Electrolyte chemistry of adaptive hydrogen bonded domains for high voltage lithium metal batteries

Domain: neuroscience
Published: 2026-02-05
Editorial outcome: major_revision

Abstract
The practical implementation of lithium metal batteries is hindered by unstable electrode-electrolyte interfaces and sluggish ion transport kinetics. Here, we report a molecular design strategy that remodels electrolyte solvation structures via the formation of hydrogen-bonded domains, thereby enhancing both the thermodynamics and interfacial dynamics of Li+ transport. Specifically, we introduce 2-cyano-N-methylacetamide, an electrochemically stable hydrogen bond donor, as a cosolvent to construct stable nanoscale hydrogen-bonded domains (<3.5 Å). 2-Cyano-N-methylacetamide generates both classical (H-bond, Hδ⁺–Oδ⁻) and nonclassical (Z-bond, Nδ⁻–Hδ⁺) hydrogen bonding, which disrupts loosely bound solvated clusters and induces tightly coordinated Li+ solvation structures. The hydrogen-bonded domains facilitate the formation of oriented fast Li+ transport channels. Accordingly, in Li||LiNi0.8Co0.1Mn0.1O2 cells cycled under demanding conditions of 4.7 V with a high areal capacity of ~3.0 mAh cm−2, the electrolyte enables a capacity retention of 78.8% after 400 cycles. In addition, a stable 4.7 V lithium metal pouch cell is demonstrated with a specific energy (based on the mass of all components) of 418.2 Wh kg−1. This work offers a useful electrolyte design principle on solvation chemistry and interfacial engineering for high-voltage lithium metal batteries.

Introduction
Increasing the charging cut-off voltage to 4.6 V or higher has become a core research focus in pursuit of lithium metal batteries (LMBs) with high specific energy exceeding 400 Wh kg−1, a key milestone towards next-generation high-performance energy storage technologies1–4. However, the strong affinity between Li+ and the solvent in the constrained battery system inhibits the Li+ desolvation process, and the disordered mass transfer and unstable interfacial reaction often lead to the degradation of the battery performance5–7. Eminently, the electrolyte is affected by the magnitude and spacing of the inter-substance forces8–10, which plays a crucial role in the battery’s operational process. In response to these challenges, considerable efforts have been made to regulate reaction kinetics by modulating the competitive Li-ion coordination environment among electrolyte components such as anions, solvent molecules, and additives11,12. For example, high-concentration electrolytes (HCEs) (Fig. 1b) have been employed to replace conventional dilute electrolytes (Fig. 1a), where enhanced ion-ion interactions promote the formation of contact ion pairs (CIPs) and aggregates (AGGs)13–15. Notably, although the large-sized solvation domains formed by abundant AGGs and CIPs are favorable for anion reduction and decomposition, their stable structure restricts Li+ movement, increasing the tortuosity of ion migration pathways and reducing the bulk phase diffusion coefficient16. In contrast, small-sized solvated domains provide continuous ion channels through dynamic ligand bond reorganization, thereby facilitating rapid Li+ transport17. Several studies in recent years have shown that localized high-concentration electrolytes (LHCEs) modulate solvation domains by regulating dipole-dipole interactions between diluents and solvents (Fig. 1c), which in turn form stable interfaces18–20. However, the disordered dipole-dipole interactions introduced by the diluents tend to disrupt the permeable three-dimensional solvation network, thereby preventing Li+ hopping21–23. Therefore, it is important to tune electrolyte molecule/ion interactions and regulate their spatial effects on solvation sheaths for realizing stable electrolytes with efficient ion-transport properties24.

Existing studies demonstrated that the interactions are usually dominated around smaller dipole-dipole forces (Fig. 1e, 0.1~5 kJ mol−1) and van der Waals forces (0.4~4 kJ mol−1)25–27. The key mechanism of how larger ion-dipole interactions (ligand bonding, Fig. 1f, 50~200 kJ mol−1) are modulated by smaller dipole-dipole forces remains unclear. Considering the enigmatic nature of H-bond interactions in electrolytes, whether relatively strong hydrogen-bond interactions (Figs. 1g, 5~100 kJ mol−1) can better regulate the strength and topology of Li+-solvents has piqued our keen interest28,29. According to molecular orbital (MO) theory, H-bond donor-acceptor interactions due to charge transfer from the Highest Occupied Molecular Orbital (HOMO) of X to the empty antibonding orbitals of σ* Y-H bonds enhance the stability of the (Xδ−–Hδ+···Yδ−) cluster28,29. With the expansion of IUPAC’s definition of the H-bond, Xδ− does not need to be very negatively charged, but only requires X–H to be at least slightly polar30. This requirement includes bonds of C–H, P–H, etc. Antipolar X–H groups can form non-classical hydrogen bonds (Z-bonds), such as C–H···X and C–H···O31. The presence of specific co-solvent molecules has been reported in literature to reshape the interfacial molecular interactions by forming Z-bonds with the solvent molecules, which greatly facilitates Li+ desolvation dynamics32,33. Despite these advances, the potential mechanisms of how to rationally regulate the co-solvent-solvent H-bond and Z-bond interactions on the behavior of the electrode-electrolyte interface are still poorly understood.

To fill this gap in understanding, we developed a strategy for high-voltage LMBs to utilize a donor-containing co-solvent to form H-bond (Hδ⁺–Oδ⁻) and Z-bond (Nδ⁻–Hδ⁺)-coupled hydrogen-bonded domains to remodel the solvation structure. The electronegativity of the ester solvent oxygen atom is weakened by H-bond and Z-bond, thus weakening the Li+-solvent force and encouraging more anions to enter the Li+ inner solvation shell (Fig. 1d), based on which inorganic solid-electrolyte interphase (SEI) and cathode-electrolyte interphase (CEI) dominated by the anionic decomposition products are preferentially formed. In addition, the construction of hydrogen-bonded domains exhibits a dynamic Li+-solvent solvation/desolvation process, significantly enhancing the Li+ transport kinetics and the interfacial desolvation kinetics. With such an electrolyte design, the Li||LiNi0.8Co0.1Mn0.1O2 (NCM811) coin cell data high-voltage of 4.7 V and 2.98 mAh cm–2 retained 78.8% capacity after 400 cycles. Furthermore, the 1.6 Ah Li||NCM811 cell achieved an specific energy of 418.2 Wh kg–1.

Results
Adaptive hydrogen-bonded domains electrolyte design strategy

In screening electrolyte co-solvents, multiple metrics were evaluated to ensure an optimal broad electrochemical window. Density functional theory (DFT) calculations of the HOMO and Lowest Unoccupied Molecular Orbital (LUMO) energies served as preliminary indicators of the redox properties of short-chain amides (Fig. 1h and Figs. S1, 2). Most of the coordination solvents in conventional solvents such as ethylene carbonate (EC), dimethyl carbonate (DMC), 1,3-dioxolane (DOL), dimethylsulfoxide (DMSO), and vinyl sulfate (DTD) are prone to oxidation due to their high HOMO energy levels. Substituting the hydrogen atom in the amide’s –NH2 group with a methyl group (e.g., FM, Ac, NMAc, and NNDMAc, the full chemical names of these compounds are provided in Fig S1) enhanced oxidation resistance34. Further modifications, such as replacing the –NH2 group with isopropyl (–CH3), trifluoromethyl (–CF3), or cyanoethyl (–CH2CN) groups (e.g., TFAc, ANAc, NNDMTFAc, ANNNAc, NMTF, and ANM, the full chemical names of these compounds are provided in Fig. S1), improved compatibility with lithium metal electrodes by increasing polarity35. Moreover, molecules bearing a −CH2CN group exhibited long-term chemical stability toward both the electrolyte and lithium metal, and no appreciable gas evolution or discoloration was observed (Fig. S3).

An initial screening of the magnitude of solvent-Li+ forces based on the molecular electrostatic potential (ESP) showed that solvents with large ESP minima (>–30 kcal mol–1), such as 1,1,2,2-tetrafluoroethyl-2,2,3,3-tetrafluoropropylether (HFE), were weakly coordinated to Li+, which weakened the electrolyte bulk phase mass transfer (Fig. 1i). In turn, solvents with small ESP minima (>–40 kcal mol–1) are not conducive to the formation of solvated structures of CIPs and AGGs. Notably, N-Methyltrifluoroacetamide (NMTF, −35.49 kcal mol–1) and 2-cyano-N-methylacetamide (ANM, −31.92 kcal mol–1) exhibited lower ESP minima compared to conventional electrolyte solvents. ANM, in particular, demonstrates a higher ESP maximum (55.30 kcal mol–1) and increased H-bond donor capacity. Reduced density gradient (RDG) analysis revealed significant weak H-bond interactions between amide clusters (N–H) and solvent molecules (O), confirming the existence of H-bond (Fig. 1j-l and Figs. S4, 5)36. Fourier transform infrared (FTIR) spectroscopy further supported these findings, showing that the –N–H– stretching frequency of pure NMTF shifts from 3342.0 cm–1 to 3334.2 cm–1, and pure ANM shifts from 3353.1 cm–1 to 3340.4 cm–1 upon addition of ethyl methyl carbonate (EMC) and fluoroethylene carbonate (FEC), indicating H-bond formation (Fig S6a, b).

The ANM exhibited a more significant blue shift than NMTF, indicating stronger H-bond. The specific atoms involved in these interactions and the strength of the H-bond were elucidated through electron density difference diagrams, revealing that the H-bond formed by ANM with EMC and FEC was considerably stronger than those formed by NMTF (Fig. S7 and Supplementary Data 1–6). Notably, the –C≡N (N) group in ANM demonstrated weaker interactions with the –CH2 (H) group in EMC, which could be attributed to forming the Z-bond. The Z-bond interaction further strengthened the H-bond between ANM and EMC, resulting in a more robust H-bond (−10.8 kcal mol–1) compared to that of NMTF (−8.8 kcal mol–1) (Fig. 1j–l). The 1H nuclear magnetic resonance (1H NMR) spectra of the electrolytes confirmed the formation of H-bonds between NMTF/ANM and the ester solvents (Fig. S8a). In the 1H NMR spectra, the formation of H-bonds reduced the shielding of the electron cloud around the hydrogen nucleus, thereby shifting the resonance absorption toward the downfield region. Notably, the −CH2− group of EMC exhibited an upfield shift in the E-F@NMTF system, whereas it shifted downfield in the E-F@ANM system. This difference arises because ANM can form Z-bond interactions with EMC, while NMTF cannot (Fig. S8b). DFT calculations confirmed the formation of Z-bond (−2.5 kcal mol–1) (Table S1). In addition, the H-bond between NMTF/ANM and the (C–F) sites in FEC were comparatively weak, owing to the greater steric hindrance of (C–F) and its lower electron density relative to (C=O).

Furthermore, quantum theory of atoms in molecules (QTAIM) calculations were employed to elucidate the detailed characteristics of the hydrogen bonds. Higher electron density (ρ) and a more negative potential energy density (V(r)), generally indicate stronger H-bonds37. In the QTAIM results, all clusters exhibited low ρ and a positive Laplacian (▽2ρ), collectively indicating weak H-bond interactions (Table S2). The ANM and EMC (N–H···O=C) interaction showed the highest ρ, the most negative V(r), and the shortest bond length, indicating that the H-bond between ANM and EMC was the strongest. In addition, relative to the NMTF and FEC cluster, the ANM and FEC cluster displayed weaker interactions. Moreover, a Z-bond (C≡N···H–C) was identified in the ANM and EMC cluster, corroborating that Z-bond coupling in the ANM and EMC system further strengthened the H-bond interaction.

The influence of H-bond on the electronic structure was also demonstrated by the reduction in the HOMO energy levels for all cluster when interacting with EMC and FEC clusters (Fig. S9). The strongest H-bonds were formed by ANM and NMTF, leading to the lowest HOMO energy levels for their respective clusters. This outcome was particularly beneficial for enhancing the antioxidant properties of the electrolyte.

To improve the conductivity of the electrolyte, a 1.2 M solution of lithium bis(fluorosulfonyl)imide (LiFSI) in an FEC and EMC mixture in a 3:7 volume ratio (denoted as E-F) was adopted (Fig. S10a–c and Table S3). Addition of NMTF and ANM (denoted as E-F@NMTF and E-F@ANM) at a concentration of 5 M to E-F resulted in an enhancement of electrical conductivity. Notably, the conductivity of the electrolyte exhibits a non-monotonic dependence on the concentration of ANM. Although an increase in co-solvent content typically leads to higher viscosity, the introduction of 5 M ANM results in the highest ionic conductivity among all tested concentrations. This is attributed to the formation of the hydrogen-bonded domains, which enhance Li+ transport kinetics despite the increased viscosity of the electrolyte. The electrochemical performance of the Li||NCM811 coin cell with various amides added to E-F at 5 M was compared in the electrochemical window of 2.8~4.5 V (Fig. S11). As expected, E-F@ANM demonstrates optimal performance. Similarly, such H-bond interactions were operative across different carbonate solvents. In the FTIR spectra of the EC-EMC-NMTF and EC-EMC-ANM mixed solvents, the N–H stretching vibration showed a pronounced blue shift, unambiguously confirming H-bond (Fig. S12a, b). Moreover, deconvolution of the PF6– Raman bands indicated that introducing NMTF or ANM decreased the fraction of solvent-separated ion pairs (SSIP) while increasing the fractions of CIP and AGG, suggesting that H-bond promoted tighter Li+–anion association (Fig. S13a–c). In Li||NCM811 cells employing the E–E@ANM electrolyte, the capacity was retained at 93.3% after 200 cycles, significantly outperforming cells using the E-E electrolyte (Fig. S14).

Electrolyte solvation structures and de-coordination mechanisms

To elucidate the specific role of NMTF/ANM co-solvents in solvent-driven structural remodeling, the solvation structures of the electrolytes were analyzed using molecular dynamics (MD) simulations (Fig. S15a–c). Anion-enriched solvates were observed in E-F@NMTF and E-F@ANM compared to E-F (Fig. S16). In E-F system, the coordination numbers of oxygen in FSI–, FEC, and EMC were 1.78, 0.21, and 3.6, respectively (Fig. S17a and Supplementary Data 7). In contrast, these values changed to 3.17, 0.18, and 1.51 with NMTF (Fig. S17b and Supplementary Data 8) and to 3.24, 0.26, and 1.46 with ANM (Fig. S17c and Supplementary Data 9), respectively. These changes indicated that NMTF and ANM had minimal involvement in the solvation structure of Li⁺, being primarily distributed around the outer region of the solvation sheath. Because the H-bond interactions between NMTF/ANM and EMC/FEC differ in strength, the introduction of NMTF reduces the proportion of both EMC and FEC molecules participating in the solvation sheath. In contrast, ANM molecules selectively decrease the number of EMC molecules while increasing the proportion of FEC molecules in the solvation sheath (Fig. 2a). Notably, E-F@ANM system further amplified this effect by coupling H-bond and Z-bond interactions, which increases the coordination of more anions with Li+. Raman spectroscopy was used to verify the differences in solvated structures. The vibrational peak of the Li+–FSI− (720~740 cm−1) was red-shifted upon adding NMTF/ANM, indicating an increase in aggregation within the electrolyte (Fig. S18a). Concurrently, the addition of NMTF/ANM was accompanied by a simultaneous blue shift of EMC molecules’ vibrational peaks (920~940 cm−1), indicating a decrease in the coordinated EMC molecules. In contrast to NMTF, the vibrational peaks of FEC molecules (860~880 cm−1) were red-shifted after the addition of ANM co-solvent, indicating an increase in the coordination of Li with FEC (Fig. S18b). Furthermore, a statistical analysis was conducted on the H-bond distributions between NMTF, ANM and EMC or FEC in the two electrolytes (Fig. 2b, c and Fig. S19a, b). The results demonstrated that the number of H-bond per compartment is significantly higher in the E-F@ANM electrolytes compared to the E-F@NMTF (Fig. S20a–d), with a distribution of N–H∙∙∙O=C bond lengths between 2.5 Å and 3.0 Å, and bond angles between 120° and 180°. The shorter the bond angle and the shorter the bond length were considered the main characteristics of strong H-bonds, this is consistent with previous analysis that E-F@ANM had stronger H-bonds38. In addition, we further quantitatively compared the change in binding energy (Eb) of the solvated structure before and after the NMTF/ANM molecule achieves the drag effect using DFT calculations (Fig. 2d). The Li+ dissolution/dissolution energies of EMC, FEC, NMTF and ANM were calculated and compared in Fig. 2e and Fig. S21. Among them, EMC had the highest dissociation energy of −25.17 kcal mol−1, followed by FEC (−20.24 kcal mol−1), and the decrease in Eb allowed the coordination reaction to take place thermodynamically. Combining the above results, we showed that ANM could form localized hydrogen-bonded domains within a distance range of less than 3.5 Å from the carbonate molecule, and the formation of hydrogen-bonded domains stabilized the solvent molecule and forms a solvation structure with an anionic inner layer and an outer layer of hydrogen-bonded domains. The formation was confirmed by quantitative analysis of the simulated percentage of MD hydrogen-bonded domains in Fig. S22a, b and peak area fitting calculations of Raman spectra in Fig. S23a–d17. Furthermore, the solvated cluster size of 8.01 Å for the E-F system (Fig. S24a) was larger than that of 7.07 Å for the E-F@ANM system (Fig. S24b). In the E-F system, Li+ forms a large number of SSIPs (Fig. 2f), and the excess solvent ratio in solvation led to parasitic side effects of the ester molecules involved in the desolvation process, causing irreversible damage at the electrolyte-electrode interface. Additionally, the smaller size of Li+ clusters could promote faster Li+ migration. Unlike LHCEs where the fence effect diluent interrupted the Li+ transport hopping mechanism, the unique hydrogen-bonded domains in E-F@ANM were ordered outside the solvated boot layer, resulting in a more relaxed Li+ fast channel, which reduced the mass-transport energy barrier in the native phase of the electrolyte (Fig. 2g and Fig. S25)39. In addition, the shortening of the Li+-Li+ mutual distance also suggested that the construction of adaptive hydrogen-bonded domains forms tighter solvated clusters for Li+ structural transport (Fig. S26)40.

Hydrogen-bonded domains anchored solvent molecules to remodel the solvation structure of electrolytes toward enhancing their antioxidant properties

To evaluate the effect of adaptive hydrogen-bonded domains on the antioxidant properties of electrolytes, the electrochemical stability window (ESW) of various electrolytes was assessed using linear scanning voltammetry (LSV) (Fig. 3a). When constrained by an oxidation current threshold of 5 µA, electrolytes E-F@NMTF and E-F@ANM exhibited oxidation potentials of 4.58 V and 5.15 V, respectively. In contrast, the current in the E-F electrolyte sharply increased around 4.0 V. To further investigate electrolyte interactions with NCM811 at high voltages, leakage current tests were conducted at 4.6 V, 4.7 V and 4.8 V (Fig. 3b). At a constant voltage of 4.8 V, cells utilizing the E-F electrolyte displayed higher leakage currents. Conversely, electrolytes E-F@NMTF and E-F@ANM maintained lower and stable leakage currents. NMTF/ANM co-solvents modulate FSI− and Li+ interactions, as evidenced by 7Li NMR spectroscopy (Fig. 3c). Additionally, previous research indicated that anion-enriched solvation sheaths can diminish the prevalence of solvent-anion clusters, thereby reducing the negative charge transfer of anions to the solvent and effectively widening the ESW41. The incorporation of NMTF/ANM co-solvents into the E-F electrolyte resulted in a noticeable upfield shift in the 7Li NMR chemical shift, moving from −1.07 ppm in the E-F electrolyte to −0.61 ppm and −0.52 ppm in the E-F@NMTF/ANM electrolyte, respectively. This low-field shift suggested a reduction in the shielding of Li+ by EMC and FEC molecules, attributed to the H-bond interactions between NMTF/ANM and the carbonyl groups of these solvent molecules.

Recognizing that ion and solvent distributions are voltage-dependent, in situ FTIR spectroscopy was employed to monitor the dynamic evolution of the electrode-electrolyte interface under varying voltages during cell charging (Fig. S27). The peaks observed around 1090~1100 cm−1 and 1140~1160 cm−1 correspond to the vibrations of free EMC and FSI− coordinated with Li+, respectively. In the E-F electrolyte, as the charging voltage increased from 3.5 V to 4.8 V, free EMC molecules began to decompose around 4.1 V, followed by FSI− decomposition starting at ~4.3 V. Subsequently, the intensities of these signals decreased rapidly, indicating significant decomposition of the E-F electrolyte (Fig. 3d). Notably, the oxidative stability of the E-F@ANM electrolyte was enhanced, as evidenced by fewer free EMC molecules (Fig. 3e and Fig. S28a) and delayed FSI− decomposition with the increase of voltage, which commenced at around 4.6 V (Fig. S28b). To visualize the evolution of free ANM, the hydrogen-bonded domains, and H-bond strength during charge–discharge, we analyzed how the full width at half maximum (FWHM) and integrated intensity of the non H-bond and H-bond N–H peaks in in situ infrared spectra varied with the applied potential (Fig. S29a, b)42. First, the free N–H peak was fitted with a single Gaussian function, whose center frequency and FWHM were determined from the FTIR spectrum of ANM. The area of the H-bonded component extracted from the measured spectrum—outside the Gaussian-defined free N–H component—quantitatively characterized the extent of the hydrogen-bonded domains. A larger FWHM together with a larger H-bond peak area indicated a stronger H-bond, and vice versa43. At open-circuit voltage (OCV), H-bond interactions were relatively weak, as reflected by the small FWHM and peak area of the H-bonded feature. Upon applying a potential, both the FWHM and the H-bonded peak area increased, indicating rapid formation of an interfacial hydrogen-bonded domain; this trend persisted up to 4.6 V. Above 4.6 V, both metrics decreased, attributable to partial solvent decomposition that weakened the H-bond. In the early stage of charging, anion decomposition reduced the interfacial negative charge; under these conditions, insufficiently desolvated anions from the bulk were less able to reach the interface, while the high potential triggered partial decomposition of EMC and FEC. This behavior persisted up to 4.8 V. Collectively, these observations confirmed the evolution of the hydrogen-bonded domains during cycling, and their distinctive solvent-dragging effect that promoted preferential anion decomposition at the interface.

This improvement is attributed to the formation of adaptive hydrogen-bonded domains with lowered HOMO energy levels by the ANM and solvent molecules. The presence of hydrogen-bonded domains reduced the contact of free anions with the solvent, allowing for a significant broadening of the ESW of the electrolyte. The change in ESP around the carbonyl oxygen atom of the ester molecule in the hydrogen-bonded domains (Fig. 3f) showed an increase in the minimum ESP value, indicating that the generated clusters possess a stronger antioxidant capacity (Fig. S30). Furthermore, DFT calculations of oxidation potentials between various clusters showed that FSI⁻ and EMC-clusters have the lowest oxidation potential. Compared with NMTF, ANM created more stable hydrogen-bonded domains with FEC and EMC molecules, thus slowing down their decomposition (Fig. 3g). Notably, the higher oxidation potentials observed for Li+ and FSI− clusters further corroborate that the solvation of the anionic bootstrap layer contributed to the improved antioxidant properties of the electrolyte (Fig. 3h).

Overall, the adaptive hydrogen-bonded domains enhance the oxidative stability of the electrolyte by orienting and anchoring the solvent molecules, thereby increasing the anion-enriched solvation structure. This effectively addresses the oxidation of the solvent after desolvation at the high-voltage positive electrode, slowing down the flow of electrons from the HOMO of the electrolyte to the positive electrode (Fig. 3i), which causes electrolyte degradation and capacity loss.

Electrolyte compatibility with lithium metal

Notably, the formation of adaptive hydrogen-bonded domains improved their wettability with lithium metal, facilitating mass transfer (Fig. S31). Importantly, the E-F system forms an unstable SEI leading to excessive dendritic growth and interfacial side reactions, whereas the E-F@ANM promoted uniform Li+ fluxes as well as homogeneous Li deposition (Fig. 4a). In situ Raman spectroscopy revealed the dynamic process of electrolyte desolvation on lithium metal surfaces (Fig. S32). In the E-F electrolyte, Raman peaks observed at 890~900 cm−1 and 930~940 cm−1 correspond to the vibrations of Li+ coordinated with FEC and EMC, respectively (Fig. 4b). The E-F@ANM electrolyte exhibited reduced EMC coordination and increased FEC coordination, which was consistent with prior findings. During discharge, the E-F electrolyte experienced significant decomposition of EMC and FEC peaks, whereas the E-F@ANM electrolyte maintains relatively stable peaks, suggesting that the solvent reduction was mitigated by the pulling effect of ANM on the ester solvent (Fig. 4c). Furthermore, the Raman peaks of FSI− (~730 cm−1) in the E-F electrolyte red-shifted as the discharge proceeded (Fig. S33a), which was attributed to the fact that the reduction of the E-F electrolyte was mainly solvent-dominated and the decreasing solvent led to the enhancement of Li+ and FSI− coordination at the interface. The E-F@ANM electrolyte effectively maintained a stable anion-enriched solvated structure, promoting the formation of an inorganic SEI primarily composed of decomposed FEC and FSI− coordinated with Li+ (Fig. S33b).

This observation was supported by depth-profiled X-ray photoelectron spectroscopy (XPS) analyzes via sputtering conducted after 50 cycles of Li||Li cells using various electrolytes. In the C 1s spectra, signals corresponding to C–C/C–H, COOR, C–O, and Li2CO3 are observed, with C–O and COOR peaks attributed to EMC and FEC decomposition products (Fig. S34a–c). Notably, the reduced intensity of C–O and COOR peaks in the E-F@NMTF and E-F@ANM electrolytes suggested decreased solvent reduction-generated organic SEI. Conversely, the O 1s spectra exhibit enhanced Li2O signals in both E-F@NMTF and E-F@ANM electrolytes, indicating increased inorganic SEI formation (Fig. S34d–f). Additionally, the F 1s and N 1s spectra display higher Li3N and LiF signals in Li metal electrodes cycled with the E-F@NMTF and E-F@ANM electrolytes compared to those cycled with the E-F electrolyte (Fig. 4d–g), aligning with trends observed in FSI− decomposition (Fig. S35a, b). The ratio of clathrate-related signals after sputtering demonstrates that the SEI formed by the E-F system was predominantly organic. In contrast, the SEI formed by the E-F@NMTF and E-F@ANM electrolytes was predominantly inorganic (Fig. S36a–c). Together, these findings suggested that adaptive hydrogen-bonded domains induce the formation of more stable and robust SEI, which could be beneficial for improving the overall performance and stability of the LMBs.

To investigate the effect of adaptive hydrogen-bonded domains on the thermodynamics of mass transport in bulk-phase electrolytes, the Li+ transference number (tLi+) was measured to assess Li+ migration within the electrolyte (Fig. S37). The tLi+ values were 0.2 for the E-F, 0.47 for the E-F@NMTF, and 0.51 for E-F@ANM electrolytes. Adaptive hydrogen-bonded domains anchoring to solvent molecules as well as constructing Li+-directed channels and lowering Li+ transference barriers results in high tLi+ values.

In addition, the effect of adaptive hydrogen-bonded domains on Li+ deposition/stripping kinetics was analyzed. To quantitatively evaluate the kinetics of Li+ desolvation and diffusion through the SEI, activation energies were determined from using electrochemical impedance spectroscopy (EIS) spectra collected between −20 and 70 °C. This analysis allowed for the determination of activation energies for charge transfer (Ea,Rct) and Li+ transport through the SEI (Ea,RSEI). The values of temperature-dependent charge transfer resistance (Rct) and SEI resistance (RSEI) were extracted from the EIS data using the equivalent circuit model depicted in Fig. S38. As shown in Fig. S39a–f, the Ea,Rct of the E-F@NMTF and E-F@ANM are 61.90 kJ mol−1 and 56.53 kJ mol−1, respectively, which are much lower than that of E-F (70.78 kJ mol−1), indicating that the adaptive hydrogen-bonded domains anchoring strategy facilitates enhanced Li+ desolvation. Meanwhile, the interfacial Li+ diffusion barrier, represented by Ea,RSEI decreased from 49.42 kJ mol−1 in E-F to 43.14 kJ mol−1 in E-F@NMTF and 38.07 kJ mol−1 in E-F@ANM, which was attributed to the formation of an inorganic component-rich SEI with a lower ion mobility barrier. The E-F@NMTF and E-F@ANM were more efficient for lithium plating and stripping due to their anion-rich solvated structures. Fig. S40a–c show the lithium plating morphology of the three electrolytes. Uneven dendritic deposition was observed in the E-F, whereas the deposition microstructure in the E-F@ANM is more uniform and compact compared to that in the E-F@NMTF. The higher coulombic efficiency (CE) values (91.7%/98.9%) of the E-F@NMTF and E-F@ANM (Fig. 4h) compared with that of the E-F (83.43%) can be attributed to the reduction of solvation energy and the inhibition of solvent decomposition in the E-F@ANM. In addition, the lithium nucleation overpotential decreased significantly in the E-F@NMTF and E-F@ANM electrolytes, indicating better interfacial kinetics (Fig. S41). The CE under full plating/stripping conditions was further evaluated at a current density of 1 mA cm−2 and an areal capacity of 1 mAh cm−2. Li||Cu cell cycling with E-F@NMTF and E-F@ANM electrolytes showed stable CE, with average values of 98.1% and 99.0%, respectively, and the E-F@ANM sustained cycling for up to 150 cycles (Fig. 4i). By contrast, the E-F electrolyte barely exceeded 95% CE, and the cell rapidly failed within 30 cycles due to a short circuit caused by lithium dendrites. Comparison of the impedance changes of Li||Cu cells after plating/stripping cycles revealed that the interfacial impedance was significantly reduced with the E-F@NMTF and E-F@ANM compared to the E-F, indicating the formation of a stable and highly ionic-conducting SEI layer and enhanced Li+ charge transport kinetics (Fig. S42). The reductive stability of various electrolytes was evaluated in Li||Cu cells by monitoring the voltage response over the 0~1 V window at a current density of 0.2 mA cm−2 (Fig. S43a). Notably, during the initial discharge to 0 V, corresponding to first-cycle SEI formation, the E-F@ANM system completed the process in the shortest time, indicating fast passivation kinetics and good reductive stability. In addition, a 24 h rest period was introduced during Li||Cu cycling to assess electrolyte corrosivity toward deposited lithium (Fig. S43b). The results showed that both E-F@ANM and E-F@NMTF mitigated lithium corrosion and retained more reversible active lithium after aging. Meanwhile, the increase in exchange current density confirmed this result (Fig. S44). As shown in Fig. 4j, the symmetric cells cycled at 10 mA cm−2 and 5 mAh cm−2 with the E-F@NMTF and E-F@ANM electrolytes exhibited long-term cycling stability above 400 h. In contrast, the polarization voltage of the E-F increased significantly after 10 h. Similar results were observed when current density and capacity were further reduced (Figs. S45–46). The Li||Li symmetric cells cycling at different current densities also demonstrated the performance of the E-F@ANM electrolyte (Fig. S47). Notably, the cycling stability achieved in this study is consistent with that previously reported for ester-based electrolytes, suggesting the lithium metal–electrolyte interface by adaptive hydrogen-bonded domains.

The in situ EIS results for the Li||Li symmetric cell indicated that during the initial cycles of the E-F system, the Rct value remained in a relatively high range, suggesting uneven lithium deposition and ongoing degradation at the electrode–electrolyte interface (Fig. S48a). In contrast, the E-F@NMTF and E-F@ANM systems promoted rapid stabilization of Rct during the initial cycles, which was attributed to the formation of a more stable SEI, thereby ensuring long-term interface stability (Fig. S48b, c). Additionally, by decomposing the high- and low-frequency components of the EIS, the relaxation time distribution (DRT) was obtained. The DRT analysis decomposed the overall impedance into six distinct relaxation time peaks (p1~p6): p1 corresponded to low contact resistance (R); p2 and p3 corresponded to SEI resistance (RSEI); p4 and p5 corresponded to charge transfer resistance (Rct); and p6 corresponded to Ohmic resistance (Rb)44. The significant p2 and p3 features observed in the E-F system indicated instability in the electrode–electrolyte interface properties. In contrast, the nearly negligible Rct and RSEI in the E-F@ANM system suggested a faster Li+ desolvation process and faster Li+ transport through the SEI.

NCM811 electrolyte interface stability

In addition to being LMA compatible, the electrolyte must also be compatible with high-voltage positive electrode materials to increase specific energy. Unfortunately, the oxidation state of Ni in NCM811 increases (e.g., Ni2+ to Ni4+) when the cell voltage increases. The high oxidation state of Ni4+ will catalyze the decomposition of the electrolyte, leading to problems such as overgrowth of the CEI and transition metal loss. To further quantify the effect of the adaptive hydrogen-bonded domains on the oxidation resistance of the solvent, the hydrogen transfer reaction energy between the solvent and the charging positive electrode (using a LiNiO2 substrate) was calculated using DFT. Comparing Fig. 5a, b, the hydrogen transfer energies of EMC (−4.88 eV) and FEC (−4.34 eV) were not only lower than those of the cosolvent ANM (−4.27 eV) but also lower than that of the ANM-EMC cluster (−4.11 eV). In addition, the cross-section of the absorbed solvent molecule relaxation structure at the positive electrode was analyzed for charge density. After the H-transfer reaction, the charge density of O atoms on the absorbing sites changed to a spherical distribution, and the change in O charge density revealed the extent of charge transfer feasibility45. As shown in Fig. 5c, the EMC molecules absorbed at the positive electrode surface exhibit stronger charge transfer activity. In contrast, the FEC molecules had a slower dehydrogenation reaction due to the spatial site resistance effect. Moreover, the anchoring of EMC (C=O) by ANM (N–H) effectively attenuated the charge density of O (Fig. 5d), which demonstrated that the introduction of ANM could effectively alleviate the high activity of EMC in the dehydrogenation reaction of NCM811. X-ray absorption near-edge structure (XANES) studies were performed to further investigate the effect of high-voltage electrolytes on the stability of the bulk positive electrode structure (Fig. S49a, b). The Ni K-edge spectra of NCM811 at a cut-off voltage of 4.7 V exhibited higher edge energies for electrodes in E-F@NMTF and E-F@ANM than for those in E-F, indicating a higher Ni oxidation state. In contrast, the lower oxidation state in the E-F electrolyte was attributed to the continuous side reactions between the ester solvent and the dissolved metal ions, which disrupted the crystal structure of NCM811 thereby hindering Ni oxidation to higher states. The extended X-ray absorption fine structure (EXAFS) spectra and k3-weighted wavelet-transformed (WT) EXAFS spectra were used to probe changes in the local coordination environment of the electrodes after extended cycling in different electrolytes46–50. Compared to fresh NCM811, wavelet-transformed EXAFS for electrodes cycled in E-F@ANM, relative to those cycled in E-F, reveal an increase in Ni–O bond strength, indicative of greater structural order, while the Ni–O bond length increased by 0.01 Å, suppressing the Jahn-Teller effect (Fig. 5e, f and Fig. S50)51,52. The synchrotron X-ray diffraction (SXRD) Rietveld refinement was performed in GSAS-II using a single rhombohedral R-3m phase (Fig. 5g, h and Fig. S51)53. The results showed that the E-F@ANM system (2.58%) exhibited a lower degree of Li/Ni cation mixing relative to the E-F system (3.73%) (Tables S4–6). The E-F@ANM electrolyte enhanced the structural integrity of the NCM811 positive electrode by promoting cation ordering within its layered structure. This improvement was evidenced by a higher I(003)/I(104) intensity ratio, indicative of reduced cation mixing and a more ordered crystal lattice.

To investigate the morphological and compositional changes of the CEI induced by the E-F@NMTF and E-F@ANM electrolytes, the NCM811 positive electrodes were analyzed using scanning electron microscopy (SEM), high-resolution transmission electron microscopy (HRTEM), XPS, and time-of-flight secondary-ion mass spectrometry (TOF-SIMS). Positive electrode morphology after cycling was examined by SEM. As shown in Fig. S52a–d and Fig. S53a–c, the positive electrode particles cycled in the E-F electrolyte were encapsulated by entangled decomposition products from side reactions. They exhibited significant fractures, which could be attributed to the collapse of the positive electrode structure and dissolution of the transition metal elements (TMEs, including Ni, Co, and Mn). In contrast, the positive electrode surfaces in the E-F@NMTF and E-F@ANM electrolytes exhibited intact secondary particles after cycling. Energy-dispersive X-ray spectroscopy (EDS) maps (Figs. S54–57 and Table S7) indicated that the introduction of NMTF reduced the content of TMEs on the positive electrode surface, particularly in the E-F@ANM system. In addition, ICP-MS of the post-cycled electrolytes showed that the contents of TMEs in the E-F@ANM system were much lower than those in the E-F system (Fig. S58). These results strongly suggested that the CEI layer derived from the electrolyte of the H-bond-anchored solvation structure strategy inhibit the solubilization of TMEs ions and effectively protects the positive electrode/electrolyte interface. In addition, with the E-F@ANM electrolyte, a thin and uniform CEI of ~3.5 nm formed on the surface of NCM811, whereas NCM811 cycled in the E-F electrolyte was encapsulated by a thick and non-uniform CEI of 12.8 nm (Fig. S59a–c). The structure and composition of the CEI formed on the surface of NCM811 in the E-F@ANM electrolyte were further analyzed by HRTEM images (Fig. S60a). LiF and Li2O nanoparticles were distributed in the outer layer of the CEI, while Li3N dominated the inner layer. This result was further verified by the corresponding HRTEM and inverse fast Fourier transform (FFT) images (Fig. S60b–d). We analyzed the composition of CEI in NCM811 after 100 cycles using depth-resolved XPS. As shown in the O 1s spectra of NCM811 in Fig. S61a–c, the RCO2Li and C=O peaks assigned to carbonate decomposition products decreased progressively from E–F, E–F@NMTF, and E–F@ANM. This indicated that the thickness of the organic CEI layer formed on these positive electrodes decreases accordingly. Meanwhile, a transition metal oxide peak (TM-O) appeared in the E-F system, indicating transition metal dissolution on the NCM811 surface. In addition, the increase in the intensities of the LiF (Fig. S62a–c) and Li3N peaks (Fig. S63a–c) indicated that the decomposition of FSI− anions on the surface of NCM811 was facilitated, which might contribute to the formation of an inorganic-rich CEI. The intensities of C–F peaks decreased in the E-F@NMTF and E-F@ANM systems, whereas LiF peaks increased. The C–F-containing products maybe generated by FSI− decomposition or attributed to the polyvinylidene fluoride (PVDF) binder. Given that the NCM811 electrodes were prepared under the same conditions, the contribution of PVDF to the C–F signal should not be significantly different in these three samples. Therefore, it could be inferred that the increase in LiF peak intensity should be attributed to the decomposition of the FSI− (Fig. S64a–c).

To further support the XPS analysis and to observe the homogeneity of Li3N/LiF in CEI, we conducted TOF-SIMS on the cycled NCM811 electrodes. As in Fig. 5i, j, C2HO− and C2H3O− fragments correspond to the decomposition products of EMC and FEC. It was observed that the organic components of the E-F system were more extensively distributed along the CEI depth than the E-F@ANM system. In addition, LiO− signal (corresponding to Li2O) was more abundant in the E-F@ANM system than the E-F system (Fig. 5l), which is attributed to decomposition of the FSI− anion. Furthermore, LiF2− (Fig. 5k) and CN− fragments (Fig. 5m) (products of LiFSI) were uniformly distributed throughout the interface (Figs. S65, 66). When LiF was concentrated in the outer layer, it protected the positive electrode interface and reduced side reactions. However, when LiF penetrated deeper into the CEI, it indicated that side reactions were occurring at the positive electrode/electrolyte interface, continuously generating LiF as a decomposition product. This aggravated collapse of the positive electrode structure and led to decrease high-voltage cycling performance. Similarly, in the E-F@ANM system, the peaks of TMEs ions were lower (Fig. S67a–d), and the dissolution of TMEs ions was reduced (Figs. S68, 69). Furthermore, to quantitatively analyze the various components in the CEI, the integrated yields of secondary-ion fragments, including LiF2− (a marker of the thickest CEI layer), from the surface to 50% depth were compared across electrolytes, and these data were used for relative quantification of degradation products (Fig. S70)54. The solvent decomposition products in the E-F@ANM system were lower than those in the E-F system, and the inorganic CEI content arising from anion decomposition products was higher in E-F. This indicated that adaptive hydrogen-bonded domains anchoring strategy reduced electrolyte decomposition by slowing down solvent decomposition under high voltage and forming a CEI rich in Li3N/LiF with high ionic conductivity and good electronic insulating properties.

The effect of the E-F@ANM induced CEI layer on the crystal structure of the protected NCM811 at a high-voltage of 4.7 V was further investigated by using in situ X-ray diffraction (XRD) to track the shifts of the (003) and (104) peaks. During the initial charging process, the (003) peak gradually moved to a lower 2θ angle and finally transforms to (003)H1, indicating the lattice extension along the c-axis. When the charging voltage reached about 4.25 V, the (003)H2 peak shifted back to a higher 2θ angle, indicating the phase transition from H2 to H355. A significant anisotropic volume change accompanies the H2-H3 phase transition, and their reversible H2-H3 phase transition can lead to the formation of microcracks at the phase boundaries of the NCM811 material56. The (003)H2-H3 peak shift at the positive electrode in the E-F system was 1.76° (Fig. 6a and Fig. S71), which was higher than the peak shift in the E-F@ANM system (1.71°) (Fig. 6b and Fig. S72). During charging, the peak of (104) shifted to a higher angle, indicating lattice contraction along the a- and c-axes. Similarly, the peak shift in the E-F@ANM system decreased from 2.03° to 1.87°. More importantly, the narrower voltage range of the H2 to H3 phase transition in the E-F@ANM demonstrates that the E-F@ANM system effectively mitigates the irreversible phase transition and improved the structural stability of the positive electrode of the NCM811 during charging and discharging. The dQ/dV contour plots during cycling further confirmed that the E-F@ANM system suppressed the irreversible phase transitions of H2-H3, H1+M and M+H2 (Fig. 6c, d)57. The above results demonstrated that the E-F@ANM derived CEI film effectively suppressed the irreversible phase transition in the positive electrode, thereby enhancing the stability of the NCM811 active material during the cycling process.

In situ EIS was used to detect the impedance changes during delithiation and lithiation in a high-voltage environment. At each equal voltage, the cell impedance of the E-F@ANM system was smaller than that of the cell using the E-F system (Fig. 6e, f). This indicated that the CEI layer of E-F@ANM has fast and stable ionic conduction efficiency. The effect of the E-F@ANM system on Li+ migration kinetics was further confirmed by relaxation time in situ DRT analysis. By mathematically deconvolving the electrochemical impedance data in the relaxation time-space, DRT can accurately distinguish between the interface resistance (Rf; Rf = RCEI + RSEI) and charge transfer resistance (Rct). In this case, the Rct of the cell was located in the low-frequency region (101 > τ > 10−1), which corresponded to the kinetics of the Faraday process associated with the dissolution and desolvation of Li+ occurring at both electrodes. The mid-frequency peaks (10−1 > τ > 10−3) can be attributed to the Rf58. The Rf decreases slightly during Li+ desolvation, corresponding to CEI formation, and then rises during discharge, reflecting CEI stripping59. The smaller Rf and Rct in the E-F@ANM compared to the cell using the E-F system suggested that Li+ transported more readily through the CEI and the Li+ desolvation was much faster (Fig. 6g, h and Fig. S73a, b). And the Rct of E-F was very large near 4.7 V, indicating that Li+ is difficult to desolvate at high-voltages. This result further demonstrated that the improved CEI, induced by the formation of adaptive hydrogen-bonded domains, promotes Li+ desolvation.

Electrochemical properties of electrolytes

To demonstrate the feasibility of the adaptive hydrogen-bonded domains anchoring strategy in high-voltage LMBs, we conducted a comparative study of the electrochemical performance of Li||NCM811 cells using different electrolytes. The cycling of Li||NCM811 cells with the E-F@NMTF and E-F@ANM electrolytes were much stable at a cut-off voltage of 4.7 V, exhibiting high capacity retention of 66.5% (127.4 mAh g−1) and 78.8% (152.4 mAh g−1) after 400 cycles with E-F@NMTF and E-F@ANM electrolytes, respectively (Fig. 7a). More importantly, when charged to 4.7 V, the E-F@ANM electrolyte enabled an initial coulombic efficiency of 89.6% for NCM811 (Fig. 7b). In contrast, the cell’s capacity with the E-F electrolyte dropped drastically to only 66 mAh g−1 after 400 cycles. In addition, the coulombic efficiency of the E-F@ANM electrolyte was stably above 99.8%, which was much higher than the average coulombic efficiency of 99.1% for the E-F electrolyte, proving that the oxidative decomposition of the E-F@ANM electrolyte was minimal (Fig. S74). Moreover, the E-F@ANM system exhibited stable rate performance (Fig. 7d), with a discharge-specific capacity of 146.6 mAh g−1 at a high C-rate of 3 A g−1. The electrolyte side-reaction rate was accelerated at high temperatures of 60 °C. The battery with the E-F@ANM electrolyte retained 79.5% of its capacity after 250 cycles (Fig. 7c and Fig. S75). In addition to the 4.7 V NCM811, we had also demonstrated the high-voltage stability of the E-F@ANM electrolyte in 4.8 V Li-rich layered oxide positive electrodes Li1.2Ni0.13Co0.13Mn0.54O2 (LRLOs) and 4.6 V LiCoO2 (LCO) (Fig. S76a–d, Fig. S77a–d and S74), and showed the stable rate performance in Li||LCO (Fig. S79). Furthermore, we assembled lithium metal pouch cells under harsh real-world conditions. To maximize the specific energy, we used a highly loaded negative electrode (20 mg cm−2), a lean electrolyte (2.4 g Ah−1), and a thin lithium negative electrode (80 μm, N/P ratio of 1.7). Here, the specific energy of the 1.6 Ah pouch cell with the E-F@ANM electrolyte was as high as 418.2 Wh kg−1 (Fig. 7e, Fig. S80 and Tables S8, 9). The voltage profile of the Li||NCM811 pouch cell with E-F@ANM electrolyte was shown in Fig. S81, where the first cycle at 20 mA g−1 produces an initial coulombic efficiency of up to 89.8%. The battery then maintained 80% capacity retention after 70 cycles of stable cycling at 20 mA g−1 charging/40 mA g−1 discharging. The capacity of the pouch cell with the E-F electrolyte decayed much faster, and the capacity dived after the 22th charging process, leading to battery failure. This indicated that severe electrolyte side reactions, negative electrode structure collapse and lithium dendrite growth occur in the E-F electrolyte under high-voltage conditions, making stabilizing LMBs for stable cycling at such high mass loading challenging. This also confirmed that the adaptive hydrogen-bonded domains anchoring strategy is conducive to improving the intrinsic stability of the electrolyte, decreasing Li+ desolvation energy, effectively forming electrode-electrolyte interfaces with low interfacial impedance and high ionic conductivity. The Ah-grade specific energy presented in this work is comparable to that reported in recent relevant studies (Fig. 7f), demonstrating potential for practical applications60–74. To further validate the practical value of our E-F@ANM system, a 2 Ah-class Gr||NCM811 full cell was tested under 4.5 V conditions (Fig. S82). With the E-F system, the capacity retention declined to 62.5% after 300 cycles. In contrast, the E-F@ANM system exhibited slow capacity decay, retaining 91.3% of its capacity after cycling, demonstrating its effective suppression of electrolyte decomposition, transition metal dissolution, and impedance growth under high-voltage conditions, while maintaining a stable interface layer. These results indicate that the E-F@ANM system is not only compatible with LMBs systems but also holds significant potential for high-voltage operation in commercial lithium-ion batteries.

Methods
Material

Battery-grade LiFSI, EMC, and FEC were obtained from Zhuhai Saiwei Electronic Materials Co. and used as-received. Polyvinylidene difluoride (PVDF) and N-methyl pyrrolidone (NMP) were obtained from N-methyl pyrrolidone (NMP, Aladdin) and used as-received. Formamide (FM), Acetamide (Ac), Trifluoroacetamide (TFAc), Cyanoacetamide (ANAc), N-Methylacetamide (NMAc), N-N-Dimethylacetamide (NNDMAc), N, N-Dimethyltrifluoroacetamide (NNDMTFAc), 2-Cyano-N,N-dimethylacetamide (ANNNAc), N-Methyl-trifluoroacetamide (NMTF) and 2-Cyano-N-methyl-acetamide (ANM) (all in battery-grade purity) were purchased from Aladdin Biochemical Technology Co (Aladdin). Thick lithium metal negative electrodes (15.6 mm diameter and 450 μm thick) and thin Li metal negative electrodes (15.6 mm diameter and 50 μm thick) were received from China Energy Lithium. Polycrystalline NCM811 particles, carbon black (Super P), Celgard 2400 battery separator, Cu collectors, etched Al collectors, 2032-type coin-cell cases, springs and spacers were purchased from Kelude Co, Ltd. positive electrodes (NCM811 positive electrode, LCO positive electrode, LRLOs positive electrode) were prepared by mixing active materials, Super P (Timcal), and polyvinylidene fluoride (PVDF 5130, MTI) in N-methyl-2-pyrrolidone (NMP, Aladdin) in a ratio of 8:1:1 to form a slurry, which was then coated onto aluminum foils. The positive electrodes were initially dried at 80 °C for 4 h and subsequently in a vacuum oven at 120 °C for 12 h. Post-dried positive electrodes were punched into circular positive electrodes disks with a diameter of 13 mm, and then transferred to an argon-filled glovebox (water and oxygen levels maintained below 0.01 ppm) for subsequent use. The mass loadings of the three as-prepared positive electrodes (NCM811, LCO, and LRLOs) were measured to be 14.9 mg cm−2, 14.7 mg cm−2, and 8.84 mg cm−2, respectively, with all positive electrodes adopting a single-side coating configuration.

Electrolytes were prepared by dissolving the lithium salt LiFSI in the chosen solvents in glass sample vials equipped with polytetrafluoroethylene (PTFE) magnetic stir bars, within a glovebox filled with inert argon gas (water and oxygen levels maintained below 0.01 ppm). The mixtures were stirred at 28 °C for 12 h to yield the electrolytes.

Electrochemical measurements

CR2032 coin cells were assembled in an argon-filled glovebox for electrochemical tests, with 50 μL of electrolyte added before crimping. The as-assembled coin cells were rested at 28 °C for 6 h prior to conducting constant current charge-discharge cycling tests. The charge/discharge intervals of coin cell Li||LCO, Li||NCM811, and Li||LRLOs are 2.8 V~4.6 V, 2.8 V~4.7 V, and 2 V~4.8 V, respectively. Li||NCM811, Li||LCO, and Li||LRLOs coin cells were first pre-cycled twice at 20, 18, and 25 mA g−1, respectively, followed by continuous cycling at 200, 180, and 250 mA g−1 (1C = 200, 180, and 250 mA g−1 for NCM811, LCO, and LRLOs, respectively). For Li||NCM811 pouch cells, charging was performed at a constant current-constant voltage (CC-CV) of 0.1 C, and discharging at a constant current (CC) of 0.2 C (NEWARE-CT4008 battery testing system). Linear sweep voltammetry (LSV) and cyclic voltammetry (CV) tests were conducted using a BioLogic instrument, while cycling tests were carried out using a Neware battery cycler (CT-4008) at 28 °C. The LSV curves were obtained using a three-electrode system at a sweep rate of 1.0 mV s–1. Platinum wire (Φ = 0.1 mm, 99%) and two lithium pieces were used as the working counter, and reference electrodes, respectively. The voltage ranges from 3.5~6.0 V vs Li/Li+ for the LSV measurement. For Li||Li symmetric cells, thick lithium metal negative electrode, separator, and thick lithium metal negative electrode were stacked together, while Cu replaced the lithium metal in Li||Cu cells.

The Tafel test was conducted on an AutoLab (PGSTAT302N) electrochemical workstation with the Li||Li symmetric cells. The electrochemical impedance spectra (EIS) measurements were conducted on an AutoLab (PGSTAT302N) electrochemical workstation over a frequency range from 100 kHz to 0.1 Hz with an amplitude of 5 mV. The measurement was conducted with 5 points per decade in logarithmic spacing.

The Li+ transference number was determined with chronoamperometry at a constant step potential of 10 mV, which is based on Li||Li symmetric cells after several cycles to form a stable SEI. Li+ transference number was calculated according to the following Eq. (1):

tLi+ = I0(ΔV - IsR0) / Is(ΔV - I0R0) (1)

Where tLi+ is the transference number, I0 and Is are the initial and steady state current; R0 and Rs are the initial and steady state resistances of the passivating layers, V is a constant applied polarizing potential.

To obtain the activation energies of Li+, the EIS tests of Li||Li symmetric cells under various temperatures were carried out. The following Eq. (2, 3) can obtain the activation energy:

1/Rct = A e(-Ea,Rct/RT) (2)

1/RSEI = A e(-Ea,RSEI/RT) (3)

Where RSEI, Rct, A and R represent the Li+ transport through the SEI resistance, charge–transfer resistance, pre–exponential constant and the standard gas constant, respectively.

The exchange current density was obtained from the Tafel Eq. (4):

η = A × (log i - log i0) (4)

where η and A represent overpotential and the kinetic constant in the diffusion–controlled surface dynamic range, respectively. And i0 represents the exchange current density. The value of i0 can be obtained from the intersection of the extrapolated linear part of the log i versus η plot with the μ = 0 line.

The modified Aurbach method was carried out in Li||Cu half cells to evaluate the average CE of different negative electrodes. Modified Aurbach’s method was applied as follows. First, 5 mAh cm–2 Li was plated on Cu and charged until 0.5 V at 0.5 mA cm–2 to passivate the Cu collector. Second, plated 5 mAh cm–2 Li at 0.5 mA cm–2 on Cu again. Third, charged to 1 mAh cm–2 at 0.5 mA cm–2 and then discharged to 1 mAh cm–2 at 0.5 mA cm–2 for 10 cycles. Last, charged to 0.5 V at 0.5 mA cm–2 and stripped all remaining Li on Cu. Then, a pre-plating capacity of Li (QT) was used as the Li source and a fixed capacity Li (QC) was used between the Li working and counter electrodes. After n cycles, the remaining Li reservoir was completely stripped to the cut-off voltage. By measuring the amount of Li remaining after cycling, the average CE over n cycles can be calculated by the Eq. (5):

CE = (nQC + QS) / (nQC + QT) (5)

Where QS corresponds to the final stripped capacity or the amount of Li remaining after n cycles. Herein, the value of QT of the pre-plating Li was 5 mAh cm–2, and the value of QC was 1 mAh cm−2. Finally, QS was measured after 10 cycles and stripping to 0.5 V. at 1 mA cm−2. The current density of the above operations was 0.5 mA cm−2.

Characterizations

At 28 °C, the ionic conductivity of the electrolyte was measured using an 856 conductivity module (from Sweden). At 28 °C, the ionic conductivity of the electrolyte was measured using an 856 conductivity module (from Sweden). For physical and chemical characterization, coin cells before and after cycling were transferred to an argon-filled glovebox and disassembled. The as-prepared electrode samples were placed in glass sample vials, vacuum-sealed for storage, and maintained in this state until transported to the testing site. X-ray photoelectron spectroscopy (XPS, Thermo Fisher, Nexsa G2) was used to analyze the Li electrode after 10 cycles. A 1 keV Ar ion gun sputtered In-depth XPS (Thermo Fisher, Escalab 250Xi, the sputtering area was 3 mm × 3 mm, and the sputtering rate was calibrated to 1.5 nm/min by thermally oxidated SiO2 (standard). Sputtering time is 400 s (10 nm), 800 s (20 nm). The field emission scanning electron microscopy (FESEM, ZEISS Gemini 500) was used to analyze the morphology of NCM811 and sodium deposition after cycling. The CEI of the solid electrolyte membrane of the NCM811 positive electrode after cycling was measured using a transmission electron microscope (TEM) operated at 100.0 kV using HT7700 (Hitachi, Tokyo, Japan). Raman spectra were obtained with inViaTM confocal Raman microscope (Qontor, RENISHAW) using a 532 nm laser source. 1H, 7Li NMR of the electrolytes were conducted on 600 MHZ NMR (AVANCE III HD). TOF-SIMS (ION-TOF.SIMS5) investigated the composition of the SEI on the Li metal negative electrode. The 30 keV Bi3+ was used as the primary-ion gun. The area of analysis is 100 μm × 100 μm, while the sputtering area is 400 μm × 400 μm. All sample preparation was done in an argon atmosphere glovebox. Using in situ Fourier transform infrared (FTIR) spectroscopy (Thermo Scientific Nicolet 6700 spectrometer) in the attenuated total reflection (ATR) mode with a wavenumber range of 500~4000 cm–1 and a resolution of 1.0 cm–1, the dynamic conditions of the electrolyte interfacial reactions. An in situ X-ray diffraction XRD measurement was performed using a Bruker D8 advance equipped with a Neware battery testing system. The measurements were conducted within a 2θ range of 10° to 80° with a step rate of 10° min–1. The amount of transition metal ions (Mn and Ni) dissolved from the cycled NCM811 positive electrode was detected by Inductively Coupled Plasma-Mass Spectrometry (ICP-OES, PerkinElmer Optima 8000, USA).

As for the local structure, the hard X-ray absorption spectroscopy (XAS) experiments with transmission mode were performed using the BL13SSW beamline in the SSRF at Shanghai, China. Ni-foils were measured simultaneously as the standard comparison sample. The ATHENA software package75 was used to process and fit the X-ray absorption near-edge structure (XANES) and extended X-ray absorption fine structure (EXAFS) data.

Synchrotron X-ray diffraction (SXRD) patterns were obtained at the beamline BL14B of the Shanghai Synchrotron Radiation Facility (SSRF) in which the X-ray wavelength was 0.6887 Å.

MD simulation

GROMACS 2021.6 package76 for molecular dynamics simulations and the Packmol package for creating the initial simulation box77. The OPLS-AA force field parameters were used to describe the components of the system. The charges of the anions and cations were scaled down by a factor of 0.8 to account for polarization and charge transfer effects78. The FSI– anions were taken from the literature79,80, while the solute molecules used the 1.2 CM5 charge obtained at the B3LYP-D3(BJ)/def2-TZVPP computational level. The system was first subjected to a 3000-step energy minimization using the steepest gradient method to eliminate unreasonable overlaps between atoms, followed by 100 ps relaxation under NVT and NPT tethered traces, respectively, and finally the whole system was simulated by sampling the whole system under NPT tethered traces for 50 ns. In the MD simulations, periodic boundary conditions were used in all three directions, and all hydrogen-containing chemical bonds (C-H, O-H) bond lengths were constrained using the LINCS algorithm with a timestep of 2.0 fs. The temperature was maintained at 298.15 K, by means of a V-rescale thermostat with a coupling time of 1.0 ps. The pressure was controlled at 1 bar using a Parrinello-Rahman thermostat with a coupling time of 2.0 ps. Long-range electrostatic interactions were handled using the Particle Mesh Ewald (PME) method, and short-range electrostatic and vdW interactions were calculated using a cut-off of 1.2 nm. Trajectories were analyzed and processed and visualized using the Gromacs and VMD 1.9.3 software packages81, respectively.

Density functional theory (DFT) calculation

Density Functional Theory (DFT) calculations were used to obtain further information on the hydrogen bonding bond energies and binding energies between the electrolyte components. DFT theoretical calculations were performed using the built-in B3LYP two-hybridization generalization of Gaussian 16 software82 in combination with DFT-D3 dispersion corrections with damping83,84. Geometry optimization of the solutes and frequency analysis was carried out using a B3LYP/def2SVP calculation setup to obtain an optimized structure and ensure that it is an energy minimum point. Finally, the B3LYP-D3(BJ)/def2TZVPP calculation level was used to obtain more accurate electron densities and energies. Calculation of hydrogen bonding bond energies based on Eq. (6):

Ebind = HAB - HA - HB (6)

Where Ebind is the bonding energy and HAB, HA and HB are the enthalpies of the AB complex and the A and B molecules, respectively.

In addition, to explore the mechanism of hydrogen bond formation, Multiwfn 3.8 software85 was used to derive the full electron density diagram of the system before and after the formation of the hydrogen bond.

The structural optimization and frequency analysis were performed by the Gaussian 16 A.03 package. The density functional theory (DFT) function based on the 6-311+G(d,p) level basis set with the use of an implicit solvent model was applied to calculate LUMO/HOMO, ESP. The binding sites were investigated by the molecular electrostatic potential (MESP) method using Multiwfn 3.8 based on wavefunction statistics. All MESP isosurface maps are extracted by visual molecular dynamics (VMD) 1.9.3 program based on the outputs of the Multiwfn analysis. The adsorption energy was calculated as following Eq. (7):

Eads = EAB - EA - EB (7)

Where EAB, EA, EB are the total energy of AB combination, the total energy of A and the total energy of B, respectively.

Theoretical calculation of the solvent dehydrogenation reaction on the surface of NCM811. All computations were carried out with VASP, where electron–ion interactions were represented using the PAW formalism. The exchange–correlation term was described by the PBE functional within the GGA framework. The plane-wave basis set was truncated at 500 eV. Self-consistent field cycles were terminated once the total-energy change dropped below 10–5 eV. Reciprocal-space integration used a 3×2×1 k-point sampling. Geometry optimizations were performed until the remaining atomic forces were reduced to under 0.03 eV Å–1 for every atom.

Solvent contributions to the oxidation potential were incorporated via a dielectric continuum description. Aniline was taken as the solvent medium with a dielectric constant of 7.3, and Eox referenced to Li/Li+ was computed following Eq. (8):

EOX = (G(M+) - G(M)) / F - 1.4 V (8)

At 298.15 K, the Gibbs free energies of the neutral and oxidized states are denoted as G(M) and G(M+) respectively, and F is the Faraday constant. Deformation charge density analysis was conducted with DMol3 (Materials Studio).

Peer review rounds
[{"round": 1, "reviewer_comments": {"reviewer_1": "The authors proposed a solvation structure engineering strategy by introducing a hydrogen bond donor co-solvent ANM to\nconstruct adaptive hydrogen-bonding domains, aiming to optimize Li⁺ transport pathways and interfacial stability in high-\nvoltage LMBs. While the study presents promising electrochemical performance under high cut-off voltages and high\ncathode loading, several critical concerns remain to be addressed or clarified before further consideration:\n1. The claimed novelty of modulating Li⁺ solvation structure via hydrogen bonding is not sufficiently justified. Similar\napproaches, especially those involving Z-bond interactions and weak hydrogen bonding networks, have been extensively\nreported in recent literature. The manuscript lacks a clear articulation of how this work surpasses the existing strategies in\nterms of design principle or performance enhancement.\n2. The selection of ANM as a co-solvent is rather conventional. The –NH– and –CN functional groups are commonly\nemployed in high-voltage electrolyte systems in LMBs. However, the manuscript does not provide any systematic structure–\nproperty correlation or design rationale that distinguishes ANM from previously reported additives with similar motifs.\n3. The strength and nature of hydrogen bonding were primarily evaluated using FTIR shifts and RDG analysis. However,\nstate-of-the-art methods such as NBO analysis and QTAIM are more appropriate for quantitatively assessing hydrogen bond\nstrength and electronic characteristics. The manuscript lacks rigorous theoretical analysis to validate and compare the\nstrength of the H-bonds formed by ANM.\n4. The concept of \"oriented Li⁺ migration pathways\" proposed in Figure 1d is attractive, but there is no direct evidence\ndemonstrating the directional transport or adaptive anchoring behavior of hydrogen bond domains.\n5. The authors use E-F as the baseline electrolyte throughout the study, but fail to include control formulations using\nstandard carbonate-based electrolytes (e.g., EC/EMC) or non-H-bond-donating cosolvents. This omission prevents a\nconclusive attribution of the observed improvements solely to the hydrogen bond engineering strategy.\n6. While XPS and TOF-SIMS depth profiling are performed, the analysis remains qualitative. There is no quantitative ratio\nanalysis of organic/inorganic components, and the spatial distribution of inorganic species in TOF-SIMS appears non-\nuniform. Therefore, the conclusion that CEI is predominantly inorganic in composition on NCM811 surfaces is not sufficiently\nsupported.\n7. The electrochemical stability window was evaluated by LSV, but the manuscript lacks complementary characterizations\nsuch as Tafel plots, interfacial impedance evolution, or full-range EIS fitting to further support the proposed interfacial\nimprovements. Additionally, in the in situ EIS/DRT analysis of Li||NCM811 cells, the attribution of mid-frequency peaks\nsolely to RCEI is debatable, as RSEI can also contribute significantly in this region.\n8. The authors claim to maximize energy density by employing high-loading cathodes (20 mg cm⁻²), lean electrolyte (2.4 g\nAh⁻¹), and a thin lithium anode (80 μm, N/P ratio = 1.7). However, such a high N/P ratio is impractical for industrial-scale\nproduction due to both economic and safety concerns associated with lithium metal. Therefore, the soft-pack pouch cell\nvalidation in its current configuration lacks practical relevance and is of limited reference value for practical applications.", "reviewer_2": "In this study, Yang et al. demonstrated that stronger hydrogen bonds and non-typical hydrogen bond interactions are crucial\nfor regulating the performance of lithium metal batteries under high-pressure conditions. The electrolyte modified by a\nhydrogen bond donor (ANM) was validated to demonstrate superior electrochemical performance compared to most\npreviously reported electrolytes. This work opens a new avenue for designing and regulating electrolyte. In summary, I\nbelieves that this paper provides valuable information for the readers of Nature Communications, which can be published\nafter fully addressing the concerns below.\nconcentration potentially increase the viscosity of the electrolyte and reduce its conductivity? Based on the authors’ in-depth\nanalysis of hydrogen bonding between solvent molecules, it is suggested that a more rational discussion regarding the\nselection of additive concentrations should be included.\nlithium metal. Can the series of solvents (TFAc, ANAc, NNDMTFAc, ANNNAc, NMTF, and ANM) mentioned by the authors\nremain stable in lithium metal batteries? It is recommended that the authors provide further results and discussion on this\npoint.\n(e.g, TFAc, ANAc, NNDMTFAc, ANNNAc, NMTF, ANM) should be clearly defined.\nelectrolyte systems? Besides, why the addition of NMTF and ANM resulted in an enhancement of electrical conductivity\n(Lines 6-8, Page 8)?\ncathodes with E-F and E-F@ANM derived CEI should be tested and compared in E-F system (e.g., XRD of such cathodes\nafter prolonged cycling).", "reviewer_3": "The authors present a compelling strategy for modulating the solvation structure of high-voltage electrolytes through the\ndeliberate exploitation of intermolecular hydrogen bonding. By combining advanced theoretical simulations with systematic\nexperimental characterization, they conclusively demonstrate that hydrogen-bonding interactions significantly enhance\nelectrolyte stability at extreme potentials (>4.5 V vs. Li+/Li), promote favorable desolvation kinetics, and stabilize critical\nelectrode-electrolyte interfaces. This rational design enables the stable operation of high-energy-density lithium metal pouch\ncells at an unprecedented 4.7 V cutoff voltage. The work provides fundamental insights into electrolyte engineering and\nrepresents a significant advancement toward next-generation battery systems. The findings are of broad interest to the\nenergy storage community and warrant publication after addressing the following revisions.\n1. The author proposes in Figure 1 that the atypical hydrogen bonds of ANM, as revealed by 1HNMR and theoretical\ncalculations, further enhance the effects of H-bond. Could the magnitude of the Z-bond (Nδ⁻–Hδ⁺) be calculated to better\nillustrate its strength relative to typical hydrogen bonds?\n2. The author conducted float charging tests on Li||NCM811 button cells to demonstrate the high-voltage stability of the\nelectrolyte, but I would still like to know if there are other experimental techniques that can further verify the reduction\nbehaviour of this electrolyte. This would help to more comprehensively evaluate the stability of the electrolyte.\n3. The author performed activation energy fitting of the Nyquist plot at (30oC~60oC) in Figure S34 and suggested\nsupplementing the impedance with a wider temperature range to more convincingly prove the kinetic phenomenon.\n4. The author demonstrates the effect of hydrogen bonding domains on lithium metal anodes in Figure 4 and suggests\nsupplementing the in situ EIS of Li||Li symmetric cells to prove the improvement of the electrolyte on the kinetics of lithium\nmetal anodes.\n5. The author mentioned the dissolution phenomenon of transition metals, but Figure S50 lacks supporting data on Co\nelement leaching. It is recommended that the author supplement the data to prove the positive effect of ANM co-solvent in\nimproving the voltage resistance performance of the electrolyte."}, "author_response": "RESPONSE TO REVIEWERS’ COMMENTS\nWe thank all the reviewers for their valuable comments on our manuscript. Their\nconstructive suggestions have certainly raised the quality of our manuscript. The main\ntext as well as the supplemental materials have been revised accordingly and\nhighlighted with YELLOW background.\nNew experiments and discussions:\n Figure 4 was revised consistent with the accompanying discussions.\n 1 new Figure and 3 new Tables were added to the SI.\n 2 new density-functional theory (DFT) calculations of the QTAIM method were\nused to perform on different hydrogen-bonded domains (Table S1-2).\n 1 new experiment was conducted to compare the contact stability of lithium metal\nand electrolyte with electrolytes containing cyanoethyl groups (Figure S3).\n 1 new set of experiments was conducted on the application of hydrogen-bonded\ndomains in EC-based electrolytes (Figure S12-14).\n 1 new data analysis on the influence of hydrogen-bonded domains on electrolyte\ndecomposition at the cathode interface during charging, dynamically characterized\nby in situ FTIR. (Figure S29).\n 1 new experiment was conducted to obtain more accurate activation energy data\nover a wider temperature range (Figure S39).\n 1 new experiment was conducted to investigate the effect of hydrogen-bonded\ndomains of the electrolyte on the corrosion resistance of lithium metals (Figure\nS43).\n 1 new experiment was conducted to investigate the protective effect of hydrogen\nbonded electrolytes on lithium metal using in situ EIS and corresponding DRT\nanalysis of Li||Li symmetric cells (Figure S48).\n 1 new experiment was conducted to supplement the quantitative analysis of Co\nelement dissolution after cycling (Figure S58).\n 1 new experiment was conducted to supplement the fast Fourier transform (FFT)\nanalysis of Li N, Li O, and LiF in CEI from HRTEM images (Figure S60).\n3 2\n 1 new data analysis comparing the total yield of secondary ion fragments in\nnormalized TOF-SIMS spectra of two electrolytes (Figure S70).\n 1 new electrochemical long-cycle performance test was conducted using Ah-\ngrade Gr||NCM811 pouch cells (Figure S82).\n 1 energy density calculation for the lithium metal pouch cell has been added\n(Table S8).\nReviewer #1 (Remarks to the Author):\nComments:\nThe authors proposed a solvation structure engineering strategy by introducing a\nhydrogen bond donor co-solvent ANM to construct adaptive hydrogen-bonding\ndomains, aiming to optimize Li+ transport pathways and interfacial stability in high-\nvoltage LMBs. While the study presents promising electrochemical performance under\nhigh cut-off voltages and high cathode loading, several concerns remain to be addressed\nor clarified:\nReply: We greatly appreciate the reviewer’s comprehensive and meticulous\nassessment of this manuscript which contributed significantly to the thoroughness of\nthe study. We have carefully revised our paper considering the reviewer’s comments\nand suggestions. We provided point-by-point replies to address the reviewer’s concerns.\nComment 1. The claimed novelty of modulating Li+ solvation structure via hydrogen\nbonding is not sufficiently justified. Similar approaches, especially those involving Z-\nbond interactions and weak hydrogen bonding networks, have been reported in recent\nliterature. The manuscript lacks a clear articulation of how this work surpasses the\nexisting strategies in terms of design principle or performance enhancement.\nReply: We thank the reviewer for the constructive comments.\nIn the current literature, several studies have utilized the nonclassical hydrogen-\nbond (Z-bond) interactions to modulate electrolyte solvation structures of electrolytes\nin lithium metal batteries, thereby enhancing battery stability and performance. For\nexample, Ren et al. proposed using molecular anchoring strategy to suppress side\nreactions of ether solvents at the electrode interface, thereby achieving high-voltage\nperformance in low-concentration electrolytes. The researchers selected fluorinated\nethers containing the –CF H group (e.g., 1,1,2,2-tetrafluoroethyl-2,2,3,3-\n2\ntetrafluoropropylether, TTE) as molecular anchors. These molecules can form Z-bonds\nwith the ether solvent 1,2-dimethoxyethane (DME), reducing the electron density on\nthe ether oxygen atoms, thereby enhancing the ether’s oxidation resistance (Nat.\nCommun. 2023, 15, 2033). Fan et al. proposed a molecular docking strategy, where Z-\nbonds are formed between fluorobenzene (FB) or dichloromethane (DCE) and 1,2-\nbis(tri-methoxy methyl)ethane (BTE) or 1,3-bis(tri-methylmethoxy)propane (BTP),\naltering the solvent docking configuration and reducing the desolvation energy of Li+\n(Nat. Chem. 2024, 16, 1427-1435). However, these existing studies generally focus on\nadjusting the solvation structure and de-coordination mechanism of electrolytes\nthrough tuning a single Z-bond. The potential mechanisms governing the rational\nregulation of co-solvent–solvent H-bond and Z-bond interactions at the electrode–\nelectrolyte interface remain insufficiently understood.\nSpecifically, the novelty of our present work is stated as follows:\n The screening of electrochemically stable hydrogen bond donor amides for\nmodulating the size and stability of hydrogen-bonded domains.\n For the first time, the contributions of H-bonds and Z-bonds to the hydrogen-\nbonding domain were distinguished, and their presence was quantified\nthrough spectroscopic characterization and computational methods.\n Hydrogen-bonded domains mediate highly stable inorganic-rich liquid-solid\ninterfaces by remodeling the solvated structure of the interface.\nThe atypical hydrogen bond (Z-bond) in our study not only acts independently but also\nenhances the formation of typical hydrogen bond (H-bond), with intermolecular\ninteractions that are far more complex than those observed in traditional studies of Z-\nbond. While the formation of Z-bonds and H-bonds can be directly determined through\nin situ Raman/FT-IR spectroscopy, this adaptive hydrogen-bonded domains still\nrequires additional characterization methods.\nOur research indicates that Z-bonds are weak interactions formed between atoms\nwith partial charges (e.g., Nδ⁻–Hδ+). Z-bonds are weaker than typical hydrogen bonds,\nwith strengths approximately one-thousandth of coordination bonds. This substantial\ndifference in bond energy raises intriguing questions: how can this weak Z-bond\nmodulate stronger coordination bonds? Can its effect be amplified by coupling to\nstronger interactions? Because of the relatively weak binding strength of the Z-bond,\nits ability to alter the solvation structure is limited and far less pronounced than that of\nstrong H-bonds.\nImportantly, our study differentiates the interactions between H-bonds and Z-\nbonds for the first time. When H-bond interactions dominate, the concomitant Z-\nbonds create adaptive hydrogen-bonded domains that stabilize the solvation\nenvironment while facilitating rapid Li+ transport. Notably, we performed DFT\nscreening of a series of amide molecules as hydrogen-bond acceptors. The reason for\nselecting amides is their relatively high stability in hydrogen bond donors for lithium-\nmetal batteries (Figure S1-5). We successfully confirmed the differences in the size of\nH-bonds formed between amides and electrolytes with different side groups (Figure\n3d-e; S17-26) through integrated infrared, Raman, and NMR spectroscopy techniques,\naccompanied with molecular dynamics simulations. Furthermore, to verify the role of\nthis hydrogen-bonded domains under charge-discharge conditions, we employed\nadvanced in situ Raman and infrared spectroscopy at the anode and cathode interfaces\nto test the hydrogen-bonded domain’s effect on electrolyte solvent decomposition and\nits de-coordination mechanism (Figure 4b-c). We also performed DFT calculations to\nassess the impact of hydrogen-bonded domains on solvent dehydrogenation at\nNCM811 cathodes and their protective role in preserving the layered structure of\nNCM811 (Figure 5a-h).\nOverall, our research elucidates the fundamental principles distinguishing H-\nbonds and Z-bonds. When both coexist, H-bond interactions dominate. Across kinetic\nstudies and interfacial analyses, hydrogen-bonded domains stabilize the solvation\nenvironment, thereby enhancing high-voltage stability and facilitating rapid Li+\ntransport in lithium-metal batteries.\nChanges made: In the revised manuscript, we added: (Page 5, line 16)\nThe presence of specific co-solvent molecules has been reported in literature to\nreshape the interfacial molecular interactions by forming Z-bonds with the solvent\nmolecules, which greatly facilitates Li+ desolvation dynamics. Despite these advances,\nthe potential mechanisms of how to rationally regulate the co-solvent-solvent H-bond\nand Z-bond interactions on the behavior of the electrode-electrolyte interface are still\npoorly understood.\nComment 2. The selection of ANM as a co-solvent is conventional. The –NH–and –\nCN functional groups are commonly employed in high-voltage electrolyte systems in\nLMBs. However, the manuscript does not provide any systematic structure-property\ncorrelation or design rationale that distinguishes ANM from previously reported\nadditives with similar motifs.\nReply: Thank you for your valuable comments. Although amide molecules exhibit\nhigh thermodynamic stability, the primary issue with the –NH– functional group in\nsolvent molecules is its presence of “active hydrogen”. Lithium metal, known for its\nlowest electrochemical potential and the strongest reducing ability, is highly sensitive\nto proton sources. When solvent molecules containing –NH– groups come into contact\nwith the lithium metal surface, the weakly acidic hydrogen atom in the N–H bond is\nreadily reduced by lithium, generating hydrogen gas (H ) and the corresponding lithium\n2\nsalts (e.g., lithium amides).\nThe reaction can be represented as: 2 R–NH–R' + 2 Li → 2 R–N⁻ Li+–R' + H ↑.\nsolvent rather than during electrochemical cycling, resulting in extremely low first-\ncycle coulombic efficiency (CE) and rapid capacity fading. The evolution of H leads\n2\nto an increase in internal pressure within the battery, causing serious safety hazards,\nsuch as swelling or even rupture. Continuous gas generation compromises the integrity\nof the SEI, preventing it from effectively passivating the lithium metal surface.\nResearchers have improved the instability of amides towards lithium metal by\nintroducing additives or adjusting the main solvent strategy (Joule 2024, 8, 2542-4351;\nACS Energy Lett. 2024, 9, 3931-3938). Additionally, the nitrogen atom in the –NH–\nfunctional group possesses a lone pair of electrons, making it a nucleophilic center that\nis susceptible to electrochemical oxidation. This oxidation process is catalyzed and\nintensified at high-voltage on the surface of cathode materials (Adv. Energy Mater. 2022,\n12, 2200139). In contrast to the one-sided disadvantages associated with the –NH–\ngroup, the –CN (nitrile) functional group exhibits a certain “duality.” It has high\noxidation stability and was once considered a candidate for high-voltage solvents.\nHowever, the –C≡N triple bond in nitrile solvents with low lowest unoccupied\nmolecular orbital (LUMO) energy levels (such as acetonitrile and propionitrile) is a\nstrong electron-withdrawing group, which results in very low LUMO energy levels for\nthese molecules, meaning that they are highly prone to electron gain and reduction (Adv.\nMater. 2024, 36, 2404271; Adv. Funct. Mater. 2025, 35, 2421329). Existing strategies\nprimarily focus on the molecular design of nitrile solvents, adjusting the regional\nelectronic cloud and polarity to slow down side reactions with lithium metal (Adv. Funct.\nMater. 2025, 35, 2421329; Angew. Chem. Int. Ed. 2024, 63, e202408902).\nIn this study, we first screened different short-chain amide molecules and\nperformed a comparative analysis of their H-bond stability using DFT calculations,\ncombined with cluster stability analysis. We found that molecules containing –NH\n2\nexhibited poor electrochemical stability due to the reactivity of their active hydrogen\nwith lithium metal batteries (Figure S1). At the same time, introducing –CF or –\n2\nresistance of the amide (Figure S2). Additionally, we have demonstrated\nexperimentally that molecules containing –CH CN exhibit good compatibility with the\n2\nelectrolyte and with lithium metal batteries (Figure S3).\nMore importantly, the design principle proposed in this study is not merely based\non the individual chemical characteristics of –NH– or –CN, but rather on the spatially\ncoordinated arrangement of these two functional groups within the solvation\nstructure of the electrolyte (Figure S4-6). This arrangement allows them to play\ncomplementary roles in the solvation structure regulation. –NH–, as a typical hydrogen\nbond donor, can form a strong hydrogen bond (–10.8 kcal·mol–1) with the solvent's\ncarbonyl oxygen, effectively stabilizing the solvation structure. –CN, as a Z-bond (Nδ⁻–\nHδ+) acceptor site, forms weak interactions (–2.5 kcal·mol–1) with the hydrogen of the\nadjacent solvent –CH – (Table S1). While the interaction is weak, it can polarize and\n2\norient the surrounding carbonyl groups, thereby enhancing the stability of the main –\nNH–···O=C hydrogen bond. Furthermore, we established a structure-property\nrelationship in our work: by quantifying the strength of the H-bond and Z-bond\ninteractions, changes in hydrogen bond stability, Raman/NMR shifts, and molecular\ndynamics simulations. We demonstrated that ANM, compared to other amide molecules,\ncan simultaneously optimize the solvation and desolvation processes and stabilize the\nelectrode/electrolyte interface.\nChanges made: In the revised manuscript, we added:\n(Page 7, line 16)\nMoreover, molecules bearing a −CH CN group exhibited long-term chemical stability\n2\ntoward both the electrolyte and lithium metal, and no appreciable gas evolution or\ndiscoloration was observed (Figure S3).\n(Page 8, line 24)\nDFT calculations confirmed the formation of Z-bond (−2.5 kcal mol–1) (Table S1). In\naddition, the H-bond between NMTF/ANM and the (C–F) sites in FEC were\ncomparatively weak, owing to the greater steric hindrance of (C–F) and its lower\nelectron density relative to (C=O).\nIn the revised supporting information, we added：\nTable S1. The magnitudes of the interactions between H-bonds, Z-bonds, and van der\nWaals forces observed in this study.\nForce Magnitude\nComplexes Bond Type Types of Forces\n(kcal mol–1)\nANM-EMC N–H∙∙∙O=C H-bond –10.8\nANM-FEC N–H∙∙∙O=C H-bond –7.4\nNMTF -EMC N–H∙∙∙O=C H-bond –8.8\nNMTF-FEC N–H∙∙∙O=C H-bond –7.6\nNMTF-FEC N–H∙∙∙F-C H-bond –4.4\nEMC-FEC / Van der Waals force –6.9\nANM-EMC C≡N∙∙∙H–C Z-bond –2.5\nFigure S3. Optical photographs of fresh electrolytes (top), and electrolytes (middle)\nand lithium metals (bottom) after immersing lithium metals in electrolytes for two\nweeks.\nComment 3. The strength and nature of hydrogen bonding were primarily evaluated\nusing FTIR shifts and RDG analysis. However, state-of-the-art methods such as NBO\nanalysis and QTAIM are more appropriate for quantitatively assessing hydrogen bond\nstrength and electronic characteristics. The manuscript lacks theoretical analysis to\nvalidate and compare the strength of the H-bonds formed by ANM.\nReply: Thanks for the careful review and constructive feedback. We employed\nFTIR and RDG analyses primarily to identify and characterize hydrogen-bond\nformation and properties. FTIR shift analysis can infer the strength of hydrogen bonds\nby observing the changes in the stretching frequency, while RDG reveals intermolecular\ninteraction regions and strengths via the electron-density distribution. The QTAIM and\nNBO analyses mentioned by the reviewer are advanced tools for the quantitative\nassessment of hydrogen bond strength and electronic properties. QTAIM can precisely\nquantify the strength of hydrogen bonds by analyzing the electron density (ρ) and\npotential energy density (V(r)) at the bond critical point (BCP), and is particularly suited\nfor revealing the electron transfer and interactions during hydrogen bond formation.\nAlthough NBO analysis can provide rich electronic structure information, it is not the\nmost direct tool for describing solvation structures and ion transport. In our study,\nRaman peak fitting was conducted for a detailed quantitative analysis, which effectively\ncharacterizes the formation of hydrogen bonds and the solvation process. MD\nsimulations demonstrate that ANM can form local hydrogen-bonded structural domains\nwithin a distance of less than 3.5 Å from the carbonate molecules (Figure S22a-b). The\nformation of these hydrogen-bonded domains stabilizes the solvent molecules and\nresults in a solvation structure featuring an anion-rich inner coordination shell and an\nouter hydrogen-bonded domains. Raman spectroscopy of both free and bound hydrogen\nbonds confirms that the hydrogen-bonded domains formed by ANM are more stable\nand more numerous, providing strong evidence for the exceptional stability of the E-\nF@ANM electrolyte (Figure S23a-d).\na b\n72.17%\n52.14%\nu.) u.)\na. a.\ny\n(\ny\n(\nn\nsit\n25.45% n\nsit\ne 22.41% e\nnt nt\nI I\n17.05%\n10.78%\nFree NMTF NMTF-EMC NMTF-FEC Free ANM ANM-EMC ANM-FEC\nFigure S22. MD simulation results revealing the percentages of free solvents and\nhydrogen-bonded domains (distance between solvents <3.5 Å) in E-F@NMTF\nelectrolyte (a) and E-F@ANM electrolyte (b).\nFigure S23. Intensities of H-bonds and non H-bonds in (a) pure NMTF solvent and E-\nF@NMTF electrolyte and (b) pure ANM solvent and E-F@ANM electrolyte.\nIntensities of non H-bond (c) and H-bond (d) (ratios calculated from the area of the\nfitted peaks).\nAdditionally, DFT calculations have provided us with an in-depth analysis of\nhydrogen bond strength, helping to quantify the bond strength and electronic\ncharacteristics through parameters such as electron density maps and bond lengths.\nConsidering that the main goal of this study is to explore the interactions between amide\ncosolvents and electrolyte carbonate solvent molecules, without delving into the\nelectronic properties of the molecules, we did not employ NBO analysis.\nWe supplemented our analysis with QTAIM calculations, which reveal the detailed\ncharacteristics of hydrogen bonds. Each H-bond has a BCP, located along the\ninteraction path between the atoms involved in the bond. QTAIM calculations are key\nindicators for quantitatively analyzing hydrogen bond strength. Higher ρ and negative\nV(r) typically indicate stronger hydrogen bonds. In the QTAIM results, all clusters\nexhibit low ρ and positive Laplacian (▽2ρ), collectively indicating the presence of weak\nhydrogen bonding interactions (Table S2). The ANM-EMC (N–H∙∙∙O=C) interaction\nshows the highest ρ, the lowest potential energy density V(r), and the shortest bond\nlength, indicating the strongest hydrogen bond formation between ANM and EMC\nmolecules. Furthermore, the weaker interaction in the ANM-FEC cluster compared to\nthe NMTF-FEC cluster is consistent with the results obtained from RDG analysis, FTIR,\nRaman spectroscopy, and MD simulations presented in this work. This result suggests\nthat in the E-F@ANM system, more FEC molecules participate in the primary solvation\nshell of Li+. Additionally, an atypical hydrogen bond (C≡N∙∙∙H–C) was identified in the\nANM-EMC cluster, which is consistent with the findings from RDG analysis and 1H\nNMR experiments. The presence of this atypical hydrogen bond further enhances the\nintermolecular H-bond interactions.\nChanges made: In the revised manuscript, we added: (Page 8, line 28)\nFurthermore, quantum theory of atoms in molecules (QTAIM) calculations were\nemployed to elucidate the detailed characteristics of the hydrogen bonds. Higher\nelectron density (ρ) and a more negative potential energy density (V(r)), generally\nindicate stronger H-bonds. In the QTAIM results, all clusters exhibited low ρ and a\npositive Laplacian (▽2ρ), collectively indicating weak H-bond interactions (Table S2).\nThe ANM and EMC (N–H···O=C) interaction showed the highest ρ, the most negative\nV(r), and the shortest bond length, indicating that the H-bond between ANM and EMC\nwas the strongest. In addition, relative to the NMTF and FEC cluster, the ANM and\nFEC cluster displayed weaker interactions. Moreover, a Z-bond (C≡N···H–C) was\nidentified in the ANM and EMC cluster, corroborating that Z-bond coupling in the\nANM and EMC system further strengthened the H-bond interaction.\nIn the revised supporting information, we added：\nTable S2. Composition of different electrolytes. The bond length and electron density\n(ρ), Laplacian (▽2ρ), and potential energy density (V(r)) at the bond critical points\n(BCPs) from each H-bond.\nQTAIM\nComplexes Bond Type Bond Length (Å)\nρ ▽2ρ V(r)\nANM-EMC N–H∙∙∙O=C 1.9616 0.0217 0.0972 ‒0.0170\nANM-FEC N–H∙∙∙O=C 2.0860 0.0169 0.0659 ‒0.0111\nNMTF-EMC N–H∙∙∙O=C 2.0209 0.0215 0.0898 ‒0.0162\nNMTF-FEC N–H∙∙∙O=C 2.0098 0.0206 0.0777 ‒0.0142\nANM-EMC C≡N∙∙∙H–C 2.7489 0.0056 0.0612 ‒0.0052\nWe have incorporated the above analysis into the revised manuscript and\nSupporting Information. Once again, we sincerely thank you for your insightful\nfeedback, which has enabled us to elaborate more comprehensively on this key issue\nand further refine the manuscript.\nComment 4. The concept of “oriented Li+ migration pathways” proposed in Figure 1d\nis attractive, but there is no direct evidence demonstrating the directional transport or\nadaptive anchoring behavior of hydrogen bond domains.\nReply: We sincerely appreciate your valuable comments regarding the concept of\n‘directional Li+ migration pathways’ in our study. In conventional dilute-salt\nelectrolytes, Li+ migrates predominantly via a vehicular mechanism, i.e., motion of the\nentire solvated cluster; when solvent–solvent and solvent–anion interactions strengthen,\nthe solvent degrees of freedom decrease and anions enter the first solvation shell,\nmaking Li+ desolvation easier (Nature 2024, 627, 101–107). During charging in the E-\nF@ANM electrolyte, the presence of self-adaptive hydrogen-bond domains renders the\ninterface simultaneously anion-rich and solvent-deficient. This is corroborated by the\ndecreased population of free EMC molecules during charging (Figure 3e and Figure\nS22b), and by the delayed onset of FSI⁻ decomposition as the voltage increases (starting\nat ~4.6 V).\nTo more intuitively visualize the evolution of free and hydrogen-bonded species,\nas well as hydrogen-bond strength, during charge and discharge, we analyzed the\npotential-dependent changes in the FWHM and integrated intensity of the non H-bonds\nand H-bonds in operando infrared spectra. (J. Energy Chem. 2023, 81, 2095–4956). We\nemployed a simple semi-analytical deconvolution of the in situ FTIR spectra. First, the\nfree N–H band was fitted with a single Gaussian function, with the center frequency\nand FWHM fixed by the IR spectrum of ANM. Beyond the Gaussian-defined free N–\nH component, the area of the hydrogen-bonded component extracted from the observed\nspectra quantifies the extent of the hydrogen-bonded domains. A larger FWHM together\nwith a larger H-bond peak area indicates stronger hydrogen bonding, and vice versa\n(RSC Adv. 2018, 8, 23481–23488). At open-circuit voltage (OCV), hydrogen-bonding\ninteractions are relatively weak, as reflected by the smaller H-bond FWHM and peak\narea. Upon applying voltage, both the H-bond FWHM and peak area increase,\nindicating rapid interfacial H-bond formation; this trend persists until ~4.5–4.6 V. At\n~4.6 V, both the H-bond FWHM and peak area begin to decrease, consistent with\nweakened hydrogen bonding caused by partial solvent decomposition (Figure S29).\nThis arises because earlier anion decomposition reduces negative interfacial charge,\nallowing insufficiently desolvated anions from the bulk to reach the interface and\ntrigger partial decomposition of carbonate solvents; this behavior persists up to ~4.8 V.\nThese observations confirm the evolving role of hydrogen-bond domains during\ncycling. The distinctive “solvent-dragging” effect promotes preferential anion\ndecomposition at the interface, consistent with the interfacial analysis presented in the\nmanuscript. Once again, we appreciate your constructive feedback on our manuscript.\nChanges made: In the revised manuscript, we added the text: (Page 15, line 9)\nTo visualize the evolution of free ANM, the hydrogen-bonded domains, and H-bond\nstrength during charge–discharge, we analyzed how the full width at half maximum\n(FWHM) and integrated intensity of the non H-bond and H-bond N–H peaks in in situ\ninfrared spectra varied with the applied potential (Figure S29a-b). First, the free N–H\npeak was fitted with a single Gaussian function, whose center frequency and FWHM\nwere determined from the FTIR spectrum of ANM. The area of the H-bonded\ncomponent extracted from the measured spectrum—outside the Gaussian-defined free\nN–H component—quantitatively characterized the extent of the hydrogen-bonded\ndomains. A larger FWHM together with a larger H-bond peak area indicated stronger\nH-bond, and vice versa. At open-circuit voltage (OCV), H-bond interactions were\nrelatively weak, as reflected by the small FWHM and peak area of the H-bonded feature.\nUpon applying a potential, both the FWHM and the H-bonded peak area increased,\nindicating rapid formation of an interfacial hydrogen-bonded domains; this trend\npersisted up to 4.6 V. Above 4.6 V, both metrics decreased, attributable to partial solvent\ndecomposition that weakened the H-bond. In the early stage of charging, anion\ndecomposition reduced the interfacial negative charge; under these conditions,\ninsufficiently desolvated anions from the bulk were less able to reach the interface,\nwhile the high potential triggered partial decomposition of EMC and FEC. This\nbehavior persisted up to 4.8 V. Collectively, these observations confirmed the evolution\nof the hydrogen-bonded domains during cycling, and its distinctive “solvent-dragging”\neffect that promoted preferential anion decomposition at the interface.\nIn the revised supporting information, we added the figure：\nFigure S29. In situ FTIR analysis of Li||NCM811 batteries using E-F@ANM\nelectrolyte to separate H-bond and non H-bond FWHM peaks (a) and the ratio of peak\narea integrals (b). (For specific peak separation methods, refer to Figure S23.)\nComment 5. The authors use E-F as the baseline electrolyte throughout the study, but\nfail to include control formulations using standard carbonate-based electrolytes (e.g.,\nEC/EMC) or non-H-bond-donating cosolvents.\nReply: We thank the reviewer for this constructive comment. In conventional\nEC/EMC electrolytes, EC predominantly governs Li+ solvation and dictates the\nchemical composition of the SEI and CEI. The excessively strong interaction between\nLi+ and EC weakens the causal relationship we aim to investigate—namely,\nrestructuring the Li+ primary solvation shell via the construction of hydrogen-bonded\ndomains. Furthermore, EC-based electrolytes tend to form inferior SEI films, limiting\ntheir applicability in high-voltage lithium metal batteries. Therefore, in the main text\nwe selected FEC/EMC (E-F) as the baseline system to eliminate EC’s strong solvation\npreference, allowing hydrogen-bond engineering to emerge as the dominant variable.\nIn addition, we employed co-solvents such as N,N-dimethyltrifluoroacetamide\n(NNDMTFAc) and 2-cyano-N,N-dimethylacetamide (ANNNAc), which possess the\nsame polar functional groups but lack N–H hydrogen-bond donors. DFT calculations\nconfirmed that neither NNDMTFAc nor ANNNAc forms strong hydrogen bonds with\nEMC or FEC. This trend was clearly reflected in the electrochemical performance\nshown in Figure S11, where co-solvents containing N–H hydrogen-bond donors\nyielded higher capacity retention after cycling compared to those lacking hydrogen-\nbond donors.\nTo directly address the reviewer’s suggestion, we have added control experiments\nbased on EC/EMC: in EC : EMC = 1:1 (v/v) containing NMTF or ANM (5 M), the N–\nH stretching vibration in the FTIR spectra exhibited a pronounced blue shift,\nunequivocally confirming the presence of hydrogen-bonding interactions (Figure\nS12a-b). These features were absent in the pure EC/EMC system, indicating that the\nobserved spectral changes originate from the hydrogen-bond donor capability of NMTF\nand ANM. Peak-fitting analysis of the Raman spectra in the PF ⁻ vibration region\n6\nrevealed that both NMTF and ANM decreased the proportion of solvent-separated ion\npairs (SSIPs) while increasing the fractions of contact ion pairs (CIPs) and aggregates\n(AGGs), indicating that hydrogen-bond domains promote tighter Li+–anion association\n(Figure S13a-c). This observation aligns with the mechanism proposed in the\nmanuscript, wherein hydrogen bonding weakens solvent–Li+ interactions and drives\nanions into the primary solvation shell. In NCM811||Li cells operated at 4.5 V, those\nemploying the E-E@ANM electrolyte achieved a capacity retention of 93.3% after 200\ncycles, markedly surpassing that of cells with E-E electrolyte (63.6%) and E-E@NMTF\nelectrolyte (75.7%) (Figure S14). This trend is consistent with the pattern observed in\nthe E-F system, and the magnitude of performance enhancement correlates positively\nwith the hydrogen-bond strength of the additive, further substantiating the pivotal role\nof hydrogen-bond domains.\nChanges made: In the revised manuscript, we added the text: (Page 10, line 1)\nSimilarly, such H-bond interactions were operative across different carbonate solvents.\nIn the FTIR spectra of the EC-EMC-NMTF and EC-EMC-ANM mixed solvents, the\nN–H stretching vibration showed a pronounced blue shift, unambiguously confirming\nH-bond (Figure S12a-b). Moreover, deconvolution of the PF – Raman bands indicated\n6\nthat introducing NMTF or ANM decreased the fraction of solvent-separated ion pairs\n(SSIP) while increasing the fractions of CIP and AGG, suggesting that H-bond\npromoted tighter Li+–anion association (Figure S13a-c). In Li||NCM811 cells\nemploying the EC-EMC (E-E)@ANM electrolyte (Table S3), the capacity was retained\nat 93.3% after 200 cycles, significantly outperforming cells using the E-E electrolyte\n(Figure S14).\nIn the revised supporting information, we added the figures：\nFigure S12. FTIR spectra region using NMTF (a) and ANM (b) co-solvents (5 mol\nL⁻1) added to EC: EMC=1 : 1 (v/v) solvent mixture.\nFigure S13. The Raman spectra of (a) EC-EMC (E-E), (b) E-E@NMTF, and (c) E-\nE@ANM electrolytes with the proportions of the three solvated structures determined\nbased on peak area integration.\nFigure S14. The long cycle performance of Li||NCM811 coin cells at 2.8–4.5 V with\ndifferent electrolytes.\nComment 6. While XPS and TOF-SIMS depth profiling are performed, the analysis\nremains qualitative. There is no quantitative ratio analysis of organic/inorganic\ncomponents, and the spatial distribution of inorganic species in TOF-SIMS appears\nnon-uniform. Therefore, the conclusion that CEI is predominantly inorganic in\ncomposition on NCM811 surfaces is not sufficiently supported.\nReply: We appreciate the reviewer’s valuable comments regarding the XPS and\nTOF-SIMS analyses in our study. Our research shows that the LiF – fragments in E-\n2\nconcentrated near the outer surface. Specifically, when LiF is concentrated in the outer\nlayer, it can protect the cathode interface and reduce side reactions. However, if LiF\npenetrates deeper into the CEI, it indicates that side reactions are occurring between the\ncathode surface and the electrolyte, continuously generating LiF as a decomposition\nproduct. This exacerbates cathode degradation and leads to poor high-rate cycling\nperformance. We compared the integrated yields of various secondary ion fragments\nfrom the surface to the 50% level cross-section of the thickest CEI fragments (LiF –)\n2\nformed in different electrolytes and used these to relatively quantify the degradation\nproducts (Figure S70) (J. Am. Chem. Soc. 2025, 147, 7, 6023–6036). The results\nconfirmed the expected trend: the solvent decomposition products in E-F@ANM are\nlower than those in the E-F electrolyte, and the inorganic CEI content of anion\ndecomposition products is higher in E-F, which is consistent with the elemental ratio\nresults from XPS (Figure S64a-c). Using high-resolution transmission electron\nmicroscopy (HRTEM), we confirmed that the CEI formed by the E-F@ANM\nelectrolyte is uniform and dense. Furthermore, we conducted a detailed study on the\nstructure and composition of the CEI formed on the NCM811 surface in the E-F@ANM\nelectrolyte. As shown in Figure S59a-c, a dense and uniform CEI layer approximately\n5 nm thick was observed. Further analysis revealed that LiF and Li O nanoparticles are\n2\ndistributed in the outer layer of the CEI, while Li N dominates the inner layer. This\n3\nresult was further validated by the corresponding HRTEM images and inverse FFT\nimages (Figure S60a-d).\nChanges made: In the revised manuscript, we added the text: (Page 25, line 21)\nThe structure and composition of the CEI formed on the surface of NCM811 in the E-\nF@ANM electrolyte were further analyzed by HRTEM images (Figure S60a). LiF and\nLi O nanoparticles were distributed in the outer layer of the CEI, while Li N dominated\n2 3\nthe inner layer. This result was further verified by the corresponding HRTEM and\ninverse fast Fourier transform (FFT) images (Figure S60b-d).\n(Page 26, line 28)\nSimilarly, in the E-F@ANM system, the peaks of TMEs ions were lower (Figure S67a-\nd), and the dissolution of TMEs ions was reduced (Figures S68-69). Furthermore, to\nquantitatively analyze the various components in the CEI, the integrated yields of\nsecondary-ion fragments, including LiF − (a marker of the thickest CEI layer), from the\n2\nsurface to 50% depth were compared across electrolytes, and these data were used for\nrelative quantification of degradation products (Figure S70). The solvent\ndecomposition products in the E-F@ANM system were lower than those in the E-F\nsystem, and the inorganic CEI content arising from anion decomposition products was\nhigher in E-F. This indicated that adaptive hydrogen-bonded domains anchoring\nstrategy reduced electrolyte decomposition by slowing down solvent decomposition\nunder high voltage and forming a CEI rich in Li N/LiF with high ionic conductivity\n3\nand good electronic insulating properties.\nIn the revised supporting information, we added the figures：\nFigure S70. A comparison of the overall yield of secondary ion fragments from the\nnormalized TOF-SIMS spectra of the two electrolytes.\nFigure S71. (a) TEM image of the NCM811 cathode after cycling with E-F@ANM,\nalong with the corresponding inverse FFT images of (b) Li O, (c) LiF, and (d) Li N.\n2 3\nComment 7. The electrochemical stability window was evaluated by LSV, but the\nmanuscript lacks complementary characterizations such as Tafel plots, interfacial\nimpedance evolution, or full-range EIS fitting to further support the proposed interfacial\nimprovements. Additionally, in the in situ EIS/DRT analysis of Li||NCM811 cells, the\nattribution of mid-frequency peaks solely to RCEI is debatable, as RSEI can also\ncontribute significantly in this region.\nReply: Thanks for valuable comments. We analyzed the Tafel curves of the Li||Li\nsymmetric cell, and the E-F@ANM system showed a lower exchange current density.\nAdditionally, we fitted the in situ EIS impedance data for R and R (fitting method in\nf ct\nFigure S38), which indicated faster charge transfer kinetics and lower interfacial\nimpedance in the E-F@ANM system (Figure S73b).\nThe DRT peak in the mid-frequency range (10⁻1~10⁻3 s) of full cell may indeed contain\nresponses from both the SEI and CEI, rather than being solely attributed to R . When\nCEI\nthe relaxation time constants of the two processes are very close, their DRT peaks may\nstill overlap. To address this, we supplemented our study with in situ EIS measurements\nand DRT analysis of Li||Li symmetric cells (Figure S39a-f). Surprisingly, the interfacial\nresistance of the lithium anode (R ) measured in the Li||Li symmetric cell was\nSEI\nsignificantly higher than the total interfacial resistance R (including contributions from\nf\nR and R ) measured in the Li||NCM811 full cell (Figure R1). Since the NCM811\nCEI SEI\ncathode stores Li+ via intercalation, while the lithium metal anode undergoes Li+\ndeposition and stripping, their mass transport mechanisms are not entirely the same.\nDuring cycling of the Li||Li symmetric cell, Li+ deposition and stripping are essentially\nspontaneous and disordered. Due to unavoidable microscopic roughness on the lithium\nmetal surface, the electric field becomes concentrated at tip regions during deposition,\ncausing Li+ to preferentially deposit there. This “tip effect” reinforces itself and induces\nlithium dendrite growth. When the lithium anode is paired with an NCM811 cathode in\na full cell, the situation changes fundamentally. The NCM811 cathode, acting as a\n“container” for Li+, provides a relatively uniform and stable Li+ source during charge-\ndischarge processes. This controlled ion flow ensures a more spatially uniform Li+ flux\nreaching the anode. Compared with the intense, uncontrolled current density at tip\nregions in symmetric cells, the current distribution in full cells is much smoother. The\nuniform ion flux and electric field distribution effectively suppress the ‘tip effect’ at the\nlithium anode surface. Li+ are no longer prone to excessive local accumulation and\ndeposition but instead tend to deposit relatively evenly and densely across the electrode\nsurface. Overall, the presence of the NCM811 cathode actively transforms the lithium\nanode interface from a high-impedance, unstable state to a low-impedance, relatively\nstable one. Although the observed mid-frequency peak is a superposition of R and\nCEI\nR , the “flattening effect” greatly reduces the magnitude of R , thereby optimizing\nSEI SEI\nthe overall interfacial performance of the full cell. We therefore conclude that R in the\nf\nfull cell is composed of both R and R , with R being the dominant controlling\nSEI CEI CEI\nfactor, consistent with our finding that the E-F@ANM electrolyte forms superior\nCEI/SEI layers at both electrodes. We again thank the reviewer for offering a valuable\nopportunity to deepen and enhance this study.\nFigure R1. Impedance DRT of Li||Li symmetrical cell and Li||NCM811 cell during the\nfirst charging cycle.\nChanges made: In the revised manuscript, we added: (Page 21, line 19)\nThe in situ EIS results for the Li||Li symmetric cell indicated that during the initial\ncycles of the E-F system, the R value remained in a relatively high range, suggesting\nct\nuneven lithium deposition and ongoing degradation at the electrode–electrolyte\ninterface (Figure S48a). In contrast, the E-F@NMTF and E-F@ANM systems\npromoted rapid stabilization of R during the initial cycles, which was attributed to the\nct\nformation of a more stable SEI, thereby ensuring long-term interface stability (Figure\nS48b-c). Additionally, by decomposing the high- and low-frequency components of the\nEIS, the relaxation time distribution (DRT) was obtained. The DRT analysis\ndecomposed the overall impedance into six distinct relaxation time peaks (p1~p6): p1\ncorresponded to low contact resistance (R ); p2 and p3 corresponded to SEI resistance\nc\n(R ); p4 and p5 corresponded to charge transfer resistance (R ); and p6 corresponded\nSEI ct\nto ohmic resistance (R ). The significant p2 and p3 features observed in the E-F system\nb\nindicated instability in the electrode–electrolyte interface properties. In contrast, the\nnearly negligible R and R in the E-F@ANM system suggested a faster Li+\nct SEI\ndesolvation process and faster Li+ transport through the SEI.\nIn the revised supporting information, we added the figure：\nFigure S48. In situ impedance spectroscopy and DRT of Li||Li symmetric cells in E-F\n(a), E-F@NMTF (b) and E-F@ANM (c) electrolytes.\nComment 8. The authors claim to maximize energy density by employing high-loading\ncathodes (20 mg cm⁻2), lean electrolyte (2.4 g Ah⁻1), and a thin lithium anode (80 μm,\nN/P ratio = 1.7). However, such a high N/P ratio is impractical for industrial-scale\nproduction due to both economic and safety concerns associated with lithium metal.\nTherefore, the soft-pack pouch cell validation in its current configuration lacks practical\nrelevance and is of limited reference value for practical applications.\nReply: We appreciate the valuable suggestions provided. Lithium metal anodes\nare often regarded as the “holy grail” of battery technology, owing to their unparalleled\ntheoretical performance. First, lithium metal possesses an exceptionally high theoretical\nspecific capacity of 3860 mAh g⁻1—over ten times that of commercial graphite anodes\n(theoretical specific capacity ~372 mAh g⁻1) (Nat. Nanotechnol. 2017, 12, 194–206). Studies have shown that simply replacing a conventional graphite anode with lithium\nmetal could increase the overall energy density of a battery by approximately 40~50%.\nSuch a substantial improvement holds transformative potential for extending the\ndriving range of electric vehicles, prolonging the operating time of portable electronics,\nand enabling other applications where weight and volume are critical constraints. For\nexample, Cai et al. reported that lithium metal batteries (LMBs) could achieve energy\ndensities of ~450 Wh kg⁻1 and 1200 Wh L⁻1, far exceeding those of current lithium-ion\nbatteries (250~300 Wh kg⁻1 and 600~700 Wh L⁻1), positioning them as a key\ntechnology for electric vehicle applications. They further identified ten critical\nrequirements—including improvements in energy density, electrolyte formulation, and\ncell pressure control—that must be addressed to meet industry needs for electric\nvehicles (Nat. Energy 2024, 9, 1199–1205). Xu et al. developed a compact ion-pair\naggregate (CIPA) electrolyte, enabling a 505.9 Wh kg⁻1 lithium metal pouch cell to\nretain 91% of its energy after 130 cycles with an average Coulombic efficiency of\n99.94%. This electrolyte resolved the interfacial stability conflict between the cathode\nand anode found in conventional electrolytes and achieved 300 stable cycles under lean-\nelectrolyte conditions (E/C = 1.25 g Ah⁻1), laying the foundation for the practical\nimplementation of high-energy-density LMBs (Nat. Energy 2024, 9, 987–998). Huang\net al. proposed stabilizing LMBs using microemulsion electrolytes that leverage liquid-\nliquid interfacial tension, thereby substantially enhancing performance. This\nmicroemulsion electrolyte enabled ampere-hour-scale pouch cells to achieve optimal\ncycling, delivering high energy densities of 531 Wh kg⁻1 and 547 Wh kg⁻1, with 81%\nand 79% capacity retention after 189 and 155 cycles, respectively (Nature 2025, 1255–\n1262). Collectively, these studies illustrate the persistent pursuit of high-energy-density\nbattery systems, highlighting that achieving “ultimate” energy density is a systems-\nlevel challenge, not solely determined by anode material selection. The reviewer has\nnoted our use of a high-loading cathode (20 mg cm⁻2), lean electrolyte (2.4 g Ah⁻1), and\nthin lithium anode (80 μm). This parameter combination was deliberately chosen to\nmaximize cell-level energy density under conditions closely approximating practical\napplications, with the understanding that reducing the N/P ratio is an essential pathway\ntoward LMB commercialization.\nAdmittedly, the commercialization of LMBs still faces significant limitations. To\nfurther substantiate the practical value of our developed electrolyte, we have\nsupplemented our study with performance evaluations in a mainstream lithium-ion\nbattery system. We selected a 2 Ah-grade Gr||NCM811 full cell—highly promising for\ncommercialization yet demanding in electrolyte requirements—and tested it at an\nelevated cutoff voltage of 4.5 V (Figure S82). With the E-F electrolyte, capacity\nretention dropped to 62.5% after 300 cycles, indicating rapid degradation due to poor\ninterfacial stability and severe side reactions under high-voltage conditions. In contrast,\nthe E-F@ANM electrolyte exhibited slow capacity decay, retaining 91.3% capacity\nafter cycling, suggesting its effectiveness in suppressing electrolyte decomposition,\ntransition-metal dissolution, and impedance growth under high voltage, while\nmaintaining a stable interfacial layer. These results indicate that the E-F@ANM\nelectrolyte is not only compatible with LMB systems but also holds significant potential\nfor enabling high-voltage operation in commercial lithium-ion batteries.\nChanges made: In the revised manuscript, we added the text: (Page 31, line 13)\nTo further validate the practical value of our E-F@ANM system, a 2 Ah-class\nGr||NCM811 full cell was tested under 4.5 V conditions (Figure S82). With the E-F\nsystem, the capacity retention declined to 62.5% after 300 cycles. In contrast, the E-\nF@ANM system exhibited slow capacity decay, retaining 91.3% of its capacity after\ncycling, demonstrating its effective suppression of electrolyte decomposition, transition\nmetal dissolution, and impedance growth under high-voltage conditions, while\nmaintaining a stable interface layer. These results indicate that the E-F@ANM system\nis not only compatible with LMBs systems but also holds significant potential for high-\nvoltage operation in commercial lithium-ion batteries.\nIn the revised supporting information, we added the figure：\nFigure S82. Long-term cycling performance of the two electrolytes in 2 Ah-capacity\nGr||NCM811 pouch cells under 0.5 C charge/1 C discharge conditions with a cutoff\nvoltage of 4.5 V.\nReviewer #2 (Remarks to the Author):\nComments:\nIn this study, Yang et al. demonstrated that stronger hydrogen bonds and non-typical\nhydrogen bond interactions are crucial for regulating the performance of lithium metal\nbatteries under high-pressure conditions, thereby filling a research gap in the study of\nhydrogen bond interactions in electrolytes. Furthermore, the electrolyte formulated by\nthe authors, 1.2 M LiFSI-5M ANM-FEC/EMC (3:7), demonstrated remarkable efficacy\nin high-pressure lithium metal batteries, exhibiting superior electrochemical\nperformance compared to most previously reported electrolytes under more stringent\ntesting conditions. The article provides a comprehensive explanation of the underlying\nprinciples and mechanisms, offering extensive theoretical support and providing\nexperimental validation of these principles. In summary, I believe that this paper\nprovides valuable information for the readers of Nature Communications and\nrecommends the following revisions.\nReply: We appreciate the reviewer’s recognition of the novelty of our hydrogen-\nbonded domains electrolyte design in high-voltage lithium metal batteries. We are also\nvery grateful to the reviewer for recognizing the attractiveness of the results and the\nthoroughness of the study. To address the reviewer’s concerns, we have provided point-\nby-point replies and carefully revised the manuscript accordingly.\nComment 1. Why did the author choose amides as hydrogen bond donors rather than\nother hydrogen bond donors?\nReply: We thank the reviewer for this advice. In this study, amide compounds\nwere selected as hydrogen bond donors instead of common alternatives such as alcohols,\ncarboxylic acids, or hydrogen fluoride (HF), owing to the multiple advantages amides\nexhibit in electrochemical systems. Firstly, the N–H bonds in amide molecules can\nform relatively moderate hydrogen bonds with the carbonyl groups (C=O) in carbonate-\nbased solvents (Figure S6b).\na b\nEMC EMC\nFEC\nFEC\nu.) u.)\na. a.\ny\n(\ny\n(\nn\nsit\nN-H NMTF n\nsit N-H\nANM\ne e\nnt nt\nI I\nH-bond H-bond\nFEC-EMC-NMTF FEC-EMC-ANM\n3450 3400 3350 3300 3250 3450 3400 3350 3300 3250\nWavenumber (cm-1) Wavenumber (cm-1)\nFigure S6. FTIR spectra in the 3250 cm–1 to 3450 cm–1 region using NMTF (a) and\nANM (b) co-solvents (5 mol L–1) added to FEC: EMC = 3 : 7 (v/v) solvent mixture.\nCompared to strong hydrogen bond donors like carboxylic acids, this interaction\nis more tunable, facilitating the construction of stable and dynamically adjustable\nsolvation structures. These hydrogen bonds exhibit excellent directionality and\nsaturation, which not only promotes the formation of well-ordered solvation structures\nbut also improves the coordination environment of anions. Secondly, compared to\nhydrogen bond donors such as HF, alcohols, and carboxylic acids, the designed ANM\n(2-cyano-N-methylacetamide) exhibits a lower highest occupied molecular orbital\n(HOMO) energy level, indicating superior oxidative stability (Figure 1i). This property\nenables ANM to maintain chemical inertness under high-voltage operating conditions,\nmaking it ideal cosolvent for high-voltage lithium metal batteries.\nFigure 1i. The screening of possible co-solvents with HOMO/LUMO energies.\nMoreover, ANM induces the formation of anion-rich, ordered solvation domains,\nwhich facilitates the desolvation process of Li+ and enhances their transport kinetics\n(Figure S18). In contrast, conventional hydrogen bond donors such as alcohols may\nlead to overly compact solvent structures, thereby hindering Li+ mobility.\nIn summary, due to their excellent oxidative stability, moderate and tunable\nhydrogen bonding capability, favorable influence on Li+ transport kinetics, and superior\ncompatibility with high-voltage cathode materials, amide-based hydrogen bond donors\nrepresent an ideal choice for developing high-performance, long-lifespan lithium metal\nbattery electrolytes. Their integrated properties enable stable electrolyte operation\nunder high voltage and significantly enhance the overall electrochemical performance.\na b\nLi-FSI− 730.4 cm−1 Coordinated FEC Coordinated EMC\nE-F@ANM\na.\nu.)\n730.1 cm−1 a.\nu.) E-F@ANM\ny\n(\ny\n(\nsit E-F@NMTF sit\nn n E-F@NMTF\ne e\nnt 727.7 cm−1 nt\nI I\nE-F\nE-F\n640 660 680 700 720 740 760 780 800 840 860 880 920 940 960\nRaman shift (cm−1) Raman shift (cm−1)\nFigure S18. Raman spectra of different electrolytes. (a) Li+ and FSI– coordination (b)\nLi+ and FEC or EMC coordination.\nComment 2. In the manuscript, the authors selected an addition concentration of 5 M/L\nfor ANM or other solvents. Would such a high concentration potentially increase the\nviscosity of the electrolyte and reduce its conductivity? Based on the authors’ in-depth\nanalysis of hydrogen bonding between solvent molecules, it is suggested that a more\nrational discussion regarding the selection of additive concentrations should be\nincluded.\nReply: We sincerely appreciate the insightful comments on our manuscript. First,\nregarding the concern that higher concentrations may increase viscosity and thus reduce\nconductivity, we fully agree that this trend generally holds true in conventional\nelectrolyte systems (Angew. Chem. Int. Ed. 2021, 60, 25624–25638). However, in the\nsolvent system developed in this study, ANM serving as a hydrogen bond donor with a\nunique molecular structure does not strictly follow the conventional dilution\nconductivity decreasing pattern. To systematically evaluate the electrolyte performance\nat different additive concentrations, we present the conductivity variation curves for co-\nsolvent concentrations ranging from 3.5 M to 6.5 M in Figure S10. The experimental\nresults show that conductivity first increases and then decreases with increasing ANM\nconcentration, reaching a peak at 5 M. Therefore, we selected 5 M as the optimal\nconcentration for the main experiments.\na b c\n-1)\nm\n9.0 8.290 -1)\nm\n-1) m\nS\nc7.5\nS\nc9.0 8.380\nS\nc9.0 8.590\nm m m\ny\n(6.0\ny\n(\ny\n(\ncti\nvit\n4.5 cti\nvit7.5\ncti\nvit7.5\nu u u\nd3.0 d d\no\nn\no\nn6.0\no\nn6.0\nc\nc1.5\nc\nc\nc\nc\nni ni ni\nI\no0.0\n0.0 0.5 1.0 1.5 2.0 2.5 3.0 I\no4.5\n3.0 3.5 4.0 4.5 5.0 5.5 6.0 6.5 7.0 I\no4.5\n3.0 3.5 4.0 4.5 5.0 5.5 6.0 6.5 7.0\nConcentration (M) Concentration (M) Concentration (M)\nFigure S10. The conductivity of electrolytes configured with different moles of LiFSI\nin FEC : EMC = 3 : 7 (by volume) (a) and with different moles of NMTF (b) and ANM\n(c) added to 1.2 M LiFSI in FEC : EMC = 3 : 7 (by volume) electrolytes.\nNotably, this observation contradicts the conventional understanding that “higher\nviscosity leads to lower conductivity.” The key to this phenomenon lies in the molecular\ninteraction characteristics of ANM. Through both theoretical and experimental\ninvestigations of intermolecular hydrogen bonding in the solvent system, we found that\nANM facilitates the formation of “anion-rich solvated hydrogen-bonded domains”\nwithin a certain concentration range. Specifically, ANM molecules utilize their N–H\ngroups as hydrogen bond donors to form directional hydrogen bond networks with C=O\ngroups in the solvent. This effectively disrupts the strong Li+–solvent interactions\ncharacteristic of conventional solvent-separated ion pairs, enhances anion participation,\nand lowers the desolvation energy barrier for Li+, thereby improving their migration\nrate. Furthermore, combined molecular dynamics simulations and FTIR/Raman\nspectroscopy analyses confirmed that at a concentration of 5 M, ANM molecules induce\nthe formation of optimal hydrogen-bonded domains and anion-dominated solvation\nstructures. This structure not only effectively modulates viscosity but also significantly\nenhances the dissociation degree of lithium salt and the overall ionic mobility,\nultimately resulting in improved bulk conductivity.\nWhen the concentration exceeds 5 M (e.g., at 5.5 M), we observed excessive cross-\nlinking of hydrogen-bonded domains, which leads to denser local structures and\nrestricted diffusion of free ions, thereby causing a decline in conductivity. Therefore,\nthe 5 M concentration represents not only the optimal point for conductivity but also\nthe best balance between hydrogen bond network order and ion transport efficiency.\nIn summary, the selection of ANM concentration was not arbitrarily determined, but\nrather based on systematic experimental screening, conductivity measurements, and\nmechanistic understanding of hydrogen bonding at the molecular level. This\nconcentration not only avoids conductivity loss but also yields the best electrochemical\nperformance. We are deeply grateful for your insightful question, which prompted us\nto provide a more thorough and robust explanation of this aspect in the manuscript.\nChanges made: In the revised manuscript, we added the text: (Page 9, line 21)\nNotably, the conductivity of the electrolyte exhibits a non-monotonic dependence on\nthe concentration of ANM. Although an increase in co-solvent content typically leads\nto higher viscosity, the introduction of 5 M ANM results in the highest ionic\nconductivity among all tested concentrations. This is attributed to the formation of the\nhydrogen-bonded domains, which enhance Li+ transport kinetics despite the increased\nviscosity of the electrolyte.\nComment 3. Generally speaking, molecules containing cyanoethyl groups tend to\nexhibit strong chemical reactivity upon contact with lithium metal. Can the series of\nsolvents (TFAc, ANAc, NNDMTFAc, ANNNAc, NMTF, and ANM) mentioned by the\nauthors remain stable in lithium metal batteries? It is recommended that the authors\nprovide further results and discussion on this point.\nReply: We sincerely appreciate the valuable suggestions. In response to your\nconcern regarding the stability of solvents such as TFAc, ANAc, NNDMTFAc,\nANNNAc, NMTF, and ANM upon contact with lithium metal, we have supplemented\nour study with additional results and presented a more comprehensive analysis.\nThe E-F electrolyte system demonstrates excellent stability in Li||Li and Li||Cu\nbatteries. We believe that the stability of ANM solvent does not stem from the inherent\ninertness of their functional groups but rather from their ability to form adaptive\nhydrogen-bonded domains with shielding effects in specific co-solvent environments,\nparticularly when coexisting with FEC and EMC. These hydrogen-bonded domains\nexhibit high spatial selectivity and temporal evolution characteristics, enabling them to\nundergo “adaptive” restructuring in response to changes in solvent distribution, electric\nfield strength, and lithium salt concentration within the system. Within the hydrogen-\nbonded domains, the cyano groups of ANM adopt a spatial configuration that positions\nthem away from the lithium metal surface. This hydrogen-bond-guided molecular\narrangement not only stabilizes the solvation structure of the electrolyte but also\neffectively blocks electron leakage pathways, thereby suppressing reductive\ndecomposition and parasitic side reactions. Consequently, undesirable side reactions at\nthe lithium metal surface are effectively mitigated. Experimental evidence from in situ\nRaman spectroscopy supports this hypothesis: in systems containing ANM, the SEI\nlayer predominantly consists of inorganic products such as LiF and Li N, indicating\n3\nthat the decomposition of organic solvents is effectively suppressed during interfacial\nformation. This stable interfacial composition arises from the spatially regulated effect\nof the adaptive hydrogen-bonded domains. This system not only alleviates the reactivity\nissues between cyano functional groups and lithium metal but also significantly\npromotes the formation of a stable and high-performance electrochemical interface.\nHerein, we have conducted additional experiments to further validate the\nelectrolyte’s stability. First, we tested the Li||Cu cells between 0~1 V at a current density\nof 0.2 mA cm–2 to monitor voltage changes, without plating Li (Figure S43a). It is\nnoteworthy that during the initial discharge to 0 V—corresponding to the first SEI\nformation process—the E-F@ANM electrolyte exhibited the shortest completion time,\nindicating the fastest passivation capability and superior resistance to reduction.\nFurthermore, a 24 h resting period was introduced during the cycling of Li||Cu cells to\nevaluate the corrosiveness of the electrolytes toward deposited Li (Figure S43b). The\nresults reveal that both E-F@ANM and E-F@NMTF exhibited reduced corrosiveness\ntoward Li, retaining a greater amount of reversible Li metal after aging.\nWe evaluated the long-term chemical stability between cyanoethyl electrolytes and\nlithium metal by immersing lithium metal in electrolytes supplemented with various\nnitrile amides (5 M) in E-F. After storage at 25 oC for two weeks, no noticeable side\nreactions were observed. These experiments further reinforce our conclusions, for\nwhich we seek your recognition.\nChanges made: In the revised manuscript, we added the text: (Page 7, line 16)\nMoreover, molecules bearing a −CH CN group exhibited long-term chemical stability\n2\ntoward both the electrolyte and lithium metal, and no appreciable gas evolution or\ndiscoloration was observed (Figure S3).\n(Page 21, line 1)\nThe reductive stability of various electrolytes was evaluated in Li||Cu cells by\nmonitoring the voltage response over the 0~1 V window at a current density of 0.2 mA\ncm−2 (Figure S43a). Notably, during the initial discharge to 0 V, corresponding to first-\ncycle SEI formation, the E-F@ANM system completed the process in the shortest time,\nindicating the fastest passivation kinetics and superior reductive stability. In addition, a\n24 h rest period was introduced during Li||Cu cycling to assess electrolyte corrosivity\ntoward deposited lithium (Figure S43b). The results showed that both E-F@ANM and\nE-F@NMTF mitigated lithium corrosion and retained more reversible active lithium\nafter aging.\nIn the revised supporting information, we added the figures：\nFigure S43. (a) Galvanostatic tests of the Li||Cu cells between 0-1 V at a current density\nof 0.2 mA cm–2. (b) LMA CE under full plating/stripping conditions at 1 mA cm–2 and\n1 mAh cm–2, with a 24-hour aging period after Li deposition.\nFigure S3. Optical photographs of fresh electrolytes (top), and electrolytes (middle)\nand lithium metals (bottom) after immersing lithium metals in electrolytes for two\nweeks.\nComment 4. The image in Figure 4 is not cross-referenced properly. Please check it.\nIn addition, the meaning of some abbreviations (e.g, TFAc, ANAc, NNDMTFAc,\nANNNAc, NMTF, ANM) should be clearly defined.\nReply: Thanks for pointing out this issue and apologize for our carelessness.\nRegarding the cross-referencing issue in Figure 4, we acknowledge that the original\nversion lacked clarity in the citation location. Furthermore, we fully agree with your\nsuggestion regarding the definition of abbreviations. Indeed, several terms were not\nclearly defined upon their first appearance, which may hinder reader comprehension.\nTo address this, we have implemented the following corrections in the revised\nmanuscript: detailed definitions of all abbreviations have been provided in Figure S1.\nFormamide Acetamide Trifluoroacetamide Cyanoacetamide\n(FM) (Ac) (TFAc) (ANAc)\nN-Methylacetamide N-N- N,N- 2-Cyano-N,N-\n(NMAc) Dimethylacetamide Dimethyltrifluoroacetamide dimethylacetamide\n(NNDMAc) (NNDMTFAc) (ANNNAc)\nN-Methyltrifluoroacetamide 2-Cyano-N-methyl-acetamide\n(NMTF) (ANM)\nFigure S1. Structural formulas of all the amide cosolvents used in the experiments as\nwell as the abbreviations.\nChanges made: In the revised manuscript, we added the text: (Page 7, line 11)\nSubstituting the hydrogen atom in the amide’s –NH group with a methyl group (e.g.,\nresistance. Further modifications, such as replacing the –NH group with isopropyl (–\n3 3 2\nNNDMTFAc, ANNNAc, NMTF, and ANM; full names provided in Figure S1),\nimproved compatibility with lithium metal electrodes by increasing polarity.\n(Page 19, line 7)\nAdditionally, the F 1s and N 1s spectra display higher Li N and LiF signals in Li metal\n3\nelectrodes cycled with the E-F@NMTF and E-F@ANM electrolytes compared to those\ncycled with the E-F electrolyte (Figure 4d-g), aligning with trends observed in FSI⁻\ndecomposition (Figure S35a-b).\nComment 5. Why E-F was selected as the baseline electrolyte. Can the design of\nhydrogen bonding strategy be feasible in other electrolyte systems? Besides, why the\naddition of NMTF and ANM resulted in an enhancement of electrical conductivity\n(Lines 6-8, Page 8)?\nReply: We sincerely appreciate the valuable comments on our manuscript. First,\nwe selected E-F as the baseline electrolyte due to its wide electrochemical window and\nexcellent stability, particularly under high-voltage conditions. In the previous work\n(Small 2024, 20, 2311961), E-F demonstrated superior ionic conductivity and enhanced\nstability compared to conventional EC-based electrolytes, especially in high-voltage\nenvironments. The upper limit of an electrolyte’s electrochemical window is primarily\ndetermined by its most oxidation-susceptible component. In EC-based electrolytes, the\nstrong Li+–solvent interactions promote the formation of solvent-anion clusters, within\nwhich the anions tend to transfer partial negative charges to the solvent molecules. This\nfacilitates charge transfer and triggers oxidative decomposition reactions (Energy\nEnviron. Sci. 2022, 15, 434904361). Therefore, we employed a FEC-based electrolyte,\nas free FEC exhibits higher oxidative stability and weaker Li+-solvent coordination.\nThis strategy allows us to maximize the high-voltage tolerance of the base electrolyte\n(E-F), providing a more robust foundation for the subsequent incorporation of co-\nsolvents.\nWe have added control experiments based on EC : EMC = 1:1 (v/v) containing\nNMTF or ANM (5 M L⁻1), the N–H stretching vibration in the FTIR spectra exhibited\na pronounced blue shift, unequivocally confirming the presence of hydrogen-bond\ninteractions (Figure S12a-b). These features were absent in the pure EC/EMC system,\nindicating that the observed spectral changes originate from the hydrogen-bond donor\ncapability of NMTF and ANM. Peak-fitting analysis of the Raman spectra in the PF ⁻\n6\nvibration region revealed that both NMTF and ANM decreased the proportion of\nsolvent-separated ion pairs (SSIPs) while increasing the fractions of contact ion pairs\n(CIPs) and aggregates (AGGs), indicating that hydrogen-bonded domains promote\ntighter Li+–anion association (Figure S13a-c). This observation aligns with the\nmechanism proposed in the manuscript, wherein hydrogen bonding weakens solvent–\nLi+ interactions and drives anions into the primary solvation shell. In NCM811||Li cells\noperated at 4.5 V, those employing the E-E@ANM electrolyte achieved a capacity\nretention of 93.3% after 200 cycles, markedly surpassing that of cells with E-E\nelectrolyte (63.6%) and E-E@NMTF electrolyte (75.7%) (Figure S14). This trend is\nconsistent with the pattern observed in the E-F system, and the magnitude of\nperformance enhancement correlates positively with the hydrogen-bond strength of the\nadditive, further substantiating the pivotal role of hydrogen-bonded domains.\nMoreover, the addition of NMTF and ANM improves the solvation structure of\nthe electrolyte by establishing hydrogen-bonded domains. In particular, ANM can form\nstrong H-bonds and Z-bonds, disrupting the loose solvent-ion pairs typically observed\nin conventional solvation structures. This facilitates the formation of tighter anion–Li+\ncoordination, thereby enhancing the electrolyte’s overall ionic conductivity. These\neffects were confirmed by both conductivity measurements and molecular dynamics\nsimulations, which revealed significant improvements in ionic diffusion coefficients\nand conductivity for the E-F@ANM system.\nOnce again, we sincerely thank you for your insightful feedback, which has\nenabled us to elaborate more comprehensively on this key issue and further refine the\nmanuscript.\nChanges made: In the revised manuscript, we added the text: (Page 10, line 1)\nSimilarly, such H-bond interactions were operative across different carbonate solvents.\nIn the FTIR spectra of the EC-EMC-NMTF and EC-EMC-ANM mixed solvents, the\nN–H stretching vibration showed a pronounced blue shift, unambiguously confirming\nH-bond (Figure S12a-b). Moreover, deconvolution of the PF – Raman bands indicated\n6\nthat introducing NMTF or ANM decreased the fraction of solvent-separated ion pairs\n(SSIP) while increasing the fractions of CIP and AGG, suggesting that H-bond\npromoted tighter Li+–anion association (Figure S13a-c). In Li‖NCM811 cells\nemploying the E–E@ANM electrolyte (Table S3), the capacity was retained at 93.3%\nafter 200 cycles, significantly outperforming cells using the E-E electrolyte (Figure\nS14).\nIn the revised supporting information, we added the figures：\nFigure S12. FTIR spectra region using NMTF (a) and ANM (b) co-solvents (5 mol\nL⁻1) added to EC: EMC=1 : 1 (v/v) solvent mixture.\nFigure S13. The Raman spectra of (a) EC-EMC (E-E), (b) E-E@NMTF, and (c) E-\nE@ANM electrolytes were obtained, and the proportions of the three solvated\nstructures were determined based on peak area integration.\nFigure S14. The long cycle performance of Li||NCM811 coin cells at 2.8-4.5 V with\ndifferent electrolytes.\nComment 6. For the pouch cells used in this article, more detailed technical parameters\nneed to be provided.\nReply: Thanks for your valuable suggestions. We have added detailed\nspecification parameters of pouch cells on page 98 in the revised supporting\ninformation.\nChanges made: In the revised manuscript, we added the text: (Page 30, line 23)\nFurthermore, we assembled lithium metal pouch cells under harsh real-world\nconditions. To maximize the energy density, we used a highly loaded anode (20 mg\ncm−2), a lean electrolyte (2.4 g Ah−1), and a thin lithium anode (80 μm, N/P ratio of 1.7).\nHere, the energy density of the 1.6 Ah pouch cell with the E-F@ANM electrolyte was\nas high as 418.2 Wh kg−1 (Figure 7e, Figure S80 and Table S8).\nIn the revised supporting information, we added the table：\nTable S8. The specifications of the Li||NCM811 pouch cell.\nCell component Specification Parameters\nActive material mass loading (mg cm–2, each side) 20\nCathode (NCM811\nNumber 10\nwith Al current\ncollector) Specific areal capacity (mAh cm–2, each side) 4\nWeight (g) 8.243\nNumber 11\nAnode (Li) Thickness (μm) 80\nWeight (g) 0.637\nElectrolyte/Capacity (g Ah–1) 2.4\nElectrolyte\nWeight (g) 3.463\nSeparator (PE) Weight (g) 0.480\nPackage and tabs Weight (g) 1.710\nSize (cm × cm) 6 × 5\nN/P 1.7\nDischarge capacity (Ah) 1.6\nFull cell Mid-value voltage (V) 3.8\nDischarge energy (Wh) 6.08\nTotal weight (g) 14.533\nSpecific energy (Wh kg–1) 418.2\nComment 7. To demonstrate the positive effectiveness of E-F@ANM induced CEI on\ncrystal structure, experiments of NMC811 cathodes with E-F and E-F@ANM derived\nCEI should be tested and compared in E-F system (e.g., XRD of such cathodes after\nprolonged cycling).\nReply: We sincerely appreciate your valuable comments on our manuscript. We\nacknowledge that directly validating the CEI’s contribution to cathode stability at the\nstructural level would significantly strengthen the manuscript We consider synchrotron\nX-ray diffraction (SXRD) to offer significant advantages over conventional XRD,\nincluding higher source brightness and superior spatial and temporal resolution. These\nfeatures enable the acquisition of high signal-to-noise ratio data within shorter\ntimescales, making SXRD particularly suitable for high-precision, in situ or quasi-in\nsitu investigations of crystalline structure evolution in trace samples, complex\ninterfaces, or battery materials during cycling (Nat. Mater. 2020, 19, 1029–1036; J. Am.\nChem. Soc. 2021, 143, 19614–19627). Moreover, the tunable wavelength and\nexceptional detection sensitivity of SXRD enable the identification of subtle\nmicrostructural changes that are often beyond the reach of conventional XRD. This\nmakes SXRD especially well-suited for investigating phase transitions and lattice\ndistortions in cathode materials such as NMC during cycling (Chem. Mater. 2022, 34,\n4223–4234).\nAccordingly, we conducted SXRD measurements on NCM811 electrodes after\n200 cycles (Figures 5g-h and S51). Rietveld refinement of the SXRD data was\nperformed in GSAS-II using a fixed rhombohedral R-3m phase model. The results\nindicate a lower Li/Ni cation mixing value for the E-F@ANM system (2.58%)\ncompared to the E-F system (3.73%) (Tables S4-S6). After cycling in the conventional\nE-F electrolyte, the NMC811 electrodes exhibit a characteristic rightward shift and\nintensity attenuation of the (003) peak, indicative of partial collapse of the layered\nstructure and lattice distortion. In contrast, the NMC811 electrodes cycled in the E-\nF@ANM electrolyte show (003), (101), and other characteristic peaks that remain close\nto those of fresh electrodes, with sharper peak profiles. This suggests that the layered\nstructure is better preserved and lattice distortion is significantly mitigated. These\nresults confirm that the CEI induced by the anion-rich solvation hydrogen-bonded\ndomains constructed in the E-F@ANM electrolyte exerts a positive effect in\nsuppressing structural degradation under high-voltage conditions.\nFigure 5. Rietveld refinement results for SXRD of NCM811 after 200 cycles using E-\nF (g) and E-F@ANM (h) electrolytes. (The single-phase model used by Rietveld was\nrefined and used to significantly explore the (003) characteristic peak as well as the\n(104) characteristic peak change.)\nFigure S51. Rietveld refinement results for Synchrotron X-ray diffraction (SXRD) of\nfresh NCM811 X-ray diffraction patterns.\nWe sincerely thank you for your insightful feedback, which has enabled us to elaborate\nmore comprehensively on this key issue and further refine the manuscript.\nReviewer #3 (Remarks to the Author):\nComments:\nThe authors present a compelling strategy for modulating the solvation structure of\nhigh-voltage electrolytes through the deliberate exploitation of intermolecular\nhydrogen bonding. By combining advanced theoretical simulations with systematic\nexperimental characterization, they conclusively demonstrate that hydrogen-bonding\ninteractions significantly enhance electrolyte stability at extreme potentials (>4.5 V vs.\nLi+/Li), promote favorable desolvation kinetics, and stabilize critical electrode-\nelectrolyte interfaces. This rational design enables the stable operation of high-energy-\ndensity lithium metal pouch cells at an unprecedented 4.7 V cutoff voltage. The work\nprovides fundamental insights into electrolyte engineering and represents a significant\nadvancement toward next-generation battery systems. The findings are of broad interest\nto the energy storage community and warrant publication after addressing the following\nrevisions.\nReply: We extend our sincere gratitude to the reviewer’s diligent evaluation of our\nfindings. We are very grateful to the reviewer for highlighting the novelty of our single\nsolvent electrolyte concepts and recognizing our comprehensive analysis of the\nelectrolyte mechanism. We have added additional experimental results and figures to\nimprove the discussion.\nComment 1. The author proposes in Figure 1 that the atypical hydrogen bonds of ANM,\nas revealed by 1HNMR and theoretical calculations, further enhance the effects of H-\nbond. Could the magnitude of the Z-bond (Nδ⁻–Hδ+) be calculated to better illustrate its\nstrength relative to typical hydrogen bonds?\nReply: Thanks for your valuable suggestions. We conducted DFT calculations on\nthe Z-bond (Nδ⁻–Hδ+) formed between ANM and EMC in a mixed solvent solution\nsystem. In this system, the interaction energy of the Z-bond (C≡N···H–C) in ANM-\nEMC is –2.5 kcal·mol⁻1, while the H-bond N–H···O=C is –10.8 kcal·mol⁻1 for ANM-\nEMC, –7.4 kcal·mol⁻1 for ANM–FEC, –8.8 kcal·mol⁻1 for NMTF–EMC, and –7.6\nkcal·mol⁻1 for NMTF–FEC; the weak hydrogen bond N–H···F–C is –4.4 kcal·mol⁻1\n(Table S1). These values indicate that the Z-bond is approximately 3–5 times weaker\nthan the H-bond, which aligns perfectly with our qualitative assessment in the main text\nthat “the Z-bond is weaker”. Therefore, the Z-bond alone is not sufficient to\nsignificantly alter the solvent structure of the electrolyte; its main role is to enhance and\nstabilize the surrounding primary H-bond interactions through electronic\npolarization/orientation constraints, thus forming the coupled structural domains of Z-\nbond and H-bond presented in our study, leading to the formation of a special solvation\nstructure with an “anion-rich inner and outer layer hydrogen-bonded domains.”.\nChanges made: In the revised manuscript, we added the text: (Page 8, line 24)\nDFT calculations confirmed the formation of Z-bond (–2.5 kcal mol⁻1) (Table S1). In\naddition, the hydrogen bonds between NMTF/ANM and the (C–F) sites in FEC were\ncomparatively weak, owing to the greater steric hindrance of (C–F) and its lower\nelectron density relative to (C=O).\nIn the revised supporting information, we added the table：\nTable S1. The magnitudes of the interactions between H-bonds, Z-bonds, and van der\nWaals forces observed in this study.\nForce Magnitude\nComplexes Bond Type Types of Forces\n(kcal mol⁻1)\nANM-EMC N–H∙∙∙O=C H-bond –10.8\nANM-FEC N–H∙∙∙O=C H-bond –7.4\nNMTF -EMC N–H∙∙∙O=C H-bond –8.8\nNMTF-FEC N–H∙∙∙O=C H-bond –7.6\nNMTF-FEC N–H∙∙∙F-C H-bond –4.4\nEMC-FEC / Van der Waals force –6.9\nANM-EMC C≡N∙∙∙H–C Z-bond –2.5\nComment 2. The author conducted float charging tests on Li||NCM811 button cells to\ndemonstrate the high-voltage stability of the electrolyte, but I would still like to know\nif there are other experimental techniques that can further verify the reduction\nbehaviour of this electrolyte. This would help to evaluate the stability of the electrolyte\nmore comprehensively.\nReply: We sincerely appreciate the valuable comments on our manuscript. We\nhave conducted additional experiments to verify the redox stability of the electrolyte,\nand we hope to obtain your approval. Regarding the reduction stability, we have already\nincorporated the average Coulombic efficiency of Li||Cu cells (Figure 4h), calculated\nthe Tafel slope (Figure S38), and conducted electrochemical performance tests in the\noriginal manuscript. Here, we have conducted additional experiments to further\nvalidate the electrolyte’s stability. First, we tested the Li||Cu cells between 0~1 V at a\ncurrent density of 0.2 mA cm–2 to monitor voltage changes, without plating Li (Figure\nS43a). It is noteworthy that during the initial discharge to 0 V—corresponding to the\nfirst SEI formation process—the E-F@ANM electrolyte exhibited the shortest\ncompletion time, indicating the fastest passivation capability and superior resistance to\nreduction. Furthermore, a 12 h resting period was introduced during the cycling of\nLi||Cu cells to evaluate the corrosiveness of the electrolytes toward deposited Li\n(Figure S43b). The results reveal that both E-F@ANM and E-F@NMTF exhibited\nreduced corrosiveness toward Li, retaining a greater amount of reversible Li metal after\naging. These experiments further reinforce our conclusions, for which we seek your\nrecognition.\nChanges made: In the revised manuscript, we added the text: (Page 21, line 1)\nThe reductive stability of various electrolytes was evaluated in Li||Cu cells by\nmonitoring the voltage response over the 0~1 V window at a current density of 0.2 mA\ncm−2 (Figure S43a). Notably, during the initial discharge to 0 V, corresponding to first-\ncycle SEI formation, the E-F@ANM system completed the process in the shortest time,\nindicating the fastest passivation kinetics and superior reductive stability. In addition, a\n24 h rest period was introduced during Li||Cu cycling to assess electrolyte corrosivity\ntoward deposited lithium (Figure S43b). The results showed that both E-F@ANM and\nE-F@NMTF mitigated lithium corrosion and retained more reversible active lithium\nafter aging.\nIn the revised supporting information, we added the figure：\nFigure S43. (a) Galvanostatic tests of the Li||Cu cells between 0-1 V at a current density\nof 0.2 mA cm–2. (b) LMA CE under full plating/stripping conditions at 1 mA cm–2 and\n1 mAh cm–2, with a 24-hour aging period after Li deposition.\nComment 3. The author performed activation energy fitting of the Nyquist plot at (30\noC~60 oC) in Figure S34 and suggested supplementing the impedance with a wider\ntemperature range to more convincingly prove the kinetic phenomenon.\nReply: We sincerely appreciate the valuable comments on our manuscript. We\nhave expanded the impedance spectroscopy analysis to a broader temperature range (–\n20 oC~70 oC) in the revised manuscript. The resulting Nyquist plots and fitted data yield\nmore comprehensive and robust kinetic parameters, further substantiating the crucial\nrole of hydrogen-bonded domains in reducing interfacial energy barriers and enhancing\nlithium-ion transport efficiency. These updated results are presented in Figure S39a-f\nof the Supporting Information. We used an equivalent circuit model (R -R -R series\nb SEI ct\nfitting model, see Figure S38) to fit the Nyquist plot of the Li||Li symmetric battery,\nand extracted the activation energies of R and R under temperature dependence\nct SEI\nthrough the Arrhenius equation ln(R) = ln(A) + E /RT.\na\nThis expanded set of experiments further validates that the adaptive hydrogen-\nbonded domains induced by ANM effectively lower the desolvation energy barrier at\nthe interface and improve overall interfacial kinetic behavior. The underlying\nmechanism can be attributed to the weakened Li+–solvent interactions due to H-bond\nand Z-bond modulation, which in turn promotes a higher fraction of anion-dominated\nsolvation structures (AGG/CIP). This results in a reduced desolvation energy barrier\nand enhanced reaction kinetics. Once again, we thank you for your insightful and\nconstructive suggestions.\nChanges made: In the revised manuscript, we added the text: (Page 20, line 2)\nAs shown in Figure S39a-f, the E of the E-F@NMTF and E-F@ANM are 61.90 kJ\na,Rct\nmol−1 and 56.53 kJ mol−1, respectively, which are much lower than that of E-F (70.78\nkJ mol−1), indicating that the adaptive hydrogen-bonded domains anchoring strategy\nfacilitates enhanced Li+ desolvation. Meanwhile, the interfacial Li+ diffusion barrier,\nrepresented by E decreased from 49.42 kJ mol−1 in E-F to 43.14 kJ mol−1 in E-\na (RSEI)\nF@NMTF and 38.07 kJ mol−1 in E-F@ANM, which was attributed to the formation of\nan inorganic component-rich SEI with a lower ion mobility barrier. The E-F@NMTF\nand E-F@ANM were more efficient for lithium plating and stripping due to their anion-\nrich solvated structures.\nIn the revised supporting information, we added the figure：\nFigure S39. Nyquist plots of symmetric cells based on E-F (a), E-F@NMTF (b) and E-\nF@ANM (c) electrolytes at different temperatures (–20~70 oC) and calculated\nactivation energies of R and R for E-F (d), E-F@NMTF (e) and E-F@ANM (f).\nct SEI\nComment 4. The author demonstrates the effect of hydrogen bonding domains on\nlithium metal anodes in Figure 4 and suggests supplementing the in situ EIS of Li||Li\nsymmetric cells to prove the improvement of the electrolyte on the kinetics of lithium\nmetal anodes.\nReply: We sincerely appreciate your valuable comments on our manuscript. In\nsitu EIS measurements of Li||Li symmetric cells revealed a sharp decrease in R during\nct\nthe initial cycles in the E-F electrolyte, indicating non-uniform lithium deposition and\ncontinuous degradation of the physicochemical properties at the electrode–electrolyte\ninterface (Figure S48a-c). By contrast, the E-F@NMTF and E-F@ANM electrolytes\nmaintained stable R values during the initial cycles, attributable to the formation of a\nct\nuniform SEI, which ensured long-term interfacial stability. Moreover, the\ndecomposition of FSI⁻ at the electrode interface promoted the formation of a dense and\nuniform inorganic SEI. This suppressed interfacial side reactions, enhanced uniform\nlithium deposition, and ultimately yielded a highly reversible lithium anode. In addition,\ndeconvolution of the high- and low-frequency components of the EIS spectra yielded\nthe distribution of relaxation times (DRT). The DRT analysis decomposed the overall\nimpedance spectrum into six distinct relaxation-time peaks (p1~p6): p1 corresponds to\nlow contact resistance (R ); p2 and p3 to SEI resistance (R ); p4 and p5 to charge-\nc SEI\ntransfer resistance (R ); and p6 to ohmic resistance (R ) (J. Power Sources 2019, 425,\nct b\n121-129). The large p2 and p3 features observed in the E-F electrolyte indicate\ninstability in the electrode-electrolyte interfacial properties. In contrast, the negligible\nR and R in the E-F@ANM electrolyte suggest a faster Li+ desolvation process and\nct SEI\nmore facile Li+ transport through the SEI. The charge-transfer process at the electrode\ninterface was markedly enhanced through dynamic Li+–solvent coordination.\nChanges made: In the revised manuscript, we added the text: (Page 21, line 19)\nThe in situ EIS results for the Li||Li symmetric cell indicated that during the initial\ncycles of the E-F system, the R value remained in a relatively high range, suggesting\nct\nuneven lithium deposition and ongoing degradation at the electrode–electrolyte\ninterface (Figure S48a). In contrast, the E-F@NMTF and E-F@ANM systems\npromoted rapid stabilization of R during the initial cycles, which was attributed to the\nct\nformation of a more stable SEI, thereby ensuring long-term interface stability (Figure\nS48b-c). Additionally, by decomposing the high- and low-frequency components of the\nEIS, the relaxation time distribution (DRT) was obtained. The DRT analysis\ndecomposed the overall impedance into six distinct relaxation time peaks (p1~p6): p1\ncorresponded to low contact resistance (R ); p2 and p3 corresponded to SEI resistance\nc\n(R ); p4 and p5 corresponded to charge transfer resistance (R ); and p6 corresponded\nSEI ct\nto ohmic resistance (R ). The significant p2 and p3 features observed in the E-F system\nb\nindicated instability in the electrode–electrolyte interface properties. In contrast, the\nnearly negligible R and R in the E-F@ANM system suggested a faster Li+\nct SEI\ndesolvation process and faster Li+ transport through the SEI.\nIn the revised supporting information, we added the figure：\nFigure S48. In situ impedance spectroscopy and DRT of Li||Li symmetric cells in E-F\n(a), E-F@NMTF (b) and E-F@ANM (c) electrolytes.\nComment 5. The author mentioned the dissolution phenomenon of transition metals,\nbut Figure S50 lacks supporting data on Co element leaching. It is recommended that\nthe author supplement the data to prove the positive effect of ANM co-solvent in\nimproving the voltage resistance performance of the electrolyte.\nReply: We sincerely appreciate the valuable feedback on our manuscript. We have\nadded relevant experimental data on the dissolution of Co elements. In the revised\nSupporting Information (Figure S58), we performed a quantitative analysis of the\ndissolved transition metal elements in different electrolytes using ICP-MS. Notably,\nwhen using the E-F@ANM system, we observed a significant reduction in the\ndissolution of cobalt (Co) from the NCM811 electrode during cycling. Specifically,\ncompared to the conventional E-F electrolyte, the dissolution of cobalt in the E-\nF@ANM system was reduced by approximately 50%. This result suggests that the\nANM cosolvent enhances the interfacial stability of NCM811, reducing the dissolution\nof transition metals and slowing down side reactions in the electrolyte at the cathode,\nthereby effectively protecting the structure of the NCM811 electrode under high voltage.\nWe have expanded and strengthened the above analysis in both the revised manuscript\nand Supporting Information, explicitly highlighting the important role of the ANM\ncosolvent in the electrolyte, especially in improving the dissolution of cobalt transition\nmetals and electrolyte stability under high voltage.\nChanges made: In the revised manuscript, we added the text: (Page 25, line 18)\nIn addition, ICP-MS of the post-cycled electrolytes showed that the contents of TMEs\nin the E-F@ANM system were much lower than those in the E-F system (Figure S58).\nIn the revised supporting information, we added the figure：\nFigure S58. Transition metal ion (TM) content on the diaphragm side after 200 cycles\nin three different electrolytes using ICP-MS technology."}, {"round": 2, "reviewer_comments": {"reviewer_1": "The authors have addressed the comments well, and I think it can be considered for acceptance.", "reviewer_2": "After revisions, this work now can be considered to be accepted and published by Nature Communications."}, "author_response": ""}]

Decision letter