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. 2026 May 26;65(30):e9377760. doi: 10.1002/anie.9377760

Self‐Limiting Covalent Ligation Mechanism Enabling Anomalously High Interfacial Compatibility in Organic‐in‐Sulfide All‐Solid‐State Lithium Batteries

Yan Zhang 1,2, Yuxi Zhong 1, Ruiqi Guo 1,2, Songjie Li 1, Kaidi Kang 1,2, Bingguang Ye 1, Wei Shan 1,2, Ran Zhao 1,2,3, Xinran Wang 1,2,3,✉, Chuan Wu 1,2,3,✉, Ying Bai 1,2,3,✉
PMCID: PMC13383021  PMID: 42189699

ABSTRACT

Polymer‐in‐sulfide composite electrolytes have emerged as highly promising candidates for all‐solid‐state lithium batteries (ASSLBs) due to their on‐demand shaping and rapid ion diffusivity. However, a striking paradox arises in the case of ethylene oxide‐tethered polyacrylates (EO‐PAs): their high polarity/strong nucleophilic tendencies constitute a major threat to sulfide stability yet exhibit anomalously high polymer/sulfide compatibility in practice. The underlying mechanism remains a matter of uncertainty. Herein, we first reveal a self‐limiting covalent ligation mechanism that accounts for this compatibility paradox. Central to this principle is the identification of intimate interactions between terminal −CH3 in EO‐PAs and PS4 3− units in Li6PS5Cl, not only suppressing parasitic nucleophilic reactions by EO ligands but also enhancing air stability. The self‐limiting interface was rigorously validated by density functional theory calculations, 31P solid‐state nuclear magnetic resonance, x‐ray computed tomography, and time of flight secondary ion mass spectrometry. The robust polymer‐in‐sulfide electrolyte achieves dendrite‐free Li plating/stripping for over 1200 h at 3 mA cm−2 and delivers approximately 100% capacity retention over 1000 cycles in NCM811‐based ASSLBs. These findings elucidate the core mechanism of interface regulation and provide pivotal guidance for the development of high‐performance ASSLBs.

Keywords: interface stability, polymers, self‐limiting covalent ligation, solid electrolytes, sulfides

The compatibility paradox between polar polymers and sulfide electrolytes is resolved by a self‐limiting covalent ligation mechanism. The C−S covalent interaction between the terminal –CH3 of EO‐PA and PS4 3– in sulfide electrolyte effectively suppresses the nucleophilic attack with enhanced interfacial compatibility.

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1. Introduction

Driven by the rapid expansion of electric vehicles, portable electronics, and grid‐scale energy storage [1], there is an ever‐growing demand for the development of safe and energy‐dense batteries [2, 3]. In this context, lithium (Li) metal anodes stand out as transformative candidates, offering a 40%–50% boost in energy densities via their unparalleled theoretical capacity (3860 mAh/g) and low equivalent potential (−3.04 V vs. standard hydrogen electrode), which far surpasses state‐of‐the‐art lithium‐ion batteries (LIBs) using graphite anodes [4, 5]. Concomitantly, realizing this potential hinges on breakthroughs in electrolyte chemistry to mitigate critical issues: flammability, uncontrolled Li dendrite growth, and thermal runaway. Solid electrolytes (SEs) have risen to prominence as a linchpin solution, leveraging intrinsic flame resistance while enabling high‐energy‐density electrode materials, including Li metal anodes and high‐voltage cathodes [6, 7]. Under these circumstances, thiophosphate‐based electrolytes, such as Li₆PS₅Cl (LPSC), Li₁₀GeP₂S₁₂ (LGPS), and Li₉.₅₄Si₁.₇₄P₁.₄₄S₁₁.₇Cl₀.₃ (LSiPSC), have garnered exceptional attention for high ionic conductivity (>10−2 S cm−1), mechanical deformability and compatibility with solvent‐free fabrication (e.g., dry‐powder pressing) [8]. Despite these advantages, interfacial instability remains the Achilles’ heel of sulfide SEs [9]. At the cathode, physical contact between sulfides and oxide‐based active materials incurs the penalty of deleterious elemental interdiffusion and crosstalk, fostering space charge layer formation, transition metal dissolution, and progressive interfacial degradation [10, 11, 12]. For instance, Uchimoto et al. [13] linked the formation of cobalt sulfide byproducts to mutual diffusion between LiCoO2 and 80Li2S·20P2S5 glass sulfides, further complicating interfacial dynamics. More problematic is the incompatibility at the Li anode interface, where spontaneous reactions between Li and SEs generate kinetically/thermodynamically unstable interphases. In LPSC, decomposition yields LixP and Li2S, poor ionic conductors (< 10−7 S cm−1) which impede ion transport [14], degrade cycling capacity and reduce Coulombic efficiency [15]. Wu et al. specifically demonstrated that these byproducts at the Li₆PS₅Cl/Li interface accelerate Li filament propagation, posing significant barriers to commercialization [16].

“Organic‐in‐sulfide” SEs architecture, integrating dry‐film‐processed LPSC membranes with interfacially infiltrated polymers, has emerged as a pivotal focus in sulfide‐based ASSLBs. This design synergizes superior interfacial compatibility and retains high ionic conductivity and processability (e.g., ease of lamination, Figure 1a). For sulfide matrices, dry‐film fabrication via mechanical calendering of PTFE binders and sulfide particles enables thin films with precisely controlled thickness [17], a critical prerequisite for high‐energy‐density batteries [18]. However, the screening of organic polymers for this architecture remains challenging, primarily due to the chemical vulnerability of sulfides to nucleophilic attacks [19]. High‐polarity polymers/solvents with functional groups (e.g., nitrogen or oxygen with lone‐pair electrons) readily induce sulfide degradation [20]. According to the hard‐soft acid‐base (HSAB) theory, these electrophilic species strongly interact with P5+ cations [21], triggering collapse of PS4 3− tetrahedra and chemical degradation—severely compromising ionic diffusivity and electrochemical performance [22]. A well‐documented example is the decomposition of LPSC upon contact with poly(ethylene oxide) (PEO) [15], whose polar EO backbone attacks PS4 3− tetrahedral anions, yielding byproducts like LiCl, Li3PS4, and Li2Sx. This phenomenon can also be observed in Li7P3S11 (LPS) and Li10GeP2S12 [23]. Ghorbanzad et al. [24] employed solid‐state NMR to characterize PEO‐LPSC/LPSCB interactions, identifying ether groups’ high Li+ affinity to form P(EO)3LiCl complexes and oxide PS4 3− into P2S7 4− dimers, impeding interfacial ion migration [25]. Such incompatibility extends to solvent selection during ball milling for sulfide particle size reduction, exacerbating processing challenges.

FIGURE 1.

FIGURE 1

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Structural characterization of organic‐in‐sulfide electrolyte: (a) Schematic illustration of interface advantages in organic‐in‐sulfide ASSLBs. (b, c) Morphology and structure of flexible organic‐in‐sulfide membranes integrated with the battery assembly process. (d) 3D reconstruction images of the original LPSC membrane with corresponding 2D slice images in the x–y and y–z planes. Gray and black represent the LPSC phase and pores, respectively. Cross‐sectional SEM images and EDS mapping of P and Cl elements in (e−k) the organic‐free sulfide electrolyte and (l−r) the organic‐in‐sulfide electrolyte, highlighting the porosity distinction.

Polyacrylates (PAs) have been spotlighted for stable integration with sulfide electrolytes featuring prevalent matrices of either alkyl [26] or polar ethylene oxide (EO) side chains [27, 28]. Among these, EO‐functionalized PAs (EO‐PAs) are hypothesized to enable Li+ conduction via thermal segmental motion of comb‐like EO‐Li+ complexes [29], while their reactive C=C double bonds support in situ polymerization, a widely adopted strategy for optimizing solid‐solid interfaces [30]. Nevertheless, a long‐standing contradiction persists: analogous to PEO, EO‐induced parasitic reactions pose a threat to sulfide stability, yet EO‐PAs practically exhibit high compatibility with sulfides. The mechanistic basis for the EO‐PA compatibility paradox remains a matter of debate, necessitating urgent and in‐depth investigation.

Herein, we elucidate the enhanced interfacial compatibility of ASSLBs based on thin EO‐PAs/thiophosphate organic‐in‐sulfide SEs, while resolving the compatibility paradox through in‐depth analysis of local interactions and PS4 3− electron distributions. Using poly(ethylene glycol) methyl ether methacrylate (PEGMA) as a proof‐of‐concept polymer system, we propose and validate a self‐limiting covalent ligation mechanism, referring to CH3−PS4 3− interactions that self‐limit the undesirable EO‐PS4 3− side reactions. The dominant CH3−PS4 3− interaction was comprehensively confirmed by density functional theory (DFT) calculations, x‐ray photoelectron spectroscopy (XPS), 31P solid‐state nuclear magnetic resonance (ssNMR), x‐ray computed tomography (X‐CT), and time of flight secondary ion mass spectrometry (TOF‐SIMS). Both the electron distribution and steric hindrance contribute to preventing nucleophilic attack on PS4 3− tetrahedra, thereby stabilizing the interface. The resulting organic‐in‐sulfide SEs exhibit synergistic properties: enhanced air stability, high ionic conductivity (1.14 mS cm−1) and high Li transference number (0.85), enabling ultralong cycling stability of Li||Li symmetric batteries (>1200 h at 3 mA cm−2). When integrated into practical ASSLBs with a LiNbO3‐coated LiNi0.8Co0.1Mn0.1O2 (NCM811@LNO) cathode, the system delivers ∼100% capacity retention over 1000 cycles at 1 C (1 C = 200 mAh g−1).

2. Results and Discussion

Figure 1a and Figure S1 illustrate the structural advantages and fabrication workflow for “organic‐in‐sulfide” membranes using a solvent‐free approach. Repeated calendering of the LPSC/PTFE mixture results in a flexible sulfide electrolyte membrane [31] with mechanical robustness and a thickness of 75 µm, which can be integrated into coin batteries (Figure 1b,c) [32]. The internal structure of “organic‐in‐sulfide” membranes was investigated by X‐CT, scanning electron microscopy (SEM), and energy‐dispersive spectroscopy (EDS). Layered scanning and 3D reconstruction of X‐CT (Figure 1d) confirm the presence of a uniformly distributed permeable pore network. Random longitudinal slicing yields x−y and y−z plane cross‐sections, where bright gray regions correspond to the LPSC phase and black dots represent pores (green frame). The pores and cracks are discretely distributed throught the LPSC membrane. SEM‐EDS characterization of the LPSC membrane (Figure 1e–k) reveals a homogeneous distribution of P and Cl elements but a massive number of voids, which serve as preferential nucleation sites for Li dendrite growth via either mechanical crack propagation or stress‐driven diffusion creep mechanisms [33]. By sharp contrast, cross‐sectional SEM of organic‐in‐sulfide electrolyte (Figure 1l–r) showcases a dense texture after PEGMA percolation with 10 µm thickness (Figure 1m), the structure of which is further confirmed by EDS mapping. According to the literature, percolation architecture renders additional Li+ migration pathways through the PEGMA‐LPSC interface, and therefore increases polymer‐sulfide compatibility and interfacial stability for fast ion diffusivity [34, 35].

Organic‐in‐sulfide composite electrolytes should fulfill the requirements of fast Li+ diffusivity and compatibility against sulfide decomposition. Due to the well‐known nucleophilic attack, two key descriptors of polymers, dipole moment (μ) and donor number (DN), are accordingly focused [36]. Strong dipole‐dipole interactions (high μ) increase salt dissociation and ionic hopping, while higher DN suggests high Lewis basicity and the tendency to donate electrons to destabilize sulfides by attacking Lewis acid sites, such as P5+ in LPSC [37]. For example, benzene exhibits negligible Lewis basicity (low DN, 0.1 kcal mol−1), thus high compatibility with LPSC but limited dissociation capability for solid‐state ion diffusivity [38, 39]. By systematic examination of “DN–μ” relationships (Figure 2a and Table S1), three distinct regions are categorized: (1) a strong polarity region (e.g., DMSO and NMP solvents) with both high DN and μ, leading to high σ Li + yet strong LPSC degradation; (2) a weak polarity region, falling within the “safe window,” with low DN and μ, offering a balanced yet limited performance; (3) a medium stability—high σ Li + region, where the PEGMA paradox is located: despite its striking polarity, PEGMA demonstrates high compatibility with LPSC, challenging conventional understandings and obscuring polymer‐sulfide interactions.

FIGURE 2.

FIGURE 2

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(a) Scrutinization of dipole moment and donor number across solvents and (b) calculated electrostatic potential maps and molecular polarity index of organic solvents and polymer precursors. (c) 1H NMR spectra and (d) 31P ssNMR spectra of PEGMA monomers, PEGMA polymers, and organic‐in‐sulfide electrolyte. (e) FTIR spectra and (f) S 2p XPS spectra of LPSC and organic‐in‐sulfide electrolyte.

To rationalize this paradox, we re‐scrutinize PEGMA by electrostatic potential (ESP) simulation (Figure 2b) to estimate electron delocalization and possible polymer‐sulfide interactions [40]. Molecular polarity index (MPI) was selected, which is determined by the gap between the most negative (ESPmin) and most positive (ESPmax) surface potentials in ESP [41]. Obviously, PEGMA exhibits the lowest ESPmin and relatively high ESPmax among medium‐polarity solvents/polymers, yielding the lowest MPI (10.333 kcal mol−1) as compared with DOL, DME, and DMC. More detailed observation reveals an interesting phenomenon: the electron‐deficient region localizes on the terminal α‐carbon (−CH3), suggesting the site with top priority of nucleophilic attack to PS4 3− subunits. Positively, S−CH3 interaction may self‐limit parasitic P−S bond breakage against sulfide decomposition due to the reaction priority and steric hindrance in PEGMA. To confirm these hypotheses, a series of experiments was employed, including 1H solid‐state NMR, 31P solid‐state NMR, FTIR and XPS. For 1H NMR analysis (Figure 2c), the peak located at 1.93 ppm corresponds to the methyl 1H, while peaks at 5.58 and 6.10 ppm correspond to 1H in −C=CH2 groups. Notably, no 1H signal from the C=C backbone is observed in PEGMA and organic‐in‐sulfide SEs, indicating the complete polymerization. To further investigate the coordination environment, 31P ssNMR resonances at 85 ppm (corresponding to PS4 3– tetrahedra) are observed in both LPSC and organic‐in‐sulfide samples [42], signifying high compatibility. More importantly, an additional peak at 127.9 ppm appears only in the organic‐in‐sulfide electrolyte [27], attributable to P−(SR)x species (where R represents the alkyl backbone) [26], directly confirming the presence of C−S interactions. FTIR spectra of PEGMA, LiTFSI, and organic‐in‐sulfide electrolyte are comparatively analyzed to further corroborate changes in chemical composition and structure (Figure 2e). Fingerprint signals at 1453 and 2866 cm−1 (−CH2 bending), 1717 cm−1 (−C=O bending), and 1227 cm−1 (C−O−C stretching) are observed in both PEGMA and the composite electrolyte [43, 44]. Similarly, an obvious absorption peak at 753 cm−1 for the C−S bond is detected [45], indicating the targeted chemical interactions between PEGMA and LPSC (Table S2). XPS analysis further substantiates the formation of C−S interactions. High‐resolution S 2p spectra of LPSC and organic‐in‐sulfide electrolyte (Figure 2f) reveal strong doublet peaks at 161.5 and 162.7 eV, assigned to the P−S bond, indicative of intact LPSC. For an organic‐in‐sulfide electrolyte, two sets of double peaks at 163.2–164.5 eV and 160.1–160.6 eV are attributed to the P−S−C bonds and “free” S2–, respectively. Concurrently, high‐resolution C 1s XPS spectra of organic‐in‐sulfide electrolyte (Figure S2) also verify the formation of C−S bonds, as evidenced by the  additional peak at 285.6 eV (Table S3–S4) [46].

Figure 3a,b illustrate the self‐limiting covalent ligation mechanism in an organic‐in‐sulfide electrolyte. The characteristics of electronic distribution were obtained via electron localization function (ELF) analysis (Figure 3c−f) [47]. Electrons around the S atom (ELF value: 0.8) are significantly enhanced in composite SEs as compared to LPSC (ELF value: ∼0.5) due to the formation of C−S interactions [48]. This localized electron density provides higher resistance to electron attack by ether‐oxygen groups, as it generally delocalizes S electron densities [49]. The formation of C–S covalent bonds between the terminal –CH3 of PEGMA and the S atoms of PS4 3– units, which plays a dual role in suppressing nucleophilic attack by EO groups. First, upon C–S bond formation, electron density around the coordinated S atom becomes partially localized, reducing the electrophilicity of the adjacent P5+ center and thereby diminishing its susceptibility to further nucleophilic attack. Simultaneously, the covalently tethered polymer chain physically blocks access to the PS4 3– tetrahedron, creating a protective barrier that prevents EO groups from approaching the reactive P–S moieties. Collectively, these effects establish a self‐limiting interface where the initial C–S interaction effectively neutralizes the inherent nucleophilic threat of the EO side chains. To validate the positive effect of compatibility on solid‐state ionic diffusion, Li||Li symmetric batteries were assembled, and the Li+ migration behavior was probed by 7Li ssNMR (Figure 3g). Pristine PEGMA and LPSC materials show 7Li chemical shifts at 1.47 ppm and ‐0.91 ppm, respectively. Crucially, organic‐in‐sulfide electrolyte exhibits an additional resonance peak at 0.25 ppm, attributed to the localized Li+ at the LPSC/PEGMA interface, as also demonstrated in previous studies [50]. The ionic conductivity of organic‐in‐sulfide electrolyte was evaluated via electrochemical impedance spectroscopy (EIS) tests (Figure 3h). A low resistance of 44.5 Ω is achieved at room temperature, corresponding to a high ionic conductivity of 1.14 mS cm−1 for organic‐in‐sulfide electrolyte.

FIGURE 3.

FIGURE 3

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(a−b) “Self‐limiting covalent ligation” mechanism in organic‐in‐sulfide electrolyte. (c−f) ELF analysis before and after PEGMA incorporation. (g) 7Li magic‐angle solid‐state NMR. (h−i) Ionic conductivity of organic‐in‐sulfide electrolyte. (j) Comparison of organic‐in‐sulfide performance in the literature.

Integrated with Li5.5PS4.5Cl1.5, the composite electrolyte delivers an even higher ionic conductivity to 4.9 mS cm−1 (Figure 3h,i and Figures S3–S5). To systematically evaluate the architecture advantages of organic‐in‐sulfide electrolytes, sulfide‐in‐organic electrolytes with varying LPSC contents are compared and optimized (Figures S6–S12), which exhibit very limited chemical stability and electrochemical performance.

To further demonstrate the enhanced air stability, x‐ray diffraction (XRD) analysis was applied to monitor phase transition during air exposure (Figure S13). For pristine LPSC after 20 min of air exposure, XRD spectra reveal P2O5, Li3PS4, LiOH, and Li2O byproducts, the intensity of which increases with prolonged exposure time. In turn, LPSC characteristic peaks (e.g., Li3PS4 phase at 2θ = 27.3°) diminish accordingly [51]. Moisture and oxygen penetration through surface defects into LPSC induces bulk decomposition by hydrolysis reactions with PS4 3− units. The generated LiOH further reacts with oxygen to form an insulating Li2O2/Li2CO3 layer, blocking Li+ transport pathways and exacerbating interfacial impedance [52]. Notably, the organic‐in‐sulfide electrolyte exhibits exceptional structural stability under ambient conditions (Figure S13), and no obvious structural changes are observed even after 60 min of air exposure. To further evaluate the air stability of the “organic‐in‐sulfide” electrolyte, the solid electrolytes were placed in a sealed container filled with humid air (25°C, 20% RH). The amount of H2S released over different storage periods was used to assess the extent of the hydrolysis reaction in the sulfide electrolytes. As shown in Figure S14, the “organic‐in‐sulfide” electrolyte exhibited excellent hydrolysis resistance in humid air, with an H2S release of only 0.50 cm3 g−1 after 30 min, whereas that of LPSC reached 23.08 cm3 g−1. Furthermore, the inset of Figure S14 shows that the “organic‐in‐sulfide” electrolyte maintained its structural integrity, suggesting that the improved air stability of the sulfide electrolyte may be related to its structure. The ionic conductivity of LPSC and the “organic‐in‐sulfide” electrolyte after exposure to humid air for 30 min is presented in Figure S15. It can be observed that the “organic‐in‐sulfide” electrolyte retained a high ionic conductivity of 0.91 × 10−4 S cm−1, whereas the ionic conductivity of LPSC dropped sharply from 1.95 × 10−3 S cm−1 to 0.32 × 10−4 S cm−1. This indicates that the PEGMA protective layer structure in the “Organic‐in‐Sulfide” electrolyte significantly impedes moisture penetration and substantially suppresses the hydrolysis reaction. In contrast, the pristine LPSC membrane possessed a loose structure containing numerous voids and grain boundaries, allowing water molecules to fully contact the PS4 3– species on the membrane surface, triggering a severe hydrolysis reaction.

To explore the contribution of compatibility to stabilizing the Li anode interface, Nyquist plots of Li|organic‐in‐sulfide|Li symmetric cells were evaluated. Critical current density (CCD), which reflects interfacial stability and Li dendrite inhibition, serves as a key metric for assessing the maximum current density of ASSLBs without short‐circuiting or failure [53, 54]. Pristine LPSC exhibits a CCD value of only 0.15 mA cm−2 (Figure 4a–c), and the limitation arises from side reactions at the LPSC interface and the presence of the binder PTFE. Volume shrinkage caused by electrochemical defluorination induces interfacial microcracking, further accelerating Li dendrite growth and ultimately resulting in short‐circuits [32]. By contrast, the organic‐in‐sulfide electrolyte enhances interfacial contact, optimizes the structure of SEI, and improves compatibility with Li metal. To verify these effects, the CCD of a Li|organic‐in‐sulfide|Li symmetric battery was tested, which achieves an exceptionally high value of 3.5 mA cm−2 (Figure 4d).

FIGURE 4.

FIGURE 4

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Anodic compatibility of Li|organic‐in‐sulfide|Li battery. CCDs of (a) Li|LPSC|Li, (b) Li|PEGMA|Li, (c) Li|sulfide‐in‐organic|Li, and (d) Li|organic‐in‐sulfide|Li symmetric batteries. (e) Comparison of CCDs with other organic‐in‐sulfide outcomes. (f) Cycling profiles of Li||Li symmetrical batteries with the organic‐in‐sulfide electrolyte at a current density of 3 mA cm−2 and a capacity of 3 mAh cm−2. (g) Comparison of symmetrical battery performance with other organic‐in‐sulfide outcomes [31, 51, 52, 55, 56, 57, 58, 59, 60].

Compared to other Li anode modification approaches reported in the literature, the anode fabricated by this method demonstrates superior interfacial stability (Figure 4e). Furthermore, the overpotential changes during long‐term Li+ plating/stripping cycling were analyzed using a Li||Li symmetric battery to assess the practical effectiveness of organic‐in‐sulfide electrolyte in suppressing Li dendrite growth. As shown in Figure S17, a Li|LPSC|Li symmetric battery exhibits a polarization voltage of 52.4 mV at a current density of 0.1 mA cm−2 with an irreversible short circuit observed after 25 h due to sparse and loose Li deposition. In contrast, the initial overpotential of a Li|organic‐in‐sulfide|Li symmetric battery is 16.71 mV and stable cycling exceeding 1700 h is achieved at the current density of 1 mA cm−2 (Figure S18). Long‐term Li plating/striping behavior at a current density of 3 mA cm−2 is also demonstrated in Figure 4f, where a stable overpotential is maintained for more than 1200 h, surpassing that of sulfide‐in‐organic structure (Figure S22). It can be concluded that an organic‐in‐sulfide composite electrolyte optimizes Li deposition/dissolution behavior and facilitates the formation of a reliable SEI [61]. Moreover, Nyquist plots and related fitting data of two batteries were evaluated. The Li|LPSC|Li battery initially exhibits an impedance of 4700 Ω (Figures S20 and S21), significantly higher than that of organic‐in‐sulfide system (181 Ω). After 1500 h, the interfacial impedance of a Li|organic‐in‐sulfide|Li symmetric battery increases to only 372 Ω, confirming that detrimental and persistent interfacial reactions are effectively suppressed due to the introduction of PEGMA [62].

Post‐mortem interface detection provides more essential evidence on the origin of enhanced interfacial stability in organic‐in‐sulfide SEs. SEM imaging of cycled Li anodes in bare PEGMA, sulfide‐in‐organic, and organic‐in‐sulfide SEs demonstrates distinct SEI morphologies (Figure 5). A rough and fluffy Li surface is displayed in bare PEGMA with a significant accumulation of Li dendrites and “dead Li,” leading to substantial volume expansion to 112 µm (Figure 5a). Although the cycled sulfide‐in‐organic sample presents a smooth SEI surface with flat and dense Li deposition, the thick SEI layer of approximately 70 µm indicates very limited interface ion diffusivity (Figure S22) [55].

FIGURE 5.

FIGURE 5

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Typical SEM images of cycled Li anodes in (a) Li|PEGMA|Li, (b) Li|sulfide‐in‐organic|Li, and (c) Li|organic‐in‐sulfide|Li batteries after 100 cycles. XPS spectra of (d) S 2p and (e) P 2p regions for Li anodes in cycled Li|LPSC|Li and Li|organic‐in‐sulfide|Li batteries. (f) TOF‐SIMS 2D spectra and (g) 3D reconstructions of representative SEI species formed on the surface. (h) TOF‐SIMS depth profiles of various secondary ion species for the organic‐in‐sulfide electrolyte

As a stark contrast, a more compact and uniform Li surface is observed with an organic‐in‐sulfide electrolyte (Figure 5c and Figure S22), forming a dense Li deposition layer with a thickness of only 45 µm, which corresponds to favorable Li+ plating/stripping behavior and the formation of stable SEI. XPS was carried out to further analyze the SEI composition according to different architectures. For pristine LPSC electrolyte, the presence of Li2S (160 eV) and Li3P (130.6 eV) are clearly observed in Figure 5d,e [63], the decomposition products via side reactions with Li metal [64]. In contrast, these signals disappear dramatically in organic‐in‐sulfide electrolytes. Through TOF‐SIMS, the compositional distribution and corresponding chemical characteristics of the cycled interface were further elucidated. TOF‐SIMS chemical imaging (Figure 5f and Figure S23) and three‐dimensional reconstruction (Figure 5g) reveal strong C2H3O− signals (from PEGMA) and LiS− signals (from LPSC bulk phase), confirming the presence of a uniform PEGMA‐penetrated surface layer [65]. Notably, when LiTFSI salt in PEGMA contacts with Li metal, it induces the formation of a LiF‐rich SEI [66]. As shown in Figure 5h, LiF and CH2O3 − species dominate the Li surface layer, indicating further suppression of Li dendrite growth [67]. Meanwhile, LiS−, Cl−, and P− signals are absent in the SEI layer, suggesting effective isolation of Li metal from the LPSC electrolyte, thereby inhibiting interfacial side reactions [68].

To further illustrate the superior stability and compatibility governed by the self‐limiting covalent ligation mechanism, Li||NCM811 cells with organic‐in‐sulfide SEs were assembled and tested. An initial discharge specific capacity of 130.1 mAh g−1 was delivered at 0.1 C (1 C = 200 mAh g−1) at 25°C (Figure 6a), higher than those of sulfide‐in‐organic electrolyte (116.5 mAh g−1, Figures S24 and S25) and LPSC electrolyte (76.2 mAh g−1). A remarkable capacity of 106.2 mAh g−1 with a capacity retention of 81.5% was maintained after 100 cycles. In sharp contrast, a rapid capacity decay to 18.4 mAh g−1 after 30 cycles was observed in the bare LPSC, a fact of structural degradation and the high‐impedance passivation interface layer caused by Li contacts. Typical charge‐discharge curves are shown in Figure 6b, indicating smaller voltage hysteresis and a more stable interphase. The rate performance of Li||LCO batteries is presented in Figure 6c, demonstrating the discharge capacities of 130.6, 102.8, 96.0, 86.8, 77.4, and 73.0 mAh g−1 at 0.1, 0.2, 0.3, 0.5, 0.8, and 1 C. After 100 cycles, Li||LCO batteries exhibit high capacity retention of 91.4% at 0.1 C (1 C = 150 mAh g−1) (Figure 6d), significantly surpassing that of PEGMA and Sulfide‐in‐Organic electrolyte (Figures S26 and S27). The impedance spectra after 50 cycles (Figure S28) indicate stable charge exchange and favorable electrochemical performance. Additionally, organic‐in‐sulfide electrolytes exhibit reduced hysteresis and thus higher average voltages (Figure 6e). More importantly, organic‐in‐sulfide SEs empower outstanding cycling stability in Li||NCM811, delivering ∼100% capacity retention after 1000 cycles at 1C. Comparative analysis regarding discharge capacity and capacity retention across different SEs is presented in Figure 6g [55, 59, 69, 70, 71, 72, 73, 74, 75, 76, 77, 78, 79, 80, 81], demonstrating a significant electrochemical breakthrough in ASSLBs with self‐limiting organic‐in‐sulfide SEs.

FIGURE 6.

FIGURE 6

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Electrochemical performance of Li||NCM811@LNO and Li||LCO@LNO batteries at room temperature. (a) Galvanostatic cycling of Li|organic‐in‐sulfide|NCM811@LNO and Li|LPSC|NCM811@LNO coin‐type batteries at 0.1 C (1C = 200 mAh g−1). (b) Charge/discharge profiles of Li|organic‐in‐sulfide|NCM811@LNO battery at 0.1 C within the voltage of 2.5−4.3 V. (c) Rate performances of Li|organic‐in‐sulfide|LCO@LNO coin‐type batteries. (d) Galvanostatic cycling of Li|organic‐in‐sulfide|LCO@LNO and Li|LPSC|LCO@LNO batteries at 0.1 C (1 C = 150 mAh g−1). (e) The electrochemical behavior of ASSLBs with different solid electrolytes and electrodes has been reported [55, 59, 69, 70, 71, 72, 73, 74, 75, 76, 77, 78, 79, 80, 81]. (f) Galvanostatic cycling of Li|Organic‐in‐Sulfide|NCM811 ASSLMB at 1 C (1 C = 200 mAh g−1) (g) Comparison of specific capacity of ASSLB at different charge‐discharge rates. The size of the bubble area intuitively reflects comprehensive performance at a specific magnification [55, 59, 69, 70, 71, 72, 73, 74, 75, 76, 77, 78, 79, 80, 81].

3. Conclusion

This work presents a self‐limiting covalent ligation mechanism to resolve the long‐standing compatibility paradox between polar PEGMA polymers and sulfide electrolytes. By rationally identifying the terminal‐induced strong interactions, conventional nucleophilic attack by EO groups is self‐limited with enhanced polymer/sulfide compatibility. The unique electrolyte architecture optimizes Li+ transport and electronic insulation, achieving superior interfacial compatibility while suppressing the decomposition of the sulfide electrolyte under humid conditions (25°C, 20% RH). In addition, the “Organic‐in‐Sulfide” electrolyte generates organic‐inorganic hybrid SEI during cycling, which inhibits the nucleation and growth of Li crystals. The resulting organic‐in‐sulfide architecture combines high ionic conductivity (1.14 mS cm−1), exceptional interfacial stability (CCD: 3 mA cm−2) and ambient air stability. ASSLMBs utilizing the “ Organic‐in‐Sulfide” electrolyte demonstrate a notable capacity of 127.9 mAh g−1 after 1000 cycles at 1 C, with an impressive retention rate of 100%. This study provides insights for molecular‐level interface engineering in next‐generation solid‐state batteries, laying the foundation for the development of high‐energy density electric vehicle batteries in the future.

Author Contributions

Yan Zhang: investigation, formal analysis, visualization, writing – original draft, data curation. Yuxi Zhong: investigation, validation, formal analysis, data curation. Ruiqi Guo: methodology, validation, formal analysis. Songjie Li: software, data curation, formal analysis. Kaidi Kang: methodology, validation. Bingguang Ye: data curation, validation. Wei Shan: software, formal analysis. Ran Zhao: methodology, formal analysis. Xinran Wang: conceptualization, formal analysis, supervision, funding acquisition, writing – review and editing. Chuan Wu: conceptualization, formal analysis, supervision, funding acquisition, resources. Ying Bai: methodology, supervision, project administration, resources, funding acquisition, conceptualization.

Conflicts of Interest

The authors declare no conflicts of interest.

Supporting information

Supporting File: anie72869‐sup‐0001‐SuppMat.docx.

The authors have cited additional references within the Supporting Information.

ANIE-65-e9377760-s001.docx (19.6MB, docx)

Acknowledgments

This work was supported by the National Program of China, the National Key Research and Development Program of China (2022YFB2404400), the National Natural Science Foundation of China (22379016), the Beijing Natural Science Foundation (L247017), and the Fundamental Research Funds for the Central Universities (2024CX06061).

Contributor Information

Xinran Wang, Email: wangxinran@bit.edu.cn.

Chuan Wu, Email: chuanwu@bit.edu.cn.

Ying Bai, Email: membrane@bit.edu.cn.

Data Availability Statement

The data that support the findings of this study are available from the corresponding author upon reasonable request.

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Associated Data

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Supplementary Materials

Supporting File: anie72869‐sup‐0001‐SuppMat.docx.

The authors have cited additional references within the Supporting Information.

ANIE-65-e9377760-s001.docx (19.6MB, docx)

Data Availability Statement

The data that support the findings of this study are available from the corresponding author upon reasonable request.

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