Multicomponent solid-solution alloy negative electrode for Li-metal batteries

Domain: chemistry
Published: 2026-03-05
Editorial outcome: minor_revision

Abstract
Lithium metal batteries possess the high theoretical specific energy owing to the high theoretical specific capacity of lithium metal. However, practical implementations necessitate a reduction in both lithium mass fraction and lithium utilization within the negative electrode to enhance cycling life and suppress dendrite formation, resulting in a practical reversible capacity of only 30 – 50% of the theoretical value. Herein, we present a lithium metal-based multicomponent solid-solution alloy comprising approximately 90 wt.% lithium, with equal atomic ratios of cadmium, silver, magnesium, and aluminium constituting the remaining 10 wt.%. The increased mixing entropy enables high lithium-atom diffusivity, facilitating inward lithium transport into the metal foil rather than the surface deposition typically observed in conventional lithium metal negative electrodes, while simultaneously promoting a thermodynamically stable (110) crystal facet. These factors collectively yield a dendrite-free negative electrode with a reversible specific capacity of 3100 mAh g−1. One-ampere-hour pouch cells employing this negative electrode and a LiNi0.8Co0.1Mn0.1O2 positive electrode achieve an specific energies of 385 Wh kg−1 (based on the total mass of the pouch cell) with 82% capacity retention over 600 cycles. This achievement highlights the potential of this alloy negative electrode to enable safe and durable high-energy-density lithium metal batteries for practical applications.

Introduction
The pursuit of high-energy-density batteries for emerging electric aircraft and other advancements in the low-altitude economy has stimulated a significant increase in research on Li metal batteries (LMBs) owing to the high theoretical specific capacity (3860 mAh g−1) and the lowest redox potential (−3.04 V versus the standard hydrogen electrode) of Li metal negative electrode.1,2 Current investigations have predominantly centered on the challenges posed by uncontrolled Li dendrite growth, which results in low Coulombic efficiency (CE), poor cycle life expectancy, and potential safety hazards due to internal short-circuiting.3,4 In an effort to mitigate Li dendrite formation and extend the batteries' cycle life, a variety of extensive strategies have been employed, including the optimization of electrolytes and additives,5,6 the development of artificial protective layers,7,8 the utilization of three-dimensional host scaffolds,9 and the incorporation of Li-Ag binary solid-solution alloy.10,11

In the practical application of LMBs, it is imperative to not only address the challenges associated with Li dendrites and battery life span but also consider the Li mass fraction within the electrode (M) and Li utilization (U) during cycling. These factors directly determine the achievable practical specific capacity of the Li metal negative electrode, which can be expressed mathematically as: Practical specific capacity = 3860 mAh/g × M ×U.12 Current designs of Li metal electrodes compromise Li mass fraction, which is below 60 wt.%, or Li utilization, which is below 80 %, in an effort to extend cycling life and avoid the Li dendrite (Fig. 1a and Supplementary Table 1).11,13-22 This trade-off results in a practical capacity typically below 2000 mAh g−1, thereby diminishing the inherent advantage of Li metal's high specific capacity. Consequently, LMBs still suffered from the unsatisfactory trade-off between specific energy and cycling life (Fig. 1b).14,15,22-33 Taking a 10 Ah Li||LiNi0.8Co0.1Mn0.1O2 pouch cell as an example, the positive electrode employs a standard commercial positive electrode with an areal capacity of 3.5 mAh cm−2. If the negative electrode applies a 50 μm-thick Li metal, equaling to an areal capacity of 10 mAh cm−2, which is coated on both sides of an 8 μm-thick copper foil — the most commonly used commercial-scale Li-Cu-Li composite foil available for pouch cells — the Li mass ratio and the Li utilization within the negative electrode would be calculated to be 41 wt.% and 35%, respectively (Supplementary Table 2).34 This configuration yields a cell-level specific energy of up to 408 Wh kg−1 (based on the total mass of the pouch cell; Supplementary Fig. 1 and Supplementary Table 2). Should the Li mass fraction within the electrode and Li utilization be improved to 100 wt.% each, a higher specific energy (495 Wh kg−1, based on the total mass of the pouch cell) will be achievable (Supplementary Note 1). Therefore, the development of a dendrite-free Li metal negative electrode with a high Li mass ratio and utilization is essential for advancing the specific energy of Li metal batteries.

We hereby present a multicomponent solid-solution alloy, consisting of 90 wt.% metallic Li, referred to as Li-MSSA, with Cd, Ag, Mg, and Al constituting the remaining 10 wt.% in equal atomic ratios of 1:1:1:1. The increased mixing entropy within the Li-MSSA alloys enhances compatibility among the elements, thereby preventing phase separation and mitigating the formation of intermetallic compounds, which facilitates the establishment of a solid-solution-based metal alloy phase. The Li-MSSA presents rapid Li atom diffusion throughout the bulk of the electrode, rendering inward transfer of Li atoms into the metallic foil while concurrently directing the deposited Li toward the formation of the thermodynamically stable (110) crystal facet. These two effects lead to a reversible and dendrite-free lithiation-delithiation process of Li-MSSA, realize an combination of a high Li mass ratio and Li utilization, which leads to a reversible specific capacity of 3100 mAh g−1 with areal capacity reaching up to 20 mAh cm−2, depending on the thickness of the Li-MSSA foil. The ampere-hour level Li-metal pouch cells, comprising the configuration Li-MSSA||LiNi0.8Co0.1Mn0.1O2, achieve a balance of high specific energy and excellent cycling performance.

Results
2.1 Structure and electrochemical properties of Li-MSSA

The solid-solution-based metal alloy phase exhibits lower energy barriers for phase transition in comparison to intermetallic compounds. This characteristic results in markedly enhanced reversibility and cycling stability during the Li-alloying/dealloying process, as the solid-solution reaction accommodates a degree of stoichiometric tolerance to variations in Li atom concentrations and leads to significantly less structural alteration.10,11 Consequently, we selected metallic elements Ag, Mg, and Cd, which are recognized for their broad solid solution ranges with Li atoms, as alloying constituents.35-37 Given the diverse melting points and compatibilities of Mg, Ag, and Cd, Al was chosen as a "binder" due to its negative enthalpy of mixing,38 which promotes a high chemical affinity among distinct atomic species and guarantees the homogeneous dispersion of metallic components. The Li-MSSA was synthesized by heating a mixture of Li, Cd, Ag, Mg, and Al metallic powders at 600 °C under mechanical stirring in an argon-filled glovebox (see Methods section for detail). The resultant Li-MSSA comprises 90 wt.% metallic Li, with Cd, Ag, Mg, and Al making up the remaining 10 wt.% in an equal atomic ratio of 1:1:1:1 (Supplementary Fig. 2 and Supplementary Table 3). Atomically resolved high-angle annular dark-field scanning transmission electron microscopy (HAADF-STEM) with energy-dispersive X-ray spectroscopy (EDX) spectrum imaging (Fig. 2a and Supplementary Fig. 3) shows that Ag, Cd, Mg, and Al are co-spatial within our spatial-resolution and counting-statistics limits, with no detectable nanoscale segregation; large-area scanning electron microscopy coupled with (SEM-EDX) maps corroborate this (Supplementary Fig. 4). Because Li cannot be mapped reliably by EDX, we performed electron energy-loss spectroscopy (EELS) spectrum imaging at the Li K-edge; after standard background subtraction, the Li-K signal is uniform across the probed area, with no evidence of Li-rich/Li-poor domains down to sub-100-nanometre length scales (Fig. 2b and Supplementary Fig. 5). Taken together, these microscopy and spectroscopy results indicate that, within our nanometer-scale resolution, the delithiated Li-MSSA is compositionally uniform and behaves as a single-phase random solid solution, rather than a mixture of element-rich regions or intermetallic secondary phases. Of particular interest, as shown in Fig. 2c, a 20-meter-long Li-MSSA foil can be readily fabricated through heating-rolling process. Its thickness can be adjusted to several hundred micrometers or reduced to tens of micrometer (Supplementary Fig. 6). This characteristic enables the Li-MSSA negative electrode to align positive electrodes of varying areal capacities, thus providing a broad range for optimizing full-cell design, which is critical for battery engineering. It is noteworthy that the heating-rolling process for preparing the Li-MSSA is commonly utilized in industrial-scale metal processing, showing the feasibility of scaling the Li-MSSA technology for practical applications.

To evaluate the practical specific capacity of our electrode, Li was extracted from the Li-MSSA within an electrochemical cell comprising Cu foil as the counter electrode. We observed that a 30 μm-thick Li-MSSA can undergo reversible lithiation and delithiation at potentials of −25 mV and 25 mV, respectively (Fig. 2d). During this process, the Li atom stoichiomitry varies from Li9.9M to LiM (M stands for the Cd, Ag, Mg, and Al in the Li-MSSA), resulting in a high reversible specific capacity of 3100 mAh g−1 and an areal capacity of 6 mAh cm−2 (Fig. 2d and Supplementary Note 2). This value is within the upper range of practical values reported for lithium-metal electrodes when both lithium mass fraction and utilization are considered (Fig. 1a and Supplementary Table 1). This high reversible specific capacity corresponds to a substantial Li utilization of 89% for the Li-MSSA (Supplementary Note 3). To underscore the critical importance of high Li utilization in practical LMBs, in the scenario involving the same 30 μm-thick Li-MSSA, a reduction of the reversible areal capacity to 4 mAh cm−2 results in a decrease in Li utilization to 59%, consequently lowering the electrode’s specific capacity capacity to 2060 mAh g−1 (Supplementary Fig. 7). The practical application potential of our negative electrode is further demonstrated using Li-MSSA foils of varying thicknesses (5 to 100 μm), which achieve a consistent Li utilization of 89% alongside an elevated areal capacity of up to 20 mAh cm−2, compared with 4 mAh cm−2 typically used in Li-ion and Li-metal batteries (Fig. 2e and Supplementary Fig. 8). Therefore, the Li-MSSA electrode provides a practical route to meet diverse electrode capacity requirements for practical battery applications.

To underscore the importance of a multicomponent design — comprising at least five elements — for achieving highly reversible cycling at high Li mass ratio and high Li utilization, we prepared a comparison series of quaternary alloy negative electrodes. Specifically, we synthesized Li-MgCdAl, Li-AgCdAl, Li-AgMgAl and Li-AgMgCd. Experimental results revealed that the the Li-MSSA (3100 mAh g−1) exhibited a higher Li storage capacity compared to the Li-based quaternary alloys, achieving capacities of 2620, 2490, 2650, and 2120 mAh g−1 for Li-MgCdAl, Li-AgCdAl, Li-AgMgAl, and Li-AgMgCd, respectively (Supplementary Fig. 9). The CE, reflecting the electrochemical reversibility of the Li-MSSA, averages at 99.4 ± 0.2 % over 500 cycles without electrolyte optimization, surpassing its counterparts (Fig. 2f). Under identical conditions, a Li/Cu half-cell (prepared by laminating a Li foil onto Cu) delivers an average CE = 98.3 ± 0.5% over the first 200 cycles; thereafter the CE drops sharply, consistent with dendritic soft-shorting, supporting the practical relevance of the Li-MSSA design over pure Li. These findings illustrate the electrochemical performance of our Li-MSSA electrode, ensures a high Li mass ratio in the alloy along with high Li utilization and highly reversible cycling.

2.2 Elemental distribution and morphological evolution of Li-MSSA during lithiation-delithiation cycling

Given that the electrochemical performance of the Li-MSSA negative electrode is intricatedly linked to its lithiation behavior, we investigated the distribution of the constituent elements throughout the depth of the cycled Li-MSSA in its lithiated state. Ex-situ time-of-flight secondary ion mass spectrometry (ToF-SIMS) 3D reconstructions were acquired on Li-MSSA electrodes after 10 full delithiation-lithiation cycles at 2 mA cm−2, with 6 mAh cm−2 stripped and 6 mAh cm−2 plated per half-cycle. The ToF-SIMS 3D images (Fig. 3a) of the fully cycled Li-MSSA exhibt nearly overlapping signals corresponding to Li, Cd, Ag, Mg, and Al, thereby indicating a uniform distribution of these elements. This finding is further confirmed by large-area SEM-EDX mapping of the cycled Li-MSSA (Supplementary Fig. 10). The uniform distribution of elements within the cycled Li-MSSA closely resembles the highly reversible delithiation-lithiation behavior of the Li-MSSA characterized by a high Li mass ratio and utility, which collectively contribute to a significant breakthrough in reversible capacity and cycling stability (Fig. 2). To test whether Li-MSSA remains a solid solution after full delithiation, we also acquired atomically resolved STEM-EDX and Li-K-edge EELS on delithiated electrodes (Supplementary Fig. 11). The Ag, Cd, Mg, and Al maps remain co-spatial with no nanoscale segregation or intermetallic domains observed, and the residual Li-K signal after delithiation is likewise spatially uniform across the mapped area (Supplementary Fig. 12). In addition, we further employed hybrid Monte Carlo and molecular dynamics (MCMD) simulations to evaluate phase stability and elemental mixing during lithiation/delithiation across representative Li contents (Supplementary Figs. 13 and 14). Across multiple randomized initial configurations, 300 K trajectories (0.5 ns) show no tendency toward elemental clustering either after lithiation or after delithiation; the alloy remains compositionally uniform within the simulation cell.

Furthermore, the uniform vertical distribution of Li, Cd, Ag, Mg, and Al elements in the cycled Li-MSSA indicates that freshly deposited Li during lithiation does not accumulate on the foil surface. Rather, it is likely that Li atoms diffuse within the Li-MSSA foil, a finding that aligns with previous studies on solid-solution alloys, such as Li-Ag and Li-Mg.11,39 Given the high electronic conductivity of the Li-MSSA, Li+ from the electrolyte should undergo reduction to formulate Li atoms at the electrolyte/Li-MSSA interface prior to diffusion into the interior. The diffusivity of Li atoms within the Li-MSSA plays a critical role in facilitating the inward transfer of Li into the metal foil. Because the solid-solution alloy has a finite Li solubility, the bulk Li chemical potential varies continuously with composition and transport is therefore chemical-potential-driven.10 We therefore employed a non-ideal, chemical-potential-based diffusion formulation to fit the full voltage transients of the GITT dataset, using 0.01 C pulses of 30 s with 1800 s rests between pulses (Supplementary Fig. 15).40 Under this protocol, the fitted Li chemical diffusivities are around 6  10−9 cm2 s−1 (Supplementary Fig. 16). As expected, applying other analysis routes to the same dataset — such as the traditional GITT early-time √τ approximation, or EIS-based diffusivity extraction — yields values up to one or two orders of magnitude higher relative to our non-ideal, chemical-potential-driven fit.40 Notably, these values are higher than the reported room-temperature self-diffusion of Li in metallic Li (typically 10−11 cm2 s−1).41 Taken together, the fitted D is sufficiently high to accommodate the per-pulse Li flux by bulk insertion into Li-MSSA. We believe that the elevated diffusion coefficient of Li atoms in Li-MSSA is attributable to the diverse pathways available for Li atom diffusion, which arise from the increased configurational entropy of the Li-MSSA. With a greater variety of configurations, multiple diffusion pathways, presumably characterized by lower energy barriers, are available, thereby enhancing the kinetics of Li atom diffusion.42 Therefore, it is the high Li atom diffusion coefficient inherent in Li-MSSA that facilitates Li atom transfer into the foil, as opposed to surface plating of Li during lithiation process, despite the the framework mass constituting only ~10 wt.% of the entire electrode.

The inward transfer of Li atoms into the Li-MSSA foil effectively inhibits Li-dendrite growth, which markedly contrasts with the behavior observed in conventional Li metal. This is evidenced by the in-situ electrochemical optical microscope (Fig. 3c,d and Supplementary Movies 1 and 2). Cells configured as Li||Li and Li||Li-MSSA were assembled with an areal capacity of 5 mAh cm−2 deposited on the working electrodes. The formation of Li dendrites on the conventional Li metal foil was observed (Fig. 3c and Supplementary 17), revealing heterogeneous growth during the deposition of metallic Li. In contrast, the Li-MSSA electrode exhibited a smooth surface after lithiating, showing only variations in thickness with no uncontrollable mossy Li deposition (Fig. 3d and Supplementary Fig. 17). These results indicate that the Li-MSSA facilitates the descent of freshly generated Li atoms at the surface into the metal foil, effectively preventing surface deposition induced Li dendrite growth.

2.3 Structure evolution of Li-MSSA in the lithiation-delithiation cycles

We investigated the structural evolution of the Li-MSSA during the delithiation-lithiation process using in-situ X-ray diffraction (XRD). Our observations revealed a transformation in the Li crystallography of the Li-MSSA, shifting from (200) orientations to (110) orientations (Fig. 4, Supplementary Fig. 18, and Supplementary Note 4). Initially, the Li-MSSA exhibited the (200) and (211) peaks at 52.0° and 65.0°, respectively. As delithiation progressed during the first cycle, the intensities of these peaks gradually diminished due to the reduction in Li content (Fig. 4a and 4c). However, upon the completion of re-lithiation, the intensity of the (200) peak was significantly weaker compared to its initial state (Fig. 4a). Conversely, a diffraction peak of (110) at 36.2° emerged as the most prominet signal at the end of lithiation process (Fig. 4a). Following four cycles of delithiation-lithiation, we observed a marked transformation in the Li crystallography of the Li-MSSA, changing from predominantly (200) orientations to (110) orientations, with the (200) peak becoming undetectable (Fig. 4b). In comparison, the crystallography of the control sample, Li metal coated on a Cu foil (Li/Cu), which initially showed (110), (200), and (211) orientations, remained relatively stable even after prolonged cycling (Fig. 4c and 4d). Furthermore, the electron backscatter diffraction (EBSD) analysis elucidated that, in addition to the bulk of Li-MSSA (Fig. 4b), the surface plane of the Li-MSSA foil under went a transformation from (200) to (110) (Fig. 4e and 4f). In contrast, the Li/Cu negative electrode consistently exhibited a preference for (200) orientations (Fig. 4g and 4h).

Previous thermodynamic studies suggest that the (110) facets exhibit faster lithiation kinetics than the (200) and (211).43,44 The enhanced lithiation kinetics of our Li-MSSA is evidenced by the exchange current density measurement (Supplementary Fig. 19), whereby the Li-MSSA exhibits an exchange current density almost threefold higher (2.13 mA cm−2) than the Li/Cu negative electrode (0.77 mA cm−2). The (110) facets, which facilitate accelerated lithiation kinetics, are expected to diminish the overpotential and reduce the driving force for side reactions, consequently prolonging the cycling life of LMBs.

Unlike the widely used “lithiophilic site” strategy — where alloy additives mainly lower the Li nucleation overpotential to promote more uniform Li deposition45-47 — our alloy is designed to promote an inward-growth mode, in which Li atoms migrate into the bulk rather than accumulate at the interface10,11. In the lithiophilic approach, once a thin Li layer forms, subsequent cycling still proceeds by Li+ reduction on Li metal at the electrolyte interface; the locus of deposition does not fundamentally change. By contrast, our multicomponent solid solution operates as a single phase with finite Li solubility, so the Li chemical potential varies continuously with composition. Under plating, this establishes a chemical-potential gradient that drives freshly formed Li atoms at the surface to diffuse inward rather than accumulate at the interface — consistent with solid-solution-assisted inward growth.10,11 Kinetically, multicomponent solid-solution alloys (≥5 elements) possess high configurational entropy that induces significant local chemical and strain disorder, thereby broadening the distribution of Li site energies. This, in turn, generates multiple, percolating diffusion pathways and enhances Li mobility — our rationale for employing a multicomponent alloy.42,48,49 Consequently, our alloy sustains percolating, quasi-isotropic bulk pathways that provide sufficient subsurface transport to accommodate the interfacial Li flux, thereby preventing interfacial supersaturation. Together with the thermodynamic drive from continuous chemical potential (μ), this bulk transport relocates Li uptake into the foil interior and suppresses surface Li deposition. Consistent with this mechanism, our in-situ XRD, TOF-SIMS, and STEM-EDX/EELS indicate a single-phase solid solution across the cycling window.

2.4 Electrochemical performance of Li-MSSA in full cells

Our investigations into the Li-MSSA during the lithiation-delithiation cycles, which include an analysis of element distribution (Fig. 3a), Li atom diffusion (Supplementary Fig. 16), and structure evolution (Fig. 4), demonstrate that the Li-MSSA not only facilitates the inward-transfer of freshly generated Li atoms into the alloy metal foil, with a thickness of tens of micrometers, but also promotes the development of preferred (110) facets. These two mechanisms are, at least in principle, self-reinforcing, as they mitigate the Li-dendrite growth (Fig. 3d) and reduce side reactions between the electrode and electrolyte (Fig. 2f), thereby enabling the Li-MSSA to achieve a high specific capacity (Fig. 2d) and maintain stable cycling performance (Fig. 2f) in Li metal negative electrodes concurrently.

The Li-MSSA negative electrode exhibits stable cycling and a high capacity, making it suitable for long-term operation in practical LMBs. To facilitate industrial-level evaluation of our Li-MSSA, we designed an Ah-level Li-metal pouch cell featureing rigorous parameters, including an electrolyte usage of 2 g Ah−1, thereby transitioning the Li-MSSA negative electrode into practical applications. This Li-MSSA||LiNi0.8Co0.1Mn0.1O2 pouch cell achieves an specific energy of 385 Wh kg−1 (based on the total mass of the pouch cell) while sustaining a capacity retention exceeding 82% after 600 cycles (Fig. 5a; Supplementary Fig. 20; Supplementary Table 4). The charge-discharge curves associated with the Li-MSSA||LiNi0.8Co0.1Mn0.1O2 pouch cell demonstrate stable behavior over 600 cycles (Fig. 5b). The negligible voltage shifts (i.e., reduced polarization) indicate that the parasitic reaction on the Li-MSSA electrode are significantly mitigated.24 Furthermore, based on the performance metrics of this pouch cell (Fig. 5c), the specific energy is estimated to reach 462 Wh kg−1 (based on the total mass of the pouch cell) in a 10 Ah pouch cell (Supplementary Table 5). Li-MSSA-based metal batteries show a combined specific energy and cycling life under the tested conditions, compared with cells using conventional modified Li-metal negative electrodes (Fig. 1b and Supplementary Fig. 21). To evaluate the rate capability of Li-MSSA, we assembled Li-MSSA||LiNi0.8Co0.1Mn0.1O2 and Li/Cu||LiNi0.8Co0.1Mn0.1O2 full-cells under identical positive electrode loadings (10.8 mg cm−2, Supplementary Fig. 22). C-rates were defined with respect to the positive electrode (1 C = 209 mA g−1, based on a practical capacity of 209 mAh g−1). The Li-MSSA||LiNi0.8Co0.1Mn0.1O2 cell delivered 208 (0.1 C), 201 (0.2 C), 190 (0.5 C), 175 (1 C), 144 (2 C), 105 (4 C), and 68 mAh g−1 (6 C). By contrast, Li/Cu||LiNi0.8Co0.1Mn0.1O2 cell delivered a lower capacity of 25 mAh g−1 at 6 C under the same conditions, indicating improved high-rate performance for Li-MSSA (Supplementary Fig. 22). To assess rate-dependent structural evolution of the Li-MSSA negative electrode, we performed ex-situ STEM-HAADF imaging with high-resolution STEM-EDX mapping on Li-MSSA after rate cycling up to 6 C (Supplementary Fig. 23). Elemental maps for Ag, Cd, Mg, and Al are co-localized within our spatial resolution, with no detectable nanoscale segregation or heavy-element clustering. Collectively, the data indicate that (i) under matched full-cell conditions, Li-MSSA delivers a higher usable capacity than pure Li at elevated rates, and (ii) after high-rate cycling (up to 6 C) the negative electrode shows no evidence of rate-induced phase separation or nanoscale compositional partitioning. In conventional Li-MSSA||LiNi0.8Co0.1Mn0.1O2 full cells, the alloy negative electrode experiences a lithiation-first step upon the initial charge; Li-MSSA therefore accommodates incoming Li interstitially within the matrix. To verify that the structural evolution during charge/discharge follows the same single-phase mechanism analyzed in Figs. 2-4, we harvested Li-MSSA foils from NCM811 full cells at end of charge (lithiated) after 50 cycles and performed cross-sectional SEM-EDX mapping. As shown in Supplementary Fig. 24, Ag, Cd, Al and Mg are spatially uniform across the foil thickness. It is reasonable to assume that if Li metal was formed and deposited at the foil surface, the cross-sectional SEM-EDX would be expected to show a near-surface depletion band of Ag/Cd/Mg/Al in the immediately underlying alloy region due to dilution/redistribution. Neither signature is observed, supporting inward Li diffusion and homogeneous solid-solution lithiation rather than superficial Li plating. To further decouple positive electrode-derived Li, we rebuilt a Li-MSSA||delithiated-NCM811 full cell in which all active Li originates from the alloy (Supplementary Fig. 25). Commercial NCM811 electrodes (areal capacity 2.3 mAh cm−2) were first charged to 4.3 V vs. Li/Li+ to extract Li (Supplementary Fig. 25a). The fully delithiated NCM811 discs were then assembled opposite Li-MSSA foils (areal capacity 2.5 mAh cm−2) in fresh CR2032 cells (Supplementary Fig. 25b). Because the positive electrode now contains no Li, the cell begins at N/P = 1.1, providing a stringent platform to assess alloy utilization and cycling stability without surplus positive electrode Li. As a control, Li/Cu||delithiated-NCM811 cells were prepared with areal capacity and N/P ratio matching the Li-MSSA||delithiated-NCM811 cells. Under identical conditions, the Li-MSSA||delithiated-NCM811 cell delivered 2.3 mAh cm−2 initially and retained 2.1 mAh cm−2 after 100 cycles, whereas the Li/Cu control showed fluctuating CE and rapid capacity fade (Supplementary Fig. 25c ). These results show more stable Li utilization in Li-MSSA compared with the control and suggest that the alloy maintains stable behavior during lithiation-first operation even under low-Li-excess, high-loading conditions. In light of the potential offered by emerging high-energy-density positive electrode materials such as sulfur, Li||S batteries demonstrate promising prospects of achieving cell-level specific energies exceeding 500 Wh kg−1 (based on the total mass of the pouch cell). Further more, we have developed Li||S batteries featuring high areal capacities of 3.5 and 7.4 mAh cm−2 (Fig. 5d and 5e) alongside a low negative electrode-to- positive electrode capacity ratio of 1.1. Both batteries showed minimal capacity degradation, thereby validating the advantages of Li-MSSA and showcasing its practical viability for achieving durable high-loading Li-based cells. It should also be noted that Li-MSSA is compatible with optimized liquid and solid-state electrolytes, in addition to high-voltage high-capacity positive electrodes such as Li-rich manganese positive electrode material, thus affirming the critical role of Li-MSSA in practical batteries boasting specific energies exceeding 500 Wh kg−1 (based on the total mass of the pouch cell) and lifespans of 1000 cycles.

Methods
Material Synthesis
The Li-MSSA composite was synthesized by heating a mixture of Li, Cd, Ag, Mg, and Al metallic powders to 600 °C at a ramp rate of 5 °C min−1, followed by holding at 600 °C for 2 h under mechanical stirring in an argon-filled glovebox, where the O2 and H2O levels were maintained below 0.01 ppm. The final composite comprises 90 wt.% lithium, with Cd, Ag, Mg, and Al constituting the remaining 10 wt.% in an equimolar ratio (1:1:1:1). Subsequently, the molten mixture was rapidly cooled to −196 °C by immersion in liquid nitrogen, resulting in a dense, grey solid. This solid was then sectioned and rolled to the desired thickness in a dry room with a dew point below −40 °C. Notably, for film thicknesses below 10 μm, rolling was not feasible; instead, a melting and casting approach was employed to achieve the specified dimensions. Lithium foil (Li, 99.9%, Sigma-Aldrich), cadmium granules (Cd, ≥99%, Aladdin; purchased via the CASMART procurement platform), silver powder (Ag, ≥99.99% trace metals basis, Sigma-Aldrich), magnesium powder (Mg, ≥99.9% trace metals basis, Sigma-Aldrich), and aluminum foil (Al, ≥99.99% trace metals basis, Sigma-Aldrich) were used as received unless otherwise noted.

Structural characterization
The surface and cross-sectional morphology of the electrodes were analyzed using SEM (FEI Co., Ltd.) at an operating voltage of 10 kV. HADDF-STEM images were obtained by a JEOL JEM-ARF200F TEM/STEM instrument with a spherical aberration corrector operated at 200 KV. For Figure 2a, STEM mapping was acquired at an electron accelerating voltage of 300 kV with a beam current of 78 pA, a dwell time of 2.0 μs per pixel, and a 1024 × 1024 pixel matrix over an 8.6 × 8.6 nm field of view. The total acquisition time for the image was 205 s, corresponding to a cumulative electron dose of 1.34 × 107 e− Å−2. For Figure S11, STEM mapping was acquired at 300 kV with a beam current of 78 pA, a dwell time of 2.0 μs per pixel, and a 1024 × 1024 pixel matrix over a 6.1 × 6.1 nm field of view. The total acquisition time was 188 s, and the cumulative electron dose was 2.44 × 107 e− Å−2. For Figure S23, STEM mapping was acquired at 300 kV with a beam current of 78 pA, a dwell time of 2.0 μs per pixel, and a 1024 × 1024 pixel matrix over a 12.2 × 12.2 nm field of view. The total acquisition time was 446 s, giving a cumulative electron dose of 1.45 × 107 e− Å−2. For Li STEM-EELS mapping, data were acquired at an accelerating voltage of 300 kV with a beam current of 100 pA, a dwell time of 0.005 s per pixel, and an 86 × 80 pixel matrix over a 694.3 × 645.9 nm field of view. The total acquisition time was 34.4 s, corresponding to a cumulative electron dose of 4.79 × 102 e− Å−2. The crystal structures of the electrodes were characterized via X-ray diffraction (XRD) (Rigaku Co., Ltd.) with Cu-Kα radiation (λ = 1.5418 Å) at 40.0 kV and 1 mA. To protect moisture-sensitive samples, a custom transfer box was employed. Elemental distributions on the electrodes and depth profiles were obtained through time-of-flight secondary ion mass spectrometry (TOF-SIMS), conducted over a 50 μm × 50 μm area in positive-ion mode, utilizing a Cs⁺ sputtering gun. With an etching rate calibrated at approximately 0.4 μm per minute (based on lithium metal foil etching rate), the sample thickness was estimated by multiplying the etching rate by the total etching time. In-situ electrochemical optical microscopy was conducted in a commercial transparent cell (Beijing Scistar) with a glass window. Each cell contained 1.0 mL of electrolyte (1.0 M LiTFSI in DME:DOL, 1:1 v/v, with 2 wt.% LiNO3). Two configurations were examined: Li‖Li and Li‖Li-MSSA. In the Li‖Li-MSSA cells, the working electrode was a Li-MSSA alloy film, and a Li-metal foil served as the counter/reference; in the Li‖Li controls, a Li-metal foil was used as the working electrode. During the optical tests, an areal capacity of 5 mAh cm−2 was deposited on the working electrodes. In-situ XRD measurements were performed using a LIB-XRD cell (Beijing Scistar Technology Co., Ltd.) equipped with a beryllium window to permit X-ray penetration. The alloy electrode, with a diameter of approximately 16 mm, was cast on the backside of the beryllium window. This in-situ cell was cycled at a current density of 1 mA cm−2.

Electrochemical measurements
All electrochemical measurements were carried out in an environmental chamber maintained at 25.0 ± 0.5 °C and were recorded using LAND CT2001A and NEWARE battery testers. A commercial microporous polypropylene (PP) separator (Celgard 2400, Celgard) was used. The separator has a nominal thickness of 25 μm, a porosity of ~41%, and an average pore size of ~43 nm. For CR2032 coin cells, the separator was punched into 16 mm-diameter disks. To evaluate the capacity, CE, and cycling stability of Li-MSSA, CR 2032 coin cells were assembled with Li-MSSA or bare-Cu working electrodes (10 mm diameter), a Li-metal counter/reference (10 mm diameter, 300 μm thick), a Celgard 2400 separator, and 60 μL of 1.0 M LiTFSI in DME:DOL (1:1, v/v) containing 2 wt.% LiNO3. Polyacrylonitrile (PAN, M = 150,000, Sigma-Aldrich) and sulfur (Sigma-Aldrich) were used to synthesize sulfurized polyacrylonitrile (SPAN). PAN and sulfur were first mixed in an agate mortar at a mass ratio of 1 : 4, and then heated in an argon-filled furnace to 450 °C and held for 6 h at a ramping rate of 10 °C min−1. For the SPAN positive electrode slurry, SPAN, Super P, poly(acrylic acid) (PAA, average M ≈ 450000, Sigma-Aldrich), and guar gum (Sigma-Aldrich) were dispersed in deionized water using a high-speed homogenizer at a mass ratio of 8 : 1 : 0.5 : 0.5. The slurry was coated onto a 10 μm Al current collector and vacuum-dried at 80 °C for 12 h. The electrode loading was controlled by adjusting the coating thickness. For Li||S CR2032 cells, we used a Celgard 2400 separator and 60 μL of 2.0 M LiFSI in n-butyl methyl ether as the electrolyte. Cells were cycled at 0.06 C for high-loading positive electrodes (7.4 mAh cm−2) and 0.2 C for low-loading positive electrodes (3.5 mAh cm−2). Here, 1 C is defined as 1675 mA g−1. The pouch cells were fabricated on a semi-automated pilot line in a dry room (dew-point -55 °C, RH < 0.1 %). In the 1 Ah pouch cells we use a free-standing Li-MSSA foil as both active layer and current collector; no additional Cu substrate is employed. The foil is prepared by hot-rolling a cast Li-MSSA ingot to a thickness of 40 µm, then slit, calendared, and spot-welded to a Ni tab before stacking. The Li-MSSA negative electrode provides an areal capacity of 7.8 mAh cm−2. Commercial NCM811 positive electrode sheets were purchased from Kejing Materials Technology Co., Ltd. The positive electrode contained 96.8 wt.% NCM811. The NCM811 positive electrode was double-side coated onto 10 μm Al foil, delivering 3.5 mAh cm−2 per side (7.0 mAh cm−2 total). Positive electrode and negative electrode sheets were Z-stacked with 12 µm polypropylene separators, tab-welded (Al for the positive electrode, Ni for the negative electrode), packaged in 85 mm × 65 mm laminated pouches, vacuum-dried (60 °C, 12 h), and vacuum-injected with 1.5 mol L−1 LiFSI in 1,2-dimethoxyethane/1,1,2,2-tetrafluoroethyl-2,2,3,3-tetrafluoropropyl ether (1.2 : 3, v/v) at an electrolyte-to-capacity ratio of 2 g Ah−1. After two C/20 formation cycles, the cells were cycled between 2.6 and 4.3 V using a LAND battery tester, with a 0.1 C charge and 1 C discharge; during cycling, the cells were connected to the tester via leads and placed in an environmental chamber maintained at 25.0 ± 0.5 °C. Pouch cells were cycled under a constant external stack pressure of ~0.2 MPa using a plate-and-clamp fixture. After stacking, electrolyte was injected through the fill port and the pouch was pre-sealed after a wetting/rest period. Following formation, the pouch was vacuum-degassed, and then finally heat-sealed (with the gas pouch trimmed off).

Computational methods
We investigated the distribution trends of each component element in the unit cell of the Li-Ag-Al-Cd-Mg system under charged and discharged states using the GPUMD software package: first, we constructed single-crystal models consistent with experimental observations, featuring quasi-random element arrangements, for the charged state (with an element ratio of Li:Ag:Al:Cd:Mg = 317:8:8:8:8) and the discharged state (with an element ratio of Li:Ag:Al:Cd:Mg = 345:1:1:1:1), where each model contains a total of 9423 atoms; subsequently, we performed simulations at 300 K using the canonical ensemble hybrid Monte Carlo and molecular dynamics (MCMD) algorithm and the NEP89 potential based on neuroevolution potential architecture to achieve system equilibrium and sampling; after the simulations were completed, we determined whether the system tends to form a solid solution with uniform element distribution by analyzing the variation of potential energy with MCMD simulation time and selecting the configuration of the last simulation frame for element distribution visualization. The initial and final configurations from the MCMD simulations are provided in Supplementary Data 1.

Peer review rounds
[{"round": 1, "reviewer_comments": {"reviewer_1": "This work reports a Li metal-based multicomponent solid-solution alloy (Li-MSSA) anode that achieves a high reversible capacity of 3100 mAh g⁻¹ and enables pouch cells with an energy density of 470 Wh kg⁻¹ and 82% capacity retention over 600 cycles. The alloy design leverages high Li diffusivity and facet engineering to suppress dendrite formation and promote uniform Li deposition. While alloy-based substrates for stabilizing Li metal anodes are a well-established research direction, the reported performance is highly impressive. Therefore, it is critical to further elucidate the underlying mechanisms. I would be happy to support the publication of this work if the authors can address the following concerns.\n1. The authors claim that MSSA does not undergo phase separation. However, the in situ XRD data lack the alloy-related peaks at 32.5°, 39.1°, and 46.8°. Is there any in situ or ex situ XRD data that can confirm the alloy phase remains unaffected after both lithiation and delithiation?\n2. The HAADF-STEM image in Figure 2b and the corresponding intensity profile in Figure 2c are used to distinguish between \"heavy\" and \"light\" atomic columns. Given that HAADF-STEM contrast scales approximately with atomic number (Z), one would expect to observe three distinct levels of intensity: (1) the lowest for Li (Z = 3), (2) intermediate for Mg (Z = 12) and Al (Z = 13), and (3) the highest for Ag (Z = 47) and Cd (Z = 48). However, Figure 2c does not clearly show such a gradation. It is therefore difficult to unambiguously assign specific intensity peaks or valleys to particular atomic species, especially for light elements like Li.\n3. The schematic in Figure 2d shows Li and other atoms explicitly, but it is not evident how their positions were determined from the experimental data. It is recommended that complementary techniques such as EELS be employed to verify the presence and distribution of Li and thereby strengthen the interpretation.\n4. Figure 2h suggests that all four metal elements in MSSA are essential for performance. Could the authors provide further insights into this result? Given that these metals are known to be lithium-philic and have the potential to individually shown good cycling stability, additional discussion would help clarify the observed synergy.\n5. The lithium-philic behavior of most metals is often demonstrated under high-temperature molten conditions. Could the authors clarify how such a property is retained or manifested under room-temperature electrochemical cycling? While MSSA may exhibit strong lithium affinity in its molten state, it would be helpful to understand whether this property remains relevant during practical battery operation.", "reviewer_2": "In this paper, the authors present a multicomponent lithium metal anode with 90% lithium and four other metallic components. The authors show good cycling behavior, without significant concentration gradients that form. The concept is interesting and the characterization of the material is good. The cycling behavior is very good. Despite the promise, however, there are a number of issues that should be addressed before publication, including seemingly misleading statements regarding specific energy and insufficient comparison to control experiments. Please see notes below.\n-What is the rate behavior of the multicomponent material compared to pure Li in full cells? Rate behavior and rate effects on structural evolution are not discussed in the manuscript.\n-All details of cell cycling should be given in the captions for each figure. For instance, what is specific electrolyte used for pouch cells? Two are mentioned in methods. Authors need to be much more careful about providing all details for individual experiments – right now it is muddled and difficult to determine the details.\n-pouch cell fab and testing info is not provided in the methods.\n-important: in the abstract and conclusion, it is stated that the paper demonstrats 470 Wh/kg. This is not true, as it is only extrapolated from a 1-Ah cell to a 10 Ah-cell. Please remove this claim and please claim only accurate demonstrations.\n-the superlatives (“exceptional” “unprecedented”) sprinkled throughout the paper are unnecessary.\n-on page 7 it is stated that the capacities of the alloy surpasses that of “state of the art Li metal electrodes.” Why would this be the case? Li metal has higher theoretical capacity and can be fully stripped. This is a confusing statement.\n-on page 8 there is lots of discussion about stripping and cycling of the alloy electrodes, but they are never compared to pure Li electrodes. Direct comparison to pure Li tested under the same conditions would be very helpful for the reader to understand the true benefits of the alloy.\n-In fig 3b, why does the cell voltage decrease during GITT? Shouldn’t it be a constant plateau as in Fig 2?\n-Are all the assumptions for GITT to be used to calculate diffusion coefficients satisfied? This is generally not the case for most works, and extracted diffusion coefficients are wrong. See https://pubs.acs.org/doi/10.1021/acsaem.1c02218\n-what is the capacity stripped/plated in Fig 3a?\n-does the MSSA material remain a solid solution even as lithium is fully removed, or do compounds form? Are there other peaks that show up in XRD that are not shown in the limited 2theta range of Fig 4?\n-In the XRD data in fig 4, why does the MSSA material have a higher extent of 200 peak displayed compared to pure Li before stripping? Is this because of different heat treatments of the different materials?\n- The XRD pattern of the Li-MSSA anode in Fig. S2 shows peaks at 32.5°, 39.1°, and 46.8° that could be from the metallic compound phases. The existence of metallic compounds means that the Li-MSSA is not a uniform solid-solution. These XRD peaks are not shown in the limited 2theta range of Fig. 4. How do the metallic compound phases evolve during delithiation-lithiation cycling?\n- In full cells with LiNi0.8Co0.1Mn0.1O2 cathodes, the Li-MSSA anode will be lithiated first during the charging step of each cycle, while the characterizations and tests in Figs. 2, 3, 4 are based on the delithiation-lithiation cycling of Li-MSSA from Li88M to Li9.9M. How does the lithiation-first mode affect the structure evolution and electrochemical behavior of Li-MSSA?\n-How was the Li utilization of the Li-MSSA anode calculated when it is paired with LiNi0.8Co0.1Mn0.1O2 cathodes? This is a bit confusing.\n- For the calculation of energy densities of the 1 Ah pouch cell in Table S4, the Li-MSSA anode doesn’t have a current collector. Is just Li-MSSA foil used as the anode in the 1 Ah pouch cell? How is the thickness of the Li-MSSA anode used for the pouch cell? The details of pouch cell fabrication should be provided in the methods.\n- Line 50, “Li-Ag binary solids-solution alloy” should be “Li-Ag binary solid-solution alloy”.", "reviewer_3": "I co-reviewed this manuscript with one of the reviewers who provided the listed reports. This is part of the Nature Communications initiative to facilitate training in peer review and to provide appropriate recognition for Early Career Researchers who co-review manuscripts.", "reviewer_4": "In this paper the authors report the synthesis of a lithium alloy containing Cd, Ag, Mg and Al and test its performance as an anode material. The alloy composition developed is novel and the authors claim it to form a solid solution with a high lithium diffusivity that enables a high lithium utilisation of 89%. This alloy anode is then implemented in pouch cells with various cathode materials. The results presented in this manuscript and its supplementary information do not convincingly support all of these claims, so further evidence and clarification will be required before it can be considered for publication. The following points should be addressed:\nThe authors claim the Cd, Ag, Mg and Al all form a solid solution with the lithium and the main evidence for this is a single TEM image with a narrow field of view (Fig. 2a). The XRD pattern in Fig. S2 does not support this claim and the authors acknowledge the peaks at 32.5, 39.1 and 46.8° can be assigned to “metal alloy phases of the MSSA.” This suggests the alloy contains additional phases and does not simply form a solid solution. The additional phase(s) should be identified and microstructural characterisation should be performed with a larger-field-of-view using SEM, for example.\nAdditionally, the lithium concentration in the electrode will decrease significantly during discharge (to approximately 10 wt.% using the reported utilisation). How does this affect the phases present in the alloy and their distribution? Fig. S7 shows clear segregation of Al in the alloy after cycling.\nIt is unclear how the practical gravimetric capacity of the alloy electrode is determined. The manuscript states that “Li was extracted from the Li-MSSA within an electrochemical cell comprising Cu foil as the counter electrode,” but the voltage curves presented in Fig. 2f and S5 do not appear to be consistent with this. Please could additional details be provided to describe how these measurements were conducted?\nSimilarly, insufficient details have been provided about how the GITT measurement was conducted and how the data was analysed. The pulse current density, pulse duration and rest period duration should be stated. It would also be useful to relate the time to a composition/state-of-charge of the alloy electrode.\nFinally, the claim that an N/P ratio of approximately 1.1 is “rigorous” could be misleading. While an N/P ratio of 1.1 would be suitable for testing a lithium-ion battery, in this case both the cathode and anode are fully lithiated during cell assembly so all the lithium contained in the anode will be an excess. This is likely to inflate the cycling performance presented in Fig. 5. Experimental details describing pouch cell preparation are also missing from the methods section and it is unclear why no anode current collector has been included in the calculations in Tables S4 and S5."}, "author_response": "We have carefully considered all valuable comments from the reviewers and have made suitable revisions accordingly. The reviewers’ comments and our responses are listed as follows:\nReviewer #1\nThis work reports a Li metal-based multicomponent solid-solution alloy (Li-MSSA)\nanode that achieves a high reversible capacity of 3100 mAh g-1 and enables pouch cells\nwith an energy density of 470 Wh kg-1 and 82% capacity retention over 600 cycles. The\nalloy design leverages high Li diffusivity and facet engineering to suppress dendrite\nformation and promote uniform Li deposition. While alloy-based substrates for\nstabilizing Li metal anodes are a well-established research direction, the reported\nperformance is highly impressive. Therefore, it is critical to further elucidate the\nunderlying mechanisms. I would be happy to support the publication of this work if the\nauthors can address the following concerns.\nResponse: We express our sincere appreciation to Reviewer #1 for the meticulous\nreview of our manuscript and the invaluable comments provided. These comments have\nplayed a crucial role in the continuous improvement of our manuscript. We have\nthoughtfully revised the manuscript in direct response to these insightful suggestions.\nWe also extend our gratitude to Reviewer #1 for generously dedicating precious time\nand expertise to assessing our work.\nComment 1: The authors claim that MSSA does not undergo phase separation.\nHowever, the in situ XRD data lack the alloy-related peaks at 32.5°, 39.1°, and 46.8°.\nIs there any in situ or ex situ XRD data that can confirm the alloy phase remains\nunaffected after both lithiation and delithiation?\nResponse: Thank you for this helpful comment. After re-examining our data, we\nidentified a misassignment in the original Figure S2: the reflections near 32.5°, 39.1°,\nand 46.8° are best ascribed to segregates formed during non-equilibrium solidification\nrather than to intrinsic features of a homogeneous alloy matrix. In other words, these\nsignals are parasitic and do not represent a distinct bulk alloy phase.\nFrom a metallurgical standpoint, compositional microsegregation is expected to\nsome extent whenever solidification proceeds under non-equilibrium conditions,\nproducing localized solute fluctuations on ~10 – 100 µm length scales (Castings\nPrinciples: The New Metallurgy of Cast Metals 2nd ed, 2003). Such microsegregation\ncan be substantially reduced by post-solidification homogenization below the solidus,\n1\nfollowed by rapid cooling to retain the uniform solid solution (Porter & Easterling,\nPhase Transformations in Metals and Alloys, 1981). In our original synthesis, the Li-\nMSSA melt was simply quenched from 600 °C, which suppresses long-range diffusion\nbut can freeze-in interdendritic microsegregation. To mitigate this artefact, we re-\nprepared Li-MSSA and introduced a mild subsolidus homogenization (200 °C, ~24 h)\nbefore rapid cooling to retain the solid solution. Ex-situ XRD of the homogenized\nmaterial shows only Li-like bcc reflections at 36.2°, 52.0°, and 65.0° (Figure R1; also\nFigure S2 in the revised SI), with no additional peaks attributable to segregated alloy\nphases. XRD of the as-prepared Li-MSSA shows (200) as the most intense reflection,\nwhereas (110) is the most intense reflection for Li metal (PDF No. 1-1131). The\nenhancement of (200) reflection in Li-MSSA is explained by a {100} fiber texture\nwithin a Li-like bcc MSSA, consistent with our electron backscatter diffraction (EBSD)\npole figures (Figure 4e of the original manuscript) that show {100} fiber texture (axes\nnear <100>) and thus re-weight the (110)/(200)/(211) reflection intensities accordingly.\nNon-bcc alternatives do not simultaneously reproduce the observed 110/200/211\npositions and relative intensities under reasonable texture models. For a representative\ndelithiation composition Li AgMgCdAl (Li M), the average VEC (1.55) lies far\n39.6 9.9\nbelow the ~6.87 threshold that empirically favors bcc matrices in multicomponent\nalloys (J. Appl. Phys. 2011, 109, 103505), further supporting a Li-like bcc matrix. We\nnote that alloying may slightly shift the bcc lattice parameter relative to Li metal; this\ndoes not affect the texture-based explanation for the (200) enhancement. Taken together,\nthese results demonstrate that Li-MSSA is a Li-like bcc solid solution, and the prior\nsignals near 32.5°, 39.1°, and 46.8° reflect parasitic microsegregation rather than\nintrinsic alloy phases.\nFor the in-situ XRD, we re-extracted and compared the end-of-lithiation and end-\nof-delithiation patterns for cycles 1 and 4 (Figure R2; also Figure S18 in the revised SI).\nAs the Be-windowed cell introduces strong Be reflections and an elevated diffuse\nbackground in the 30° – 50° range, which markedly reduces the visibility of weak\nfeatures; under this geometry even the intense (200) near ~52.0° in ex-situ XRD is\nattenuated, so any much weaker signals at 32.5°, 39.1°, and 46.8° are not discernible.\nBeyond the principal bcc-type reflections highlighted in the main text (36.2°, 52.0°,\n65.0°), we observe in some scans (e.g., cycle 4), broad features at 34.9°, 43.1°, and\n57.3°. Owing to the multicomponent nature of the alloy, these features cannot be\nunambiguously assigned to specific intermetallic phases at present. Given the\n2\ncombinatorial phase space of multicomponent alloys, future efforts could employ\nmachine-learning-assisted phase fingerprinting (e.g., simulated pattern libraries\ncombined with Bayesian or neural-network classifiers), together with CALPHAD-\nguided searches and pair-distribution-function analyses, to enable more robust phase\nidentification. In this work, our focus is the design concept of a multicomponent solid-\nsolution alloy that leverages a continuous chemical-potential gradient in the bulk and\nhigh configurational entropy to provide multiple, percolating diffusion pathways,\nthereby enhancing Li mobility and enabling stable cycling under high Li mass fraction\nand high Li utilization conditions. Determining if the minor features are bona fide\nsecondary phases and, if so, whether they influence long-term cycling remains an\nimportant direction for future work.\nIn response to the reviewer’s suggestion, we have updated the relevant data and\nrefined the discussion accordingly in the revised SI:\n“We re-extracted and compared the end-of-lithiation and end-of-delithiation\npatterns for cycles 1 and 4. Beyond the principal bcc-type reflections (36.2°, 52.0°,\n65.0°), we observe in some scans (e.g., cycle 4), broad features at 34.9°, 43.1°, and\n57.3°. Owing to the multicomponent nature of the alloy, these features cannot be\nunambiguously assigned to specific intermetallic phases at present. Given the\ncombinatorial phase space of multicomponent alloys, future efforts could employ\nmachine-learning-assisted phase fingerprinting (e.g., simulated pattern libraries\ncombined with Bayesian or neural-network classifiers), together with CALPHAD-\nguided searches and pair-distribution-function analyses, to enable more robust phase\nidentification. Determining if the minor features are bona fide secondary phases and,\nif so, whether they influence long-term cycling remains an important direction for future\nwork.” (Page 19 in the revised SI)\n“From a metallurgical standpoint, compositional microsegregation is expected to\nsome extent whenever solidification proceeds under non-equilibrium conditions,\nproducing localized solute fluctuations on ~10 – 100 µm length scales.3 Such\nmicrosegregation can be substantially reduced by post-solidification homogenization\nbelow the solidus, followed by rapid cooling to retain the uniform solid solution.4 In\nour original synthesis, the Li-MSSA melt was simply quenched from 600 °C, which\nsuppresses long-range diffusion but can freeze-in interdendritic microsegregation. To\nmitigate this artefact, we re-prepared Li-MSSA and introduced a mild subsolidus\nhomogenization (200 °C, ~24 h) before rapid cooling to retain the solid solution. Ex-\n3\nsitu XRD of the homogenized material shows only Li-like bcc reflections at 36.2°, 52.0°,\nand 65.0°, with no additional peaks attributable to segregated alloy phases (Figure S2).\nXRD of the as-prepared Li-MSSA shows (200) as the most intense reflection, whereas\n(110) is the most intense reflection for Li metal. The enhancement of (200) reflection in\nLi-MSSA is explained by a {100} fiber texture within a Li-like bcc MSSA, consistent\nwith our EBSD pole figures (Figure 4e in the main text) that show {100} fiber texture\n(axes near <100>) and thus re-weight the (110)/(200)/(211) reflection intensities\naccordingly. Non-bcc alternatives do not simultaneously reproduce the observed\n110/200/211 positions and relative intensities under reasonable texture models. For a\nrepresentative delithiation composition Li AgMgCdAl (Li M), the average VEC\n39.6 9.9\n(1.55) lies far below the ~6.87 threshold that empirically favors bcc matrices in\nmulticomponent alloys,5 further supporting a Li-like bcc matrix. We note that alloying\nmay slightly shift the bcc lattice parameter relative to Li metal; this does not affect the\ntexture-based explanation for the (200) enhancement. Previous literatures show that\nalloying elements robustly alter deformation/recrystallization pathways — via solute\ndrag, precipitation/particle-stimulated nucleation (PSN), and stored-energy gradients\n— and thereby reshape texture and the resulting XRD peak intensities across cubic and\nhcp alloys; for example, Al-Mg alloys containing Sc/Zr form dispersoid-driven PSN\nthat changes recrystallization textures, and Mg alloys with rare-earth (or Ca/Y)\nadditions systematically weaken or re-orient basal textures while promoting non-basal\ncomponents.6-10 Therefore, we conclude that the enhanced (200) reflection in the XRD\nof Li-MSSA is best attributed to alloying-driven texture evolution within a Li-like bcc\nsolid solution.” (Page 33 in the revised SI)\nFigure R1. XRD patterns of the Li-MSSA, re-annealed Li-MSSA, and metallic Li\n(PDF No. 1-1131).\ninitial, delithiated, and lithiated. Black arrowheads mark reflections from the Be\nwindow; blue arrowheads indicate the principal Li-like bcc reflections. Green\narrowheads denote intermittent features (unassigned) observed in some scans (e.g., near\n34.9°, 43.1°, and 57.3°).\nComment 2: The HAADF-STEM image in Figure 2b and the corresponding intensity\nprofile in Figure 2c are used to distinguish between \"heavy\" and \"light\" atomic columns.\nGiven that HAADF-STEM contrast scales approximately with atomic number (Z),\none would expect to observe three distinct levels of intensity: (1) the lowest for Li (Z = 3),\n(2) intermediate for Mg (Z = 12) and Al (Z = 13), and (3) the highest for Ag (Z = 47)\nand Cd (Z = 48). However, Figure 2c does not clearly show such a gradation. It is\ntherefore difficult to unambiguously assign specific intensity peaks or valleys to\nparticular atomic species, especially for light elements like Li.\nResponse: We appreciate the reviewer’s careful reading. Our intent in Figure 2b and\n2c of the original manuscript was not to unambiguously assign individual peaks/valleys\nto specific elements, but to assess chemical uniformity at the atomic-column level.\nIn HAADF-STEM, the bright and dim spots in Figure 2b of the original manuscript\nare not single atoms but atomic columns imaged along a <110> zone axis in high-angle\nannular dark-field STEM; the integrated intensity of each column follows the well-\nestablished proportionality I ∝ ∑nZ1.7, where I is the background-subtracted integrated\ni i\nHAADF intensity of a column, n denotes the number of atoms of species i along that\ni\ncolumn within the imaged thickness, and Z is the atomic number of species i\ni\n(Ultramicroscopy 1991, 37, 14-38). Because Li-MSSA is a random solid-solution alloy,\nevery column contains a statistical mixture of Li, Mg, Al, Ag, and Cd, so the measured\nintensities form a continuous distribution rather than three discrete levels. Columns that\nare relatively rich in Ag/Cd appear as brighter maxima, whereas those enriched in\ntypically falling below both the detector’s noise floor and normal specimen-thickness\nvariations — and is therefore effectively invisible in HAADF images (Phys. Rev. Lett.\n2008, 100, 206101). Accordingly, the line profile in Figure 2c of the original manuscript\nshould not be interpreted as element-resolved “levels”; rather, it reflects a random\nintensity distribution that is fully consistent with random mixing and shows no evidence\nof heavy-element clustering.\nTo substantiate compositional uniformity of the Li-MSSA electrode at the atomic\nscale, we acquired atomically resolved STEM-EDX maps (Figure R3; see also Figure\n2a in the revised manuscript and Figure S3 in the revised SI). The Ag, Cd, Mg, and Al\nsignals are co-spatial within the spatial resolution with no detectable nanoscale\nsegregation. As expected, Li cannot be mapped reliably by EDX. We therefore\nperformed spectrum-imaging EELS at the Li-K edge to assess the Li distribution; after\nstandard background subtraction, the Li-K signal is uniform across the probed area with\nno evidence of Li-rich or Li-poor domains (Figure R4; see also Figure 2b in the revised\nmanuscript and Figure S5 in the revised SI). Taken together, these results support that\nthe Li-MSSA is a single-phase random solid solution, rather than a mixture of element-\nrich regions or intermetallic secondary phase.\nWe added the relevant discussions in the revised manuscript:\n“Atomically resolved high-angle annular dark-field scanning transmission electron\nmicroscopy (HAADF-STEM) with energy-dispersive X-ray spectroscopy (EDX)\nspectrum imaging (Fig. 2a and Supplementary Fig. 3) shows that Ag, Cd, Mg, and Al\nare co-spatial within our spatial-resolution and counting-statistics limits, with no\ndetectable nanoscale segregation; large-area scanning electron microscopy coupled\nwith (SEM-EDX) maps corroborate this (Supplementary Fig. 4). Because Li cannot be\nmapped reliably by EDX, we performed electron energy-loss spectroscopy (EELS)\nspectrum imaging at the Li K-edge; after standard background subtraction, the Li-K\nsignal is uniform across the probed area, with no evidence of Li-rich/Li-poor domains\ndown to sub-100-nanometre length scales (Fig. 2b and Supplementary Fig. 5). Taken\ntogether, these microscopy and spectroscopy results indicate that, within our\nnanometer-scale resolution, the delithiated Li-MSSA is compositionally uniform and\nbehaves as a single-phase random solid solution, rather than a mixture of element-rich\nregions or intermetallic secondary phases.” (Line 9, Page 8)\nEDX elemental maps for Ag, Mg, Cd, and Al, showing a homogeneous distribution\nwith no detectable nanoscale segregation. (b) Integrated EDX spectrum (0-8 keV) from\nthe region in (a). The Mo-L features in the spectrum arise from the molybdenum support\ngrid used to mount the TEM specimen (not from the alloy). The trace F signal may be\nattributed to sample preparation rather than to the alloy. Specifically, the bulk alloy was\nhot-rolled into thin foils under a fluorinated antistatic protective film, and the foils were\nthen dispersed/sonicated in DME for TEM mounting. We infer that tribochemical\ntransfer from the protective film during rolling, combined with subsequent solvent\nexposure, introduced a small amount of adventitious fluorine (Nat. Commun. 2023, 14,\n5678; Nano-Micro Lett. 2023, 15, 235).\nFigure R4. (a) HAADF-STEM image of the Li-MSSA in Li-rich state. (b) EELS Li-K\nelemental map from the same region acquired by spectrum imaging, showing a uniform\nLi distribution. (c) Background-subtracted EELS spectrum with the Li-K edge near 60\n7\neV, confirming the presence of Li in the matrix.\nComment 3: The schematic in Figure 2d shows Li and other atoms explicitly, but it is\nnot evident how their positions were determined from the experimental data. It is\nrecommended that complementary techniques such as EELS be employed to verify the\npresence and distribution of Li and thereby strengthen the interpretation.\nResponse: We thank the reviewer and agree that a schematic should not imply atom-\nby-atom identification. Figure 2d of the original manuscript is a conceptual illustration\nof a random solid solution (colors indicate species, not experimentally determined per-\natom positions). In multicomponent solid-solution alloy, the experimentally accessible\nquantity is statistical occupancy and chemical uniformity (and, if present, short-range\norder) rather than the exact coordinates of each atom. To avoid any potential\nmisinterpretation, we have removed Figure 2d in the revision.\nTo strengthen our interpretation and address the reviewer’s suggestion, we carried\nout complementary EELS at the Li-K edge. Atom-column-resolved Li mapping,\nhowever, requires ultrathin, exceptionally stable lamellae and long dwell times to\nachieve adequate counts. Lithium is highly reactive and beam-sensitive; even under\nlow-dose conditions, the Li-K signal rapidly degrades and evolves spectrally due to\nknock-on damage and radiolysis, as widely reported (Science 2017, 358, 506-510; Adv.\nMater. 2020, 32, 1908293). These constraints preclude acquiring reliable atomic-\nresolution Li-K spectrum images without introducing irradiation-induced artifacts. To\nmaximize spatial resolution while preserving sample integrity, we acquired Li-K-edge\nEELS spectrum-imaging at short dwell times and low dose. Across fields of view a few\nhundred nanometres wide, the Li-K maps are homogeneous; no composition contrast\nexceeding counting-statistics noise is detected (Figure R4; see also Figure 2b in the\nrevised manuscript and Figure S5 in the revised SI). These measurements directly verify\nthe presence of Li and its spatial uniformity on sub-100-nm length scales, thereby\nsupporting that Li-MSSA is a single-phase, randomly mixed solid solution, rather than\na mixture of element-rich regions or intermetallic secondary phases.\nThe relevant discussion has been added to the revised manuscript as shown in the\nresponse to Comment 2.\nComment 4: Figure 2h suggests that all four metal elements in MSSA are essential for\nperformance. Could the authors provide further insights into this result? Given that\n8\nthese metals are known to be lithium-philic and have the potential to individually shown\ngood cycling stability, additional discussion would help clarify the observed synergy.\nResponse: We appreciate the critical comments from Reviewer #1. In designing the\nalloy, our primary objective was to maximize the active Li mass fraction at the anode\nwhile preserving a single-phase solid solution throughout cycling, because lithium is\nthe capacity-bearing species. This necessitates constituents that exhibit wide solid-\nsolution ranges with Li (e.g., Mg, Ag, Cd) so that Li can be stored in the bulk rather\nthan plated at the interface. Practically, however, differences in melting points, mutual\nsolubilities, and reduction potentials of Mg, Ag, Cd elements make such\nmulticomponent alloys difficult to homogenize. Guided by the mixing free energy,\n∆𝐺 = ∆𝐻 −𝑇∆𝑆 , we therefore introduced Al as an enthalpic stabilizer\n𝑚𝑖𝑥 𝑚𝑖𝑥 𝑚𝑖𝑥\n(negative ∆𝐻 ; Nature 2024, 625, 697-702), thereby lowering ∆𝐺 and enabling\n𝑚𝑖𝑥 𝑚𝑖𝑥\na homogeneous multicomponent solid solution. This thermodynamic stabilization\nunderpins our choice of the quaternary composition and does not rely on the nominal\n“lithiophilicity” of any single metal.\nUnlike the widely used “lithiophilic site” strategy — where alloy additives mainly\nlower the Li nucleation overpotential to promote more uniform Li deposition (Adv.\nMater. 2024, 36, 2308257; Nat. Commun. 2024, 15, 1354; Adv. Mater. 2024, 36,\n2404689) — our alloy is designed to promote an inward-growth mode, in which Li\natoms migrate into the bulk rather than accumulate at the interface (J. Am. Chem. Soc.\n2023, 145, 24775-24784; J. Am. Chem. Soc. 2020, 142, 8818-8826). In the lithiophilic\napproach, once a thin Li layer forms, subsequent cycling still proceeds by Li+ reduction\non Li metal at the electrolyte interface; the locus of deposition does not fundamentally\nchange. By contrast, our multicomponent solid solution operates as a single phase with\nfinite Li solubility, so the Li chemical potential varies continuously with composition.\nUnder plating, this establishes a chemical-potential gradient that drives freshly formed\nLi atoms at the surface to diffuse inward rather than accumulate at the interface —\nconsistent with solid-solution-assisted inward growth (J. Am. Chem. Soc. 2023, 145,\n24775-24784; J. Am. Chem. Soc. 2020, 142, 8818-8826). Kinetically, multicomponent\nsolid-solution alloys (≥5 elements) possess high configurational entropy that induces\nsignificant local chemical and strain disorder, thereby broadening the distribution of Li\nsite energies. This, in turn, generates multiple, percolating diffusion pathways and\nenhances Li mobility — our rationale for employing a multicomponent alloy (J. Am.\nChem. Soc. 2024, 146, 21320-21334; Adv. Mater. 2024, 36, 2314054; Adv. Mater. 2024,\n9\n36, 2409278). Consequently, our alloy sustains percolating, quasi-isotropic bulk\npathways that provide sufficient subsurface transport to accommodate the interfacial Li\nflux, thereby preventing interfacial supersaturation. Together with the thermodynamic\ndrive from continuous chemical potential ( ), this bulk transport relocates Li uptake\nLi\ninto the foil interior and suppresses surface deposition. Consistent with this mechanism,\nour in-situ XRD, TOF-SIMS, and STEM-EDX/EELS indicate a single-phase solid\nsolution across the cycling window.\nTo further substantiate this point, we have added the corresponding discussion to\nthe main text:\n“Unlike the widely used “lithiophilic site” strategy — where alloy additives mainly\nlower the Li nucleation overpotential to promote more uniform Li deposition45-47 — our\nalloy is designed to promote an inward-growth mode, in which Li atoms migrate into\nthe bulk rather than accumulate at the interface10,11. In the lithiophilic approach, once\na thin Li layer forms, subsequent cycling still proceeds by Li+ reduction on Li metal at\nthe electrolyte interface; the locus of deposition does not fundamentally change. By\ncontrast, our multicomponent solid solution operates as a single phase with finite Li\nsolubility, so the Li chemical potential varies continuously with composition. Under\nplating, this establishes a chemical-potential gradient that drives freshly formed Li\natoms at the surface to diffuse inward rather than accumulate at the interface —\nconsistent with solid-solution-assisted inward growth.10,11 Kinetically, multicomponent\nsolid-solution alloys (≥5 elements) possess high configurational entropy that induces\nsignificant local chemical and strain disorder, thereby broadening the distribution of Li\nsite energies. This, in turn, generates multiple, percolating diffusion pathways and\nenhances Li mobility — our rationale for employing a multicomponent alloy.42,48,49\nConsequently, our alloy sustains percolating, quasi-isotropic bulk pathways that\nprovide sufficient subsurface transport to accommodate the interfacial Li flux, thereby\npreventing interfacial supersaturation. Together with the thermodynamic drive from\ncontinuous chemical potential ( ), this bulk transport relocates Li uptake into the foil\nLi\ninterior and suppresses surface deposition. Consistent with this mechanism, our in-situ\nXRD, TOF-SIMS, and STEM-EDX/EELS indicate a single-phase solid solution across\nthe cycling window.” (Line 7, Page 16)\nhigh-temperature molten conditions. Could the authors clarify how such a property is\nretained or manifested under room-temperature electrochemical cycling? While MSSA\nmay exhibit strong lithium affinity in its molten state, it would be helpful to understand\nwhether this property remains relevant during practical battery operation.\nResponse: We appreciate the reviewer’s insightful question. In molten Li,\n“lithiophilicity” is typically quantified by the static contact angle 𝜃 of liquid Li on a\nsubstrate, which — via the Young-Dupré relation — reflects the balance of the solid-\nliquid interfacial free energy 𝛾 , the solid-vapor 𝛾 , and the liquid-vapor (surface\n𝑆𝐿 𝑆𝑉\ntension) 𝛾 . Lowering 𝛾 and/or 𝛾 increases cos𝜃, transforming Li from a bead\n𝐿𝑉 𝑆𝐿 𝐿𝑉\n(non-wetting) to a thin film (wetting). In practice, enhanced wetting is achieved by\ndepositing ultrathin lithiophilic coatings — e.g., Al O , Si, Mg, Ge, ZnO, or surface-\n2 3\ntethered organic functional groups — that form Li-O, Li-Si, or Li-M interfacial bonds;\nthe improved wetting arises from the thermodynamic favorability of interfacial bond\nformation, with a negative Gibbs free-energy change (∆𝐺 < 0) as the quantitative\ncriterion (Nat. Nanotech. 2016, 11, 626-632; Proc. Natl. Acad. Sci. U.S.A. 2016, 113,\n2862-2867; Proc. Natl. Acad. Sci. U.S.A. 2017, 114, 3584-3589; Adv. Energy Mater.\n2018, 8, 1701963; Energy Storage Mater. 2018, 14, 345-350; Nat. Commun. 2019, 10,\n4930).\nUnder room-temperature operation, lithiophilicity is defined as an electrode’s\nability to achieve a low nucleation overpotential, a high and spatially uniform nuclei\ndensity, and a smooth deposit morphology, thereby guiding Li to deposit in an ordered,\nuniform manner. Most hosts with weak Li-substrate interactions produce disordered\nnucleation/growth; a genuinely lithiophilic interface promotes order via three coupled\neffects: (i) favorable interfacial thermodynamics (interfacial reactions/alloying that\nreduce the solid-liquid interfacial free energy and the nucleation free-energy barrier),\n(ii) sufficient adatom binding to stabilize nascent Li without pinning, and (iii)\nappropriately low surface-diffusion barriers so that adatoms can redistribute into low-\nenergy configurations at the applied current. Together these factors convert interfacial\n“chaos” into more ordered deposition.\nOur system does not rely on conventional “lithiophilic-site” interfacial deposition.\nInstead, our alloy is designed to promote an inward-growth mode, in which Li atoms\nmigrate into the bulk rather than accumulate at the interface. Because the Li-based\n11\nsolid-solution alloy has a finite yet appreciable Li solubility, the Li chemical potential\nvaries smoothly with composition inside the solid, providing a persistent\nthermodynamic drive for freshly formed Li at the surface to diffuse inward.\nConsequently, the dominant Li flux is redirected from the electrolyte/electeode\ninterface into the alloy: Li penetrates tens of micrometres beneath the surface via solid-\nsolution insertion within the matrix, rather than accumulating as a metallic overlayer.\nThis endogenous inward-growth pathway effectively suppresses surface dendrites.\nIn this paper, the authors present a multicomponent lithium metal anode with 90%\nlithium and four other metallic components. The authors show good cycling behavior,\nwithout significant concentration gradients that form. The concept is interesting and the\ncharacterization of the material is good. The cycling behavior is very good. Despite the\npromise, however, there are a number of issues that should be addressed before\npublication, including seemingly misleading statements regarding specific energy and\ninsufficient comparison to control experiments. Please see notes below.\nResponse: We sincerely thank Reviewer #2 for the careful evaluation of our manuscript\nand the insightful comments provided. These suggestions have been instrumental in\nimproving the quality of our work. We have carefully revised the manuscript in\nresponse to the reviewer’s feedback and greatly appreciate the time and expertise\ndedicated to this review.\nComment 1: What is the rate behavior of the multicomponent material compared to\npure Li in full cells? Rate behavior and rate effects on structural evolution are not\ndiscussed in the manuscript.\nResponse: We thank the reviewer for raising this important point. To evaluate the rate\ncapability of Li-MSSA, we assembled Li-MSSA||NCM811 and Li/Cu||NCM811 cells\nunder identical cathode loadings (10.8 mg cm-2). C-rates were defined with respect to\nthe cathode (1 C = 209 mA g-1, based on a practical capacity of 209 mAh g-1). The\nLi-MSSA||NCM811 cell delivered 208 (0.1 C), 201 (0.2 C), 190 (0.5 C), 175 (1 C), 144 (2\nC), 105 (4 C), and 68 mAh g-1 (6 C). By contrast, Li/Cu||NCM811 showed markedly\npoorer performance, reaching only 25 mAh g-1 at 6 C, which underscores the Li-MSSA\nadvantage at high rates (Figure R5; also Figure S21 in the revised SI).\nTo assess rate-dependent structural evolution of the Li-MSSA anode, we performed\nex-situ STEM-HAADF imaging with high-resolution STEM-EDX mapping on Li-\nMSSA after rate cycling up to 6 C (Figure R6; also Figure S22 in the revised SI).\nElemental maps for Ag, Cd, Mg, and Al are co-localized within our spatial resolution,\nwith no detectable nanoscale segregation or heavy-element clustering. Collectively, the\ndata indicate that (i) under matched full-cell conditions, Li-MSSA delivers a higher\nusable capacity than pure Li at elevated rates, and (ii) after high-rate cycling (up to 6 C)\nthe anode shows no evidence of rate-induced phase separation or nanoscale\ncompositional partitioning.\naspect in the main text:\n“To evaluate the rate capability of Li-MSSA, we assembled Li-\nMSSA||LiNi Co Mn O and Li/Cu||LiNi Co Mn O full-cells under identical\n0.8 0.1 0.1 2 0.8 0.1 0.1 2\ncathode loadings (10.8 mg cm-2, Supplementary Fig. 21). C-rates were defined with\nrespect to the cathode (1 C = 209 mA g-1, based on a practical capacity of 209 mAh g-\n1). The Li-MSSA||LiNi Co Mn O cell delivered 208 (0.1 C), 201 (0.2 C), 190 (0.5\n0.8 0.1 0.1 2\nC), 175 (1 C), 144 (2 C), 105 (4 C), and 68 mAh g-1 (6 C). By contrast,\nLi/Cu||LiNi Co Mn O showed markedly poorer performance, reaching only 25\n0.8 0.1 0.1 2\nmAh g-1 at 6 C, which underscores the Li-MSSA advantage at high rates\n(Supplementary Fig. 21). To assess rate-dependent structural evolution of the Li-MSSA\nanode, we performed ex-situ STEM-HAADF imaging with high-resolution STEM-EDX\nmapping on Li-MSSA after rate cycling up to 6 C (Supplementary Fig. 22). Elemental\nmaps for Ag, Cd, Mg, and Al are co-localized within our spatial resolution, with no\ndetectable nanoscale segregation or heavy-element clustering. Collectively, the data\nindicate that (i) under matched full-cell conditions, Li-MSSA delivers a higher usable\ncapacity than pure Li at elevated rates, and (ii) after high-rate cycling (up to 6 C) the\nanode shows no evidence of rate-induced phase separation or nanoscale compositional\npartitioning.” (Line 21, Page 18)\nrates. Both cells use the same NCM811 cathode loading (10.8 mg cm-2); C-rates are\ndefined with respect to the cathode (1 C = 209 mA g-1). (b) Voltage profiles of Li-\nMSSA||NCM811 (b) and Li/Cu||NCM811 (c) at selected C-rates, plotted versus\ngravimetric capacity (mAh g-1).\nFigure R6. (a) HAADF-STEM image of Li-MSSA after rate cycling up to 6 C, with\ncorresponding EDX elemental maps for Ag, Mg, Cd, and Al, showing a homogeneous\ndistribution with no detectable nanoscale segregation. (b) Integrated EDX spectrum (0-\n8 keV) from the region in (a). The Mo-L features in the spectrum arise from the\n15\nmolybdenum support grid used to mount the TEM specimen (not from the alloy). The\nF emission from fluorinated SEI species formed by electrolyte decomposition.\nComment 2: All details of cell cycling should be given in the captions for each figure.\nFor instance, what is specific electrolyte used for pouch cells? Two are mentioned in\nmethods. Authors need to be much more careful about providing all details for\nindividual experiments - right now it is muddled and difficult to determine the details.\nResponse: We appreciate the reviewer’s careful reading and fully agree that each figure\ncaption must contain all cycling details needed for reproducibility.\nFor Figure 2f and 2h of the original manuscript (half-cell capacity and CE),\nCR2032 coin cells were assembled using Li-MSSA or bare-Cu working electrodes (10\nmm diameter), a Li-metal counter/reference electrode (10 mm diameter, 300 µm thick),\na Celgard 2400 separator, and 60 µL per cell of 1.0 M LiTFSI in DME:DOL (1:1, v/v)\nwith 2 wt.% LiNO . For Figure 3a of the original manuscript (TOF-SIMS), CR2032\n3\ncoin cells were assembled with Li-MSSA as working electrode (10 mm diameter), a Li-\nmetal counter/reference (10 mm diameter, 300 µm thick), a Celgard 2400 separator, and\n60 µL per cell of 1.0 M LiTFSI in DME:DOL (1:1, v/v) with 2 wt.% LiNO . The Li-\nand 6 mAh cm-2 plated per half-cycle, for a total of 10 full cycles. After cycling, the\ncells were opened in the lithiated state and subjected to ex-situ TOF-SIMS analysis. For\nFigure 3d and 3e of the original manuscript, in-situ electrochemical optical microscopy\nwas conducted in a commercial transparent cell (Beijing Scistar) with a glass window.\nEach cell contained 1.0 mL of electrolyte (1.0 M LiTFSI in DME:DOL, 1:1 v/v, with 2\nwt.% LiNO ). Two configurations were examined: Li‖Li and Li‖Li-MSSA. In the Li‖Li-\nserved as the counter/reference; in the Li‖Li controls, a Li-metal foil was used as the\nworking electrode. During the optical tests, an areal capacity of 5 mAh cm-2 was\ndeposited on the working electrodes. For Figure 4a-d of the original manuscript (in-situ\nXRD), measurements used a commercial LIB-XRD cell (Beijing Scistar) with a Be\nwindow; the working electrode was an alloy film (16 mm diameter) cast on the window,\na Li-metal foil (16 mm diameter, 300 µm thick) served as counter/reference, the cell\ncontained 100 µL of 1.0 M LiTFSI in DOL:DME (1:1, v/v) with 2 wt.% LiNO . This\n3\nin-situ cell was cycled at a current density of 1 mA cm-2. For Figure 5a of the original\nmanuscript (pouch-cell testing), pouch cells were filled with 1.5 M LiFSI in DME:TTE\n16\n(1.2:3, v/v) at E/C = 2 g Ah-1 and cycled at 25 °C with 0.1 C charge/1 C discharge (C-\nrate referenced to the NCM811 cathode). For Figure 5d of the original manuscript (Li-\nS CR2032 cells), we employed a Celgard 2400 separator and 60 μL of 2.0 M LiFSI in\nn-butyl methyl ether as the electrolyte. The cells were cycled on a LAND tester at 25 °C,\nwith a rate of 0.06 C for high-loading cathodes (7.4 mAh cm-2) and 0.2 C for low-\nloading cathodes (3.5 mAh cm-2).\nFor clarity, all electrochemical testing parameters are now stated in the figure\ncaptions.\nComment 3: pouch cell fab and testing info is not provided in the methods.\nResponse: We appreciate the reviewer’s attention to detail and have expanded the\nMethods to comprehensively describe pouch-cell fabrication and evaluation.\n“In the 1 Ah pouch cells we use a free-standing Li-MSSA foil as both active layer\nand current collector; no additional Cu substrate is employed. The foil is prepared by\nhot-rolling a cast Li-MSSA ingot to a thickness of 40 µm, then slit, calendared, and\nspot-welded to a Ni tab before stacking. The pouch cells were fabricated on a semi-\nautomated pilot line in a dry room (dew-point -55 °C, RH < 0.1 %). The NCM811\ncathode slurry was double-side coated onto 10 µm Al foil, delivering 3.5 mAh cm-2 per\nside (7.0 mAh cm-2 total). We use a free-standing Li-MSSA foil that serves as both the\nactive layer and the current collector; no Cu substrate is employed. The foil is produced\nby hot-rolling a cast Li-MSSA ingot to 40 µm, then slit, calendared, and spot-welded to\na Ni tab before stacking. The Li-MSSA anode provides an areal capacity of 7.8 mAh\ncm-2. Cathode and anode sheets were Z-stacked with 12 µm polypropylene separators,\ntab-welded (Al for the cathode, Ni for the anode), packaged in 85 mm × 65 mm\nlaminated pouches, vacuum-dried (60 °C, 12 h), and vacuum-injected with 1.5 mol L-1\nLiFSI in 1,2-dimethoxyethane/1,1,2,2-tetrafluoroethyl-2,2,3,3-tetrafluoropropyl ether\n(1.2 : 3, v/v) at an electrolyte-to-capacity ratio of 2 g Ah-1. After two C/20 formation\ncycles, the cells were charged at 0.1 C and discharged at 1 C between 2.6 V and 4.3 V\non a LAND tester at 25 °C.” (Line 23, Page 23)\nComment 4: -important: in the abstract and conclusion, it is stated that the paper\ndemonstrats 470 Wh/kg. This is not true, as it is only extrapolated from a 1-Ah cell to\na 10 Ah-cell. Please remove this claim and please claim only accurate demonstrations.\nResponse: We thank the reviewer for pointing out this. We have removed this number\n17\nfrom both the Abstract and the Conclusion.\nComment 5: -the superlatives (“exceptional” “unprecedented”) sprinkled throughout\nthe paper are unnecessary.\nResponse: We have removed all superlative terms throughout the manuscript.\nComment 6: -on page 7 it is stated that the capacities of the alloy surpasses that of\n“state of the art Li metal electrodes.” Why would this be the case? Li metal has higher\ntheoretical capacity and can be fully stripped. This is a confusing statement.\nResponse: We apologize for the ambiguous wording. Our intention was not to imply\nthat the alloy exceeds the theoretical capacity of bulk Li metal (3860 mAh g-1). Rather,\nour point is that it delivers a higher practical, reversible capacity than most state-of-the-\nart modified Li-metal anodes once the penalties of lithium mass fraction (M) and\nutilization (U) are accounted for. For instance, many recent studies focus on electrolyte\nengineering to stabilize metallic Li. In such designs, the lithium mass fraction is\nnominally 100 wt.%, yet only a fraction of the Li inventory is cycled, leading to\ngravimetric capacities well below the theoretical limit. As one example, Jie et al. (Nat.\nEnergy 2024, 9, 987-998) reported a compact ion-pair aggregate electrolyte enabling\nstable pouch cells. In a representative configuration pairing a double-sided 8.2 mAh\ncm-2 cathode with a 60 µm Li foil (12.3 mAh cm-2 Li), the Li utilization is ~67%,\ncorresponding to a practical anode-specific capacity of ~2590 mAh g-1. Alternatively,\nstructural strategies such as 3D current collectors can improve Li stability but often\nreduce the Li mass fraction significantly. For example, Qingshuai Xu et al. designed a\nLi-carbon composite anode via cold-rolling (Nature 2025, 637, 339-346). When the\nmass of the carbon scaffold is included, the overall Li mass fraction drops to ~43 wt.%,\nagain yielding a capacity of ~1660 mAh g-1.\nIn short, most advanced Li-metal architectures sacrifice either M (< 60 wt.%) or U\n(< 80%) to mitigate dendrite growth and improve cycling, resulting in practical\ncapacities typically below 2000 mAh g-1. By contrast, our multicomponent solid-\nsolution alloy retains both a high Li mass fraction (90 wt.%) and a high utilization\n(89%), delivering a reversible capacity of 3100 mAh g-1. This value approaches the\ntheoretical limit of pure lithium and, when both lithium mass fraction and utilization\nare considered, exceeds the practical values reported for current lithium-metal\nelectrodes (see Table S1 in the original SI).\nrevised the relevant description:\n“This value approaches the theoretical limit of pure lithium and, when both lithium\nmass fraction and utilization are considered, exceeds the practical values reported for\ncurrent lithium-metal electrodes (Fig. 1a and Supplementary Table 1).” (Line 13, Page\n9)\nComment 7: -on page 8 there is lots of discussion about stripping and cycling of the\nalloy electrodes, but they are never compared to pure Li electrodes. Direct comparison\nto pure Li tested under the same conditions would be very helpful for the reader to\nunderstand the true benefits of the alloy.\nResponse: We thank the reviewer for this valuable suggestion. We have added CE data\nfor a pure Li/Cu half-cell tested under identical conditions to those used for the alloy\nelectrodes. As now shown in Figure R7 (also shown as Figure 2f in the revised\nmanuscript), the Li/Cu cell delivers an average CE of 98.3 ± 0.5% over the first 200\ncycles at 1 mA cm-2 with an areal capacity of 1 mAh cm-2; thereafter, the CE drops\nsharply, consistent with dendritic shorting. In contrast, the Li-MSSA alloy electrode\nmaintains a higher and more stable CE over prolonged cycling, highlighting the\npractical advantages of the alloy design compared to pure Li.\nWe have updated the relevant data and revised the main-text discussion\naccordingly:\n“Under identical conditions, a Li/Cu half-cell (prepared by laminating a Li foil\nonto Cu) delivers an average CE = 98.3 ± 0.5% over the first 200 cycles; thereafter the\nCE drops sharply, consistent with dendritic soft-shorting, highlighting the practical\nadvantage of the Li-MSSA design over pure Li.” (Line 12, Page 10)\nFigure R7. Coulombic-efficiency comparison for Li-MSSA, representative Li-based\nquaternary alloy anodes, and Li metal, cycled at 1 mA cm-2 with an areal capacity of 1\nmAh cm-2.\nbe a constant plateau as in Fig. 2?\nResponse: We thank the reviewer for this insightful comment. In our GITT protocol,\n0.01 C/120 s pulses were applied with 900 s rests between pulses. During each current-\non period, the measured cell voltage equals the equilibrium potential plus ohmic and\nkinetic/concentration polarizations. The pulse induces an initial iR drop and a time-\ndependent diffusion overpotential, producing a slight monotonic drift before the current\nis switched off; the voltage then relaxes back toward the quasi-OCV during the rest\nperiod. Meanwhile, in our prior work, the GITT experiments primarily sampled the\nearly stage of lithiation — i.e., a magnified segment of the discharge curve in Figure 2f\nof the original manuscript — and only briefly entered the subsequent voltage plateau.\nAs a result, the intra-pulse trace in Figure 3b of the original manuscript appears as a\ngentle slope rather than a flat plateau. By contrast, Figure 2f of the original manuscript\nshows conventional galvanostatic cycling (no pulses): under constant current the\noverpotential is quasi-steady, and the profile is plotted on a coarser time/voltage scale,\nso the same composition window appears as a nearly flat plateau.\nTo address the reviewer’s concern, we performed extended-window GITT over a\nbroader lithiation range, applying 0.01 C pulses for 30 s with 1800 s rests between\npulses. The resulting profile (Figure R8; also Figure S15 in the revised SI) exhibits a\nmore pronounced plateau region while preserving the expected intra-pulse drift and\npost-pulse relaxation.\nFigure R8. GITT curve for calculating the Li atom diffusion coefficient. The inset\nenlarges the initial 0 – 15 h segment.\nComment 9: -Are all the assumptions for GITT to be used to calculate diffusion\ncoefficients satisfied? This is generally not the case for most works, and extracted\ndiffusion coefficients are wrong. See https://pubs.acs.org/doi/10.1021/acsaem.1c02218.\nResponse: We thank the reviewer for raising this important point. We fully agree with\n20\nthe reviewer that diffusivities extracted using the conventional GITT treatment are\nreliable only when the short-time, semi-infinite-diffusion assumptions are satisfied. In\nour case, taking the Li-MSSA foil thickness as the characteristic diffusion length (L =\n30 µm) and our initial estimate 𝐷≈10-7 cm2 s-1, the diffusion penetration depth for a\nτ=120 s pulse is 𝛿 = √𝜋𝐷𝜏 =61 μm, giving 𝛿/𝐿 =2≫0.2. Thus the short-time\ncriterion is violated, the classical √𝜏 relation does not apply under our conditions, and\nthe apparent 𝐷 from that treatment would be biased (overestimated).\nIn response, we remeasured GITT and reanalyzed the data using the finite-length\nspherical diffusion model recommended by Horner et al. (ACS Appl. Energy Mater.\n2021, 4, 11460-11469), i.e., direct pulse fitting of the entire voltage transient rather than\nrelying on an early √𝜏 window. Because the solid-solution alloy has a finite Li\nsolubility, the bulk Li chemical potential varies continuously with composition and\ntransport is therefore chemical-potential-driven (J. Am. Chem. Soc. 2023, 145, 24775-\n24784). We therefore employed a non-ideal, chemical-potential-based diffusion\nformulation to fit the full voltage transients of the new GITT dataset, using 0.01 C\npulses of 30 s with 1800 s rests between pulses (Figure R8; also Figure S15 in the\nrevised SI). Following the reviewer-recommended analysis framework, which was\noriginally implemented for a one-dimensional single-particle model, we adapted the\ngeometry to our one-dimensional slab electrode: the radial coordinate 𝑟 is replaced\nby the thickness coordinate 𝑥, the particle radius 𝑅 by the foil thickness 𝐿, and\n𝑃\ndiffusion is assumed to occur only along the thickness direction. The governing\nequation is thus the 1D slab analogue of Fick’s law with a thermodynamic factor (non-\nideal transport), solved with finite-length boundary conditions and fitted to the entire\ncurrent-voltage transient rather than an early √𝜏 segment. From fits to the new dataset,\nthe extracted Li chemical diffusivity remains 6  10-9 cm2 s-1 (Figure R9; also Figure\nS16 in the revised SI). As expected, applying other analysis routes to the same dataset\n— such as the traditional GITT early-time √𝜏 approximation, or EIS-based diffusivity\nextraction — yields values up to one or two orders of magnitude higher relative to our\nnon-ideal, chemical-potential-driven fit (ACS Appl. Energy Mater. 2021, 4, 11460-11469).\nNotably, these values are higher than the reported room-temperature self-diffusion of\nLi in metallic Li (typically 10-11 cm2 s-1, Glass Phys. Chem. 2010, 36, 570-574). Taken\ntogether, the fitted 𝐷 is sufficiently high to accommodate the per-pulse Li flux by bulk\ninsertion into Li-MSSA.\n“Because the solid-solution alloy has a finite Li solubility, the bulk Li chemical\npotential varies continuously with composition and transport is therefore chemical-\npotential-driven.10 We therefore employed a non-ideal, chemical-potential-based\ndiffusion formulation to fit the full voltage transients of the GITT dataset, using 0.01 C\npulses of 30 s with 1800 s rests between pulses (Supplementary Fig. 15).40 Under this\nprotocol, the fitted Li chemical diffusivities are around 6  10-9 cm2 s-1 (Supplementary\nFig. 16). As expected, applying other analysis routes to the same dataset — such as the\ntraditional GITT early-time √𝜏 approximation, or EIS-based diffusivity extraction —\nyields values up to one or two orders of magnitude higher relative to our non-ideal,\nchemical-potential-driven fit.40 Notably, these values are higher than the reported\nroom-temperature self-diffusion of Li in metallic Li (typically 10-11 cm2 s-1).41 Taken\ntogether, the fitted 𝐷 is sufficiently high to accommodate the per-pulse Li flux by bulk\ninsertion into Li-MSSA.” (Line 3, Page 13)\nFigure R9. Comparison of diffusion coefficients extracted from galvanostatic\nintermittent titration transients using two models: Ideal model and Non-ideal model.\nComment 10: -what is the capacity stripped/plated in Fig 3a?\nResponse: We appreciate the reviewer’s request for clarification. In Figure 3a of the\noriginal manuscript the Li-MSSA electrode was galvanostatically cycled at 2 mA cm-2,\nstripping and plating 6 mAh cm-2 per cycle (i.e., 6 mAh cm-2 Li removed, then 6 mAh\ncm-2 redeposited) for a total of 10 full cycles.\nComment 11: -does the MSSA material remain a solid solution even as lithium is fully\nremoved, or do compounds form? Are there other peaks that show up in XRD that are\nnot shown in the limited 2theta range of Fig 4?\nResponse: We are grateful to the reviewer for this insightful question. Within our nm-\n22\nscale spatial and ns-scale temporal windows, the delithiated Li-MSSA behaves as a\nsingle-phase solid solution. We do note intermittent XRD features in some scans during\ncycling; however, owing to the alloy’s multicomponent complexity, we do not assign\nthese features at this stage.\nTo directly determine whether Li-MSSA remains a solid solution after full\ndelithiation, we acquired atomically resolved STEM-EDX spectrum-imaging on\ndelithiated electrodes (Figure R10; also Figure S11 in the revised SI). The Ag, Cd, Mg,\nand Al maps are co-spatial across the probed regions, with no nanoscale segregation or\nintermetallic domains observed. To assess Li specifically, we performed spectrum-\nimaging EELS at the Li-K edge; the residual Li-K signal after delithiation is spatially\nuniform across the mapped area, with no evidence of Li-rich or Li-poor domains at\nhundreds-of-nanometers length scale (Figure R11; also Figure S12 in the revised SI).\nIn addition, we further employed hybrid Monte Carlo and molecular dynamics (MCMD)\nsimulations to evaluate phase stability and elemental mixing during\nlithiation/delithiation across representative Li contents (Figures R12 and R13; see also\nFigure 3b in the revised manuscript and Figures S13 and S14 in the revised SI). Across\nmultiple randomized initial configurations, 300 K trajectories (0.5 ns) show no\ntendency toward elemental clustering either after lithiation or after delithiation; the\nalloy remains compositionally uniform within the simulation cell. The combined\nmicroscopy/spectroscopy and simulation evidence indicates that, within our nanometer\nspatial and nanosecond temporal resolution, the delithiated Li-MSSA remains\ncompositionally uniform.\nWe reprocessed and extended the in-situ XRD scans of the delithiated electrode for\ncycles 1 and 4 (Figure R14; also Figure S16 in the revised SI). Beyond the principal\nbcc-type reflections emphasized in the main text (36.2°, 52.0°, 65.0°), some scans (e.g.,\ncycle 4) exhibit weak, broad, and intermittent features near 34.9°, 43.1°, and 57.3°.\nGiven the multicomponent nature of the alloy, these features cannot be assigned\nunambiguously to specific intermetallic phases and are left unassigned in the present\nstudy. Looking forward, the combinatorial phase space of multicomponent alloys\nsuggests that machine-learning-assisted phase fingerprinting (e.g., simulated pattern\nlibraries combined with Bayesian/NN classifiers), together with CALPHAD-guided\nsearches and pair-distribution-function analyses, could enable more robust\nidentification. Our focus here is the design concept of a multicomponent solid-solution\nalloy that leverages a continuous chemical-potential gradient in the bulk and high\n23\nconfigurational entropy to provide multiple, percolating diffusion pathways, thereby\nenhancing Li mobility and enabling stable cycling at high Li mass fraction and high Li\nutilization. Determining whether the features correspond to bona fide secondary phases,\nand whether they influence long-term cycling, is an important direction for future work.\nTaken together, the STEM-EDX/EELS, MCMD, and XRD results indicate that,\nwithin our nanometer spatial and nanosecond temporal resolution, the delithiated\nelectrode behaves as a single-phase solid solution, even though occasional secondary\nfeatures can appear at the macroscopic diffraction level. These clarifications are now\nmade explicit in the main text, and all datasets are provided in the SI:\n“To test whether Li-MSSA remains a solid solution after full delithiation, we also\nacquired atomically resolved STEM-EDX and Li-K-edge EELS on delithiated\nelectrodes (Supplementary Fig. 11). The Ag, Cd, Mg, and Al maps remain co-spatial\nwith no nanoscale segregation or intermetallic domains observed, and the residual Li-\nK signal after delithiation is likewise spatially uniform across the mapped area\n(Supplementary Fig. 12). In addition, we further employed hybrid Monte Carlo and\nmolecular dynamics (MCMD) simulations to evaluate phase stability and elemental\nmixing during lithiation/delithiation across representative Li contents (Supplementary\nFigs. 13 and 14). Across multiple randomized initial configurations, 300 K trajectories\n(0.5 ns) show no tendency toward elemental clustering either after lithiation or after\ndelithiation; the alloy remains compositionally uniform within the simulation cell.”\n(Line 9, Page 12)\n“We re-extracted and compared the end-of-lithiation and end-of-delithiation\npatterns for cycles 1 and 4. Beyond the principal bcc-type reflections (36.2°, 52.0°,\n65.0°), we observe in some scans (e.g., cycle 4), broad features at 34.9°, 43.1°, and\n57.3°. Owing to the multicomponent nature of the alloy, these features cannot be\nunambiguously assigned to specific intermetallic phases at present. Given the\ncombinatorial phase space of multicomponent alloys, future efforts could employ\nmachine-learning-assisted phase fingerprinting (e.g., simulated pattern libraries\ncombined with Bayesian or neural-network classifiers), together with CALPHAD-\nguided searches and pair-distribution-function analyses, to enable more robust phase\nidentification. Determining if the minor features are bona fide secondary phases and,\nif so, whether they influence long-term cycling remains an important direction for future\nwork.” (Page 19 in the revised SI)\ncorresponding EDX elemental maps for Ag, Mg, Cd, and Al, showing a homogeneous\ndistribution with no detectable nanoscale segregation. (b) Integrated EDX spectrum (0-\n8 keV) from the region in (a). The Mo-L features in the spectrum arise from the\nmolybdenum support grid used to mount the TEM specimen (not from the alloy). The\nF emission from fluorinated SEI species formed by electrolyte decomposition.\nFigure R11. (a) HAADF-STEM image of the Li-MSSA in the delithiated state. (b)\nEELS Li-K elemental map from the same region acquired by spectrum imaging,\nshowing a uniform Li distribution. (c) Background-subtracted EELS spectrum with the\nLi-K edge near 60 eV, confirming the presence of Li in the matrix.\nMCMD simulation of a 9423-atom Li AgMgCdAl (Li M) cell at 300 K; data are\n352 88\nsampled every 0.5 ps. (b) Equilibrated snapshot of the supercell (Li, green; Ag, grey;\nAl, blue; Cd, orange; Mg, magenta; periodic boundary conditions). (c) Element-\nresolved projections highlighting the spatial distributions of Ag, Al, Cd, and Mg,\nrespectively.\nFigure R13. (a) Time evolution of the potential and kinetic energies during a hybrid\nMCMD simulation of a 9423-atom Li AgMgCdAl (Li M) cell at 300 K; data are\n39.6 9.9\nsampled every 0.5 ps. (b) Equilibrated snapshot of the supercell (Li, green; Ag, grey;\nAl, blue; Cd, orange; Mg, magenta; periodic boundary conditions). (c-f) Element-\nresolved projections highlighting the spatial distributions of Ag, Al, Cd, and Mg,\nrespectively, showing homogeneous dispersion in the Li matrix with no obvious\nclustering or nanoscale segregation at the simulation length scale.\n— initial, delithiated, and lithiated. Black arrowheads mark reflections from the Be\nwindow; blue arrowheads indicate the principal Li-like bcc reflections. Green\narrowheads denote intermittent features (unassigned) observed in some scans (e.g., near\n34.9°, 43.1°, and 57.3°).\nComment 12: -In the XRD data in fig 4, why does the MSSA material have a higher\nextent of 200 peak displayed compared to pure Li before stripping? Is this because of\ndifferent heat treatments of the different materials?\nResponse: We thank the reviewer for the insightful question. We attribute the higher\nextent of the (200) peak in Li-MSSA to alloying-induced changes in crystallographic\ntexture within a Li-like bcc matrix. This is consistent with our EBSD pole figures\n(Figure 4e of the original manuscript), which reveal a {100} fiber texture (axes near\n<100>) in Li-MSSA; such a texture naturally amplifies (200) and redistributes intensity\naway from (110)/(211), matching the 110/200/211 pattern in XRD. Previous literatures\nshow that alloying elements robustly alter deformation/recrystallization pathways —\nvia solute drag, precipitation/particle-stimulated nucleation (PSN), and stored-energy\ngradients — and thereby reshape texture and the resulting XRD peak intensities across\ncubic and hcp alloys; for example, Al-Mg alloys containing Sc/Zr form dispersoid-\ndriven PSN that changes recrystallization textures, and Mg alloys with rare-earth (or\nCa/Y) additions systematically weaken or re-orient basal textures while promoting non-\nbasal components (Mater. Sci. Technol. 2016, 32, 1303-1315; Materials 2021, 14, 5516;\nMater. Sci. Eng. A 2000, 277, 102-113; Scr. Mater. 2010, 63, 725-730; Metall. Mater.\nTrans. A, 2013, 45A, 3205-3212). Therefore, we conclude that the enhanced (200)\nintensity in Li-MSSA is best attributed to alloying-driven texture evolution within a Li-\nlike bcc solid solution.\nWe have added the discussion accordingly in the revised SI:\n27\n“XRD of the as-prepared Li-MSSA shows (200) as the most intense reflection,\nwhereas (110) is the most intense reflection for Li metal. The enhancement of (200)\nreflection in Li-MSSA is explained by a {100} fiber texture within a Li-like bcc MSSA,\nconsistent with our EBSD pole figures (Figure 4e in the main text) that show {100} fiber\ntexture (axes near <100>) and thus re-weight the (110)/(200)/(211) reflection\nintensities accordingly. Non-bcc alternatives do not simultaneously reproduce the\nobserved 110/200/211 positions and relative intensities under reasonable texture\nmodels. For a representative delithiation composition Li AgMgCdAl (Li M), the\n39.6 9.9\naverage VEC (1.55) lies far below the ~6.87 threshold that empirically favors bcc\nmatrices in multicomponent alloys,5 further supporting a Li-like bcc matrix. We note\nthat alloying may slightly shift the bcc lattice parameter relative to Li metal; this does\nnot affect the texture-based explanation for the (200) enhancement. Previous literatures\nshow that alloying elements robustly alter deformation/recrystallization pathways —\nvia solute drag, precipitation/particle-stimulated nucleation (PSN), and stored-energy\ngradients — and thereby reshape texture and the resulting XRD peak intensities across\ncubic and hcp alloys; for example, Al-Mg alloys containing Sc/Zr form dispersoid-\ndriven PSN that changes recrystallization textures, and Mg alloys with rare-earth (or\nCa/Y) additions systematically weaken or re-orient basal textures while promoting non-\nbasal components.6-10 Therefore, we conclude that the enhanced (200) reflection in the\nXRD of Li-MSSA is best attributed to alloying-driven texture evolution within a Li-like\nbcc solid solution.” (Page 33 in the revised SI)\nComment 13: - The XRD pattern of the Li-MSSA anode in Fig. S2 shows peaks at\n32.5°, 39.1°, and 46.8° that could be from the metallic compound phases. The existence\nof metallic compounds means that the Li-MSSA is not a uniform solid-solution. These\nXRD peaks are not shown in the limited 2theta range of Fig. 4. How do the metallic\ncompound phases evolve during delithiation-lithiation cycling?\nResponse: Thank you for this helpful comment. After re-examining our data, we\nidentified a misassignment in the original Figure S2: the reflections near 32.5°, 39.1°,\nand 46.8° are best ascribed to segregates formed during non-equilibrium solidification\nrather than to intrinsic features of a homogeneous alloy matrix. In other words, these\nsignals are parasitic and do not represent a distinct bulk alloy phase.\nFrom a metallurgical standpoint, compositional microsegregation is expected to\nsome extent whenever solidification proceeds under non-equilibrium conditions,\n28\nproducing localized solute fluctuations on ~10 – 100 µm length scales (Castings\nPrinciples: The New Metallurgy of Cast Metals 2nd ed, 2003). Such microsegregation\ncan be substantially reduced by post-solidification homogenization below the solidus,\nfollowed by rapid cooling to retain the uniform solid solution (Porter & Easterling,\nPhase Transformations in Metals and Alloys, 1981). In our original synthesis, the Li-\nMSSA melt was simply quenched from 600 °C, which suppresses long-range diffusion\nbut can freeze-in interdendritic microsegregation. To mitigate this artefact, we re-\nprepared Li-MSSA and introduced a mild subsolidus homogenization (200 °C, ~24 h)\nbefore rapid cooling to retain the solid solution. Ex-situ XRD of the homogenized\nmaterial shows only Li-like bcc reflections at 36.2°, 52.0°, and 65.0° (Figure R15; also\nFigure S2 in the revised SI), with no additional peaks attributable to segregated alloy\nphases.\nFor the in-situ XRD, we re-extracted and compared the end-of-lithiation and end-\nof-delithiation patterns for cycles 1 and 4 (Figure R14; also Figure S18 in the revised\nSI). As the Be-windowed cell introduces strong Be reflections and an elevated diffuse\nbackground in the 30° – 50° range, which markedly reduces the visibility of weak\nfeatures; under this geometry even the intense (200) near ~52.0° in ex-situ XRD is\nattenuated, so any much weaker signals at 32.5°, 39.1°, and 46.8° are not discernible.\nBeyond the principal bcc-type reflections highlighted in the main text (36.2°, 52.0°,\n65.0°), we observe in some scans (e.g., cycle 4), broad features at 34.9°, 43.1°, and\n57.3°. Owing to the multicomponent nature of the alloy, these features cannot be\nunambiguously assigned to specific intermetallic phases at present. Given the\ncombinatorial phase space of multicomponent alloys, future efforts could employ\nmachine-learning-assisted phase fingerprinting (e.g., simulated pattern libraries\ncombined with Bayesian or neural-network classifiers), together with CALPHAD-\nguided searches and pair-distribution-function analyses, to enable more robust phase\nidentification. In this work, however, our focus is the design concept of a\nmulticomponent solid-solution alloy that leverages a continuous chemical-potential\ngradient in the bulk and high configurational entropy to provide multiple, percolating\ndiffusion pathways, thereby enhancing Li mobility and enabling stable cycling under\nhigh Li mass fraction and high Li utilization conditions. Determining if the minor\nfeatures are bona fide secondary phases and, if so, whether they influence long-term\ncycling remains an important direction for future work.\nTo further substantiate this point, we have added the corresponding discussion to\n29\nthe main text:\n“From a metallurgical standpoint, compositional microsegregation is expected to\nsome extent whenever solidification proceeds under non-equilibrium conditions,\nproducing localized solute fluctuations on ~10 – 100 µm length scales.3 Such\nmicrosegregation can be substantially reduced by post-solidification homogenization\nbelow the solidus, followed by rapid cooling to retain the uniform solid solution.4 In\nour original synthesis, the Li-MSSA melt was simply quenched from 600 °C, which\nsuppresses long-range diffusion but can freeze-in interdendritic microsegregation. To\nmitigate this artefact, we re-prepared Li-MSSA and introduced a mild subsolidus\nhomogenization (200 °C, ~24 h) before rapid cooling to retain the solid solution. Ex-\nsitu XRD of the homogenized material shows only Li-like bcc reflections at 36.2°, 52.0°,\nand 65.0°, with no additional peaks attributable to segregated alloy phases (Figure S2).”\n(Page 33 in the revised SI)\nFigure R15. XRD patterns of re-annealed Li-MSSA and metallic Li (PDF No. 1-1131).\nComment 14: - In full cells with LiNi Co Mn O cathodes, the Li-MSSA anode\n0.8 0.1 0.1 2\nwill be lithiated first during the charging step of each cycle, while the characterizations\nand tests in Figs. 2, 3, 4 are based on the delithiation-lithiation cycling of Li-MSSA\nfrom Li M to Li M. How does the lithiation-first mode affect the structure evolution\n88 9.9\nand electrochemical behavior of Li-MSSA?\nResponse: We thank the reviewer for this insightful comment. We agree that in\nconventional Li-MSSA||NCM811 full cells the alloy anode undergoes an initial\nlithiation-first step upon charge. Because Li-MSSA remains a bcc-like solid solution in\nthe Li-rich regime, it accommodates incoming Li interstitially within the same matrix;\nconsequently, additional uptake proceeds via bulk solid-solution insertion rather than\nby nucleation of a discrete plated-Li layer at the interface. Consequently, the anode’s\nstructural evolution during charge/discharge is governed by the same single-phase\n30\nmechanism analyzed in Figures 2-4 in the original manuscript.\nTo substantiate this, we harvested Li-MSSA foils from NCM811 full cells at end-\nof-charge (lithiated) after 50 cycles and performed cross-sectional SEM-EDX mapping.\nAs shown in Figure R16 (also Figure S23 in the revised SI), Ag, Cd, Al, and Mg are\nspatially uniform across the foil thickness. It is reasonable to assume that if Li metal\nwas formed and deposited at the foil surface, the cross-sectional SEM-EDX would be\nexpected to show a near-surface depletion band of Ag/Cd/Mg/Al in the immediately\nunderlying alloy region due to dilution/redistribution. Neither signature is observed,\nsupporting inward Li diffusion and homogeneous solid-solution lithiation rather than\nsuperficial Li plating.\nTo further decouple cathode-derived Li, we rebuilt a Li-MSSA||delithiated-\nNCM811 full cell in which all active Li originates from the alloy in anode. The\nNCM811 electrodes (areal capacity 2.3 mAh cm-2) were first charged to 4.3 V vs. Li/Li+\nto extract Li (Figure R17a; also Figure S24a in the revised SI). The fully delithiated\nNCM811 discs were then assembled opposite to Li-MSSA foils (areal capacity 2.5 mAh\ncm-2) in fresh CR2032 cells (Figure R17b; also Figure S24b in the revised SI). Because\nthe cathodes now contain no Li, the cells begin at N/P = 1.1, providing a stringent\nplatform to assess alloy utilization and cycling stability without surplus cathode Li. As\na control, Li/Cu||delithiated-NCM811 cells were prepared, with both the areal capacity\nand N/P matching the Li-MSSA cells. Under identical conditions, the Li-\nMSSA||delithiated-NCM811 cell delivered 2.3 mAh cm-2 initially and retained 2.1 mAh\ncm-2 after 100 cycles, whereas the Li/Cu||delithiated-NCM811 control cells showed\nfluctuating CE and rapid capacity fade (Figure R17c; also Figure S24c in the revised\nSI). These results highlight the superior stability of Li utilization in Li-MSSA and\nindicate that the alloy’s behavior remains robust during lithiation-first operation even\nunder low-Li-excess, high-loading conditions.\nWe have added relevant discussion in the revised manuscript:\n“In conventional Li-MSSA||LiNi Co Mn O full cells, the alloy anode\n0.8 0.1 0.1 2\nexperiences a lithiation-first step upon the initial charge; Li-MSSA therefore\naccommodates incoming Li interstitially within the matrix. To verify that the structural\n31\nevolution during charge/discharge follows the same single-phase mechanism analyzed\nin Figs. 2-4, we harvested Li-MSSA foils from NCM811 full cells at end of charge\n(lithiated) after 50 cycles and performed cross-sectional SEM-EDX mapping. As shown\nin Supplementary Fig. 23, Ag, Cd, Al and Mg are spatially uniform across the foil\nthickness. It is reasonable to assume that if Li metal was formed and deposited at the\nfoil surface, the cross-sectional SEM-EDX would be expected to show a near-surface\ndepletion band of Ag/Cd/Mg/Al in the immediately underlying alloy region due to\ndilution/redistribution. Neither signature is observed, supporting inward Li diffusion\nand homogeneous solid-solution lithiation rather than superficial Li plating. To further\ndecouple cathode-derived Li, we rebuilt a Li-MSSA||delithiated-NCM811 full cell in\nwhich all active Li originates from the alloy (Supplementary Fig. 24). Commercial\nNCM811 electrodes (areal capacity 2.3 mAh cm-2) were first charged to 4.3 V vs. Li/Li+\nto extract Li (Supplementary Fig. 24a). The fully delithiated NCM811 discs were then\nassembled opposite Li-MSSA foils (areal capacity 2.5 mAh cm-2) in fresh CR2032 cells\n(Supplementary Fig. 24b). Because the cathode now contains no Li, the cell begins at\nN/P = 1.1, providing a stringent platform to assess alloy utilization and cycling stability\nwithout surplus cathode Li. As a control, Li/Cu||delithiated-NCM811 cells were\nprepared with areal capacity and N/P ratio matching the Li-MSSA||delithiated-\nNCM811 cells. Under identical conditions, the Li-MSSA||delithiated-NCM811 cell\ndelivered 2.3 mAh cm-2 initially and retained 2.1 mAh cm-2 after 100 cycles, whereas\nthe Li/Cu control showed fluctuating CE and rapid capacity fade (Supplementary Fig.\n24c). These results highlight the superior stability of Li utilization in Li-MSSA and\nindicate that the alloy’s behavior remains robust during lithiation-first operation even\nunder low-Li-excess, high-loading conditions.” (Line 12, Page 19)\nFigure R16. (a) Cross-sectional scanning electron micrograph of Li-MSSA foils\nharvested from Li-MSSA∥NCM811 full cells at end-of-charge (lithiated state) after 50\ncycles, together with the corresponding cross-sectional EDX elemental maps. (b) Area-\nintegrated EDX spectrum (0-8 keV) acquired from the region in (a). Peaks for Mg-Kα,\nadventitious carbon and surface oxidation. A weak S-Kα (2.31 keV) signal originates\nfrom trace contamination.\nFigure R17. (a) Charge/discharge curve of NCM811 cathode. The areal capacity of the\ncathode was measured to be 2.3 mA h cm-2 by delithiation of the NCM811 electrode.\n(b) Charge/discharge curve of the Li-MSSA anode. The Li-MSSA anode with areal\ncapacity of 2.5 mA h cm-2 was used as the anode to ensure the anode/cathode capacity\nratio to be 1.1 in the Li-MSSA||delithiated-NCM811. (c) Cycling performance of Li-\nMSSA and Li/Cu anodes paired with a delithiated-NCM811 cathode under a 0.5 C\ncharge/1 C discharge protocol; all full cells were balanced at N/P = 1.1.\nComment 15: -How was the Li utilization of the Li-MSSA anode calculated when it is\npaired with LiNi Co Mn O cathodes? This is a bit confusing.\n0.8 0.1 0.1 2\nResponse: We thank the reviewer for pointing out the potential confusion. We agree\nthat describing N/P ≈ 1.1 as “rigorous” in a cell where the as-received NCM811 cathode\nalready contains Li can be misleading, because any Li stored in the Li-MSSA anode is\nthen, by definition, excess relative to the cathode inventory. This shorthand is common\nin the Li-metal literatures for Li-foil||NCM811 cells (where N/P is quoted from the\nanode side), but it does not by itself provide a rigorous Li-utilization metric for the\nanode in a Li-containing-cathode configuration (Nat. Energy 2024, 9, 987-998; Nature\n2025, 643, 1255-1262; Nat. Energy 2025. doi.org/10.1038/s41560-025-01842-5). We\nhave revised the main text to avoid this wording.\nS4, the Li-MSSA anode doesn’t have a current collector. Is just Li-MSSA foil used as\nthe anode in the 1 Ah pouch cell? How is the thickness of the Li-MSSA anode used for\nthe pouch cell? The details of pouch cell fabrication should be provided in the methods.\nResponse: Thank you for raising this point. In the 1 Ah pouch cells we use a free-\nstanding Li-MSSA foil as both active layer and current collector; no additional Cu\nsubstrate is employed. The foil is prepared by hot-rolling a cast Li-MSSA ingot to a\nthickness of 40 µm, then slit, calendared, and spot-welded to a Ni tab before stacking.\nThe pouch cells were fabricated on a semi-automated pilot line in a dry room (dew-\npoint -55 °C, RH < 0.1 %). The NCM811 cathode slurry was double-side coated onto\n10 µm Al foil, delivering 3.5 mAh cm-2 per side (7.0 mAh cm-2 total), while the Li-\nMSSA anode provided 7.8 mAh cm-2. Cathode and anode sheets were Z-stacked with\n12 µm polypropylene separators, tab-welded (Al for the cathode, Ni for the anode),\npackaged in 85 mm × 65 mm laminated pouches, vacuum-dried (60 °C, 12 h), and\nvacuum-injected with 1.5 mol L-1 LiFSI in 1,2-dimethoxyethane/1,1,2,2-\ntetrafluoroethyl-2,2,3,3-tetrafluoropropyl ether (1.2:3, v/v) at an electrolyte-to-capacity\nratio of 2 g Ah-1. After two C/20 formation cycles, the cells were charged at 0.1 C and\ndischarged at 1 C between 2.6 V and 4.3 V on a LAND tester at 25 °C.\nThese points are now made explicit in the revised manuscript:\n“In the 1 Ah pouch cells we use a free-standing Li-MSSA foil as both active layer\nand current collector; no additional Cu substrate is employed. The foil is prepared by\nhot-rolling a cast Li-MSSA ingot to a thickness of 40 µm, then slit, calendared, and\nspot-welded to a Ni tab before stacking. The pouch cells were fabricated on a semi-\nautomated pilot line in a dry room (dew-point -55 °C, RH < 0.1 %). The NCM811\ncathode slurry was double-side coated onto 10 µm Al foil, delivering 3.5 mAh cm-2 per\nside (7.0 mAh cm-2 total). We use a free-standing Li-MSSA foil that serves as both the\nactive layer and the current collector; no Cu substrate is employed. The foil is produced\nby hot-rolling a cast Li-MSSA ingot to 40 µm, then slit, calendared, and spot-welded to\na Ni tab before stacking. The Li-MSSA anode provides an areal capacity of 7.8 mAh\ncm-2. Cathode and anode sheets were Z-stacked with 12 µm polypropylene separators,\ntab-welded (Al for the cathode, Ni for the anode), packaged in 85 mm × 65 mm\nlaminated pouches, vacuum-dried (60 °C, 12 h), and vacuum-injected with 1.5 mol L-1\nLiFSI in 1,2-dimethoxyethane/1,1,2,2-tetrafluoroethyl-2,2,3,3-tetrafluoropropyl ether\n(1.2 : 3, v/v) at an electrolyte-to-capacity ratio of 2 g Ah-1. After two C/20 formation\n34\ncycles, the cells were charged at 0.1 C and discharged at 1 C between 2.6 V and 4.3 V\non a LAND tester at 25 °C.” (Line 23, Page 23)\nComment 17: -Line 50, “Li-Ag binary solids-solution alloy” should be “Li-Ag binary\nsolid-solution alloy”\nResponse: Thank you for noticing; we have corrected the phrase to “Li-Ag binary\nsolid-solution alloy” in the revised manuscript.\nI co-reviewed this manuscript with one of the reviewers who provided the listed reports.\nThis is part of the Nature Communications initiative to facilitate training in peer review\nand to provide appropriate recognition for Early Career Researchers who co-review\nmanuscripts.\nResponse: We would like to extend our heartfelt gratitude to Reviewer #3 for the\nmeticulous review of our manuscript. We also wish to convey our appreciation to\nReviewer #3 for generously dedicating precious time and expertise to assess our work.\nIn this paper the authors report the synthesis of a lithium alloy containing Cd, Ag, Mg\nand Al and test its performance as an anode material. The alloy composition developed\nis novel and the authors claim it to form a solid solution with a high lithium diffusivity\nthat enables a high lithium utilisation of 89%. This alloy anode is then implemented in\npouch cells with various cathode materials. The results presented in this manuscript and\nits supplementary information do not convincingly support all of these claims, so\nfurther evidence and clarification will be required before it can be considered for\npublication. The following points should be addressed:\nResponse: We express our sincere appreciation to Reviewer #4 for the meticulous\nreview of our manuscript and the invaluable comments provided. These comments have\nplayed a crucial role in the continuous improvement of our manuscript. We have\nthoughtfully revised the manuscript in direct response to these insightful suggestions.\nWe also extend our gratitude to Reviewer #4 for generously dedicating precious time\nand expertise to assessing our work.\nComment 1: The authors claim the Cd, Ag, Mg and Al all form a solid solution with\nthe lithium and the main evidence for this is a single TEM image with a narrow field of\nview (Fig. 2a). The XRD pattern in Fig. S2 does not support this claim and the authors\nacknowledge the peaks at 32.5, 39.1 and 46.8° can be assigned to “metal alloy phases\nof the MSSA.” This suggests the alloy contains additional phases and does not simply\nform a solid solution. The additional phase(s) should be identified and\nmicrostructural characterisation should be performed with a larger-field-of-view using\nSEM, for example.\nResponse: Thank you for this helpful comment. After re-examining our data, we\nidentified a misassignment in the original Figure S2: the reflections near 32.5°, 39.1°,\nand 46.8° are best ascribed to segregates formed during non-equilibrium solidification\nrather than to intrinsic features of a homogeneous alloy matrix. In other words, these\nsignals are parasitic and do not represent a distinct bulk alloy phase.\nFrom a metallurgical standpoint, compositional microsegregation is expected to\nsome extent whenever solidification proceeds under non-equilibrium conditions,\n37\nproducing localized solute fluctuations on ~10 – 100 µm length scales (Castings\nPrinciples: The New Metallurgy of Cast Metals 2nd ed, 2003). Such microsegregation\ncan be substantially reduced by post-solidification homogenization below the solidus,\nfollowed by rapid cooling to retain the uniform solid solution (Porter & Easterling,\nPhase Transformations in Metals and Alloys, 1981). In our original synthesis, the Li-\nMSSA melt was simply quenched from 600 °C, which suppresses long-range diffusion\nbut can freeze-in interdendritic microsegregation. To mitigate this artefact, we re-\nprepared Li-MSSA and introduced a mild subsolidus homogenization (200 °C, ~24 h)\nbefore rapid cooling to retain the uniform solid solution. Ex-situ XRD of the homogenized\nmaterial shows only Li-like bcc reflections at 36.2°, 52.0°, and 65.0° (Figure R18; also\nFigure S2 in the revised SI), with no additional peaks attributable to segregated alloy\nphases.\nTo address the request for larger fields of view, we now provide large-area SEM\nand SEM-EDX of the original alloys in Figure R19 (also Figure S4 in the revised SI).\nBefore cycling, the electrode surface is smooth with no obvious particulate features,\nand EDX maps show uniform element distributions. After cycling, particle-like deposits\nbecome visible on the surface (Figure R20; also Figure S10 in the revised SI) —\nmorphologically consistent with our prior observations for solid-solution alloys under\nplating/stripping (J. Am. Chem. Soc. 2020, 142, 8818-8826) — while SEM-EDX maps\nstill indicate uniform distributions of Ag, Cd, Mg, and Al over the mapped areas.\nIn response to the reviewer’s suggestion, we have updated the relevant data and\nrefined the discussion accordingly in the revised manuscript and SI:\n“Atomically resolved high-angle annular dark-field scanning transmission electron\nmicroscopy (HAADF-STEM) with energy-dispersive X-ray spectroscopy (EDX)\nspectrum imaging (Fig. 2a and Supplementary Fig. 3) shows that Ag, Cd, Mg, and Al\nare co-spatial within our spatial-resolution and counting-statistics limits, with no\ndetectable nanoscale segregation; large-area scanning electron microscopy coupled\nwith (SEM-EDX) maps corroborate this (Supplementary Fig. 4).” (Line 9, Page 8)\n“This finding is further confirmed by large-area SEM-EDX mapping of the cycled\nLi-MSSA (Supplementary Fig. 10).” (Line 4, Page 12)\n“From a metallurgical standpoint, compositional microsegregation is expected to\nsome extent whenever solidification proceeds under non-equilibrium conditions,\nproducing localized solute fluctuations on ~10 – 100 µm length scales.3 Such\nmicrosegregation can be substantially reduced by post-solidification homogenization\nbelow the solidus, followed by rapid cooling to retain the uniform solid solution.4 In\nour original synthesis, the Li-MSSA melt was simply quenched from 600 °C, which\nsuppresses long-range diffusion but can freeze-in interdendritic microsegregation. To\n38\nmitigate this artefact, we re-prepared Li-MSSA and introduced a mild subsolidus\nhomogenization (200 °C, ~24 h) before rapid cooling to retain the uniform solid solution. Ex-\nsitu XRD of the homogenized material shows only Li-like bcc reflections at 36.2°, 52.0°,\nand 65.0°, with no additional peaks attributable to segregated alloy phases (Figure S2).”\n(Page 33 in the revised SI)\nFigure R18. XRD patterns of re-annealed Li-MSSA and metallic Li (PDF No. 1-1131).\nFigure R19. (a) SEM image and corresponding EDX elemental maps for Mg, Cd, Al,\nand Ag of the as-prepared Li-MSSA alloy. (b) Area-integrated EDX spectrum (0-8 keV)\nfrom the region in (a). Peaks for C-K and O-K arise from adventitious carbon and\nsurface oxides formed by air exposure (surface oxidation), rather than bulk constituents.\nof the Li-MSSA electrode in the lithiated state after 10 cycles at 2 mA cm-2 with 6 mAh\ncm-2 per half-cycle. (b) Area-integrated EDX spectrum (0-8 keV) from the region in (a).\nPeaks for Mg-Kα, Al-Kα, Ag-L, and Cd-Lα/Lβ match the alloy composition; C-K and\nO-K arise from adventitious carbon and surface oxides formed upon brief air exposure.\nA F-Kα (0.677 keV) signal can originate from SEI, while the Si-Kα (~1.74 keV) signal\narises from Si-bearing substrate/adhesive contamination.\nComment 2: Additionally, the lithium concentration in the electrode will decrease\nsignificantly during discharge (to approximately 10 wt.% using the reported utilisation).\nHow does this affect the phases present in the alloy and their distribution? Fig. S7 shows\nclear segregation of Al in the alloy after cycling.\nResponse: We thank the reviewer for highlighting this point. Even when the lithium\ncontent is reduced to ~10 wt.% at the discharge endpoints, the bulk alloy remains a\nsingle-phase, Li-like bcc solid solution. Regarding Figure S7 in the original manuscript,\nthe apparent Al enrichment in the initial EDX map most likely arose from measurement\nartifacts (substrate pickup) and/or sample-preparation effects; after remounting on a Cu\nholder and re-acquiring the maps with appropriate background correction, the Al\ndistribution is uniform within noise.\nTo directly test whether deep delithiation induces compositional partitioning, we\nacquired atomically resolved STEM-EDX spectrum-imaging on delithiated electrodes\n(Figure R21; also Figure S11 in the revised SI). In the probed regions, the Ag, Cd, Mg,\nand Al maps are co-spatial, with no nanoscale segregation or intermetallic domains\nobserved. Because Li cannot be quantified reliably by EDX, we additionally performed\n40\nspectrum-imaging EELS at the Li-K edge (Figure R22; also Figure S12 in the revised\nSI), which yields Li maps that are spatially uniform over hundreds of nanometres, with\nno Li-rich/Li-poor domains detectable within counting-statistics-limited sensitivity.\nIn parallel, hybrid Monte Carlo and molecular dynamics (MCMD) simulations starting\nfrom randomized atomic configurations for the delithiated state (300 K, 0.5 ns) show\nno tendency toward elemental clustering (Figure R23; see also Figure 3b in the revised\nmanuscript and Figure S14 in the revised SI). While the MCMD timescale is necessarily\nlimited, the convergent microscopy/spectroscopy and simulation evidence is consistent\nwith a random solid-solution manifold; i.e., the delithiated Li-MSSA remains\ncompositionally uniform rather than decomposing into element-rich regions or\nintermetallic secondary phases.\nFigure S7 in the original SI was intended to demonstrate that the lithiated Li-MSSA\nelectrode retains a uniform elemental distribution after cycling. Consistently, in Figure\n3a of the original manuscript we assessed through-thickness composition in the lithiated\nstate by ex situ ToF-SIMS: the 3D reconstructions of fully cycled Li-MSSA show nearly\noverlapping Li, Cd, Ag, Mg, and Al signals, indicating a uniform distribution. In this\nrevision, after switching to a Cu holder and preparing fresh specimens, we further added\ncross-sectional SEM-EDS mapping (Mg, Cd, Al, Ag) of lithiated Li-MSSA. As shown\nin Figure R24 (see also Figure S23 in the revised SI), Ag, Cd, Al, and Mg are spatially\nuniform across the foil thickness. In addition, MCMD simulations were employed to\nevaluate phase stability and elemental mixing during lithiation across representative Li\ncontents (Figure R25; see also Figure 3b in the revised manuscript and Figure S13 in\nthe revised SI); across multiple randomized initial configurations, 300 K trajectories\n(0.5 ns) show no tendency toward elemental clustering after lithiation, and the alloy\nremains compositionally uniform within the simulation cell.\nTaken together, the STEM-EDX/EELS, ToF-SIMS/SEM-EDX, and MC/MD\nresults indicate that both the lithiated and deeply delithiated electrodes are single-phase\nsolid solutions without detectable elemental partitioning within our sensitivity. These\npoints are now made explicit in the main text and SI:\n“To test whether Li-MSSA remains a solid solution after full delithiation, we also\nacquired atomically resolved STEM-EDX and Li-K-edge EELS on delithiated\nelectrodes (Supplementary Fig. 11). The Ag, Cd, Mg, and Al maps remain co-spatial\nwith no nanoscale segregation or intermetallic domains observed, and the residual Li-\n(Supplementary Fig. 12). In addition, we further employed hybrid Monte Carlo and\nmolecular dynamics (MCMD) simulations to evaluate phase stability and elemental\nmixing during lithiation/delithiation across representative Li contents (Supplementary\nFigs. 13 and 14). Across multiple randomized initial configurations, 300 K trajectories\n(0.5 ns) show no tendency toward elemental clustering either after lithiation or after\ndelithiation; the alloy remains compositionally uniform within the simulation cell.”\n(Line 9, Page 12)\n“As shown in Supplementary Fig. 23, Ag, Cd, Al and Mg are spatially uniform\nacross the foil thickness. It is reasonable to assume that if Li metal was formed and\ndeposited at the foil surface, the cross-sectional SEM-EDX would be expected to show\na near-surface depletion band of Ag/Cd/Mg/Al in the immediately underlying alloy\nregion due to dilution/redistribution. Neither signature is observed, supporting inward\nLi diffusion and homogeneous solid-solution lithiation rather than superficial Li\nplating.” (Line 18, Page 19)\nFigure R21. (a) HAADF-STEM image of Li-MSSA in the delithiated state, with\ncorresponding EDX elemental maps for Ag, Mg, Cd, and Al, showing a homogeneous\ndistribution with no detectable nanoscale segregation. (b) Integrated EDX spectrum (0-\n8 keV) from the region in (a). The Mo-L features arise from Mo in the specimen holder\n(holder peaks). The ~0.68 keV feature has two contributions: (i) a Si-escape peak of\nfluorinated SEI species formed by electrolyte decomposition.\nFigure R22. (a) HAADF-STEM image of the Li-MSSA in the delithiated state. (b)\nEELS Li-K elemental map from the same region acquired by spectrum imaging,\nshowing a uniform Li distribution. (c) Background-subtracted EELS spectrum with the\nLi-K edge near 60 eV, confirming the presence of Li in the matrix.\nFigure R23. (a) Time evolution of the potential and kinetic energies during a hybrid\nMC/MD simulation of a 9423-atom Li AgMgCdAl (Li M) cell at 300 K; data are\n39.6 9.9\nsampled every 0.5 ps. (b) Equilibrated snapshot of the supercell (Li, green; Ag, grey;\nAl, blue; Cd, orange; Mg, magenta; periodic boundary conditions). (c-f) Element-\nresolved projections highlighting the spatial distributions of Ag, Al, Cd, and Mg,\nrespectively, showing homogeneous dispersion in the Li matrix with no obvious\nclustering or nanoscale segregation at the simulation length scale.\nharvested from Li-MSSA∥NCM811 full cells at end-of-charge (lithiated state) after 50\ncycles, together with the corresponding cross-sectional EDX elemental maps. (b) Area-\nintegrated EDX spectrum (0-8 keV) acquired from the region in (a). Peaks for Mg-Kα,\nAl-Kα, Ag-L, and Cd-Lα/Lβ match the alloy composition; C-K and O-K arise from\nadventitious carbon and surface oxidation. A weak S-Kα (2.31 keV) signal originates\nfrom trace contamination.\nFigure R25. (a) Time evolution of the potential and kinetic energies during a hybrid\nMC/MD simulation of a 9423-atom Li AgMgCdAl (Li M) cell at 300 K; data are\n352 88\nsampled every 0.5 ps. (b) Equilibrated snapshot of the supercell (Li, green; Ag, grey;\nAl, blue; Cd, orange; Mg, magenta; periodic boundary conditions). (c) Element-\nresolved projections highlighting the spatial distributions of Ag, Al, Cd, and Mg,\nrespectively.\nComment 3: It is unclear how the practical gravimetric capacity of the alloy electrode\n44\nis determined. The manuscript states that “Li was extracted from the Li-MSSA within\nan electrochemical cell comprising Cu foil as the counter electrode,” but the voltage\ncurves presented in Fig. 2f and S5 do not appear to be consistent with this. Please could\nadditional details be provided to describe how these measurements were conducted?\nResponse: We realised the reviewer’s confusion stemmed from the way the original\ncurves in Figures 2f and S5 were displayed: charge and discharge were over-laid into a\nsingle closed loop, making it hard to tell which portion represented Li extraction from\nthe alloy and which represented Li replating. We have therefore re-plotted the data as\nseparate charge (stripping) and discharge (plating) traces, as shown in the updated\nfigure above (Figures R26 and R27; also Figure 2d in the revised manuscript and Figure\nS8 in the revised SI).\nSince freshly prepared Li M foil is in the lithiated state with pre-stored Li, to\n88\nmeasure the practical gravimetric capacity of the Li M, a pre-delithiation process was\n88\nfirst performed. This pre-delithiation process was carried out at the constant current to\nreach a voltage of 0.1 V, which yielded a capacity. When charging the Li M to Li M,\n88 9.9\nwhich means that the number of the lithium atom extracted from the fresh Li M is (88-\n88\n9.9=78.1), since the theoretical specific capacity of Li metal is 3860 mA h g-1, the\ntheoretical capacity of Li M when charged to Li M can be calculated to be:\n88 9.9\nTheoretical capacity\n= 3860 mA h g−1\n78.1×M\nLi\n×\n88×M +0.25×(M +M +M +M )\nLi Mg Ag Cd Al\nM , M , M , M , and M are the atomic weights of Li, Mg, Ag, Cd, and Al.\nLi Mg Ag Cd Al\nThese points are now made explicit in the revised SI:\n“Since freshly prepared Li M foil is in the lithiated state with pre-stored Li, to\n88\nmeasure the practical gravimetric capacity of the Li M, a pre-delithiation process was\n88\nfirst performed. This pre-delithiation process was carried out at the constant current to\nreach a voltage of 0.1 V, which yielded a capacity. When charging the Li M to Li M,\n88 9.9\nwhich means that the number of the lithium atom extracted from the fresh Li M is (88-\n88\n9.9=78.1), since the theoretical specific capacity of Li metal is 3860 mA h g-1, the\ntheoretical capacity of Li M when charged to Li M can be calculated to be:\n88 9.9\n45\n𝑇ℎ𝑒𝑜𝑟𝑒𝑡𝑖𝑐𝑎𝑙 𝑐𝑎𝑝𝑎𝑐𝑖𝑡𝑦\n= 3860 𝑚𝐴 ℎ 𝑔−1\n78.1×𝑀\n𝐿𝑖\n×\n88×𝑀 +0.25×(𝑀 +𝑀 +𝑀 +𝑀 )\n𝐿𝑖 𝑀𝑔 𝐴𝑔 𝐶𝑑 𝐴𝑙\n𝑀 , 𝑀 , 𝑀 , 𝑀 , and 𝑀 are the atomic weights of Li, Mg, Ag, Cd, and\n𝐿𝑖 𝑀𝑔 𝐴𝑔 𝐶𝑑 𝐴𝑙\nAl.” (Page 34 in the revised SI)\nFigure R26. Voltage profiles of the Li-MSSA anode acquired at 2 mA cm-2, which\ndelivered a gravimetric capacity of 3100 mAh g-1 with an areal capacity of 6 mAh cm-\n2.\nFigure R27. Voltage profiles of Li-MSSA composite electrodes with capacities of (a)\n1, (b) 5, (c) 10, and (d) 20 mAh cm-2.\nComment 4: Similarly, insufficient details have been provided about how the GITT\nmeasurement was conducted and how the data was analysed. The pulse current density,\npulse duration and rest period duration should be stated. It would also be useful to relate\nthe time to a composition/state-of-charge of the alloy electrode.\nResponse: We thank the reviewer for this helpful comment. Because Li-MSSA exhibits\nfinite Li solubility, the bulk lithium chemical potential varies continuously with\ncomposition, and transport is therefore driven by chemical-potential gradients (J. Am.\nChem. Soc. 2023, 145, 24775-24784). Accordingly, we use a non-ideal, chemical-\npotential-based diffusion formulation to fit the full GITT voltage transients\n46\nrecommended by Horner et al. (ACS Appl. Energy Mater. 2021, 4, 11460-11469). The\nGITT protocol employed 0.01 C, 30 s current pulses with 1800 s rests between pulses\n(Figure R28; also Figure S15 in the revised SI). While Horner et al. implemented their\nmethod for a one-dimensional single-particle (spherical) model, our system is a one-\ndimensional slab electrode; thus, we replace the radial coordinate 𝑟 with the thickness\ncoordinate 𝐿, the particle radius 𝑅 with the foil thickness 𝐿, and assume diffusion\n𝑃\noccurs solely along the thickness direction. The governing equation is thus the 1D slab\nanalogue of Fick’s law with a thermodynamic factor (non-ideal transport), solved with\nfinite-length boundary conditions and fitted to the entire current-voltage transient. From\nfits to the this dataset, the extracted Li chemical diffusivity remains 6  10-9 cm2 s-1\n(Figure R29; also Figure S16 in the revised SI). As expected, applying other analysis\nroutes to the same dataset — such as the traditional GITT early-time √𝜏\napproximation, or EIS-based diffusivity extraction — yields value up to one or two\norder of magnitude higher relative to our non-ideal, chemical-potential-driven fit (ACS\nAppl. Energy Mater. 2021, 4, 11460-11469). Notably, these values are higher than\nreported room-temperature self-diffusion of Li in metallic Li (typically 10-11 cm2 s-1,\nGlass Phys. Chem. 2010, 36, 570-574). Taken together, the fitted 𝐷 is sufficiently\nhigh to accommodate the per-pulse Li flux by bulk insertion into Li-MSSA.\nIn response to the request, we reprocessed the dataset and now report the chemical\ndiffusion coefficient 𝐷 as a function of composition/SOC. The text, figures, and\nconclusions have been updated accordingly.\n“Because the solid-solution alloy has a finite Li solubility, the bulk Li chemical\npotential varies continuously with composition and transport is therefore chemical-\npotential-driven.10 We therefore employed a non-ideal, chemical-potential-based\ndiffusion formulation to fit the full voltage transients of the GITT dataset, using 0.01 C\npulses of 30 s with 1800 s rests between pulses (Supplementary Fig. 15).40 Under this\nprotocol, the fitted Li chemical diffusivities are around 6  10-9 cm2 s-1 (Supplementary\nFig. 16). As expected, applying other analysis routes to the same dataset — such as the\ntraditional GITT early-time √𝜏 approximation, or EIS-based diffusivity extraction —\nyields values up to one or two orders of magnitude higher relative to our non-ideal,\nchemical-potential-driven fit.40 Notably, these values are higher than the reported\nroom-temperature self-diffusion of Li in metallic Li (typically 10-11 cm2 s-1).41 Taken\ntogether, the fitted 𝐷 is sufficiently high to accommodate the per-pulse Li flux by bulk\n47\ninsertion into Li-MSSA.” (Line 3, Page 13)\nFigure R28. GITT curve for calculating the Li atom diffusion coefficient. The inset\nenlarges the initial 0 – 15 h segment.\nFigure R29. Comparison of diffusion coefficients extracted from galvanostatic\nintermittent titration transients using two models: Ideal model and Non-ideal model.\nComment 5: Finally, the claim that an N/P ratio of approximately 1.1 is “rigorous”\ncould be misleading. While an N/P ratio of 1.1 would be suitable for testing a lithium-\nion battery, in this case both the cathode and anode are fully lithiated during cell\nassembly so all the lithium contained in the anode will be an excess. This is likely to\ninflate the cycling performance presented in Fig. 5. Experimental details describing\npouch cell preparation are also missing from the methods section and it is unclear why\nno anode current collector has been included in the calculations in Tables S4 and S5.\nResponse: We thank the reviewer for pointing out the potential confusion. We agree\nthat describing N/P ≈ 1.1 as “rigorous” in a cell where the as-received NCM811 cathode\nalready contains Li can be misleading, because any Li stored in the Li-MSSA anode is\nthen, by definition, excess relative to the cathode inventory. This shorthand is common\nin the Li-metal literatures for Li-foil║NCM811 cells (where N/P is quoted from the\nanode side), but it does not by itself provide a rigorous Li-utilization metric for the\nanode in a Li-containing-cathode configuration (Nat. Energy 2024, 9, 987-998; Nature\n2025, 643, 1255-1262; Nat. Energy 2025. doi.org/10.1038/s41560-025-01842-5).\nNCM811 full cell in which all active Li originates from the alloy in anode. The\nNCM811 electrodes (areal capacity 2.3 mAh cm-2) were first charged to 4.3 V vs. Li/Li+\nto extract Li (Figure R30a; also Figure S24a in the revised SI). The fully delithiated\nNCM811 discs were then assembled opposite to Li-MSSA foils (areal capacity 2.5 mAh\ncm-2) in fresh CR2032 cells (Figure R30b; also Figure S24b in the revised SI). Because\nthe cathodes now contain no Li, the cells begin at N/P = 1.1, providing a stringent\nplatform to assess alloy utilization and cycling stability without surplus cathode Li. As\na control, Li/Cu||delithiated-NCM811 cells were prepared, with both the areal capacity\nand N/P matching the Li-MSSA cells. Under identical conditions, the Li-\nMSSA||delithiated-NCM811 cell delivered 2.3 mAh cm-2 initially and retained 2.1 mAh\ncm-2 after 100 cycles, whereas the Li/Cu||delithiated-NCM811 control cells showed\nfluctuating CE and rapid capacity fade (Figure R30c; also Figure S24c in the revised\nSI). These results highlight the superior stability of Li utilization in Li-MSSA and\nindicate that the alloy’s behavior remains robust during lithiation-first operation even\nunder low-Li-excess, high-loading conditions.\nIn the 1 Ah pouch cells we use a free-standing Li-MSSA foil as both active layer\nand current collector; no additional Cu substrate is employed. The foil is prepared by\nhot-rolling a cast Li-MSSA ingot to a thickness of 40 µm, then slit, calendared, and\nspot-welded to a Ni tab before stacking. The pouch cells were fabricated on a semi-\nautomated pilot line in a dry room (dew-point -55 °C, RH < 0.1 %). The NCM811\ncathode slurry was double-side coated onto 10 µm Al foil, delivering 3.5 mAh cm-2 per\nside (7.0 mAh cm-2 total), while the Li-MSSA anode provided 7.8 mAh cm-2. Cathode\nand anode sheets were Z-stacked with 12 µm polypropylene separators, tab-welded (Al\nfor the cathode, Ni for the anode), packaged in 85 mm × 65 mm laminated pouches,\nvacuum-dried (60 °C, 12 h), and vacuum-injected with 1.5 mol L-1 LiFSI in 1,2-\ndimethoxyethane/1,1,2,2-tetrafluoroethyl-2,2,3,3-tetrafluoropropyl ether (1.2 : 3, v/v)\nat an electrolyte-to-capacity ratio of 2 g Ah-1. After two C/20 formation cycles, the cells\nwere charged at 0.1 C and discharged at 1 C between 2.6 V and 4.3 V on a LAND tester\nat 25 °C.\nWe have added relevant discussion in the revised manuscript:\n“In conventional Li-MSSA||LiNi Co Mn O full cells, the alloy anode\n0.8 0.1 0.1 2\nexperiences a lithiation-first step upon the initial charge; Li-MSSA therefore\naccommodates incoming Li interstitially within the matrix. To verify that the structural\n49\nevolution during charge/discharge follows the same single-phase mechanism analyzed\nin Figs. 2-4, we harvested Li-MSSA foils from NCM811 full cells at end of charge\n(lithiated) after 50 cycles and performed cross-sectional SEM-EDX mapping. As shown\nin Supplementary Fig. 23, Ag, Cd, Al and Mg are spatially uniform across the foil\nthickness. It is reasonable to assume that if Li metal was formed and deposited at the\nfoil surface, the cross-sectional SEM-EDX would be expected to show a near-surface\ndepletion band of Ag/Cd/Mg/Al in the immediately underlying alloy region due to\ndilution/redistribution. Neither signature is observed, supporting inward Li diffusion\nand homogeneous solid-solution lithiation rather than superficial Li plating. To further\ndecouple cathode-derived Li, we rebuilt a Li-MSSA||delithiated-NCM811 full cell in\nwhich all active Li originates from the alloy (Supplementary Fig. 24). Commercial\nNCM811 electrodes (areal capacity 2.3 mAh cm-2) were first charged to 4.3 V vs. Li/Li+\nto extract Li (Supplementary Fig. 24a). The fully delithiated NCM811 discs were then\nassembled opposite Li-MSSA foils (areal capacity 2.5 mAh cm-2) in fresh CR2032 cells\n(Supplementary Fig. 24b). Because the cathode now contains no Li, the cell begins at\nN/P = 1.1, providing a stringent platform to assess alloy utilization and cycling stability\nwithout surplus cathode Li. As a control, Li/Cu||delithiated-NCM811 cells were\nprepared with areal capacity and N/P ratio matching the Li-MSSA||delithiated-\nNCM811 cells. Under identical conditions, the Li-MSSA||delithiated-NCM811 cell\ndelivered 2.3 mAh cm-2 initially and retained 2.1 mAh cm-2 after 100 cycles, whereas\nthe Li/Cu control showed fluctuating CE and rapid capacity fade (Supplementary Fig.\n24c). These results highlight the superior stability of Li utilization in Li-MSSA and\nindicate that the alloy’s behavior remains robust during lithiation-first operation even\nunder low-Li-excess, high-loading conditions.” (Line 12, Page 19)\n“In the 1 Ah pouch cells we use a free-standing Li-MSSA foil as both active layer\nand current collector; no additional Cu substrate is employed. The foil is prepared by\nhot-rolling a cast Li-MSSA ingot to a thickness of 40 µm, then slit, calendared, and\nspot-welded to a Ni tab before stacking. The pouch cells were fabricated on a semi-\nautomated pilot line in a dry room (dew-point -55 °C, RH < 0.1 %). The NCM811\ncathode slurry was double-side coated onto 10 µm Al foil, delivering 3.5 mAh cm-2 per\nside (7.0 mAh cm-2 total). We use a free-standing Li-MSSA foil that serves as both the\nactive layer and the current collector; no Cu substrate is employed. The foil is produced\nby hot-rolling a cast Li-MSSA ingot to 40 µm, then slit, calendared, and spot-welded to\na Ni tab before stacking. The Li-MSSA anode provides an areal capacity of 7.8 mAh\n50\ncm-2. Cathode and anode sheets were Z-stacked with 12 µm polypropylene separators,\ntab-welded (Al for the cathode, Ni for the anode), packaged in 85 mm × 65 mm\nlaminated pouches, vacuum-dried (60 °C, 12 h), and vacuum-injected with 1.5 mol L-1\nLiFSI in 1,2-dimethoxyethane/1,1,2,2-tetrafluoroethyl-2,2,3,3-tetrafluoropropyl ether\n(1.2 : 3, v/v) at an electrolyte-to-capacity ratio of 2 g Ah-1. After two C/20 formation\ncycles, the cells were charged at 0.1 C and discharged at 1 C between 2.6 V and 4.3 V\non a LAND tester at 25 °C.” (Line 23, Page 23)\nFigure R35. (a) Charge/discharge curve of NCM811 cathode. The areal capacity of the\ncathode was measured to be 2.3 mA h cm-2 by delithiation of the NCM811 electrode.\n(b) Charge/discharge curve of the Li-MSSA anode. The Li-MSSA anode with areal\ncapacity of 2.5 mA h cm-2 was used as the anode to ensure the anode/cathode capacity\nratio to be 1.1 in the Li-MSSA||delithiated-NCM811. (c) Cycling performance of Li-\nMSSA and Li/Cu anodes paired with a delithiated-NCM811 cathode under a 0.5 C\ncharge/1 C discharge protocol; all full cells were balanced at N/P = 1.1.\nsuitable revisions accordingly. The reviewers’ comments and our responses are listed\nas follows:\nReviewer #1\nUpon revision with additional experimental evidence, the manuscript is dramatically\nimproved. I do not have any further questions. It is recommended that the authors\nspecify the testing conditions for low-dose STEM, including the dose rate and other\nrelevant parameters, for future reference.\nResponse: We sincerely thank the reviewer for the positive evaluation of our revised\nmanuscript and for the constructive suggestion. In response, we have now explicitly\nspecified the experimental conditions used for low-dose STEM in the Methods section,\nas follows:\n“For Figure 2a, STEM mapping was acquired at an electron accelerating voltage\nof 300 kV with a beam current of 78 pA, a dwell time of 2.0 μs per pixel, and a 1024 ×\n1024 pixel matrix over an 8.6 × 8.6 nm field of view. The total acquisition time for the\nimage was 205 s, corresponding to a cumulative electron dose of 1.34 × 107 e- Å-2. For\nFigure S11, STEM mapping was acquired at 300 kV with a beam current of 78 pA, a\ndwell time of 2.0 μs per pixel, and a 1024 × 1024 pixel matrix over a 6.1 × 6.1 nm field\nof view. The total acquisition time was 188 s, and the cumulative electron dose was 2.44\n× 107 e- Å-2. For Figure S22, STEM mapping was acquired at 300 kV with a beam\ncurrent of 78 pA, a dwell time of 2.0 μs per pixel, and a 1024 × 1024 pixel matrix over\na 12.2 × 12.2 nm field of view. The total acquisition time was 446 s, giving a cumulative\nelectron dose of 1.45 × 107 e- Å-2. For Li STEM-EELS mapping, data were acquired at\nan accelerating voltage of 300 kV with a beam current of 100 pA, a dwell time of 0.005\ns per pixel, and an 86 × 80 pixel matrix over a 694.3 × 645.9 nm field of view. The total\nacquisition time was 34.4 s, corresponding to a cumulative electron dose of 4.79 × 102\ne- Å-2.” (Line 17, Page 23)\nComment 1: However, I do have some concerns about the response to comment 2 of\nReferee #4. Please see my comments below: The authors appear to replace previously\nreported data by attributing the earlier observations to “measurement artifacts (substrate\npickup) and/or sample-preparation effects.” This explanation is overly general and\nscientifically insufficient. Simply invoking generic artifacts does not justify the removal\nor revision of prior results, nor does it offer a mechanistic explanation for the originally\n1\nobserved Al segregation. The manuscript must clearly articulate why Al segregation\nwas detected before, how these alleged artifacts specifically produced such signatures,\nand what experimental evidence demonstrates that the new data are more reliable.\nResponse: We thank Reviewer #1 for the careful reading and agree that our previous\nreply to Comment 2 of Referee #4 was not sufficiently specific. In particular, our\nwording could be interpreted as dismissing the originally reported Al contrast in Figure\nS7 of the original SI as “generic artifacts”, and we apologize for the lack of clarity.\nIn the original SI, the SEM-EDS map in Figure S7 was acquired over a relatively\nlarge field of view (tens of micrometres) and appeared to show a non-uniform Al\ndistribution, with an enhanced Al signal concentrated near the lower/underside surface.\nUpon re-examining the original acquisition conditions, we identified a specific,\nexperiment-related cause that can generate such an apparent Al-rich signature. In the\noriginal experiment, the Li-MSSA foil was mounted on an SEM stub made of Al, and\nthe region selected for mapping was located close to the sample edge. Under this\nconfiguration, the electron interaction volume and scattered electrons near the edge of\nthe sample can partially extend beyond the foil and interact with the underlying Al\nmount, contributing additional Al-K X-ray counts to the detected spectrum. Moreover,\nthe edge geometry and partial shadowing can modify the effective X-ray take-off\nconditions and background, which can further exaggerate the apparent Al contrast\nrelative to the other metals. Therefore, the apparent Al enrichment in the original Figure\nS7 is due to a mounting- and geometry-induced Al contribution, rather than the Al\nsegregation in the alloy.\nIn the previous round of revision, we added Figure S23 in the revised SI. These\nSEM-EDS maps were acquired on freshly prepared specimens mounted on a holder\nmade of Cu, with mapping regions deliberately selected away from the sample edge.\nUnder these controlled conditions, the SEM-EDS maps show nearly uniform\ndistributions of Ag, Cd, Al, and Mg across the foil thickness, and the pronounced Al-\nenhanced contrast seen previously is no longer observed. This comparison supports our\nconclusion that the Al-stub/edge-geometry configuration in the original measurement\nartificially amplified the apparent Al signal.\nComment 2: In addition, the STEM evidence presented is not convincing enough to\nsupport the claim that no aggregation occurs. STEM EDX maps alone do not confirm\natomic-scale homogeneity, and the HAADF-STEM images shown cannot resolve\n2\natomic distribution. Moreover, the regions selected for mapping are extremely localized,\nmaking it difficult to assess whether the conclusions are representative of the entire\nsample.\nResponse: We thank the reviewer for raising this important point regarding the\ninterpretation and representativeness of the STEM data. We agree that, strictly speaking,\nlocalized STEM-EDS maps cannot prove absolute atomic-scale homogeneity across the\nentire electrode.\nIn terms of methodology, the STEM data were acquired in an aberration-corrected\ninstrument with a probe size of ~1 Å and a scan step comparable to the interatomic\nspacing, so that the EDS spectrum-imaging is effectively atomic-column-resolved: each\npixel integrates the signal from a small number of neighbouring atomic columns (J.\nElectron Microsc. 2009, 58, 87-97; J. Microsc. 2015, 259, 165-172; Ultramicroscopy\n2017, 176, 52-62). The HAADF-STEM images themselves are not intended to\ndistinguish individual Ag, Cd, Mg and Al atoms by Z-contrast (their Z values are too\nclose for that), but rather to provide the structural framework (bcc lattice, absence of\nsecondary phases) within which the simultaneously acquired STEM-EDS maps report\nthe local composition.\nWe also recognise the reviewer’s concern regarding locality and representativeness.\nIn practice, assessing compositional uniformity across orders of magnitude in length\nscale inevitably involves a trade-off between spatial resolution and sampling volume.\nSTEM-EDS/EELS offers the highest spatial resolution but is necessarily local; SEM-\nEDS surveys much larger areas (tens of micrometres) but cannot resolve atomic-scale\ndistributions; and ToF-SIMS provides three-dimensional compositional information\nacross larger volumes, albeit with its own matrix effects and quantification limits.\nAccordingly, no single technique can, by itself, prove absolute homogeneity across the\nentire electrode in a statistically exhaustive manner, as this would require\nsimultaneously high chemical sensitivity, a very large sampling volume, and atomic-\nscale spatial resolution. In our study, within the probed STEM volumes, the Ag, Cd, Mg\nand Al signals are co-spatial within our spatial and counting-statistics limits, and no\nresolvable nanoscale clusters or secondary intermetallic domains are observed.\nCrucially, we do not rely on STEM alone to argue for bulk uniformity. Instead, we\ncombine complementary evidence across multiple length scales. STEM-EDS/EELS\nprobes the nanoscale; SEM-EDS and ToF-SIMS provide mesoscale and three-\ndimensional compositional mapping; and MC/MD simulations show no tendency\n3\ntoward bulk clustering in defect-free supercells. Taken together, these results support a\nconsistent picture: Li-MSSA behaves predominantly as a single bcc solid solution on\nthe bulk scale within the sensitivity and detection limits of the techniques employed.\nIn this work, we have therefore made a best-effort, multi-length-scale assessment using\ncomplementary approaches (STEM–EDS/EELS, SEM–EDS, ToF-SIMS, and MC/MD),\nwhich collectively supports our conclusions within their respective detection limits.\nFurther improvements in macroscopic statistical representativeness enabled by\nemerging high-throughput or correlative 3D techniques will be valuable in future work.\nComment 3: Finally, although the authors provide multiple characterizations (STEM-\nEDX/EELS, ToF-SIMS/SEM-EDX, MC/MD simulations), these results still do not\nanswer a critical mechanistic question: when the lithium content decreases to ~10 wt.%\nupon discharge, why does the alloy structure remain unaffected? What structural\nevolution actually occurs at low Li content, and how do these characterizations support\nthe robustness of the proposed alloying mechanism?\nResponse: We thank the reviewer for raising this important mechanistic question. We\nagree that our original discussion focused on what our characterizations show (i.e., that\nthe alloy remains a single bcc phase down to low Li content), but did not sufficiently\nexplain why this structure is so robust when the Li content decreases to ~10 wt.%.\nWe note that wt.% is not the most direct metric for discussing lattice occupancy\nand phase stability in Li-containing alloys, because Li has a much lower atomic mass\nthan the Ag/Cd/Mg/Al atoms. For questions related to whether the bcc framework is\nretained upon delithiation, the atomic fraction is the more relevant descriptor. In our\ncompositional notation, the fully lithiated alloy is approximately Li AgMgCdAl, i.e.,\nstoichiometry is Li AgMgCdAl. Li takes 90.9 at.% of the alloy and remains the\n39.6\ndominant specie by atom, outnumbering the sum of Ag, Mg, Cd, and Al atoms by\napproximately an order of magnitude. We therefore view Li-MSSA as a Li-rich bcc\nsolid solution, in which Li atoms form a bcc host lattice while the minority Ag, Cd, Mg\nand Al atoms occupy substitutional sites and interact strongly with Li and with each\nother. This picture is consistent with our in-situ XRD results, which show a single set\nof bcc reflections throughout the investigated Li-content range. These multicomponent\ninteractions lower the mixing enthalpy and, together with the large configurational\nentropy associated with distributing four different metals on the same sublattice,\n4\nstabilise a disordered bcc solid solution over a broad Li-content window. Within this\npicture, delithiation from Li AgMgCdAl to Li AgMgCdAl does not require a\n352 39.6\nreconstructive transformation into separate intermetallic phases. Instead, Li is removed\nfrom an already-stable multicomponent bcc host, and the structural response is\ndominated by changes in Li occupancy and local bond lengths/strain, plus limited\nsegregation at extended defects, rather than by wholesale breakdown of the bcc\nframework.\nSuch behaviour is consistent with prior observations on intercalation and alloy\nelectrodes, where Li insertion/extraction is accommodated within a single-phase host\nby modest lattice distortions and bond-length changes, rather than repeated\ntransformations between distinct bulk phases. For example, zero-strain Li Ti O\n4 5 12\noperates via a single-phase solid-solution mechanism in which Li storage is buffered\nby subtle internal structural adjustments while the cubic framework is preserved (J.\nElectrochem. Soc. 1995, 142, 1431-1435), and Li-Mg alloy anodes provide a concrete\nexample of Li-rich bcc solid-solution hosts: a single bcc Li-Mg phase extends from ~30\nto nearly 100 at.% Li, allowing Li plating/stripping to be accommodated by continuous,\nrelatively stress-free volume changes within one phase instead of cycling through\nmultiple intermetallics (J. Power Sources 2007, 174, 810-812). High-entropy and other\nmulticomponent solid-solution electrodes likewise demonstrate that configurational\nentropy and multicomponent mixing can stabilise a single crystalline or amorphous host\nstructure during repeated (de)lithiation, suppressing phase separation and mitigating\nvolume change (Nano-Micro Lett. 2025, 17, 22). Within this thermodynamic\nframework, the different characterizations in our work can be viewed as complementary\nevidence that Li-MSSA follows the same general paradigm: Li content varies\nsubstantially, but within a robust single-phase solid-solution manifold.\nWe also wish to emphasise that the primary focus of this work is the design concept\nof a multicomponent solid-solution alloy that leverages a continuous chemical-potential\ngradient in the bulk and high configurational entropy to create multiple, percolating\ndiffusion pathways. This design enhances Li mobility and enables stable cycling under\nhigh Li mass fraction and high Li utilisation conditions. We acknowledge that the\ndeeper physical origin of why the bcc solid-solution manifold remains stable down to\nvery low Li content, including the detailed free-energy landscape and defect chemistry\nas a function of composition, is an important question that is not yet quantitatively\nresolved here. A complete mapping of these thermodynamic and defect-chemical\n5\nfactors across all compositions would be substantial and is beyond the scope of the\npresent study. Nevertheless, we fully agree that such mechanistic understanding is\nvaluable for future materials design, and we highlight it as a key direction for follow-\nup work building on the present results.\nWe added the relevant discussions in the revised manuscript:\n“In parallel, it will be important to elucidate the physical origin of the bcc solid-\nsolution stability down to low Li content, through quantitative mapping of the\ncomposition-dependent free-energy landscape and defect/segregation energetics.”\n(Line 24, Page 21)\nI have examined the revised manuscript, and I commend the authors on doing a very\nnice job on comprehensively addressing all reviewer comments. The paper is clearly\nimproved, as there were a number of issues before. I believe it is acceptable for\npublication.\nResponse: We sincerely thank the reviewer for the careful assessment and the positive\nrecommendation for acceptance. We appreciate the constructive feedback provided in\nthe earlier round, which helped us improve the clarity and rigor of the manuscript. We\nare grateful for the reviewer’s supportive comments and have no further changes to\nrequest.\nI co-reviewed this manuscript with one of the reviewers who provided the listed reports.\nThis is part of the Nature Communications initiative to facilitate training in peer review\nand to provide appropriate recognition for Early Career Researchers who co-review\nmanuscripts.\nResponse: We thank Reviewer #3 for a careful and constructive evaluation. We are\ngrateful for the time and expertise invested, which have helped improve the clarity and\nrigor of the manuscript.\n8"}]

Decision letter