Customized composition of lithium metal solid-electrolyte interphase by electric field modulation of anion motion direction

Domain: neuroscience
Published: 2026-02-04
Editorial outcome: major_revision

Abstract
In high-voltage lithium-metal battery systems, electrolyte characteristics are crucial for achieving the optimal balance between non-flammability and battery performance. Herein, an electrolyte system based on triethyl phosphate solvent with lithium oxide difluoroborate, lithium tetrafluoroborate and lithium nitrate as solutes is proposed. TEP, as the sole solvent, ensures the inherent non-flammability of the electrolyte, while the solutes LiODFB, LiBF4 and LiNO3, benefit from the different binding energies of the three anions and lithium ions, to optimize the structure of the solvation shell and direction of anion movement, thereby forming a favorable interfacial phase. Consequently, solid electrolyte interphase enriched with B-O and Li3N in the inner layer and LiF in the outer layer is formed, which improves the stability and reversibility of the lithium metal negative electrode. The Li||NCM811 cell with the as-prepared electrolyte can be stably cycled for 600 cycles at a high cut-off voltage of 4.5 V with a capacity retention of 90.19%. Even at 60 °C for 600 cycles, the capacity retention rate remains at 81.18%. This work demonstrates an effective strategy for the design of high-voltage non-flammable LMBs.

Introduction
Lithium-ion batteries (LIBs) are experiencing a surge in demand for portable electronics, electric vehicles, and grid storage applications. Currently, LIBs based on LiCoO2/LiMnO2/LiFePO4 as a positive electrode and graphite as a negative electrode are approaching the upper limit of specific energy (~300 Wh kg−1). Employing high-energy-density positive electrode materials, such as NCM (LiNixMn1-xO2, M=Mn, Co, x > 0.8), in conjunction with lithium metal negative electrodes is one of the most effective avenues to achieve greater specific energy (500 Wh kg−1). However, traditional carbonate-based electrolytes are not suitable for high-voltage lithium-metal batteries (LMBs) due to their poor thermal stability, flammability, and incompatibility with highly reactive lithium negative electrodes and high-voltage positive electrodes. Therefore, the development of electrolytes that can meet the requirements of high specific energy and high safety is of vital importance and urgency.

In recent years, organophosphorus-based electrolytes (e.g., trimethyl phosphate (TMP), triethyl phosphate (TEP), and dimethyl methylphosphonate (DMMP)) have emerged as promising candidates for LMBs due to their wide electrochemical window, temperature operating range, low viscosity, and solubility. Nevertheless, in previous reports, batteries assembled with phosphate ester-based electrolytes have demonstrated relatively short cycle life and limited cut-off voltage. This limitation is primarily attributed to the incompatibility between organic phosphates with metallic lithium, which impedes the formation of a stable solid electrolyte interphase (SEI) on the lithium metal surface. Mixed salt system electrolytes with rich anionic functional groups can enhance the stability of the interface layer by adjusting the composition of the interface layer, which offers a promising solution to the aforementioned issues. From a commercial perspective, a non-flammable electrolyte comprising solely a phosphate ester as a solvent and a suitable quantity of mixed lithium salts, offering minimal costs, stable physicochemical properties, and electrochemical stabilities, would represent an optimal electrolyte system for realizing high-voltage LMBs.

Herein, an electrolyte system is proposed to construct the optimal interfacial phase on the surface of lithium metal. The system employs TEP as the sole solvent to ensure the intrinsic non-flammable property of the electrolyte, while lithium oxide difluoroborate (LiODFB), Lithium tetrafluoroborate (LiBF4), and lithium nitrate (LiNO3) serve as solutes to facilitate the formation of an interfacial phase. Notably, the distinct binding energies between lithium salt anions and lithium ions (Li+) have contributed to the optimization of the solvation shell structure. Owing to the strong binding energy of ODFB− and NO3− anions with Li+, they will migrate along the direction of the electric field, aggregate on the lithium metal surface, and enter the interior of the solvation shell, which will preferentially decompose to form the inner interfacial layer of B–O and Li3N. Since the binding energy of the BF4− anion to Li+ is relatively weak, it moves away from the lithium metal surface and is distributed in the outer layer of the solvation shell structure, resulting in a LiF-rich outer layer of SEI. More importantly, the formed three-phase interfacial layer serve distinct functions. LiF enhances the hardness and stability of the interfacial layer, while Li3N boosts the ionic conductivity and facilitates ion transport. The B–O species, which is a glassy component with low crystallinity, facilitates the homogeneity of the ionic flux and improves the flexibility of the SEI. The Li||NCM811 full cell assembled with the electrolyte system can stably cycle 600 times at a high cut-off voltage of 4.5 V with a capacity retention of 90.19%. Even after 600 cycles at 60 °C, the capacity retention remains at 81.18%. Moreover, the Li||NCM811 pouch cell achieved a high specific energy (based on the mass of the whole cell) of 430.51 Wh kg−1, representing one of the highest levels among lithium metal batteries utilizing non-flammable electrolytes.

Results
Design principle of the three-phase interfacial layer
To comprehensively evaluate the electrochemical stability and ionic conductivity of electrolytes, linear sweep voltammetry (LSV) and electrochemical impedance spectroscopy (EIS) were used to systematically screen electrolytes prepared by different electrolytes, salt additions, and chemical formulations. Initially, a comprehensive evaluation of electrolyte systems with a total ion concentration of 1.2 M and varying molar ratios of lithium salts was conducted. The electrolyte system with a lithium salt molar ratio of 1:1:1 exhibited the highest ionic conductivity and the widest electrochemical window, as shown in Fig. S1a, b. Subsequently, further investigations were carried out on electrolyte systems with different concentrations of lithium salts at the 1:1:1 molar ratio, with a particular focus on their Li+ conductivity. Interestingly, comparisons revealed that the 0.5 M-DBN electrolyte (a concentration of 0.5 M for each of LiODFB, LiBF4, and LiNO3, denoted as 0.5 M-DBN) achieved the highest ionic conductivity (Fig. S1c, d). Therefore, the 0.5 M-DBN electrolyte was selected as the subject for further study.

After identifying a relatively suitable additive, further investigations into the decomposition mechanisms of the three lithium salts in the 0.5 M-DBN electrolyte were conducted using all-atom molecular dynamics (MD) simulations under a specific electric field condition. Figure 1a, b illustrates the final atomic distributions at the interface of lithium metal and electrolyte in the presence and absence of an electric field. In general, the presence of an electric field induces positively charged ions or particles to migrate in the same direction as the applied field (Fig. S13), whereas negatively charged ions typically migrate in the opposite direction. Different from expectations but interesting, opposite results can be observed. As illustrated in Fig. 1c, d, the distributions of ODFB− and NO3− anions migrate towards the negative electrode side in response to an electric field with direction illustrated in the figure. This anomalous migration pattern is supposed to be explained by the strong binding energy between the ODFB− and NO3− anions and Li+. As the anions are dragged by lithium ions moving toward the lithium surface, the probability that these two anions appear on the surface increases (Fig. 1g). Consequently, at the negative electrode side, the ODFB− and NO3− anions are able to accept electrons from the lithium surface, resulting in the formation of a layer of Li3N- and B–O-rich SEIs. In contrast, the BF4− anion moves in the direction opposite to the electric field lines due to its relatively weaker binding energy with Li+ (Fig. 1e). This lower binding energy of BF4− with Li+ implies a reduced solvation energy, and the increased availability of free ions improves the overall reaction kinetics of the electrolyte. Furthermore, the TEP migrates counter to the direction of the electric field, which significantly mitigates the possibility of TEP electrolyte reduction at the negative electrode interface (Fig. 1f). Meanwhile, the DFT calculations (Figs. S14 and 15) indicate that lithium metal exhibits a preferential adsorption toward LiNO3 and LiODFB. Figure 1h and Fig. S16 show that the ODFB− and NO3− anions are strongly clustered around Li+ with radial distribution functions (RDFs) peak positions at 1.95 Å and 1.94 Å, respectively, while the BF4− anion is located at the farthest distance from the lithium-ion with a peak position of 2.72 Å. The solvation structure and synergistic migration mode of the 0.5 M-DBN electrolyte lead to preferential decomposition of LiNO3 and LiODFB on the lithium negative electrode surface, while in the meantime, LiBF4 undergoes partial decomposition, thus forming Li3N- and B–O-rich inner SEIs and LiF-rich outer SEIs. The decomposition energy and order of the ternary salts were further investigated through Density Functional Theory (DFT) calculations. With high electron affinity of LiODFB (−2.58 eV) and LiNO3 (−1.53 eV), a B- and N-rich SEI layer is formed on the lithium metal negative electrode, while LiBF4 has trace decomposition to form a F-rich components for the SEI (Fig. 1j). Consequently, differences in binding energies between lithium ions and anions are able to locate an unequal position of these three anions in the electrolyte with field as well as distribution in the solvation shell, triggering the formation of an inner SEI layer consisting of Li3N and B–O and an outer SEI layer consisting of LiF, which greatly overcome the solid limitation of the obstacle to form a stabilized SEI on the lithium surface by organophosphate esters (Fig. 1j). Meanwhile, experimental verification using other combinations of strong and weak coordinating lithium salts further revealed the same SEI formation pattern, confirming the universality and effectiveness of this strategy (Fig. S17).

Interfacial chemistry on the negative electrode side of lithium metal
The above MD simulation and DFT calculation have obtained the preliminary concept of the electrolyte negative electrode interface, which provides theoretical guidance for the subsequent experimental design. Here, the SEI layer has been explored using in situ Fourier Transform Infrared spectroscopy, X-ray photoelectron spectroscopy (XPS), Atomic Force Microscopy (AFM), and Transmission Electron Microscopy (TEM). To monitor the SEI formation process, in situ infrared spectroscopy was used to study the chemical changes at the interface between the electrolyte and the lithium metal. Figure S18 shows the electrochemical performance recorded during the plating and stripping cycles of a Li||Cu cell in the in situ FTIR test. As shown in Fig. 2a and Fig. S20, during the lithium deposition stage (i.e., the desolvation process), the pronounced depletion peaks at 977.90 cm−1 (P–O–C) and 1255.31 cm−1 (P=O) correspond to the gradual consumption of solvated TEP, while the emerging peak at 958.62 cm−1 (P–O–C) indicates the release of desolvated solvent molecules at the electrode interface. Figure S19 presents the in situ infrared spectra of the 1 M LiPF6/EC–DMC electrolyte system along with the corresponding spectral analysis. Furthermore, by comparing the ratio of characteristic peak areas at the end of lithium deposition and during the stripping stage, we quantitatively analyzed the interfacial behavior of surface anions. Notably, PF6− (which weakly coordinates with Li+) exhibits significantly higher reversibility than ODFB−, indicating that ODFB− is more extensively consumed in interfacial reactions (Fig. 2b). These results are consistent with the theoretical calculations, indicating that the stronger binding affinity of the ODFB− anion drives it closer to the lithium surface, where it undergoes more intensive decomposition reactions and contributes to the formation of the inner interfacial layer. More interestingly, an enhanced peak at 837 cm−1 associated with LiF was observed, further confirming the formation of an inorganic-rich interphase (Fig. S20b).

To further illustrate the nanostructure of SEI, the sputtering XPS was utilized to investigate Li||NCM811 lithium negative electrode side after 100 cycles. In the F1s spectra of Fig. 2c, the LiF component is dominantly found in the outer layers of the SEI. In contrast, by comparing the B1s and N1s spectra, an increased peak width and intensity of B–O and Li3N are found with longer etching time, which indicates that they are mainly distributed in the inner part of SEI layer (Fig. 2d, e). LiF could increase the stiffness of the surface and stability of the interfacial layer, while Li3N could simultaneously enhance the ionic conductivity and facilitate ion transport (Table S2). The B–O related substances in this layer is a glassy-type matter with low crystallinity, leading to homogeneity flux. In contrast, the SEI layer formed by utilizing 1 M LiPF6 EC/DMC electrolyte consists mainly of LiF as an inorganic material (Fig. S22a). Due to the low ionic conductivity of SEI, the transport of lithium ions is hindered, resulting in uneven deposition on the surface of the lithium negative electrode, and ultimately the formation of lithium with highly porous morphology and whisker structure. The Derjaguin–Müller–Toporov (DMT) modulus was subsequently quantified by Atomic Force Microscopy (AFM) to assess the impact of introducing B–O bonding to change the mechanical stability of SEI. Average DMT of as-formed B–O-rich SEI films by 0.5 M-DBN electrolyte was 2.12 GPa, which was 1.44 times higher than that of SEI films formed by the 1.5 M-N (1.5 M of LiNO3 of TEP, noted as 1.5 M-N) electrolyte (the average DMT of 1.5 M-N was 1.47 GPa) (Fig. 2f and Fig. S23). Therefore, B–O-containing SEI with high DMT modulus enables better adaptability to volume changes during the lithium plating and stripping process and thereby dismisses lithium deformation and negative electrode fracture. Figure 2g and S24 show the atomic ratios of several elements under different etching times. It is evident that P content in the SEI layer is only 7% after 100 seconds of etching, while the B content is as high as 17% (Fig. 2g). This suggests that the SEI layer in the 0.5 M–DBN electrolyte was formed mainly by decomposition of lithium salts, which correspond to a higher proportion of C and O elements in this layer, rather than the solvent TEP itself. The in situ decomposition of LiBF4 would provide a high content of F atoms in this layer, which was not found to be as high as B by the estimation of the XPS result. Therefore, a large amount of B–O comes from the preferred and expected decomposition of LiODFB instead of other salts, consistent with the theory suggested by simulation. In contrast, the commercial electrolytes are usually adapting salt of LiPF6 which does not contain C and O elements, and thus the layer of SEI with high C and O ratios are caused by the decomposition of EC and DMC solvents (Figs. S25 and S26). Overall, for the SEI formed in the 0.5 M-DBN electrolyte, LiF has high interfacial rigidity to prevent the growth of lithium dendrites, and compounds such as Li3N have high Li+ ionic conductivity, which helps to improve the ionic transport properties of the SEI. Such multifunctional inorganic compounds together with B–O bonded highly flexible substances form a robust SEI with high ionic conductivity and high flexibility, and the resulting SEI achieves dense bulk deposition of lithium and high plating/stripping efficiency.

To further elucidate the spatial distribution of the tri-phase SEI formed in the 0.5 M-DBN electrolyte, Time-of-flight Secondary Ion Mass Spectrometry (TOF-SIMS) was employed to provide a three-dimensional visualization of the lithium negative electrode structure evolution (Fig. 2h, i). Interestingly, LiF particles are primarily concentrated in the outer SEI, while Li3N and B–O species are mainly found in the inner SEI. Furthermore, the presence of this bilayer SEI structure was corroborated by cryo-electron microscopy observations (Fig. S29). Overall, the nonuniform distribution of the three added anions and their specific solvation sheaths is evident, which allows for controlled decomposition sequences, as anticipated.

Deposition of lithium metal negative electrode
To verify that the formation of a three-phase interfacial layer plays a key role in achieving uniform lithium deposition, the surface of the lithium negative electrode was characterized. Figure 3a shows the Coulombic efficiency (CE) of a Li||Cu half-cell assembled with the 0.5 M–DBN electrolyte to be approximately 97.22% after 100 cycles. For comparison, the CE of Li plating/stripping with the 0.75 M–DB (0.75 M LiODFB and LiBF4 in TEP, denoted as 0.75 M–DB), 1.5 M–N, and 1 M LiPF6/EC–DMC electrolytes are approximately 91.45%, 87.05%, and 84.52%, respectively. In addition, a Li||Cu symmetric cell assembled with the 0.5 M–DBN electrolyte demonstrated a higher CE of 99.18% after 10 Li plating/stripping cycles at a current density of 0.25 mA cm−2 and with an area capacity of 0.5 mAh cm−2, compared to Li||Cu symmetric cells assembled with 0.75 M–DB (94.53%), 1.5 M–N (87.89%), and 1 M LiPF6/EC–DMC (88.34%) electrolytes (Fig. S30b), which further emphasizes the benefits of three-phase SEI formation. Furthermore, the incorporation of alternative flame-retardant additives into the 1 M LiPF6 EC/DMC electrolyte resulted in markedly lower Coulombic efficiency compared to the 0.5 M-DBN-based system, thereby reinforcing the effectiveness of the proposed approach (Fig. S30a). In Fig. 3b and Fig. S31, it can be observed that cells using the 0.75 M-DBN, 1.5 M-N, and 1 M LiPF6 EC/DMC electrolytes experienced short-circuits stopping after cycling for 680 h, 40 h, and 840 h, respectively, at a current density of 0.25 mA cm−2 and an areal capacity of 0.5 mAh cm−2. In contrast to these three conditions, a battery assembled with the 0.5 M–DBN electrolyte exhibited a longer cycle life exceeding 1000 h for Li plating/stripping, mainly due to the effective suppression of lithium dendrite growth. Similar tendencies were also observed at higher current densities, with a current density of 1 mA cm−2 and an areal capacity of 1 mAh cm−2 (Fig. S32). The electrochemical performance demonstrates that the 0.5 M-DBN electrolyte is functioning in enhancing the reversibility and enlarging the length of Li plating/stripping. This improvement is attributed to the formation of a robust interfacial protective layer on the lithium metal negative electrode by the tri-salt.

To directly observe lithium deposition in the electrolyte, Scanning Electron Microscopy (SEM), Scanning Probe Microscopy (SPM), and in situ electrochemical optical microscopy were used to characterize the surface of the lithium negative electrode. Figure 4c, d displays the morphology of lithium negative electrodes after plating in 0.5 M-DBN and 1 M LiPF6 EC/DMC electrolytes. The formation of lithium dendrites and the porous surface morphology are associated with low CE values and with instability in Li plating/stripping. In contrast, 0.5 M-DBN electrolyte provides dense and uniform lithium deposition, which is consistent with the battery’s high CE and stable Li plating/stripping performance. Similar results were also observed for Li||NCM811 cells after 100 cycles (Fig. S33). Furthermore, the growth of lithium whisker-liked dendrite leads to a considerable rise in porosity. This tendency is apparent in the cross-sectional SEM images, as shown by the thicknesses of the deposited lithium layers, which were 28.20 μm and 47.20 μm for cells cycled in the 0.5 M–DBN electrolyte and the commercial electrolyte, respectively (Fig. 3e, f). As shown in Fig. 3g–h and Fig. S34, compared with the conventional 1 M LiPF6 EC/DMC electrolyte system, the use of the 0.5 M-DBN electrolyte significantly reduces the thickness of the deposited layer on the lithium metal negative electrode surface, while promoting a denser and bulk-like lithium deposition morphology. Potential map analysis by scanning probes suggests that the potential distribution of the 0.5 M-DBN electrolyte is uniform. This uniform and flat in situ potential distribution prevents the formation of lithium dendrites and facilitates the uniform deposition of Li+, where Li deposition occurs at relatively low potential. Consequently, the use of the 0.5 M–DBN electrolyte produces a dense and flattened lithium layer with minimal dendrite formation, which thus improves the CE.

Lithium deposition was monitored in a battery by 0.5 M-DBN and 1 M LiPF6 EC/DMC electrolytes at a current density of 1 mA cm−2 using in situ electrochemical microscopy (Fig. 3i). The growth of dendritic crystals was observed during the deposition process using a commercial electrolyte, accumulating in the formation of a large dendritic crystal on the lithium surface after only 5 h plating. In contrast, adaptation of the 0.5 M-DBN electrolyte was found to effectively inhibit the growth of dendritic crystals on the lithium surface. After deposition of 5 mAh cm−2, uniform lithium deposition on the lithium surface was still observed. This uniform deposition was attributed to the formation of a highly ion-conductive SEI layer on the lithium surface, which provides high ionic conductivity at the interface and promotes uniform Li+ deposition.

Interfacial chemistry on the positive electrode side of NCM811
In addition to the interface layer on the negative electrode side, the nature of the cathode-electrolyte interphase (CEI) layer is equally important, which affects the stability and electrochemical properties of the battery, especially at high potentials. A stable, relatively complete, and uniform CEI can ensure the stability of the structure and composition of the positive electrode interface, prevent unexpected electron transfer on the positive electrode/electrolyte interface, and hinder the side reaction between the positive electrode and electrolyte. In a rigorous float test at 4.6 V with an 8 h retention time, the cell with the 0.5 M-DBN electrolyte had a leakage current density of 2.37 μA cm−2, which was only 23% of the leakage current density observed with the 1 M LiPF6 EC/DMC electrolyte (10.17 μA cm−2) (Fig. 4a). The float test with 0.5 M-DBN electrolyte showed a lower leakage current density, which means fewer side reactions occur on the positive electrode surface. At high potentials, parasitic reactions between the delithiated positive electrode and the electrolyte accelerate the dissolution of transition metals and structural degradation, which is considered one of the major failure mechanisms of high-potential lithium-ion batteries. Successfully decreasing the likelihood of side reactions between the 0.5 M–DBN electrolyte and the delithiated positive electrode effectively suppresses the dissolution of transition metals (Fig. 4b and Fig. S35), which is expected to maintain the stability of the cell’s performance.

X-ray diffraction (XRD) is a valuable and widely used technique for monitoring high-nickel positive electrode materials. XRD characterization of positive electrodes removed from cells using the 0.5 M-DBN and 1 M LiPF6 EC/DMC electrolytes was conducted (Fig. 4c) after 100 cycles. The XRD analysis indicates that the peak positions of the positive electrode samples are essentially unchanged after cycling with the 0.5 M–DBN electrolyte. For comparison, the positive electrodes of batteries using commercial electrolytes show significant peak shifts, indicating a partial transformation of the interfacial layered structure into an irreversible rock-salt phase. The structural evolution of the NCM811-positive electrode has been studied by in situ XRD measurements during charge and discharge with various electrolytes. Peaks and Bragg angles of the (003) and (101) crystal planes of the NCM811-positive electrode in batteries using the commercial electrolyte demonstrate significant changes during the phase transition from the H1 phase to the M phase below ~3.8 V (Fig. 4e and Fig. S36). It was observed that the (003) and (101) crystal planes of the NCM811-positive electrode were in the H2 phase, while the (101) crystal planes of the NCM811-positive electrode were in the H2 phase as well. After transitioning to the H2 phase, subsequently the (003) peak shifts to the left (low angle) and then the right (high angle), indicating a phase transition from the H2 phase to the H3 phase, and also a decrease in layer spacing. In the cell using 0.5 M-DBN electrolyte, the (003) diffraction peak of the NCM811-positive electrode exhibits only a slight shift of 0.72°, compared to 0.82° for the NCM811-positive electrode in the 1 M LiPF6 EC/DMC electrolyte, indicating that the CEI formed in the 0.5 M-DBN electrolyte effectively suppresses the variation in the lattice parameter c. Furthermore, the smaller displacement of the (101) peak of the NCM811-positive electrode of the battery using 0.5 M-DBN electrolyte during the charging process is related to the lattice parameter a (Fig. 4d and Fig. S37). These findings confirm that the NCM811-positive electrode in the cell with 0.5 M-DBN electrolyte sustains a more reversible structural evolution and remains stable during cycling. In addition, SEM images revealed significant cracks on the positive electrode of the battery using the commercial electrolyte, whereas no cracks were observed on that of the battery using the 0.5 M–DBN electrolyte (Fig. S38).

To investigate why the 0.5 M–DBN electrolyte is a better match for the NCM811-positive electrode than the commercial electrolyte, we characterized the interfacial layer at the positive electrode. The CEI of the NCM811-positive electrode material was directly observed using high-resolution transmission electron microscopy (HRTEM) after cycling. The NCM811-positive electrode surface of the battery using the 0.5 M–DBN electrolyte was coated with an approximately 3.83 nm thick layer of amorphous CEI. The surrounding near-surface laminar structure remained, indicating almost no phase transition after cycling. Conversely, the NCM811-positive electrode surface of battery with commercial electrolyte exhibited a 15.8 nm thick amorphous CEI layer after cycling (Fig. 4f). Figure 4g shows the XPS O 1s spectra of the positive electrode surface of the battery using the commercial electrolyte, indicating that the inner and outer layers of the CEI are composed of organic matter with strong C–O and C=O signals. However, the XPS pattern of the positive electrode surface of the battery using the 0.5 M–DBN electrolyte showed opposite results and also exhibited a lattice oxygen peak in NCM811 upon initial detection, suggesting the formation of a thinner CEI (Fig. 4h). We observed that in the 0.5 M-DBN electrolyte, the inner layer of the CEI exhibited a lower carbon content than the outer layer, in contrast to the commercial electrolyte, which showed the opposite trend (Fig. S39). Figure S40 displays the F 1s spectra, which also confirm that the LiF content on the positive electrode surface of the battery using the 0.5 M–DBN electrolyte is higher than that of the battery using the commercial electrolyte. This verified the formation of inorganic-rich CEIs. Characteristics of CEIs show that the 0.5 M–DBN electrolyte not only has high oxidation resistance but also in situ forms a protective inorganic-rich CEI through the decomposition of the selected lithium salts, thereby lessening the direct contact between the electrolyte and the NCM811-positive electrode and effectively preventing further decomposition of the electrolyte. Thus, the rapid embedding and detachment of lithium ions is enabled. As expected, the inorganic-rich CEI layer formed by the 0.5 M–DBN electrolyte inhibit the side reaction between the electrolyte and the nickel-rich positive electrode, resulting in improved performance at a high cut-off voltage.

Li||NCM811 full cell performance
The compatibility of the 0.5 M–DBN electrolyte with the high-voltage positive electrode was evaluated using a highly active NCM811-positive electrode. Figure 5a–c shows the enhanced performance of the lithium metal full cell with the 0.5 M–DBN electrolyte compared to that with the EC/DMC-based electrolyte. In particular, the 0.5 M-DBN electrolyte demonstrated a high capacity of 205 mAh g−1 (based on the mass of the positive electrode), with a capacity retention rate of 87.90% after 900 cycles at a cut-off voltage of 4.3 V and a 0.2 A g−1. While the capacity of the cell with the commercial electrolyte was only 30.20%. At a higher cut-off voltage of 4.5 V, Fig. 5d shows the improved high-voltage performance of Li||NCM811 cells with the 0.5 M–DBN electrolyte compared to 1 M LiPF6 EC/DMC. Cycling of the as-made Li||NCM811 cells with 0.5 M-DBN was stable, with 90.19% capacity retention and 99.6% average CE after 600 cycles. In comparison, the commercial electrolyte showed only 48.51% capacity retention and a CE below 99% at this voltage. The large voltage drop observed at the beginning of the discharge process in Fig. 5e, f could be attributed to the overpotential of the positive electrode and the increase in SEI thickness on the lithium-metal negative electrode during the cycling process. Based on a successful and suitable adaptation of the NCM811-positive electrode at both 4.3 V and 4.5 V, the designed electrolyte is also compatible with Li||NCM811 cells at a 4.6 V cut-off voltage. With the 0.5 M–DBN electrolyte, a capacity of 173.3 mAh g−1 (based on the mass of the positive electrode) with a retention of 78.65% was obtained after 400 cycles, and a smaller voltage decay compared to that of the commercial electrolyte was also observed (Fig. S43). Obviously, in contrast, only 61.3% of the capacity was maintained in the 1 M LiPF6 EC/DMC electrolyte after 400 cycles. To further understand the performance of 0.5 M-DBN, batteries under high temperatures (60 °C) were assessed. Figure 5g–i shows the stability of the Li||NCM811 full cells at 60 °C using 0.5 M-DBN and 1 M LiPF6 EC/DMC electrolytes at 0.2 A g−1. The capacity retention of the cell with 0.5 M-DBN was 81.18% with an average CE at 99.98% after 600 cycles under 60 °C and almost no voltage drop. In contrast, the Li||NCM811 full cell with 1 M LiPF6 EC/DMC electrolyte showed a rapid capacity decay during 120 cycles and unstable Coulombic efficiency.

To further elucidate the high-voltage stability of the positive electrode with the 0.5 M-DBN electrolyte and the delayed capacity decay mainly caused by phase transitions during cycling, the dQ/dV curves were evaluated at 0.1 A g−1/4.5 V (Fig. 5j and Fig. S46). The dQ/dV curve of the battery with the positive electrode in the commercial electrolyte showed serious capacity loss and voltage decay, with its peak intensity decreasing severely during cycling. Furthermore, Galvanostatic Intermittent Titration Technique (GITT) measurements were conducted on a button cell cycled for 100 cycles (Fig. 6k). Due to phase transitions, structural damage, and a more resistive interphase layer, the Li+ diffusion coefficient on the positive electrode side was found to be much higher in the 0.5 M–DBN electrolyte than in the commercial electrolyte, suggesting a lower barrier inhibiting Li+ from moving across the CEI formed by the 0.5 M–DBN electrolyte. Based on the aforementioned results, Fig. 5m compares the performance of 0.5 M-DBN and 1 M LiPF6 EC/DMC electrolyte systems in six aspects. Except for lower ionic conductivity, the 0.5 M-DBN electrolyte offers attractive features over commercial carbonate-based electrolytes, including higher safety, a wider potential window, higher cycling stability for lithium-metal batteries, higher Coulombic efficiency for Li||Cu cells, and broader operating conditions at elevated temperatures.

Pouch cell performance and safety assessment
Early-stage research on high-voltage Li-ion batteries concentrated on low-loading positive electrodes with excessive electrolyte and sometimes neglected the impact of realistic operational conditions during cycling. In this study, a practical Li||NCM811 pouch cell was fabricated using a high-loading positive electrode (22 mg cm−2) and a restricted amount of electrolyte (electrolyte-to-capacity ratio of 1.75 g Ah−1) to implement the designed gradient SEI. The specific energy of this pouch cell was calculated by the mass of all components (Table S5). The assembled Li||NCM811 cell exhibited a specific energy (based on the mass of the whole cell) of 430.51 Wh kg−1, calculated based on the mass of the full cell, and maintained a stable capacity with 84.86% retention after 83 charging/discharging cycles at 0.02 A g−1, confirming the performance of the three-phase SEI design (Fig. 6a, b).

In this work, utilization of 0.5 M-DBN is expected to improve the flame-retardant properties of the electrolyte. The heat resistance of the post-cycled lithium negative electrode and NCM811-positive electrode was measured by Differential Scanning Calorimetry (DSC). As shown in Fig. 6c, a heat absorption peak at 187 °C could be observed on both the Cycled Li/0.5 M-DBN and Cycled Li/1 M LiPF6 EC/DMC, mainly caused by lithium. Moreover, the commercial electrolyte showed a large exothermic peak with a heat release of 1470.23 J g−1 immediately following the lithium absorption peak, which could be attributed to the reaction between highly active lithium flakes and the electrolyte. In contrast, the exothermic peak of the Cycled Li/0.5 M-DBN mixture was obviously delayed to 263 °C, and the exothermic was reduced to 178.03 J g−1. On the positive electrode side, Fig. 6d shows that the exothermic heat of the commercial electrolyte in contact with the fully charged NCM811-positive electrode is 1161.66 J g−1, while the exothermic heat of the 0.5 M-DBN electrolyte is only 421.58 J g−1, showing the same trend of higher stability of 0.5 M-DBN at the negative electrode. The measured exothermic data showed that dangerous reactions at both the positive and negative electrodes were effectively suppressed by 0.5 M-DBN.

The hazards of pouch cells usually arise from possible and dangerous reactions between deep-delithiated positive electrodes and the electrolyte, especially between highly oxidizing Ni sites in the positive electrode and organic compounds in the electrolyte, which inhibit the practical application of NCM811-positive electrodes. To further demonstrate the successful formation of a safe CEI on the positive electrode side, we examined the structural changes in the fully charged positive electrode from cells with 0.5 M-DBN and 1 M LiPF6 EC/DMC electrolytes by in situ elevated XRD. Notably, as shown in Fig. 6e, f, the 0.5 M-DBN electrolyte increased the critical transition temperature from 170 to 200 °C, effectively delaying the onset of structural collapse and oxygen release from the NCM811-positive electrode under high-temperature conditions, which corresponds to higher overall battery safety. Enhanced thermal stability on the positive electrode side results from a robust CEI facilitated by the 0.5 M-DBN. Combustion tests were conducted on 0.5 M-DBN and 1 M LiPF6 EC/DMC electrolytes, as depicted in Fig. 6g, illustrating an obvious difference. Commercial electrolytes are highly flammable and can be easily ignited with a lighter. In contrast, the 0.5 M-DBN electrolyte produces only a small amount of smoke when exposed to a flame, demonstrating its safety attributes.

In summary, TEP is utilized as the solvent to ensure the intrinsic non-flammability of the electrolyte, while LiODFB, LiBF4, and LiNO3 serve as solutes to enhance the formation of an optimal interfacial phase. The different binding energies between the added anions and the lithium ions refine the position of the triple salt aggregation in the battery system and the solvation sheath. By calculation and simulation, the ODFB− and NO3− anions, due to their strong binding energies to the lithium ions, are more uniformly distributed on the lithium surface and also get access to the inner part of the solvation sheath layer, where they are preferentially decomposed. Due to the relatively weak binding energy of the BF4− anion to the lithium-ion, it is predominantly distributed in the outer layer of the dissolution sheath, where it decomposes to a small extent to form LiF. In addition, the weak binding energy results in a relatively low dissolution energy, which accelerates the reaction kinetics of the electrolyte. Combined with the NCM811-positive electrode, the Li||NCM811 full cell can withstand 600 cycles at a high cut-off voltage of 4.5 V with 90.19% capacity retention. Even after 600 cycles at 60 °C, capacity retention is 81.18%. Moreover, the pouch cell utilizing the 0.5 M-DBN non-flammable electrolyte attained a high specific energy of 430.51 Wh kg−1, based on the mass of the whole cell, along with a prolonged cycle life of 83 cycles. This ternary salt engineering strategy offers an approach to the design of high-voltage, non-flammable LMB.

Methods
Materials
The following materials were procured from Shanghai Aladdin Bio-Chem Technology Co., Ltd.: Lithium difluoro(oxalato)borate (LiDFOB, ≥99%), Lithium Tetrafluoroborate (LiBF4, 98%), Lithium nitrate (LiNO3, 99%), Lithium Hexafluorophosphate (LiPF6, 97%) Triethyl phosphate (TEP, 99.5%), Ethylene carbonate (EC, >99.5%), Dimethyl carbonate (DMC, >99%), and N-methyl-2-pyrrolidone (NMP, >99.5%). The active material, LiNi0.8Co0.1Mn0.1O2 (NCM811) powder, polyvinylidene fluoride (PVDF, molecular weight ~1,100,000, purity ≥99.5%), Super P Li (particle size: 40–50 nm, purity ≥99.5%; powder conductivity: 10–15 S cm−1), and carbon-coated aluminum current collectors (thickness: 15 μm) were all supplied by Keruode Co., Ltd. The polypropylene (PP) separator used in this study was obtained from the Guangdong Canrd New Energy Technology Co., Ltd. It has a thickness of 16 ± 2 μm, lateral dimensions of 88 mm × 10 m, a porosity of 38–40%, and a puncture strength of ≥280 gf. Lithium metal (battery grade, ≥99.9% purity, metal basis, 15.6 mm in diameter, 100 μm) was purchased from Innochem. The materials received were immediately used to prepare the corresponding electrolytes and handled in an Ar-filled glove box (water and oxygen content <1.0 ppm).

Electrolyte and negative electrode material preparation
(1) 0.5 M-DBN electrolyte: 0.5 M LiDFOB, 0.5 M LiBF4, 0.5 M LiNO3 salt was dissolved in blended TEP. Then put it on a stirring table and stir it until it is clear and transparent to obtain 0.5 M-DBN electrolyte. (2) 0.75 M-DB electrolyte: 0.75 M LiDFOB, 0.75 M LiBF4 salt was dissolved in blended TEP. Then put it on a stirring table and stir it until it is clear and transparent to obtain 0.75 M-DB electrolyte. (3) 1.5 M-N electrolyte: 1.5 M LiNO3 salt was dissolved in blended TEP. Then put it on a stirring table and stir it until it is clear and transparent to obtain 1.5 M-N electrolyte. (4) 1 M LiPF6 EC/DMC electrolyte: 1 M LiPF6 salt was dissolved in blended ethylene carbonate (EC) and dimethyl carbonate (DMC) solvents with a volume ratio of 1:1. The NCM811 electrode was obtained by coating the slurry of NCM811 powders, acetylene black and poly(vinylidene difluoride) with a weight ratio of 8:1:1 in N-methyl-2-pyrrolidone solvent onto an aluminum current collector. Dry the slurry for 8 h at 80 °C and punch into disks with a diameter of 14 mm. The average mass loading of NCM811 is 3–5 mg cm−2.

Cell assembling
All coin cells were assembled in the Argon-filled glove box with a water and oxygen content of less than 0.1 ppm. Li||Li CR-2032 type coin cells were assembled in the following order: positive electrode shell, Li-metal sheet, Celgard PP separator (16 mm), Li-metal sheet, gaskets, splinter, and negative electrode shell. Li||Cu CR-2032 type coin half-cells were assembled in the following order: positive electrode shell, Cu foil (14 mm), Celgard PP separator (16 mm), Li-metal sheet, gaskets, splinter, and negative electrode shell. Both Li||Li symmetrical cells and Li||Cu half-cells were added with 60 μL electrolyte. Li||NCM full cell assembly followed the order: positive electrode case, NCM positive electrode disk, Celgard PP separator (16 mm), Li-metal sheet, gaskets, splinter, and negative electrode shell (a total of 60 μL electrolyte was added to the cell).

The Li‖NCM811 pouch cell (10.0 × 7.0 cm) was assembled in a dry room with a dew point 25 °C. The cell was constructed by stacking eight double-sided coated positive electrode layers (NCM811, areal loading: 22 mg cm−2 per side) and nine metallic lithium negative electrode layers (50 μm thickness on Cu collector) in a staggered sequence, separated by a polyethylene separator. This configuration corresponded to a total theoretical capacity of ~3.0 Ah. The tabs for the current collection were made of aluminum for the cathode and copper for the anode. An electrolyte was injected into the dry cell by weight prior to the final heat sealing, with a controlled electrolyte-to-capacity (E/C) ratio of 1.75 g Ah−1 (total electrolyte weight: 5.25 g). The cell was then subjected to a vacuum soaking period to ensure complete wetting of the electrodes and separator. Any residual gas generated during the subsequent formation process was evacuated under vacuum before the pouch was hermetically final-sealed to ensure cell compactness. The final cell delivered a discharge capacity of 3.0 Ah with an average voltage of 3.8 V, resulting in a Specific Energy of 430.5 Wh kg−1 (total cell mass: 26.48 g). For the cycling performance evaluation, all cells were tested under a fixture applying an external stack pressure of 0.5 MPa to maintain optimal interfacial contact and electrode stability.

Depth profiling XPS
XPS test samples were prepared using Li||NCM811 Li-Metal full cells with different electrolytes (0.5 M-DBN and 1 M LiPF6 EC/DMC) (cycling at a multiplicity of 0.1 C and 10 cycles). The post-cycling lithium negative electrode and NCM811-positive electrode were then subjected to XPS testing. XPS data were collected using Thermo Scientific K-Alpha+. The vacuum in the analysis chamber was ~2 × 10−7 mbar, X-ray source: monochromatic Al Kα source, energy: 1486.6 eV, voltage: 12 kV, beam current: 6 mA, analyzer scanning mode: CAE, work function: 0.5 mA cm−2 and 0.25 mA cm−2: CAE, work function: 4.2 eV. All peaks were calibrated with the C-C bond peak as a reference. The peak was set to 284.8 eV. In addition, high-resolution scans were performed for the following regions: B 1s, C 1s, F 1s, N 1s, O 1s, and P 2p. Depth profiling XPS was performed at an etch rate of 0.2 nm/s.

In situ Fourier transform infrared spectroscopy
We executed in situ FTIR spectroscopy experiments on an IR Tracer-100. To establish a good electrical contact and to ensure an intimate interface between the cathode surface and the germanium crystal, we placed a copper mesh as a working electrode on the surface of the germanium crystal and used a lithium metal foil as a counter electrode and a reference electrode, which were attached to a stainless steel (SS316) spacer. We used an electrolyte containing 0.5 M-DBN and 1 M LiPF6 EC/DMC. To enhance the intensity of the FTIR signal, we used SS316 rods to press the lid onto the coin cell when it was closed, ensuring good electrical contact between the electrodes. After completing the assembly, we let the cell sit for 6 h. We then performed in situ electrostatic discharge and charging experiments using an IVIUM electrochemical workstation. At a current density of 0.5 mA cm−2, we performed 2 h of in situ cell discharge and 2 h of charging and collected FTIR data in real-time to record changes during cycling.

In situ X-ray diffraction
structural changes during electrolyte charging and discharging. Specifically, a thin aluminum window (12 µm thick) was fixed as a transparent window on the cathode side of the cell. The cathode was mounted at the bottom of the cell with the active material facing up. A uniform drop of 40 μL of electrolyte was placed on a glass fiber filter separator (GF/A, Whatman). We then performed in situ electrostatic discharge and charging experiments using the IVIUM electrochemical workstation. We performed charging and discharging at a multiplicity of 0.5 C and collected XRD data in real-time to document changes during cycling.

In situ microscopic imaging
LIB-MS-R Ltd. provides the LIB-MS-R device, which is specialized equipment designed to simulate and observe the growth and corresponding changes of lithium dendrites in lithium-ion button batteries. By precisely controlling the experimental conditions and parameters, the device provides a reliable experimental platform for studying the formation of dendrites in lithium-ion batteries, the growth mechanism, and the correlation with battery performance.

Transmission electron microscope and scanning electron microscope sample preparation process
For postmortem analysis, all cyclic cathodes and anodes are removed directly from the cell. They were washed in ethylene glycol dimethyl ether and required drying for at least 12 h after the washing was completed. The NCM cathodes used for postmortem analysis were obtained after 100 cycles at 1 C multiplicity with different electrolyte conditions at a temperature of 25 °C. SEM samples were obtained by lithium deposition at 25 °C (deposition temperature of 25 °C, deposition current density, and capacity of 0.5 mA cm−2, respectively). The morphology and particle size of the samples were examined by FESEM. TEM and HR-TEM samples were obtained by lithium deposition at 25 °C (deposition temperature of 25 °C, deposition current density, and capacity of 0.01 mA cm−2 and 0.1 mAh cm−2, respectively). A commercially available perforated carbon TEM mesh was used as the working electrode for lithium metal plating in the button cell device. After lithium metal deposition and disassembly of the button cell, the excess electrolyte was removed from the TEM grids by using a double-sided suction cloth in an argon-filled glove box, followed by liquid nitrogen vitrification to prevent exposure to air.

Electrochemical testing
The electrochemical data for this study were collected using CR-2032 coin cell. These batteries were assembled in an Ar-filled glove box with oxygen and water concentrations maintained below 0.01 ppm. Ambient data points were obtained by storing the batteries in a Neware thermostat box during testing at 25 °C and 60 °C, respectively. Before testing, the cells were equilibrated at 25 °C for at least 10 h. The 25 °C potential tests were conducted using a Neware BTS 4000 system, while the electrostatic potential tests were performed using a CHI-760E instrument. Electrochemical impedance spectroscopy (EIS) measurements were performed over a frequency range from 1 MHz to 0.01 Hz using a 5 mV AC excitation signal. The data acquisition density was set to 12 points per frequency decade. Prior to the EIS measurements, the cells were held at open-circuit voltage (OCV) for 1000 s to ensure voltage stabilization. Cyclic voltammetry (CV) was performed at a scan rate of 0.1 mV s−1 using a Chenhua electrochemical workstation. For the linear sweep voltammetry (LSV) test, carbon-coated aluminum foil (c-Al foil) served as the working electrode, and a scan rate of 1 mV s−1 was used. The cell was cycled from 2 V to 6 V (for Li/Li+), followed by stable cathodic scans as LSV curves for each lithium-aluminum-foil cell under different electrolytes.

A custom two-electrode coin cell configuration was employed to determine the ionic conductivity of the electrolyte. In this setup, two stainless steel electrodes were symmetrically arranged with 0.027-inch-thick PTFE spacers placed between them. A fiberglass separator fully saturated with the electrolyte was confined within the cell gasket to maintain a well-defined and constant electrode contact area. The ionic resistance (R) of the electrolyte was obtained via electrochemical impedance spectroscopy (EIS). The conductivity (σ) was then calculated using the expression σ = L/(A*R), where L denotes the distance between the electrodes and A represents the known contact area of the electrode surface. All measurements were performed inside a JHY-H-50L thermal chamber across a temperature range from 0 to 80 °C. Prior to each impedance measurement, the cell was held at the target temperature for 1 h to ensure thermal equilibrium.

The cationic transference number (tLi+) for each electrolyte was evaluated via potentiostatic polarization, a widely adopted electrochemical technique. Measurements were conducted on a CHI-760E workstation using a symmetric Li||Li cell assembled with five layers of Celgard separator. A fixed voltage of 10 mV was applied over 10,000 s to monitor current evolution. Two distinct current readings were recorded: the initial current (I0), corresponding to the total ionic flux under uniform cation distribution, and the steady-state current (Iss), which reflects contributions solely from Li+ migration after polarization. Electrochemical impedance spectroscopy (EIS) was performed immediately before and after the polarization step to extract the corresponding cell resistances (R0 and Rss). The cation transference number was then determined according to the expression tLi+ = I_SS(ΔV - I_0 R_0) / I_0(ΔV - I_SS R_SS), where ΔV denotes the applied polarization potential (10 mV).

Calculation details
The molecular dynamics (MD) simulations were carried out using the GROMACS 2024 package. Atomic interactions within the systems were modeled with the OPLS-AA force field, while cross-interaction LJ parameters were determined via the Lorentz-Berthelot mixing rules. Electrolyte molar ratios, taken from the experimental section of this study, are provided for the four MD models in Table S1. Periodic boundary conditions were applied in all three directions. Long-range electrostatic interactions were computed using the Particle Mesh Ewald (PME) method with a real-space cut-off of 1.2 nm, which was also used as the cut-off distance for short-range van der Waals interactions. The simulation systems were initialized using Packmol. After energy minimization, all models underwent a two-step equilibration protocol: a 2 ns NPT ensemble simulation followed by a 15 ns NVT ensemble simulation. Subsequent production runs were performed for 5 ns under NVT conditions with a 2 fs timestep, and coordinates were recorded every 0.2 ps. Temperature was held constant at 298.15 K using the Nosé-Hoover thermostat, while pressure was maintained at 1 bar through the Berendsen coupling method during NPT equilibration. To enable comparative analysis, an external electric field of −1.0 V/nm was applied along the z axis during NVT simulations. Although this field strength is higher than typical experimental values, such magnitudes are widely adopted in computational studies to generate observable responses within practical simulation timescales. Structural visualization was conducted with VMD, and detailed simulation datasets are provided in Supplementary Data 1.

Adsorption energies were investigated through density functional theory (DFT) calculations performed with the Vienna Ab initio Simulation Package (VASP). The Perdew–Burke–Ernzerhof generalized gradient approximation (PBE-GGA) was adopted for the exchange–correlation functional. Electron–ion interactions were treated using the projector augmented-wave (PAW) method, and a plane-wave cut-off energy of 450 eV was applied throughout. To mimic the influence of voltage, uniform external electric fields of 0, −0.1, and −1 V/nm were introduced along the vacuum direction. Structural optimizations were carried out until the Hellmann–Feynman forces fell below 0.02 eV Å−1 and the energy variation between successive steps was less than 10−5 eV. Dispersion corrections were included via Grimme’s DFT-D3 scheme to properly describe van der Waals interactions in the adsorption models. Transition states were located using the nudged elastic band (NEB) method available in VASP. The adsorption energy for each system was evaluated according to the relation E(binding) = E(A) + E(B) − E(AB), where E(A) and E(B) represent the single-point energies of the isolated substrates, and E(AB) denotes the total energy of the adsorbed system. The detailed outcomes of the DFT calculations can be found in Supplementary Data 2.

Peer review rounds
[{"round": 1, "reviewer_comments": {"reviewer_1": "This manuscript presents an electrolyte formulation by mixing three different lithium salts with a flame-retardant solvent. The\nauthors claim that the proposed electrolyte exhibits superior electrochemical properties compared to a conventional 1 M\nLiPF₆/EC-DMC electrolyte, attributing this enhancement to the cooperative effect of the three anions in forming a robust\nprotective layer on the electrode surface. However, the manuscript contains numerous deficiencies, with significant\nshortcomings in the analysis used to support its conclusions. Consequently, the reviewer finds that this work does not meet\nthe high standards of Nature Communications. The key issues are as follows:\n1. The manuscript adopts 1 M LiPF₆/EC-DMC as the benchmark electrolyte. However, it is well known that this electrolyte\nundergoes severe degradation when paired with a lithium metal anode. This choice appears to be a deliberate attempt to\nexaggerate the performance of the proposed electrolyte. To substantiate the claim of superior performance, a comparison\nshould be made against state-of-the-art imide-based electrolytes.\n2. The study employs molecular dynamics simulations of the Li electrode–electrolyte interface under an applied electric field,\nyet several methodological issues undermine the reliability of these calculations. First, the interaction between lithium metal\nand the electrolyte cannot be accurately described using a standard classical force field (OPLS-AA), as it severely\nunderestimates adsorption strength compared to first-principles calculations. Additionally, an excessively high electric field\n(-1.0 V/nm) is applied to the electrode surface, leading to an unrealistic accumulation of Li ions at the interface, which likely\naffects the anion distribution as well. Since the proposed mechanistic interpretation relies solely on simulation results, further\nefforts are required to ensure an accurate representation of ion distributions at the Li metal surface.\n3. The proposed mechanism is based on differences in ion distributions within a sub-nanometer scale. However, the spatial\nscale of the solid electrolyte interphase (SEI) typically ranges from 10 to 100 nm. Thus, the claim that such minor initial\nvariations directly translate into the experimentally observed SEI composition lacks sufficient justification.\n4. LiNO₃ is widely recognized as an additive that stabilizes the SEI layer, improving both film stability and battery lifespan.\nThe inclusion of LiNO₃ in the electrolyte and its reported effects do not present a novel finding.\n5. If the cooperative effect of three anions is to be claimed as a central finding, it is necessary to consider a broader range of\nanionic species to evaluate the generality of the proposed mechanism. Without such validation, the novelty of the proposed\nstrategy remains questionable.", "reviewer_2": "1. The article is interesting, and reveals a layered SEI formed by ternary lithium salts: an inner layer composed of Li3N and\nB-O, and an outer layer of LiF, achieving a comprehensive performance of high ionic conductivity, mechanical strength, and\ntoughness for the SEI. The results in the manuscript are very meaningful for the development of lithium metal batteries, and I\nwould like to try a industrial demo according to the reported cell design after the manuscript, of course, will be published.\n2. According to the results of the radial distribution function, regardless of whether an electric field is applied, TEP still\npredominates on the lithium metal surface. Following the authors' design logic, higher lithium salt concentrations may result\nin better SEI formation efficiency. The understanding of the optimized 0.5M-DBN should be addressed.\n3. In line 165, \"electrical conductivity\" is suggested to be more accurately changed to \"ionic conductivity\".\n4. The experiment uses the electrolyte of LiPF6/carbonate as the reference, while the electrolyte of TEP/three salts is a\ncompletely different recipe. It would be more convincing if a comparison of carbonate/three salts with LiPF6/carbonate will\nbe adopted.\n5. It is unclear that the first appearance of 1.5 M-N in line 171 while its explanation is in line 213. This should be revised for\nclarity.\n6. Are the three salt concentrations in 0.5-DBN all 0.5 mol? The 1.5M-N electrolyte is noted later as 1.5 mmol, which makes\nthe comparison of SEI strength in line 171 meaningless.\n7. In line 212, is the salt concentration also in mmol? Comparing to the total salt concentration of 1.5 mol in 0.5M-DBN, it is\nalso not appropriate.\n8. Figure S1b shows that the ionic conductivity of 0.4 M-DBM is 4.92 mS/cm, but in Figure S1c, the ionic conductivity of 0.4\nM-DBM is shown as 3.86 mS/cm. Please confirm the actual ionic conductivity of the electrolyte.\n9. In Figure 2a, there is a typo in \"In-suit Fourier-transform Infrared (FTIR) spectra\". The spectra provided in the article are\ndifficult to be interpreted in terms of the trends of the signal peaks. It would be more helpful to present the data as Contour\nplots.\n10. The initial discharge capacity of 2.8-4.5 V (Figure 5b & Figure S39) at 0.5 C is lower than that of 2.8-4.3 V (Figure 5a)\nat 1 C. How to understand?\n11. In Figure 5l \"DLi+ of NCM523 cathode\" should be \"DLi+ of NCM811 cathode\"?\n12. The author claims that the GED of the pouch cell is 452.47 Wh/kg, the battery design in detail needs to be clarified.\n13. TEP based electrolyte could be retardant at a content of 5-10% wt. The molecule itself has a high reactivity with lithium\nmetal, which has been reflected in the optimized coulombic efficiency of only 97.5%. This coulombic efficiency makes it\ndifficult for practical applications in lithium metal cells. Does the author compare TEP to other retardant solvents?\n14. Fig \"S33\" in line 327 should actually be Fig S37. Please confirm.", "reviewer_3": "This manuscript reports a non-flammable phosphate-based electrolyte for lithium metal batteries using a ternary-salt\ncombination of LiDFOB, LiNO3, and LiBF4 in triethyl phosphate (TEP). The electrolyte facilitates the formation of a solid\nelectrolyte interphase (SEI) enriched with B-O and Li3N in the inner layer and LiF in the outer layer, contributing to improved\nstability and reversibility of the lithium metal anode. While the study presents comprehensive characterization, its novelty\ndoes not appear to lie in new material design or good performance. A similar electrolyte formulation of 1 M LiDFOB together\nwith 0.2 M LiNO3 dissolved in TEP, ethylene carbonate (EC) and dimethyl carbonate (DMC), with better Li metal stability\n(Coulombic efficiency of ~98% vs. 97.2% in this work), has been reported several years ago in the Journal of\nElectrochemical Society (J. Electrochem. Soc., 166, A2523). Moreover, some other phosphate electrolytes in literatures have\nalso shown comparable and higher CE with non-flammable feature (e.g., Batteries & Supercaps, 2023, 6, e202200453;\nBatteries & Supercaps, 5, e202100407; Energy Storage Materials, 2024, 71, 103603; ACS Nano 2023, 17, 23, 24227–24241).\nHowever, this work may potentially provide some insights in the mechanistic understanding, if some unclear points can be clarified.\nMajor comments:\n1. The key finding of this work is the contrasting response of DFOB-/NO3- anions vs. BF- anion under electric field, where\nthe former migrate towards Li metal anode while the latter move away, as evidenced by molecular dynamic (MD)\nsimulations. This phenomenon is thought to drive the preferential reduction of DFOB- and NO3- anions to form an SEI with\nB-O and Li3N-rich inner layer and LiF-rich outer layer.\nIf this phenomenon truly occurs in lithium metal batteries, it would be interesting. However, there are several unclear points.\nThe first is concerning the model of MD calculations. Does it consider the effect of electric double layer in the calculation?\nIf not, the electric double layer itself has a species distribution, which might violate the calculation results in the paper, i.e.,\nDFOB-, NO3- and BF4- anions do not move in behaviors predicted by the MD simulation results. Second, since the Li+ ion\nis coordinated with DFOB-, NO3- or BF4- anion to form the charge-neutral ion pair, why can the anions in the ion pair move\ntowards the Li metal anode under electric field? Third, what’s the meaning of y axis in Figure 1c-1f? Does it reflect the real\ncontent of anions and solvents close to Li(100)? If it is, then why all anions and solvents just sum up to ~10%? Fourth, what’s\nthe rational of choosing Li(100) for calculation rather than Li(110), the thermodynamically most stable plane, or other Li\nplanes? The fifth point, the statistcal distribution of DFOB-, NO3- and BF4- does not differ too much in Figure 1c-1f, ~1.0 for\nthe former peak and 0.8 for the BF4- peak. It is not a significant difference for the amount of DFOB-, NO3- and BF4- on Li\nsurface.\n2. For the in-situ FTIR test which can potentially validate the MD simulation results, there are also many confusing issues.\nThe manuscript writes that “during plating, …Notably, the reversibility of BF₄⁻ is significantly higher compared to that of\nODFB⁻, indicating that ODFB⁻ is more extensively consumed”. This phenomenon is difficult to interpret, since the electrode\npolarization should be negative for the plating process in Li/Cu cell, which will expel the negatively charged anions from Li\nmetal surface instead of increasing their concentration. On the other point of view, the decrease in DFPB- intensity might be\nattributed to the negative electric field rather than the consumption of DFPB-.\nA more fatal defect of the FTIR result is the strong intensity of BF4- on the Li surface, which is completely contradictory to the\nMD results in Figure 1 which claims that BF4- can hardly approach the Li metal and thus weakens its reduction. The\ndiscrepancy between the FTIR observation and the MD simulation highlights a significant gap in the study's coherence.\nAlso, the LiF and free TEP peaks are almost not visible, whose change cannot be discerned during test. This is problematic\nbecause these peaks are expected to be critical indicators of the formation and evolution of the SEI, and their absence\nraises doubts about the accuracy of the in-situ FTIR analysis or the interpretation of the spectra.\n3. The formation mechanism of dual-layered SEI is not clear.\nSince DFOB- can be decomposed to form B-O species in the inner SEI, why can’t we detect large amount of LiF in the inner\nlayer, given that DFOB- can donate element F during decomposition?\n4. What’s the role BF4- plays in this electrolyte? From the paper, it seems that the binary combination of DFOB- and NO3-\ncan have a positive effect on the Li metal anode.\n5. The reported Coulombic efficiency values of Li/Li symmetic cells are very confusing and potentially misleading. The\nmanuscript states that “the Li||Li symmetric battery assembled with 0.5 M-DBN electrolyte demonstrated a higher CE of\n99.18% after 10 cycles at a current density of 0.25 mA cm-2 and with an area capacity of 0.5 mAh cm-2 than the Li||lithium\nsymmetric batteries assembled with 0.75 M-DB (94.53%), 1.5 M-N (87.89%), and 1M LiPF6 EC/DMC (88.34%) electrolytes”.\nHowever, since Li reserivor is super excesive in the Li/Li cell, how can the authors calculate the CE? This expression may\nnot be considered professional or appropriate in a Li metal battery study.\n6. The cathode degradation comparsion with XRD characterizations (Figure 4c-4e) does not perform in the perfectly same\ncondition for the reported electrolyte and base electrolyte. The LiPF6 EC/DMC electrolyte has a higher ionic conductivity,\nthus enabling higher initial discharge capacity for Li/NMC811 cell. This corresponds to higher utilization rate of Li in the\nNMC811 cathode under the same cut-off voltage. Since it is widely accepted that the degradation of high-Ni cathode is\nclosely linked to the Li de-intecalation ratio, the higher Li de-intecalation for cathode in the base LiPF6 EC/DMC electrolyte\ncan leads to faster cathode degradation.\nConsidering that XRD peak shift is comparable for both the reported electrolyte and base electrolyte, it is unclear that the\ncathode degradation is caused by the electrolyte itself or merely due to the different Li deintecalation ratio.\n7. The signal of LixC on the NMC811 cathode side is interesting. How can such a highly reducing material exist in the high\nvoltage environment?\n8. The calculation method of energy density of pouch cell is vague. Does it contain the weight of Al package? It is\nchallenging to obtain such a high energy density with a relative high E/C ratio of 1.75 g/Ah. If the Al package weight was\nexcluded, what’s the real energy density for the cell?\nMinor comments\n1. Some other confusing points. For instance, “the height of layer deposited at the anode side is greater and the clump-like\ndendritic are larger when using the 0.5 M-DBN electrolyte than that of battery with 1 M LiPF6 EC/DMC electrolyte.” This\nmeans that the Li metal is more porous in the reported electrolyte, which is not favorable.\n“the leakage current density observed with the 1 M LiPF6 EC/DMC electrolyte (10.17 mA cm-2)”. mA level leakage current is\nso large, which can be larger than the current for cell test.\n2. “relatively complete and uniform CEI can … prevent unexpected electron transfer on the cathode/electrolyte interface”.\nFirst, what’s a complete SEI and unexpected electron transfer? Second, why a complete CEI can prevent unexpected\nelectron transfer?\n3. Some typo and grammar errors, e.g., “Successful decreasing the possibility of the parasitic reaction”"}, "author_response": "Thank you for your email informing us of your editorial decision on our manuscript\n(NCOMMS-25-16384A). According to the instructions in your email, we have\ncarefully considered the helpful criticism from the referees and revised the manuscript\naccordingly. We provide a detailed response to the referees' criticism below. The\nchanges made during revision are highlighted in yellow and the file is uploaded\nseparately for review only.\nReviewer #1 (Remarks to the Author):\nThis manuscript presents an electrolyte formulation by mixing three different lithium\nsalts with a flame-retardant solvent. The authors claim that the proposed electrolyte\nexhibits superior electrochemical properties compared to a conventional 1 M LiPF /EC-\nanions in forming a robust protective layer on the electrode surface. However, the\nmanuscript contains numerous deficiencies, with significant shortcomings in the\nanalysis used to support its conclusions. Consequently, the reviewer finds that this work\ndoes not meet the high standards of Nature Communications. The key issues are as\nfollows:\nResponse: Many thanks for the referee’s valuable suggestion, especially, the insightful\ncomments for further improving the quality of the manuscript. We have carefully\naddressed the issues as follows.\n1. The manuscript adopts 1 M LiPF /EC-DMC as the benchmark electrolyte. However,\n6\nit is well known that this electrolyte undergoes severe degradation when paired with a\nlithium metal anode. This choice appears to be a deliberate attempt to exaggerate the\nperformance of the proposed electrolyte. To substantiate the claim of superior\nperformance, a comparison should be made against state-of-the-art imide-based\nelectrolytes.\nResponse: We deeply appreciate the reviewer's meticulous assessment and insightful\nfeedback. We selected the commercial LiPF /EC-DMC electrolyte as the baseline\n6\nbecause it has been widely adopted in the industry and serves as a standard reference\nfor evaluating electrolyte performance. To further investigate the effects of different\nelectrolyte systems, we also conducted comparative experiments using lithium\nsulfonylimide-based electrolytes, as suggested. These systems include 1.5 M\nLiFSI/TEP, 1.5 M LiTFSI/TEP, 1.5 M LiFSI/EC-DMC, and 1.5 M LiTFSI/EC-DMC.\nOur findings revealed that all four systems suffered from inadequate cycling\nperformance. Specifically, 1.5 M LiFSI/EC-DMC exhibited voltage stagnation during\nthe initial charge, likely due to severe oxidative decomposition (Figure caption1a).\nAlthough 1.5 M LiTFSI/EC-DMC completed charge–discharge cycles, it rapidly lost\ncapacity within 100 cycles (Figure caption1b). Both 1.5 M LiFSI/TEP and 1.5 M\nLiTFSI/TEP electrolytes delivered low initial capacities and experienced rapid capacity\ndecay (Figure caption1c and d).\nThese observations suggest that LiTFSI and LiFSI electrolytes promote aluminum\ncurrent collector corrosion: TFSI⁻ ions aggressively attack the Al O layer, causing\n2 3\nlocalized pitting (Nat Commun 15, 2922 (2024)), while FSI⁻ ions at high voltages\npromote AlF formation and induce mechanical failure within the passivation\n3\nlayer (Journal of Applied Electrochemistry 29, 1053–1061 (1999)). Moreover, LSV\nanalyses indicated that the oxidation thresholds of these electrolytes are below 4.8 V,\ninherently limiting their high-voltage compatibility and contributing to poor overall\nelectrochemical performance (Figure caption1e).\nFigure caption1. (a) Time–voltage profile of the first cycle for the 1.5 M LiFSI/EC-\nDMC electrolyte. (b) Cycling performance of the 1.5 M LiTFSI/EC-DMC electrolyte\nunder 4.3 V 1C and 4.5 V 0.5C conditions. (c) Cycling performance of the 1.5 M\nLiFSI/TEP electrolyte under 4.3 V 1C and 4.5 V 0.5C conditions. (d) Cycling\nperformance of the 1.5 M LiTFSI/TEP electrolyte under 4.3 V 1C and 4.5 V 0.5C\nconditions. (e) Linear sweep voltammetry (LSV) curves of the four electrolytes.\n2. The study employs molecular dynamics simulations of the Li electrode–electrolyte\ninterface under an applied electric field, yet several methodological issues undermine\nthe reliability of these calculations. First, the interaction between lithium metal and the\nelectrolyte cannot be accurately described using a standard classical force field (OPLS-\nAA), as it severely underestimates adsorption strength compared to first-principles\ncalculations. Additionally, an excessively high electric field (-1.0 V/nm) is applied to\nthe electrode surface, leading to an unrealistic accumulation of Li ions at the interface,\nwhich likely affects the anion distribution as well. Since the proposed mechanistic\ninterpretation relies solely on simulation results, further efforts are required to ensure\nan accurate representation of ion distributions at the Li metal surface.\nResponse: We sincerely thank the reviewer for their thorough evaluation of our work.\nThe reviewer correctly pointed out that the OPLS-AA force field may have certain\nlimitations in describing the interactions at the lithium metal/electrolyte interface.\nHowever, the main objective of this study is to investigate the migration behavior of\nanions in the electrolyte under realistic operating voltages, rather than to accurately\nmodel interfacial behavior between lithium metal and the electrolyte. Therefore, the use\nof the OPLS-AA force field remains appropriate and justified for the purpose of this\nwork.\nWe thank the reviewer for their insightful comments regarding the electric field strength\n(−1.0 V/nm) used in our simulations. We acknowledge that this value exceeds the\ntypical electric field strengths encountered under experimental conditions. However, as\ncommonly adopted in molecular dynamics (MD) simulations (see RSC Adv., 2018, 8,\n5255–5267; Phys. Chem. Chem. Phys., 2023, 25, 23937–23953; Advanced Powder\nMaterials, 2025, 100296), the application of a relatively high electric field is a necessary\nstrategy to achieve statistically significant results within the limitations of finite\nsimulation boxes and nanosecond timescales. It is important to emphasize that while\nthe absolute magnitude of the field is elevated, it does not compromise the core\nconclusion of our study—that is, the relative migration behavior of different anions\nunder an applied electric field, which remains the central scientific question addressed\nin this work.\nWe sincerely thank the reviewer for their insightful and thorough evaluation of our\nstudy. We fully understand the reviewer’s concern regarding the mechanistic\nexplanation being primarily based on theoretical simulations. In fact, our investigation\nis built upon a tightly integrated framework that combines both theoretical modeling\nand extensive experimental validation. In our preliminary literature review, we\nobserved that different lithium salt anions exhibit significantly varied binding energies\nwith Li⁺—for instance, LiODFB and LiBOB show stronger binding affinities, while\nLiBF and LiPF exhibit relatively weaker interactions. This observation inspired us to\n4 6\ninvestigate how such disparities manifest in practical battery systems. Through\nmolecular dynamics (MD) simulations, we found that these binding energy differences\nindeed influence the migration behavior of anions. Strongly coordinating anions such\nas NO ⁻ and ODFB⁻ tend to migrate in conjunction with Li⁺ toward the anode interface,\n3\nthereby theoretically favoring the formation of an inner SEI layer enriched with Li N\n3\nand B–O species. In contrast, weakly coordinating anions like BF₄⁻ tend to remain\nfurther from the electrode surface, contributing instead to the formation of a LiF-rich\nouter SEI. To experimentally confirm this SEI formation mechanism, we employed in\nsitu infrared spectroscopy. As shown in Figures 2a and S16, we quantitatively analyzed\nthe interfacial behavior of surface anions by comparing the ratios of characteristic peak\nareas at the end of lithium deposition and during the stripping stage. Notably, PF −\n6\n(which has a weaker coordination with Li⁺) exhibited significantly higher reversibility,\nwhile ODFB⁻ underwent more pronounced consumption, indicating its stronger\nparticipation in interfacial reactions (Figure 2b). This result is consistent with our\ntheoretical calculations and supports the hypothesis that the stronger binding affinity of\nODFB⁻ facilitates its migration toward the lithium surface, where it undergoes intensive\ndecomposition and contributes to the formation of the inner SEI layer. Subsequently,\nwe employed X-ray photoelectron spectroscopy (XPS) and time-of-flight secondary ion\nmass spectrometry (TOF-SIMS) to validate the predicted compositional distribution\nwithin the SEI. To further probe the nanoscale structure of the solid electrolyte\ninterphase, we performed sputtering XPS analysis on the Li||NCM811 interface after\n100 cycles. As illustrated in the F 1s spectrum (Figure 2a), LiF is predominantly\ndistributed in the outer SEI. In contrast, the B 1s and N 1s spectra exhibit broader and\nmore intense peaks for B–O and Li N species with increasing sputtering time,\n3\nindicating their localization within the inner SEI (Figures 2b–c). To further elucidate\nthe spatial distribution of the triphasic SEI formed in the 0.5 M-DBN electrolyte, TOF-\nSIMS analysis was conducted to visualize the structural evolution of the lithium anode\nin three dimensions (Figures 2h-i). Interestingly, LiF species were found primarily in\nthe outer SEI layer, whereas Li N and B–O components were concentrated in the inner\n3\nregion. In summary, our proposed mechanism is supported by a coherent and robust\nchain of evidence—from theoretical predictions to systematic experimental\nvalidation—demonstrating the rationality and reliability of our mechanistic model.\nFigure 1a-f. Snapshots of the lithium-metal-anode-electrolyte interfacial system without (a)\nand with (b) an electric field. The large arrow indicates the direction of the external electric\nfield. Quantity distribution in the z-direction for each substance. (c)NO −, (d)ODFB−, (e)BF −\n3 4\nand (f)TEP.\nFigure 2a-b. Contour plots of In-suit Fourier-transform Infrared (FTIR) spectra of the\nelectrode/electrolyte interfaces in a 0.5 M-DBN electrolyte (a) and their content variations (b).\nFigure S16. Contour plots of Fourier-transform infrared (FTIR) spectra of the\nelectrode/electrolyte interfaces in a 1 M LiPF EC/DMC.\nmetal anodes recovered from full cell cycled with 0.5 M-DBN electrolyte at sputtering times of\n0, 50, and 100 s, respectively.\nFigure 2h-i. Time-of-flight Secondary Ion Mass Spectrometry (TOF-SIMS) depth\ndistributions (h) and three-dimensional images of various species (i) within a 0.5 M-DBN\nelectrolyte.\n3. The proposed mechanism is based on differences in ion distributions within a sub-\nnanometer scale. However, the spatial scale of the solid electrolyte interphase (SEI)\ntypically ranges from 10 to 100 nm. Thus, the claim that such minor initial\nvariations directly translate into the experimentally observed SEI composition lacks sufficient\njustification.\nResponse: We are grateful to the reviewer for the constructive comments and for\nhighlighting critical questions regarding our proposed mechanism. The reviewer rightly\nnoted that our mechanism fundamentally depends on the nanoscale modulation of\ninterfacial ion distribution. This mechanism relies on the variation in binding energy\nbetween lithium ions (Li+) and different anions, specifically ODFB−, NO ⁻, and BF −.\n3 4\nDue to the significantly higher binding energies of Li+ with ODFB⁻ and NO − compared\n3\nto BF −, Li+ tends to migrate alongside ODFB− and NO − under the influence of an\n4 3\nelectric field, resulting in their accumulation near the lithium metal surface. In contrast,\nBF −, with its weaker binding, remains more dispersed and distant from the interface.\nelectrolyte interphase (SEI), where the inner layer is enriched with B–O bonds and\nLi N, while the outer layer primarily consists of LiF. We would like to emphasize that\n3\nthis effect is not restricted to the initial stages but persists throughout battery operation.\nSince the electrolyte consistently contains all three anions and the electric field\ncontinuously drives ion migration, the selective distribution process remains active\nthroughout. As shown by the in-situ infrared spectroscopy in Figures 2a-b and S18,\nLiODFB is continuously and irreversibly consumed on the copper current collector side\nduring the discharge stage. Therefore, the binding-energy-driven distribution plays a\nsustained and pivotal role in steering the long-term evolution of SEI composition and\nstructure. We thank the reviewer again for enhancing the rigor of our explanation.\nFigure 2a-b. Contour plots of In-suit Fourier-transform Infrared (FTIR) spectra of the\nelectrode/electrolyte interfaces in a 0.5 M-DBN electrolyte (a) and their content variations (b).\nFigure S18. (a) Contour plots of Fourier-transform infrared (FTIR) spectra of the\nelectrode/electrolyte interfaces in a 0.5 M-DBN. (b) LiF content variation in in situ\nFTIR.\n4. LiNO is widely recognized as an additive that stabilizes the SEI layer, improving\n3\nboth film stability and battery lifespan. The inclusion of LiNO in the electrolyte and\n3\nits reported effects do not present a novel finding.\nResponse: We sincerely thank the reviewer for the careful and thoughtful review of\nour manuscript. As the reviewer rightly pointed out, the addition of LiNO and its effect\n3\non SEI formation has been previously reported and is not itself a novel finding.\nHowever, the novelty of our work lies not in the use of LiNO as an additive, but in the\n3\ninterfacial regulation mechanism we propose, which is based on the binding energy\ndifferences between Li⁺ and various anions. Specifically, our key finding is that Li+\nexhibits significantly stronger binding with ODFB− and NO − than with BF ⁻ (Li+–\n3 4\nODFB−/NO − > Li+–BF −), which leads to a spatially selective distribution of anions\n3 4\nunder an applied electric field. That is: (1) ODFB− and NO −, with stronger binding\n3\nenergies, tend to co-migrate with Li⁺ and accumulate near the lithium metal surface; (2)\nBF −, with weaker binding, remains away from the electrode interface. This dynamic\n4\nion distribution enables the formation of a gradient SEI structure—an inner layer rich\nin B–O bonds and Li N, and an outer layer primarily composed of LiF. This mechanism\n3\nis fundamentally different from previously reported approaches that rely solely on the\naddition of LiNO (e.g., Angew. Chem. Int. Ed. 2022, 61, e202206682; Angew. Chem.\ncompetitive ion distribution. In contrast, our work introduces a new design concept\nbased on binding-energy-driven ion selection, offering a theoretical foundation and\npractical strategy for precise interfacial engineering in lithium metal batteries. We\nappreciate the reviewer’s comments, which helped us to clarify the innovation of our\nstudy.\n5. If the cooperative effect of three anions is to be claimed as a central finding, it is\nnecessary to consider a broader range of anionic species to evaluate the generality of\nthe proposed mechanism. Without such validation, the novelty of the proposed strategy\nremains questionable.\nResponse: We sincerely appreciate the reviewer’s thorough evaluation of our study. As\nrightly pointed out, the original manuscript did not adequately address the generality of\ndifferent anions. In the revision, we we conducted two additional control experiments\nto validate the broader applicability of our proposed strategy according to the\nreviewer’s suggestion:\nIn the first set, a 0.5 M LiBOB + 0.5 M LiFSI TEP-based electrolyte was employed\n(LiBOB exhibits a higher binding energy than LiFSI). According to our theoretical\npredictions, LiBOB is expected to reside closer to the lithium surface and preferentially\ndecompose, forming a B–O-rich inner SEI layer, while the weaker-binding LiFSI\ncontributes to a LiF-rich outer SEI. Sputtering XPS analysis of the SEI composition\n(Fig. X) fully confirms the presence of this layered structure.\nIn the second set, a 0.5 M LiNO + 0.5 M LiTFSI TEP electrolyte was used (LiNO has\n3 3\na higher binding energy than LiTFSI). Theoretical predictions suggest that LiNO\n3\npreferentially decomposes on the Li surface, forming a Li O-rich inner SEI (Li O was\n2 2\nselected as a marker to exclude the influence of nitrogen from LiTFSI), while LiTFSI\nforms a LiF-dominated outer SEI. Sputtering XPS results (Fig. Y) again confirm this\nspatial distribution.\nThese findings collectively demonstrate that the anion gradient distribution mechanism\nwe proposed is not a system-specific phenomenon, but rather a general interfacial\nregulation principle. This discovery offers critical theoretical support for designing\nhigh-performance SEIs via strategic anion selection. Please refer to the description in\nlines 223–226 of the manuscript: “Meanwhile, experimental verification using other\ncombinations of strong and weak coordinating lithium salts further revealed the same\nSEI formation pattern, confirming the universality and effectiveness of this strategy\n(Fig. S15).”\nFigure caption (a) B1s and (b) F1s XPS spectra of the SEI formed in the 0.5 M LiBOB\n+ 0.5 M LiFSI TEP electrolyte. (c) O1s and (d) F1s XPS spectra of the SEI formed in\nthe 0.5 M LiNO + 0.5 M LiTFSI TEP electrolyte.\n1. The article is interesting, and reveals a layered SEI formed by ternary lithium salts:\nan inner layer composed of Li N and B-O, and an outer layer of LiF, achieving a\n3\ncomprehensive performance of high ionic conductivity, mechanical strength, and\ntoughness for the SEI. The results in the manuscript are very meaningful for the\ndevelopment of lithium metal batteries, and I would like to try a industrial demo\naccording to the reported cell design after the manuscript, of course, will be published.\nResponse: Many thanks for your very positive evaluation on the topic of our research\nand the support for publication, especially, the constructive comments and insightful\nsuggestions for further improving the quality of the manuscript.\n2. According to the results of the radial distribution function, regardless of whether an\nelectric field is applied, TEP still predominates on the lithium metal surface. Following\nthe authors' design logic, higher lithium salt concentrations may result in better SEI\nformation efficiency. The understanding of the optimized 0.5 M-DBN should be\naddressed.\nResponse: We sincerely thank the reviewer for taking the time to read our manuscript\nand for the valuable comments. As the reviewer insightfully pointed out, TEP (triethyl\nphosphate) dominates the interfacial layer on the lithium metal surface regardless of the\npresence of an applied electric field. This observation can be attributed to the\noverwhelming molar ratio of TEP in the electrolyte composition. As shown in Table S1,\nthe TEP content reaches 5.9 mmol per 1 mL of electrolyte, while the total concentration\nof the three lithium salts (LiODFB, LiBF , and LiNO ) is only 0.5 mmol. However,\n4 3\ndespite its quantitative dominance, TEP exhibits a significantly higher LUMO energy\nlevel (Figure 1i) compared to the anions of the lithium salts. This electronic structure\nfeature ensures that the lithium salts are preferentially reduced over TEP, leading to an\nSEI film that is predominantly composed of lithium salt decomposition products rather\nthan solvent-derived species. Regarding the reviewer’s comment on the relationship\nbetween lithium salt concentration and SEI formation efficiency, we fully agree that a\nhigher salt concentration can enhance the formation of the SEI. Nevertheless, we\nemphasize that excessively high concentrations introduce two key drawbacks: (1) a\ndecrease in ionic conductivity and (2) a substantial increase in viscosity, both of which\nseverely impair ion transport in the electrolyte. Our systematic screening experiments\n(Figure S1) demonstrate that under a fixed salt ratio (LiODFB: LiBF : LiNO = 1:1:1,\n4 3\nmol/mol/mol), the 0.5 M DBN-based electrolyte exhibits the highest ionic conductivity\n(5.13 mS/cm). Further increasing the concentration does not improve performance;\non the contrary, it leads to reduced conductivity and ultimately degrades the cycling\nstability of the battery.\nFigure1i. LUMO energy of neutral of EC, DMC, TEP, LiPF , LiODFB, LiNO ,\n6 3\nLIBF .\n(a) Oxidative stability of various electrolytes in various proportional formulations at a\ntotal concentration of 1.2 M. (b) Ionic conductivity of various electrolytes in different\nproportions of formulations with a total concentration of 1.2 M. (c) The Nyquist plots\nof XMol-DBN electrolytes. (d) The Li+ conductivity of XMol-DBN electrolytes.\nTable S1. Detailed components of the four MD models\nTEP LiODFB LiBF LiNO\n4 3\nComponents\n(mmol) (mmol) (mmol) (mmol)\n0.5 M LiODFB+0.5 M LiBF\n4\n5.9 0.5 0.5 0.5\n+0.5 M LiNO in TEP\n3. In line 165, \"electrical conductivity\" is suggested to be more accurately changed to\n\"ionic conductivity\".\nResponse: We thank the reviewer for the careful review of our manuscript. In response\nto the reviewer’s suggestion, we have revised the sentence at line 165 of the manuscript\nas follows: “Due to the low ionic conductivity of SEI, the transport of lithium ions is\nhindered, resulting in uneven deposition on the surface of the lithium anode, and\nultimately the formation of lithium with highly porous morphology and whisker\nstructure.”\n4. The experiment uses the electrolyte of LiPF /carbonate as the reference, while the\n6\nelectrolyte of TEP/three salts is a completely different recipe. It would be more\nconvincing if a comparison of carbonate/three salts with LiPF /carbonate will be\n6\nadopted.\nResponse: We sincerely thank the reviewer for the thorough evaluation of our\nmanuscript. Due to the extremely low solubility of lithium nitrate in commercial\ncarbonate solvents (approximately 0.01 mg/mL), these solvents were not used as the\nelectrolyte medium in this study. In response to the reviewer’s comment, we prepared\na control electrolyte composed of 0.75 M LiODFB/EC-DMC and 0.75 M LiBF /EC-\n3\ntherefore not included in the control formulation). Comparative experiments were\nconducted under two different testing conditions: 4.3 V at 1C and 4.5 V at 0.5C.\nAs shown in the figures caption2, under the 4.3 V 1C condition, the 0.5 M-DBN\nelectrolyte exhibited excellent cycling stability, retaining 87.90% of its initial capacity\nafter 900 cycles. In contrast, the 0.75 M LiODFB + 0.75 M LiBF EC/DMC electrolyte\n4\nretained only 36.97% of its capacity. Under the higher voltage condition (4.5 V), the\n0.5 M-DBN electrolyte still maintained 90.91% capacity retention after 600 cycles.\nThese comparative results clearly demonstrate the superior cycling performance of the\n0.5 M-DBN electrolyte.\nFigure caption2. Comparison of the cycling performance of the two electrolytes under\n(a) 4.3 V at 1C and (b) 4.5 V at 0.5C conditions.\n5. It is unclear that the first appearance of 1.5 M-N in line 171 while its explanation is\nin line 213. This should be revised for clarity.\nResponse: We thank the reviewer for the thorough review of our manuscript. In\naccordance with the reviewer’s suggestion, we have revised the description of the 1.5\nM-N electrolyte and relocated it to the appropriate position at line 171. The revised text\nis as follows: “Average DMT of as-formed B-O-rich SEI films by 0.5 M-DBN electrolyte\nwas 2.12 GPa, which was 1.44 times higher than that of SEI films formed by the 1.5M-\nN (1.5 mmol of LiNO of TEP, noted as 1.5 M-N) electrolyte (the average DMT of 1.5\n6. Are the three salt concentrations in 0.5-DBN all 0.5 mol? The 1.5M-N electrolyte is\nnoted later as 1.5 mmol, which makes the comparison of SEI strength in line 171\nmeaningless.\nResponse: We sincerely thank the reviewer for the careful and detailed review of our\nmanuscript. In response to the comment, we have revised the description of the 1.5 M-\nN electrolyte at line 171 to explicitly state that its total lithium salt concentration is 1.5\nmol L−1. Additionally, as noted at line 89, the total lithium salt concentration of the 0.5\nM-DBN electrolyte is also 1.5 mol L−1. This consistency in concentration ensures the\nscientific validity of the SEI performance comparison. The revised sentence in the\nmanuscript now reads as follows: “Average DMT of as-formed B-O-rich SEI films by\n0.5 M-DBN electrolyte was 2.12 GPa, which was 1.44 times higher than that of SEI\nfilms formed by the 1.5M-N (1.5 mmol of LiNO of TEP, noted as 1.5 M-N) electrolyte\n3\n(the average DMT of 1.5 M-N was 1.47 GPa) (Figure 2f and S19).”\n7. In line 212, is the salt concentration also in mmol? Comparing to the total salt\nconcentration of 1.5 mol in 0.5M-DBN, it is also not appropriate.\nResponse: We thank the reviewer for the careful evaluation of our manuscript. In\nresponse to the suggestion, we have standardized the concentration units for all\nelectrolyte formulations throughout the manuscript, replacing the original millimole\n(mmol) expressions with molar (M) units. The revised description at line 213 now reads\nas follows: “For comparison, CE of cycling with the 0.75 M-DB (0.75 M LiODFB and\nLiBF of TEP, noted as 0.75 M-DB), 1.5 M-N, and 1M LiPF EC/DMC electrolytes were\n4 6\napproximately 91.45%, 87.05%, and 84.52%.”\n8. Figure S1b shows that the ionic conductivity of 0.4 M-DBM is 4.92 mS/cm, but in\nFigure S1c, the ionic conductivity of 0.4 M-DBM is shown as 3.86 mS/cm. Please\nconfirm the actual ionic conductivity of the electrolyte.\nResponse: We sincerely thank the reviewer for the detailed and constructive comments.\nIn response, we have re-measured the ionic conductivity of all electrolyte samples,\nincluding 0.4 M-DBN, using the same batch of cells under identical experimental\nconditions. The corrected data (presented in Supporting Information Figure S1)\nshow good consistency across all samples. We sincerely apologize for the data deviation\ncaused by previously inconsistent testing conditions. The updated results further\nsupport the reliability of the original conclusions.\nFigure S1. Systematic screening of the electrolyte formulation with an optimum ratio.\n(a) Oxidative stability of various electrolytes in various proportional formulations at a\ntotal concentration of 1.2 M. (b) Ionic conductivity of various electrolytes in different\nproportions of formulations with a total concentration of 1.2 M. (c) The Nyquist plots\nof XMol-DBN electrolytes. (d) The Li+ conductivity of XMol-DBN electrolytes.\nA comprehensive comparison of the oxidative stability and Li+ conductivity of different\nelectrolytes with different ratio formulations, with a total concentration of 1.2 M in (a)\nand (b), revealed that the electrolyte with a lithium salt ratio of 1:1:1 exhibited the most\npromising results. Further comparison of the Li+ conductivity of the XM-DBN\nelectrolyte revealed that 0.5M-DBN was selected as the optimal electrolyte formulation.\n9. In Figure 2a, there is a typo in \"In-suit Fourier-transform Infrared (FTIR) spectra\".\nThe spectra provided in the article are difficult to be interpreted in terms of the trends\nof the signal peaks. It would be more helpful to present the data as Contour plots.\nResponse: We sincerely thank the reviewer for the important suggestion regarding data\nvisualization. As rightly pointed out, contour plots can offer advantages in illustrating\nthe trends of spectral features. However, the in situ infrared data obtained in our study\nmay not be well-suited for contour representation due to baseline inconsistencies across\ndifferent time points. As shown in Figure 10, the initial segment of the data (highlighted\nby the dashed box) exhibits notably different baseline levels. If plotted as contour maps,\nthese inconsistencies could mask subtle but important features, such as the LiF-related\npeak at 833 cm⁻¹. Although we have performed minimal baseline correction using the\nlowest intensity as reference, further aggressive normalization to force all spectra onto\na unified baseline would distort peak intensities and potentially compromise the\nreliability of the results. Moreover, as demonstrated in Figure 11, we have attempted to\ngenerate a contour plot for the in situ infrared data. Notably, due to the aforementioned\nbaseline variations, downward-directed generation peaks could not be accurately\ncaptured, further limiting the effectiveness of this approach.\nIn contrast, the line plots we used are better suited for preserving the integrity of the\noriginal data. They clearly present intensity differences over time, capture the evolution\nof negative peaks, and effectively reveal trends in weaker spectral features. Particularly\nfor analyzing how individual peak shapes change over time, line plots offer superior\nresolution and clarity.\nFigure caption 3. (a) Stacked in situ infrared spectra of the 0.5 M-DBN electrolyte. (b)\nContour plot of the in situ infrared spectra for the 0.5 M-DBN electrolyte.\n10. The initial discharge capacity of 2.8-4.5 V (Figure 5b & Figure S39) at 0.5 C is\nlower than that of 2.8-4.3 V (Figure 5a) at 1 C. How to understand?\nResponse: We sincerely thank the reviewer for the careful review of our manuscript\nand the valuable comments. Regarding the mechanism underlying this phenomenon,\nwe acknowledge that it may involve the coupling of multiple factors. However, we\nbelieve that the most critical factor is a partially reversible structural transformation of\nthe cathode material during cycling. Both theoretical analysis and experimental\nevidence suggest that this structural evolution is closely related to the dynamic behavior\nof lithium vacancies within the material. Specifically, during the charging process,\nextensive extraction of lithium ions leads to the formation of a significant number of\nlithium vacancies, rendering the cathode material thermodynamically unstable. This\neffect is particularly pronounced under high cutoff voltage conditions, where deeper\ndelithiation further promotes vacancy formation. Such structural changes have two\nmajor consequences: (1) The accumulation of lithium vacancies substantially slows\ndown the lithiation kinetics during the subsequent discharge process, leading to\nincreased overpotential at the end of the discharge half-cycle. This manifests as a sharp\nvoltage drop and rapid capacity fade. (2) When kinetic conditions permit, the cathode\nmaterial undergoes structural reconstruction to eliminate these vacancies. This\nvacancy-healing process is supported by the gradual capacity recovery observed in\nFigure 5d and Figure S39 (ACS Energy Letters 2019, 4 (8), 1902–1906).\nFigure 5d. Cycling performance of Li||NCM811 cell at 0.5C, 4.5 V.\nFigure S39. (a) Cycling performance and CE of NCM 811 cathode in 0.5 M-DBN\nand 1 M LiPF EC/DMC electrolytes at 1 C multiplicity with up to 4.5 V.\n11. In Figure 5l \"𝐷 of NCM523 cathode\" should be \"𝐷 of NCM811 cathode\"?\n𝐿𝑖+ 𝐿𝑖+\nResponse: We thank the reviewer for the careful review of our manuscript. In\naccordance with the reviewer’s suggestion, we have corrected “NCM523” to “NCM811”\nin the caption of Figure 5i. The corresponding statement in the manuscript text (line\n213) has also been revised accordingly: “(k) The GITT technique was used to measure\nthe charge/discharge voltage profiles and (l) corresponding D of NCM811 cathode\nLi+\nafter 100 cycles with different electrolytes.”\n12. The author claims that the GED of the pouch cell is 452.47 Wh/kg, the battery\ndesign in detail needs to be clarified.\nResponse: We sincerely thank the reviewer for the rigorous and constructive evaluation\nof our work. We apologize for the oversight during manuscript proofreading, in which\nthe weight of the aluminum-laminated pouch was inadvertently excluded from the\nenergy density calculation of the pouch cell. We have now re-evaluated the relevant\ndata, and the corrected value (as presented in Supplementary Table S5) shows that the\nsystem still achieves a high energy density of 430.51 Wh kg−1, with the total mass of\nthe aluminum pouch included. This corrected result continues to demonstrate a\nsignificant advantage over other reported non-flammable electrolyte systems (see\ncomparative data in Fig. 6c), thereby reinforcing the innovative value of our work in\nsimultaneously enhancing both energy density and safety.\nFigure 5c. In both published literature and this work, lithium metal pouch cells with\nhigh energy density and non-flammable electrolytes exhibit state-of-the-art\nperformance.\nTable S5. Pouch cell energy density calculation\nParameter Value\nDischarge Capacity 221 mAh g−1\nArea weight 22 mg cm−2\nCathode (NCM811 with\nActive material ratio 97%\nAl\nArea capacity 4.86 mAh cm−2\ncurrent collector)\nNumber of layers 8\nWeight 15.13 g\nLi thickness 50 μm\nLi anode Number of layers 9\nWeight 1.92 g\nCu thickness 8 μm\nCollector (Cu)\nWeight 1.93 g\nE/C ratio 1.75 g Ah−1\nElectrolyte\nWeight 5.25\ntype Polyethylene\nSeparator\nWeight 0.5 g\nWeight 1.75 g\nPackage and tabs\nSize 7 cm*10 cm\nPouch cell Discharge capacity 3 Ah\nAverage voltage 3.8 V\nWeight 26.48 g\nEnergy density 430.51 Wh kg−1\nSpecific energy = discharge energy (Wh) / total weight (kg). This is a more precise\ncalculation method compared to the method: Specific energy=discharge capacity (Ah)\n× mid-value voltage (V) / total weight (kg).\n13. TEP based electrolyte could be retardant at a content of 5-10% wt. The molecule\nitself has a high reactivity with lithium metal, which has been reflected in the optimized\ncoulombic efficiency of only 97.5%. This coulombic efficiency makes it difficult for\npractical applications in lithium metal cells. Does the author compare TEP to other\nretardant solvents?\nResponse: Thank you for the reviewer’s thorough evaluation and valuable comments\non our work. In response to your concern, we would like to clarify the following:\nalthough the addition of 5–10% TEP can significantly enhance the flame retardancy of\nthe electrolyte, the presence of 90–95% flammable solvents still poses potential safety\nrisks. To address the reactivity between TEP and lithium metal, we introduced lithium\nsalts with low LUMO energy levels, which preferentially decompose to form a stable\nanion-derived interphase. This interphase effectively prevents direct contact between\nTEP and lithium metal. While there remains room for further improvement in\nCoulombic efficiency, the achieved performance is already at an advanced level\ncompared to previously reported phosphate ester-rich electrolyte systems (Sustainable\nEnergy Fuels, 2022, 6, 1281–1288; ACS Energy Lett. 2025, 10, 4, 1700–1711).\nIn addition, we conducted a systematic experimental study to compare the performance\nof different flame retardants. Three commonly used flame retardants (TMP, TEP, and\nPFPN) were individually added to a commercial electrolyte (1 M LiPF in EC/DMC,\n6\ndenoted as S) in volumes of 50 μL and 100 μL. Their Coulombic efficiencies (CE) were\nevaluated using the Aurbach method. The results showed that the S + 100 μL TEP\nsystem exhibited the highest CE among the flame-retardant-modified electrolytes,\nreaching 94.63%, which is an improvement over the baseline commercial electrolyte\n(88.34%) (Figure caption4a). However, this value still falls short of the CE achieved by\nthe 0.5 M-DBN electrolyte (99.18%) (Figure caption4b). The superior performance of\nthe 0.5 M-DBN system is primarily attributed to the synergistic decomposition of the\nthree lithium salts, which forms a highly stable triphasic interphase with enhanced\ninterfacial stability and Li⁺ transport properties. Please refer to the description in lines\n223–226 of the manuscript:“Furthermore, the incorporation of alternative flame-\nretardant additives into the 1 M LiPF EC/DMC electrolyte resulted in markedly lower\nthe effectiveness of the proposed approach (Figure S27).”\nTherefore, this fully flame-retardant electrolyte system—composed of pure TEP as the\nsolvent and three lithium salts as solutes—features a uniquely structured interfacial\nlayer that not only mitigates the corrosive effects of the flame retardant on the lithium\nmetal anode but also enables high Coulombic efficiency.\nFigure caption4. (a) Coulombic efficiencies (CE) of commercial electrolytes\ncontaining TMP, TEP, and PFPN additives measured by the Aurbach method. (b)\nCoulombic efficiency was determined by a modified Aurbach method at 25 °C and 0.25\nmA cm−2 current density.\n14. Fig \"S33\" in line 327 should actually be Fig S37. Please confirm.\nResponse: We thank the reviewer for the careful review of our manuscript. The\nreference to “Fig. S33” in the main text (line 327) has been corrected to “Fig. S37.” In\naddition, the corresponding description at line 213 has been revised accordingly:\n“Figure S37 displays the F1s spectra, which also confirms that the LiF content on the\ncathode surface of battery with 0.5 M-DBN electrolyte is higher than that of with the\ncommercial electrolyte. This verified the formation of inorganic-rich CEIs.”\nReviewer 3\nThis manuscript reports a non-flammable phosphate-based electrolyte for lithium metal\nbatteries using a ternary-salt combination of LiDFOB, LiNO3, and LiBF4 in triethyl\nphosphate (TEP). The electrolyte facilitates the formation of a solid electrolyte\ninterphase (SEI) enriched with B-O and Li3N in the inner layer and LiF in the outer\nlayer, contributing to improved stability and reversibility of the lithium metal anode.\nWhile the study presents comprehensive characterization, its novelty does not appear\nto lie in new material design or good performance. A similar electrolyte formulation of\nand dimethyl carbonate (DMC), with better Li metal stability (Coulombic efficiency of\n~98% vs. 97.2% in this work), has been reported several years ago in the Journal of\nElectrochemical Society (J. Electrochem. Soc., 166, A2523). Moreover, some other\nphosphate electrolytes in literatures have also shown comparable and higher CE with\nnon-flammable feature (e.g., Batteries & Supercaps, 2023, 6, e202200453; Batteries &\nSupercaps, 5, e202100407; Energy Storage Materials, 2024, 71, 103603; ACS Nano\n2023, 17, 23, 24227–24241). However, this work may potentially provide some insights\nin the mechanistic understanding, if some unclear points can be clarified.\nResponse: We sincerely thank the reviewer for the thoughtful evaluation and valuable\nsuggestions regarding our work. With respect to the research field of stabilizing TEP-\nbased electrolytes through lithium salt regulation, we believe that our study differs\nsignificantly from the references cited by the reviewer. A comparative analysis between\nour system and the five referenced works (J. Electrochem. Soc., 166, A2523; Batteries\n& Supercaps 2023, 6, e202200453; Batteries & Supercaps 2023, 5, e202100407;\nEnergy Storage Materials 2024, 71, 103603; ACS Nano 2023, 17, 23, 24227–24241) is\nprovided below to highlight the distinctiveness and innovation of our approach:\nWe note that Reference 1 (J. Electrochem. Soc., 166, A2523) has two key limitations:\n(1) The employed TEP/EC/DMC co-solvent system (8.4:8.4:83.2 v/v) remains\ncarbonate-dominated (91.6% by volume), which cannot ensure full non-flammability\nof the electrolyte. (2) The study only tested the electrolyte with an LFP cathode (2.0–\n4.0 V), without verifying compatibility with high-voltage, high-nickel cathodes. In\ncontrast, our work achieves three major breakthroughs: (1) Development of a fully non-\nflammable, TEP-based electrolyte system without carbonate solvents; (2) Successful\ncompatibility with NCM811 cathodes over a wide voltage window (2.8–4.6 V); (3) A\nnovel voltage-driven anion migration mechanism enabling the formation of a gradient\ninterfacial layer. It is particularly worth noting that the interfacial stabilization in\nReference 1 relies mainly on the preferential reduction of LiNO , whereas our study\n3\nestablishes a fundamentally distinct theoretical framework based on anion migration\nand selective enrichment under electric fields. \nReference 2 (Batteries & Supercaps 2023, 6, e202200453) employed a\nTEP/LiNO /FEC ternary system, where the electrode–electrolyte interface was\n3\nregulated through the synergistic decomposition of FEC and LiNO additives. Although\n3\nboth this study and ours focus on interfacial engineering, the underlying regulation\nstrategies differ fundamentally. Specifically, Reference 2 relies on a static, additive-\nassisted approach, where interfacial layers are chemically formed through the\nintroduction of functional additives. In contrast, our work introduces a novel voltage-\ndriven dynamic interfacial construction mechanism, wherein anions undergo in situ\nelectrochemical migration and gradient decomposition under an electric field, leading\nto the spontaneous formation of a structured interphase. This dynamic,\nelectrochemically-driven strategy stands in sharp contrast to the static, additive-based\nmethod adopted in Reference 2, underscoring the originality and advancement of our\nproposed interfacial design concept.\nReference 3 (Batteries & Supercaps 2023, 5, e202100407) reported a high-dilution\nelectrolyte system by combining TEP with fluorobenzene (FB) as a diluent and using\nLiFSI as the lithium salt. In this work, TEP serves as the primary solvent, while FB is\nintroduced to mitigate the high viscosity and low ionic conductivity typically associated\nwith highly concentrated electrolytes. The study focuses on solvation structure\nregulation, aiming to form an anion-dominated solvation sheath and enhance interfacial\nstability through solvent–salt interactions. In contrast, our study employs a binding\nenergy–driven strategy involving the synergistic use of LiODFB, LiBF , and LiNO .\n4 3\nBy leveraging the differences in binding energies between Li⁺ and various anions, we\nachieve selective interfacial enrichment and controlled SEI composition. Therefore, the\ncore mechanism and scientific rationale of Reference 3 are fundamentally different\nfrom ours.\nReference 4 (Energy Storage Materials 2024, 71, 103603) proposed a TEP-based\nelectrolyte system incorporating a ternary lithium salt combination of LiTFSI, LiDFOB,\nand LiNO . While both this study and ours utilize multi-salt strategies to stabilize\n3\nphosphate-based electrolytes, the underlying mechanisms differ fundamentally.\nReference 4 stabilizes the electrode–electrolyte interface primarily through the\nsynergistic effect of decomposition products derived from the three lithium salts. In\ncontrast, our work for the first time uncovers the differentiated electrochemical\nmigration behavior of DFOB−, NO −, and BF − anions under an applied electric field.\n3 4\nThis selective migration leads to spatially controlled reduction, resulting in the\nformation of a gradient SEI structure composed of a B–O/Li N-rich inner layer and a\nReference 4 are fundamentally different from those of our study.\nReference 5 (ACS Nano 2023, 17, 23, 24227–24241) achieved a synergistic regulation\nof the 1.5 M LiTFSI/TEP electrolyte by incorporating dual additives, LiNO and\n3\ndimethoxyethane (DME). This approach led to the formation of a thin, dense, and\nrobust electrode/electrolyte interface on both the lithium anode and high-nickel cathode\nsurfaces, enhancing the compatibility between the electrodes and electrolyte. Unlike\nour study, which focuses on synergistic regulation of lithium salts to optimize the\ninterfacial composition, this work relies on the cooperative addition of solvents and\nlithium salts to enable high-voltage lithium-metal batteries. Therefore, Reference 5 and\nour study differ fundamentally in their approach and underlying mechanisms.\n1. The key finding of this work is the contrasting response of DFOB−/NO − anions vs.\n4\nthe latter move away, as evidenced by molecular dynamic (MD) simulations. This\nphenomenon is thought to drive the preferential reduction of DFOB− and NO − anions\n3\nto form an SEI with B-O and Li N-rich inner layer and LiF-rich outer layer.\nHowever, there are several unclear points. The first is concerning the model of MD\ncalculations. Does it consider the effect of electric double layer in the calculation? If\nnot, the electric double layer itself has a species distribution, which might violate the\ncalculation results in the paper, i.e., DFOB−, NO − and BF − anions do not move in\n3 4\nbehaviors predicted by the MD simulation results. Second, since the Li+ ion is\ncoordinated with DFOB−, NO − or BF − anion to form the charge-neutral ion pair, why\n3 4\ncan the anions in the ion pair move towards the Li metal anode under electric field?\nThird, what’s the meaning of y axis in Figure 1c-1f? Does it reflect the real content of\nanions and solvents close to Li(100)? If it is, then why all anions and solvents just sum\nup to ~10%? Fourth, what’s the rational of choosing Li(100) for calculation rather than\nLi(110), the thermodynamically most stable plane, or other Li planes? The fifth point,\nthe statistcal distribution of DFOB−, NO − and BF − does not differ too much in Figure\n3 4\n1c-1f, ~1.0 for the former peak and 0.8 for the BF − peak. It is not a significant\n4\ndifference for the amount of DFOB−, NO − and BF − on Li surface.\n3 4\nResponse: We sincerely thank the reviewer for the thorough and careful evaluation of\nour manuscript. The reviewer has raised five critical questions, which we address in\ndetail below.\nIn our calculations, the presence of the electric double layer was fully considered. The\nelectrochemical double layer refers to a nanoscale structure formed by charge\nseparation at the electrode/electrolyte interface, comprising the charge distribution on\nthe electrode surface and the counterion arrangement in the electrolyte. By introducing\nan external electric field in the simulations, the anode interface attracts Li+ ions under\nthe influence of the field, leading to a specific ion aggregation near the anode. This\nbehavior can, to a certain extent, be regarded as representing the formation of an electric\ndouble layer.\nSecond, even though Li+ forms electrically neutral ion pairs with DFOB−, NO −, or\n4\nfield. This phenomenon arises because Li+ ions must reach the anode to accept electrons\nand be deposited. When the binding between Li⁺ and the anion is strong, the anion is\npartially dragged along by the migrating Li+, resulting in its transport toward the anode\n(Energy Storage Materials,2022,45,1-13).\nRegarding the third question, we sincerely apologize for the oversight in labeling the\ny-axis. After careful verification, we confirm that the y-axis in Figures 1c–1f and Figure\nS13 should represent the number density, with the correct label being “Number density\n(N nm−3).” This parameter reflects the number of particles per unit volume and is\ntherefore capable of accurately indicating the local content of anions and solvent\nmolecules near the Li(100) surface. It is also important to note that, as number density\nis a particle-based quantity, the individual component curves (anions, solvents, etc.) are\nplotted independently. Thus, the sum of their values at each position does not\nnecessarily equal a meaningful or fixed total percentage.\nFigure 1c-f. Quantity distribution in the z-direction for each substance. (c)NO −, (d)ODFB−,\n3\n(e)BF − and (f)TEP.\nelectric field.\nRegarding the fourth point, we appreciate the reviewer’s valuable suggestion. It is\nindeed well established that the Li(110) surface is the most thermodynamically stable\nfacet and plays a significant role in promoting uniform lithium deposition. However,\nwe would like to clarify that the molecular dynamics simulations in this study focus\nprimarily on the migration behavior of anions in the electrolyte under an applied electric\nfield, rather than on the lithium deposition process itself. Due to this difference in\nresearch objectives, the impact of surface orientation on our conclusions is relatively\nlimited. It is also worth noting that the use of the Li(100) surface in molecular dynamics\nsimulations of lithium metal interfaces has precedent in the literature (e.g., Energy\nStorage Materials, 2022, 45, 1–13). Our choice of surface is thus consistent with prior\nstudies, further supporting the validity of our modeling approach.\nRegarding the fifth question, we acknowledge the minor differences in the statistical\ndistributions of DFOB−, NO −, and BF − in Figure 1c-1f, where the peak value for\n3 4\nDFOB− is 1.0, and for BF ⁻ is 0.8 N nm−3. At the Li surface, there is no significant\n4\ndifference in the quantities of DFOB−, NO −, and BF −. However, it is important to note\n3 4\nthat our calculations were conducted in a relatively small molecular box. While the peak\ndifference is only 0.2 N nm−3, if this were extended to the actual electrode-electrolyte\ninterface, the number of these species would show a substantial difference.\n2. For the in-situ FTIR test which can potentially validate the MD simulation results,\nthere are also many confusing issues. The manuscript writes that “during plating,\n…Notably, the reversibility of BF − is significantly higher compared to that of ODFB⁻,\n4\nindicating that ODFB− is more extensively consumed”. This phenomenon is difficult to\ninterpret, since the electrode polarization should be negative for the plating process in\nLi/Cu cell, which will expel the negatively charged anions from Li metal surface instead\nof increasing their concentration. On the other point of view, the decrease in DFPB−\nintensity might be attributed to the negative electric field rather than the consumption\nof DFPB−.\nA more fatal defect of the FTIR result is the strong intensity of BF − on the Li surface,\n4\nwhich is completely contradictory to the MD results in Figure 1 which claims that BF −\n4\ncan hardly approach the Li metal and thus weakens its reduction. The discrepancy\nbetween the FTIR observation and the MD simulation highlights a significant gap in\nthe study's coherence.\nResponse: We sincerely thank the reviewer for their in-depth evaluation and\nconstructive comments. In response to the concern regarding the accuracy of peak\nassignment and the interpretation of the reaction mechanism, we have thoroughly re-\nexamined our experimental data and refined the corresponding analyses. The core\nrevisions are detailed as follows:\nAfter recalibrating the in-situ infrared spectroscopy data, we have corrected the peak\nassignments: the peak at 833.24 cm−1 corresponds to the characteristic vibration of LiF;\nthe peaks at 958.62 cm−1 and 977.90 cm−1 are attributed to free TEP and coordinated\nTEP (P–O–C stretching vibration), respectively; the peak at 1026.13 cm−1 is ascribed\nto coordinated LiODFB (Li+–ODFB− complex, denoted as coor ODFB); and the\n1253.73 cm−1 peak arises from the symmetric stretching vibration of the P=O bond in\nTEP. It is important to note that the in-situ FTIR measurements were performed in\ntransmission mode, where the direction of spectral peaks correlates with changes in\nspecies concentration: downward peaks indicate an increase in content, while upward\npeaks reflect consumption.\nIn the Li||Cu half-cell system, the first stage (0–120 min) corresponds to Li⁺ stripping\nfrom the lithium metal anode and its insertion into the copper current collector, while\nthe second stage (120–240 min) involves the oxidation of metallic Li on the copper\nmesh (Li → Li+ + e−) followed by Li⁺ migration and deposition onto the lithium side.\nSpectral analysis shows a continuous decrease in the intensity of the coordinated TEP\npeak (977.90 cm−1), accompanied by an enhancement of the free TEP signal (958.62\ncm−1), directly evidencing Li⁺ desolvation at the electrode interface, with dissociation\nof the TEP coordination structure and release of free molecules.\nNotably, the consumption peak of coordinated LiODFB (1026.13 cm−1) retains 72.62%\nof its final delithiation intensity by the end of the stripping phase. If, as the reviewer\nsuggested, the disappearance of the coor ODFB− signal were merely caused by\nelectrostatic repulsion after desolvation, a recovery (i.e., weakening of the consumption\npeak) should be observed during the stripping stage. However, this contradicts the\nexperimental results. Therefore, we propose that the irreversible consumption of\nLiODFB is more plausibly attributed to electron transfer reactions of ODFB⁻ anions at\nhighly reductive potentials (ODFB− + e− → decomposition products),\nwhich\nsubsequently participate in the formation of anion-derived solid electrolyte interphase\n(SEI).\nFor comparison, we also investigated the conventional 1 M LiPF /EC–DMC electrolyte\n6\nsystem. The results indicate that LiPF -coordinated peaks exhibit highly reversible\n6\nbehavior throughout cycling (reversibility >92%), in sharp contrast to the LiODFB\nsystem. This confirms that strongly coordinating anions (e.g., ODFB−) tend to undergo\nirreversible decomposition and contribute to SEI formation.\nWe have revised and improved the textual descriptions according to the reviewer’s\ncomments, as detailed below: \"As shown in Figure 2a and Figure S17, during the\nlithium deposition stage (i.e., the desolvation process), the pronounced consumption\npeak at 977.90 cm−1 corresponds to the gradual depletion of solvated TEP, while the\nemerging peak at 958.62 cm−1 indicates the release of desolvated solvent molecules at\nthe electrode interface14. Figure S16 presents the in situ infrared spectra of the 1 M\nLiPF EC/DMC electrolyte system. Furthermore, by comparing the ratio of\n6\ncharacteristic peak areas at the end of lithium deposition and during the stripping stage,\nwe quantitatively analyzed the interfacial behavior of surface anions. Notably, PF −\n6\n(which weakly coordinates with Li+) exhibits significantly higher reversibility than\nODFB−, indicating that ODFB− is more extensively consumed in interfacial reactions\n(Figure 2b). This observation aligns well with the theoretical calculations, suggesting\nthat ODFB− preferentially decomposes and contributes to the formation of the inner\ninterfacial layer. More interestingly, an enhanced peak at 833 cm−1 associated with LiF\nwas observed, further confirming the formation of an inorganic-rich interphase (Figure\nS17)25.\" \nFigure 2a-b. Contour plots of In-suit Fourier-transform Infrared (FTIR) spectra of the\nelectrode/electrolyte interfaces in a 0.5 M-DBN electrolyte (a) and their content variations (b).\nAlso, the LiF and free TEP peaks are almost not visible, whose change cannot be\ndiscerned during test. This is problematic because these peaks are expected to be critical\nindicators of the formation and evolution of the SEI, and their absence raises doubts\nabout the accuracy of the in-situ FTIR analysis or the interpretation of the spectra.\nResponse: We sincerely thank the reviewer for the careful evaluation and valuable\nsuggestions regarding our work. In response to the concern about the visibility of the\nIR characteristic peaks for LiF and free TEP, we have conducted a detailed re-\nexamination. Our analysis reveals that the original spectra had relatively broad line\nwidths, which led to insufficient resolution at the key positions of 833.24 cm⁻¹ (LiF)\nand 958.62 cm⁻¹ (free TEP). To more clearly present these two characteristic peaks, we\ncompared the FTIR spectra at two representative time points: 0 min (initial state) and\n240 min (final state) (Figure caption5). The absorption signals at the aforementioned\nwavenumbers are clearly observable in this comparison, confirming the presence of LiF\nand free TEP.\nFigure caption5. In-situ infrared spectra at 0 min and 240 min.\n3. The formation mechanism of dual-layered SEI is not clear.\nSince DFOB− can be decomposed to form B-O species in the inner SEI, why can’t we\ndetect large amount of LiF in the inner layer, given that DFOB− can donate element F\nduring decomposition?\nResponse: We sincerely thank the reviewer for their thorough evaluation of our\nmanuscript. The central mechanism of our work lies in the differences in binding\nenergies between lithium ions (Li⁺) and various anions (ODFB−, NO −, and BF −).\n3 4\nSpecifically, Li⁺ exhibits significantly stronger binding with ODFB− and NO −\n3\ncompared to BF₄⁻. Under an applied electric field, Li⁺ preferentially co-migrates with\nthe more strongly bound ODFB− and NO − anions, leading to their enrichment near the\n3\nlithium metal surface, while BF −, due to its weaker interaction with Li⁺, tends to be\n4\nexcluded from the interface. This dynamic ion distribution ultimately results in the\nformation of a gradient-structured SEI layer, characterized by a B–O bond- and Li N-\n3\nrich inner region and an LiF-dominated outer region.\nWe further appreciate the reviewer’s insightful comment regarding the predominance\nof B–O species rather than LiF in the inner SEI layer derived from the decomposition\nof the ODFB⁻ anion. This observation can be attributed to the relatively low bond\ndissociation energy of the B–O bond in the DFOB⁻ anion (210.7 kJ mol−1), which\ndecreases further to −50.1 kJ mol−1 in the DFOB• radical state (Nature Energy 2018, 3,\n739–746). The significantly weakened B–O bond makes it more susceptible to\nelectrochemical cleavage during interfacial reactions, thereby favoring the formation of\nB–O-rich species in the inner SEI. Consequently, the SEI is dominated by ODFB−-\nderived B–O components rather than LiF in its inner region. This mechanism is in\nexcellent agreement with previously reported decomposition pathways of the DFOB−\nanion (Nature Energy 2018, 3, 739–746).\n4. What’s the role BF − plays in this electrolyte? From the paper, it seems that the binary\n4\ncombination of DFOB− and NO − can have a positive effect on the Li metal anode.\nour electrolyte design, the BF ⁻ anion—owing to its relatively weak interaction with\nLiF-rich outer SEI layer, while its high degree of dissociation significantly enhances\nthe ionic conductivity of the electrolyte. It is worth emphasizing that although the\nNO ⁻/ODFB⁻ binary system can effectively stabilize the lithium metal interface, it also\n3\nleads to increased viscosity (up to 13.4 cP in the 0.75 M-DN system) and reduced ionic\nconductivity (3.68 mS cm⁻¹). In contrast, the introduction of LiBF in the 0.5 M-DBN\n4\nsystem enables a synergistic effect through ion coordination, which not only maintains\ninterfacial stability but also improves the ionic conductivity to 5.13 mS cm⁻¹ (an\nincrease of 39%) and reduces viscosity to 9.44 cP (a 30% decrease) (Figure caption6).\nThese results clearly demonstrate the dual enhancement role of LiBF in both interfacial\n4\nmodification and electrolyte physicochemical optimization.\nFigure caption6. (a) The ionic conductivity of the 0.5 M-DBN and 0.75 M-DN\nelectrolytes. (b) The viscosity of TEP, 0.5 M-DBN, and 0.75 M-DN electrolytes\n5. The reported Coulombic efficiency values of Li/Li symmetic cells are very confusing\nand potentially misleading. The manuscript states that “the Li||Li symmetric battery\nassembled with 0.5 M-DBN electrolyte demonstrated a higher CE of 99.18% after 10\ncycles at a current density of 0.25 mA cm−2 and with an area capacity of 0.5 mAh cm−2\nthan the Li||lithium symmetric batteries assembled with 0.75 M-DB (94.53%), 1.5 M-\nN (87.89%), and 1M LiPF EC/DMC (88.34%) electrolytes”.\ncalculate the CE? This expression may not be considered professional or appropriate in\na Li metal battery study.\nResponse: We sincerely thank the reviewer for the insightful comments. The reviewer’s\nconcern regarding the reported Coulombic efficiency (CE) is entirely reasonable. We\ndeeply apologize for the mistake in our original manuscript, where the “Li||Cu cell” was\nincorrectly described as a “Li||Li cell.” This has now been corrected.\nTo clarify, the CE measurements in our study were performed using Li||Cu half-cells to\nevaluate the reversibility of lithium plating/stripping in the investigated electrolyte\nsystem. Specifically, we adopted the Aurbach method (Adv. Energy Mater. 2018,\n8,1702097, Electrochimica Acta 1997, 42, 697-718), which enables a more accurate\nassessment of electrolyte reversibility. Initially, lithium was deposited at a current\ndensity of 0.25 mA cm⁻2 until a predetermined capacity (Qₜ) was reached, forming a\nlithium reservoir on the copper substrate. This was followed by n galvanostatic\nplating/stripping cycles (n = 10 in our case) using a current density of 0.25 mA cm⁻²\nand a fixed capacity of 0.5 mAh cm⁻2 per cycle. The remaining lithium was\nsubsequently stripped at 0.25 mA cm⁻2 up to 1 V to determine the residual capacity (Qₛ).\nThe average Coulombic efficiency (𝐶𝐸 ) was then calculated using the following\n𝑎𝑣𝑔−𝑛\nequation:\n𝑛𝑄 +𝑄\n𝐶𝐸 = \n𝐶 𝑠\n𝑎𝑣𝑔−𝑛 𝑛𝑄 +𝑄\n𝐶 𝑡\nWhere Q is the plating/stripping capacity for each of the n cycles (fixed at 0.5 mAh\nc\ncm⁻2), Q is the initial lithium reservoir capacity, and Q is the final stripped capacity\nt s\nfrom the copper substrate.\nWe have revised the relevant statements in the manuscript according to the reviewer’s\nsuggestions. The detailed modifications can be found in Lines 214–219 of the revised\nmanuscript: \"In addition, the Li||Cu symmetric battery assembled with 0.5 M-DBN\nelectrolyte demonstrated a higher CE of 99.18% after 10 cycles at a current density of\n0.25 mA cm−2 and with an area capacity of 0.5 mAh cm−2 than the Li||Cu symmetric\nbatteries assembled with 0.75 M-DB (94.53%), 1.5 M-N (87.89%), and 1M LiPF\nthe three-phase SEI formation.\"\n6. The cathode degradation comparsion with XRD characterizations (Figure 4c-4e)\ndoes not perform in the perfectly same condition for the reported electrolyte and base\nelectrolyte. The LiPF EC/DMC electrolyte has a higher ionic conductivity,\n6\nthus enabling higher initial discharge capacity for Li/NMC811 cell. This corresponds to\nhigher utilization rate of Li in the NMC811 cathode under the same cut-off voltage.\nSince it is widely accepted that the degradation of high-Ni cathode is closely linked to\nthe Li de-intecalation ratio, the higher Li de-intecalation for cathode in the base LiPF6\nEC/DMC electrolyte can leads to faster cathode degradation.\nResponse: We sincerely thank the reviewer for the valuable comments. Regarding the\ncomparison of in situ XRD data, we have unified all charge–discharge curves to the\nsame vertical scale for better consistency, as shown in Figure 4c. Additionally, the\nphase transition angles of the NCM811 cathode are clearly labeled in Figures 4d and 4e to\nfacilitate a more direct comparison of structural evolution.\nWe fully agree that ionic conductivity can influence the initial discharge capacity. In\naddition, the initial capacity is also affected by other factors, such as the loss of active\nlithium caused by the formation of SEI/CEI and the oxidation stability of the electrolyte.\nNotably, although the charging time in the 0.5 M-DBN electrolyte is longer (7057 s vs.\n6381 s for 1 M LiPF₆ EC/DMC), which theoretically should result in a higher degree\nof delithiation, the XRD peak shifts of the (003) and (101) planes in the NCM811\ncathode are actually smaller when using the 0.5 M-DBN electrolyte (Figure caption 7).\nThis result indicates that the NCM811 cathode cycled in the 0.5 M-DBN electrolyte\nundergoes a more reversible structural evolution and maintains structural integrity\nduring cycling.\nIn summary, although the 0.5 M-DBN electrolyte delivers a higher discharge capacity,\nthe superior CEI formed effectively suppresses structural phase transitions in NCM811,\nthereby enhancing its long-term cycling stability.\nFigure caption 7. Charge–discharge curves of the NCM811 cathode using (a) 0.5 M\nDBN and (b) 1 M LiPF EC/DMC electrolytes, plotted on the same Y-axis scale with\n6\ncharging and discharging times labeled.\n7. The signal of LixC on the NMC811 cathode side is interesting. How can such a highly\nreducing material exist in the high voltage environment?\nResponse: We sincerely thank the reviewer for raising this critical and insightful point.\nIn response to your comment, we have carefully re-examined the XPS spectra and the\ncorresponding peak deconvolution procedures. Upon thorough evaluation, we\nacknowledge that the previously assigned Li C signal was likely the result of a\nx\nmisinterpretation during the C 1s peak fitting process. To address this, we adopted a\nmore rigorous deconvolution strategy and reanalyzed the original data. The updated\nresults no longer support the presence of a significant Li C phase on the surface of the\nx\nNMC811 cathode.\nWe have accordingly revised the relevant figure and accompanying descriptions in the\nmanuscript. The detailed modifications can be found in Lines [344–346] of the revised\nmanuscript: “We observed that in the 0.5 M-DBN electrolyte, the inner layer of the CEI\nexhibited a lower carbon content than the outer layer, in contrast to the commercial\nelectrolyte, which showed the opposite trend (Fig. S36).”\nFigure S36. C1s XPS spectra of NCM811 cathode in 0.5 M-DBN (a) and 1 M LiPF\n8. The calculation method of energy density of pouch cell is vague. Does it contain the\nweight of Al package? It is challenging to obtain such a high energy density with a\nrelative high E/C ratio of 1.75 g/Ah. If the Al package weight was excluded, what’s the\nreal energy density for the cell?\nResponse: We sincerely thank the reviewer for the rigorous and constructive evaluation\nof our work. We apologize for the oversight during manuscript proofreading, in which\nthe weight of the aluminum-laminated pouch was inadvertently excluded from the\nenergy density calculation of the pouch cell. We have now re-evaluated the relevant\ndata, and the corrected value (as presented in Supplementary Table S5) shows that the\nsystem still achieves a high energy density of 430.51 Wh kg−2, with the total mass of\nthe aluminum pouch included. This corrected result continues to demonstrate a\nsignificant advantage over other reported non-flammable electrolyte systems (see\ncomparative data in Fig. 6c), thereby reinforcing the innovative value of our work in\nsimultaneously enhancing both energy density and safety.\nTable S5. Pouch cell energy density calculation\nParameter Value\nDischarge Capacity 221 mAh g−1\nArea weight 22 mg cm−2\nCathode (NCM811 with\nActive material ratio 97%\nAl\nArea capacity 4.86 mAh cm−2\ncurrent collector)\nNumber of layers 8\nWeight 15.13 g\nLi thickness 50 μm\nLi anode Number of layers 9\nWeight 1.92 g\nCu thickness 8 μm\nCollector (Cu)\nWeight 1.93 g\nE/C ratio 1.75 g Ah−1\nElectrolyte\nWeight 5.25\ntype Polyethylene\nSeparator\nWeight 0.5 g\nWeight 1.75 g\nPackage and tabs\nSize 7 cm*10 cm\nDischarge capacity 3 Ah\nAverage voltage 3.8 V\nPouch cell\nWeight 26.48 g\nEnergy density 430.51 Wh kg−1\nSpecific energy = discharge energy (Wh) / total weight (kg). This is a more precise\ncalculation method compared to the method: Specific energy=discharge capacity (Ah)\n× mid-value voltage (V) / total weight (kg).\nFigure 5c. In both published literature and this work, lithium metal pouch cells with\nhigh energy density and non-flammable electrolytes exhibit state-of-the-art\nperformance.\nMinor comments\n1. Some other confusing points. For instance, “the height of layer deposited at the anode\nside is greater and the clump-like dendritic are larger when using the 0.5 M-DBN\nelectrolyte than that of battery with 1 M LiPF EC/DMC electrolyte.” This means that\n6\nthe Li metal is more porous in the reported electrolyte, which is not favorable.\nResponse: We sincerely thank the reviewer for the careful and insightful review of our\nmanuscript. Due to our oversight, there was an inaccurate description regarding the\nphenomenon observed in the Scanning Probe Microscopy (SPM) analysis. We\napologize for this mistake and have revised the relevant statements accordingly. The\ndetailed modifications can be found in Lines 241–245 of the revised manuscript: \"As\nshown in Figure 3g–h and Figure S31, compared with the conventional 1 M LiPF\nthe thickness of the deposited layer on the lithium metal anode surface, while promoting\na denser and bulk-like lithium deposition morphology. \"\n“the leakage current density observed with the 1 M LiPF EC/DMC electrolyte (10.17\n6\nmA cm-2)”. mA level leakage current is so large, which can be larger than the current\nfor cell test.\nResponse: We sincerely thank the reviewer for the careful reading and valuable\ncomments. We would like to express our sincere apologies for the unit error in the\nmanuscript, where “μA cm⁻²” was mistakenly written as “mA cm⁻².” This was indeed\nan oversight during proofreading. The error has now been corrected in the revised\nmanuscript at line 278. The corrected statement reads as follows: “ In a rigorous float\ntest at 4.6 V with an 8-hour retention time, the cell with the 0.5 M-DBN electrolyte had\na leakage current density of 2.37 μA cm−2, which was only 23% of the leakage current\ndensity observed with the 1 M LiPF EC/DMC electrolyte (10.17 μA cm−2) (Figure 4a).”\n6\n2. “relatively complete and uniform CEI can … prevent unexpected electron transfer\non the cathode/electrolyte interface”. First, what’s a complete SEI and unexpected\nelectron transfer? Second, why a complete CEI can prevent unexpected electron\ntransfer?\nResponse: We sincerely thank the reviewer for the thorough evaluation of our work.\nRegarding the complete cathode–electrolyte interphase (CEI), its essential\ncharacteristics lie in the formation of a stable, uniform, and\nelectrochemically/mechanically inert protective layer. The term \"unexpected electron\ntransfer\" specifically refers to the phenomenon of electron tunneling. Although the\nelectrolyte is intrinsically an ionic conductor, when the CEI contains structural defects\n(e.g., pores or cracks), electrons may tunnel from the cathode through the CEI into the\nelectrolyte. This leads to undesirable oxidation reactions—such as the decomposition\nof EC solvent (EC → CO + byproducts) (Adv. Funct. Mater. 2023, 33, 2210465).\n2\n3. Some typo and grammar errors, e.g., “Successful decreasing the possibility of the\nparasitic reaction”\nResponse: We sincerely appreciate the reviewer’s professional evaluation and valuable\nsuggestions. We have thoroughly revised the manuscript for grammar and language to\nensure clarity, accuracy, and consistency in academic expression.\nThank you for your email informing us of your editorial decision on our manuscript\n(NCOMMS-25-16384A). According to the instructions in your email, we have\ncarefully considered the helpful criticism from the referees and revised the manuscript\naccordingly. We provide a detailed response to the referees' criticism below. The\nchanges made during revision are highlighted in yellow and the file is uploaded\nseparately for review only.\nReviewer #1 (Remarks to the Author):\nWhile the authors have addressed some initial concerns, critical issues remain\nunresolved, particularly regarding the mechanistic interpretation and its theoretical\nfoundation. Without substantial revisions, the manuscript is not suitable for publication.\nResponse: Many thanks for the referee’s valuable suggestion, especially, the insightful\ncomments for further improving the quality of the manuscript. We have carefully\naddressed the issues as follows.\n1. The use of the OPLS-AA force field for lithium metal, combined with a strong\nelectric field of -1.0 V/nm, remains problematic. Although the authors cite precedent,\nthey do not demonstrate that the observed ion migration trends persist under more\nrealistic conditions. As the force field does not capture specific adsorption at the Li\nsurface, the model likely exaggerates interfacial differences under the applied field.\nMore seriously, despite the strong electric field, the anion distributions are noisy, and\nthe field-induced effects cannot be clearly distinguished, raising serious concerns about\nthe validity of the proposed mechanism. Furthermore, the SEI depth profiles, which\nthe authors emphasize as key experimental evidence, reflect cumulative decomposition and\ntransport processes over extended cycling and therefore do not directly validate the\nsimulated initial anion distributions.\nResponse: We sincerely thank the reviewer for the thorough evaluation of our\nmanuscript. Regarding the concern about employing the OPLS-AA force field to\ndescribe the interaction between lithium metal and lithium ions in lithium-metal-based\nsystems, we have given this issue further consideration and performed additional\nvalidation calculations. Molecular dynamics (MD) simulations using the OPLS-AA\nforce field in Gromacs confirmed that this force field indeed incorporates the interaction\nbetween lithium metal and lithium ions. Specifically, the extracted interaction energies\nwere 320.56 kJ/mol (≈3.32 eV) under zero external electric field and 410.39 kJ/mol (\n≈4.25 eV) under an applied external electric field. To further validate the reliability of\nthe OPLS-AA force field description, we conducted density functional theory (DFT)\ncalculations of the adsorption energies of three lithium salts—LiBF , LiODFB, and\n3\nand −1 V/nm. Notably, under 0 V/nm and −1 V/nm (corresponding to the MD\nsimulation conditions), the average adsorption energies of the three salts were 3.57 eV\nand 4.47 eV, respectively. These DFT results are in good agreement with the MD-\nderived interaction energies (3.32 eV and 4.25 eV), indicating that the OPLS-AA force\nfield can reasonably capture the interaction between lithium metal and lithium ions to\na certain extent. Moreover, the DFT results clearly reveal (Figure S14-S15) that the\nadsorption energies of LiODFB and LiNO₃ on lithium metal surfaces are significantly\nhigher than that of LiBF . Therefore, within the system, LiODFB and LiNO are more\n4 3\nlikely to be enriched at the lithium metal surface. Noted. In lines 112–113, the\nmanuscript states: “Meanwhile, the DFT calculations (Fig. S14-15) indicate that\nlithium metal exhibits a preferential adsorption toward LiNO and LiODFB.”\nnote that OPLS-AA is able to reasonably describe the interactions between the Li metal\nanode and Li⁺ ions. Specifically, in the absence of an applied electric field, the Li metal–\nLi⁺ interaction energy predicted by this force field is approximately 320.56 kJ/mol (≈\n3.32 eV), which is very close to the average adsorption energy (3.57 eV) we obtained\nfor the three lithium salts under the same conditions. Similarly, under an applied electric\nfield, the interaction energy is about 418.39 kJ/mol (≈ 4.33 eV), again in good\nagreement with the average adsorption energy (4.47 eV) from our DFT calculations.\nGiven the close correspondence between the key interaction energies obtained via\nOPLS-AA and the DFT-calculated average adsorption energies—both with and without\nan electric field—we are confident that the MD simulation results based on this force\nfield are accurate and reliable.\nFurthermore, following the reviewers' suggestions, we conducted experiments to\ndirectly verify the distribution of the initial anions. First, we employed quantitative\nnuclear magnetic resonance (NMR) spectroscopy to monitor the variation in the\nconcentrations of two lithium salts on the anode side during charging. Specifically, we\nassembled a Li–Li symmetric cell and applied a constant current of 0.2 mA.\nFigureS19a-b illustrate the schematic diagram of the experimental setup, which\nsimulates the charging process of a full cell. After approximately 10 min and 30 min of\ncurrent application, 0.3 mL of electrolyte was collected from the vicinity of the lithium\ndeposition electrode and mixed with 0.3 mL of deuterated chloroform (CDCl ). The\n3\nresulting solutions were analyzed by ¹¹B NMR spectroscopy, and the relative contents\nof the two anions were determined by integrating the corresponding peaks.As shown in\nFigures S19c-d, with increasing lithiation time, the concentration of ODFB⁻ anions\nnear the lithium-deposition side (anode side) gradually increased, whereas the\nconcentration of BF ⁻ anions decreased. This trend is consistent with the predictions\n4\nfrom our molecular dynamics (MD) simulations.\nWe then further examined the nanostructure of the solid electrolyte interphase (SEI)\nat the electrode/electrolyte interface using cryogenic transmission electron microscopy\n(Cryo-TEM) (Figure S29). The SEI was found to feature an outer layer rich in LiF and\nan inner layer containing Li N and B–O species, which provides additional evidence\n3\nsupporting our design concept.\nAs the reviewers' comment, the sputtering XPS depth profiling primarily reflect\nthe cumulative decomposition and transport processes of the electrolyte after prolonged\ncycling, and therefore cannot directly verify the initial anion distribution predicted by\nthe simulations. To address this, we supplemented our study with depth-profiling XPS\nanalyses after different numbers of cycles (1, 20, and 100) (Figure R1). The results\nshowed that, regardless of the number of cycles, the anode interface consistently\nmaintained a nanostructure with an LiF-rich outer layer and an Li N/B–O-containing\n3\ninner layer. This clearly demonstrates that the SEI structure formed in our experimental\nsamples retains excellent integrity throughout cycling.\nFigure S14. The adsorption energy calculations of three lithium salts — LiBF (a, d,\n4\ng), LiODFB (b, e, h), and LiNO (c, f, i) — under three applied electric field conditions:\n3\nFigure S15. Adsorption energies of different lithium salts under various voltage\nconditions\nFigure S19. (a, b) Experimental setup for ion migration studies using 0.5 M-DBN\nelectrolyte. (c) Quantitative NMR results showing the electrolyte composition near the\nlithium deposition electrode. (d) Relative compositional changes in the electrolyte near\nthe lithium deposition electrode.\nFigure S29. (a) Cryo-TEM image of the SEI formed on lithium metal using the 0.5 M-\nDBN electrolyte; the colored rectangles indicate the presence of LiF (yellow), Li N\n3\n(blue), and LiBO (maroon). The marked regions in the cryo-TEM image (a) are\n2\nenlarged in (b–d), showing the corresponding nanocrystals of LiF (b), Li N (c),\ninterface in Li||NCM811 cells using the 0.5 M-DBN electrolyte after 1 cycle (a, b),\n20\ncycles (c, d), and 100 cycles (e, f).\n2. The proposed link between the initial anion distribution and the long-term SEI\nlayering remains speculative. Continued LiODFB consumption does not ensure\nsustained interfacial enrichment. Over extended cycling, differences in anion transport\nproperties are likely to dominate, and without direct evidence, the proposed mechanism\nis not convincingly supported.\nResponse: We sincerely thank the reviewer for the detailed review of our manuscript.\nAs the reviewer rightly pointed out, the growth of the solid electrolyte interphase (SEI)\nis indeed a continuous process. However, the consumption of the lithium salt (in this\ncase, lithium difluoro(oxalato)borate, LiODFB) does not occur extensively or\nuniformly throughout cycling. The critical stage for SEI formation is during the initial\nfew cycles of battery operation, during which LiODFB participates in reactions to form\na uniform and stable SEI layer.\nGiven the electronically insulating nature of SEI, once a complete initial layer is\nestablished, electrons can no longer penetrate the interface to trigger further\ndecomposition of the lithium salt. As a result, in subsequent cycles, the evolution of the\nSEI is primarily driven by the volume changes associated with lithium plating/stripping,\nwhich may locally damage the SEI. In such cases, the electrolyte decomposes at these\ndamaged sites to form new SEI layers for “self-repair.”\nThe SEI formed in our study demonstrates excellent elasticity, allowing it to\neffectively accommodate volume fluctuations during cycling and thereby significantly\nreduce lithium salt consumption from repeated SEI rupture and repair (Figure R2).\nFurthermore, we conducted XPS characterization of the anode interface after different\ncycle numbers, and the results show that the SEI composition and structure remain\nlargely consistent throughout cycling. This provides strong evidence supporting our\nconclusion that once the initial SEI is stably formed, the subsequent lithium salt\nconsumption is substantially reduced (Figure R1).\nFigure R2. DMT modulus distribution of 1.5 M-N (a) and 0.5 M-DBN (b) SEI.\nFigure R1. F 1s (a, c, e) and N 1s (b, d, f) XPS spectra of the lithium metal anode\ninterface in Li||NCM811 cells using the 0.5 M-DBN electrolyte after 1 cycle (a, b),\n20\ncycles (c, d), and 100 cycles (e, f).\n3. The two additional binary systems (LiDOB + LiFSI and LiNO + LiTFSI) provide\n3\nsome insight but are insufficient to establish the generality of the proposed design\nprinciple. Attributing inner-layer formation to LiDOB or LiNO decomposition based\n3\nsolely on Li O detection is unconvincing, as Li O can also result from the\n2 2\ndecomposition of imide-based salts such as LiFSI or LiTFSI. This overlap in\ndecomposition pathways limits the interpretability of the experimental evidence.\nWithout further mechanistic clarification, the claim of a general design principle\nremains premature.\nResponse: We thank the reviewer for the careful evaluation of our manuscript. We fully\nagree with the reviewer’s opinion that, in our previous study involving the LiNO +\n2 3\nmay be inappropriate.\nTo further validate the universality of the proposed mechanism—namely, that\ndifferences in anion binding energy drive their competitive adsorption and preferential\ndecomposition near the lithium metal surface—we specifically designed two entirely\nnew electrolyte systems, distinct from those used in previous studies:\n1. 0.5 M LiBOB / 0.5 M LiPF in TEP: In this system, the binding energy of\n6\n2. 0.5 M LiNO / 0.5 M LiPF in TEP: In this system, the binding energy of LiNO\n3 6 3\nis higher than that of LiPF .\ninterface to investigate the SEI structures formed in each system:\n For the LiBOB/LiPF 6 system: According to our theory, LiBOB, having a\nhigher binding energy, preferentially enriches near the lithium metal surface and\ndecomposes first, forming a B–O-rich inner SEI. In contrast, LiPF , with a lower\n6\nbinding energy, tends to form an outer SEI enriched in LiF. This layered SEI\nstructure was confirmed by depth-profiling XPS analysis, which revealed a B–\nO-rich inner layer and a LiF-rich outer layer (Figure R3a-b).\n For the LiNO 3 /LiPF 6 system: Similarly, based on our proposed mechanism,\nLiNO (higher binding energy) decomposes preferentially near the lithium\n3\nsurface, resulting in a significantly higher Li N content in the inner SEI\n3\ncompared to the outer layer. Meanwhile, LiPF contributes mainly to the\n6\nformation of a LiF-rich outer SEI. The depth-profiling XPS results of this\nsystem were also consistent with the predicted spatial distribution (Figure R3c-\nd).\nOverall, these experimental results provide strong evidence that the anion gradient\ndistribution mechanism we have uncovered is not a phenomenon specific to a particular\nelectrolyte system, but rather a universally applicable interfacial regulation principle.\nFigure R3. (a) F1s and (b) B1s XPS spectra of the SEI formed in the 0.5 M LiBOB +\n0.5 M LiPF TEP electrolyte. (c) F1s and (d) N1s XPS spectra of the SEI formed in the\n6\n0.5 M LiNO + 0.5 M LiPF TEP electrolyte.\n3 6\nReviewer #3 (Remarks to the Author):\nThe authors provide detailed response to the reviewer’s comment. The “voltage-driven\nanion migration” strategy was further discussed in details in the revised manuscript,\nwhich may promote the development of phosphate electrolyte. Some remaining points\nneed to be addressed.\nResponse: Many thanks for the referee’s valuable suggestion, especially, the insightful\ncomments for further improving the quality of the manuscript. We have carefully\naddressed the issues as follows.\n1. The experimental evidence of dual-layer SEI is not convincing. XPS was used in the\nmanuscript to verify the dual-layer feature of SEI, which is a key finding of this work.\nHowever, XPS results can be misleading sometime, where the sputtering is not\nhomogeneous among different components, undermining the claim of dual-layer SEI\nformation. To clearly demonstrate the dual-layer feature of SEI, cryogenic electron\ncharacterization is suggested.\nResponse: We appreciate the reviewer’s careful evaluation of our manuscript. As\nreviewer rightly pointed out, XPS analysis may suffer from reliability issues due to\npossible artifacts introduced during sputtering. To address this concern, we followed\nthe reviewer’s suggestion and employed cryogenic transmission electron microscopy\n(Cryo-TEM) for further characterization.\nAs shown in Figure S29a, the SEI layer formed using the 0.5 M-DBN electrolyte\nexhibits a dense, inorganic-rich morphology on the lithium metal surface. High-\nresolution analysis identified several crystalline components within the SEI, including\nthe LiF (111) plane with a lattice spacing of 0.233 nm, the Li N (220) plane (0.195 nm),\n3\nand the LiBO (022) plane (0.252 nm), as displayed in Figures S29b–d. These\n2\nobservations reveal a distinct layered structure: the inner SEI region is primarily\ncomposed of Li₃N and LiBO , while the outer layer is dominated by LiF. This spatial\n2\ndistribution is fully consistent with our earlier XPS depth profiling and TOF-SIMS\nresults. Noted. In lines 204–205, the manuscript states: “Furthermore, the presence of\nthis bilayer SEI structure was corroborated by cryo-electron microscopy observations\n(Figures S29).”\nSuch a stratified SEI structure offers synergistic advantages: the outer LiF layer\nprovides high interfacial rigidity that effectively suppresses lithium dendrite growth,\nwhereas inner compounds like Li N exhibit high Li⁺ conductivity, enhancing ionic\n3\ntransport through the SEI. Additionally, the presence of B–O bonded components\ncontributes to mechanical flexibility. Together, these features enable the formation of a\nrobust SEI with both high ionic conductivity and mechanical adaptability, facilitating\ndense lithium deposition and significantly improving plating/stripping efficiency.\nFigure S29. (a) Cryo-TEM image of the SEI formed on lithium metal using the 0.5 M-\nDBN electrolyte; the colored rectangles indicate the presence of LiF (yellow), Li N\n3\n(blue), and LiBO (maroon). The marked regions in the cryo-TEM image (a) are\n2\nenlarged in (b–d), showing the corresponding nanocrystals of LiF (b), Li N (c),\n2\n2. The applicability of the so-called “voltage-driven anion migration” strategy in Li\nmetal full cell should be fully charified. This manuscript reports a voltage-based\nstrategy to regulate SEI. However, it is known that Li chemically reacts with electrolyte\nto form an SEI before cycling, i.e., before applying the electric field. Such a chemically\nformed SEI may substantially affect the “voltage-driven anion migration” strategy to\nregulate the SEI structure. The compasion between reported electrolyte and traditional\ncarbonate electrolyte is not very convincing.\nResponse: We sincerely thank the reviewer for the thoughtful and detailed comments.\nAs correctly noted, prior to battery cycling, lithium metal spontaneously reacts with the\nelectrolyte to form a chemically driven solid electrolyte interphase (SEI) that is\nextremely thin, structurally loose, and disordered (ACS Energy Letters, 2024, 9(11),\n5268–5278). Although this chemically formed SEI can offer an initial barrier against\nfurther electrolyte decomposition, its high porosity and limited blocking ability result\nin poor stability, rendering it vulnerable to subsequent electrochemical reconstruction.\nWith the application of external potential during charge/discharge processes, further\nelectrochemical reduction of the electrolyte occurs, leading to the formation of a second\nSEI layer on top of the initial chemical SEI. This electrochemically dominated SEI is\ntypically denser and more functionalized. The electrochemical formation stage largely\ndetermines the final thickness, compactness, and composition of the SEI, which in turn\ncritically influences long-term cycling performance (Batteries & Supercaps, 2021, 4,\n909; Small, 2020, 16(23), e2000756; Energy Environ. Sci., DOI:\n10.1039/D5EE01030F).\nAs the reviewers' suggestion, we prepared a control electrolyte system consisting\nof 0.75 M LiODFB/EC-DMC and 0.75 M LiBF /EC-DMC (it should be noted that\n3\nsolvent). The comparative experiments were conducted under two different conditions:\n4.3 V at 1C and 4.5 V at 0.5C. As shown in Figure R4, under the 4.3 V, 1C condition,\nthe 0.5 M-DBN electrolyte exhibited excellent cycling stability, retaining 87.90% of its\ninitial capacity after 900 cycles. In contrast, the capacity retention of the 0.75 M\nLiODFB + 0.75 M LiBF /EC-DMC electrolyte was only 36.97%. Under the higher\n4\nvoltage condition of 4.5 V, the 0.5 M-DBN electrolyte maintained 90.91% capacity\nretention after 600 cycles.\nFurthermore, to investigate the influence of different electrolyte systems,\ncomparative experiments were performed using lithium sulfonylimide-based\nelectrolytes, including 1.5 M LiFSI/TEP, 1.5 M LiTFSI/TEP, 1.5 M LiFSI/EC-DMC,\nand 1.5 M LiTFSI/EC-DMC.\nThe results showed that all four systems exhibited insufficient cycling\nperformance. Specifically, 1.5 M LiFSI/EC-DMC displayed a voltage plateau\nstagnation during the initial charging stage, likely due to severe oxidative\ndecomposition (Figure R5a). Although 1.5 M LiTFSI/EC-DMC completed charge–\ndischarge cycles, its capacity rapidly decayed within 100 cycles (Figure R5b). Both\n1.5 M LiFSI/TEP and 1.5 M LiTFSI/TEP showed low initial capacities and rapid\ncapacity fading (Figures R5c-d).\nThese comparative results clearly demonstrate the significant advantages of the\n0.5 M-DBN electrolyte in cycling stability and high-voltage tolerance, strongly\nvalidating our electrolyte design strategy.\nFigure R4. Comparison of the cycling performance of the two electrolytes under (a)\n4.3 V at 1C and (b) 4.5 V at 0.5C conditions.\nFigure R5. (a) Time–voltage profile of the first cycle for the 1.5 M LiFSI/EC-DMC\nelectrolyte. (b) Cycling performance of the 1.5 M LiTFSI/EC-DMC electrolyte under\n4.3 V 1C and 4.5 V 0.5C conditions. (c) Cycling performance of the 1.5 M LiFSI/TEP\nelectrolyte under 4.3 V 1C and 4.5 V 0.5C conditions. (d) Cycling performance of the\n1.5 M LiTFSI/TEP electrolyte under 4.3 V 1C and 4.5 V 0.5C conditions. (e) Linear\nsweep voltammetry (LSV) curves of the four electrolytes.\n3. The FTIR results need more detailed discussion. Please provide references for each\npeak’s attribution in the profiles. In particular, the peak position of LiF needs to be\nchecked. Meanwhile, the curve change during charge and discharge process is not very\nclear, since most of curves overlap in Fig. 2a.\nResponse: We sincerely thank the reviewer for the careful evaluation of our manuscript.\nIn light of your professional guidance, we have we have conducted a more detailed\ndiscussion of the FTIR results in the revised manuscript “As shown in Figures 2a and\nS20, during the lithium deposition stage (i.e., the desolvation process), the pronounced\ndepletion peaks at 977.90 cm−1 (P–O–C) and 1255.31 cm−1 (P=O) correspond to the\ngradual consumption of solvated TEP, while the emerging peak at 958.62 cm−1 (P–O–\nC) indicates the release of desolvated solvent molecules at the electrode interface14.\nFigure S19 presents the in-situ infrared spectra of the 1 M LiPF /EC–DMC electrolyte\n6\nsystem along with the corresponding spectral analysis.” in Lines 148-154. Furthermore,\nthe assignments of each peak in the in situ FTIR spectra and the corresponding\nreferences are as follows: the characteristic peak of LiF appears at 837 cm⁻¹ (Reference:\nNat. Commun. 13, 1398 (2022)); the peaks corresponding to P–O–C, B–O–R, and P=O\nare located at approximately 980 cm⁻¹, 1020 cm⁻¹, and 1260 cm⁻¹, respectively\n(Reference: Energy Storage Materials 42, 628–635 (2021)).\nRegarding the reviewer’s concern about the insufficiently clear spectral changes\nin the in situ FTIR curves, we fully understand this issue. We previously attempted to\npresent the data as contour maps to better illustrate the trends; however, the outcome\nwas unsatisfactory. This was mainly due to baseline inconsistencies in the current in\nsitu FTIR dataset (as indicated by the dashed box region in Figure R6a, where\nsubstantial baseline variations exist during the initial charge/discharge stages). If we\nforcibly normalize the baseline to generate contour plots, subtle signals—such as the\nLiF peak at ~833 cm⁻¹—would be obscured. Additionally, the baseline mismatch leads\nto difficulty in visualizing downward-oriented (i.e., product) peak evolutions in the\ncontour format (Figure R6b). It is worth noting that the data presented in Figure 7b\nhave already undergone a minimal baseline correction. More aggressive baseline\nnormalization could distort peak intensities and compromise data reliability.\nNevertheless, we have also provided individual spectra for the lithiation and delithiation\nstages below the full charge/discharge curves (Figure R7). These plots allow for more\ndistinct observation of the evolution of key characteristic peaks.\nFigure R6. (a) Stacked in situ infrared spectra of the 0.5 M-DBN electrolyte. (b)\nContour plot of the in situ infrared spectra for the 0.5 M-DBN electrolyte.\nFigure R7. Stacked plots of the in-situ infrared spectra during the plating (a) and\nstripping (b) stages.\n4. The formation mechanism of dual-layer SEI remains not very clear. It was found that\nDFOB- can be decomposed to form B-O species in the inner SEI. In principle, DFOB-\ncan decompose to form LiF in the meantime, but the inner SEI lacks LiF. Why can’t we\ndetect large amount of LiF in the inner layer?\nResponse: Thank you to the reviewer for the careful evaluation of our manuscript. We\nfully understand the concern that LiODFB decomposition might theoretically generate\nLiF. However, based on bond dissociation energy analysis, the B–O bond in the\ndifluoro(oxalato)borate anion (DFOB–) of LiODFB has a much lower bond energy\n(210.7 kJ mol–1), and in its radical form (DFOB•), the B–O bond energy can be as low\nas –50.1 kJ mol–1 (references: Nat. Energy 3, 739–746 (2018); ACS Appl. Mater.\nInterfaces 2018, 10 (3), 2469–2479). This suggests that the B–O bond is highly\nsusceptible to cleavage, especially in the radical state, leading to the formation of B–O-rich decomposition products in the SEI.\nIn contrast, the B–F bond possesses significantly higher bond dissociation\nenergy—575.3 kJ mol–1 in the anion and 236.4 kJ mol–1 in the radical—which is far\ngreater than that of the B–O bond. Therefore, the decomposition of LiODFB\npreferentially involves B–O bond cleavage and does not produce LiF in large quantities.\nTo validate this mechanism, we formulated a 1.5 M LiODFB in TEP electrolyte\nand performed sputtering XPS depth profiling of the Li metal interface. The results\nrevealed that B–O species are enriched in the inner SEI layer, while LiF is\npredominantly located in the outer SEI region. This spatial distribution pattern is fully\nconsistent with the decomposition pathway we proposed based on bond energy\ndifferences (Figure R8).\nFigure R8. (a) F1s and (b) B1s XPS spectra of the SEI formed in the 1.5M LiODFB\nTEP electrolyte.\n5. In Fig. 1f, it can be seen that the solvent TEP dominates on the Li surface. Since all\nof electrolyte components, including TEP, BF4-, DFOB- and NO3-, can decompose on\nLi surface, more TEP corresponds to more severe TEP decomposition. How can those\nminor species (anions) on Li surface can protect Li metal from reactive phosphate?\nResponse: We thank the reviewer for their careful evaluation of our manuscript. As\nshown in Table S1, TEP is indeed the dominant component in the electrolyte, and as\nthe reviewer rightly pointed out, it constitutes the major species at the interface.\nHowever, since the LUMO energy levels of the anions are lower than that of TEP,\nthe anions preferentially undergo decomposition at the interface, leading to the\nformation of an inorganic-rich interfacial layer. This layer effectively prevents direct\ncontact between TEP and the lithium metal, thereby reducing the reactivity of TEP.\nTable S1. Detailed components of the four MD models\nComponents TEP LiODFB LiBF LiNO\n4 3\n0.5 M LiODFB+0.5 M LiBF\n4\n5.9 0.5 0.5 0.5\n+0.5 M LiNO in TEP (mmol)\nReviewer #1 (Remarks to the Author):\nI thank the authors for their comprehensive revisions and additional analyses, which\nhave considerably strengthened the manuscript. In addition, I would ask the authors to\naddress the following points.\nResponse: Many thanks for the referee’s valuable suggestion, especially, the insightful\ncomments for further improving the quality of the manuscript. We have carefully\naddressed the issues as follows.\n1. The finding that OPLS-AA can describe the interaction between Li metal and Li ions\nwith accuracy comparable to DFT is indeed surprising. Could the authors clarify\nwhether the Lennard–Jones parameters for Li metal were specifically adjusted in this\nwork, or whether they were taken directly from previous literature? For the sake of\nreproducibility and transparency, it is essential that the full set of Lennard–Jones\nparameters used in the simulations be included in the Supplementary Information.\nFurthermore, the computational details of the DFT calculations performed for\nvalidation are not currently provided in the Supplementary Information. These details\nmust be explicitly included to allow the results to be properly assessed and reproduced.\nResponse: We thank the reviewer for pointing out the need for clarification regarding\nthe force field parameters and DFT calculation details. To ensure reproducibility and\ntransparency, we have included in the revised Supplementary Information (SI) a\ncomplete list and explanation of all Lennard–Jones (LJ) parameters and charge\nassignments used in this work, along with the full computational settings and key input\nfiles employed in the validating DFT calculations. Noted SI. In lines 122–136, the\nmanuscript states: “The density functional theory (DFT) calculation implemented in the\nVienna Ab initio simulation package (VASP) was conducted to study adsorption energy.\nThe exchange-correlation potential was described by using the generalized gradient\napproximation of Perdew-Burke-Ernzerhof (GGA-PBE). The projector augmented-\nwave (PAW) method was employed to treat interactions between ion cores and valence\nelectrons. We simulated the voltage effect by applying an external electric field of 0,\n−0.1 and −1 V/nm along the vacuum direction. The plane-wave cutoff energy was fixed\nto 450 eV. Given structural models were relaxed until the Hellmann–Feynman forces\nsmaller than −0.02 eV Å−1 and the change in energy smaller than 10−5 eV was attained.\nGrimme’s DFT-D3 methodology was used to describe the dispersion interactions\namong all the atoms in adsorption models. For searching the transition states, we\nemployed the nudged elastic band (NEB) method as implemented in VASP. The\nadsorption energy of these systems was calculated via following formula:\nE(binding) = E(A) + E(B) - E(AB).\nIn which, E(binding) is the adsorption energy, E(A), E(B), E(AB) are the single-\npoint energy of the various substrates.”\nTable S6. Lennard-Jones parameters[S18-S19]\natom σ (nm) ε (KJ/mol) Charge (e)\nLi+\nLi+ 0.21645 0.0764793 0.80\nBF −\nF 0.31181 0.25104 −0.3655\nF 0.31181 0.25104 −0.3655\nF 0.31181 0.25104 −0.3655\nF 0.31181 0.25104 −0.3655\nNO −\nO 0.286 0.87864 −0.4784\nO 0.286 0.87864 −0.4784\nO 0.286 0.87864 −0.4784\nDFOB−\nC 0.39967 0.359824 0.6604\nC 0.39967 0.359824 0.6604\nO 0.296 0.87864 −0.5376\nO 0.296 0.87864 −0.5376\nOS 0.300 0.71128 −0.5950\nOS 0.300 0.71128 −0.5950\nB 0.364 0.75312 1.0392\nF 0.312 0.255224 −0.4474\nF 0.312 0.255224 −0.4474\nTEP\nP 0.374 0.8368 2.1376\nO 0.296 0.87864 −0.9188\nO1 0.290 0.58576 −0.6499\nO1 0.290 0.58576 −0.6803\nO1 0.290 0.58576 −0.7062\nC 0.350 0.276144 −0.0101\nC 0.350 0.276144 −0.2477\nC 0.350 0.276144 −0.0357\nC 0.350 0.276144 −0.2471\nC 0.350 0.276144 −0.0016\nC 0.350 0.276144 −0.3275\nH 0.250 0.12552 0.1151\nH 0.250 0.12552 0.1151\nH 0.250 0.12552 0.1046\nH 0.250 0.12552 0.1046\nH 0.250 0.12552 0.1046\nH 0.250 0.12552 0.1043\nH 0.250 0.12552 0.1043\nH 0.250 0.12552 0.1069\nH 0.250 0.12552 0.1069\nH 0.250 0.12552 0.1069\nH 0.250 0.12552 0.1349\nH 0.250 0.12552 0.1349\nH 0.250 0.12552 0.1147\nH 0.250 0.12552 0.1147\nH 0.250 0.12552 0.1147\nLi_subs\nLi 0.280 9.00 0\n2. I remain unconvinced by the manuscript's conclusion that differences in anion\nadsorption strength within 1 nm of the Li metal surface directly determine the SEI\nstructure extending over tens of nanometers. Even if the adsorbed anions decompose\nduring the initial stages of SEI formation, subsequent processes such as ion transport\nare also involved in SEI growth. Therefore, it does not seem reasonable to directly link\nthe initial local adsorption state to the final SEI architecture. A clear rationale is required\nto bridge this scale gap and substantiate the proposed mechanism.\nResponse: We sincerely thank the reviewer for the careful evaluation of our manuscript\nand for the insightful and constructive comments. As you have correctly pointed out,\nthere exists a complex physicochemical process bridging the initial atomic-scale\nadsorption events and the subsequent mesoscopic formation of the SEI, and that directly\ncorrelating the two indeed requires a more rigorous justification. In our work, we\nemphasized that the initial anion adsorption and decomposition behavior establishes the\ncritical starting conditions for heterogeneous SEI nucleation and growth. We agree with\nyour viewpoint that a complete mechanism must incorporate the subsequent processes.\nNevertheless, different anion adsorption strengths result in initial SEIs with distinctly\ndifferent chemical compositions and properties. Specifically, due to the stronger\nadsorption affinity and lower LUMO levels of LiODFB and LiNO , their\n3\ndecomposition preferentially forms an inorganic-rich and compact inner layer, whereas\nthe weaker adsorption of LiBF instead favors the formation of an outer LiF layer. In\n4\nthe subsequent cycling process, the distribution state of anions is preserved, such that\nthe newly formed SEI in the fractured regions retains the original compositional\ncharacteristics: a B–O- and Li₃N-rich inner layer, and a LiF-dominated outer layer. This\nmechanism effectively ensures the overall uniformity and structural stability of the SEI.\nAt the same time, we acknowledge that our theoretical simulations were conducted\nin relatively small simulation cells. If extended to the scale of a full practical cell, such\nconcentration gradients would likely become more pronounced, thereby further\nenhancing the stratified SEI structure. Furthermore, our experimental\ncharacterizations—including sputtering XPS (Figure 2c-e), TOF-SIMS (Figure 2h-i),\nand Cryo-TEM (Figure S29)—have confirmed the successful construction of the\nenvisioned triple-layer SEI and demonstrated that it remains stable throughout extended\ncycling (Figure R1).\nFigure 2c-e. XPS depth profiles were determined of (c) F1s, (d) N1s, and (e) B1s of\nlithium metal anodes recovered from full cell cycled with 0.5 M-DBN electrolyte at\nsputtering times of 0, 50, and 100 s, respectively.\nFigure 2h-i. Time-of-flight Secondary Ion Mass Spectrometry (TOF-SIMS) depth\ndistributions (h) and three-dimensional images of various species (i) within a 0.5 M-\nDBN electrolyte.\nFigure S29. (a) Cryo-TEM image of the SEI formed on lithium metal using the 0.5 M-\nDBN electrolyte; the colored rectangles indicate the presence of LiF (yellow), Li N\n3\n(blue), and LiBO (maroon). The marked regions in the cryo-TEM image (a) are\n2\nenlarged in (b–d), showing the corresponding nanocrystals of LiF (b), Li N (c),\ninterface in Li||NCM811 cells using the 0.5 M-DBN electrolyte after 1 cycle (a, b),\n20\ncycles (c, d), and 100 cycles (e, f).\nMinor comment: Please change the frame color in Fig. S29b to yellow.\nResponse: We are grateful to the reviewer for the insightful comments. In accordance\nwith the suggestion, the frame color in Figure S29b has been corrected.\nFigure S29. (a) Cryo-TEM image of the SEI formed on lithium metal using the 0.5 M-\nDBN electrolyte; the colored rectangles indicate the presence of LiF (yellow), Li N\n3\n(blue), and LiBO (maroon). The marked regions in the cryo-TEM image (a) are\n2\nenlarged in (b–d), showing the corresponding nanocrystals of LiF (b), Li N (c),\n2"}, {"round": 2, "reviewer_comments": {"reviewer_1": "I thank the authors for their comprehensive revisions and additional analyses, which have considerably strengthened the\nmanuscript. In addition, I would ask the authors to address the following points.\n1. The finding that OPLS-AA can describe the interaction between Li metal and Li ions with accuracy comparable to DFT is\nindeed surprising. Could the authors clarify whether the Lennard–Jones parameters for Li metal were specifically adjusted in\nthis work, or whether they were taken directly from previous literature? For the sake of reproducibility and transparency, it is\nessential that the full set of Lennard–Jones parameters used in the simulations be included in the Supplementary\nInformation. Furthermore, the computational details of the DFT calculations performed for validation are not currently\nprovided in the Supplementary Information. These details must be explicitly included to allow the results to be properly\nassessed and reproduced.\n2. I remain unconvinced by the manuscript's conclusion that differences in anion adsorption strength within 1 nm of the Li\nmetal surface directly determine the SEI structure extending over tens of nanometers. Even if the adsorbed anions\ndecompose during the initial stages of SEI formation, subsequent processes such as ion transport are also involved in SEI\ngrowth. Therefore, it does not seem reasonable to directly link the initial local adsorption state to the final SEI architecture.\nA clear rationale is required to bridge this scale gap and substantiate the proposed mechanism.\nMinor comment: Please change the frame color in Fig. S29b to yellow.", "reviewer_2": "I have no more questions.", "reviewer_3": "The authors provide detailed response to the reviewer’s comment. The “voltage-driven anion migration” strategy was further\ndiscussed in details in the revised manuscript, which may promote the development of phosphate electrolyte. Some\nremaining points need to be addressed.\n1. The experimental evidence of dual-layer SEI is not convincing. XPS was used in the manuscript to verify the dual-layer\nfeature of SEI, which is a key finding of this work. However, XPS results can be misleading sometime, where the sputtering\nis not homogeneous among different components, undermining the claim of dual-layer SEI formation. To clearly demonstrate\nthe dual-layer feature of SEI, cryogenic electron characterization is suggested.\n2. The applicability of the so-called “voltage-driven anion migration” strategy in Li metal full cell should be fully charified.\nThis manuscript reports a voltage-based strategy to regulate SEI. However, it is known that Li chemically reacts with\nelectrolyte to form an SEI before cycling, i.e., before applying the electric field. Such a chemically formed SEI may\nsubstantially affect the “voltage-driven anion migration” strategy to regulate the SEI structure. The compasion between\nreported electrolyte and traditional carbonate electrolyte is not very convincing.\n3. The FTIR results need more detailed discussion. Please provide references for each peak’s attribution in the profiles.\nIn particular, the peak position of LiF needs to be checked. Meanwhile, the curve change during charge and discharge process\nis not very clear, since most of curves overlap in Fig. 2a.\n4. The formation mechanism of dual-layer SEI remains not very clear. It was found that DFOB- can be decomposed to form\nB-O species in the inner SEI. In principle, DFOB- can decompose to form LiF in the meantime, but the inner SEI lacks LiF.\nWhy can’t we detect large amount of LiF in the inner layer?\n5. In Fig. 1f, it can be seen that the solvent TEP dominates on the Li surface. Since all of electrolyte components, including\nTEP, BF4-, DFOB- and NO3-, can decompose on Li surface, more TEP corresponds to more severe TEP decomposition.\nHow can those minor species (anions) on Li surface can protect Li metal from reactive phosphate?"}, "author_response": "Thank you for your email informing us of your editorial decision on our manuscript\n(NCOMMS-25-16384A). According to the instructions in your email, we have\ncarefully considered the helpful criticism from the referees and revised the manuscript\naccordingly. We provide a detailed response to the referees' criticism below. The\nchanges made during revision are highlighted in yellow and the file is uploaded\nseparately for review only.\nReviewer #1 (Remarks to the Author):\nWhile the authors have addressed some initial concerns, critical issues remain\nunresolved, particularly regarding the mechanistic interpretation and its theoretical\nfoundation. Without substantial revisions, the manuscript is not suitable for publication.\nResponse: Many thanks for the referee’s valuable suggestion, especially, the insightful\ncomments for further improving the quality of the manuscript. We have carefully\naddressed the issues as follows.\n1. The use of the OPLS-AA force field for lithium metal, combined with a strong\nelectric field of -1.0 V/nm, remains problematic. Although the authors cite precedent,\nthey do not demonstrate that the observed ion migration trends persist under more\nrealistic conditions. As the force field does not capture specific adsorption at the Li\nsurface, the model likely exaggerates interfacial differences under the applied field.\nMore seriously, despite the strong electric field, the anion distributions are noisy, and\nthe field-induced effects cannot be clearly distinguished, raising serious concerns about\nthe validity of the proposed mechanism. Furthermore, the SEI depth profiles, which\nthe authors emphasize as key experimental evidence, reflect cumulative decomposition and\ntransport processes over extended cycling and therefore do not directly validate the\nsimulated initial anion distributions.\nResponse: We sincerely thank the reviewer for the thorough evaluation of our\nmanuscript. Regarding the concern about employing the OPLS-AA force field to\ndescribe the interaction between lithium metal and lithium ions in lithium-metal-based\nsystems, we have given this issue further consideration and performed additional\nvalidation calculations. Molecular dynamics (MD) simulations using the OPLS-AA\nforce field in Gromacs confirmed that this force field indeed incorporates the interaction\nbetween lithium metal and lithium ions. Specifically, the extracted interaction energies\nwere 320.56 kJ/mol (≈3.32 eV) under zero external electric field and 410.39 kJ/mol (\n≈4.25 eV) under an applied external electric field. To further validate the reliability of\nthe OPLS-AA force field description, we conducted density functional theory (DFT)\ncalculations of the adsorption energies of three lithium salts—LiBF , LiODFB, and\n3\nand −1 V/nm. Notably, under 0 V/nm and −1 V/nm (corresponding to the MD\nsimulation conditions), the average adsorption energies of the three salts were 3.57 eV\nand 4.47 eV, respectively. These DFT results are in good agreement with the MD-\nderived interaction energies (3.32 eV and 4.25 eV), indicating that the OPLS-AA force\nfield can reasonably capture the interaction between lithium metal and lithium ions to\na certain extent. Moreover, the DFT results clearly reveal (Figure S14-S15) that the\nadsorption energies of LiODFB and LiNO₃ on lithium metal surfaces are significantly\nhigher than that of LiBF . Therefore, within the system, LiODFB and LiNO are more\n4 3\nlikely to be enriched at the lithium metal surface. Noted. In lines 112–113, the\nmanuscript states: “Meanwhile, the DFT calculations (Fig. S14-15) indicate that\nlithium metal exhibits a preferential adsorption toward LiNO and LiODFB.”\nnote that OPLS-AA is able to reasonably describe the interactions between the Li metal\nanode and Li⁺ ions. Specifically, in the absence of an applied electric field, the Li metal–\nLi⁺ interaction energy predicted by this force field is approximately 320.56 kJ/mol (≈\n3.32 eV), which is very close to the average adsorption energy (3.57 eV) we obtained\nfor the three lithium salts under the same conditions. Similarly, under an applied electric\nfield, the interaction energy is about 418.39 kJ/mol (≈ 4.33 eV), again in good\nagreement with the average adsorption energy (4.47 eV) from our DFT calculations.\nGiven the close correspondence between the key interaction energies obtained via\nOPLS-AA and the DFT-calculated average adsorption energies—both with and without\nan electric field—we are confident that the MD simulation results based on this force\nfield are accurate and reliable.\nFurthermore, following the reviewers' suggestions, we conducted experiments to\ndirectly verify the distribution of the initial anions. First, we employed quantitative\nnuclear magnetic resonance (NMR) spectroscopy to monitor the variation in the\nconcentrations of two lithium salts on the anode side during charging. Specifically, we\nassembled a Li–Li symmetric cell and applied a constant current of 0.2 mA.\nFigureS19a-b illustrate the schematic diagram of the experimental setup, which\nsimulates the charging process of a full cell. After approximately 10 min and 30 min of\ncurrent application, 0.3 mL of electrolyte was collected from the vicinity of the lithium\ndeposition electrode and mixed with 0.3 mL of deuterated chloroform (CDCl ). The\n3\nresulting solutions were analyzed by ¹¹B NMR spectroscopy, and the relative contents\nof the two anions were determined by integrating the corresponding peaks.As shown in\nFigures S19c-d, with increasing lithiation time, the concentration of ODFB⁻ anions\nnear the lithium-deposition side (anode side) gradually increased, whereas the\nconcentration of BF ⁻ anions decreased. This trend is consistent with the predictions\n4\nfrom our molecular dynamics (MD) simulations.\nWe then further examined the nanostructure of the solid electrolyte interphase (SEI)\nat the electrode/electrolyte interface using cryogenic transmission electron microscopy\n(Cryo-TEM) (Figure S29). The SEI was found to feature an outer layer rich in LiF and\nan inner layer containing Li N and B–O species, which provides additional evidence\n3\nsupporting our design concept.\nAs the reviewers' comment, the sputtering XPS depth profiling primarily reflect\nthe cumulative decomposition and transport processes of the electrolyte after prolonged\ncycling, and therefore cannot directly verify the initial anion distribution predicted by\nthe simulations. To address this, we supplemented our study with depth-profiling XPS\nanalyses after different numbers of cycles (1, 20, and 100) (Figure R1). The results\nshowed that, regardless of the number of cycles, the anode interface consistently\nmaintained a nanostructure with an LiF-rich outer layer and an Li N/B–O-containing\n3\ninner layer. This clearly demonstrates that the SEI structure formed in our experimental\nsamples retains excellent integrity throughout cycling.\nFigure S14. The adsorption energy calculations of three lithium salts — LiBF (a, d,\n4\ng), LiODFB (b, e, h), and LiNO (c, f, i) — under three applied electric field conditions:\n3\nFigure S15. Adsorption energies of different lithium salts under various voltage\nconditions\nFigure S19. (a, b) Experimental setup for ion migration studies using 0.5 M-DBN\nelectrolyte. (c) Quantitative NMR results showing the electrolyte composition near the\nlithium deposition electrode. (d) Relative compositional changes in the electrolyte near\nthe lithium deposition electrode.\nFigure S29. (a) Cryo-TEM image of the SEI formed on lithium metal using the 0.5 M-\nDBN electrolyte; the colored rectangles indicate the presence of LiF (yellow), Li N\n3\n(blue), and LiBO (maroon). The marked regions in the cryo-TEM image (a) are\n2\nenlarged in (b–d), showing the corresponding nanocrystals of LiF (b), Li N (c),\ninterface in Li||NCM811 cells using the 0.5 M-DBN electrolyte after 1 cycle (a, b),\n20\ncycles (c, d), and 100 cycles (e, f).\n2. The proposed link between the initial anion distribution and the long-term SEI\nlayering remains speculative. Continued LiODFB consumption does not ensure\nsustained interfacial enrichment. Over extended cycling, differences in anion transport\nproperties are likely to dominate, and without direct evidence, the proposed mechanism\nis not convincingly supported.\nResponse: We sincerely thank the reviewer for the detailed review of our manuscript.\nAs the reviewer rightly pointed out, the growth of the solid electrolyte interphase (SEI)\nis indeed a continuous process. However, the consumption of the lithium salt (in this\ncase, lithium difluoro(oxalato)borate, LiODFB) does not occur extensively or\nuniformly throughout cycling. The critical stage for SEI formation is during the initial\nfew cycles of battery operation, during which LiODFB participates in reactions to form\na uniform and stable SEI layer.\nGiven the electronically insulating nature of SEI, once a complete initial layer is\nestablished, electrons can no longer penetrate the interface to trigger further\ndecomposition of the lithium salt. As a result, in subsequent cycles, the evolution of the\nSEI is primarily driven by the volume changes associated with lithium plating/stripping,\nwhich may locally damage the SEI. In such cases, the electrolyte decomposes at these\ndamaged sites to form new SEI layers for “self-repair.”\nThe SEI formed in our study demonstrates excellent elasticity, allowing it to\neffectively accommodate volume fluctuations during cycling and thereby significantly\nreduce lithium salt consumption from repeated SEI rupture and repair (Figure R2).\nFurthermore, we conducted XPS characterization of the anode interface after different\ncycle numbers, and the results show that the SEI composition and structure remain\nlargely consistent throughout cycling. This provides strong evidence supporting our\nconclusion that once the initial SEI is stably formed, the subsequent lithium salt\nconsumption is substantially reduced (Figure R1).\nFigure R2. DMT modulus distribution of 1.5 M-N (a) and 0.5 M-DBN (b) SEI.\nFigure R1. F 1s (a, c, e) and N 1s (b, d, f) XPS spectra of the lithium metal anode\ninterface in Li||NCM811 cells using the 0.5 M-DBN electrolyte after 1 cycle (a, b),\n20\ncycles (c, d), and 100 cycles (e, f).\n3. The two additional binary systems (LiDOB + LiFSI and LiNO + LiTFSI) provide\n3\nsome insight but are insufficient to establish the generality of the proposed design\nprinciple. Attributing inner-layer formation to LiDOB or LiNO decomposition based\n3\nsolely on Li O detection is unconvincing, as Li O can also result from the\n2 2\ndecomposition of imide-based salts such as LiFSI or LiTFSI. This overlap in\ndecomposition pathways limits the interpretability of the experimental evidence.\nWithout further mechanistic clarification, the claim of a general design principle\nremains premature.\nResponse: We thank the reviewer for the careful evaluation of our manuscript. We fully\nagree with the reviewer’s opinion that, in our previous study involving the LiNO +\n2 3\nmay be inappropriate.\nTo further validate the universality of the proposed mechanism—namely, that\ndifferences in anion binding energy drive their competitive adsorption and preferential\ndecomposition near the lithium metal surface—we specifically designed two entirely\nnew electrolyte systems, distinct from those used in previous studies:\n1. 0.5 M LiBOB / 0.5 M LiPF in TEP: In this system, the binding energy of\n6\n2. 0.5 M LiNO / 0.5 M LiPF in TEP: In this system, the binding energy of LiNO\n3 6 3\nis higher than that of LiPF .\ninterface to investigate the SEI structures formed in each system:\n For the LiBOB/LiPF 6 system: According to our theory, LiBOB, having a\nhigher binding energy, preferentially enriches near the lithium metal surface and\ndecomposes first, forming a B–O-rich inner SEI. In contrast, LiPF , with a lower\n6\nbinding energy, tends to form an outer SEI enriched in LiF. This layered SEI\nstructure was confirmed by depth-profiling XPS analysis, which revealed a B–\nO-rich inner layer and a LiF-rich outer layer (Figure R3a-b).\n For the LiNO 3 /LiPF 6 system: Similarly, based on our proposed mechanism,\nLiNO (higher binding energy) decomposes preferentially near the lithium\n3\nsurface, resulting in a significantly higher Li N content in the inner SEI\n3\ncompared to the outer layer. Meanwhile, LiPF contributes mainly to the\n6\nformation of a LiF-rich outer SEI. The depth-profiling XPS results of this\nsystem were also consistent with the predicted spatial distribution (Figure R3c-\nd).\nOverall, these experimental results provide strong evidence that the anion gradient\ndistribution mechanism we have uncovered is not a phenomenon specific to a particular\nelectrolyte system, but rather a universally applicable interfacial regulation principle.\nFigure R3. (a) F1s and (b) B1s XPS spectra of the SEI formed in the 0.5 M LiBOB +\n0.5 M LiPF TEP electrolyte. (c) F1s and (d) N1s XPS spectra of the SEI formed in the\n6\n0.5 M LiNO + 0.5 M LiPF TEP electrolyte.\n3 6\nReviewer #3 (Remarks to the Author):\nThe authors provide detailed response to the reviewer’s comment. The “voltage-driven\nanion migration” strategy was further discussed in details in the revised manuscript,\nwhich may promote the development of phosphate electrolyte. Some remaining points\nneed to be addressed.\nResponse: Many thanks for the referee’s valuable suggestion, especially, the insightful\ncomments for further improving the quality of the manuscript. We have carefully\naddressed the issues as follows.\n1. The experimental evidence of dual-layer SEI is not convincing. XPS was used in the\nmanuscript to verify the dual-layer feature of SEI, which is a key finding of this work.\nHowever, XPS results can be misleading sometime, where the sputtering is not\nhomogeneous among different components, undermining the claim of dual-layer SEI\nformation. To clearly demonstrate the dual-layer feature of SEI, cryogenic electron\ncharacterization is suggested.\nResponse: We appreciate the reviewer’s careful evaluation of our manuscript. As\nreviewer rightly pointed out, XPS analysis may suffer from reliability issues due to\npossible artifacts introduced during sputtering. To address this concern, we followed\nthe reviewer’s suggestion and employed cryogenic transmission electron microscopy\n(Cryo-TEM) for further characterization.\nAs shown in Figure S29a, the SEI layer formed using the 0.5 M-DBN electrolyte\nexhibits a dense, inorganic-rich morphology on the lithium metal surface. High-\nresolution analysis identified several crystalline components within the SEI, including\nthe LiF (111) plane with a lattice spacing of 0.233 nm, the Li N (220) plane (0.195 nm),\n3\nand the LiBO (022) plane (0.252 nm), as displayed in Figures S29b–d. These\n2\nobservations reveal a distinct layered structure: the inner SEI region is primarily\ncomposed of Li₃N and LiBO , while the outer layer is dominated by LiF. This spatial\n2\ndistribution is fully consistent with our earlier XPS depth profiling and TOF-SIMS\nresults. Noted. In lines 204–205, the manuscript states: “Furthermore, the presence of\nthis bilayer SEI structure was corroborated by cryo-electron microscopy observations\n(Figures S29).”\nSuch a stratified SEI structure offers synergistic advantages: the outer LiF layer\nprovides high interfacial rigidity that effectively suppresses lithium dendrite growth,\nwhereas inner compounds like Li N exhibit high Li⁺ conductivity, enhancing ionic\n3\ntransport through the SEI. Additionally, the presence of B–O bonded components\ncontributes to mechanical flexibility. Together, these features enable the formation of a\nrobust SEI with both high ionic conductivity and mechanical adaptability, facilitating\ndense lithium deposition and significantly improving plating/stripping efficiency.\nFigure S29. (a) Cryo-TEM image of the SEI formed on lithium metal using the 0.5 M-\nDBN electrolyte; the colored rectangles indicate the presence of LiF (yellow), Li N\n3\n(blue), and LiBO (maroon). The marked regions in the cryo-TEM image (a) are\n2\nenlarged in (b–d), showing the corresponding nanocrystals of LiF (b), Li N (c),\n2\n2. The applicability of the so-called “voltage-driven anion migration” strategy in Li\nmetal full cell should be fully charified. This manuscript reports a voltage-based\nstrategy to regulate SEI. However, it is known that Li chemically reacts with electrolyte\nto form an SEI before cycling, i.e., before applying the electric field. Such a chemically\nformed SEI may substantially affect the “voltage-driven anion migration” strategy to\nregulate the SEI structure. The compasion between reported electrolyte and traditional\ncarbonate electrolyte is not very convincing.\nResponse: We sincerely thank the reviewer for the thoughtful and detailed comments.\nAs correctly noted, prior to battery cycling, lithium metal spontaneously reacts with the\nelectrolyte to form a chemically driven solid electrolyte interphase (SEI) that is\nextremely thin, structurally loose, and disordered (ACS Energy Letters, 2024, 9(11),\n5268–5278). Although this chemically formed SEI can offer an initial barrier against\nfurther electrolyte decomposition, its high porosity and limited blocking ability result\nin poor stability, rendering it vulnerable to subsequent electrochemical reconstruction.\nWith the application of external potential during charge/discharge processes, further\nelectrochemical reduction of the electrolyte occurs, leading to the formation of a second\nSEI layer on top of the initial chemical SEI. This electrochemically dominated SEI is\ntypically denser and more functionalized. The electrochemical formation stage largely\ndetermines the final thickness, compactness, and composition of the SEI, which in turn\ncritically influences long-term cycling performance (Batteries & Supercaps, 2021, 4,\n909; Small, 2020, 16(23), e2000756; Energy Environ. Sci., DOI:\n10.1039/D5EE01030F).\nAs the reviewers' suggestion, we prepared a control electrolyte system consisting\nof 0.75 M LiODFB/EC-DMC and 0.75 M LiBF /EC-DMC (it should be noted that\n3\nsolvent). The comparative experiments were conducted under two different conditions:\n4.3 V at 1C and 4.5 V at 0.5C. As shown in Figure R4, under the 4.3 V, 1C condition,\nthe 0.5 M-DBN electrolyte exhibited excellent cycling stability, retaining 87.90% of its\ninitial capacity after 900 cycles. In contrast, the capacity retention of the 0.75 M\nLiODFB + 0.75 M LiBF /EC-DMC electrolyte was only 36.97%. Under the higher\n4\nvoltage condition of 4.5 V, the 0.5 M-DBN electrolyte maintained 90.91% capacity\nretention after 600 cycles.\nFurthermore, to investigate the influence of different electrolyte systems,\ncomparative experiments were performed using lithium sulfonylimide-based\nelectrolytes, including 1.5 M LiFSI/TEP, 1.5 M LiTFSI/TEP, 1.5 M LiFSI/EC-DMC,\nand 1.5 M LiTFSI/EC-DMC.\nThe results showed that all four systems exhibited insufficient cycling\nperformance. Specifically, 1.5 M LiFSI/EC-DMC displayed a voltage plateau\nstagnation during the initial charging stage, likely due to severe oxidative\ndecomposition (Figure R5a). Although 1.5 M LiTFSI/EC-DMC completed charge–\ndischarge cycles, its capacity rapidly decayed within 100 cycles (Figure R5b). Both\n1.5 M LiFSI/TEP and 1.5 M LiTFSI/TEP showed low initial capacities and rapid\ncapacity fading (Figures R5c-d).\nThese comparative results clearly demonstrate the significant advantages of the\n0.5 M-DBN electrolyte in cycling stability and high-voltage tolerance, strongly\nvalidating our electrolyte design strategy.\nFigure R4. Comparison of the cycling performance of the two electrolytes under (a)\n4.3 V at 1C and (b) 4.5 V at 0.5C conditions.\nFigure R5. (a) Time–voltage profile of the first cycle for the 1.5 M LiFSI/EC-DMC\nelectrolyte. (b) Cycling performance of the 1.5 M LiTFSI/EC-DMC electrolyte under\n4.3 V 1C and 4.5 V 0.5C conditions. (c) Cycling performance of the 1.5 M LiFSI/TEP\nelectrolyte under 4.3 V 1C and 4.5 V 0.5C conditions. (d) Cycling performance of the\n1.5 M LiTFSI/TEP electrolyte under 4.3 V 1C and 4.5 V 0.5C conditions. (e) Linear\nsweep voltammetry (LSV) curves of the four electrolytes.\n3. The FTIR results need more detailed discussion. Please provide references for each\npeak’s attribution in the profiles. In particular, the peak position of LiF needs to be\nchecked. Meanwhile, the curve change during charge and discharge process is not very\nclear, since most of curves overlap in Fig. 2a.\nResponse: We sincerely thank the reviewer for the careful evaluation of our manuscript.\nIn light of your professional guidance, we have we have conducted a more detailed\ndiscussion of the FTIR results in the revised manuscript “As shown in Figures 2a and\nS20, during the lithium deposition stage (i.e., the desolvation process), the pronounced\ndepletion peaks at 977.90 cm−1 (P–O–C) and 1255.31 cm−1 (P=O) correspond to the\ngradual consumption of solvated TEP, while the emerging peak at 958.62 cm−1 (P–O–\nC) indicates the release of desolvated solvent molecules at the electrode interface14.\nFigure S19 presents the in-situ infrared spectra of the 1 M LiPF /EC–DMC electrolyte\n6\nsystem along with the corresponding spectral analysis.” in Lines 148-154. Furthermore,\nthe assignments of each peak in the in situ FTIR spectra and the corresponding\nreferences are as follows: the characteristic peak of LiF appears at 837 cm⁻¹ (Reference:\nNat. Commun. 13, 1398 (2022)); the peaks corresponding to P–O–C, B–O–R, and P=O\nare located at approximately 980 cm⁻¹, 1020 cm⁻¹, and 1260 cm⁻¹, respectively\n(Reference: Energy Storage Materials 42, 628–635 (2021)).\nRegarding the reviewer’s concern about the insufficiently clear spectral changes\nin the in situ FTIR curves, we fully understand this issue. We previously attempted to\npresent the data as contour maps to better illustrate the trends; however, the outcome\nwas unsatisfactory. This was mainly due to baseline inconsistencies in the current in\nsitu FTIR dataset (as indicated by the dashed box region in Figure R6a, where\nsubstantial baseline variations exist during the initial charge/discharge stages). If we\nforcibly normalize the baseline to generate contour plots, subtle signals—such as the\nLiF peak at ~833 cm⁻¹—would be obscured. Additionally, the baseline mismatch leads\nto difficulty in visualizing downward-oriented (i.e., product) peak evolutions in the\ncontour format (Figure R6b). It is worth noting that the data presented in Figure 7b\nhave already undergone a minimal baseline correction. More aggressive baseline\nnormalization could distort peak intensities and compromise data reliability.\nNevertheless, we have also provided individual spectra for the lithiation and delithiation\nstages below the full charge/discharge curves (Figure R7). These plots allow for more\ndistinct observation of the evolution of key characteristic peaks.\nFigure R6. (a) Stacked in situ infrared spectra of the 0.5 M-DBN electrolyte. (b)\nContour plot of the in situ infrared spectra for the 0.5 M-DBN electrolyte.\nFigure R7. Stacked plots of the in-situ infrared spectra during the plating (a) and\nstripping (b) stages.\n4. The formation mechanism of dual-layer SEI remains not very clear. It was found that\nDFOB- can be decomposed to form B-O species in the inner SEI. In principle, DFOB-\ncan decompose to form LiF in the meantime, but the inner SEI lacks LiF. Why can’t we\ndetect large amount of LiF in the inner layer?\nResponse: Thank you to the reviewer for the careful evaluation of our manuscript. We\nfully understand the concern that LiODFB decomposition might theoretically generate\nLiF. However, based on bond dissociation energy analysis, the B–O bond in the\ndifluoro(oxalato)borate anion (DFOB–) of LiODFB has a much lower bond energy\n(210.7 kJ mol–1), and in its radical form (DFOB•), the B–O bond energy can be as low\nas –50.1 kJ mol–1 (references: Nat. Energy 3, 739–746 (2018); ACS Appl. Mater.\nInterfaces 2018, 10 (3), 2469–2479). This suggests that the B–O bond is highly\nsusceptible to cleavage, especially in the radical state, leading to the formation of B–O-rich decomposition products in the SEI.\nIn contrast, the B–F bond possesses significantly higher bond dissociation\nenergy—575.3 kJ mol–1 in the anion and 236.4 kJ mol–1 in the radical—which is far\ngreater than that of the B–O bond. Therefore, the decomposition of LiODFB\npreferentially involves B–O bond cleavage and does not produce LiF in large quantities.\nTo validate this mechanism, we formulated a 1.5 M LiODFB in TEP electrolyte\nand performed sputtering XPS depth profiling of the Li metal interface. The results\nrevealed that B–O species are enriched in the inner SEI layer, while LiF is\npredominantly located in the outer SEI region. This spatial distribution pattern is fully\nconsistent with the decomposition pathway we proposed based on bond energy\ndifferences (Figure R8).\nFigure R8. (a) F1s and (b) B1s XPS spectra of the SEI formed in the 1.5M LiODFB\nTEP electrolyte.\n5. In Fig. 1f, it can be seen that the solvent TEP dominates on the Li surface. Since all\nof electrolyte components, including TEP, BF4-, DFOB- and NO3-, can decompose on\nLi surface, more TEP corresponds to more severe TEP decomposition. How can those\nminor species (anions) on Li surface can protect Li metal from reactive phosphate?\nResponse: We thank the reviewer for their careful evaluation of our manuscript. As\nshown in Table S1, TEP is indeed the dominant component in the electrolyte, and as\nthe reviewer rightly pointed out, it constitutes the major species at the interface.\nHowever, since the LUMO energy levels of the anions are lower than that of TEP,\nthe anions preferentially undergo decomposition at the interface, leading to the\nformation of an inorganic-rich interfacial layer. This layer effectively prevents direct\ncontact between TEP and the lithium metal, thereby reducing the reactivity of TEP.\nTable S1. Detailed components of the four MD models\nComponents TEP LiODFB LiBF LiNO\n4 3\n0.5 M LiODFB+0.5 M LiBF\n4\n5.9 0.5 0.5 0.5\n+0.5 M LiNO in TEP (mmol)\nReviewer #1 (Remarks to the Author):\nI thank the authors for their comprehensive revisions and additional analyses, which\nhave considerably strengthened the manuscript. In addition, I would ask the authors to\naddress the following points.\nResponse: Many thanks for the referee’s valuable suggestion, especially, the insightful\ncomments for further improving the quality of the manuscript. We have carefully\naddressed the issues as follows.\n1. The finding that OPLS-AA can describe the interaction between Li metal and Li ions\nwith accuracy comparable to DFT is indeed surprising. Could the authors clarify\nwhether the Lennard–Jones parameters for Li metal were specifically adjusted in this\nwork, or whether they were taken directly from previous literature? For the sake of\nreproducibility and transparency, it is essential that the full set of Lennard–Jones\nparameters used in the simulations be included in the Supplementary Information.\nFurthermore, the computational details of the DFT calculations performed for\nvalidation are not currently provided in the Supplementary Information. These details\nmust be explicitly included to allow the results to be properly assessed and reproduced.\nResponse: We thank the reviewer for pointing out the need for clarification regarding\nthe force field parameters and DFT calculation details. To ensure reproducibility and\ntransparency, we have included in the revised Supplementary Information (SI) a\ncomplete list and explanation of all Lennard–Jones (LJ) parameters and charge\nassignments used in this work, along with the full computational settings and key input\nfiles employed in the validating DFT calculations. Noted SI. In lines 122–136, the\nmanuscript states: “The density functional theory (DFT) calculation implemented in the\nVienna Ab initio simulation package (VASP) was conducted to study adsorption energy.\nThe exchange-correlation potential was described by using the generalized gradient\napproximation of Perdew-Burke-Ernzerhof (GGA-PBE). The projector augmented-\nwave (PAW) method was employed to treat interactions between ion cores and valence\nelectrons. We simulated the voltage effect by applying an external electric field of 0,\n−0.1 and −1 V/nm along the vacuum direction. The plane-wave cutoff energy was fixed\nto 450 eV. Given structural models were relaxed until the Hellmann–Feynman forces\nsmaller than −0.02 eV Å−1 and the change in energy smaller than 10−5 eV was attained.\nGrimme’s DFT-D3 methodology was used to describe the dispersion interactions\namong all the atoms in adsorption models. For searching the transition states, we\nemployed the nudged elastic band (NEB) method as implemented in VASP. The\nadsorption energy of these systems was calculated via following formula:\nE(binding) = E(A) + E(B) - E(AB).\nIn which, E(binding) is the adsorption energy, E(A), E(B), E(AB) are the single-\npoint energy of the various substrates.”\nTable S6. Lennard-Jones parameters[S18-S19]\natom σ (nm) ε (KJ/mol) Charge (e)\nLi+\nLi+ 0.21645 0.0764793 0.80\nBF −\nF 0.31181 0.25104 −0.3655\nF 0.31181 0.25104 −0.3655\nF 0.31181 0.25104 −0.3655\nF 0.31181 0.25104 −0.3655\nNO −\nO 0.286 0.87864 −0.4784\nO 0.286 0.87864 −0.4784\nO 0.286 0.87864 −0.4784\nDFOB−\nC 0.39967 0.359824 0.6604\nC 0.39967 0.359824 0.6604\nO 0.296 0.87864 −0.5376\nO 0.296 0.87864 −0.5376\nOS 0.300 0.71128 −0.5950\nOS 0.300 0.71128 −0.5950\nB 0.364 0.75312 1.0392\nF 0.312 0.255224 −0.4474\nF 0.312 0.255224 −0.4474\nTEP\nP 0.374 0.8368 2.1376\nO 0.296 0.87864 −0.9188\nO1 0.290 0.58576 −0.6499\nO1 0.290 0.58576 −0.6803\nO1 0.290 0.58576 −0.7062\nC 0.350 0.276144 −0.0101\nC 0.350 0.276144 −0.2477\nC 0.350 0.276144 −0.0357\nC 0.350 0.276144 −0.2471\nC 0.350 0.276144 −0.0016\nC 0.350 0.276144 −0.3275\nH 0.250 0.12552 0.1151\nH 0.250 0.12552 0.1151\nH 0.250 0.12552 0.1046\nH 0.250 0.12552 0.1046\nH 0.250 0.12552 0.1046\nH 0.250 0.12552 0.1043\nH 0.250 0.12552 0.1043\nH 0.250 0.12552 0.1069\nH 0.250 0.12552 0.1069\nH 0.250 0.12552 0.1069\nH 0.250 0.12552 0.1349\nH 0.250 0.12552 0.1349\nH 0.250 0.12552 0.1147\nH 0.250 0.12552 0.1147\nH 0.250 0.12552 0.1147\nLi_subs\nLi 0.280 9.00 0\n2. I remain unconvinced by the manuscript's conclusion that differences in anion\nadsorption strength within 1 nm of the Li metal surface directly determine the SEI\nstructure extending over tens of nanometers. Even if the adsorbed anions decompose\nduring the initial stages of SEI formation, subsequent processes such as ion transport\nare also involved in SEI growth. Therefore, it does not seem reasonable to directly link\nthe initial local adsorption state to the final SEI architecture. A clear rationale is required\nto bridge this scale gap and substantiate the proposed mechanism.\nResponse: We sincerely thank the reviewer for the careful evaluation of our manuscript\nand for the insightful and constructive comments. As you have correctly pointed out,\nthere exists a complex physicochemical process bridging the initial atomic-scale\nadsorption events and the subsequent mesoscopic formation of the SEI, and that directly\ncorrelating the two indeed requires a more rigorous justification. In our work, we\nemphasized that the initial anion adsorption and decomposition behavior establishes the\ncritical starting conditions for heterogeneous SEI nucleation and growth. We agree with\nyour viewpoint that a complete mechanism must incorporate the subsequent processes.\nNevertheless, different anion adsorption strengths result in initial SEIs with distinctly\ndifferent chemical compositions and properties. Specifically, due to the stronger\nadsorption affinity and lower LUMO levels of LiODFB and LiNO , their\n3\ndecomposition preferentially forms an inorganic-rich and compact inner layer, whereas\nthe weaker adsorption of LiBF instead favors the formation of an outer LiF layer. In\n4\nthe subsequent cycling process, the distribution state of anions is preserved, such that\nthe newly formed SEI in the fractured regions retains the original compositional\ncharacteristics: a B–O- and Li₃N-rich inner layer, and a LiF-dominated outer layer. This\nmechanism effectively ensures the overall uniformity and structural stability of the SEI.\nAt the same time, we acknowledge that our theoretical simulations were conducted\nin relatively small simulation cells. If extended to the scale of a full practical cell, such\nconcentration gradients would likely become more pronounced, thereby further\nenhancing the stratified SEI structure. Furthermore, our experimental\ncharacterizations—including sputtering XPS (Figure 2c-e), TOF-SIMS (Figure 2h-i),\nand Cryo-TEM (Figure S29)—have confirmed the successful construction of the\nenvisioned triple-layer SEI and demonstrated that it remains stable throughout extended\ncycling (Figure R1).\nFigure 2c-e. XPS depth profiles were determined of (c) F1s, (d) N1s, and (e) B1s of\nlithium metal anodes recovered from full cell cycled with 0.5 M-DBN electrolyte at\nsputtering times of 0, 50, and 100 s, respectively.\nFigure 2h-i. Time-of-flight Secondary Ion Mass Spectrometry (TOF-SIMS) depth\ndistributions (h) and three-dimensional images of various species (i) within a 0.5 M-\nDBN electrolyte.\nFigure S29. (a) Cryo-TEM image of the SEI formed on lithium metal using the 0.5 M-\nDBN electrolyte; the colored rectangles indicate the presence of LiF (yellow), Li N\n3\n(blue), and LiBO (maroon). The marked regions in the cryo-TEM image (a) are\n2\nenlarged in (b–d), showing the corresponding nanocrystals of LiF (b), Li N (c),\ninterface in Li||NCM811 cells using the 0.5 M-DBN electrolyte after 1 cycle (a, b),\n20\ncycles (c, d), and 100 cycles (e, f).\nMinor comment: Please change the frame color in Fig. S29b to yellow.\nResponse: We are grateful to the reviewer for the insightful comments. In accordance\nwith the suggestion, the frame color in Figure S29b has been corrected.\nFigure S29. (a) Cryo-TEM image of the SEI formed on lithium metal using the 0.5 M-\nDBN electrolyte; the colored rectangles indicate the presence of LiF (yellow), Li N\n3\n(blue), and LiBO (maroon). The marked regions in the cryo-TEM image (a) are\n2\nenlarged in (b–d), showing the corresponding nanocrystals of LiF (b), Li N (c),\n2"}, {"round": 3, "reviewer_comments": {"reviewer_1": "I thank the authors for their comprehensive revisions and additional analyses, which have considerably strengthened the\nmanuscript. In addition, I would ask the authors to address the following points.\n1. The finding that OPLS-AA can describe the interaction between Li metal and Li ions with accuracy comparable to DFT is\nindeed surprising. Could the authors clarify whether the Lennard–Jones parameters for Li metal were specifically adjusted in\nthis work, or whether they were taken directly from previous literature? For the sake of reproducibility and transparency, it is\nessential that the full set of Lennard–Jones parameters used in the simulations be included in the Supplementary\nInformation. Furthermore, the computational details of the DFT calculations performed for validation are not currently\nprovided in the Supplementary Information. These details must be explicitly included to allow the results to be properly\nassessed and reproduced.\n2. I remain unconvinced by the manuscript's conclusion that differences in anion adsorption strength within 1 nm of the Li\nmetal surface directly determine the SEI structure extending over tens of nanometers. Even if the adsorbed anions\ndecompose during the initial stages of SEI formation, subsequent processes such as ion transport are also involved in SEI\ngrowth. Therefore, it does not seem reasonable to directly link the initial local adsorption state to the final SEI architecture.\nA clear rationale is required to bridge this scale gap and substantiate the proposed mechanism.\nMinor comment: Please change the frame color in Fig. S29b to yellow.", "reviewer_3": "This reviewer has no other questions. The revised manuscript can be published."}, "author_response": "Thank you for your email informing us of your editorial decision on our manuscript\n(NCOMMS-25-16384A). According to the instructions in your email, we have\ncarefully considered the helpful criticism from the referees and revised the manuscript\naccordingly. We provide a detailed response to the referees' criticism below. The\nchanges made during revision are highlighted in yellow and the file is uploaded\nseparately for review only.\nReviewer #1 (Remarks to the Author):\nWhile the authors have addressed some initial concerns, critical issues remain\nunresolved, particularly regarding the mechanistic interpretation and its theoretical\nfoundation. Without substantial revisions, the manuscript is not suitable for publication.\nResponse: Many thanks for the referee’s valuable suggestion, especially, the insightful\ncomments for further improving the quality of the manuscript. We have carefully\naddressed the issues as follows.\n1. The use of the OPLS-AA force field for lithium metal, combined with a strong\nelectric field of -1.0 V/nm, remains problematic. Although the authors cite precedent,\nthey do not demonstrate that the observed ion migration trends persist under more\nrealistic conditions. As the force field does not capture specific adsorption at the Li\nsurface, the model likely exaggerates interfacial differences under the applied field.\nMore seriously, despite the strong electric field, the anion distributions are noisy, and\nthe field-induced effects cannot be clearly distinguished, raising serious concerns about\nthe validity of the proposed mechanism. Furthermore, the SEI depth profiles, which\nthe authors emphasize as key experimental evidence, reflect cumulative decomposition and\ntransport processes over extended cycling and therefore do not directly validate the\nsimulated initial anion distributions.\nResponse: We sincerely thank the reviewer for the thorough evaluation of our\nmanuscript. Regarding the concern about employing the OPLS-AA force field to\ndescribe the interaction between lithium metal and lithium ions in lithium-metal-based\nsystems, we have given this issue further consideration and performed additional\nvalidation calculations. Molecular dynamics (MD) simulations using the OPLS-AA\nforce field in Gromacs confirmed that this force field indeed incorporates the interaction\nbetween lithium metal and lithium ions. Specifically, the extracted interaction energies\nwere 320.56 kJ/mol (≈3.32 eV) under zero external electric field and 410.39 kJ/mol (\n≈4.25 eV) under an applied external electric field. To further validate the reliability of\nthe OPLS-AA force field description, we conducted density functional theory (DFT)\ncalculations of the adsorption energies of three lithium salts—LiBF , LiODFB, and\n3\nand −1 V/nm. Notably, under 0 V/nm and −1 V/nm (corresponding to the MD\nsimulation conditions), the average adsorption energies of the three salts were 3.57 eV\nand 4.47 eV, respectively. These DFT results are in good agreement with the MD-\nderived interaction energies (3.32 eV and 4.25 eV), indicating that the OPLS-AA force\nfield can reasonably capture the interaction between lithium metal and lithium ions to\na certain extent. Moreover, the DFT results clearly reveal (Figure S14-S15) that the\nadsorption energies of LiODFB and LiNO₃ on lithium metal surfaces are significantly\nhigher than that of LiBF . Therefore, within the system, LiODFB and LiNO are more\n4 3\nlikely to be enriched at the lithium metal surface. Noted. In lines 112–113, the\nmanuscript states: “Meanwhile, the DFT calculations (Fig. S14-15) indicate that\nlithium metal exhibits a preferential adsorption toward LiNO and LiODFB.”\nnote that OPLS-AA is able to reasonably describe the interactions between the Li metal\nanode and Li⁺ ions. Specifically, in the absence of an applied electric field, the Li metal–\nLi⁺ interaction energy predicted by this force field is approximately 320.56 kJ/mol (≈\n3.32 eV), which is very close to the average adsorption energy (3.57 eV) we obtained\nfor the three lithium salts under the same conditions. Similarly, under an applied electric\nfield, the interaction energy is about 418.39 kJ/mol (≈ 4.33 eV), again in good\nagreement with the average adsorption energy (4.47 eV) from our DFT calculations.\nGiven the close correspondence between the key interaction energies obtained via\nOPLS-AA and the DFT-calculated average adsorption energies—both with and without\nan electric field—we are confident that the MD simulation results based on this force\nfield are accurate and reliable.\nFurthermore, following the reviewers' suggestions, we conducted experiments to\ndirectly verify the distribution of the initial anions. First, we employed quantitative\nnuclear magnetic resonance (NMR) spectroscopy to monitor the variation in the\nconcentrations of two lithium salts on the anode side during charging. Specifically, we\nassembled a Li–Li symmetric cell and applied a constant current of 0.2 mA.\nFigureS19a-b illustrate the schematic diagram of the experimental setup, which\nsimulates the charging process of a full cell. After approximately 10 min and 30 min of\ncurrent application, 0.3 mL of electrolyte was collected from the vicinity of the lithium\ndeposition electrode and mixed with 0.3 mL of deuterated chloroform (CDCl ). The\n3\nresulting solutions were analyzed by ¹¹B NMR spectroscopy, and the relative contents\nof the two anions were determined by integrating the corresponding peaks.As shown in\nFigures S19c-d, with increasing lithiation time, the concentration of ODFB⁻ anions\nnear the lithium-deposition side (anode side) gradually increased, whereas the\nconcentration of BF ⁻ anions decreased. This trend is consistent with the predictions\n4\nfrom our molecular dynamics (MD) simulations.\nWe then further examined the nanostructure of the solid electrolyte interphase (SEI)\nat the electrode/electrolyte interface using cryogenic transmission electron microscopy\n(Cryo-TEM) (Figure S29). The SEI was found to feature an outer layer rich in LiF and\nan inner layer containing Li N and B–O species, which provides additional evidence\n3\nsupporting our design concept.\nAs the reviewers' comment, the sputtering XPS depth profiling primarily reflect\nthe cumulative decomposition and transport processes of the electrolyte after prolonged\ncycling, and therefore cannot directly verify the initial anion distribution predicted by\nthe simulations. To address this, we supplemented our study with depth-profiling XPS\nanalyses after different numbers of cycles (1, 20, and 100) (Figure R1). The results\nshowed that, regardless of the number of cycles, the anode interface consistently\nmaintained a nanostructure with an LiF-rich outer layer and an Li N/B–O-containing\n3\ninner layer. This clearly demonstrates that the SEI structure formed in our experimental\nsamples retains excellent integrity throughout cycling.\nFigure S14. The adsorption energy calculations of three lithium salts — LiBF (a, d,\n4\ng), LiODFB (b, e, h), and LiNO (c, f, i) — under three applied electric field conditions:\n3\nFigure S15. Adsorption energies of different lithium salts under various voltage\nconditions\nFigure S19. (a, b) Experimental setup for ion migration studies using 0.5 M-DBN\nelectrolyte. (c) Quantitative NMR results showing the electrolyte composition near the\nlithium deposition electrode. (d) Relative compositional changes in the electrolyte near\nthe lithium deposition electrode.\nFigure S29. (a) Cryo-TEM image of the SEI formed on lithium metal using the 0.5 M-\nDBN electrolyte; the colored rectangles indicate the presence of LiF (yellow), Li N\n3\n(blue), and LiBO (maroon). The marked regions in the cryo-TEM image (a) are\n2\nenlarged in (b–d), showing the corresponding nanocrystals of LiF (b), Li N (c),\ninterface in Li||NCM811 cells using the 0.5 M-DBN electrolyte after 1 cycle (a, b),\n20\ncycles (c, d), and 100 cycles (e, f).\n2. The proposed link between the initial anion distribution and the long-term SEI\nlayering remains speculative. Continued LiODFB consumption does not ensure\nsustained interfacial enrichment. Over extended cycling, differences in anion transport\nproperties are likely to dominate, and without direct evidence, the proposed mechanism\nis not convincingly supported.\nResponse: We sincerely thank the reviewer for the detailed review of our manuscript.\nAs the reviewer rightly pointed out, the growth of the solid electrolyte interphase (SEI)\nis indeed a continuous process. However, the consumption of the lithium salt (in this\ncase, lithium difluoro(oxalato)borate, LiODFB) does not occur extensively or\nuniformly throughout cycling. The critical stage for SEI formation is during the initial\nfew cycles of battery operation, during which LiODFB participates in reactions to form\na uniform and stable SEI layer.\nGiven the electronically insulating nature of SEI, once a complete initial layer is\nestablished, electrons can no longer penetrate the interface to trigger further\ndecomposition of the lithium salt. As a result, in subsequent cycles, the evolution of the\nSEI is primarily driven by the volume changes associated with lithium plating/stripping,\nwhich may locally damage the SEI. In such cases, the electrolyte decomposes at these\ndamaged sites to form new SEI layers for “self-repair.”\nThe SEI formed in our study demonstrates excellent elasticity, allowing it to\neffectively accommodate volume fluctuations during cycling and thereby significantly\nreduce lithium salt consumption from repeated SEI rupture and repair (Figure R2).\nFurthermore, we conducted XPS characterization of the anode interface after different\ncycle numbers, and the results show that the SEI composition and structure remain\nlargely consistent throughout cycling. This provides strong evidence supporting our\nconclusion that once the initial SEI is stably formed, the subsequent lithium salt\nconsumption is substantially reduced (Figure R1).\nFigure R2. DMT modulus distribution of 1.5 M-N (a) and 0.5 M-DBN (b) SEI.\nFigure R1. F 1s (a, c, e) and N 1s (b, d, f) XPS spectra of the lithium metal anode\ninterface in Li||NCM811 cells using the 0.5 M-DBN electrolyte after 1 cycle (a, b),\n20\ncycles (c, d), and 100 cycles (e, f).\n3. The two additional binary systems (LiDOB + LiFSI and LiNO + LiTFSI) provide\n3\nsome insight but are insufficient to establish the generality of the proposed design\nprinciple. Attributing inner-layer formation to LiDOB or LiNO decomposition based\n3\nsolely on Li O detection is unconvincing, as Li O can also result from the\n2 2\ndecomposition of imide-based salts such as LiFSI or LiTFSI. This overlap in\ndecomposition pathways limits the interpretability of the experimental evidence.\nWithout further mechanistic clarification, the claim of a general design principle\nremains premature.\nResponse: We thank the reviewer for the careful evaluation of our manuscript. We fully\nagree with the reviewer’s opinion that, in our previous study involving the LiNO +\n2 3\nmay be inappropriate.\nTo further validate the universality of the proposed mechanism—namely, that\ndifferences in anion binding energy drive their competitive adsorption and preferential\ndecomposition near the lithium metal surface—we specifically designed two entirely\nnew electrolyte systems, distinct from those used in previous studies:\n1. 0.5 M LiBOB / 0.5 M LiPF in TEP: In this system, the binding energy of\n6\n2. 0.5 M LiNO / 0.5 M LiPF in TEP: In this system, the binding energy of LiNO\n3 6 3\nis higher than that of LiPF .\ninterface to investigate the SEI structures formed in each system:\n For the LiBOB/LiPF 6 system: According to our theory, LiBOB, having a\nhigher binding energy, preferentially enriches near the lithium metal surface and\ndecomposes first, forming a B–O-rich inner SEI. In contrast, LiPF , with a lower\n6\nbinding energy, tends to form an outer SEI enriched in LiF. This layered SEI\nstructure was confirmed by depth-profiling XPS analysis, which revealed a B–\nO-rich inner layer and a LiF-rich outer layer (Figure R3a-b).\n For the LiNO 3 /LiPF 6 system: Similarly, based on our proposed mechanism,\nLiNO (higher binding energy) decomposes preferentially near the lithium\n3\nsurface, resulting in a significantly higher Li N content in the inner SEI\n3\ncompared to the outer layer. Meanwhile, LiPF contributes mainly to the\n6\nformation of a LiF-rich outer SEI. The depth-profiling XPS results of this\nsystem were also consistent with the predicted spatial distribution (Figure R3c-\nd).\nOverall, these experimental results provide strong evidence that the anion gradient\ndistribution mechanism we have uncovered is not a phenomenon specific to a particular\nelectrolyte system, but rather a universally applicable interfacial regulation principle.\nFigure R3. (a) F1s and (b) B1s XPS spectra of the SEI formed in the 0.5 M LiBOB +\n0.5 M LiPF TEP electrolyte. (c) F1s and (d) N1s XPS spectra of the SEI formed in the\n6\n0.5 M LiNO + 0.5 M LiPF TEP electrolyte.\n3 6\nReviewer #3 (Remarks to the Author):\nThe authors provide detailed response to the reviewer’s comment. The “voltage-driven\nanion migration” strategy was further discussed in details in the revised manuscript,\nwhich may promote the development of phosphate electrolyte. Some remaining points\nneed to be addressed.\nResponse: Many thanks for the referee’s valuable suggestion, especially, the insightful\ncomments for further improving the quality of the manuscript. We have carefully\naddressed the issues as follows.\n1. The experimental evidence of dual-layer SEI is not convincing. XPS was used in the\nmanuscript to verify the dual-layer feature of SEI, which is a key finding of this work.\nHowever, XPS results can be misleading sometime, where the sputtering is not\nhomogeneous among different components, undermining the claim of dual-layer SEI\nformation. To clearly demonstrate the dual-layer feature of SEI, cryogenic electron\ncharacterization is suggested.\nResponse: We appreciate the reviewer’s careful evaluation of our manuscript. As\nreviewer rightly pointed out, XPS analysis may suffer from reliability issues due to\npossible artifacts introduced during sputtering. To address this concern, we followed\nthe reviewer’s suggestion and employed cryogenic transmission electron microscopy\n(Cryo-TEM) for further characterization.\nAs shown in Figure S29a, the SEI layer formed using the 0.5 M-DBN electrolyte\nexhibits a dense, inorganic-rich morphology on the lithium metal surface. High-\nresolution analysis identified several crystalline components within the SEI, including\nthe LiF (111) plane with a lattice spacing of 0.233 nm, the Li N (220) plane (0.195 nm),\n3\nand the LiBO (022) plane (0.252 nm), as displayed in Figures S29b–d. These\n2\nobservations reveal a distinct layered structure: the inner SEI region is primarily\ncomposed of Li₃N and LiBO , while the outer layer is dominated by LiF. This spatial\n2\ndistribution is fully consistent with our earlier XPS depth profiling and TOF-SIMS\nresults. Noted. In lines 204–205, the manuscript states: “Furthermore, the presence of\nthis bilayer SEI structure was corroborated by cryo-electron microscopy observations\n(Figures S29).”\nSuch a stratified SEI structure offers synergistic advantages: the outer LiF layer\nprovides high interfacial rigidity that effectively suppresses lithium dendrite growth,\nwhereas inner compounds like Li N exhibit high Li⁺ conductivity, enhancing ionic\n3\ntransport through the SEI. Additionally, the presence of B–O bonded components\ncontributes to mechanical flexibility. Together, these features enable the formation of a\nrobust SEI with both high ionic conductivity and mechanical adaptability, facilitating\ndense lithium deposition and significantly improving plating/stripping efficiency.\nFigure S29. (a) Cryo-TEM image of the SEI formed on lithium metal using the 0.5 M-\nDBN electrolyte; the colored rectangles indicate the presence of LiF (yellow), Li N\n3\n(blue), and LiBO (maroon). The marked regions in the cryo-TEM image (a) are\n2\nenlarged in (b–d), showing the corresponding nanocrystals of LiF (b), Li N (c),\n2\n2. The applicability of the so-called “voltage-driven anion migration” strategy in Li\nmetal full cell should be fully charified. This manuscript reports a voltage-based\nstrategy to regulate SEI. However, it is known that Li chemically reacts with electrolyte\nto form an SEI before cycling, i.e., before applying the electric field. Such a chemically\nformed SEI may substantially affect the “voltage-driven anion migration” strategy to\nregulate the SEI structure. The compasion between reported electrolyte and traditional\ncarbonate electrolyte is not very convincing.\nResponse: We sincerely thank the reviewer for the thoughtful and detailed comments.\nAs correctly noted, prior to battery cycling, lithium metal spontaneously reacts with the\nelectrolyte to form a chemically driven solid electrolyte interphase (SEI) that is\nextremely thin, structurally loose, and disordered (ACS Energy Letters, 2024, 9(11),\n5268–5278). Although this chemically formed SEI can offer an initial barrier against\nfurther electrolyte decomposition, its high porosity and limited blocking ability result\nin poor stability, rendering it vulnerable to subsequent electrochemical reconstruction.\nWith the application of external potential during charge/discharge processes, further\nelectrochemical reduction of the electrolyte occurs, leading to the formation of a second\nSEI layer on top of the initial chemical SEI. This electrochemically dominated SEI is\ntypically denser and more functionalized. The electrochemical formation stage largely\ndetermines the final thickness, compactness, and composition of the SEI, which in turn\ncritically influences long-term cycling performance (Batteries & Supercaps, 2021, 4,\n909; Small, 2020, 16(23), e2000756; Energy Environ. Sci., DOI:\n10.1039/D5EE01030F).\nAs the reviewers' suggestion, we prepared a control electrolyte system consisting\nof 0.75 M LiODFB/EC-DMC and 0.75 M LiBF /EC-DMC (it should be noted that\n3\nsolvent). The comparative experiments were conducted under two different conditions:\n4.3 V at 1C and 4.5 V at 0.5C. As shown in Figure R4, under the 4.3 V, 1C condition,\nthe 0.5 M-DBN electrolyte exhibited excellent cycling stability, retaining 87.90% of its\ninitial capacity after 900 cycles. In contrast, the capacity retention of the 0.75 M\nLiODFB + 0.75 M LiBF /EC-DMC electrolyte was only 36.97%. Under the higher\n4\nvoltage condition of 4.5 V, the 0.5 M-DBN electrolyte maintained 90.91% capacity\nretention after 600 cycles.\nFurthermore, to investigate the influence of different electrolyte systems,\ncomparative experiments were performed using lithium sulfonylimide-based\nelectrolytes, including 1.5 M LiFSI/TEP, 1.5 M LiTFSI/TEP, 1.5 M LiFSI/EC-DMC,\nand 1.5 M LiTFSI/EC-DMC.\nThe results showed that all four systems exhibited insufficient cycling\nperformance. Specifically, 1.5 M LiFSI/EC-DMC displayed a voltage plateau\nstagnation during the initial charging stage, likely due to severe oxidative\ndecomposition (Figure R5a). Although 1.5 M LiTFSI/EC-DMC completed charge–\ndischarge cycles, its capacity rapidly decayed within 100 cycles (Figure R5b). Both\n1.5 M LiFSI/TEP and 1.5 M LiTFSI/TEP showed low initial capacities and rapid\ncapacity fading (Figures R5c-d).\nThese comparative results clearly demonstrate the significant advantages of the\n0.5 M-DBN electrolyte in cycling stability and high-voltage tolerance, strongly\nvalidating our electrolyte design strategy.\nFigure R4. Comparison of the cycling performance of the two electrolytes under (a)\n4.3 V at 1C and (b) 4.5 V at 0.5C conditions.\nFigure R5. (a) Time–voltage profile of the first cycle for the 1.5 M LiFSI/EC-DMC\nelectrolyte. (b) Cycling performance of the 1.5 M LiTFSI/EC-DMC electrolyte under\n4.3 V 1C and 4.5 V 0.5C conditions. (c) Cycling performance of the 1.5 M LiFSI/TEP\nelectrolyte under 4.3 V 1C and 4.5 V 0.5C conditions. (d) Cycling performance of the\n1.5 M LiTFSI/TEP electrolyte under 4.3 V 1C and 4.5 V 0.5C conditions. (e) Linear\nsweep voltammetry (LSV) curves of the four electrolytes.\n3. The FTIR results need more detailed discussion. Please provide references for each\npeak’s attribution in the profiles. In particular, the peak position of LiF needs to be\nchecked. Meanwhile, the curve change during charge and discharge process is not very\nclear, since most of curves overlap in Fig. 2a.\nResponse: We sincerely thank the reviewer for the careful evaluation of our manuscript.\nIn light of your professional guidance, we have we have conducted a more detailed\ndiscussion of the FTIR results in the revised manuscript “As shown in Figures 2a and\nS20, during the lithium deposition stage (i.e., the desolvation process), the pronounced\ndepletion peaks at 977.90 cm−1 (P–O–C) and 1255.31 cm−1 (P=O) correspond to the\ngradual consumption of solvated TEP, while the emerging peak at 958.62 cm−1 (P–O–\nC) indicates the release of desolvated solvent molecules at the electrode interface14.\nFigure S19 presents the in-situ infrared spectra of the 1 M LiPF /EC–DMC electrolyte\n6\nsystem along with the corresponding spectral analysis.” in Lines 148-154. Furthermore,\nthe assignments of each peak in the in situ FTIR spectra and the corresponding\nreferences are as follows: the characteristic peak of LiF appears at 837 cm⁻¹ (Reference:\nNat. Commun. 13, 1398 (2022)); the peaks corresponding to P–O–C, B–O–R, and P=O\nare located at approximately 980 cm⁻¹, 1020 cm⁻¹, and 1260 cm⁻¹, respectively\n(Reference: Energy Storage Materials 42, 628–635 (2021)).\nRegarding the reviewer’s concern about the insufficiently clear spectral changes\nin the in situ FTIR curves, we fully understand this issue. We previously attempted to\npresent the data as contour maps to better illustrate the trends; however, the outcome\nwas unsatisfactory. This was mainly due to baseline inconsistencies in the current in\nsitu FTIR dataset (as indicated by the dashed box region in Figure R6a, where\nsubstantial baseline variations exist during the initial charge/discharge stages). If we\nforcibly normalize the baseline to generate contour plots, subtle signals—such as the\nLiF peak at ~833 cm⁻¹—would be obscured. Additionally, the baseline mismatch leads\nto difficulty in visualizing downward-oriented (i.e., product) peak evolutions in the\ncontour format (Figure R6b). It is worth noting that the data presented in Figure 7b\nhave already undergone a minimal baseline correction. More aggressive baseline\nnormalization could distort peak intensities and compromise data reliability.\nNevertheless, we have also provided individual spectra for the lithiation and delithiation\nstages below the full charge/discharge curves (Figure R7). These plots allow for more\ndistinct observation of the evolution of key characteristic peaks.\nFigure R6. (a) Stacked in situ infrared spectra of the 0.5 M-DBN electrolyte. (b)\nContour plot of the in situ infrared spectra for the 0.5 M-DBN electrolyte.\nFigure R7. Stacked plots of the in-situ infrared spectra during the plating (a) and\nstripping (b) stages.\n4. The formation mechanism of dual-layer SEI remains not very clear. It was found that\nDFOB- can be decomposed to form B-O species in the inner SEI. In principle, DFOB-\ncan decompose to form LiF in the meantime, but the inner SEI lacks LiF. Why can’t we\ndetect large amount of LiF in the inner layer?\nResponse: Thank you to the reviewer for the careful evaluation of our manuscript. We\nfully understand the concern that LiODFB decomposition might theoretically generate\nLiF. However, based on bond dissociation energy analysis, the B–O bond in the\ndifluoro(oxalato)borate anion (DFOB–) of LiODFB has a much lower bond energy\n(210.7 kJ mol–1), and in its radical form (DFOB•), the B–O bond energy can be as low\nas –50.1 kJ mol–1 (references: Nat. Energy 3, 739–746 (2018); ACS Appl. Mater.\nInterfaces 2018, 10 (3), 2469–2479). This suggests that the B–O bond is highly\nsusceptible to cleavage, especially in the radical state, leading to the formation of B–O-rich decomposition products in the SEI.\nIn contrast, the B–F bond possesses significantly higher bond dissociation\nenergy—575.3 kJ mol–1 in the anion and 236.4 kJ mol–1 in the radical—which is far\ngreater than that of the B–O bond. Therefore, the decomposition of LiODFB\npreferentially involves B–O bond cleavage and does not produce LiF in large quantities.\nTo validate this mechanism, we formulated a 1.5 M LiODFB in TEP electrolyte\nand performed sputtering XPS depth profiling of the Li metal interface. The results\nrevealed that B–O species are enriched in the inner SEI layer, while LiF is\npredominantly located in the outer SEI region. This spatial distribution pattern is fully\nconsistent with the decomposition pathway we proposed based on bond energy\ndifferences (Figure R8).\nFigure R8. (a) F1s and (b) B1s XPS spectra of the SEI formed in the 1.5M LiODFB\nTEP electrolyte.\n5. In Fig. 1f, it can be seen that the solvent TEP dominates on the Li surface. Since all\nof electrolyte components, including TEP, BF4-, DFOB- and NO3-, can decompose on\nLi surface, more TEP corresponds to more severe TEP decomposition. How can those\nminor species (anions) on Li surface can protect Li metal from reactive phosphate?\nResponse: We thank the reviewer for their careful evaluation of our manuscript. As\nshown in Table S1, TEP is indeed the dominant component in the electrolyte, and as\nthe reviewer rightly pointed out, it constitutes the major species at the interface.\nHowever, since the LUMO energy levels of the anions are lower than that of TEP,\nthe anions preferentially undergo decomposition at the interface, leading to the\nformation of an inorganic-rich interfacial layer. This layer effectively prevents direct\ncontact between TEP and the lithium metal, thereby reducing the reactivity of TEP.\nTable S1. Detailed components of the four MD models\nComponents TEP LiODFB LiBF LiNO\n4 3\n0.5 M LiODFB+0.5 M LiBF\n4\n5.9 0.5 0.5 0.5\n+0.5 M LiNO in TEP (mmol)\nReviewer #1 (Remarks to the Author):\nI thank the authors for their comprehensive revisions and additional analyses, which\nhave considerably strengthened the manuscript. In addition, I would ask the authors to\naddress the following points.\nResponse: Many thanks for the referee’s valuable suggestion, especially, the insightful\ncomments for further improving the quality of the manuscript. We have carefully\naddressed the issues as follows.\n1. The finding that OPLS-AA can describe the interaction between Li metal and Li ions\nwith accuracy comparable to DFT is indeed surprising. Could the authors clarify\nwhether the Lennard–Jones parameters for Li metal were specifically adjusted in this\nwork, or whether they were taken directly from previous literature? For the sake of\nreproducibility and transparency, it is essential that the full set of Lennard–Jones\nparameters used in the simulations be included in the Supplementary Information.\nFurthermore, the computational details of the DFT calculations performed for\nvalidation are not currently provided in the Supplementary Information. These details\nmust be explicitly included to allow the results to be properly assessed and reproduced.\nResponse: We thank the reviewer for pointing out the need for clarification regarding\nthe force field parameters and DFT calculation details. To ensure reproducibility and\ntransparency, we have included in the revised Supplementary Information (SI) a\ncomplete list and explanation of all Lennard–Jones (LJ) parameters and charge\nassignments used in this work, along with the full computational settings and key input\nfiles employed in the validating DFT calculations. Noted SI. In lines 122–136, the\nmanuscript states: “The density functional theory (DFT) calculation implemented in the\nVienna Ab initio simulation package (VASP) was conducted to study adsorption energy.\nThe exchange-correlation potential was described by using the generalized gradient\napproximation of Perdew-Burke-Ernzerhof (GGA-PBE). The projector augmented-\nwave (PAW) method was employed to treat interactions between ion cores and valence\nelectrons. We simulated the voltage effect by applying an external electric field of 0,\n−0.1 and −1 V/nm along the vacuum direction. The plane-wave cutoff energy was fixed\nto 450 eV. Given structural models were relaxed until the Hellmann–Feynman forces\nsmaller than −0.02 eV Å−1 and the change in energy smaller than 10−5 eV was attained.\nGrimme’s DFT-D3 methodology was used to describe the dispersion interactions\namong all the atoms in adsorption models. For searching the transition states, we\nemployed the nudged elastic band (NEB) method as implemented in VASP. The\nadsorption energy of these systems was calculated via following formula:\nE(binding) = E(A) + E(B) - E(AB).\nIn which, E(binding) is the adsorption energy, E(A), E(B), E(AB) are the single-\npoint energy of the various substrates.”\nTable S6. Lennard-Jones parameters[S18-S19]\natom σ (nm) ε (KJ/mol) Charge (e)\nLi+\nLi+ 0.21645 0.0764793 0.80\nBF −\nF 0.31181 0.25104 −0.3655\nF 0.31181 0.25104 −0.3655\nF 0.31181 0.25104 −0.3655\nF 0.31181 0.25104 −0.3655\nNO −\nO 0.286 0.87864 −0.4784\nO 0.286 0.87864 −0.4784\nO 0.286 0.87864 −0.4784\nDFOB−\nC 0.39967 0.359824 0.6604\nC 0.39967 0.359824 0.6604\nO 0.296 0.87864 −0.5376\nO 0.296 0.87864 −0.5376\nOS 0.300 0.71128 −0.5950\nOS 0.300 0.71128 −0.5950\nB 0.364 0.75312 1.0392\nF 0.312 0.255224 −0.4474\nF 0.312 0.255224 −0.4474\nTEP\nP 0.374 0.8368 2.1376\nO 0.296 0.87864 −0.9188\nO1 0.290 0.58576 −0.6499\nO1 0.290 0.58576 −0.6803\nO1 0.290 0.58576 −0.7062\nC 0.350 0.276144 −0.0101\nC 0.350 0.276144 −0.2477\nC 0.350 0.276144 −0.0357\nC 0.350 0.276144 −0.2471\nC 0.350 0.276144 −0.0016\nC 0.350 0.276144 −0.3275\nH 0.250 0.12552 0.1151\nH 0.250 0.12552 0.1151\nH 0.250 0.12552 0.1046\nH 0.250 0.12552 0.1046\nH 0.250 0.12552 0.1046\nH 0.250 0.12552 0.1043\nH 0.250 0.12552 0.1043\nH 0.250 0.12552 0.1069\nH 0.250 0.12552 0.1069\nH 0.250 0.12552 0.1069\nH 0.250 0.12552 0.1349\nH 0.250 0.12552 0.1349\nH 0.250 0.12552 0.1147\nH 0.250 0.12552 0.1147\nH 0.250 0.12552 0.1147\nLi_subs\nLi 0.280 9.00 0\n2. I remain unconvinced by the manuscript's conclusion that differences in anion\nadsorption strength within 1 nm of the Li metal surface directly determine the SEI\nstructure extending over tens of nanometers. Even if the adsorbed anions decompose\nduring the initial stages of SEI formation, subsequent processes such as ion transport\nare also involved in SEI growth. Therefore, it does not seem reasonable to directly link\nthe initial local adsorption state to the final SEI architecture. A clear rationale is required\nto bridge this scale gap and substantiate the proposed mechanism.\nResponse: We sincerely thank the reviewer for the careful evaluation of our manuscript\nand for the insightful and constructive comments. As you have correctly pointed out,\nthere exists a complex physicochemical process bridging the initial atomic-scale\nadsorption events and the subsequent mesoscopic formation of the SEI, and that directly\ncorrelating the two indeed requires a more rigorous justification. In our work, we\nemphasized that the initial anion adsorption and decomposition behavior establishes the\ncritical starting conditions for heterogeneous SEI nucleation and growth. We agree with\nyour viewpoint that a complete mechanism must incorporate the subsequent processes.\nNevertheless, different anion adsorption strengths result in initial SEIs with distinctly\ndifferent chemical compositions and properties. Specifically, due to the stronger\nadsorption affinity and lower LUMO levels of LiODFB and LiNO , their\n3\ndecomposition preferentially forms an inorganic-rich and compact inner layer, whereas\nthe weaker adsorption of LiBF instead favors the formation of an outer LiF layer. In\n4\nthe subsequent cycling process, the distribution state of anions is preserved, such that\nthe newly formed SEI in the fractured regions retains the original compositional\ncharacteristics: a B–O- and Li₃N-rich inner layer, and a LiF-dominated outer layer. This\nmechanism effectively ensures the overall uniformity and structural stability of the SEI.\nAt the same time, we acknowledge that our theoretical simulations were conducted\nin relatively small simulation cells. If extended to the scale of a full practical cell, such\nconcentration gradients would likely become more pronounced, thereby further\nenhancing the stratified SEI structure. Furthermore, our experimental\ncharacterizations—including sputtering XPS (Figure 2c-e), TOF-SIMS (Figure 2h-i),\nand Cryo-TEM (Figure S29)—have confirmed the successful construction of the\nenvisioned triple-layer SEI and demonstrated that it remains stable throughout extended\ncycling (Figure R1).\nFigure 2c-e. XPS depth profiles were determined of (c) F1s, (d) N1s, and (e) B1s of\nlithium metal anodes recovered from full cell cycled with 0.5 M-DBN electrolyte at\nsputtering times of 0, 50, and 100 s, respectively.\nFigure 2h-i. Time-of-flight Secondary Ion Mass Spectrometry (TOF-SIMS) depth\ndistributions (h) and three-dimensional images of various species (i) within a 0.5 M-\nDBN electrolyte.\nFigure S29. (a) Cryo-TEM image of the SEI formed on lithium metal using the 0.5 M-\nDBN electrolyte; the colored rectangles indicate the presence of LiF (yellow), Li N\n3\n(blue), and LiBO (maroon). The marked regions in the cryo-TEM image (a) are\n2\nenlarged in (b–d), showing the corresponding nanocrystals of LiF (b), Li N (c),\ninterface in Li||NCM811 cells using the 0.5 M-DBN electrolyte after 1 cycle (a, b),\n20\ncycles (c, d), and 100 cycles (e, f).\nMinor comment: Please change the frame color in Fig. S29b to yellow.\nResponse: We are grateful to the reviewer for the insightful comments. In accordance\nwith the suggestion, the frame color in Figure S29b has been corrected.\nFigure S29. (a) Cryo-TEM image of the SEI formed on lithium metal using the 0.5 M-\nDBN electrolyte; the colored rectangles indicate the presence of LiF (yellow), Li N\n3\n(blue), and LiBO (maroon). The marked regions in the cryo-TEM image (a) are\n2\nenlarged in (b–d), showing the corresponding nanocrystals of LiF (b), Li N (c),\n2"}]

Decision letter