of the long-τ P2 intensity gradually increases more than 30 times (from 17.4 to 649 Ω), indicating a substantial growth in resistance originating from the LYC/C composite. Integrating this peak yields the corresponding absolute resistance, which grows markedly from 68.8 to 1944 Ω. During de-lithiation, however, the P2 peak intensity rapidly collapses toward its pristine magnitude once the potential exceeds 0.5 V. In contrast to the stepwise impedance buildup observed during lithiation, the recovery occurs in a single sharp transition, implying that Li + extraction from LYC does not experience the same kinetic constraints that limit insertion. A similar evolution and recovery pattern repeats over multiple cycles as the impedance response follows a consistent Li + insertion-extraction pathway ( Supplementary Fig. 13 ). The kinetic asymmetry between lithiation and de-lithiation can be well captured by evaluating the effective Li + diffusivity using Galvanostatic Intermittent Titration Technique (GITT, Fig. 3c ). In this method, small current pulses introduce incremental Li⁺ and the subsequent relaxation isolates kinetic overpotentials from steady-state changes 41-42 , allowing the effective diffusivity to be determined for insertion and extraction (detailed step provided in Supplementary Fig. 14 ). Under lithiation, the effective Li⁺ diffusivity starts on the order of 10 -8 cm 2 s -1 (marked as Stage 1) but gradually falls by nearly three orders of magnitude to 10 -11 cm 2 s -1 around 0 V (Stage 3). Over this range, the open-circuit voltage (OCV) changes are minor, while the bias overpotential grows strongly as kinetic barrier for further insertion increases. By contrast, the effective diffusivity stays near 10 -8 to 10 -9 cm 2 s -1 through most of the charge process (Stage 4) and decreases only near the end of charge (Stage 5), where the OCV shifts as extractable Li gets depleted. Therefore, the kinetic penalty for removing Li is far smaller over most of the composition range than for inserting Li into an already heavily lithiated framework. One mechanistic interpretation is that lithiation proceeds heterogeneously, forming a Li-rich region near particle surfaces while the interior remains relatively Li-poor (Stage 2). As lithiation advances, incoming Li + must traverse this increasingly Li-filled region, where ionic transport is reduced due to fewer available vacancies/interstitial pathways, thereby amplifying concentration gradients and polarization. This transport contrast can promote an effective ‘shrinking-core’-like behavior even though the bulk voltage response appears largely continuous. Upon de-lithiation, Li can be removed from the Li-rich near-surface region first (Stage 4), rapidly reopening percolating diffusion pathways and restoring lower resistance. Taking the electrochemical observations together with spectroscopy data, it becomes clear that LYC can incorporate additional Li + without major structural changes, while maintaining its predominant ionic conduction. Such a response suggests that lithiation does not readily progress into a significant phase decomposition, nor does it trigger a transition to a mixed ionic-electronic state. The outcome fundamentally differs from the common perception based on the theory that LYC should fully reduce to metallic Y beyond its thermodynamic limit, which would form an electronically conductive phase driving short-circuiting. Instead, because the lithiated LYC preserves its ionic dominance with minimal electronic leakage, it can form an interphase that is kinetically self-limiting under these conditions, analogous in function (though not necessarily composition) to a Solid Electrolyte Interphase (SEI) in liquid electrolyte systems. By suppressing further decomposition, this mode of reduction provides a mechanistic basis for enabling LYC to stably interface with low-voltage anodes beyond its nominal thermodynamic limit without undergoing runaway breakdown. 2.4 Dynamic Redox Activities Enabling Low-voltage Anodes LYC is formally predicted to be unstable below its reduction onset of 0.62 V vs. Li/Li + 17 . Most prior demonstrations therefore paired LYC with a Li-In anode at the cell level, whose main operating potential lies near the same voltage range 5,19,23 . However, the elevated anode potential sacrifices a large portion of cell operation voltage window and significantly limits the achievable energy density. On the other hand, our findings cohesively direct that the electrolyte can take up additional Li through a dynamic redox. Reversible lithiation enables LYC to kinetically buffer a reducing environment without decomposition and to operate against low-voltage anodes beyond its nominal threshold ( Fig. 4a ). When the operational window of LYC can be extended, several attractive low-voltage, high-capacity anode chemistries become accessible ( Fig. 4b ). For instance, Li-Ag alloy delivers a capacity (670 mAh g -1 ) comparable to Li-In while offering a broader voltage span (0.175 V vs. Li/Li + 43 ). Alloy anode such as Li-Mg further enlarges both specific capacity (2150 mAh g -1 ) and operating voltage (0.03 V 43 ), which effectively boost the energy density of LYC-based systems. However, it should be noted that ionic-conductivity loss becomes substantial at deep lithiation. An anode that is too reducing may still impose large interfacial resistance and introduce polarization that compromises performance. These considerations highlight the need to select an anode whose operation potential maintains a balanced and manageable interfacial kinetics. To evaluate how LYC behaves when paired with different anodes, we selected three representative electrode materials spanning progressively lower potentials: Li-In, operating within LYC’s nominal stability around 0.6 V; Li-Si, which provides a moderately reducing environment between 0.2 to 0.4 V; and Li metal, presenting the strongest reducing condition at 0 V vs. Li/Li + . Symmetric cells (M | LYC | M) were assembled using each alloy (M) and characterized under both storage and cycling conditions to determine which anodes can realistically operate with LYC. Under open-circuit voltage (OCV) storage, the Li-In cell exhibits the most stable response based on its thermodynamic compatibility ( Fig. 4c ). Its Nyquist impedance spectra resolve into three distinct semi-circles, attributed to ion conduction within LYC (> 10 6 Hz), ionic transport across the Li/LYC interface (~ 10 4 Hz), and a diffusive process in the alloy electrode (< 10 2 Hz), respectively 44 . Over 1000 min of storage, no additional features appear and the overall magnitude remains constant, assuring that Li-In anode does not introduce any reducing drive toward LYC. The Li-Si cell also maintains a relatively stable impedance profile during storage. Both the magnitude and spectral shape remain unchanged except for the low-frequency electrode region, suggesting that exposure to a Li-Si anode does not trigger reduction or the formation of a resistive interphase during that period. In contrast, the Li metal cell behaves differently from the outset. Its initial interfacial impedance at mid-frequency is substantially higher than that of Li-In and Li-Si, which originates from strong lithiation at the immediate interface forming a highly resistive region. Interestingly, the interfacial resistance then reaches a steady value after a mild increase. The absence of continued impedance growth may indicate that the reaction remains confined to the interfacial vicinity rather than extending into the bulk electrolyte. Under galvanostatic cycling, the contrasts among the three systems sharpen ( Fig. 4d ). The thermodynamically-stable Li-In symmetric cell maintains long-term operation with minimal polarization buildup over 1000 hours. In contrast, the Li-Si cell shows a noticeable rise in polarization during the first ~100 h, followed by a gradual relaxation. By around 300 h, the cell settles into a lower and more stable voltage window and then maintains the magnitude during prolonged cycling over 1000 hours. The Li cell initially follows a similar trend to Li-Si, and the polarization increases over the first few tens of hours. However, the cell voltage collapses abruptly within the first 100 h, consistent with internal short-circuit formation. The cell eventually reaches a complete hard short at around 515 h, as evidenced by the voltage dropping to near-zero level without recovery. To identify the origin of the polarization, we collected EIS spectra in parallel with cycling ( Fig. 4e ). Equivalent-circuit fitting separated the Li/LYC interfacial resistance (R int ) and allowed us to track its evolution (fitting details in Supplementary Fig. 15 ). In the Li-In cell, R int starts at 29.1 Ω and remains in a narrow range, decreasing slightly to 20.2 Ω after 900 h of cycling. The nearly constant interfacial resistance indicates that no additional reaction pathways are activated and that the interfacial chemistry remains preserved over extended cycling. The Li-Si cell shows a more complex impedance response. As cycling begins, an additional intermediate-frequency contribution (~ 10 5 Hz) newly appears between the electrolyte and interfacial resistances and rises to 101 Ω at 100 h. With continued cycling, the feature gradually diminishes and the overall impedance stabilizes at a lower level. Impedance elements in this frequency range are commonly assigned to thin resistive layers forming at electrode-electrolyte contacts, often called as “surface film” resistance 44-45 . In liquid systems, the contribution corresponds to SEI-related components 46 . By analogy, we infer that the impedance contribution observed here can be interpreted as a resistive lithiation layer that develops at the LYC/Li-Si interface. In the early stage of cycling, LYC near the Li-Si contact is rapidly driven into a highly lithiated state, which appears as an increased film-type resistance. With repeated lithiation and de-lithiation, continued Li exchange gradually redistributes this highly lithiated region over a larger thickness, reducing its local blocking effect ( Fig. 4d ). In line with this picture, the film resistance (R film ) decreases from 101 Ω to 39.6 Ω by 300 h and remains largely unchanged thereafter. More importantly, the bulk impedance of LYC shows no significant change over time as the lithiation process stays confined to a near-surface region. This localized redox buffering preserves the bulk electrolyte and enables stable galvanostatic cycling of the Li-Si anode with LYC for more than 1000 h. Compared to the other two cells, the Li metal cell already shows a strong drop in both bulk and interfacial resistances within the first 100 h. The rapid decrease may originate since the Li/LYC interface is initially highly resistive and non-uniform, and the current concentrates at a few locally less resistive spots where Li begins to plate in an inhomogeneous manner. These plated regions then grow into filaments that extend through the LYC layer and first establish soft short-circuit pathways, which eventually develop into a hard short (> 515 h). The concurrent impedance collapse and voltage failure shows that Li metal does not form a uniform, self-limiting interphase as in the Li-Si cell. Instead, its strong reducing nature, together with the high initial interfacial resistance, drives intrusive Li propagation across LYC. 2.5 Full Cell Demonstration of the Dynamic Stability Symmetric cells illustrate how LYC responds to different reducing conditions ( Fig. 5a ). At the Li-In interface, LYC stays within its thermodynamic stability window and the interfacial structure remains essentially unchanged. At intermediate potentials with Li-Si, Li uptake converts only a thin surface layer of LYC into a lithiated phase that acts as a self-limiting redox buffer and keeps the interface kinetically stable. At the most reducing limit with Li metal, lithiation extends much deeper state, which makes the interfacial region more resistive and laterally uneven. The current is then funneled through a few lower-resistance spots, where Li develops filaments and eventually penetrates the electrolyte. Together, these three responses define the anode potential window in which LYC can be used reliably in practical full-cell configurations. Using this interfacial map as a guide, we assembled LYC-based full cells with three anode chemistries and a high-capacity NMC811 (LiNi 0.8 Mn 0.1 Co 0.1 O 2 ) composite cathode. The NMC811/LYC/C composite was prepared by low-energy ball-milling to ensure intimate mixing between the active material and conductors. When paired with a Li-In anode, the cell initially delivers around 180 mAh g -1 at 0.1C, confirming the full-cell architecture can access most of the cathode capacity. This high level of cathode utilization is largely preserved at higher current and the cell provides in average 127.7 mAh g -1 at 0.5C with high average CE of 99 % over 500 cycles ( Fig. 5b ). The Li-In full cell defines a near-equilibrium baseline that represents the thermodynamically compatible regime. However, the trade-off is clear: its operating voltage stays below about 3.7 V, which in turn caps the energy density even though the capacity remains high. Moving into the dynamically stabilized regime, the Li-Si full cell initially delivers a high capacity of 155 mAh g -1 at 0.1C. The net capacity is slightly lower than the Li-In cell, but the lower Si anode potential allows us to extend the upper cutoff voltage to 4.1 V (4.3 V vs. Li/Li + ). Interestingly, the early-cycle specific capacity at 0.33 C starts around 110 mAh g -1 . However, the capacity gradually increases with continued cycling and climbs to over 140 mAh g -1 by about 300 cycles ( Fig. 5c ). The slow activation is consistent with the interfacial reorganization scenario inferred from the Li-Si symmetric cells, where the interface starts with relatively high resistance and then gradually stabilizes as Li redistributes within the lithiated LYC layer. It is also important to note that our Li-Si full cells are operated with a tightly limited N/P capacity ratio. Before assembling full cells, we independently evaluated the practical capacity of the pre-lithiated Si anode and obtained about 300 mAh g -1 ( Supplementary Fig. 16 ). Based on this value, the N/P capacity ratio was set to 1.2. In such a tightly limited N/P configuration, the fact that the cells cycle stably for 500 cycles at 0.33 C with an average CE of about 99% indicates that continuous LYC decomposition is not the dominant process. Instead, it supports the presence of a thin lithiated LYC interphase that acts as a self-limiting redox buffer. At the most reducing limit, Li metal based full cells do not reach similar performance or stability as expected. The initial discharge capacity at 0.1 C is already limited to about 110 mAh g -1 and shows large polarization ( Fig. 5d ). With further cycling at 0.33 C, the cells degrade rapidly and soon display over-charging behavior, which is characteristic of internal short-circuit formation. This outcome is consistent with the deep-lithiation regime, where the interfacial region gains excessive resistance, the current focuses into a few weak spots, and Li filaments penetrate through the electrolyte at the full-cell level. Collectively, the symmetric and full-cell results show that LYC is not governed by a single, fixed stability window set only by thermodynamics. Its effective working range is instead controlled by how far lithiation extends from the anode into LYC and by how the resulting lithiated region balances redox buffering and charge transport kinetics. These trends redefine and extend a practical anode potential window for LYC and point to a broader design principle for redox-active halide electrolytes. Anodes that induce controlled, partial lithiation can support self-limiting interphases and extend the usable voltage range, whereas pushing the electrolyte into deep reduction promotes current localization, filament growth, and eventual failure. 3. Conclusion State-of-the-art solid electrolytes now offer fast Li + transport comparable to liquid electrolytes, yet many still struggle when placed against anodes. In many electrolyte families, entering the reducing regime is widely treated as direct phase decomposition, which narrows viable separator choices. Under this assumption, solid-state cell designs have largely relied on protective interlayers or elevated-potential alloy anodes, rather than freely adopting low-voltage, high-capacity anodes. On the other hand, redox-active halide electrolytes invite a different view of reductive stability. Reduction can proceed through a structurally-intact lithiation process and be reverted on de-lithiation, rather than immediately driving phase decomposition. The key question becomes no longer where reduction starts, but how it evolves. Once paired with a low-potential anode, the same question becomes interfacial. The focus shifts to how the redox response develops at the contact, whether it stays confined and reversible, and how that interfacial evolution changes charge-transport kinetics and long-term stability at the anode interface. In this work, we show that reduction in trivalent halide LYC does not proceed as a single, irreversible failure. When the potential is driven below the nominal reduction onset (0.62 V vs Li/Li + ), a reversible electrochemical response is activated, in which reduction at 0.25 V can be recovered by subsequent re-oxidation near 0.6 V. Such electrochemical response suggests that LYC exhibits intrinsic redox activity beyond its nominal stability limit and that the onset of reduction does not directly imply loss of electrolyte function. The observed reversibility becomes particularly advantageous when LYC is paired with low-potential anodes. Rather than triggering uncontrolled degradation, the reversible reduction allows the interface to accommodate partial lithiation and reorganize into a buffered interfacial state, mitigating continuous electrolyte reduction at the anode contact. The interfacial behavior of LYC depends strongly on the reducing strength of the anode. Li-In anodes operating within the thermodynamic stability window preserve near equilibrium interfacial kinetics. Moderately reducing Li-Si anodes induce partial lithiation that remains spatially confined and self-limiting, allowing the interface to stabilize through redox buffering rather than progressive degradation. In contrast, strongly reducing anodes like pure Li drive excessive lithiation, leading to increased interfacial resistance, lateral heterogeneity, and current focusing that ultimately drive short-circuit failure. Taken together, the results establish that practical operation of trivalent halide LYC is governed by a balance between reversible reduction and anode reducing strength at the interface. Importantly, the present study demonstrates, to our knowledge, that stable electrolyte operation can be achieved beyond the conventional thermodynamic reduction window through interfacially confined and kinetically regulated redox processes. By clarifying how redox reversibility, transport kinetics, and anode potential jointly control interfacial stability, this work extends the viable anode space for halide electrolytes and provides a mechanistic basis for rational anode matching and separator design in halide-based solid-state batteries. 4. Methods Material Synthesis and Basic Characterization Li 3 YCl 6 (LYC) solid electrolyte was synthesized via a mechanochemical route under an argon atmosphere. LiCl (99.995%, Thermo Scientific) and YCl 3 (99.99%, Thermo Scientific) were weighed in stoichiometric ratios to a total mass of 1.5 g and loaded into an airtight 50 ml ZrO 2 grinding jar (Retsch) together with 12 ZrO 2 balls (10 mm diameter, MSE PRO) inside an Ar-filled glovebox. Mechanochemical synthesis was carried out using a high-energy ball mill (Emax, Retsch) operated at 600 rpm for a total net milling time of 36 h, employing a cycle of 5 min milling followed by 1 min rest. High-energy milling was used to drive the solid-state reaction and phase formation of LYC. To ensure homogeneous mixing and prevent local overheating, the milling process was periodically paused, and midway through the synthesis the jar was opened inside the glovebox to manually redistribute powder adhered to the jar walls. After milling, the as-synthesized LYC powder was lightly hand-ground in an agate mortar to break up soft agglomerates and stored in an Ar-filled glovebox (O 2 < 0.2 ppm, VAC) without exposure to ambient air. After synthesis, the powder morphology was characterized by scanning electron microscopy (SEM) using a field-emission SEM (JEOL 7500F) operated at an accelerating voltage of 15 kV. Model Cell Preparation and Testing LYC/C composite powders were prepared by mixing the as-synthesized LYC with carbon black (Super C65, TIMCAL) in an 8:2 weight ratio. A total of 0.2 g of the mixture was loaded together with 6 ZrO 2 balls and ball-milled at 300 rpm for 50 min using the same milling interval and the same high-energy ball mill (Emax, Retsch). The procedure ensured homogeneous dispersion of the conductive additive without inducing additional structural changes in LYC. All composite preparation steps were performed under an Ar atmosphere. Indium (In, Thermo Scientfic, 0.127 mm thick, 99.99%) and Lithium (Li, Sigma Aldrich, 99.9%) metal foils were used as received to prepare Li-In alloy anodes. Lithium and indium were first combined at a Li:In molar ratio of 3:7. The metals were stacked together and mechanically alloyed by repeated roll-and-press processing using stainless steel rods inside an Ar-filled glovebox. The stacked metals were continuously rolled and pressed to promote intimate contact and interdiffusion, and the process was repeated until a visually uniform and mechanically homogeneous Li-In alloy foil was obtained. Cell assembly was carried out by sequential pressing in a pressure-controlled and airtight split coin cell holder (PSC-10N, MTI Corporation). For the SS | LYC | Li-In cell, 100 mg of LYC powder was loaded into the air-tight cell holder and pressed at 300 MPa for 5 min using a hydraulic press (YLJ-15, MTI Corporation). A 10 mm diameter stainless steel rod integrated in the cell holder was used as the current collector on one side, while an 8 mm diameter Li-In alloy foil (0.1 mm thickness) was placed on the opposite side of the electrolyte to complete the cell stack. After assembly, the cell was secured using a compression jig (EQ-JIG-2, MTI Corporation) and operated under a constant stack pressure of approximately 8 MPa, as determined from the applied force and monitored with an integrated pressure sensor. For the LYC/C | LYC | Li–In cell, the electrolyte pellet was prepared following the same procedure and under the same pressing conditions as described above. Subsequently, an LYC/C composite layer with a mass loading of 10 mg was applied onto the electrolyte surface and pressed at 150 MPa for 1 min. An 8 mm diameter Li–In alloy foil (0.1 mm thickness) was then placed on the opposite side of the electrolyte to complete the cell stack. The assembled cell was secured and operated under the same constant stack pressure (~8 MPa) using the same compression jig and pressure monitoring method as described above. Cyclic voltammetry (CV) measurements were performed using the SS | LYC | Li-In cell at a scan rate of 0.1 mV s -1 . For the first CV cycle, the potential sweep was initiated from the open-circuit voltage (OCV) and scanned toward the lower cutoff voltage, where the sweep direction was reversed. From the second cycle onward, CV scans were conducted between a fixed upper cutoff voltage of 2.0 V vs. Li/Li + (1.38 V vs. Li-In) and the designated lower cutoff voltage. The lower cutoff voltage was progressively decreased from 0.62 V to 0 V vs. Li/Li + to probe electrochemical activity beyond the nominal reduction limit of LYC. Multiple consecutive cycles were collected at each cutoff voltage to evaluate the evolution and reversibility of the redox response. Galvanostatic cycling was conducted at current densities of 0.1 and 0.2 mA cm -2 . Cells were cycled between either 0-2.5 V or 0.1-2 V vs. Li/Li⁺, depending on the specific experiment. Coulombic efficiency was calculated from the ratio of charge to discharge capacity for each cycle. All measurements were carried out under a constant stack pressure of approximately 8 MPa. In situ electrochemical impedance spectroscopy (EIS) measurements were performed on the same LYC/C | LYC | Li-In cell. Impedance spectra were collected at selected potentials during galvanostatic cycling over a frequency range of 7 MHz to 0.1 Hz with an AC amplitude of 10 mV. The EIS signal was recorded with six points per decade using a logarithmic scale. The raw impedance data were processed using Python-based DRTtools 40 to perform Distribution of Relaxation Times (DRT) analysis. A first-order derivative regularization (regularization parameter of 0.005) and a FWHM coefficient of 0.5 were applied to stabilize the solution. DRT intensities were first mapped as a function of potential for each state of charge, and the absolute resistance for each contribution was subsequently calculated by integrating the DRT peak over its associated relaxation time range. Galvanostatic Intermittent Titration Technique (GITT) experiments were conducted on the same LYC/C | LYC | Li-In cell. Small current pulses of 0.1 mA cm -2 were applied for 10 min, followed by 10 min relaxation periods to allow the potential to reach a quasi-steady state. Voltages were recorded during the pulses and relaxation periods, and the effective Li⁺ diffusivity was calculated using the established GITT relation, taking into account the slope of the quasi-steady-state voltage and the instantaneous transient voltage after each current pulse. Data processing and diffusivity calculations were performed using Python scripts developed in-house, ensuring consistent analysis across all pulses and cycles. High Energy X-ray Characterization Synchrotron-based structural and spectroscopic characterization was performed on LYC/C electrode materials recovered after electrochemical testing. All samples were prepared and handled in powder form. After cycling, the LYC/C electrode layer was mechanically separated from the cell stack and collected intact, and a portion of the recovered electrode powder was used directly for synchrotron measurements. High-resolution powder X-ray diffraction (XRD) was carried out at Beamline 2-1 of the Stanford Synchrotron Radiation Lightsource (SSRL), SLAC National Accelerator Laboratory. The recovered LYC/C powders were loaded into glass capillaries inside an Ar-filled glovebox and sealed to prevent air exposure. Diffraction data were acquired at room temperature using monochromatic X-rays, with the wavelength specified in the main text. The capillary geometry enabled bulk-sensitive measurements while minimizing preferred orientation effects. The XRD patterns were analyzed using the HighScore Plus software. Background subtraction was first applied using a polynomial fitting function, followed by peak identification based on reference patterns of LYC. Phase analysis was conducted to confirm the absence of secondary crystalline phases within the detection limit of the measurement. Lattice parameters were then calculated from the identified peak positions using least-squares fitting, allowing comparison of structural changes before and after electrochemical cycling. X-ray absorption spectroscopy (XAS) measurements were conducted at SSRL to probe the local electronic structure and coordination environment of Y and Cl in LYC. Y K-edge XAS was performed at Beamline 2-3, and Cl K-edge XAS was collected at Beamline 14-3b. For XAS measurements, the recovered LYC/C powders were transferred into airtight sample holders inside an Ar-filled glovebox. The sealed holders were transported to the beamlines and loaded directly into the measurement chambers without air exposure. All XAS spectra were collected at room temperature in fluorescence mode, with energy calibration performed using reference standards measured concurrently. X-ray absorption spectroscopy (XAS) data were processed using the Athena software package. Energy calibration was performed using simultaneously measured reference foils, followed by normalization and background subtraction using standard pre-edge and post-edge fitting procedures. White-line features at the Y K-edge and Cl K-edge were analyzed by integrating the normalized absorption intensity over a defined energy window to track relative changes in unoccupied electronic states associated with electrochemical cycling. Elemental and Chemical State Analysis The Li content of cycled LYC/C electrode powders was quantified using Inductively Coupled Plasma Optical Emission Spectroscopy (ICP-OES, Agilent 5800 Series). Recovered electrode powders were carefully separated from the cell stack to collect only the active LYC/C material, then digested in a mixture of concentrated hydrochloric acid (HCl, 37 %, ACS Grade, Sigma Aldrich) and nitric acid (HNO 3 , 68-70 vol%, ACS Grade, VWR Chemicals) at 60 °C for 12 h to ensure complete dissolution. The resulting solutions were diluted with deionized water to a final acid concentration of 2 % for measurement. Calibration standards for Li and Y were prepared from commercial stock solutions (1000 mg l -1 in 2 % HNO 3 , Sigma Aldrich) and serially diluted to generate reference points. Emission intensities at selected wavelengths (Li: 274.119, 323.263, 460.289, 610.365, 670.783 nm; Y: 324.228, 360.074, 361.104, 371.029, 377.433 nm) were measured, and calibration curves were constructed by linear regression of standard signals versus known concentrations. Sample emission intensities were converted to concentrations (ppm) using these curves, and Li/Y stoichiometry of each sample was determined by converting concentrations to molar amounts and normalizing to Y content, enabling a direct comparison with electrochemically estimated Li insertion. X-ray Photoelectron Spectroscopy (XPS) measurements were performed using a Thermo-Fisher K-Alpha Plus spectrometer equipped with a monochromatic Al Kα X-ray source (1486 eV). The cycled LYC/C electrodes were carefully recovered from the cell stack, keeping the electrode structure largely intact, and mounted in airtight sample holders inside an Ar-filled glovebox to minimize air exposure before transfer to the spectrometer. All measurements were conducted under high vacuum (<10 -8 Torr). Survey scans were first collected to identify surface species, followed by high-resolution scans of the Y 3d, Cl 2p, C 1s, and Li 1s core levels at a step size of 0.1 eV to monitor changes in elemental valence states during Li insertion and extraction. All XPS spectra were processed using CasaXPS. Background signals were subtracted using a Shirley-type baseline to remove inelastic contributions. High-resolution peaks were fitted with mixed Gaussian-Lorentzian (GL (30)) line shapes. Initial binding energy calibration was performed using the C 1s hydrocarbon peak at 284.8 eV 48 , followed by secondary calibration based on the Cl 2p peak to correct for systematic peak shifts. Quantitative fitting of the Y 3d peaks allowed determination of the relative contributions of Y 3+ and partially reduced Y n+ species, enabling tracking of valence state evolution during Li insertion and extraction. Atomistic Simulations The workflow to identify energy minimized configurations for different x values for convex hull calculations is as follows: first, 10,000 LYC configurations at each x value are initialized with the knowledge of the crystallographic sites of Y and Li and their fractional occupancies, using the CIF file provided by Schlem et al. 33 . For high enough Li values (x > 4), the interstitial tetrahedral sites were considered as the secondary occupation sites for Li in initial configuration generation. To perform greatly expedited molecular dynamics simulations and geometry optimization without compromising the accuracy, MLIP-based calculations were performed with PET-OAM-XL 35 , the accuracy of which for material systems is supported by its current top rank in the Matbench Discovery benchmark 49 . Using PET-OAM-XL, we first performed total energy evaluations for all of the randomly initialized configurations. We then selected 5 structures with the lowest total energies and another 5 structures at random from the rest of the structures for each x value. We then performed NVT ensemble MD simulations for each of the 10 systems with 2 fs timesteps for 100 ps at 400 K. i-PI was used for the MD simulations 50 , and the stochastic velocity rescaling (SVR) thermostat 51 was employed with a tau value of 50 fs. initial, final, and 4 intermediate configurations sampled at regular intervals were taken for geometry optimization in atomic simulation environment (ASE) 52 . Out of the 60 configurations in total, 5 lowest energy configurations were selected for further density functional theory (DFT)-based geometry optimization using the Vienna ab initio Software Package (VASP) v. 6.4.3 53 . In both geometry optimization calculations, maximum force component threshold was set to 0.025 eV Å -1 . The DFT calculations were performed under the projector-augmented-wave (PAW) formalism 54-55 with the Perdew-Burke-Ernzerhof (PBE) functional 56 , and used a kinetic energy cutoff of 600 eV and energy convergence threshold of 1e-6 eV in the self-consistency field (SCF) cycles. Gamma-centered k-point grid was defined with a spacing of 0.25 Å -1 , and electronic occupation was described using Gaussian smearing with a sigma value of 0.05 eV. The convex hull diagram was constructed using the lowest energy configurations after the DFT geometry optimization. RDF analysis was performed using the trajectories from the MLIP-based MD simulations. Symmetric and Full-cell Preparation and Testing For symmetric cells, LYC pellets were first prepared by pressing as-synthesized LYC powder following the same procedure as described above. Li-In alloy anodes were prepared as previously described, with 8 mm diameter discs placed on both sides of the LYC pellet. Li-Si alloy powders were synthesized starting from silicon powder (Si, 325 mesh, 99%, Sigma Aldrich) and lithium metal at a 1:1 molar ratio (Si:Li). Lithium was brought into contact with the silicon powder on a hot plate at 300 °C, where it melted and infiltrated the Si particles. During lithiation, the mixture gradually changed to a lighter color and expanded in volume, forming a partially lithiated Li-Si composite powder. The reaction proceeded for 30 min to ensure thorough lithiation, after which the resulting powder was cooled to room temperature and handled under an Ar atmosphere to prevent oxidation. For each Li-Si symmetric cell, 10 mg of this Li-Si powder was loaded onto both sides of the LYC pellet and pressed at 150 MPa for 1 min per side. Pure Li metal anodes were prepared by rolling as-received Li foil using stainless steel rods inside an Ar-filled glovebox until a uniform thickness of approximately 100 µm was achieved, and then cut into 8 mm diameter discs for placement on both sides of the LYC pellet. All symmetric cells (Li-In, Li-Si, and Li) were assembled in the same pressure-controlled and airtight split coin cell holders (10 mm diameter, PSC-10N, MTI Corporation) and compressed to a stack pressure of 16 MPa using a compression jig (EQ-JIG-2, MTI Corporation). Open-circuit storage impedance was recorded for 1000 min, with Nyquist spectra collected every 10 min (100 cycles) over a frequency range of 7 MHz to 0.1 Hz with a 10 mV AC amplitude. The EIS signal was recorded with six points per decade using a logarithmic scale. Galvanostatic cycling was then performed at 0.25 mA cm -2 with 1 h per half-cycle (plating or stripping), during which in-situ impedance measurements were collected every 100 cycles over the same frequency range with the same AC amplitude. EIS spectra were fitted using EC-lab software to extract interfacial and bulk resistance values. Full cells were assembled using the LYC electrolyte and a high-capacity LiNi 0.8 Mn 0.1 Co 0.1 O 2 single-crystal cathode (sc-NMC811, Targray). The cathode composite was prepared by low-energy ball-milling the sc-NMC811, LYC, and carbon black (Super P C65) in a weight ratio of 58:37:5 using the same ball mill machine (Emax, Retsch) at 300 rpm for 50 min, ensuring homogeneous mixing without inducing structural changes. The cells were assembled using sequential pressing of the following: 100 mg of LYC powder (300 MPa for 5 min, net thickness between 400-450 µm), followed by 15.3 mg cm -2 of cathode composite (150 MPa for 1 min. The cathode active loading is 8.8 mg cm -2 which corresponds to ~1.8 mAh cm -2 . For full cells with Li-In anodes, 10 mm diameter Li-In alloy discs prepared as previously described were placed on the opposite side of the LYC pellet. For full cells with Li-Si anodes, 8 mg of the previously synthesized Li-Si powder (Li:Si = 1:1 molar ratio) was loaded onto the LYC pellet and pressed at 150 MPa for 1 min. For full cells with pure Li metal anodes, 10 mm diameter Li foil discs (100 µm thick) were placed on the other side of the LYC pellet. All cells were assembled in the split coin cell holders (12 mm diameter, ) and a screw compression jig (EQ-YLJ-SP, MTI Corporation) compressed to a stack pressure of 8 MPa. Galvanostatic cycling was conducted in an Ar-filled glovebox (O 2 < 0.2 ppm, VAC) at ambient temperature (20±5 °C), with the current rate calculated based on the total cathode capacity of 200 mAh g -1 . Declarations Author Contributions H.C., S.Y., and G.C. conceived the idea for the project. G.C. supervised the entire research. H.C. and S.Y. performed the electrochemical and X-ray measurements, and analyzed the experimental data. S.C. conducted the computational modeling. P.D. helped with the analysis of absorption and photoelectron measurements. P.P.P. assisted with the analysis of the diffraction measurements. H.C. wrote the initial draft of the manuscript, and G.C. revised the manuscript. All authors discussed the results, contributed to the content, and reviewed the final manuscript. Conflict of Interest The authors declare no conflict of interest Acknowledgements The authors thank Dr. Kevin Stone (BL 2-1), and Dr. Sam Webb (BL 2-3 & BL 14-3b) at the Stanford Synchrotron Radiation Lightsource (SSRL) for their assistance with the high-energy XRD and XAS measurements. Use of SSRL, SLAC National Accelerator Laboratory was supported by the U.S. Department of Energy, Office of Science, Office of Basic Energy Sciences under Contract No. DE-AC02-76SF00515. Studies at the Molecular Foundry were supported by the Office of Science, Office of Basic Energy Sciences, of the U.S. Department of Energy under Contract No. DE-AC02-05CH11231. This work was supported by the Assistant Secretary for Energy Efficiency and Renewable Energy, Office of Vehicle Technologies, of the U.S. Department of Energy under Contract No. DE-AC02-05CH11231. References Janek, J. & Zeier, W. G. A solid future for battery development. Nat. Energy 1 , 16141 (2016). Janek, J. & Zeier, W. G. Challenges in speeding up solid-state battery development. Nat. Energy 8 , 230–240 (2023). Li, X., Liang, J., Yang, X., Adair, K. R., Wang, C., Zhao, F. & Sun, X. Progress and perspectives on halide lithium conductors for all-solid-state lithium batteries. Energy Environ. Sci. 13 , 1429–1461 (2020). Tanaka, Y., Ueno, K., Mizuno, K., Takeuchi, K., Asano, T. & Sakai, A. New oxyhalide solid electrolytes with high lithium ionic conductivity >10 mS cm-1 for all-solid-state batteries. Angew. Chem. Int. Ed. 62 , e202217581 (2023). Asano, T., Sakai, A., Ouchi, S., Sakaida, M., Miyazaki, A. & Hasegawa, S. Solid halide electrolytes with high lithium-ion conductivity for application in all-solid-state batteries. Adv. Mater. 30 , 1803075 (2018). Ginnings, D. C. & Phipps, T. E. Temperature–conductance curves of solid salts. III. Halides of lithium. J. Am. Chem. Soc. 52 , 1340–1345 (1930). Weppner, W. & Huggins, R. A. Ionic conductivity of alkali metal chloroaluminates. Phys. Lett. A 58 , 245–248 (1976). Kamaya, N., Homma, K., Yamakawa, Y., Hirayama, M., Kanno, R., Yonemura, M., Kamiyama, T., Kato, Y., Hama, S., Kawamoto, K. & Mitsui, A. A lithium superionic conductor. Nat. Mater. 10 , 682–686 (2011). Adeli, P., Bazak, J. D., Park, K. H., Kochetkov, I., Huq, A., Goward, G. R. & Nazar, L. F. Boosting solid-state diffusivity and conductivity in lithium superionic argyrodites by halide substitution. Angew. Chem. Int. Ed. 58 , 8681 (2019). Bouguern, M. D., Ningappa, N. G., Vishweswariah, K., Kumar, A. M. R., Kanno, R. & Zaghib, K. Comparative advances in sulfide and halide electrolytes for commercialization of all-solid-state lithium batteries. Adv. Mater. 38 , 2, e13255 (2026). Kim, K., Park, D., Jung, H.-G., Chung, K. Y., Shim, J. H., Wood, B. C. & Yu, S. Material design strategy for halide solid electrolytes Li 3 MX 6 (X = Cl, Br, and I) for all-solid-state high-voltage Li-ion batteries. Chem. Mater. 33 , 3669–3677 (2021). Jun, K., Wei, G., Yang, X., Chen, Y. & Ceder, G. Exploring the soft cradle effect and ionic transport mechanisms in the LiMXCl 4 superionic conductor family. Matter 8 , 102001 (2025). Yu, R., Wang, C., Duan, H., Jiang, M., Zhang, A., Fraser, A., Zuo, J., Wu, Y., Sun, Y., Zhao, Y., Liang, J., Fu, J., Deng, S., Ren, Z., Li, G., Huang, H., Li, R., Chen, N., Wang, J., Li, X., Singh, C. V. & Sun, X. Manipulating charge-transfer kinetics of lithium-rich layered oxide cathodes in halide all-solid-state batteries. Adv. Mater. 35 , 2207234 (2023). Wang, K., Ren, Q., Gu, Z., Duan, C., Wang, J., Zhu, F., Fu, Y., Hao, J., Zhu, J., He, L., Wang, C.-W., Lu, Y., Ma, J. & Ma, C. A cost-effective and humidity-tolerant chloride solid electrolyte for lithium batteries. Nat. Commun. 12 , 4410 (2021). Zhou, L., Zuo, T.-T., Kwok, C. Y., Kim, S. Y., Assoud, A., Zhang, Q., Janek, J. & Nazar, L. F. High areal capacity, long cycle life 4 V ceramic all-solid-state Li-ion batteries enabled by chloride solid electrolytes. Nat. Energy 7 , 83–93 (2022). Jiang, M., Mukherjee, S., Chen, Z. W., Chen, L. X., Li, M. L., Xiao, H. Y., Gao, C. & Singh, C. V. Materials perspective on new lithium chlorides and bromides: insights into thermo-physical properties. Phys. Chem. Chem. Phys. 22 , 22758–22767 (2020). Wang, S., Bai, Q., Nolan, A. M., Liu, Y., Gong, S., Sun, Q. & Mo, Y. Lithium chlorides and bromides as promising solid-state chemistries for fast ion conductors with good electrochemical stability. Angew. Chem. Int. Ed. 58 , 8039–8043 (2019). Kim, J. S., Jung, S., Kwak, H., Han, Y., Kim, S., Lim, J., Lee, Y. M. & Jung, Y. S. Synergistic halide–sulfide hybrid solid electrolytes for Ni-rich cathode design guided by digital twin for all-solid-state Li batteries. Energy Storage Mater. 55 , 193–204 (2023). Kim, S. Y., Cha, H., Kostecki, R. & Chen, G. Composite cathode design for high-energy all-solid-state lithium batteries with long cycle life. ACS Energy Lett. 8 , 521–528 (2023). Yang, S., Kim, S. Y. & Chen, G. Halide superionic conductors for all-solid-state batteries: Effects of synthesis and composition on lithium-ion conductivity. ACS Energy Lett. 9 (5), 2212–2221 (2024). Morino, Y., Aoki, N. & Nakamoto, M. Revealing the reductive decomposition at the interface of Li 3 YCl 6 –lithium metal. ACS Appl. Mater. Interfaces 17 , 45744–45751 (2025). Ji, W., Zheng, D., Zhang, X., Ding, T. & Qu, D. A kinetically stable anode interface for Li 3 YCl 6 -based all-solid-state lithium batteries. J. Mater. Chem. A 9 , 15012–15018 (2021). Hu, L., Zhu, J., Duan, C., Zhu, J., Wang, J., Wang, K., Gu, Z., Xi, Z., Hao, J., Chen, Y., Ma, J., Liu, J.-X. & Ma, C. Revealing the Pnma crystal structure and ion-transport mechanism of the Li 3 YCl 6 solid electrolyte. Cell Rep. Phys. Sci. 4 , 101428 (2023). Liu, Z., Zhang, G., Pepas, J., Ma, Y. & Chen, H. Li 2 FeCl 4 as a cost-effective and durable cathode for solid-state Li-ion batteries. ACS Energy Lett. 9 , 5464-5470 (2024). Zhang, G., Liu, Z., Ma, Y., Pepas, J., Bai, J., Zhong, H., Tang, Y. & Chen, H. Li 2.9 Fe 0.9 Zr 0.1 Cl 6 as redox-active catholyte for solid-state Li-ion batteries. Chem. Mater. 36 , 10104-10112 (2024). Song, Z., Dai, Y., Wang, T., Yu, Q., Ye, X., Wang, L., Zhang, Y., Wang, S. & Luo, W. An active halide catholyte boosts the extra capacity for all-solid-state batteries. Adv. Mater. 36 , 2405277 (2024). Wang, K., Gu, Z., Xi, Z., Hu, L. & Ma, C. Li 3 TiCl 6 as ionic conductive and compressible positive electrode active material for all-solid-state lithium-based batteries. Nat. Commun. 14 , 1396 (2023). Sun, G., Song, Z., Dai, Y., Xiao, Z., Zheng, X. & Luo, W. Redox-active halide catholytes for solid-state lithium batteries. Adv. Sci. 12 , 47, e14215 (2025). Ridley, P., Duong, G., Ko, S. L., Oh, J. A. S., Deysher, G., Griffith, K. J. & Meng, Y. S. Tailoring chloride solid electrolytes for reversible redox. J. Am. Chem. Soc. 147 , 19508–19519 (2025). Schwietert, T. K., Vasileiadis, A. & Wagemaker, M. First-principles prediction of the electrochemical stability and reaction mechanisms of solid-state electrolytes. JACS Au 1 , 1488–1496 (2021). Cheng, Z., Zhao, W., Wang, Q., Zhao, C., Lavrinenko, A. K., Vasileiadis, A., Landgraf, V., Bannenberg, L., Li, Y., Liang, J., Liu, M., Ganapathy, S. & Wagemaker, M. Beneficial redox activity of halide solid electrolytes empowering high-performance anodes in all-solid-state batteries. Nat. Mater. 24 , 1763–1772 (2025). Van der Ven, A., Bhattacharya, J. & Belak, A. A. Understanding Li diffusion in Li-intercalation compounds. Acc. Chem. Res. 46 , 1216–1225 (2013). Schlem, R., Banik, A., Ohno, S., Suard, E. & Zeier, W. G. Insights into the lithium sub-structure of superionic conductors Li 3 YCl 6 and Li 3 YBr 6 . Chem. Mater. 33 , 327–337 (2021). Taminato, S., Yonemura, M., Shiotani, S., Kamiyama, T., Torii, S., Nagao, M., Ishikawa, Y., Mori, K., Fukunaga, T., Onodera, Y., Naka, T., Morishima, M., Ukyo, Y., Adipranoto, D. S., Arai, H., Uchimoto, Y., Ogumi, Z., Suzuki, K., Hirayama, M. & Kanno, R. Real-time observations of lithium battery reactions—operando neutron diffraction analysis during practical operation. Sci. Rep. 6 , 28843 (2016). Bigi, F., Pegolo, P., Mazitov, A. & Ceriotti, M. Pushing the limits of unconstrained machine-learned interatomic potentials. arXiv :2601.16195 (2026). Kawaguchi, T., Fukuda, K. & Matsubara, E. Site- and phase-selective x-ray absorption spectroscopy based on phase-retrieval calculation. J. Phys.: Condens. Matter 29 , 113002 (2017). Wang, X., Yang, Z., Li, N., Wu, K., Gao, K., Zhao, E., Han, S. & Guo, W. Influence mechanism of interfacial oxidation of Li 3 YCl 6 solid electrolyte on reduction potential. Chem. Eur. J. 30 , e202303884 (2024). Vasquez, R. P. XPS core level spectra of some Y compounds. J. Electron Spectrosc. Relat. Phenom. 50 , 167–170 (1990). Wu, J.-F. & Guo, X. Origin of the low grain boundary conductivity in lithium ion conducting perovskites: Li 3x La 0.67−x TiO 3 . Phys. Chem. Chem. Phys. 19, 5880–5886 (2017). Wan, T. H., Saccoccio, M., Chen, C. & Ciucci, F. Influence of the discretization methods on the distribution of relaxation times deconvolution: Implementing radial basis functions with DRTtools. Electrochim. Acta 184 , 483–499 (2015). Weppner, W. & Huggins, R. A. Determination of the kinetic parameters of mixed-conducting electrodes and application to the system Li 3 Sb. J. Electrochem. Soc. 124 , 1569–1578 (1977). Deiss, E. Spurious chemical diffusion coefficients of Li+ in electrode materials evaluated with GITT. Electrochim. Acta 50 , 2927–2932 (2005). Li, J., Su, H., Liu, Y., Zhong, Y., Wang, X. & Tu, J. Li alloys in all solid-state lithium batteries: A review of fundamentals and applications. Electrochem. Energy Rev. 7 , 18 (2024). Irvine, J. T. S., Sinclair, D. C. & West, A. R. Electroceramics: Characterization by impedance spectroscopy. Adv. Mater. 2 , 132–138 (1990). Drvarič Talian, S., Brutti, S., Navarra, M. A., Moškon, J. & Gaberšček, M. Impedance spectroscopy applied to lithium battery materials: Good practices in measurements and analyses. Energy Storage Mater. 69 , 103413 (2024). Zhang, S. S., Xu, K. & Jow, T. R. EIS study on the formation of solid electrolyte interface in Li-ion battery. Electrochim. Acta 51 , 1636–1640 (2006). Lewis, J. A., Cavallaro, K. A., Liu, Y. & McDowell, M. T. The promise of alloy anodes for solid-state batteries. Joule 6 , 1418–1430 (2022). Chu, H., Noh, H., Kim, Y.-J., Yuk, S., Lee, J.-H., Lee, J., Kwack, H., Kim, Y., Yang, D.-K. & Kim, H.-T. Achieving three-dimensional lithium sulfide growth in lithium-sulfur batteries using high-donor-number anions. Nat. Commun. 10 , 188 (2019). Riebesell, J., Goodall, R. E. A., Benner, P., Chiang, Y., Deng, B., Ceder, G., Asta, M., Lee, A. A., Jain, A. & Persson, K. A. A framework to evaluate machine learning crystal stability predictions. Nat. Mach. Intell. 7 , 836–847 (2025). Litman, Y., Kapil, V., Feldman, Y. M. Y., Tisi, D., Begušić, T., Fidanyan, K., Fraux, G., Higer, J., Kellner, M., Li, T. E., Pós, E. S., Stocco, E., Trenins, G., Hirshberg, B., Rossi, M. & Ceriotti, M. i-PI 3.0: A flexible and efficient framework for advanced atomistic simulations. J. Chem. Phys. 161 , 062504 (2024). Bussi, G., Donadio, D. & Parrinello, M. Canonical sampling through velocity rescaling. J. Chem. Phys. 126 , 014101 (2007). Larsen, A. H., Mortensen, J. J., Blomqvist, J., Castelli, I. E., Christensen, R., Dułak, M., Friis, J., Groves, M. N., Hammer, B. & Hargus, C. The atomic simulation environment—a Python library for working with atoms. J. Phys. Condens. Matter 29 , 273002 (2017). Kresse, G. & Furthmüller, J. Efficient iterative schemes for ab initio total-energy calculations using a plane-wave basis set. Phys. Rev. B 54 , 11169-11186 (1996). Blöchl, P. E. Projector augmented-wave method. Phys. Rev. B 50 , 17953 (1994). Kresse, G. & Joubert, D. From ultrasoft pseudopotentials to the projector augmented-wave method. Phys. Rev. B 59 , 1758 (1999). Perdew, J. P., Burke, K. & Ernzerhof, M. Generalized gradient approximation made simple. Phys. Rev. Lett. 77 , 3865 (1996). Additional Declarations There is NO Competing Interest. Supplementary Files LxYCSI.docx Supporting information Cite Share Download PDF Status: Under Review Version 1 posted You are reading this latest preprint version Research Square lets you share your work early, gain feedback from the community, and start making changes to your manuscript prior to peer review in a journal. As a division of Research Square Company, we’re committed to making research communication faster, fairer, and more useful. We do this by developing innovative software and high quality services for the global research community. Our growing team is made up of researchers and industry professionals working together to solve the most critical problems facing scientific publishing. Also discoverable on Platform About Our Team In Review Editorial Policies Advisory Board Help Center Resources Author Services Accessibility API Access RSS feed Manage Cookie Preferences © Research Square 2026 | ISSN 2693-5015 (online) Privacy Policy Terms of Service Do Not Sell My Personal Information {"props":{"pageProps":{"initialData":{"identity":"rs-8928802","acceptedTermsAndConditions":true,"allowDirectSubmit":false,"archivedVersions":[],"articleType":"Article","associatedPublications":[],"authors":[{"id":603857014,"identity":"5605b413-dadd-4d8c-9b5b-5b860b59708e","order_by":0,"name":"Guoying Chen","email":"data:image/png;base64,iVBORw0KGgoAAAANSUhEUgAAAZAAAAAyAQMAAABI0h/eAAAABlBMVEX///8AAABVwtN+AAAACXBIWXMAAA7EAAAOxAGVKw4bAAAAqUlEQVRIiWNgGAWjYFACHiCusEjgI1HLGYkENtK0MLaRokXe/ezBx4XzJPLYGNgvPuYhRovhmbxk45nbJIrZGHiKjYnTMoPHTJp3m0RiGwNPmuQMIrWY/+adQ4oWeQkeM2beBpAW9mMSH4jRYsCTYyzNcwzoF2YeZgOitMi3nzH8zFNjk8fP3v7wQQJRthyAsZh5DIjRALSlAc5kf0CcllEwCkbBKBhxAACSOifqw3gKRQAAAABJRU5ErkJggg==","orcid":"https://orcid.org/0000-0002-3218-2609","institution":"Lawrence Berkeley National Laboratory","correspondingAuthor":true,"prefix":"","firstName":"Guoying","middleName":"","lastName":"Chen","suffix":""},{"id":603857015,"identity":"6fafbdff-355c-4ff9-abd1-6f3b8140f5a5","order_by":1,"name":"Hyunwon Chu","email":"","orcid":"","institution":"Lawrence Berkeley National Laboratory","correspondingAuthor":false,"prefix":"","firstName":"Hyunwon","middleName":"","lastName":"Chu","suffix":""},{"id":603857016,"identity":"db5b88f4-7374-4f09-b165-9c672ffd86a5","order_by":2,"name":"Shuhao Yang","email":"","orcid":"","institution":"Lawrence Berkeley National Laboratory","correspondingAuthor":false,"prefix":"","firstName":"Shuhao","middleName":"","lastName":"Yang","suffix":""},{"id":603857017,"identity":"1489df0c-02fa-4bca-940f-055e03ff3d9f","order_by":3,"name":"Sanggyu Chong","email":"","orcid":"https://orcid.org/0000-0002-6948-1602","institution":"Korea Advanced Institute of Science and Technology (KAIST)","correspondingAuthor":false,"prefix":"","firstName":"Sanggyu","middleName":"","lastName":"Chong","suffix":""},{"id":603857018,"identity":"5cd6e7e5-ba6c-46a7-986f-34ec478c3af4","order_by":4,"name":"Pravin Didwal","email":"","orcid":"","institution":"Lawrence Berkeley National Labratory","correspondingAuthor":false,"prefix":"","firstName":"Pravin","middleName":"","lastName":"Didwal","suffix":""},{"id":603857019,"identity":"8b415f07-5005-4199-aafb-63ef14ef339d","order_by":5,"name":"Partha Paul","email":"","orcid":"","institution":"University of Manchester","correspondingAuthor":false,"prefix":"","firstName":"Partha","middleName":"","lastName":"Paul","suffix":""}],"badges":[],"createdAt":"2026-02-20 20:00:46","currentVersionCode":1,"declarations":"","doi":"10.21203/rs.3.rs-8928802/v1","doiUrl":"https://doi.org/10.21203/rs.3.rs-8928802/v1","draftVersion":[],"editorialEvents":[],"editorialNote":"","failedWorkflow":false,"files":[{"id":105567044,"identity":"3e51299b-95f5-445c-988d-434e52a9d431","added_by":"auto","created_at":"2026-03-27 12:58:08","extension":"png","order_by":1,"title":"Figure 1","display":"","copyAsset":false,"role":"figure","size":354358,"visible":true,"origin":"","legend":"\u003cp\u003e\u003cstrong\u003eRedox Activity of LYC Within and Beyond its Nominal Reduction Limit. a, \u003c/strong\u003eCyclic Voltammetry (CV) of SS | LYC | Li-In cells collected at a scan rate of 0.1 mV s\u003csup\u003e-1\u003c/sup\u003e while progressively lowering the cutoff voltage from 0.62 V to 0 V vs Li/Li\u003csup\u003e+\u003c/sup\u003e. The thermodynamic stability window of LYC is shaded in green, and the extended reducing range is shaded in yellow. The low-voltage reduction feature is marked as R1 and the corresponding oxidation feature as O1. CV scans from the 1st to 5th cycles are shown, with full-cycle data provided in the Supplementary Information. \u003cstrong\u003eb, \u003c/strong\u003eGalvanostatic discharge and charge profiles of LYC/C composite electrodes measured at 0.1 mA cm\u003csup\u003e-2\u003c/sup\u003e with a lower cutoff of 0 V. Specific discharge and charge capacities for each cycle are shown in the inset. \u003cstrong\u003ec, \u003c/strong\u003eComparison between Li stoichiometry changes estimated from electrochemical capacity and Li/Y ratios measured by Inductively Coupled Plasma-Optical Emission Spectroscopy (ICP-OES) for samples collected after selected discharge and charge states. \u003cstrong\u003ed, \u003c/strong\u003eVoltage profile evolution under extended cycling at 0.2 mA cm\u003csup\u003e-2\u003c/sup\u003e (left), along with long-term cycling of LYC/C composite electrodes at 0.1 mA cm\u003csup\u003e-2\u003c/sup\u003e (top right) and 0.2 mA cm\u003csup\u003e-2\u003c/sup\u003e (bottom right) with a lower cutoff of 0.1 V. Average coulombic efficiencies (CE) for each condition are shown in the insets.\u003cbr\u003e\n\u003c/p\u003e","description":"","filename":"1.png","url":"https://assets-eu.researchsquare.com/files/rs-8928802/v1/9347b2ef2a8d5c38e843c7dc.png"},{"id":105566968,"identity":"4af154e4-df83-4ec4-8d65-06234f4c5219","added_by":"auto","created_at":"2026-03-27 12:57:52","extension":"png","order_by":2,"title":"Figure 2","display":"","copyAsset":false,"role":"figure","size":626694,"visible":true,"origin":"","legend":"\u003cp\u003e\u003cstrong\u003eStructural and Chemical State Evolution of LYC During Reversible Li Exchange. a, \u003c/strong\u003eMicroscopy image of the secondary-particle morphology of LYC (left) and its crystal structure (trigonal, space group P3̅m1) showing each elemental sites at the assigned Wyckoff positions (middle). Li sites are rendered as pink and orange spheres, Y as blue spheres, and Cl as gray spheres, with partial occupancies depicted by partially filled circles. A zoomed-in view (right) highlights the local coordination environments of the Li(6h), Li(6g), and tetrahedral vacant sites. \u003cstrong\u003eb, \u003c/strong\u003eCycling curve of LYC/C powders (left) and corresponding synchrotron X-ray Diffraction (XRD) patterns collected at selected voltages (right). The 2θ axis is plotted using the synchrotron wavelength. The region between 12° and 17° in 2θ is magnified to display peak-position changes, and reflections corresponding to each electrochemical state are labeled. \u003cstrong\u003ec, \u003c/strong\u003eY K-edge X-ray Absorption Spectroscopy (XAS) spectra showing the near-edge region and extended fine structure oscillations for different electrochemical states. A zoomed near-edge plot compares half-maximum edge positions across charge and discharge states. Cl K-edge XAS spectra are shown for selected states together with reference spectra of LiCl and YCl\u003csub\u003e3\u003c/sub\u003e. \u003cstrong\u003ed, \u003c/strong\u003eY 3d and Cl 2p X-ray Photoelectron Spectroscopy (XPS) core-level spectra for pristine, discharged, and charged samples. The spectra show the Y 3d\u003csub\u003e5/2\u003c/sub\u003e and Y 3d\u003csub\u003e3/2\u003c/sub\u003e components and the Cl 2p doublet, with fitted envelopes provided for each state. The relative populations of the Y\u003csup\u003e3+\u003c/sup\u003e and partially reduced Y\u003csup\u003en+\u003c/sup\u003e components are calculated from integrated peak areas and shown in the insets. The XPS survey scan and Li 1s spectra are separately provided in Supplementary Information.\u003c/p\u003e","description":"","filename":"2.png","url":"https://assets-eu.researchsquare.com/files/rs-8928802/v1/1ce50177b8bf28a479d552ba.png"},{"id":105532013,"identity":"a42a31a7-da49-4c22-9682-5f7902d1ac1d","added_by":"auto","created_at":"2026-03-27 06:14:00","extension":"png","order_by":3,"title":"Figure 3","display":"","copyAsset":false,"role":"figure","size":434102,"visible":true,"origin":"","legend":"\u003cp\u003e\u003cstrong\u003eIon Transport Kinetics of LYC During Lithiation and De-lithiation. a, \u003c/strong\u003eIn-situ Electrochemical Impedance Spectroscopy (EIS) collected during discharge and charge of an LYC/C | LYC | Li-In cell operated at 0.1 mA cm\u003csup\u003e-2\u003c/sup\u003e. Nyquist spectra are shown for selected voltages, and the frequency corresponding to the top of the semicircle at 1 V (0.64 Hz) is tracked across all states. The associated voltage profiles are provided in the Supplementary Information. \u003cstrong\u003eb, \u003c/strong\u003eDistribution of Relaxation Times (DRT) map generated from the impedance spectra across the full voltage window. The mapped intensities are plotted as a function of relaxation time (x-axis) and potential (y-axis), with dashed lines marking the P1 and P2 features. Raw DRT spectra for each voltage step are included in the Supplementary Information. \u003cstrong\u003ec, \u003c/strong\u003eGalvanostatic Intermittent Titration Technique (GITT) results collected using 0.1 mA cm\u003csup\u003e-2\u003c/sup\u003e current pulses applied for 10 mins followed by 10-min rest periods. The figure shows the full GITT voltage trace (black), the extracted OCV profile formed by connecting steady-state voltages (gray dashed line), and the bias overpotential formed by connecting transient voltages (blue line). The corresponding effective Li\u003csup\u003e+\u003c/sup\u003e diffusivity calculated from each pulse is plotted on the right axis. Numbered points in the figure indicate specific states during cycling: (1) pristine, (2) intermediate stage of lithiation, (3) fully lithiated, (4) intermediate stage of de-lithiation, and (5) fully de-lithiated. Steady-state potential changes (ΔE\u003csub\u003es\u003c/sub\u003e), bias potential transients (ΔE\u003csub\u003et\u003c/sub\u003e), and IR-drop values used for the calculation are provided in the Supplementary Information.\u003c/p\u003e","description":"","filename":"3.png","url":"https://assets-eu.researchsquare.com/files/rs-8928802/v1/34c77f3904e35c5b104dd18f.png"},{"id":105567449,"identity":"6564ef43-29f0-493b-8202-acc75b7e526a","added_by":"auto","created_at":"2026-03-27 12:59:37","extension":"png","order_by":4,"title":"Figure 4","display":"","copyAsset":false,"role":"figure","size":456172,"visible":true,"origin":"","legend":"\u003cp\u003e\u003cstrong\u003eKinetic Stability of LYC with Different Low-voltage Anodes. a,\u003c/strong\u003e Schematic comparison between an inert solid electrolyte (SE, left) and a redox-active SE (right). The inert SE schematic illustrates reductive decomposition when operated below its allowed voltage window. The redox-active SE schematic shows lithiation occurring at the anode surface, which effectively extends the operable voltage window. \u003cstrong\u003eb, \u003c/strong\u003eOperating potentials and specific capacities of representative alloy anodes plotted relative to the thermodynamic limit of LYC (gray dotted horizontal line). Each dot displays the potential ranges for Li-alloying materials and Li metal together with their theoretical capacities after \u003csup\u003e43,47\u003c/sup\u003e. The star mark is applied to the Li-In anode because only about half of its theoretical capacity falls within the voltage range that matches the thermodynamic stability of LYC. \u003cstrong\u003ec, \u003c/strong\u003eImpedance evolution during 1000 min of open-circuit storage for symmetric M | LYC | M cells (M = Li-In, Li-Si, Li). Nyquist spectra collected every 10 min are overlaid. Frequency regions corresponding to the LYC electrolyte contribution (LYC), the M/LYC interfacial process (Int.), and the alloy electrode process (Elec.) are annotated. \u003cstrong\u003ed, \u003c/strong\u003eGalvanostatic cycling curves for the three symmetric cells operated at 0.25 mA cm\u003csup\u003e-2\u003c/sup\u003e for 1 h per half-cycle under a constant stack pressure of 16 MPa. Voltage profiles are magnified at 515 h to compare cycling behavior and to highlight the short-circuit event in the Li cell at the corresponding time. An accompanying illustration is included to depict the lithiation level evolution at the interface. \u003cstrong\u003ee, \u003c/strong\u003eSelected Nyquist plots at 100 h and 900 h of cycling for the Li-In, Li-Si, and Li cells. The top-of-semicircle frequencies associated with each impedance contribution are marked in the insets. Nyquist spectra for all intermediate time points are provided in the Supplementary Information, along with detailed fitting results and the corresponding equivalent-circuit diagrams for all spectra.\u003c/p\u003e","description":"","filename":"4.png","url":"https://assets-eu.researchsquare.com/files/rs-8928802/v1/82d4e4a3a067c59a1d7cb979.png"},{"id":105566848,"identity":"2cfd419a-234c-40cf-9f5b-4538c89ab220","added_by":"auto","created_at":"2026-03-27 12:57:32","extension":"png","order_by":5,"title":"Figure 5","display":"","copyAsset":false,"role":"figure","size":568524,"visible":true,"origin":"","legend":"\u003cp\u003e\u003cstrong\u003eFull-cell Performance of LYC Paired with Low-voltage Anodes. a, \u003c/strong\u003eSchematic of the LYC-based full-cell configuration using NMC cathode composites and alloy anodes (left). Illustration of the expected interfacial response of LYC when paired with Li-In, Li-Si, and Li metal electrodes (right). The interfacial diagrams outline the behavior of LYC under each low-voltage anode and depict possible evolution pathways during cycling. \u003cstrong\u003eb,\u003c/strong\u003e Cycling data of NMC811 | LYC | Li-In full cells. Voltage profiles from the first five 0.1C cycles (left) and representative 0.5C cycles shown every 100 cycles (middle). Cycling performance at 0.5C is plotted with discharge capacities, charge capacities, and Coulombic efficiencies (CE, right), and the average CE over 500 cycles is provided in the inset. \u003cstrong\u003ec,\u003c/strong\u003e Cycling data of NMC811 | LYC | Li-Si full cells. The first five 0.1C voltage profiles are shown (left), and the 0.33C cycling voltage profiles are displayed at 100-cycle intervals (middle). Cycling performance at 0.33C is plotted with discharge and charge capacities along with the corresponding CE (right), and the average CE over 500 cycles is provided in the inset. \u003cstrong\u003ed,\u003c/strong\u003e Cycling data of NMC811 | LYC | Li full cells. Voltage profiles for the initial five cycles at 0.1C are presented (left), and the 0.33C voltage traces are shown at 5-cycle increments (right). Cycling performance at 0.33C includes plotted discharge and charge capacities together with the corresponding CE until failure.\u003c/p\u003e","description":"","filename":"5.png","url":"https://assets-eu.researchsquare.com/files/rs-8928802/v1/bb9562624bd03694a937c0a5.png"},{"id":105570346,"identity":"4cbab8b8-2e3d-4ed9-aa41-53bed117e439","added_by":"auto","created_at":"2026-03-27 13:16:13","extension":"pdf","order_by":0,"title":"","display":"","copyAsset":false,"role":"manuscript-pdf","size":3309221,"visible":true,"origin":"","legend":"","description":"","filename":"manuscript.pdf","url":"https://assets-eu.researchsquare.com/files/rs-8928802/v1/7cb7962f-8edf-4501-b747-c5a5c086885a.pdf"},{"id":105566822,"identity":"cc460fb6-cffc-46c8-8ae3-136685c2a9d2","added_by":"auto","created_at":"2026-03-27 12:57:27","extension":"docx","order_by":1,"title":"","display":"","copyAsset":false,"role":"supplement","size":4082450,"visible":true,"origin":"","legend":"Supporting information","description":"","filename":"LxYCSI.docx","url":"https://assets-eu.researchsquare.com/files/rs-8928802/v1/a961d0185a6e97bbe02896c3.docx"}],"financialInterests":"There is \u003cb\u003eNO\u003c/b\u003e Competing Interest.","formattedTitle":"Interfacial Redox Buffering Stabilizes Halide Solid Electrolytes Against Low-Potential Anodes","fulltext":[{"header":"1. Introduction","content":"\u003cp\u003eThe search for solid electrolytes (SEs) that combine high ionic conductivity and stable electrode interfaces remains a central challenge in solid-state battery (SSB) research\u003csup\u003e1-2\u003c/sup\u003e. Recent advances in halide electrolytes have begun to address both requirements simultaneously, as a new generation of chloride or bromide-based materials has been shown to exhibit room-temperature ionic conductivities around 1 mS cm\u003csup\u003e-1\u003c/sup\u003e \u003csup\u003e3\u003c/sup\u003e and, in some cases, exceeding 10 mS cm\u003csup\u003e-1\u003c/sup\u003e \u003csup\u003e4\u003c/sup\u003e, while maintaining strong compatibility with high-voltage (HV) lithium-ion battery (LIB) cathodes\u003csup\u003e5\u003c/sup\u003e. For many years, halide compounds were explored as Li-ion conductors\u003csup\u003e6-7\u003c/sup\u003e but attracted limited attention because their ionic conductivities lagged behind those of sulfide-based electrolytes\u003csup\u003e8-9\u003c/sup\u003e. This perception has shifted only recently with the discovery of highly conductive halide electrolytes, repositioning halides from a niche materials class to prime candidates for practical all-solid-state batteries\u003csup\u003e10\u003c/sup\u003e. Beyond their high ionic conductivity, halide electrolytes offer a distinct advantage in oxidative stability around 4 V vs. Li/Li\u003csup\u003e+\u003c/sup\u003e \u003csup\u003e11,12\u003c/sup\u003e, based on the intrinsic robustness of the halide anion framework. Compared to oxide or sulfide counterparts, the chloride sublattice lacks easily oxidizable species, minimizing parasitic interfacial reactions at high-voltage oxide cathodes\u003csup\u003e13\u003c/sup\u003e and enabling simplified composite cathode designs without protective layers\u003csup\u003e5,14-15\u003c/sup\u003e.\u003c/p\u003e\n\u003cp\u003eWithin this family, Li\u003csub\u003e3\u003c/sub\u003eYCl\u003csub\u003e6\u003c/sub\u003e (LYC) has been widely investigated as a promising electrolyte for all-solid-state battery designs targeting high energy density. LYC offers high ionic conductivity (0.51 mS cm\u003csup\u003e-1\u003c/sup\u003e at room temperature\u003csup\u003e5\u003c/sup\u003e), good processability (Poisson’s ratio of 0.274\u003csup\u003e16\u003c/sup\u003e), and excellent oxidation stability (up to 4.21 V vs. Li/Li\u003csup\u003e+ 17\u003c/sup\u003e). These attributes have driven extensive studies of LYC-based full cells in combination with conventional LIB cathodes and even high voltage electrodes\u003csup\u003e13,18-19\u003c/sup\u003e, where the electrolyte remains largely chemically inert. However, LYC, like other metal halide electrolytes, has often been regarded as intrinsically unstable when paired with low-voltage anodes. Thermodynamic calculations predict that LYC enters a thermodynamically reducing regime below 0.62 V vs Li/Li\u003csup\u003e+ 17\u003c/sup\u003e, a trend that is consistent with electrochemical observations\u003csup\u003e20-21\u003c/sup\u003e. Such behavior has been taken as indicative of phase decomposition driven by reduction of the cationic sublattice. Specifically, reduction of Y\u003csup\u003e3+\u003c/sup\u003e is predicted to promote the formation of reduced yttrium chlorides (Y\u003csub\u003e2\u003c/sub\u003eCl\u003csub\u003e3\u003c/sub\u003e) and, under sufficiently reducing conditions, metallic Y (Y\u003csup\u003e0\u003c/sup\u003e) accompanied by LiCl-rich phases\u003csup\u003e17\u003c/sup\u003e. Importantly, the decomposition-based interpretation is not unique to LYC that most transition-metal halides are assumed to always decompose into electronically conductive and ionically blocking products beyond its thermodynamic limit\u003csup\u003e12\u003c/sup\u003e. Because of this prevailing assumption, the low-potential behavior of halide electrolytes has rarely been examined systematically, except a few studies focused on direct Li-metal contact\u003csup\u003e21-22\u003c/sup\u003e. Instead, most LYC-based solid-state batteries have been designed to avoid any reductive conditions below its intrinsic threshold, mostly by pairing the electrolyte with elevated-potential anodes such as Li-In alloys\u003csup\u003e5,19,23\u003c/sup\u003e. While this strategy ensures stable operation within the thermodynamic window, the high anode potential substantially narrows the accessible cell voltage (\u0026lt; 4 V for Li-In full cells)\u003csup\u003e5\u003c/sup\u003e and limits the achievable energy density. While adding a sulfide-based interlayer has been demonstrated with some success, the complexity of adding another layer limit its practical use. As such, the perceived reduction instability of halide electrolytes has effectively acted as an inherent design constraint.\u003c/p\u003e\n\u003cp\u003eAgainst this long-standing view, recent works have started to re-examine whether reduction of halide electrolytes must necessarily result in catastrophic breakdown. A growing body of literature reported that redox activity within halide frameworks can proceed in a reversible and non-destructive manner when mediated by multivalent cations such as Fe (Li\u003csub\u003e2\u003c/sub\u003eFeCl\u003csub\u003e4\u003c/sub\u003e\u003csup\u003e24\u003c/sup\u003e and Li\u003csub\u003e3\u003c/sub\u003eFeCl\u003csub\u003e6\u003c/sub\u003e\u003csup\u003e25\u003c/sup\u003e), V (Li\u003csub\u003e3\u003c/sub\u003eVCl\u003csub\u003e6\u003c/sub\u003e\u003csup\u003e26\u003c/sup\u003e), and Ti (Li\u003csub\u003e3\u003c/sub\u003eTiCl\u003csub\u003e6\u003c/sub\u003e\u003csup\u003e27\u003c/sup\u003e). In these systems, reduction does not immediately disrupt the host structure but instead occurs through well-defined redox processes that preserve the integrity of the halide lattice\u003csup\u003e28\u003c/sup\u003e. Such reversible redox behavior has been demonstrated most clearly on the cathode side, where transition-metal redox in the halide framework falls within the typical operating potential window of oxide cathodes (i.e., Li\u003csub\u003e2\u003c/sub\u003eFeCl\u003csub\u003e4\u003c/sub\u003e at ~3.6 V vs. Li/Li\u003csup\u003e+\u003c/sup\u003e \u003csup\u003e24\u003c/sup\u003e). The dynamic redox activity of the halide electrolyte contributes an additional 20-50% reversible capacity to cathodes while maintaining structural integrity under extended cycling conditions\u003csup\u003e28\u003c/sup\u003e. At a broader level, similar behavior has also been reported in Na\u003csup\u003e+\u003c/sup\u003e-conducting halide electrolytes\u003csup\u003e29\u003c/sup\u003e, further underscoring that reduction does not universally induce irreversible decomposition of halide-based electrolytes. While redox behavior has primarily been explored to extend cathode capacity, a few studies have newly extended this redox-based view to anode-relevant voltages using lower-redox-potential cations\u003csup\u003e30\u003c/sup\u003e. For instance, in mixed-halide Li\u003csub\u003e3\u003c/sub\u003eYCl\u003csub\u003e3\u003c/sub\u003eBr\u003csub\u003e3\u003c/sub\u003e (LYCB), stable lithiation/de-lithiation has been observed below the nominal stability window (\u0026lt; 0.6 V) driven by reversible Y redox\u003csup\u003e31\u003c/sup\u003e. When incorporated into anodic composites, the low-voltage redox activity enables extended cycling of red phosphorus (RP) anodes\u003csup\u003e31\u003c/sup\u003e. Nevertheless, in these anode “composite” configurations, the redox activity remains confined within the electrode region and is leveraged mainly for anode material activation and stabilization. A key remaining question is how the same redox behavior manifests when the halide electrolyte itself serves as a stand-alone separator and directly interfaces with a low-potential anode in a practical solid-state battery configuration.\u003c/p\u003e\n\u003cp\u003eIn this work, we show that LYC is not limited by its thermodynamic stability window, but instead supports reversible Li\u003csup\u003e+\u003c/sup\u003e insertion/extraction that expands its operating range toward low-potential anodes. When the potential is pushed below the nominal reduction limit, we observe that a reversible reduction-oxidation response emerges in LYC framework with the main lithiation feature centered at 0.25 V vs. Li/Li\u003csup\u003e+\u003c/sup\u003e. After the initial activation, the process becomes highly reversible and delivers a stable cycling capacity of ~ 150 mAh g\u003csup\u003e-1\u003c/sup\u003e with a high Coulombic efficiency (CE) over 99 %. Synchrotron-based characterization across selected discharge/charge states verifies that the Li uptake and release proceed without loss of framework integrity. Computational modeling further confirms the structural stability of the accessible lithiation states and elucidates the local coordination changes during lithiation. Benefiting from LYC’s intrinsic redox, we next evaluate whether LYC can function as a stand-alone separator against low-potential anodes. \u003cem\u003eIn-situ\u003c/em\u003e tracking of charge-transport kinetics indicates that lithiation of LYC forms a self-limiting interphase that buffers further reduction. Separator-level application, however, requires careful control over an operating lithiation range, since deeper lithiation introduces excessive ionic polarization. Based on this mechanistic understanding, anode-side behavior of LYC was compared across three representative electrode materials spanning progressively lower potentials: Li-In alloy (within the thermodynamic window), Li-Si alloy (partial lithiation), and Li metal (deep lithiation). Access to the reversible Li-exchange regime enables stable operation of halide-based Si full cells (LiNi\u003csub\u003e0.8\u003c/sub\u003eMn\u003csub\u003e0.1\u003c/sub\u003eCo\u003csub\u003e0.1\u003c/sub\u003eO\u003csub\u003e2\u003c/sub\u003e (NMC) | LYC | Li-Si) for over 500 cycles without any protective layers. This protection-free compatibility directly allows the lower Si anode potential to be utilized so that the cell operating voltage can be extended up to 4.1 V with NMC electrodes (compared to 3.7 V for Li-In cells). In contrast, Li metal drives excessive lithiation at the interface, creating a highly resistive contact that promotes current localization and rapid failure, defining the practical limit of redox buffering that can be utilized in LYC-based cells. Overall, the results establish reversible Li exchange as a practical lever to extend halide separators beyond their nominal reduction limits, opening a viable route to pair LYC with low-potential, high-capacity anodes in coating-free full cells without sacrificing long-term stability.\u003cstrong\u003e\u003cbr clear=\"all\"\u003e\u0026nbsp;\u003c/strong\u003e\u003c/p\u003e"},{"header":"2. Results \u0026 Discussion","content":"\u003ch2\u003e2.1 Reversible Li Exchange Beyond Thermodynamic Limit\u003c/h2\u003e\n\u003cp\u003eBefore examining the low-voltage electrochemical behavior of the LYC electrolyte, we first confirmed the structural and morphological uniformity of as-synthesized powder to ensure that the observed redox features originate inherently. The LYC electrolyte was prepared by high-energy ball milling, and the resulting particle morphology consists of sub-micron primary particles that aggregate into secondary clusters extending to several tens of microns (\u003cstrong\u003eSupplementary Fig. 1\u003c/strong\u003e), which is typical of halide electrolytes prepared by mechanochemical synthesis routes\u003csup\u003e20\u003c/sup\u003e. The X-ray Diffraction (XRD) pattern displays reflections consistent with the targeted LYC phase and does not show detectable impurity peaks. The samples\u0026rsquo; homogeneity is essential because any redox activity discussed later should be attributed to the intrinsic LYC rather than to secondary phases or processing byproducts.\u003c/p\u003e\n\u003cp\u003eAfter confirming the phase purity, we performed cyclic voltammetry (CV) at progressively lower cutoff voltages to determine whether LYC engages in a reversible electrochemical process (\u003cstrong\u003eFig. 1a\u003c/strong\u003e). A Stainless Steel (SS) | LYC | Lithium-Indium (Li-In) cell was used, where the Li-In alloy prevents direct chemical reactivity at the counter electrode and the SS plate minimizes capacitive artifacts in the working-electrode response. LYC is predicted to have a nominal reduction limit at 0.62 V vs. Li/Li\u003csup\u003e+\u003c/sup\u003e, associated with the Y\u003csup\u003e3+\u003c/sup\u003e/Y\u003csup\u003e0\u003c/sup\u003e redox potential\u003csup\u003e17\u003c/sup\u003e. Potentials below this range are shown by theory to drive thermodynamically favorable reduction reactions\u003csup\u003e17,21\u003c/sup\u003e, although the extent to which these reactions proceed under operating conditions remains an experimental question. When the lower cutoff is held at 0.62 V, the current response remains essentially featureless, as expected within its thermodynamic stability window. As the cutoff is gradually lowered, however, a new reduction feature starts to develop. At 0.2 V cutoff, a distinct cathodic peak (R1) becomes clearly visible, and extending the cutoff to 0 V sharpens this feature into a well-defined peak centered at 0.25 V. The systematic emergence of the peak only under deeper reducing conditions suggests activation of an additional electrochemical process beyond the thermodynamic window.\u003c/p\u003e\n\u003cp\u003eReversing the potential scan provides further insight into whether LYC participates in a reversible electrochemical process. Interestingly, we observed an oxidation peak (O1) emerges near 0.7 V on the return sweep, forming a redox pair with the low-voltage reduction (R1) feature. In the 1st cycle, the CE ranges from 30 to 45 % depending upon the cutoff voltages, implying that a portion of the initial reduction introduces irreversible processes. However, the efficiency increases steeply with continued cycling, surpassing 90 % by the 2nd cycle and approaching 99 % by the 20th CV cycle (\u003cstrong\u003eSupplementary Fig. 2\u003c/strong\u003e). The rapid improvement supports that, after an initial conditioning step, LYC begins to follow a predominantly reversible redox pathway rather than sustained, cumulative side reactions dominating the response. To additionally verify whether the observed redox activity comes from LYC itself rather than other electrode components, we also varied cell configurations to include or exclude conductive carbon (C) additives (\u003cstrong\u003eSupplementary Fig. 3\u003c/strong\u003e). Regardless of configuration, the same reduction feature near 0.3 V and oxidation feature near 0.7 V were observed. The reproducibility across different electrode setups supports that the dominant redox features originate from LYC rather than the current collector or carbon additive.\u003c/p\u003e\n\u003cp\u003eSuch redox behavior becomes more evident under galvanostatic cycling, which allows a more quantitative assessment of its magnitude and reversibility. For these measurements, the cell configuration was slightly adjusted to utilize an LYC/C composite electrode so that the active mass could be precisely defined for capacity determination. Under a constant-current discharge to 0 V, the LYC/C composite delivers an initial capacity of 162 mAh g\u003csup\u003e-1\u003c/sup\u003e (\u003cstrong\u003eFig. 1b\u003c/strong\u003e). The subsequent first-charge capacity reaches 113 mAh g\u003csup\u003e-1\u003c/sup\u003e, corresponding to an initial CE of about 70 %. The lower efficiency reflects the initial irreversible portion of the reaction that was also observed in the CV sweeps. With continued cycling, however, the efficiency rapidly increases and exceeds 97 % by the 10th cycle (\u003cstrong\u003eSupplementary Fig. 4\u003c/strong\u003e). The main contributors to the high capacity can be more clearly identified when the voltage profiles are converted to differential capacity (dQ/dE) plots (\u003cstrong\u003eSupplementary Fig. 4\u003c/strong\u003e). During reduction, two distinct processes appear: a prominent peak near 0.4 V and a second rise below 0.1 V. The higher-voltage peak closely aligns with the R1 reduction feature observed in the CV scans, whereas the low-voltage contribution near 0 V likely reflects a decomposition-related process that lowers CE during the early galvanostatic cycles. This secondary pathway is far less apparent in the CV measurements, likely because the SS electrode offers limited interfacial contact area. In contrast, the composite electrode used for galvanostatic operation presents a much larger active surface, which accentuates the surface-related decomposition contribution. Even with this additional feature, CV and galvanostatic data together show that once LYC passes through its initial chemical adjustment, the predominant electrochemical response settles into a highly reversible regime when the lower cutoff is limited to avoid the most reducing potentials.\u003c/p\u003e\n\u003cp\u003eThe voltage profile of the LYC/C composite under galvanostatic discharge shows a broad sloping region often linked to solid-solution-type insertion in Li-ion electrode materials\u003csup\u003e32\u003c/sup\u003e. The similar response motivated us to examine whether LYC undergoes measurable changes in Li stoichiometry through reversible lithiation and de-lithiation. To evaluate dynamic Li-ion exchange in LYC, we first quantified the number of electrons transferred from the measured capacity and converted the value into the equivalent amount of inserted Li\u003csup\u003e+\u003c/sup\u003e. The electrochemically estimated Li content was then compared with the Li/Y ratio obtained from Inductively Coupled Plasma-Optical Emission Spectroscopy (ICP-OES) analysis of cycled samples (\u003cstrong\u003eFig. 1c\u003c/strong\u003e). The initial discharge capacity of 162 mAh g\u003csup\u003e-1\u003c/sup\u003e corresponds to an estimated 1.95 mol of inserted Li\u003csup\u003e+\u003c/sup\u003e per mol of LYC. When the discharged sample was analyzed, the ICP measurement gives the Li/Y ratio of 5.09. Assuming no Y loss during cycling, the ratio corresponds to a LYC stoichiometry of Li\u003csub\u003e5.09\u003c/sub\u003eYCl\u003csub\u003e6\u003c/sub\u003e. The measured composition is in close agreement with the expected value of Li\u003csub\u003e4.95\u003c/sub\u003eYCl\u003csub\u003e6\u003c/sub\u003e derived from electrochemical insertion of ~1.95 mol of Li. Furthermore, the Li content decreases upon the 1st charge, and the subsequent discharge step re-introduces additional Li⁺ in a manner that closely follows the capacity-based Li estimate. Overall, the electrochemical capacity aligns with reversible changes in Li stoichiometry, indicating that the redox response originates from Li-exchange within the LYC framework.\u003c/p\u003e\n\u003cp\u003eThe reversible Li exchange within the LYC framework remains stable in multiple insertion and extraction cycles. Since deep discharge near 0 V led to a current rise that likely initiated minor decomposition and a small CE loss, we narrowed the operation potential window and set the lower cutoff at 0.1 V. As the majority of the capacity arises from the plateau region between 0.2 and 0.3 V, the discharge capacity of the LYC/C composite stayed around 150 mAh g\u003csup\u003e-1\u003c/sup\u003e while maintaining high reversibility over 50 cycles with an average CE of 99.63 % (\u003cstrong\u003eFig. 1d\u003c/strong\u003e). Cycling at higher current resulted in some capacity reduction due to increased polarization, yet the cell still delivered about 110 mAh g\u003csup\u003e-1\u003c/sup\u003e with an average CE near 99 %. The sustained efficiency shows that the activated electrochemical process continues to operate in a reversible manner well beyond LYC\u0026rsquo;s thermodynamic stability window. The extended reversibility of lithiation and de-lithiation demonstrates that LYC repeatedly accommodates Li\u003csup\u003e+\u003c/sup\u003e without loss of function, prompting a detailed examination of how its lattice and phase stability evolve as Li content varies.\u003c/p\u003e\n\u003ch2\u003e2.2 Structural Integrity and Redox-Driven Chemical Changes\u003c/h2\u003e\n\u003cp\u003eThe crystal structure of Li\u003csub\u003e3\u003c/sub\u003eYCl\u003csub\u003e6\u003c/sub\u003e (trigonal, space group \u003cem\u003eP3̅m1\u003c/em\u003e) shows that Li⁺ ions intrinsically occupy the Wyckoff 6g and 6h positions with occupancies of 100 and 50 %, respectively\u003csup\u003e33\u003c/sup\u003e. A simple stoichiometric consideration indicates that these two intrinsic Li sites can accommodate an additional 1.5 mol of Li\u003csup\u003e+\u003c/sup\u003e, allowing lithiation up to approximately Li\u003csub\u003e4.5\u003c/sub\u003eYCl\u003csub\u003e6\u003c/sub\u003e. Prior studies have also reported that LYC contains tetrahedral voids that can serve as additional insertion sites\u003csup\u003e5,33\u003c/sup\u003e (\u003cstrong\u003eFig. 2a\u003c/strong\u003e). Structurally, therefore, LYC possesses sufficient Li-accessible sites to support dynamic Li exchange without requiring a phase transformation.\u003c/p\u003e\n\u003cp\u003eTo examine how the LYC lattice responds during reversible Li exchange, we recovered LYC/C powders cycled to different voltages and performed synchrotron XRD to examine structural evolution. Across all states of charge, the major reflections match exclusively to the LYC phase, and no crystalline decomposition products such as Y metal, LiCl, or YCl\u003csub\u003e3\u003c/sub\u003e are detected within the XRD detection limits (\u003cstrong\u003eSupplementary Fig. 5\u003c/strong\u003e). More importantly, the diffraction patterns exhibit clear and systematic peak shifts as a function of the electrochemical state (\u003cstrong\u003eFig. 2b\u003c/strong\u003e). In pristine LYC, three representative reflections appear in the 2\u0026theta; range of 12-17\u0026deg;, corresponding to the (002), (301), and (112) planes. During the 1st discharge, no noticeable peak shift was observed up to 0.5 V, which aligns with the electrochemical response where lithiation begins near 0.3 V. Upon subsequent discharge to 0 V, however, the (112) reflection shifts to higher 2\u0026theta; from 15.72\u0026deg; to 16.22\u0026deg;, and comparable shifts occur for the (002) and (301) peaks. The shift to higher 2\u0026theta; is consistent with a decrease in d-spacing (lattice contraction). Using the Bragg\u0026rsquo;s equation, we find that the c-axis lattice parameter decreases from 6.07 \u0026Aring; to 5.93 \u0026Aring; during the 1st discharge, bringing roughly 2 % shrinkage in unit-cell volume (\u003cstrong\u003eSupplementary Fig. 6\u003c/strong\u003e). The trend largely mirrors the early-stage lithiation behavior of layered oxide cathode materials such as LiNi\u003csub\u003ex\u003c/sub\u003eMn\u003csub\u003ey\u003c/sub\u003eCo\u003csub\u003e1-x-y\u003c/sub\u003eO\u003csub\u003e2\u003c/sub\u003e (NMC), where lithiation induces c-axis contraction from reduced repulsion between neighboring anion planes\u003csup\u003e34\u003c/sup\u003e. Notably, the contraction is fully reversible upon de-lithiation. When charged back to 1 V, the diffraction peaks of LYC/C return to lower 2\u0026theta; and restore the lattice parameters close to their pristine values. Such reversible lattice breathing persists over the initial five cycles we examined. The structural stability is further corroborated by atomistic simulation with Machine-Learning Interatomic Potential (MLIP, PET-OAM-XL\u003csup\u003e35\u003c/sup\u003e) and Density Functional Theory (DFT) calculations, where structural relaxations confirm that the LYC networks are reasonably retained across the lithiation range of Li\u003csub\u003e3\u003c/sub\u003e to Li\u003csub\u003e5\u003c/sub\u003e (\u003cstrong\u003eSupplementary Fig. 7\u003c/strong\u003e). Throughout all conditions examined, both spectroscopic and computational analyses show no evidence of significant phase decomposition. reassuring that the LYC framework remains structurally coherent and resilient during repeated lithiation and de-lithiation.\u003c/p\u003e\n\u003cp\u003eBuilding on this bulk-scale structural analysis, we then investigated whether the local bonding environment exhibits the same level of stability using X-ray Absorption Spectroscopy (XAS). XAS data are commonly interpreted by separating the near-edge (XANES) region, which reflects oxidation-state and electronic changes, from the extended fine structure (EXAFS) region, which captures the local coordination and bond environment\u003csup\u003e21,36\u003c/sup\u003e. When comparing the EXAFS region of the Y K-edge across different states of charge, the spectra overlap almost completely as the local coordination around Y remains largely unchanged during discharge and charge (\u003cstrong\u003eFig. 2c\u003c/strong\u003e). In contrast, the XANES region captures clear signatures of Y-centered redox activity. The half-maximum edge position shifts by 1.1 eV toward lower energy upon 1st discharge (1D, from 17043.6 to 17042.5 eV), showing that Y\u003csup\u003e3+\u003c/sup\u003e in pristine LYC is partially reduced during lithiation. In the subsequent charging step (1C), the edge returns to its initial position, a clear indication that the Y valence state is restored based on the highly reversible redox process. The similar shift pattern recurs over multiple cycles and continues through at least the 10th discharge. Independent analysis of the white-line peak maximum shows a comparable energy evolution as the half-maximum edge positions (\u003cstrong\u003eSupplementary Fig. 8\u003c/strong\u003e): the peak moves to lower energy during discharge and shifts back to higher energy upon charge. Comparison with reference compounds YCl\u003csub\u003e3\u003c/sub\u003e (Y\u003csup\u003e3+\u003c/sup\u003e) and Y\u003csub\u003e2\u003c/sub\u003eCl\u003csub\u003e3\u003c/sub\u003e (Y\u003csup\u003e1.5+\u003c/sup\u003e) suggests that the discharged state approaches an average Y valence near +1, consistent with the ICP-derived stoichiometry of Li\u003csub\u003e5.09\u003c/sub\u003eYCl\u003csub\u003e6\u003c/sub\u003e (Y\u003csup\u003e0.91+\u003c/sup\u003e). We note that a quantitative oxidation-state estimate depends on reference selection and fitting model. These XAS results, from both EXAFS and XANES regions, confirm that LYC undergoes a controlled Y-centered redox process rather than a decomposition-driven transition. The Cl K-edge spectra also support this understanding. Across all states, the Cl K-edge data retain its overall line shape and edge position, resembling a stable mixture of Li-Cl and Y-Cl coordination motifs. Complementing the XAS measurements, radial distribution function (RDF) analysis of Molecular Dynamics (MD) simulations with the MLIP model confirms that the local coordination environments of Li-Cl and Y-Cl experience no severe changes across varying lithiation levels (\u003cstrong\u003eSupplementary Fig. 9\u003c/strong\u003e). The absence of coordination changes indicates that chloride ions do not participate in the redox process and that the halide sublattice remains intact even under deep lithiation.\u003c/p\u003e\n\u003cp\u003eX-ray Photoelectron Spectroscopy (XPS) measurements further monitor the surface evolution of elemental valence states during Li exchange (\u003cstrong\u003eFig. 2d\u003c/strong\u003e). In the Y 3d core-level spectra, pristine LYC/C shows only Y\u003csup\u003e3+\u003c/sup\u003e presence with the Y 3d\u003csub\u003e5/2\u003c/sub\u003e peak centered around 159.2 eV. This value is consistent with reported binding energies for trivalent Y in halide environments\u003csup\u003e37\u003c/sup\u003e. After the 1st discharge, an additional component emerges at lower binding energy near 157.7 eV, which indicates the formation of partially reduced Y states (Y\u003csup\u003en+\u003c/sup\u003e, n \u0026lt; 3, exact peak position provided in \u003cstrong\u003eSupplementary Fig. 10\u003c/strong\u003e). Reference binding energies place metallic Y\u003csup\u003e0\u003c/sup\u003e near 155.8 eV and intermediate-valence Y\u003csup\u003e1.5+\u003c/sup\u003e species (such as Y\u003csub\u003e2\u003c/sub\u003eS\u003csub\u003e3\u003c/sub\u003e) around 157.4 eV\u003csup\u003e38\u003c/sup\u003e. The new feature at the discharged states remains well above the energy range characteristic of Y\u003csup\u003e0\u003c/sup\u003e, suggesting that reduction in LYC does not progress all the way toward metal formation. Instead, its position aligns with a partially reduced Y state expected within the lithiated framework. Quantitative fitting shows that the Y\u003csup\u003e3+\u003c/sup\u003e to Y\u003csup\u003en+\u003c/sup\u003e ratio shifts to roughly 30:70 after the 1st discharge. Upon charging, the ratio partially recovers to 61:39, showing substantial reversible reoxidation. During the 2nd discharge, the reduced component increases again in a reproducible manner, matching the cycle-to-cycle evolution and reversible Y-centered redox behavior observed in XAS. Throughout all states of charge, the Cl 2p spectra remain essentially unchanged in both binding energy and line shape, reaffirming that the chlorine sublattice maintains chemical stability during cycling. Together, these structural and spectroscopic results cohesively support that LYC accommodates Li\u003csup\u003e+\u003c/sup\u003e insertion through reversible changes in its Y-centered electronic states while preserving the integrity of the halide framework.\u003c/p\u003e\n\u003ch2\u003e2.3 Ion Transport Kinetics Across Lithiation States\u003c/h2\u003e\n\u003cp\u003eAlthough the structural framework remains intact, lithiation inevitably reduces the population of vacant Li sites in the LYC lattice, which can be expected to lower ionic conductivity as Li content increases. Computational Li-Y RDF profiles reveal that lithium atoms increasingly occupy shorter-distance sites starting around the Li\u003csub\u003e4\u003c/sub\u003e level (\u003cstrong\u003eSupplementary Fig. 9\u003c/strong\u003e), indicating the progressive filling of tetrahedral interstitials that normally function as conduction pathways\u003csup\u003e17\u003c/sup\u003e. Understanding how ionic transport evolves under progressive lithiation is therefore important for describing the electrolyte\u0026rsquo;s behavior when in contact with low-voltage anodes. To trace how transport kinetics change, we collected the \u003cem\u003ein-situ\u003c/em\u003e Electrochemical Impedance Spectroscopy (EIS) response of an LYC/C | LYC | Li-In cell across different lithiation states (\u003cstrong\u003eFig. 3a\u003c/strong\u003e).\u003c/p\u003e\n\u003cp\u003eDuring discharge, we observe that the impedance shows a systematic rise in the low-frequency region. The response is dominated by a semicircle centered around 0.64 Hz, which reflects ion-transport processes in the LYC/C composite\u003csup\u003e31\u003c/sup\u003e. Once the potential enters the main lithiation regime below 0.3 V, the impedance begins to increase and continues to grow toward 0 V. The same evolution appears in the Bode plot and in the single-frequency transients collected at 0.64 Hz (\u003cstrong\u003eSupplementary Fig. 11\u003c/strong\u003e). The resistance increase during lithiation can be interpreted through progressive filling of Li vacancies. As fewer vacancies become available, the number of mobile carriers declines and Li\u003csup\u003e+\u003c/sup\u003e transport through the composite becomes increasingly hindered at deep lithiation states. It is also worth noting that the impedance of LYC/C increases rather than decreases. In systems where an ionic conductor develops substantial electronic conductivity, the added electronic pathway often lowers the overall impedance\u003csup\u003e39\u003c/sup\u003e. Here, however, the increase in resistance implies that any electronic conduction that may develop does not compensate for the loss of Li\u003csup\u003e+\u003c/sup\u003e carriers. Therefore, the charge transport in lithiated LYC is still primarily dominated by ionic motion and does not show evidence of an electronically percolating pathway that dominates the impedance response over the examined window. On the other hand, as the cell enters the charge step, the low-frequency impedance begins to decrease once the potential rises above roughly 0.5 V. The resistance drop reflects the reopening of vacancy sites that were filled at high Li content. By the time the cell approaches 1 V, the semicircle associated with the LYC/C contribution returns to values similar to the pristine state, showing a near-complete recovery of optimal ionic transport pathways.\u003c/p\u003e\n\u003cp\u003eDistribution of Relaxation Times (DRT) analysis provides a clearer quantitative picture of how the impedance changes arise\u003csup\u003e40\u003c/sup\u003e. In the pristine LYC/C cell, the DRT spectrum contains three distinguishable peaks (\u003cstrong\u003eSupplementary Fig. 12\u003c/strong\u003e). As discharge proceeds, these features reorganize into two dominant peaks: a short time-constant (\u0026tau;) component near 10\u003csup\u003e-6\u003c/sup\u003e s associated with bulk electrolyte conduction (P1), and a long time-constant contribution near 10\u003csup\u003e0\u003c/sup\u003e s that reflects ionic transport in the LYC/C electrode (P2). To visualize how the kinetic signatures evolve across the full voltage range, the extracted DRT intensities (\u0026gamma;) were mapped as a function of potential (\u003cstrong\u003eFig. 3b\u003c/strong\u003e). With increasing lithiation, we find that the amplitude of the long-\u0026tau; P2 intensity gradually increases more than 30 times (from 17.4 to 649 \u0026Omega;), indicating a substantial growth in resistance originating from the LYC/C composite. Integrating this peak yields the corresponding absolute resistance, which grows markedly from 68.8 to 1944 \u0026Omega;. During de-lithiation, however, the P2 peak intensity rapidly collapses toward its pristine magnitude once the potential exceeds 0.5 V. In contrast to the stepwise impedance buildup observed during lithiation, the recovery occurs in a single sharp transition, implying that Li\u003csup\u003e+\u003c/sup\u003e extraction from LYC does not experience the same kinetic constraints that limit insertion. A similar evolution and recovery pattern repeats over multiple cycles as the impedance response follows a consistent Li\u003csup\u003e+\u003c/sup\u003e insertion-extraction pathway (\u003cstrong\u003eSupplementary Fig. 13\u003c/strong\u003e).\u003c/p\u003e\n\u003cp\u003eThe kinetic asymmetry between lithiation and de-lithiation can be well captured by evaluating the effective Li\u003csup\u003e+\u003c/sup\u003e diffusivity using Galvanostatic Intermittent Titration Technique (GITT, \u003cstrong\u003eFig. 3c\u003c/strong\u003e). In this method, small current pulses introduce incremental Li⁺ and the subsequent relaxation isolates kinetic overpotentials from steady-state changes\u003csup\u003e41-42\u003c/sup\u003e, allowing the effective diffusivity to be determined for insertion and extraction (detailed step provided in \u003cstrong\u003eSupplementary Fig. 14\u003c/strong\u003e). Under lithiation, the effective Li⁺ diffusivity starts on the order of 10\u003csup\u003e-8\u003c/sup\u003e cm\u003csup\u003e2\u003c/sup\u003e s\u003csup\u003e-1\u003c/sup\u003e (marked as Stage 1) but gradually falls by nearly three orders of magnitude to 10\u003csup\u003e-11\u003c/sup\u003e cm\u003csup\u003e2\u003c/sup\u003e s\u003csup\u003e-1\u003c/sup\u003e around 0 V (Stage 3). Over this range, the open-circuit voltage (OCV) changes are minor, while the bias overpotential grows strongly as kinetic barrier for further insertion increases. By contrast, the effective diffusivity stays near 10\u003csup\u003e-8\u003c/sup\u003e to 10\u003csup\u003e-9\u003c/sup\u003e cm\u003csup\u003e2\u003c/sup\u003e s\u003csup\u003e-1\u003c/sup\u003e through most of the charge process (Stage 4) and decreases only near the end of charge (Stage 5), where the OCV shifts as extractable Li gets depleted. Therefore, the kinetic penalty for removing Li is far smaller over most of the composition range than for inserting Li into an already heavily lithiated framework. One mechanistic interpretation is that lithiation proceeds heterogeneously, forming a Li-rich region near particle surfaces while the interior remains relatively Li-poor (Stage 2). As lithiation advances, incoming Li\u003csup\u003e+\u003c/sup\u003e must traverse this increasingly Li-filled region, where ionic transport is reduced due to fewer available vacancies/interstitial pathways, thereby amplifying concentration gradients and polarization. This transport contrast can promote an effective \u0026lsquo;shrinking-core\u0026rsquo;-like behavior even though the bulk voltage response appears largely continuous. Upon de-lithiation, Li can be removed from the Li-rich near-surface region first (Stage 4), rapidly reopening percolating diffusion pathways and restoring lower resistance.\u003c/p\u003e\n\u003cp\u003eTaking the electrochemical observations together with spectroscopy data, it becomes clear that LYC can incorporate additional Li\u003csup\u003e+\u003c/sup\u003e without major structural changes, while maintaining its predominant ionic conduction. Such a response suggests that lithiation does not readily progress into a significant phase decomposition, nor does it trigger a transition to a mixed ionic-electronic state. The outcome fundamentally differs from the common perception based on the theory that LYC should fully reduce to metallic Y beyond its thermodynamic limit, which would form an electronically conductive phase driving short-circuiting. Instead, because the lithiated LYC preserves its ionic dominance with minimal electronic leakage, it can form an interphase that is kinetically self-limiting under these conditions, analogous in function (though not necessarily composition) to a Solid Electrolyte Interphase (SEI) in liquid electrolyte systems. By suppressing further decomposition, this mode of reduction provides a mechanistic basis for enabling LYC to stably interface with low-voltage anodes beyond its nominal thermodynamic limit without undergoing runaway breakdown.\u003c/p\u003e\n\u003ch2\u003e2.4 Dynamic Redox Activities Enabling Low-voltage Anodes\u003c/h2\u003e\n\u003cp\u003eLYC is formally predicted to be unstable below its reduction onset of 0.62 V vs. Li/Li\u003csup\u003e+ 17\u003c/sup\u003e. Most prior demonstrations therefore paired LYC with a Li-In anode at the cell level, whose main operating potential lies near the same voltage range\u003csup\u003e5,19,23\u003c/sup\u003e. However, the elevated anode potential sacrifices a large portion of cell operation voltage window and significantly limits the achievable energy density. On the other hand, our findings cohesively direct that the electrolyte can take up additional Li through a dynamic redox. Reversible lithiation enables LYC to kinetically buffer a reducing environment without decomposition and to operate against low-voltage anodes beyond its nominal threshold (\u003cstrong\u003eFig. 4a\u003c/strong\u003e). When the operational window of LYC can be extended, several attractive low-voltage, high-capacity anode chemistries become accessible (\u003cstrong\u003eFig. 4b\u003c/strong\u003e). For instance, Li-Ag alloy delivers a capacity (670 mAh g\u003csup\u003e-1\u003c/sup\u003e) comparable to Li-In while offering a broader voltage span (0.175 V vs. Li/Li\u003csup\u003e+ 43\u003c/sup\u003e). Alloy anode such as Li-Mg further enlarges both specific capacity (2150 mAh g\u003csup\u003e-1\u003c/sup\u003e) and operating voltage (0.03 V\u003csup\u003e43\u003c/sup\u003e), which effectively boost the energy density of LYC-based systems. However, it should be noted that ionic-conductivity loss becomes substantial at deep lithiation. An anode that is too reducing may still impose large interfacial resistance and introduce polarization that compromises performance. These considerations highlight the need to select an anode whose operation potential maintains a balanced and manageable interfacial kinetics.\u003c/p\u003e\n\u003cp\u003eTo evaluate how LYC behaves when paired with different anodes, we selected three representative electrode materials spanning progressively lower potentials: Li-In, operating within LYC\u0026rsquo;s nominal stability around 0.6 V; Li-Si, which provides a moderately reducing environment between 0.2 to 0.4 V; and Li metal, presenting the strongest reducing condition at 0 V vs. Li/Li\u003csup\u003e+\u003c/sup\u003e. Symmetric cells (M | LYC | M) were assembled using each alloy (M) and characterized under both storage and cycling conditions to determine which anodes can realistically operate with LYC. Under open-circuit voltage (OCV) storage, the Li-In cell exhibits the most stable response based on its thermodynamic compatibility (\u003cstrong\u003eFig. 4c\u003c/strong\u003e). Its Nyquist impedance spectra resolve into three distinct semi-circles, attributed to ion conduction within LYC (\u0026gt; 10\u003csup\u003e6\u003c/sup\u003e Hz), ionic transport across the Li/LYC interface (~ 10\u003csup\u003e4\u003c/sup\u003e Hz), and a diffusive process in the alloy electrode (\u0026lt; 10\u003csup\u003e2\u003c/sup\u003e Hz), respectively\u003csup\u003e44\u003c/sup\u003e. Over 1000 min of storage, no additional features appear and the overall magnitude remains constant, assuring that Li-In anode does not introduce any reducing drive toward LYC. The Li-Si cell also maintains a relatively stable impedance profile during storage. Both the magnitude and spectral shape remain unchanged except for the low-frequency electrode region, suggesting that exposure to a Li-Si anode does not trigger reduction or the formation of a resistive interphase during that period. In contrast, the Li metal cell behaves differently from the outset. Its initial interfacial impedance at mid-frequency is substantially higher than that of Li-In and Li-Si, which originates from strong lithiation at the immediate interface forming a highly resistive region. Interestingly, the interfacial resistance then reaches a steady value after a mild increase. The absence of continued impedance growth may indicate that the reaction remains confined to the interfacial vicinity rather than extending into the bulk electrolyte. Under galvanostatic cycling, the contrasts among the three systems sharpen (\u003cstrong\u003eFig. 4d\u003c/strong\u003e). The thermodynamically-stable Li-In symmetric cell maintains long-term operation with minimal polarization buildup over 1000 hours. In contrast, the Li-Si cell shows a noticeable rise in polarization during the first ~100 h, followed by a gradual relaxation. By around 300 h, the cell settles into a lower and more stable voltage window and then maintains the magnitude during prolonged cycling over 1000 hours. The Li cell initially follows a similar trend to Li-Si, and the polarization increases over the first few tens of hours. However, the cell voltage collapses abruptly within the first 100 h, consistent with internal short-circuit formation. The cell eventually reaches a complete hard short at around 515 h, as evidenced by the voltage dropping to near-zero level without recovery.\u003c/p\u003e\n\u003cp\u003eTo identify the origin of the polarization, we collected EIS spectra in parallel with cycling (\u003cstrong\u003eFig. 4e\u003c/strong\u003e). Equivalent-circuit fitting separated the Li/LYC interfacial resistance (R\u003csub\u003eint\u003c/sub\u003e) and allowed us to track its evolution (fitting details in \u003cstrong\u003eSupplementary Fig. 15\u003c/strong\u003e). In the Li-In cell, R\u003csub\u003eint\u003c/sub\u003e starts at 29.1 \u0026Omega; and remains in a narrow range, decreasing slightly to 20.2 \u0026Omega; after 900 h of cycling. The nearly constant interfacial resistance indicates that no additional reaction pathways are activated and that the interfacial chemistry remains preserved over extended cycling. The Li-Si cell shows a more complex impedance response. As cycling begins, an additional intermediate-frequency contribution (~ 10\u003csup\u003e5\u003c/sup\u003e Hz) newly appears between the electrolyte and interfacial resistances and rises to 101 \u0026Omega; at 100 h. With continued cycling, the feature gradually diminishes and the overall impedance stabilizes at a lower level. Impedance elements in this frequency range are commonly assigned to thin resistive layers forming at electrode-electrolyte contacts, often called as \u0026ldquo;surface film\u0026rdquo; resistance\u003csup\u003e44-45\u003c/sup\u003e. In liquid systems, the contribution corresponds to SEI-related components\u003csup\u003e46\u003c/sup\u003e. By analogy, we infer that the impedance contribution observed here can be interpreted as a resistive lithiation layer that develops at the LYC/Li-Si interface. In the early stage of cycling, LYC near the Li-Si contact is rapidly driven into a highly lithiated state, which appears as an increased film-type resistance. With repeated lithiation and de-lithiation, continued Li exchange gradually redistributes this highly lithiated region over a larger thickness, reducing its local blocking effect (\u003cstrong\u003eFig. 4d\u003c/strong\u003e). In line with this picture, the film resistance (R\u003csub\u003efilm\u003c/sub\u003e) decreases from 101 \u0026Omega; to 39.6 \u0026Omega; by 300 h and remains largely unchanged thereafter. More importantly, the bulk impedance of LYC shows no significant change over time as the lithiation process stays confined to a near-surface region. This localized redox buffering preserves the bulk electrolyte and enables stable galvanostatic cycling of the Li-Si anode with LYC for more than 1000 h. Compared to the other two cells, the Li metal cell already shows a strong drop in both bulk and interfacial resistances within the first 100 h. The rapid decrease may originate since the Li/LYC interface is initially highly resistive and non-uniform, and the current concentrates at a few locally less resistive spots where Li begins to plate in an inhomogeneous manner. These plated regions then grow into filaments that extend through the LYC layer and first establish soft short-circuit pathways, which eventually develop into a hard short (\u0026gt; 515 h). The concurrent impedance collapse and voltage failure shows that Li metal does not form a uniform, self-limiting interphase as in the Li-Si cell. Instead, its strong reducing nature, together with the high initial interfacial resistance, drives intrusive Li propagation across LYC.\u003c/p\u003e\n\u003ch2\u003e2.5 Full Cell Demonstration of the Dynamic Stability\u003c/h2\u003e\n\u003cp\u003eSymmetric cells illustrate how LYC responds to different reducing conditions (\u003cstrong\u003eFig. 5a\u003c/strong\u003e). At the Li-In interface, LYC stays within its thermodynamic stability window and the interfacial structure remains essentially unchanged. At intermediate potentials with Li-Si, Li uptake converts only a thin surface layer of LYC into a lithiated phase that acts as a self-limiting redox buffer and keeps the interface kinetically stable. At the most reducing limit with Li metal, lithiation extends much deeper state, which makes the interfacial region more resistive and laterally uneven. The current is then funneled through a few lower-resistance spots, where Li develops filaments and eventually penetrates the electrolyte. Together, these three responses define the anode potential window in which LYC can be used reliably in practical full-cell configurations.\u003c/p\u003e\n\u003cp\u003eUsing this interfacial map as a guide, we assembled LYC-based full cells with three anode chemistries and a high-capacity NMC811 (LiNi\u003csub\u003e0.8\u003c/sub\u003eMn\u003csub\u003e0.1\u003c/sub\u003eCo\u003csub\u003e0.1\u003c/sub\u003eO\u003csub\u003e2\u003c/sub\u003e) composite cathode. The NMC811/LYC/C composite was prepared by low-energy ball-milling to ensure intimate mixing between the active material and conductors. When paired with a Li-In anode, the cell initially delivers around 180 mAh g\u003csup\u003e-1\u003c/sup\u003e at 0.1C, confirming the full-cell architecture can access most of the cathode capacity. This high level of cathode utilization is largely preserved at higher current and the cell provides in average 127.7 mAh g\u003csup\u003e-1\u003c/sup\u003e at 0.5C with high average CE of 99 % over 500 cycles (\u003cstrong\u003eFig. 5b\u003c/strong\u003e). The Li-In full cell defines a near-equilibrium baseline that represents the thermodynamically compatible regime. However, the trade-off is clear: its operating voltage stays below about 3.7 V, which in turn caps the energy density even though the capacity remains high. Moving into the dynamically stabilized regime, the Li-Si full cell initially delivers a high capacity of 155 mAh g\u003csup\u003e-1\u003c/sup\u003e at 0.1C. The net capacity is slightly lower than the Li-In cell, but the lower Si anode potential allows us to extend the upper cutoff voltage to 4.1 V (4.3 V vs. Li/Li\u003csup\u003e+\u003c/sup\u003e). Interestingly, the early-cycle specific capacity at 0.33 C starts around 110 mAh g\u003csup\u003e-1\u003c/sup\u003e. However, the capacity gradually increases with continued cycling and climbs to over 140 mAh g\u003csup\u003e-1\u003c/sup\u003e by about 300 cycles (\u003cstrong\u003eFig. 5c\u003c/strong\u003e). The slow activation is consistent with the interfacial reorganization scenario inferred from the Li-Si symmetric cells, where the interface starts with relatively high resistance and then gradually stabilizes as Li redistributes within the lithiated LYC layer. It is also important to note that our Li-Si full cells are operated with a tightly limited N/P capacity ratio. Before assembling full cells, we independently evaluated the practical capacity of the pre-lithiated Si anode and obtained about 300 mAh g\u003csup\u003e-1\u003c/sup\u003e (\u003cstrong\u003eSupplementary Fig. 16\u003c/strong\u003e). Based on this value, the N/P capacity ratio was set to 1.2. In such a tightly limited N/P configuration, the fact that the cells cycle stably for 500 cycles at 0.33 C with an average CE of about 99% indicates that continuous LYC decomposition is not the dominant process. Instead, it supports the presence of a thin lithiated LYC interphase that acts as a self-limiting redox buffer. At the most reducing limit, Li metal based full cells do not reach similar performance or stability as expected. The initial discharge capacity at 0.1 C is already limited to about 110 mAh g\u003csup\u003e-1\u003c/sup\u003e and shows large polarization (\u003cstrong\u003eFig. 5d\u003c/strong\u003e). With further cycling at 0.33 C, the cells degrade rapidly and soon display over-charging behavior, which is characteristic of internal short-circuit formation. This outcome is consistent with the deep-lithiation regime, where the interfacial region gains excessive resistance, the current focuses into a few weak spots, and Li filaments penetrate through the electrolyte at the full-cell level.\u003c/p\u003e\n\u003cp\u003eCollectively, the symmetric and full-cell results show that LYC is not governed by a single, fixed stability window set only by thermodynamics. Its effective working range is instead controlled by how far lithiation extends from the anode into LYC and by how the resulting lithiated region balances redox buffering and charge transport kinetics. These trends redefine and extend a practical anode potential window for LYC and point to a broader design principle for redox-active halide electrolytes. Anodes that induce controlled, partial lithiation can support self-limiting interphases and extend the usable voltage range, whereas pushing the electrolyte into deep reduction promotes current localization, filament growth, and eventual failure.\u003c/p\u003e"},{"header":"3. Conclusion","content":"\u003cp\u003eState-of-the-art solid electrolytes now offer fast Li\u003csup\u003e+\u003c/sup\u003e transport comparable to liquid electrolytes, yet many still struggle when placed against anodes. In many electrolyte families, entering the reducing regime is widely treated as direct phase decomposition, which narrows viable separator choices. Under this assumption, solid-state cell designs have largely relied on protective interlayers or elevated-potential alloy anodes, rather than freely adopting low-voltage, high-capacity anodes. On the other hand, redox-active halide electrolytes invite a different view of reductive stability. Reduction can proceed through a structurally-intact lithiation process and be reverted on de-lithiation, rather than immediately driving phase decomposition. The key question becomes no longer where reduction starts, but how it evolves. Once paired with a low-potential anode, the same question becomes interfacial. The focus shifts to how the redox response develops at the contact, whether it stays confined and reversible, and how that interfacial evolution changes charge-transport kinetics and long-term stability at the anode interface.\u003c/p\u003e\n\u003cp\u003eIn this work, we show that reduction in trivalent halide LYC does not proceed as a single, irreversible failure. When the potential is driven below the nominal reduction onset (0.62 V vs Li/Li\u003csup\u003e+\u003c/sup\u003e), a reversible electrochemical response is activated, in which reduction at 0.25 V can be recovered by subsequent re-oxidation near 0.6 V. Such electrochemical response suggests that LYC exhibits intrinsic redox activity beyond its nominal stability limit and that the onset of reduction does not directly imply loss of electrolyte function. The observed reversibility becomes particularly advantageous when LYC is paired with low-potential anodes. Rather than triggering uncontrolled degradation, the reversible reduction allows the interface to accommodate partial lithiation and reorganize into a buffered interfacial state, mitigating continuous electrolyte reduction at the anode contact.\u003c/p\u003e\n\u003cp\u003eThe interfacial behavior of LYC depends strongly on the reducing strength of the anode. Li-In anodes operating within the thermodynamic stability window preserve near equilibrium interfacial kinetics. Moderately reducing Li-Si anodes induce partial lithiation that remains spatially confined and self-limiting, allowing the interface to stabilize through redox buffering rather than progressive degradation. In contrast, strongly reducing anodes like pure Li drive excessive lithiation, leading to increased interfacial resistance, lateral heterogeneity, and current focusing that ultimately drive short-circuit failure. Taken together, the results establish that practical operation of trivalent halide LYC is governed by a balance between reversible reduction and anode reducing strength at the interface. Importantly, the present study demonstrates, to our knowledge, that stable electrolyte operation can be achieved beyond the conventional thermodynamic reduction window through interfacially confined and kinetically regulated redox processes. By clarifying how redox reversibility, transport kinetics, and anode potential jointly control interfacial stability, this work extends the viable anode space for halide electrolytes and provides a mechanistic basis for rational anode matching and separator design in halide-based solid-state batteries.\u003c/p\u003e"},{"header":"4. Methods","content":"\u003cp\u003e\u003cem\u003eMaterial Synthesis and Basic Characterization\u003c/em\u003e\u003c/p\u003e\n\u003cp\u003eLi\u003csub\u003e3\u003c/sub\u003eYCl\u003csub\u003e6\u003c/sub\u003e (LYC) solid electrolyte was synthesized via a mechanochemical route under an argon atmosphere. LiCl (99.995%, Thermo Scientific) and YCl\u003csub\u003e3\u003c/sub\u003e (99.99%, Thermo Scientific) were weighed in stoichiometric ratios to a total mass of 1.5 g and loaded into an airtight 50 ml ZrO\u003csub\u003e2\u003c/sub\u003e grinding jar (Retsch) together with 12 ZrO\u003csub\u003e2\u003c/sub\u003e balls (10 mm diameter, MSE PRO) inside an Ar-filled glovebox. Mechanochemical synthesis was carried out using a high-energy ball mill (Emax, Retsch) operated at 600 rpm for a total net milling time of 36 h, employing a cycle of 5 min milling followed by 1 min rest. High-energy milling was used to drive the solid-state reaction and phase formation of LYC. To ensure homogeneous mixing and prevent local overheating, the milling process was periodically paused, and midway through the synthesis the jar was opened inside the glovebox to manually redistribute powder adhered to the jar walls. After milling, the as-synthesized LYC powder was lightly hand-ground in an agate mortar to break up soft agglomerates and stored in an Ar-filled glovebox (O\u003csub\u003e2\u003c/sub\u003e \u0026lt; 0.2 ppm, VAC) without exposure to ambient air. After synthesis, the powder morphology was characterized by scanning electron microscopy (SEM) using a field-emission SEM (JEOL 7500F) operated at an accelerating voltage of 15 kV.\u0026nbsp;\u003c/p\u003e\n\u003cp\u003e\u003cem\u003eModel Cell Preparation and Testing\u003c/em\u003e\u003c/p\u003e\n\u003cp\u003eLYC/C composite powders were prepared by mixing the as-synthesized LYC with carbon black (Super C65, TIMCAL) in an 8:2 weight ratio. A total of 0.2 g of the mixture was loaded together with 6 ZrO\u003csub\u003e2\u003c/sub\u003e balls and ball-milled at 300 rpm for 50 min using the same milling interval and the same high-energy ball mill (Emax, Retsch). The procedure ensured homogeneous dispersion of the conductive additive without inducing additional structural changes in LYC. All composite preparation steps were performed under an Ar atmosphere. Indium (In, Thermo Scientfic, 0.127 mm thick, 99.99%) and Lithium (Li, Sigma Aldrich, 99.9%) metal foils were used as received to prepare Li-In alloy anodes. Lithium and indium were first combined at a Li:In molar ratio of 3:7. The metals were stacked together and mechanically alloyed by repeated roll-and-press processing using stainless steel rods inside an Ar-filled glovebox. The stacked metals were continuously rolled and pressed to promote intimate contact and interdiffusion, and the process was repeated until a visually uniform and mechanically homogeneous Li-In alloy foil was obtained. Cell assembly was carried out by sequential pressing in a pressure-controlled and airtight split coin cell holder (PSC-10N, MTI Corporation). For the SS | LYC | Li-In cell, 100 mg of LYC powder was loaded into the air-tight cell holder and pressed at 300 MPa for 5 min using a hydraulic press (YLJ-15, MTI Corporation). A 10 mm diameter stainless steel rod integrated in the cell holder was used as the current collector on one side, while an 8 mm diameter Li-In alloy foil (0.1 mm thickness) was placed on the opposite side of the electrolyte to complete the cell stack. After assembly, the cell was secured using a compression jig (EQ-JIG-2, MTI Corporation) and operated under a constant stack pressure of approximately 8 MPa, as determined from the applied force and monitored with an integrated pressure sensor. For the LYC/C | LYC | Li\u0026ndash;In cell, the electrolyte pellet was prepared following the same procedure and under the same pressing conditions as described above. Subsequently, an LYC/C composite layer with a mass loading of 10 mg was applied onto the electrolyte surface and pressed at 150 MPa for 1 min. An 8 mm diameter Li\u0026ndash;In alloy foil (0.1 mm thickness) was then placed on the opposite side of the electrolyte to complete the cell stack. The assembled cell was secured and operated under the same constant stack pressure (~8 MPa) using the same compression jig and pressure monitoring method as described above.\u003c/p\u003e\n\u003cp\u003eCyclic voltammetry (CV) measurements were performed using the SS | LYC | Li-In cell at a scan rate of 0.1 mV s\u003csup\u003e-1\u003c/sup\u003e. For the first CV cycle, the potential sweep was initiated from the open-circuit voltage (OCV) and scanned toward the lower cutoff voltage, where the sweep direction was reversed. From the second cycle onward, CV scans were conducted between a fixed upper cutoff voltage of 2.0 V vs. Li/Li\u003csup\u003e+\u003c/sup\u003e (1.38 V vs. Li-In) and the designated lower cutoff voltage. The lower cutoff voltage was progressively decreased from 0.62 V to 0 V vs. Li/Li\u003csup\u003e+\u003c/sup\u003e to probe electrochemical activity beyond the nominal reduction limit of LYC. Multiple consecutive cycles were collected at each cutoff voltage to evaluate the evolution and reversibility of the redox response. Galvanostatic cycling was conducted at current densities of 0.1 and 0.2 mA cm\u003csup\u003e-2\u003c/sup\u003e. Cells were cycled between either 0-2.5 V or 0.1-2 V vs. Li/Li⁺, depending on the specific experiment. Coulombic efficiency was calculated from the ratio of charge to discharge capacity for each cycle. All measurements were carried out under a constant stack pressure of approximately 8 MPa. In situ electrochemical impedance spectroscopy (EIS) measurements were performed on the same LYC/C | LYC | Li-In cell. Impedance spectra were collected at selected potentials during galvanostatic cycling over a frequency range of 7 MHz to 0.1 Hz with an AC amplitude of 10 mV. The EIS signal was recorded with six points per decade using a logarithmic scale. The raw impedance data were processed using Python-based DRTtools\u003csup\u003e40\u003c/sup\u003e to perform Distribution of Relaxation Times (DRT) analysis. A first-order derivative regularization (regularization parameter of 0.005) and a FWHM coefficient of 0.5 were applied to stabilize the solution. DRT intensities were first mapped as a function of potential for each state of charge, and the absolute resistance for each contribution was subsequently calculated by integrating the DRT peak over its associated relaxation time range. Galvanostatic Intermittent Titration Technique (GITT) experiments were conducted on the same LYC/C | LYC | Li-In cell. Small current pulses of 0.1 mA cm\u003csup\u003e-2\u003c/sup\u003e were applied for 10 min, followed by 10 min relaxation periods to allow the potential to reach a quasi-steady state. Voltages were recorded during the pulses and relaxation periods, and the effective Li⁺ diffusivity was calculated using the established GITT relation, taking into account the slope of the quasi-steady-state voltage and the instantaneous transient voltage after each current pulse. Data processing and diffusivity calculations were performed using Python scripts developed in-house, ensuring consistent analysis across all pulses and cycles.\u003c/p\u003e\n\u003cp\u003e\u003cem\u003eHigh Energy X-ray Characterization\u003c/em\u003e\u003c/p\u003e\n\u003cp\u003eSynchrotron-based structural and spectroscopic characterization was performed on LYC/C electrode materials recovered after electrochemical testing. All samples were prepared and handled in powder form. After cycling, the LYC/C electrode layer was mechanically separated from the cell stack and collected intact, and a portion of the recovered electrode powder was used directly for synchrotron measurements. High-resolution powder X-ray diffraction (XRD) was carried out at Beamline 2-1 of the Stanford Synchrotron Radiation Lightsource (SSRL), SLAC National Accelerator Laboratory. The recovered LYC/C powders were loaded into glass capillaries inside an Ar-filled glovebox and sealed to prevent air exposure. Diffraction data were acquired at room temperature using monochromatic X-rays, with the wavelength specified in the main text. The capillary geometry enabled bulk-sensitive measurements while minimizing preferred orientation effects. The XRD patterns were analyzed using the HighScore Plus software. Background subtraction was first applied using a polynomial fitting function, followed by peak identification based on reference patterns of LYC. Phase analysis was conducted to confirm the absence of secondary crystalline phases within the detection limit of the measurement. Lattice parameters were then calculated from the identified peak positions using least-squares fitting, allowing comparison of structural changes before and after electrochemical cycling. X-ray absorption spectroscopy (XAS) measurements were conducted at SSRL to probe the local electronic structure and coordination environment of Y and Cl in LYC. Y K-edge XAS was performed at Beamline 2-3, and Cl K-edge XAS was collected at Beamline 14-3b. For XAS measurements, the recovered LYC/C powders were transferred into airtight sample holders inside an Ar-filled glovebox. The sealed holders were transported to the beamlines and loaded directly into the measurement chambers without air exposure. All XAS spectra were collected at room temperature in fluorescence mode, with energy calibration performed using reference standards measured concurrently. X-ray absorption spectroscopy (XAS) data were processed using the Athena software package. Energy calibration was performed using simultaneously measured reference foils, followed by normalization and background subtraction using standard pre-edge and post-edge fitting procedures. White-line features at the Y K-edge and Cl K-edge were analyzed by integrating the normalized absorption intensity over a defined energy window to track relative changes in unoccupied electronic states associated with electrochemical cycling.\u003c/p\u003e\n\u003cp\u003e\u003cem\u003eElemental and Chemical State Analysis\u003c/em\u003e\u003c/p\u003e\n\u003cp\u003eThe Li content of cycled LYC/C electrode powders was quantified using Inductively Coupled Plasma Optical Emission Spectroscopy (ICP-OES, Agilent 5800 Series). Recovered electrode powders were carefully separated from the cell stack to collect only the active LYC/C material, then digested in a mixture of concentrated hydrochloric acid (HCl, 37 %, ACS Grade, Sigma Aldrich) and nitric acid (HNO\u003csub\u003e3\u003c/sub\u003e, 68-70 vol%, ACS Grade, VWR Chemicals) at 60 \u0026deg;C for 12 h to ensure complete dissolution. The resulting solutions were diluted with deionized water to a final acid concentration of 2 % for measurement. Calibration standards for Li and Y were prepared from commercial stock solutions (1000 mg l\u003csup\u003e-1\u003c/sup\u003e in 2 % HNO\u003csub\u003e3\u003c/sub\u003e, Sigma Aldrich) and serially diluted to generate reference points. Emission intensities at selected wavelengths (Li: 274.119, 323.263, 460.289, 610.365, 670.783 nm; Y: 324.228, 360.074, 361.104, 371.029, 377.433 nm) were measured, and calibration curves were constructed by linear regression of standard signals versus known concentrations. Sample emission intensities were converted to concentrations (ppm) using these curves, and Li/Y stoichiometry of each sample was determined by converting concentrations to molar amounts and normalizing to Y content, enabling a direct comparison with electrochemically estimated Li insertion. X-ray Photoelectron Spectroscopy (XPS) measurements were performed using a Thermo-Fisher K-Alpha Plus spectrometer equipped with a monochromatic Al K\u0026alpha; X-ray source (1486 eV). The cycled LYC/C electrodes were carefully recovered from the cell stack, keeping the electrode structure largely intact, and mounted in airtight sample holders inside an Ar-filled glovebox to minimize air exposure before transfer to the spectrometer. All measurements were conducted under high vacuum (\u0026lt;10\u003csup\u003e-8\u003c/sup\u003e Torr). Survey scans were first collected to identify surface species, followed by high-resolution scans of the Y 3d, Cl 2p, C 1s, and Li 1s core levels at a step size of 0.1 eV to monitor changes in elemental valence states during Li insertion and extraction. All XPS spectra were processed using CasaXPS. Background signals were subtracted using a Shirley-type baseline to remove inelastic contributions. High-resolution peaks were fitted with mixed Gaussian-Lorentzian (GL (30)) line shapes. Initial binding energy calibration was performed using the C 1s hydrocarbon peak at 284.8 eV\u003csup\u003e48\u003c/sup\u003e, followed by secondary calibration based on the Cl 2p peak to correct for systematic peak shifts. Quantitative fitting of the Y 3d peaks allowed determination of the relative contributions of Y\u003csup\u003e3+\u003c/sup\u003e and partially reduced Y\u003csup\u003en+\u003c/sup\u003e species, enabling tracking of valence state evolution during Li insertion and extraction.\u003c/p\u003e\n\u003cp\u003e\u003cem\u003eAtomistic Simulations\u003c/em\u003e\u003c/p\u003e\n\u003cp\u003eThe workflow to identify energy minimized configurations for different x values for convex hull calculations is as follows: first, 10,000 LYC configurations at each x value are initialized with the knowledge of the crystallographic sites of Y and Li and their fractional occupancies, using the CIF file provided by Schlem \u003cem\u003eet al.\u003c/em\u003e\u003csup\u003e33\u003c/sup\u003e. For high enough Li values (x \u0026gt; 4), the interstitial tetrahedral sites were considered as the secondary occupation sites for Li in initial configuration generation. To perform greatly expedited molecular dynamics simulations and geometry optimization without compromising the accuracy, MLIP-based calculations were performed with PET-OAM-XL\u003csup\u003e35\u003c/sup\u003e, the accuracy of which for material systems is supported by its current top rank in the Matbench Discovery benchmark\u003csup\u003e49\u003c/sup\u003e. Using PET-OAM-XL, we first performed total energy evaluations for all of the randomly initialized configurations. We then selected 5 structures with the lowest total energies and another 5 structures at random from the rest of the structures for each x value. We then performed NVT ensemble MD simulations for each of the 10 systems with 2 fs timesteps for 100 ps at 400 K. i-PI was used for the MD simulations\u003csup\u003e50\u003c/sup\u003e, and the stochastic velocity rescaling (SVR) thermostat\u003csup\u003e51\u003c/sup\u003e was employed with a tau value of 50 fs. initial, final, and 4 intermediate configurations sampled at regular intervals were taken for geometry optimization in atomic simulation environment (ASE)\u003csup\u003e52\u003c/sup\u003e. Out of the 60 configurations in total, 5 lowest energy configurations were selected for further density functional theory (DFT)-based geometry optimization using the Vienna ab initio Software Package (VASP) v. 6.4.3\u003csup\u003e53\u003c/sup\u003e. In both geometry optimization calculations, maximum force component threshold was set to 0.025 eV \u0026Aring;\u003csup\u003e-1\u003c/sup\u003e. The DFT calculations were performed under the projector-augmented-wave (PAW) formalism\u003csup\u003e54-55\u003c/sup\u003e with the Perdew-Burke-Ernzerhof (PBE) functional\u003csup\u003e56\u003c/sup\u003e, and used a kinetic energy cutoff of 600 eV and energy convergence threshold of 1e-6 eV in the self-consistency field (SCF) cycles. Gamma-centered k-point grid was defined with a spacing of 0.25 \u0026Aring;\u003csup\u003e-1\u003c/sup\u003e, and electronic occupation was described using Gaussian smearing with a sigma value of 0.05 eV. The convex hull diagram was constructed using the lowest energy configurations after the DFT geometry optimization. RDF analysis was performed using the trajectories from the MLIP-based MD simulations.\u003c/p\u003e\n\u003cp\u003e\u003cem\u003eSymmetric and Full-cell Preparation and Testing\u003c/em\u003e\u003c/p\u003e\n\u003cp\u003eFor symmetric cells, LYC pellets were first prepared by pressing as-synthesized LYC powder following the same procedure as described above. Li-In alloy anodes were prepared as previously described, with 8 mm diameter discs placed on both sides of the LYC pellet. Li-Si alloy powders were synthesized starting from silicon powder (Si, 325 mesh, 99%, Sigma Aldrich) and lithium metal at a 1:1 molar ratio (Si:Li). Lithium was brought into contact with the silicon powder on a hot plate at 300 \u0026deg;C, where it melted and infiltrated the Si particles. During lithiation, the mixture gradually changed to a lighter color and expanded in volume, forming a partially lithiated Li-Si composite powder. The reaction proceeded for 30 min to ensure thorough lithiation, after which the resulting powder was cooled to room temperature and handled under an Ar atmosphere to prevent oxidation. For each Li-Si symmetric cell, 10 mg of this Li-Si powder was loaded onto both sides of the LYC pellet and pressed at 150 MPa for 1 min per side. Pure Li metal anodes were prepared by rolling as-received Li foil using stainless steel rods inside an Ar-filled glovebox until a uniform thickness of approximately 100 \u0026micro;m was achieved, and then cut into 8 mm diameter discs for placement on both sides of the LYC pellet. All symmetric cells (Li-In, Li-Si, and Li) were assembled in the same pressure-controlled and airtight split coin cell holders (10 mm diameter, PSC-10N, MTI Corporation) and compressed to a stack pressure of 16 MPa using a compression jig (EQ-JIG-2, MTI Corporation). Open-circuit storage impedance was recorded for 1000 min, with Nyquist spectra collected every 10 min (100 cycles) over a frequency range of 7 MHz to 0.1 Hz with a 10 mV AC amplitude. The EIS signal was recorded with six points per decade using a logarithmic scale. Galvanostatic cycling was then performed at 0.25 mA cm\u003csup\u003e-2\u003c/sup\u003e with 1 h per half-cycle (plating or stripping), during which in-situ impedance measurements were collected every 100 cycles over the same frequency range with the same AC amplitude. EIS spectra were fitted using EC-lab software to extract interfacial and bulk resistance values. Full cells were assembled using the LYC electrolyte and a high-capacity LiNi\u003csub\u003e0.8\u003c/sub\u003eMn\u003csub\u003e0.1\u003c/sub\u003eCo\u003csub\u003e0.1\u003c/sub\u003eO\u003csub\u003e2\u003c/sub\u003e single-crystal cathode (sc-NMC811, Targray). The cathode composite was prepared by low-energy ball-milling the sc-NMC811, LYC, and carbon black (Super P C65) in a weight ratio of 58:37:5 using the same ball mill machine (Emax, Retsch) at 300 rpm for 50 min, ensuring homogeneous mixing without inducing structural changes. The cells were assembled using sequential pressing of the following: 100 mg of LYC powder (300 MPa for 5 min, net thickness between 400-450 \u0026micro;m), followed by 15.3 mg cm\u003csup\u003e-2\u003c/sup\u003e of cathode composite (150 MPa for 1 min. The cathode active loading is 8.8 mg cm\u003csup\u003e-2\u003c/sup\u003e which corresponds to\u0026nbsp;~1.8 mAh cm\u003csup\u003e-2\u003c/sup\u003e. For full cells with Li-In anodes, 10 mm diameter Li-In alloy discs prepared as previously described were placed on the opposite side of the LYC pellet. For full cells with Li-Si anodes, 8 mg of the previously synthesized Li-Si powder (Li:Si = 1:1 molar ratio) was loaded onto the LYC pellet and pressed at 150 MPa for 1 min. For full cells with pure Li metal anodes, 10 mm diameter Li foil discs (100 \u0026micro;m thick) were placed on the other side of the LYC pellet. All cells were assembled in the split coin cell holders (12 mm diameter, ) and a screw compression jig (EQ-YLJ-SP, MTI Corporation) compressed to a stack pressure of 8 MPa. Galvanostatic cycling was conducted in an Ar-filled glovebox (O\u003csub\u003e2\u003c/sub\u003e \u0026lt; 0.2 ppm, VAC) at ambient temperature (20\u0026plusmn;5 \u0026deg;C), with the current rate calculated based on the total cathode capacity of 200 mAh g\u003csup\u003e-1\u003c/sup\u003e.\u003c/p\u003e"},{"header":"Declarations","content":"\u003cp\u003eAuthor Contributions\u003c/p\u003e\n\u003cp\u003eH.C., S.Y., and G.C. conceived the idea for the project. G.C. supervised the entire research. H.C. and S.Y. performed the electrochemical and X-ray measurements, and analyzed the experimental data. S.C. conducted the computational modeling. P.D. helped with the analysis of absorption and photoelectron measurements. P.P.P. assisted with the analysis of the diffraction measurements. H.C. wrote the initial draft of the manuscript, and G.C. revised the manuscript. All authors discussed the results, contributed to the content, and reviewed the final manuscript.\u003c/p\u003e\n\u003cp\u003eConflict of Interest\u003c/p\u003e\n\u003cp\u003eThe authors declare no conflict of interest\u003c/p\u003e\n\u003cp\u003eAcknowledgements\u003c/p\u003e\n\u003cp\u003eThe authors thank Dr. Kevin Stone (BL 2-1), and Dr. Sam Webb (BL 2-3 \u0026amp; BL 14-3b) at the Stanford Synchrotron Radiation Lightsource (SSRL) for their assistance with the high-energy XRD and XAS measurements. Use of SSRL, SLAC National Accelerator Laboratory was supported by the U.S. Department of Energy, Office of Science, Office of Basic Energy Sciences under Contract No. DE-AC02-76SF00515. Studies at the Molecular Foundry were supported by the Office of Science, Office of Basic Energy Sciences, of the U.S. Department of Energy under Contract No. DE-AC02-05CH11231. This work was supported by the Assistant Secretary for Energy Efficiency and Renewable Energy, Office of Vehicle Technologies, of the U.S. Department of Energy under Contract No. DE-AC02-05CH11231.\u003c/p\u003e"},{"header":"References","content":"\u003col\u003e\n\u003cli\u003eJanek, J. \u0026amp; Zeier, W. G. A solid future for battery development. \u003cem\u003eNat. Energy\u003c/em\u003e \u003cstrong\u003e1\u003c/strong\u003e, 16141 (2016).\u003c/li\u003e\n\u003cli\u003eJanek, J. \u0026amp; Zeier, W. G. Challenges in speeding up solid-state battery development. \u003cem\u003eNat. Energy\u003c/em\u003e \u003cstrong\u003e8\u003c/strong\u003e, 230\u0026ndash;240 (2023).\u003c/li\u003e\n\u003cli\u003eLi, X., Liang, J., Yang, X., Adair, K. R., Wang, C., Zhao, F. \u0026amp; Sun, X. Progress and perspectives on halide lithium conductors for all-solid-state lithium batteries. \u003cem\u003eEnergy Environ. Sci.\u003c/em\u003e \u003cstrong\u003e13\u003c/strong\u003e, 1429\u0026ndash;1461 (2020).\u003c/li\u003e\n\u003cli\u003eTanaka, Y., Ueno, K., Mizuno, K., Takeuchi, K., Asano, T. \u0026amp; Sakai, A. New oxyhalide solid electrolytes with high lithium ionic conductivity \u0026gt;10 mS cm-1 for all-solid-state batteries. \u003cem\u003eAngew. Chem. Int. Ed.\u003c/em\u003e \u003cstrong\u003e62\u003c/strong\u003e, e202217581 (2023).\u003c/li\u003e\n\u003cli\u003eAsano, T., Sakai, A., Ouchi, S., Sakaida, M., Miyazaki, A. \u0026amp; Hasegawa, S. Solid halide electrolytes with high lithium-ion conductivity for application in all-solid-state batteries. \u003cem\u003eAdv. Mater.\u003c/em\u003e \u003cstrong\u003e30\u003c/strong\u003e, 1803075 (2018).\u003c/li\u003e\n\u003cli\u003eGinnings, D. C. \u0026amp; Phipps, T. E. Temperature\u0026ndash;conductance curves of solid salts. III. Halides of lithium. \u003cem\u003eJ. Am. Chem. Soc.\u003c/em\u003e \u003cstrong\u003e52\u003c/strong\u003e, 1340\u0026ndash;1345 (1930).\u003c/li\u003e\n\u003cli\u003eWeppner, W. \u0026amp; Huggins, R. A. Ionic conductivity of alkali metal chloroaluminates. \u003cem\u003ePhys. Lett. A\u003c/em\u003e \u003cstrong\u003e58\u003c/strong\u003e, 245\u0026ndash;248 (1976).\u003c/li\u003e\n\u003cli\u003eKamaya, N., Homma, K., Yamakawa, Y., Hirayama, M., Kanno, R., Yonemura, M., Kamiyama, T., Kato, Y., Hama, S., Kawamoto, K. \u0026amp; Mitsui, A. A lithium superionic conductor. \u003cem\u003eNat. Mater.\u003c/em\u003e \u003cstrong\u003e10\u003c/strong\u003e, 682\u0026ndash;686 (2011).\u003c/li\u003e\n\u003cli\u003eAdeli, P., Bazak, J. D., Park, K. H., Kochetkov, I., Huq, A., Goward, G. R. \u0026amp; Nazar, L. F. Boosting solid-state diffusivity and conductivity in lithium superionic argyrodites by halide substitution. \u003cem\u003eAngew. Chem. Int. Ed.\u003c/em\u003e \u003cstrong\u003e58\u003c/strong\u003e, 8681 (2019).\u003c/li\u003e\n\u003cli\u003eBouguern, M. D., Ningappa, N. G., Vishweswariah, K., Kumar, A. M. R., Kanno, R. \u0026amp; Zaghib, K. Comparative advances in sulfide and halide electrolytes for commercialization of all-solid-state lithium batteries. \u003cem\u003eAdv. Mater.\u003c/em\u003e \u003cstrong\u003e38\u003c/strong\u003e, 2, e13255 (2026).\u003c/li\u003e\n\u003cli\u003eKim, K., Park, D., Jung, H.-G., Chung, K. Y., Shim, J. H., Wood, B. C. \u0026amp; Yu, S. Material design strategy for halide solid electrolytes Li\u003csub\u003e3\u003c/sub\u003eMX\u003csub\u003e6\u003c/sub\u003e (X = Cl, Br, and I) for all-solid-state high-voltage Li-ion batteries. \u003cem\u003eChem. Mater.\u003c/em\u003e \u003cstrong\u003e33\u003c/strong\u003e, 3669\u0026ndash;3677 (2021).\u003c/li\u003e\n\u003cli\u003eJun, K., Wei, G., Yang, X., Chen, Y. \u0026amp; Ceder, G. Exploring the soft cradle effect and ionic transport mechanisms in the LiMXCl\u003csub\u003e4\u003c/sub\u003e superionic conductor family. \u003cem\u003eMatter\u003c/em\u003e \u003cstrong\u003e8\u003c/strong\u003e, 102001 (2025).\u003c/li\u003e\n\u003cli\u003eYu, R., Wang, C., Duan, H., Jiang, M., Zhang, A., Fraser, A., Zuo, J., Wu, Y., Sun, Y., Zhao, Y., Liang, J., Fu, J., Deng, S., Ren, Z., Li, G., Huang, H., Li, R., Chen, N., Wang, J., Li, X., Singh, C. V. \u0026amp; Sun, X. Manipulating charge-transfer kinetics of lithium-rich layered oxide cathodes in halide all-solid-state batteries. \u003cem\u003eAdv. Mater.\u003c/em\u003e \u003cstrong\u003e35\u003c/strong\u003e, 2207234 (2023).\u003c/li\u003e\n\u003cli\u003eWang, K., Ren, Q., Gu, Z., Duan, C., Wang, J., Zhu, F., Fu, Y., Hao, J., Zhu, J., He, L., Wang, C.-W., Lu, Y., Ma, J. \u0026amp; Ma, C. A cost-effective and humidity-tolerant chloride solid electrolyte for lithium batteries. \u003cem\u003eNat. Commun.\u003c/em\u003e \u003cstrong\u003e12\u003c/strong\u003e, 4410 (2021).\u003c/li\u003e\n\u003cli\u003eZhou, L., Zuo, T.-T., Kwok, C. Y., Kim, S. Y., Assoud, A., Zhang, Q., Janek, J. \u0026amp; Nazar, L. F. High areal capacity, long cycle life 4 V ceramic all-solid-state Li-ion batteries enabled by chloride solid electrolytes. \u003cem\u003eNat. Energy\u003c/em\u003e \u003cstrong\u003e7\u003c/strong\u003e, 83\u0026ndash;93 (2022).\u003c/li\u003e\n\u003cli\u003eJiang, M., Mukherjee, S., Chen, Z. W., Chen, L. X., Li, M. L., Xiao, H. Y., Gao, C. \u0026amp; Singh, C. V. Materials perspective on new lithium chlorides and bromides: insights into thermo-physical properties. \u003cem\u003ePhys. Chem. Chem. Phys.\u003c/em\u003e \u003cstrong\u003e22\u003c/strong\u003e, 22758\u0026ndash;22767 (2020). \u003c/li\u003e\n\u003cli\u003eWang, S., Bai, Q., Nolan, A. M., Liu, Y., Gong, S., Sun, Q. \u0026amp; Mo, Y. Lithium chlorides and bromides as promising solid-state chemistries for fast ion conductors with good electrochemical stability. \u003cem\u003eAngew. Chem. Int. Ed.\u003c/em\u003e \u003cstrong\u003e58\u003c/strong\u003e, 8039\u0026ndash;8043 (2019).\u003c/li\u003e\n\u003cli\u003eKim, J. S., Jung, S., Kwak, H., Han, Y., Kim, S., Lim, J., Lee, Y. M. \u0026amp; Jung, Y. S. Synergistic halide\u0026ndash;sulfide hybrid solid electrolytes for Ni-rich cathode design guided by digital twin for all-solid-state Li batteries. \u003cem\u003eEnergy Storage Mater.\u003c/em\u003e \u003cstrong\u003e55\u003c/strong\u003e, 193\u0026ndash;204 (2023).\u003c/li\u003e\n\u003cli\u003eKim, S. Y., Cha, H., Kostecki, R. \u0026amp; Chen, G. Composite cathode design for high-energy all-solid-state lithium batteries with long cycle life. \u003cem\u003eACS Energy Lett.\u003c/em\u003e \u003cstrong\u003e8\u003c/strong\u003e, 521\u0026ndash;528 (2023).\u003c/li\u003e\n\u003cli\u003eYang, S., Kim, S. Y. \u0026amp; Chen, G. Halide superionic conductors for all-solid-state batteries: Effects of synthesis and composition on lithium-ion conductivity. \u003cem\u003eACS Energy Lett.\u003c/em\u003e \u003cstrong\u003e9 \u003c/strong\u003e(5), 2212\u0026ndash;2221 (2024).\u003c/li\u003e\n\u003cli\u003eMorino, Y., Aoki, N. \u0026amp; Nakamoto, M. Revealing the reductive decomposition at the interface of Li\u003csub\u003e3\u003c/sub\u003eYCl\u003csub\u003e6\u003c/sub\u003e\u0026ndash;lithium metal. \u003cem\u003eACS Appl. Mater. Interfaces\u003c/em\u003e \u003cstrong\u003e17\u003c/strong\u003e, 45744\u0026ndash;45751 (2025).\u003c/li\u003e\n\u003cli\u003eJi, W., Zheng, D., Zhang, X., Ding, T. \u0026amp; Qu, D. A kinetically stable anode interface for Li\u003csub\u003e3\u003c/sub\u003eYCl\u003csub\u003e6\u003c/sub\u003e-based all-solid-state lithium batteries. \u003cem\u003eJ. Mater. Chem. A\u003c/em\u003e \u003cstrong\u003e9\u003c/strong\u003e, 15012\u0026ndash;15018 (2021).\u003c/li\u003e\n\u003cli\u003eHu, L., Zhu, J., Duan, C., Zhu, J., Wang, J., Wang, K., Gu, Z., Xi, Z., Hao, J., Chen, Y., Ma, J., Liu, J.-X. \u0026amp; Ma, C. Revealing the Pnma crystal structure and ion-transport mechanism of the Li\u003csub\u003e3\u003c/sub\u003eYCl\u003csub\u003e6\u003c/sub\u003e solid electrolyte. \u003cem\u003eCell Rep. Phys. Sci.\u003c/em\u003e \u003cstrong\u003e4\u003c/strong\u003e, 101428 (2023).\u003c/li\u003e\n\u003cli\u003eLiu, Z., Zhang, G., Pepas, J., Ma, Y. \u0026amp; Chen, H. Li\u003csub\u003e2\u003c/sub\u003eFeCl\u003csub\u003e4\u003c/sub\u003e as a cost-effective and durable cathode for solid-state Li-ion batteries. \u003cem\u003eACS Energy Lett.\u003c/em\u003e \u003cstrong\u003e9\u003c/strong\u003e, 5464-5470 (2024).\u003c/li\u003e\n\u003cli\u003eZhang, G., Liu, Z., Ma, Y., Pepas, J., Bai, J., Zhong, H., Tang, Y. \u0026amp; Chen, H. Li\u003csub\u003e2.9\u003c/sub\u003eFe\u003csub\u003e0.9\u003c/sub\u003eZr\u003csub\u003e0.1\u003c/sub\u003eCl\u003csub\u003e6\u003c/sub\u003e as redox-active catholyte for solid-state Li-ion batteries. \u003cem\u003eChem. Mater.\u003c/em\u003e \u003cstrong\u003e36\u003c/strong\u003e, 10104-10112 (2024).\u003c/li\u003e\n\u003cli\u003eSong, Z., Dai, Y., Wang, T., Yu, Q., Ye, X., Wang, L., Zhang, Y., Wang, S. \u0026amp; Luo, W. An active halide catholyte boosts the extra capacity for all-solid-state batteries. \u003cem\u003eAdv. Mater.\u003c/em\u003e \u003cstrong\u003e36\u003c/strong\u003e, 2405277 (2024).\u003c/li\u003e\n\u003cli\u003eWang, K., Gu, Z., Xi, Z., Hu, L. \u0026amp; Ma, C. Li\u003csub\u003e3\u003c/sub\u003eTiCl\u003csub\u003e6\u003c/sub\u003e as ionic conductive and compressible positive electrode active material for all-solid-state lithium-based batteries. \u003cem\u003eNat. Commun.\u003c/em\u003e \u003cstrong\u003e14\u003c/strong\u003e, 1396 (2023).\u003c/li\u003e\n\u003cli\u003eSun, G., Song, Z., Dai, Y., Xiao, Z., Zheng, X. \u0026amp; Luo, W. Redox-active halide catholytes for solid-state lithium batteries. \u003cem\u003eAdv. Sci.\u003c/em\u003e \u003cstrong\u003e12\u003c/strong\u003e, 47, e14215 (2025).\u003c/li\u003e\n\u003cli\u003eRidley, P., Duong, G., Ko, S. L., Oh, J. A. S., Deysher, G., Griffith, K. J. \u0026amp; Meng, Y. S. Tailoring chloride solid electrolytes for reversible redox. \u003cem\u003eJ. Am. Chem. Soc.\u003c/em\u003e \u003cstrong\u003e147\u003c/strong\u003e, 19508\u0026ndash;19519 (2025).\u003c/li\u003e\n\u003cli\u003eSchwietert, T. K., Vasileiadis, A. \u0026amp; Wagemaker, M. First-principles prediction of the electrochemical stability and reaction mechanisms of solid-state electrolytes. \u003cem\u003eJACS Au\u003c/em\u003e \u003cstrong\u003e1\u003c/strong\u003e, 1488\u0026ndash;1496 (2021).\u003c/li\u003e\n\u003cli\u003eCheng, Z., Zhao, W., Wang, Q., Zhao, C., Lavrinenko, A. K., Vasileiadis, A., Landgraf, V., Bannenberg, L., Li, Y., Liang, J., Liu, M., Ganapathy, S. \u0026amp; Wagemaker, M. Beneficial redox activity of halide solid electrolytes empowering high-performance anodes in all-solid-state batteries. \u003cem\u003eNat. Mater.\u003c/em\u003e \u003cstrong\u003e24\u003c/strong\u003e, 1763\u0026ndash;1772 (2025).\u003c/li\u003e\n\u003cli\u003eVan der Ven, A., Bhattacharya, J. \u0026amp; Belak, A. A. Understanding Li diffusion in Li-intercalation compounds. \u003cem\u003eAcc. Chem. Res.\u003c/em\u003e \u003cstrong\u003e46\u003c/strong\u003e, 1216\u0026ndash;1225 (2013).\u003c/li\u003e\n\u003cli\u003eSchlem, R., Banik, A., Ohno, S., Suard, E. \u0026amp; Zeier, W. G. Insights into the lithium sub-structure of superionic conductors Li\u003csub\u003e3\u003c/sub\u003eYCl\u003csub\u003e6\u003c/sub\u003e and Li\u003csub\u003e3\u003c/sub\u003eYBr\u003csub\u003e6\u003c/sub\u003e. \u003cem\u003eChem. Mater.\u003c/em\u003e \u003cstrong\u003e33\u003c/strong\u003e, 327\u0026ndash;337 (2021).\u003c/li\u003e\n\u003cli\u003eTaminato, S., Yonemura, M., Shiotani, S., Kamiyama, T., Torii, S., Nagao, M., Ishikawa, Y., Mori, K., Fukunaga, T., Onodera, Y., Naka, T., Morishima, M., Ukyo, Y., Adipranoto, D. S., Arai, H., Uchimoto, Y., Ogumi, Z., Suzuki, K., Hirayama, M. \u0026amp; Kanno, R. Real-time observations of lithium battery reactions\u0026mdash;operando neutron diffraction analysis during practical operation. \u003cem\u003eSci. Rep.\u003c/em\u003e \u003cstrong\u003e6\u003c/strong\u003e, 28843 (2016).\u003c/li\u003e\n\u003cli\u003eBigi, F., Pegolo, P., Mazitov, A. \u0026amp; Ceriotti, M. Pushing the limits of unconstrained machine-learned interatomic potentials. \u003cem\u003earXiv\u003c/em\u003e:2601.16195 (2026).\u003c/li\u003e\n\u003cli\u003eKawaguchi, T., Fukuda, K. \u0026amp; Matsubara, E. Site- and phase-selective x-ray absorption spectroscopy based on phase-retrieval calculation. \u003cem\u003eJ. Phys.: Condens. Matter\u003c/em\u003e \u003cstrong\u003e29\u003c/strong\u003e, 113002 (2017).\u003c/li\u003e\n\u003cli\u003eWang, X., Yang, Z., Li, N., Wu, K., Gao, K., Zhao, E., Han, S. \u0026amp; Guo, W. Influence mechanism of interfacial oxidation of Li\u003csub\u003e3\u003c/sub\u003eYCl\u003csub\u003e6\u003c/sub\u003e solid electrolyte on reduction potential. \u003cem\u003eChem. Eur. J.\u003c/em\u003e \u003cstrong\u003e30\u003c/strong\u003e, e202303884 (2024).\u003c/li\u003e\n\u003cli\u003eVasquez, R. P. XPS core level spectra of some Y compounds. \u003cem\u003eJ. Electron Spectrosc. Relat. Phenom.\u003c/em\u003e \u003cstrong\u003e50\u003c/strong\u003e, 167\u0026ndash;170 (1990).\u003c/li\u003e\n\u003cli\u003eWu, J.-F. \u0026amp; Guo, X. Origin of the low grain boundary conductivity in lithium ion conducting perovskites: Li\u003csub\u003e3x\u003c/sub\u003eLa\u003csub\u003e0.67\u0026minus;x\u003c/sub\u003eTiO\u003csub\u003e3\u003c/sub\u003e. \u003cem\u003ePhys. Chem. Chem. Phys.\u003c/em\u003e 19, 5880\u0026ndash;5886 (2017).\u003c/li\u003e\n\u003cli\u003eWan, T. H., Saccoccio, M., Chen, C. \u0026amp; Ciucci, F. Influence of the discretization methods on the distribution of relaxation times deconvolution: Implementing radial basis functions with DRTtools. \u003cem\u003eElectrochim. Acta\u003c/em\u003e \u003cstrong\u003e184\u003c/strong\u003e, 483\u0026ndash;499 (2015).\u003c/li\u003e\n\u003cli\u003eWeppner, W. \u0026amp; Huggins, R. A. Determination of the kinetic parameters of mixed-conducting electrodes and application to the system Li\u003csub\u003e3\u003c/sub\u003eSb. \u003cem\u003eJ. Electrochem. Soc.\u003c/em\u003e \u003cstrong\u003e124\u003c/strong\u003e, 1569\u0026ndash;1578 (1977).\u003c/li\u003e\n\u003cli\u003eDeiss, E. Spurious chemical diffusion coefficients of Li+ in electrode materials evaluated with GITT. \u003cem\u003eElectrochim. Acta\u003c/em\u003e \u003cstrong\u003e50\u003c/strong\u003e, 2927\u0026ndash;2932 (2005).\u003c/li\u003e\n\u003cli\u003eLi, J., Su, H., Liu, Y., Zhong, Y., Wang, X. \u0026amp; Tu, J. Li alloys in all solid-state lithium batteries: A review of fundamentals and applications. \u003cem\u003eElectrochem. Energy Rev.\u003c/em\u003e \u003cstrong\u003e7\u003c/strong\u003e, 18 (2024).\u003c/li\u003e\n\u003cli\u003eIrvine, J. T. S., Sinclair, D. C. \u0026amp; West, A. R. Electroceramics: Characterization by impedance spectroscopy. \u003cem\u003eAdv. Mater.\u003c/em\u003e \u003cstrong\u003e2\u003c/strong\u003e, 132\u0026ndash;138 (1990).\u003c/li\u003e\n\u003cli\u003eDrvarič Talian, S., Brutti, S., Navarra, M. A., Mo\u0026scaron;kon, J. \u0026amp; Gaber\u0026scaron;ček, M. Impedance spectroscopy applied to lithium battery materials: Good practices in measurements and analyses. \u003cem\u003eEnergy Storage Mater.\u003c/em\u003e \u003cstrong\u003e69\u003c/strong\u003e, 103413 (2024).\u003c/li\u003e\n\u003cli\u003eZhang, S. S., Xu, K. \u0026amp; Jow, T. R. EIS study on the formation of solid electrolyte interface in Li-ion battery. \u003cem\u003eElectrochim. Acta\u003c/em\u003e \u003cstrong\u003e51\u003c/strong\u003e, 1636\u0026ndash;1640 (2006).\u003c/li\u003e\n\u003cli\u003eLewis, J. A., Cavallaro, K. A., Liu, Y. \u0026amp; McDowell, M. T. The promise of alloy anodes for solid-state batteries. \u003cem\u003eJoule\u003c/em\u003e \u003cstrong\u003e6\u003c/strong\u003e, 1418\u0026ndash;1430 (2022).\u003c/li\u003e\n\u003cli\u003eChu, H., Noh, H., Kim, Y.-J., Yuk, S., Lee, J.-H., Lee, J., Kwack, H., Kim, Y., Yang, D.-K. \u0026amp; Kim, H.-T. Achieving three-dimensional lithium sulfide growth in lithium-sulfur batteries using high-donor-number anions. \u003cem\u003eNat. Commun.\u003c/em\u003e \u003cstrong\u003e10\u003c/strong\u003e, 188 (2019).\u003c/li\u003e\n\u003cli\u003eRiebesell, J., Goodall, R. E. A., Benner, P., Chiang, Y., Deng, B., Ceder, G., Asta, M., Lee, A. A., Jain, A. \u0026amp; Persson, K. A. A framework to evaluate machine learning crystal stability predictions. \u003cem\u003eNat. Mach. Intell.\u003c/em\u003e \u003cstrong\u003e7\u003c/strong\u003e, 836\u0026ndash;847 (2025).\u003c/li\u003e\n\u003cli\u003eLitman, Y., Kapil, V., Feldman, Y. M. Y., Tisi, D., Begu\u0026scaron;ić, T., Fidanyan, K., Fraux, G., Higer, J., Kellner, M., Li, T. E., P\u0026oacute;s, E. S., Stocco, E., Trenins, G., Hirshberg, B., Rossi, M. \u0026amp; Ceriotti, M. i-PI 3.0: A flexible and efficient framework for advanced atomistic simulations. \u003cem\u003eJ. Chem. Phys.\u003c/em\u003e \u003cstrong\u003e161\u003c/strong\u003e, 062504 (2024).\u003c/li\u003e\n\u003cli\u003eBussi, G., Donadio, D. \u0026amp; Parrinello, M. Canonical sampling through velocity rescaling. \u003cem\u003eJ. Chem. Phys.\u003c/em\u003e \u003cstrong\u003e126\u003c/strong\u003e, 014101 (2007).\u003c/li\u003e\n\u003cli\u003eLarsen, A. H., Mortensen, J. J., Blomqvist, J., Castelli, I. E., Christensen, R., Dułak, M., Friis, J., Groves, M. N., Hammer, B. \u0026amp; Hargus, C. The atomic simulation environment\u0026mdash;a Python library for working with atoms. \u003cem\u003eJ. Phys. Condens. Matter\u003c/em\u003e \u003cstrong\u003e29\u003c/strong\u003e, 273002 (2017).\u003c/li\u003e\n\u003cli\u003eKresse, G. \u0026amp; Furthm\u0026uuml;ller, J. Efficient iterative schemes for ab initio total-energy calculations using a plane-wave basis set. \u003cem\u003ePhys. Rev. B\u003c/em\u003e \u003cstrong\u003e54\u003c/strong\u003e, 11169-11186 (1996).\u003c/li\u003e\n\u003cli\u003eBl\u0026ouml;chl, P. E. Projector augmented-wave method. \u003cem\u003ePhys. Rev. B\u003c/em\u003e \u003cstrong\u003e50\u003c/strong\u003e, 17953 (1994).\u003c/li\u003e\n\u003cli\u003eKresse, G. \u0026amp; Joubert, D. From ultrasoft pseudopotentials to the projector augmented-wave method. \u003cem\u003ePhys. Rev. B\u003c/em\u003e \u003cstrong\u003e59\u003c/strong\u003e, 1758 (1999).\u003c/li\u003e\n\u003cli\u003ePerdew, J. P., Burke, K. \u0026amp; Ernzerhof, M. Generalized gradient approximation made simple. \u003cem\u003ePhys. Rev. Lett.\u003c/em\u003e \u003cstrong\u003e77\u003c/strong\u003e, 3865 (1996).\u003c/li\u003e\n\u003c/ol\u003e"}],"fulltextSource":"","fullText":"","funders":[],"hasAdminPriorityOnWorkflow":false,"hasManuscriptDocX":true,"hasOptedInToPreprint":true,"hasPassedJournalQc":"","hasAnyPriority":true,"hideJournal":false,"highlight":"","institution":"","isAcceptedByJournal":false,"isAuthorSuppliedPdf":false,"isDeskRejected":"","isHiddenFromSearch":false,"isInQc":false,"isInWorkflow":false,"isPdf":false,"isPdfUpToDate":true,"isWithdrawnOrRetracted":false,"journal":{"display":true,"email":"[email protected]","identity":"nature-portfolio","isNatureJournal":true,"hasQc":false,"allowDirectSubmit":false,"externalIdentity":"","sideBox":"","snPcode":"","submissionUrl":"","title":"Nature Portfolio","twitterHandle":"","acdcEnabled":false,"dfaEnabled":false,"editorialSystem":"ejp","reportingPortfolio":"","inReviewEnabled":true,"inReviewRevisionsEnabled":false},"keywords":"","lastPublishedDoi":"10.21203/rs.3.rs-8928802/v1","lastPublishedDoiUrl":"https://doi.org/10.21203/rs.3.rs-8928802/v1","license":{"name":"CC BY 4.0","url":"https://creativecommons.org/licenses/by/4.0/"},"manuscriptAbstract":"\u003cp\u003eThe practical adoption of solid-state batteries requires solid electrolytes that sustain stable interfaces with electrodes. Conventionally, electrolyte stability has been defined by its thermodynamic limits, assuming that operating beyond these boundaries causes irreversible decomposition. Such perspectives have deterred the integration of many solid electrolytes with low-potential anodes, leaving their behavior beyond the limit largely unexplored. Here, we demonstrate that the trivalent metal halide Li\u003csub\u003e3\u003c/sub\u003eYCl\u003csub\u003e6\u003c/sub\u003e (LYC) can be electrochemically driven below its calculated reduction limit while exhibiting a predominantly reversible redox response. Lowering the potential below the thermodynamic window activates a reversible lithiation process consistent with Y-centered redox, while the halide framework is retained. At the anode interface, this reversibility enables LYC to accommodate low-potential anodes by forming a self-limiting lithiation layer that kinetically suppresses continued reduction. Comparative anode screening with Li-In, Li-Si, and Li metal identifies practical operating window below the nominal reduction threshold, within which LYC maintains both phase retention and ion transport with manageable impedance growth. Utilizing the reversible regime, we realize coating-free LYC full cells with an NMC811 cathode and a Li-Si anode that cycle stably for over 500 cycles at high capacity (~ 140 mAh g\u003csup\u003e-1\u003c/sup\u003e) and extended cell voltages (\u0026gt; 4 V), expanding the practical anode selection beyond conventional thermodynamic-window constraints and enabling high-performance all-solid-state cells.\u003c/p\u003e","manuscriptTitle":"Interfacial Redox Buffering Stabilizes Halide Solid Electrolytes Against Low-Potential Anodes","msid":"","msnumber":"","nonDraftVersions":[{"code":1,"date":"2026-03-27 06:13:53","doi":"10.21203/rs.3.rs-8928802/v1","editorialEvents":[],"status":"published","journal":{"display":true,"email":"[email protected]","identity":"nature-communications","isNatureJournal":true,"hasQc":false,"allowDirectSubmit":false,"externalIdentity":"NCOMMS","sideBox":"Learn more about [Nature Communications](http://www.nature.com/ncomms/)","snPcode":"","submissionUrl":"https://mts-ncomms.nature.com/","title":"Nature Communications","twitterHandle":"","acdcEnabled":true,"dfaEnabled":true,"editorialSystem":"ejp","reportingPortfolio":"Nature Communications","inReviewEnabled":true,"inReviewRevisionsEnabled":false}}],"origin":"","ownerIdentity":"9ec49c53-284e-415d-81f3-c0a8ea50402c","owner":[],"postedDate":"March 27th, 2026","published":true,"recentEditorialEvents":[],"rejectedJournal":[],"revision":"","amendment":"","status":"under-review","subjectAreas":[{"id":64257017,"name":"Physical sciences/Materials science/Materials for energy and catalysis/Batteries"},{"id":64257018,"name":"Physical sciences/Chemistry/Electrochemistry/Batteries"}],"tags":[],"updatedAt":"2026-03-27T06:13:53+00:00","versionOfRecord":[],"versionCreatedAt":"2026-03-27 06:13:53","video":"","vorDoi":"","vorDoiUrl":"","workflowStages":[]},"version":"v1","identity":"rs-8928802","journalConfig":"researchsquare"},"__N_SSP":true},"page":"/article/[identity]/[[...version]]","query":{"redirect":"/article/rs-8928802","identity":"rs-8928802","version":["v1"]},"buildId":"XKTyCvWXoU3ODBz1xrDgd","isFallback":false,"isExperimentalCompile":false,"dynamicIds":[84888],"gssp":true,"scriptLoader":[]}