De-solvation of heteroalkali cations enabling stable solid electrolyte interphase for dendrite-free lithium metal batteries

Domain: chemistry
Published: 2025-12-17
Editorial outcome: major_revision

Abstract
Li metal is highly sought after as a negative electrode material due to its high specific capacity and low electrode potential. However, the brittle solid electrolyte interphase and undesirable coulombic efficiency have severely hindered its commercial application. Here we propose a formation mechanism of stable solid electrolyte interphase based on the de-solvation of heteroalkali cations rather than the simple electrostatic shielding effect. Cesium trifluoroacetate (CsTFA), characterized by high electron-donor anion and weak solvent-binding heteroalkali cation, was incorporated into commercial electrolytes to foster an inorganic-rich solid electrolyte interphase. The de-solvation of absorbed Cs+ dominates the initial solid electrolyte interphase formation on the Li surface. Simultaneously, the preferential reduction of TFA- promotes the enrichment of LiF within the solid electrolyte interphase. Owing to the synergistic effect of CsTFA, Li||Cu half cells deliver a high CE of 99.77%, and Li||LFP full cells exhibit satisfactory stability over 300 cycles with a 94.3% capacity retention at a negative/positive electrode capacity ratio of 1.36. Moreover, the added CsTFA in conventional ester electrolyte demonstrates improved stability of Li||NCM811 full cells with an 80% capacity retention over 222 cycles at a negative/positive electrode capacity ratio of 1.

Introduction
Lithium metal, with high theoretical specific capacity (3860 mAh g−1) and low negative potential (−3.04 V vs. SHE), is an ideal negative electrode material for the ever-growing electrochemical energy storage systems1–3. However, the commercialization of Li metal batteries (LMBs) is impeded by safety and efficiency issues stemming from the highly reactive nature of Li4–6. The initially formed solid electrolyte interphase (SEI) is fragile and prone to fracture when undergoing displacements during cycling. This fresh fracture increases the contact surface area of the electrolyte and electrode, leading to the reactive consumption of active Li, which in turn promotes the growth of dendritic Li, ultimately resulting in low coulomb efficiency (CE) during Li plating and stripping7. Particularly and recently, many attractive strategies around the basic components of LMBs have been proposed to address the tricky Li negative electrode, including artificial SEI layer8,9, modified separator10,11, electrolyte regulation7,12 and Li-host construction13,14. While all of these are effective in improving electrochemical performance, there is still the unavoidable sacrifice of excess residual Li during full cell operation to compensate for the depletion of active Li13–15, leading to a high negative/positive electrode capacity (N/P) ratio, increased cost, and reduced energy density7,16. In addition, the sharp growth of dendrites and the inactive Li deterioration pose potential safety hazards17,18. Therefore, developing strategies based on the current industry for achieving stable cycling of lean-Li full cells remains a significant challenge.

As the flowing blood inside LMBs, electrolyte has a great impact on CE19. Newly reported advanced electrolytes all have high CE values but typically rely on high salt concentrations20, expensive diluents21, and hybridized solvent molecules or salts22. Utilizing additive salts in commercial electrolytes at a normal concentration to achieve high CE serves as a more cost-effective approach23. The reduction-sensitive anions in additive salts were routinely derived into inorganic species such as LiF, Li3P and Li2O, enhancing the physicochemical properties and interfacial compatibility of SEI with Li, facilitating dendrite-free Li deposition24. However, the role of the cationic portion in additives has been less explored in depth. Since Zhang et al.25 reported the electrostatic shielding effect of added cations in electrolyte, several later research studies all have followed that26–30. Particularly, the mechanism underlying this electrostatic shielding was not well-substantiated and remains unclear. Despite ongoing research, the specific influence of cationic additives in the electrolyte remains elusive. Recent studies on cation effects have indicated that smaller cations, possessing more concentrated charges, can form larger solvation shells through the encapsulation by multiple layers of solvent molecules. Thus, as the radii of alkali metal cations increase, their binding energy to solvents and the size of their solvation shells decreases31,32. Consequently, among the feasible alkali cations (Li+, Na+, K+ and Cs+), those with larger radii display more pronounced solvent-phobic characteristics and allow them to approach the electrode surface more closely33, which brings us a bit of inspiration for the exploration of hetero-cation additives.

Herein, a bifunctional additive with high electron-donor anion and weak solvent-binding hetero cation was introduced into conventional electrolytes to achieve stable cycling and dendrite-free LMBs at ultra-low N/P ratios. CsTFA emerges as a preferred choice due to the weak binding energy of solvated Cs+, which results in a reduced coordination number with solvent and a compact solvation shell. This configuration facilitates the incorporation of anions into the internal solvation shell and enhances the absorption of solvated Cs+ onto the electrode surface. Unlike the conventional de-solvation process of Li+, which often leads to an unstable SEI. By employing electrochemical in-situ infrared spectroscopy (FT-IR) characterization, we propose here that the de-solvation of added Cs+ predominantly influences the formation of a stable SEI during the initial reductive process. Concurrently, the re-solvation of Cs+ competes with the nearby solvated Li+ by dragging solvent and anions from it, leading to an increased release of de-solvated Li+. The extracted anions then engage in the solvated Cs+, thereby participating in the enhanced formation of SEI. Subsequent Li deposition was based on the electrochemical reduction of de-solvated Li+ facilitated by the de-solvating Cs+ network (Fig. 1). Besides, the reduction-sensitive TFA- with a low lowest unoccupied molecular orbital (LUMO) energy level was preferentially decomposed to promote the enrichment of LiF in SEI34–36. In addition, the salt containing TFA- anion with low lattice enthalpy exhibits easy solubility in aprotic solvents37,38. Compositional and physicochemical analysis of the initially formed SEI indicate that the active Li tends to deposit towards larger bulk particles beneath a robust and LiF-rich SEI. Consequently, high Li||Cu coulombic efficiency, CE of 99.17% and CE of 99.77%, can be obtained in the optimized electrolyte. Notably, the second-cycle CE shoots straight up to 99.36% with a fast-activated first cycle (initial CE of 93.15%). For practical application, Li||LFP full cell with a N/P of 1.36 demonstrates stable cycling over 300 cycles, with 94.3% capacity retention and 99.79% average CE. Under lean-Li conditions, the optimized electrolyte enables well operation of Li||LFP full cells with a range of N/P ratios from 0.107 to 0.4 (158 cycles with 90% retention and 99.9% average CE for 0.415 N/P, and 44 cycles for 0.107 N/P). Interestingly, the relationship between cycling number and N/P ratio could be linearly fitted for prediction. Moreover, in commercial ester system, improved cycling stability of 4.3V Li||NCM811 full cells can also be achieved (80% retention for 1 N/P after 222 cycles).

Results
Bulk electrolyte solvation chemistry
Both cation and anion of alkali-metal-salt additives have a significant impact on the electrochemical performance of electrolyte. Notably, the cation’s radius was inversely proportional to its binding energy with solvent molecules; thus, a larger radius corresponds to a weaker solvating capability31,32,39–41. Among available alkali metals (Li, Na, K and Cs), the large Cs+ of 1.67 Å exhibits the lowest binding energy to solvent and the greatest de-solvating ability (Fig. 2a and Supplementary data 1). This suggests that the solvated Cs+ adsorbed on the Li negative electrode surface by electric field can attract less solvent but more anions to participate in the reductive reactions during discharge, thus dominating the original inorganic-rich SEI formation. The Gutmann donor number (DN) of an anion indicates its solvating capacity, and those with high DN values were prone to attract electron-deficient cations42,43. TFA- with a high DN value of 34.0 kcal mol−1 (Fig. 2b and Supplementary data 1) and good solubility (Supplementary Figs. 1–3 and Supplementary Table 1) emerges as a potential anion candidate. The strong coordination ability of TFA- with cations reduces the coordination number of solvent molecules involved in the inner solvation shell, thereby preferentially participating in the reduction reaction and exhibiting a higher CE (Supplementary Figs. 4, 5)38,44. Synergistically, CsTFA was introduced as an electrolyte additive into commercial electrolytes to modulate the solvation structure for achieving stable cycling of lean-Li full cells at ultra-low N/P ratios. The concentration of the CsTFA additive in the baseline electrolyte (BE) was optimized initially. According to the Li||Cu protocol (Supplementary Fig. 6), an additive concentration of 0.05 M CsTFA in BE electrolyte demonstrates the most extended cyclic stability and the highest CE. Therefore, 0.05 M has been set as the benchmark for assessing other available alkali-metal-cation trifluoroacetate additives. Commercial 1M LiTFSI in 1,2-dimethoxyethane (DME)/1,3-Dioxolane (DOL) electrolyte, devoid of the commonly used LiNO3 additive, was selected as the sieving electrolyte (SE) to amplify the impact of additives considering the significant cyclic fluctuation and low CE within 30 cycles (0.5 mA cm−2 with a fixed capacity of 1 mAh cm−2, Supplementary Fig. 7). The electrolytes with the addition of 0.05 M LiTFA, NaTFA, KTFA and CsTFA exhibit improved properties. The relatively worse performance of SE+Na and SE+K electrolytes maybe attributed to the co-deposition during Li plating, as suggested by the Nernst equation (Supplementary Table 2). Due to the interfacial de-solvation of Cs+, SE+Cs demonstrates greater cyclic performance and higher Aurbach CE compared to SE+Li (Supplementary Figs. 7, 8). The influence of different additives on the Li negative electrode follows a similar tendency in the commercial BE electrolyte supplemented with additional 0.1 M LiNO3 (Supplementary Figs. 9, 10). Consequently, the subsequent research all focuses on the optimized BE electrolyte with the incorporation of 0.05 M CsTFA.

Classical molecular dynamics (MD) simulations were conducted to clarify the effect of additive on the solvation structure of bulk BE+Cs and BE electrolytes (Fig. 2c–h and Supplementary data 3). Analysis of the radial distribution functions (RDF, denoted as g(r)), particularly for the Li+-O(TFA-) and Li+-O(NO3-) interactions within the BE+Cs electrolyte (Fig. 2g), reveals that anions TFA- and NO3- were more likely to associate with Li+ over solvent molecules in the innermost solvation shell. This preference was quantified by the calculated coordination numbers of Li+ with various species in the BE+Cs electrolyte, including 0.378 for DOL, 1.061 for DME, 0.480 for TFSI-, 0.085 for NO3-, and 0.048 for TFA- (detailed in Supplementary Table 3). Moreover, Supplementary Fig. 11 and Supplementary Table 4 show the calculated solvation structure of Cs+ in BE+Cs. The RDF (g(r)) of Cs+-O(TFA-) and Cs+-O(TFSI-) interactions demonstrate a more pronounced and closer association with Cs+ (around 2.7 Å) over solvent molecules. And the calculated coordination numbers of Cs+ with solvent, 0.228 for DOL and 0.335 for DME, were smaller than those for Li+. The peak located at −35.9 ppm in 133Cs nuclear magnetic resonance (NMR) spectra demonstrate the successful involvement of Cs in solvation structure (Supplementary Fig. 12)45,46. These results indicate that fewer solvent molecules but more anions (TFSI- and TFA-) tend to coordinate with Cs+ in the innermost solvation shell. The reduced involvement of solvent molecules during Li plating suggests the formation of a more inorganic-rich and stiffer SEI. Various characterization techniques were performed to investigate the chemical properties of the electrolyte. Ex-situ FT-IR spectra show the reactions between Li+, anions, and solvent molecules (Supplementary Figs. 13–17). With CsTFA addition, TFA- partially participates into the solvent shell of Li+ once dissociation. The spectra of BE+Cs and SE+Cs electrolytes both show characteristic peaks at 1705 and 1720 cm−1 (C=O asymmetric stretching frequency), corresponding to the strong interaction between Li+ and carbonyl of TFA-47. Due to the strong binding energy of Li+ to TFA- and NO3- (Fig. 2k and Supplementary data 1), 7Li nuclear magnetic resonance (NMR) spectra shift to low field with a higher chemical shift value (Fig. 2j). The deshielding effect in 7Li NMR and the apparent blueshift at the characteristic peak of TFSI- in Raman spectra (located at 735–755 cm−1, S-N-S stretching, Fig. 2i) were rationally expected with the addition of CsTFA. Density functional theory (DFT) calculation and frontier molecular orbital theory were implemented to estimate the solvation structure and redox reactivity of the electrolytes. As shown in Fig. 2l and Supplementary data 2, both CsTFA and LiTFA possess relatively high highest occupied molecular orbital (HOMO) energy levels, suggesting that TFA- in the innermost solvation shell was prone to oxidatively decompose on the positive electrode surface, forming a fluorine-rich positive electrode electrolyte interphase34. Furthermore, attributed to the strongly polar CF3- group, TFA- owns relatively low LUMO and high reduction potential (Supplementary Fig. 18 and Supplementary data 2, 4) compared to other solvent and anionic components, indicating that TFA- near the negative electrode surface was more sensitive to electrochemical reduction, thus facilitating the generation of high-quality SEI41, which was demonstrated by the distinct peak at 1.0–1.2 V in linear sweep voltammetry analysis (LSV, Supplementary Fig. 19)36,48.

In-situ visualization of de-solvation
To explore the specific mechanism by which CsTFA addition influences the de-solvation process and the initial SEI formation, electrochemical in-situ FT-IR spectroscopy was employed to monitor the dynamic variation of the interfacial electrolyte configuration during real-time Li deposition (Fig. 3a)49,50. The in-situ raw FT-IR spectra and the corresponding differential A(t)-A(t0) spectra for both BE+Cs and BE electrolytes were presented in Fig. 3d and Supplementary Fig. 20. As de-solvation proceeds, de-solvated solvents were released and gradually accumulate at the interphase. Synchronously, anions were repelled to maintain charge balance once Li+ begins to plate. Therefore, the absorption bands corresponding to free solvent (+, marked light green) appear as upward peaks, whereas those associated with Li+-solvent (-, marked light yellow) and anion (-, marked light gray) species exhibit downward dips (Supplementary Table 5 for peaks assignments). The intensities of the corresponding peaks change in tandem. During Li plating in the BE electrolyte (Fig. 3d, bottom), a clear shift in relative absorbances from Li+-solvent (-) to free-solvent (+) was observed, along with significant changes in TFSI (-) absorbances, indicating a pronounced solvent-related reaction and a decreased content of anions at the interphase. In contrast, the BE+Cs electrolyte (Fig. 3d, top) shows weaker and less variable solvent-related absorbances, with smaller and more moderate changes in TFSI (-) absorbances, consistent with the data from Raman and 7Li NMR spectra, suggesting that there were less solvent and more anions at the interphase in the presence of Cs+. Besides, the appearance of TFA (-) absorbance in the raw spectra and the related variable peak intensities in the differential spectra demonstrate its role in enhancing SEI enriched in LiF. Consequently, the initial plating process in BE electrolyte follows the conventional de-solvation and electroreduction of solvated Li+, leading to the formation of an anion-lean and solvent-rich SEI (Fig. 3c). However, integrating the weak solvent-binding behavior of Cs+ and the different results from in-situ FT-IR characterization, it’s proposed that the initial anion-rich and solvent-lean SEI was primarily shaped by the de-solvation of adsorbed Cs+, which involves a higher concentration of anions and a reduced solvent content (Fig. 3b). The de-solvated but undeposited Cs+ induces an increased release of de-solvated Li+ by dragging solvent from solvated Li+, and the subsequent de-solvation and electroreduction of de-solvated Li+ facilitate Li deposition.

Electrode compatibility and interfacial kinetics
Li||Cu half cells were assembled to assess the Li plating/stripping efficiency. As shown in Fig. 4a, the commercial LiNO3-containing BE electrolyte displays an inconspicuous average CE of ~98% and fails within 120 cycles. In contrast, improved cycling stability and a high CE of about 99.17% over 300 cycles could be obtained in BE+Cs with the addition of bifunctional CsTFA. This enhancement was attributed to the high-quality SEI formed by the de-solvation of weak solvent-binding Cs+ during the initial cycle. Specifically, the second-cycle CE of BE+Cs shoots straight up to 99.36%, and the average CE for the following four cycles reaches 99.32% (Fig. 4b). Furthermore, the voltage-capacity profiles (Supplementary Fig. 21), in conjunction with the anodic cyclic voltammetry (CV, Fig. 4d and Supplementary Fig. 22), consistently indicate that the plated Li in BE+Cs electrolyte contains a much higher content of active Li.

The BE+Cs electrolyte also outperforms in CE across various deposition amounts and rates (Supplementary Figs. 23–25). Using a modified Aurbach protocol method3, the modified BE+Cs electrolyte yields a high CE reaching up to 99.77% (Fig. 4b). In both CE test configurations conducted (CE Cycling and CE Aurbach), the BE+Cs electrolyte demonstrates excellent Li-utilization efficiency consistently, which stands out compared to other electrolyte systems reported in the literature, including both conventional normal concentration and high concentration/locally high concentration electrolyte systems (Fig. 4i and Supplementary Tables 6, 7). Furthermore, the compatibility of the electrolytes with Li metal electrode was evaluated through galvanostatic and rate performance measurements of Li||Li symmetric cells. As shown in Fig. 4e and Supplementary Fig. 26a, the incorporation of CsTFA significantly enhances the cycling stability and reduces the overpotential under specific cycling protocols (3 mA cm−2, 1 mAh cm−2 and 5 mA cm−2, 1 mAh cm−2). Notably, the BE+Cs electrolyte can sustain a stable cyclic calendar lifetime of 2000 h and 600 h under these conditions, respectively. Meanwhile, the Li||Li symmetric cells containing the BE+Cs electrolyte exhibited improved rate performance and diminished room-temperature electrochemical impedance both before and after the rate test, underscoring the great capabilities of the BE+Cs electrolyte (Supplementary Fig. 26b–d). To further explore the interfacial kinetics between the electrolyte and Li negative electrode, the exchange current density (j0) and activation energy (Ea) were measured using the Li||Li symmetric cell model51. Figure 4f illustrates the Tafel plots and the calculated j0, and Supplementary Fig. 27 displays the corresponding voltage-current curves. The j0 value fitted to the curve slope for BE+Cs electrolyte (1.82 mA cm−2) was much higher than that for BE electrolyte (1.65 mA cm−2), indicating that the processes of charge transfer were accelerated at the interphase with the addition of CsTFA. The Ea obtained from temperature-dependent electrochemical impedance spectroscopy (EIS) measurements by linear fitting of ln(1/R) vs. 1/T was also investigated to support this conclusion (Fig. 4g, h and Supplementary Fig. 28)52. The calculated Ea value of BE+Cs (41.66 kJ/mol) was lower than that of BE (46.93 kJ/mol), implying a lower energy barrier for Li deposition and faster electrochemical kinetics in the BE+Cs electrolyte.

Since the CE of Li plating/stripping was correlated with the electrodeposition state, scanning electron microscopy (SEM) was recorded to investigate the morphology and thickness of freshly deposited Li in different electrolytes (Supplementary Fig. 29). With a specific areal capacity of 5 mAh cm−2 under a current density of 0.5 mA cm−2, the Li||Cu cells were terminated at 0 V prior to disassembly and postmortem electrode measurements. The top-sectional view reveals that Li deposition in the SE electrolyte appears as fragmented, bean-like particles with conspicuous gaps, which promote the generation of “dead” Li and exacerbate the irreversible decomposition of the electrolyte due to the increased contact surface area. Besides, loose dendritic Li deposition can be observed in BE electrolyte with additional 0.1 M LiNO3. In contrast, the SE+Cs electrolyte, after the addition of CsTFA, results in Li growing towards larger bulk particles with reduced voids. Notably, the optimized BE+Cs electrolyte facilitates dense Li deposition in the teens of microns range and virtually void-free. This ideal deposition morphology mitigates side reactions caused by over-exposure of freshly active Li metal to electrolyte and inhibits the formation of “dead” Li during stripping, thus contributing to a desired high CE value53. In conjunction with the corresponding electrochemical deposition curves (Supplementary Fig. 30), the initial nucleation overpotential upon galvanostatic plating decrease from 82.2 mV in SE and 65.3 mV in BE to 36.2 mV in SE+Cs and 31.4 mV in BE+Cs, respectively. Cross-sectional SEM images with a fixed capacity of 1.0, 2.0, 3.0 and 5.0 mAh cm−2 were also conducted to investigate the Li densification relative to the reaction surface (Supplementary Fig. 31). It was revealed that the thickness of deposited Li decreases from 38.8 µm in SE and 29.3 µm in BE to 35.7 µm in SE+Cs and 25.1 µm in BE+Cs, respectively. These results clearly demonstrate that the addition of CsTFA leads to improved deposition morphology and reduced nucleation barriers.

Interfacial physicochemistry of SEI
Depth-profiling X-ray photoelectron spectroscopy (XPS) was further carried out on freshly deposited Li to investigate the chemical composition of different SEI across the thickness direction. From the curves of elemental atomic concentration versus sputtering time and the high-resolution XPS depth profiles of F 1s, O 1s, C 1s Li 1s and N 1s peaks (Fig. 5a–d and Supplementary Figs. 32, 33), there were a decreased content of C-F, RCOOLi, C-C/C-H and an increased amount of LiF, Li2O, LiNxOy with the addition of CsTFA (comparing BE+Cs to BE, and SE+Cs to SE)7,54. This indicates that the original SEI was enriched with fluorine and oxygen while being depleted of organic components due to the interfacial de-solvation of absorbed Cs+, which inhibits solvent decomposition and promotes effective anions decomposition, including TFA-, TFSI- and NO3-55.

Cryogenic transmission electron microscopy (Cryo-TEM) was employed to directly visualize the microstructure and chemical composition of the freshly formed SEI. Figure 5g, h exhibit the low magnification cryo-TEM images of plated Li in both BE+Cs and BE electrolytes. In the BE electrolyte, small-sized and inhomogeneous Li aggregates can be observed, with apparent ruptures depicted by yellow dashed circles in Fig. 5h. In contrast, the optimized BE+Cs electrolyte promotes large and chunky Li deposition on the Cu mesh, consistent with the SEM image results. High-resolution cryo-TEM images reveal distinct SEI with visible boundaries in both electrolytes (areas with brighter diffraction lining). The SEI thickness in BE electrolyte was approximately 25 nm, slightly thicker than the 17 nm observed in BE+Cs electrolyte. Specifically, the marginal region of SEI in BE electrolyte shows a messy, rippled appearance composed of small Li flakes (depicted by translucent blue line in Fig. 5f), whereas the BE+Cs electrolyte shows a flat SEI edge line (depicted by translucent red line in Fig. 5e). High-resolution TEM (HR-TEM) images and fast Fourier transform (FFT) patterns (Supplementary Figs. 34, 35) confirm the crystalline phases in the inner SEI regions of the BE+Cs electrolyte as LiF, Li2O, Li3N and LiNOx, with lattice fringes of 2.3, 2.66, 2.76 and 3.06 Å, respectively24,54. A homogeneous SEI with an amorphous organic outer layer and an inorganic inner layer exhibits great mechanical stability and ideal ionic conductivity, which were crucial for alleviating volume expansion during cycling. However, the formed SEI in BE electrolyte presents a disordered component distribution, with crystalline Li2CO3 (lattice fringes of 2.83 Å) found near the messy fringe area. This Li2CO3 originates from solvent decomposition and exhibits poor ionic conductivity, thus leading to uneven Li+ flux and escalating dendritic growth of Li56.

The detailed composition and structure of SEI in BE+Cs and BE electrolytes were further investigated by time-of-flight secondary ion mass spectrometer (TOF-SIMS) analysis (Fig. 5i, j and Supplementary Figs. 36–40). Detected secondary ion fragments such as LiF-, LiF2-, LiO-, LiO2-, Li3N-, LiN-, CO3-, C2HO- during negative mode sputtering, were characteristic of LiF, Li2O, Li3N, Li2CO3 and organic species, respectively. Figure 5i and Supplementary Fig. 36 display the trace compositional variability in quantity along with sputtering time. Notably, in BE+Cs electrolyte, the intensity of LiF- was much higher than that of other species throughout the entire sputtering process, indicating an enrichment of inorganic LiF in the inner SEI, as proved by depth-profiling XPS spectra. The intensity signal of C2HO- fades gradually after an etching depth of 100 s, suggesting that organic components were primarily located in the outermost SEI. Additionally, a significant Li-enriched layer was detected at an etching depth of about 200 s. By comparing these results with those of BE electrolyte (Supplementary Fig. 36c–h), we observed that the SEI in BE+Cs electrolyte contains a higher concentration of inorganic components (LiF, Li2O and Li3N) and exhibits more active Li beneath the SEI layer. In contrast, the SEI in BE electrolyte was enriched with organic components and a higher amount of Li2CO3, which was generally considered detrimental to SEI stability. The specific hierarchical structure of SEI was more clearly shown by 3D depth profiles (Fig. 4j and Supplementary Figs. 37–40). These TOF-SIMS results were consistent with those obtained from cryo-TEM and XPS, implying that CsTFA induces a higher LiF content in the SEI.

Atomic force microscopy (AFM) was utilized to assess the mechanical properties of SEI (Fig. 5k–m and Supplementary Fig. 41). The surface height mode reveals a flatter and smoother SEI surface in the BE+Cs electrolyte with a roughness (Rq) of 213 nm, contrasting with the pronounced undulations of 377 nm Rq observed in the BE electrolyte. Adequate mechanical strength of the SEI was pivotal in preventing interfacial breakage during cycling57. The average interfacial mechanical strength significantly increased from 3.36 GPa in BE to 10.75 GPa in BE+Cs due to the addition of CsTFA, signifying a robust SEI.

Stringent electrochemical evaluations of full cells
In order to verify the high compatibility of BE+Cs electrolyte with Li negative electrode, Li||LFP full-cell protocols (2.8–4.0 V) with an N/P ratio of 1.36 were assembled for testing by pairing a quantized Li foil negative electrode (20 µm in thickness, 4 mAh cm−2) with a quantized LFP positive electrode (17.3 mg cm−2, 2.94 mAh cm−2) and injecting a precise dosage of 30 µl electrolyte per cell (10.2 µl mAh−1). Figure 6a, c exhibit the cycling performance (0.2 C formation then 1 C cycle) and rate performance (0.2, 0.5, 1, 2, 3 and 1 C) of the Li||LFP full cells, respectively, and the corresponding charge/discharge curves were recorded in the Supplementary Figs. 42, 43. The capacity of full cells in BE electrolyte decays after 100 cycles and rapidly declines after 160 cycles, attributing to the depletion of active Li or electrolyte. However, the full cells with BE+Cs electrolyte maintain stable operation over 300 cycles, with a high reversible capacity of 146 mAh g−1 at 1 C, a capacity retention of 94.3%, and an average CE of 99.79% over the whole full-cell calendar lifespan (Fig. 6a, b). Moreover, the full cells with BE+Cs electrolyte exhibit great rate capacities of 158, 154, 148, 139 and 128 mAh g−1 at 0.2, 0.5, 1, 2 and 3 C, respectively (Fig. 6c, d). Even after reverting the current density back to 1 C, the cells could maintain stable cycling for an extended lifespan. The reduced internal resistance and improved interfacial kinetics by the added CsTFA were evidenced by the results of EIS tests during battery cycling process (Supplementary Fig. 44). The benefit of synergistic CsTFA additive was further demonstrated in lean-Li||LFP full cells with ultra-low N/P ratios ranging from 0.1 to 0.4. The limited Li negative electrodes were prepared by pre-depositing specific amounts of Li (0.4, 0.8 and 1.5 mAh cm−2) on Cu substrates in BE+Cs electrolyte. The corresponding plating curves, optical photographs and SEM images of the pre-deposited Li-Cu electrodes were shown in Supplementary Fig. 45. Then the harvested Li-Cu electrodes were paired with LFP positive electrodes of varying mass loadings to meet the range of ultra-low N/P ratios (Supplementary Table 8). Notably, the Li||LFP full cells with a N/P ratio of 0.415 exhibit great cycling stability, maintaining operation over 158 cycles with 90% capacity retention and an average CE above 99.9% during the initial 147 cycles (Fig. 6e). Even at much lower N/P ratios, the Li||LFP full cells also operate stably (Supplementary Figs. 46–49). In particular, at a squeezed N/P ratio of 0.107, the full cells still sustain a high CE of 99.9% over 28 cycles (Fig. 6f). The calculated Li inventory loss rate (LILR) values of the lean-Li||LFP full cells range from 0.138% to 0.395%, demonstrating the high stability and compatibility of the electrolyte with Li negative electrode58. In the anode-free protocol (no residual Li on the current collector), the Cu||LFP full cells possess a capacity retention of 40.5% after 200 cycles with a high average CE of 99.4% and an average Li inventory retention rate (LIRR) of 99.436% (Fig. 6g and Supplementary Fig. 50)59. After analyzing the cycling results of the lean-Li metal full cells (Fig. 6i, j), we establish two linear correlations between the cycle number (CN) and the N/P ratio, corresponding to the points when the capacity retention (CR) drops to 90% (CNCR≥90%) and when the coulomb efficiency (CE) falls to 99.7% (CNCE≥99.7%). The relationship were described by the following equations:

CNCR≥90% = 356 × N/P + 4.49 (1)

CNCE≥99.7% = 360.2 × N/P − 16.2 (2)

The corresponding CN increases linearly with the N/P ratio under lean-Li conditions, suggesting that the lifespan of full cells with pre-deposited Li can be predicted ahead. The electrochemical performance of electrolyte was reflected in the slope of the linear fit. Notably, the Y-axis intercept in Eq. (1) represents the best-effort cycling lifespan of the anode-free protocol, and the X-axis intercept in Eq. (2) can be referenced as the minimum amount of irreversible Li necessary for the cell operation. Our study outperforms recent work on Li metal batteries with ultra-low N/P ratios (<4), as shown in Supplementary Fig. 51 and Supplementary Table 9.

Interestingly, the application of CsTFA in ester system for high-voltage positive electrode Li metal batteries was also evaluated. As displayed in Supplementary Fig. 52, the CsTFA has a saturated solubility of 0.025 M in the commercial EC/DEC/FEC system, nevertheless, the formulated electrolyte with the addition of 0.05 M remains opaque after 24 h. The difference at characteristic peak (1705–1720 cm−1) in ex-situ FT-IR spectroscopy confirms the dissociation of CsTFA in the electrolyte (Supplementary Fig. 53). DFT calculation for the LiPF6-ester system indicates that large Cs+ also has a small binding energy to ester solvent molecules, and TFA- owns the lowest LUMO energy level and the strongest binding energy with Li+ (Supplementary Fig. 54). The improved Li||Cu and Li||Li half-cell performance and reduced EIS impedance in BCE+Cs electrolyte demonstrate its improved compatibility with Li negative electrode (Supplementary Figs. 55–57). High-voltage full-cell protocols (3.0–4.3 V, N/P of 1) were assembled by pairing Li foil negative electrode (20 µm in thickness, 4 mAh cm−2) with NCM811 positive electrode (20.7 mg cm-2, 4 mAh cm−2) and injecting a precise dosage of 30 µl electrolyte per cell (7.5 µl mAh−1). The Li||NCM811 full cells with BCE+Cs electrolyte exhibit stable cycling over 222 cycles with 80% capacity retention and good rate performance even at large mass loadings (Supplementary Fig. 58). These outcomes reinforce the viability of incorporating CsTFA into commercial ester-based electrolytes.

Methods
Materials and preparation of electrolytes
LiPF6 (99%), LiTFSI (99%), LiBF4 (99%), LiClO4 (99%), LiFSI (99%), LiOTF (99%), LiNO3 (99%), LiTFA (99%), NaTFA (99%), KTFA (99%) and CsTFA (99%) were all purchased from Adamas Reagent. The ether electrolytes (1 M LiTFSI in DOL/DME = 1:1 vol and 1 M LiTFSI in DOL/DME = 1:1 vol with 0.1 M LiNO3) were obtained from Zhangjiagang Guotai-Huarong New Chemical Materials Corporation. The ester electrolyte (1 M LiPF6 in EC/DEC = 1:1 vol and 1 M LiPF6 in EC/DEC/FEC = 42.5:42.5:5 vol) were obtained from Dodochem. All the additives and electrolytes were stored and used in the glove box (O2 <0.01 ppm, H2O <0.01 ppm) without any purification. The LiFePO4 (LFP) powder, LiNi0.8Co0.1Mn0.1 (NCM811) powder, polyvinylidene fluoride binder (PVDF, HSV900), Super-P carbon black, copper foil (9 µm in thickness) and aluminum foil (16 µm in thickness) were obtained from Shenzhen-Kejing Corporation. N-methylpyrrolidone (NMP, >99.0%) was purchased from Aladdin. Li foils (20 µm and 450 µm in thickness, Li content ≥99.9%) were purchased from China Energy Lithium Co. Celgard 2500 (porosity: 41%; average pore size: 0.043 µm; thickness: 25 µm), 2032-type coin-cell cases, springs and spacers were obtained from Canrd Technology Co. Ltd. LFP positive electrodes with a series of loadings (from 13.9 to 22.5 mg cm−2, in inextricably) were prepared by casting the slurry consisting of LFP, Super-P and PVDF with a weight ratio of 85:7.5:7.5 in NMP onto aluminium foil. NCM811 positive electrodes (20.7 mg cm−2) were similar to LFP, except that the mass ratio was changed to 92:4:4.

Positive electrodes preparation and electrochemical measurements
CR2032 coin cells were assembled in an Ar-filled glove box (O2 <0.01 ppm, H2O <0.01 ppm) to test the basic electrochemical properties of the electrolytes. All electrolyte preparations were carried out in the same glovebox. To prepare the additive-modified ether electrolyte, 0.05 M trifluoroacetate salts (LiTFA, NaTFA, KTFA or CsTFA) of 0.05 M were dissolved into SE or BE, denoted as SE+Li, SE+Na, SE+K, SE+Cs, BE+Li, BE+Na, BE+K and BE+Cs, respectively. Moreover, a range of concentrations of CsTFA were dissolved into BE, denoted as BE+Cs 0.005, BE+Cs 0.01 and BE+Cs 0.1. 0.05 M different anionic salts (LiPF6, LiTFSI, LiBF4, LiClO4, LiFSI, LiOTF, LiNO3 and LiTFA) were dissolved into SE or BE, denoted as SE+LiPF6, SE+LiTFSI, SE+LiBF4, SE+LiClO4, SE+LiFSI, SE+LiOTF, SE+LiNO3, SE+LiTFA, BE+LiPF6, BE+LiTFSI, BE+LiBF4, BE+LiClO4, BE+LiFSI, BE+LiOTF, BE+LiNO3 and BE+LiTFA. To prepare the additive-modified carbonate electrolyte, 0.05 M CsTFA was dissolved into SCS or BCE, denoted as SCE+Cs, BCE+Cs. The amount of electrolyte per cell was controlled at 30 µL, and Celgard 2500 was used as the separator.

Galvanostatic charge/discharge tests were carried out using a Neware battery cycler (CT-4008) at 30 °C in a climatic chamber. For Li||Cu coin cells, the cycle CE was measured by plating/stripping 1, 2, 3 or 5 mAh cm−2 Li repeatedly at 0.5 mA cm−2 with a charge cut-off voltage of 2 V. The modified Aurbach CE test was as follows: (1) perform one initial formation cycle with Li deposition of 5 mAh cm−2 on Cu under 0.5 mA cm−2 current density and stripping to 2 V; (2) subsequently deposit 5 mAh cm−2 Li as a total Li reservoir (QT); (3) repeatedly strip/deposit Li of 1 mAh cm−2 (QC) under 0.5 mA cm−2 for 10 cycles; (4) strip all the residual Li (QS) to 2 V under 0.5 mA cm−2 exhaustively. The CE Aurbach was calculated according to the following equation:

CE Aurbach = (10 × QC + QS) / (10 × QC + QT)

Cyclic voltammetry (CV) and linear sweep voltammetry (LSV) measurements of Li||Cu half cells were conducted by a CHI 760E electrochemical workstation with different voltage ranges respectively, the testing temperature was set at 30 °C.

Li||Li symmetrical cells were assembled with two identical Li foils (450 µm in thickness, 10 mm in diameter) and cycled at specified current density and capacity in each half cycle. Cyclic capabilities of symmetric batteries were measured at 3 and 5 mA cm−2 current density under a fixed capacity of 1 mAh cm−2. Rate capabilities of symmetric batteries were tested at 0.5, 1, 2, 3, 5 and 1 mA cm−2 of 1 mAh cm−2. Tafel plots of Li||Li half cells were obtained from a CHI 760E electrochemical workstation by linear sweep voltammetry measurement at a scanning rate of 10 mV s−1 within the voltage range of −0.12–0.12 V, the testing temperature was set at 30 °C. The exchange current density j0 was calculated from the Tafel equation:

η = A × (log j − log j0)

where η and j were the overpotential and current, respectively, and A represents the kinetic constant in the diffusion-controlled surface dynamic range that could be acquired after fitting the data. The value of j0 can be obtained from the intersection of the extrapolated linear part of the log j versus η plot with the μ = 0 line.

Electrochemical impedance spectroscopy (EIS) measurements of Li||Li half cells and Li||LFP full cells were conducted by a CHI 760E electrochemical workstation from 100 kHz to 100 mHz at the open-circuit potential with an amplitude of 10 mV after 30 min of stabilisation at open-circuit voltage. Activation energy measurements were tested by using Li||Li symmetrical cells through EIS measurements in an incubator under different temperatures from 5 to 30 °C with 5 °C separation. The values of Li deposition activation energy (Ea) were calculated by the following equation that could be extracted from the slope plot of log RSEI vs. inverse temperature (1/T).

T = RSEI = A exp(−Ea / RT)

where Ea was the activation energy, T was the absolute temperature, R was the gas constant (8.314 J·mol−1 K−1), RSEI was the interfacial Li+ transfer resistance obtained from the EIS plots, and A was the pre-exponential factor.

The Li||LFP full cells were assembled by matching LFP positive electrode (17.3 mg cm−2) with Li foil (20 µm in thickness), followed by activating at 0.2 C (current density: 1 C = 3.3 mA cm−2, 170 mAh g−1 for LFP) for 2 cycles and charging/discharging at 1.0 C in subsequent cycles. Each Li||LFP full cell was ejected with 30 µl of ether electrolyte and the voltage window was set within 2.8–4.0 V. The Li||NCM811 full cells were assembled by matching NCM811 positive electrode (20.7 mg cm-2) with Li foil (20 µm in thickness), followed by activating at 0.2 C (current density: 1 C = 4.1 mA cm−2, 200 mAh g−1 for NCM811 at 4.3 V) for 2 cycles, then charged at 0.33 C and discharged at 1.0 C in subsequent cycles. Each Li||NCM811 full cell was ejected with 30 µl of carbonate electrolyte and the voltage window was set within 3.0–4.3 V.

The rate capabilities of LFP or NCM811 full cells were performed by charging/discharging 6 cycles at progressively increasing current densities, respectively.

For the assembly of Li||LFP full cells at ultra-low N/P ratios, some quantities of 0.4, 0.8 and 1.5 mAh cm−2 Li were pre-electrodeposited onto the bare Cu foil (14 mm in diameter) through the Li||Cu cell protocol in advance, which were paired with LFP positive electrodes of different mass loadings (from 13.9 to 22.5 mg cm−2, the specific ratios were shown in Supplementary Table 6) to meet the set N/P ratios. Li||LFP full cells at ultra-low N/P ratios were activated at 0.2 C for 2 cycles, followed by charging/discharging at 0.5 C in subsequent cycles. The electrolyte dosage for each cell was controlled at 30 µl and the voltage window was set within 2.8–4.0 V.

The Li inventory loss rate (LILR) in Li||LFP full cells was calculated by the following equation58:

LILR = 100 × (1 − (Qn/Qt)1/n)

where Qn was the discharge capacity after n cycles, Qt was the discharge capacity at the transition point (cycle number = t) and LILR was an averaged percentage loss of Li inventory per cycle.

The Li inventory retention rate (LIRR) in Cu||LFP full cells was calculated by the following equation59:

LIRR = 100 × (Qn/Qint)1/n

where Qn was the discharge capacity after n cycles, Qint was the discharge capacity of the first cycle and LILR was an averaged percentage retention of Li inventory per cycle.

Materials characterization
SEM characterization is conducted on a Hitachi SU8010 scanning electron microscope to obtain the morphologies of the samples. Raman tests are carried out on NTEGRA Spectra AFM Raman Confocal SNOM instrument with laser wavelength of 633 nm. Ex-situ FT-IR spectra is collected on a Spectrum Two LiTa spectrometer FT-IR instrument (PerkinElmer, America), covering 4000–450 cm−1 at 1.0 cm−1 resolution. Atomic force microscopy (AFM) is carried out on a Bruker Dimension Fastscan microscope in peak force quantitative nanomechanics mode, and Young’s moduli are determined using the Derjaguin-Muller-Toporov model, the samples sent for AFM tests were obtained by depositing 1.5 mAh cm−2 at 0.5 mA cm−2 on Cu after cycling at 1 mA cm−2 and 1 mAh cm-2 for 5 cycles at 30 °C. In-situ FT-IR spectra is collected on a Nicolet iS50 spectrometer FT-IR instrument (Thermo Fisher Scientific, America) fitted with a single-reflection diamond attenuated total reflection (ATR) cell. A liquid-nitrogen-cooled MCT detector is employed at a 60° incidence angle with 4 cm−1 spectral resolution. A lithium rod is positioned flush against the diamond window as the working electrode, while a lithium ring and lithium wire are served as the working electrode and reference electrode, respectively. A CHI 660E electrochemistry workstation is used for potential control and current measurement. Li is plated at −0.1 V for 1200 s with tested electrolyte potentiostatically. 7Li and 133Cs nuclear magnetic resonance (NMR) spectra are obtained on a BRUKER AVANCE III HD 500 Nuclear Magnetic Resonance spectrometer at 30 °C. Cryogenic-transmission electron microscopy (Cryo-TEM): the samples were obtained by depositing 0.2 mAh cm−2 at 1 mA cm−2 on copper grids after cycling at 1 mA cm−2 and 1 mAh cm−2 for 5 cycles. The Li-deposited copper grids are placed into a four hole TEM sample cassette and sealed under argon gas. It is then immersed in liquid nitrogen and disassembled (ensuring isolation from water and oxygen) and transferred to a TEM column for testing. The information resolution of the instrument is 0.23 nm and the acceleration voltage is 300 kV.

Argon ion etching-assisted XPS spectra are received using a Thermo Fisher Scientific Model K-Alpha spectrometer equipped with Al Kα radiation (1486.6 eV). Time-of-flight secondary-ion mass spectrometry (TOF-SIMS) is performed on a PHI nano TOF III instrument. A 3 kV, 2 nA Bi3+ beam rastered over 100 µm is served as the analysis probe, and a 2 kV, 100 nA Ar⁺ beam covering 400×400 µm² erode the sample for depth profiling. The erosion rate is calibrated at ~9.16 nm min⁻¹ on SiO2. The samples sent for XPS and TOF-SIMS tests were obtained by depositing 1.5 mAh cm−2 at 0.5 mA cm−2 on Cu after cycling at 1 mA cm−2 and 1 mAh cm−2 for 5 cycles at 30 °C.

Nernst equation:

E = E0 − (RT/tF) ln (α reduction / α oxidation)

where R denotes the universal gas constant (8.314 J K−1 mol−1), T stands for the absolute temperature, α represents the chemical activity of each species (α reduction for the reduced form, α oxidation for the oxidized form). F is the Faraday constant (9.6485×104 J V−1 mol−1), t is the number of moles of electrons exchanged. For dilute solutions, activity α can be approximated by concentration c, yielding the simplified format 298.15 K:

E = E0 − (0.0592/t) log (1 / α oxidation)

Theoretical calculations
Classical molecular dynamics (MD) calculations are performed by using Forcite package in Materials Studio software 2020. The electrolyte modes of BE+Cs and BE are simulated by adopting amorphous cells with 56.2×56.2×56.2 Å3 and 56.4×56.4×56.4 Å3 linear dimensions, corresponding to salt-solvent ratios of 100:10:5:470:470 for 1.0 M LiTFSI+0.1 M LiNO3+0.05 M CsTFA-DOL/DME, and 100:10:470:470 for 1.0 M LiTFSI+0.1 M LiNO3-DOL/DME, respectively. With a fixed timestep of 1.0 fs (femtosecond), the statistical averages are calculated from trajectories of at least 50 ps (picosecond) in length after 0.1 ps equilibration steps. The system temperature is set to 298 K with Nose thermostat. All production simulations are carried out in isothermal-isovolumetric (NVT) ensemble using the COMPASS III force field. DFT calculations are performed by DMol3 module Materials Studio software 2020. Geometries optimizations and electrostatic potential are performed with parameter hybrid method using the LEE-Yang-Parr correlation functional (B3LYP). Convergence is enforced as the maximum force falls below 2.0×10−3 Ha Å−1 (1 Ha = 27.21 eV), the total-energy change below 1.0×10−5 Ha, and the largest displacement below 5.0×10−3 Å. An all-electron DNP basis set was employed together with a 0.005 au smearing to accelerate SCF convergence. Calculations are conducted with an all-electron DNP basis set with a 0.005 au smear for faster convergence. The values of binding energies (Ebinding) of Li+ with anion or solvent are calculated by the following formula:

Ebinding = Etotal − E1 − E2

where E is the energy of the Li+-solvent complex, single solvent, respectively.

Reduction potentials are performed by Gaussian 16 package from M05-2X calculations within the SMD implicit solvation model (acetone, ε=20.49). The Li+-complex geometry optimizations are performed with the pre-equilibrium structure at the B3LYP density functional and 6-31+G** basis set. The reduction potentials (Ereduction) of complex within the SMD implicit solvation are calculated by the following formula60–62:

Ereduction = −[ΔGgas(M−) − ΔGgas(M0) + ΔGsol(M−) − ΔGsol(M0)] / F − 1.4

where ΔGgas(M-), ΔGgas(M0) are the free energies of reduced and neutral complexes in gas-phase at 298.15 K. ΔGsol(M-), ΔGsol(M0) are the free energies of reduced and neutral complexes within the SMD implicit solvation model. F is the Faraday constant (9.6485×104 J V−1 mol−1).

Peer review rounds
[{"round": 1, "reviewer_comments": {"reviewer_1": "In this manuscript, the authors propose that Cs+, which is widely recognized for its ability to induce electrostatic shielding effects, can also facilitate the formation of LiF-rich SEI through its de-solvation and re-solvation behavior at the interface. They claimed that “solvated Cs+ adsorbed on the Li anode surface by electric field can attract less solvent but more anions to participate in the reductive reactions during discharge, thus dominating the original inorganic-rich SEI formation.” However, this “new” role of Cs+ has already been revealed in a recent study by BNL (Nat. Commun. 14, 8414 (2023)). The BNL work proposed that “When the Cs+ is present, its large size (Li+ radius: 0.59 Å; Cs+: 1.67 Å) greatly weakens the solvation interaction between DME solvent and Cs+ cation. Also because of its large size, Cs+ can be stabilized in a larger solvation sheath than that for the case of Li+. As a result, more anions can enter the solvation sheath and contribute to the interphase formation.” The mechanistic insights provided in this study is highly similar to the claim in the work by BNL. While it is possible that the authors independently arrived at the idea of Cs+ playing a novel role beyond electrostatic shielding (without referencing the BNL work), it is unfortunate that the current presentation of their findings can be regarded as a follow-up to the earlier study by BNL. One difference between this work and the BNL study is the use of TFA- as the anion in the additive, rather than NO3-. However, the selection of anions with high DN as additives is already a well-established strategy in the field of LMBs. While the authors may argue that TSA- is a superior anion additive compared to NO3-, citing its higher DN and lower lowest LUMO energy level, the performance gap between these two anions is likely limited from a theoretical perspective as follows. (I). Although TSA- has a significantly higher DN (34) than NO3- (22.2), the binding energies of Li+-TSA- (-6.98 eV) and Li+-NO3- (-6.91 eV) are comparable. This suggests that the ability of these anions to coordinate with Li+ ions is similar. (II). The authors’ DFT calculations show that the LUMO energies of all the anions studied are lower than those of the organic solvents, indicating that all anions will be reduced prior to the solvents. Therefore, comparing the LUMO energies of the anions is not meaningful for predicting their reduction behavior. Moreover, the LUMO concept itself is not a reliable descriptor for accurately predicting reduction potentials, as highlighted in prior literature (Energy Environ. Sci., 2018, 11, 2306). Given these considerations, I regret to say that I cannot recommend this manuscript to be published in Nature Communications.", "reviewer_2": "In this work, the authors propose the use of a low ion-pairing salt - CsTFA to foster a LiF-rich solid electrolyte interphase (SEI) that enables improved stripping and plating of the Li-metal anode. The authors show that the Coulombic efficiency (CE) of Li stripping/plating is improved as the alkali metal ion size increases. The authors demonstrate that the large size of Cs+-ion facilitates faster desolvation and plates with a compact inorganic-rich SEI through various characterization techniques. The cell performance is under different amounts of Li excess (varying N/P) highlights the drastic benefits of introducing the CsTFA additive. The correlation and linear fit of Li-excess to cycle life is interesting as well. This manuscript can be accepted after some major revisions based on the following comments -\n1. Previous literature reports have shown the benefits of adding Cs+ cation. Why did the authors choose TFA- as the preferred anion? Why not other high DN/ weakly pairing/ F-rich anions such as bromide, Iodide, tetrafluoroborate, nitrate? It is important to show experimentally the effect of the anion across donor numbers as this is something that the authors claim using theory. Aurbach CE tests with various anions is very critical to highlight the importance of TFA- anion.\n2. Part of the reason for using Cs+ is to also form a good seeding layer for Li deposition. It would be useful to see cross-section SEMs of the anodes with increasing amount of Li-deposition that highlight the critical role played by Cs+.\n3. It is great that the authors have used TOF-SIMS to highlight the depth-profile of various SEI components with BE+Cs electrolyte. however, it would be useful and important to contrast the SEI difference between BE and BE+Cs electrolytes.\n4. While the authors calculate the N/P using only the excess lithium available at the anode, this is not the right definition. The accepted definition is the total lithium capacity in the cell (N) versus the practical positive electrode capacity (P) [https://advanced.onlinelibrary.wiley.com/doi/full/10.1002/aesr.202000110]. Therefore, an anode-free cell would have an N/P of 1. This way of calculating also clearly represents the percent of Li-excess in the cell. It is highly recommended that the authors use this method of calculating N/P.\n5. If the Li-excess is calculated as described above, it'll also be very interesting to calculate the Li inventory loss rate. This is a great way to showcase the high CE of the electrolytes.\n6. It is interesting that there is capacity fade observed in anode-free cells without the loss in CE. How can the authors explain this?\n7. There are several typos and grammatical errors in the manuscript. It is recommended that they are addressed prior to publication.", "reviewer_3": "In this work, the authors have proposed and validated an innovative cation-desolvation mechanism involved in SEI formation in lithium-metal batteries. The proposed methodology demonstrates excellent electrochemical performance at low N/P ratios, which holds particular promise for practical energy density improvements. The mechanistic insights into how heteroalkali cations modulate solvation structures provide a valuable framework for advancing electrolyte additive design. Overall, I think this work would broaden the perspective on interfacial processes and the desolvation of electrodes. The presented results are interesting and significant. The manuscript is well-written and informative; however, I would like to offer some suggestions for minor revisions to further improve its clarity and impact.\n1. The authors have emphasized the relationship between the cation desolvation and SEI generation, how to define the desolvation and SEI generation? And what's the subordination between them?\n2. How to understand the statement in the second paragraph of Introduction: “those with larger radii display more pronounced solvent-phobic characteristics and allow them to approach the electrode surface more closely”? Indeed, why cations of larger radii facilitate closer proximity to the electrode surface? The statement here is vague, and a brief explanation of the relevant reason would be beneficial for the readers.\n3. Suggestions for rephrasing: in the first paragraph of Introduction, the sentence “ultimately leading to dendritic growth and low coulombic efficiency (CE) during Li plating and stripping” might be revised for better precision. Specifically, “dendritic growth” and “low coulombic efficiency” are not parallel concepts but rather sequential consequences. I would recommend rephrasing this sentence to more accurately reflect the progression of these phenomena.\n4. In Fig. 1: the legend mentions the role of the “electric field”, but this term is not clearly defined in Fig. 1. Adding a brief description or annotation of the “electric field” in the figure would help the readers better understand its significance in the context of the illustration.\n5. Suggestions for revision of the legends: in Fig. 2b, the legend mentions “binding energies to Na+”, which should be revised to “binding energies to Li+” to match the context of this work. Additionally, the legends (a-h) of Fig. 6 appear to be incorrect, likely due to a rearrangement of the figure without corresponding updates to the legends. I would suggest to double-check and revise these legends to ensure consistency and accuracy.\n6. In Fig. 5: in diagrams (e) and (f), the labels for the thickness of the SEI layer are inconsistent. Specifically, “SEI layer ~25nm” in Fig. 5f and “SEI layer ~ 17nm” in Fig. 5e. Please revise them into an identical format.\n7. Electrochemical impedance data of full cell: this work presents the excellent full-cell performance for the extreme anode-poor case, but it lacks data on the electrochemical impedance variation of full cells. This information would be highly valuable in substantiating the claimed improvement in the interface stability by the additives. I would suggest the authors consider supplementing additional work with impedance analyses to provide a more comprehensive understanding of the electrochemical behavior and the interface stability of their system."}, "author_response": "Response to referees’ comments\nDear Referees:\nThank you very much for your precious time in reviewing our manuscript (Manuscript\nnumber: NCOMMS-25-05148-T) and providing quite valuable comments. We have read\nyour comments carefully and tried our best to revise manuscript accordingly. We think all\ncomments are quite helpful and important to raise the quality of the revised manuscript. All\ncomments have been point-by-point replied as below and the revised parts have been\nhighlighted in the revised manuscript with YELLOW background. We sincerely look forward\nto the reviewer's approval and expect this revised manuscript can reach the high standard\nof“Nature communication”.\nReviewers’ comments:\nReviewer #1 (Remarks to the Author):\nIn this manuscript, the authors propose that Cs+, which is widely recognized for its ability\nto induce electrostatic shielding effects, can also facilitate the formation of LiF-rich SEI\nthrough its de-solvation and re-solvation behavior at the interface. They claimed that\n“solvated Cs+ adsorbed on the Li anode surface by electric field can attract less solvent\nbut more anions to participate in the reductive reactions during discharge, thus dominating\nthe original inorganic-rich SEI formation.”\nResponse: Thanks for the careful review of the manuscript and pertinent comments!\nHowever, this “new” role of Cs+ has already been revealed in a recent study by BNL (Nat.\nCommun. 14, 8414 (2023)). The BNL work proposed that “When the Cs+ is present, its\nlarge size (Li+ radius: 0.59 Å; Cs+: 1.67 Å) greatly weakens the solvation interaction\nbetween DME solvent and Cs+ cation. Also because of its large size, Cs+ can be stabilized\nin a larger solvation sheath than that for the case of Li+. As a result, more anions can enter\nthe solvation sheath and contribute to the interphase formation.” The mechanistic insights\nprovided in this study is highly similar to the claim in the work by BNL.\nWhile it is possible that the authors independently arrived at the idea of Cs+ playing a novel\nrole beyond electrostatic shielding (without referencing the BNL work), it is unfortunate that\nthe current presentation of their findings can be regarded as a follow-up to the earlier study\nby BNL.\nOne difference between this work and the BNL study is the use of TFA- as the anion in the\nadditive, rather than NO3-. However, the selection of anions with high DN as additives is\nalready a well-established strategy in the field of LMBs. While the authors may argue that\nTSA- is a superior anion additive compared to NO3-, citing its higher DN and lower lowest\nLUMO energy level, the performance gap between these two anions is likely limited from\na theoretical perspective as follows. (I). Although TSA- has a significantly higher DN (34)\nthan NO3- (22.2), the binding energies of Li+-TSA- (-6.98 eV) and Li+-NO3- (-6.91 eV) are\ncomparable. This suggests that the ability of these anions to coordinate with Li+ ions is\nsimilar. (II). The authors’ DFT calculations show that the LUMO energies of all the anions\nstudied are lower than those of the organic solvents, indicating that all anions will be\nreduced prior to the solvents. Therefore, comparing the LUMO energies of the anions is\nnot meaningful for predicting their reduction behavior. Moreover, the LUMO concept itself\nis not a reliable descriptor for accurately predicting reduction potentials, as highlighted in\nprior literature (Energy Environ. Sci., 2018, 11, 2306).\nGiven these considerations, I regret to say that I cannot recommend this manuscript to be\npublished in Nature Communications.\nResponse: Thank you for your careful reviewing and insightful comments. We appreciate\nthe foundational contributions of prior research, including the classical electrostatic\nshielding effect proposed by Zhang et al (J. Am. Chem. Soc. 135, 4450−4456 (2013)) and\nthe innovative non-LiF but CsFSI-dominating SEI theory proposed by Hu et al (Nat\nCommun 14, 8414 (2023)).\nIn our opinion, BNL's work leveraged CsNO as an additive to simultaneously create stable\n3\ninterfaces on both the Li anode and the NCM cathode, thereby implementing a dual-\nprotection strategy. More in-depth, their research delves deeper into the use of advanced\ninterfacial characterization techniques, such as X-ray fluorescence, to demonstrate that\nCsFSI is the predominant component in the SEI, distinct from the low-ionic-conductivity LiF\nand is never been reported. This novel CsFSI-dominated SEI imparts fast-charging\ncapabilities to the battery, enabling both fast charging and extended cycle life for NCM\nbatteries in their work. In short, the CsFSI-dominated SEI, which is free of LiF, is the\nhighlight of their research. Our work, however, has focused on the detailed de-solvation\nprocess of hetero-cations and its impact on the SEI formation on the Li anode. We have\nemployed in situ spectroscopic techniques to thoroughly investigate and visualize this\nunique de-solvation process. And our research represents a synergy between performance\noptimization and mechanistic understanding, rather than being antagonistic to BNL's work.\nIn fact, many optimized properties and phenomena cannot be fully explained by a single\nmechanism alone. When it comes to explaining the improvement in battery performance\nattributed to Cs cations, it is not appropriate to solely credit the mechanism proposed by\nBNL. Our analysis of the de-solvation process may also play a significant role. The\nmechanisms we propose could be complementary to those of BNL, without any inherent\nconflict. In reality, BNL's approach was more result-oriented and did not deeply explore the\nsolvation/de-solvation process of the large cation Cs+. In contrast, our work is more\nprocess-oriented, focusing on the initial de-solvation process that occurs during battery\noperation. We believe that our work will provide a more comprehensive understanding of\nthe underlying mechanisms and contribute to the development of more efficient and stable\nbattery systems.\nWhile we acknowledge the significant contributions of BNL, we would like to address an\nambiguity in the reviewer's citation of BNL's work. Specifically, it should be noted that large\ncations, due to their more concentrated charges and weaker interactions with solvent\nmolecules, form solvated structures with fewer solvent molecules, fewer solvation layers\nand a smaller overall size of solvated structure. This phenomenon has been well-\ndocumented in several studies (Chem catalysis 2, 1267-1276 (2022); J. Am. Chem. Soc.\n144, 1589-1602 (2022); Phys. Chem. Chem. Phys. 22 16301-16313 (2020)).\nIn contrast to BNL's work, which employed the NO - anion primarily to function on the\n3\ncathode electrode, we have selected TFA- to predominantly act on the anode electrode.\nOur strategic choice aims to enhance the process of SEI formation by attracting more\nanions, including TFA-, to participate in the reduction process associated with Cs+ de-\nsolvation.\nOur rationale for selecting TFA- as the preferred anion is multifaceted, and the ultrahigh\nDN value is not the only reason why we chose it. Firstly, TFA- is a highly polar, all-fluorinated\nanion that exhibits an ultrahigh electron-donating capacity and DN value among available\nLi salts. This enables TFA- to interact strongly with Li⁺, allowing it to preferentially enter the\ninner solvation shell (PNAS 121 (5), e2316914121 (2024); Small 20, 2408164 (2024)).\nAdditionally, we have considered the potential inorganic derivatives of TFA- that may form\nduring the initial reduction process, such as LiF and Li O. These compounds are known to\n2\nenhance the compatibility of the Li anode and improve the stability of the SEI properties\n(Nat Energy 9, 817-827 (2024); Nat Commun 12, 186 (2021)). Moreover, TFA- salts are\ncharacterized by their low lattice enthalpy, as reported in the literature (J. Phys. Chem. B\n116, 36, 11323-11331 (2012); J. Comput. Chem. 22, 827-834 (2001)). This property\ncontributes to their excellent solubility in aprotic solvents, which is another critical factor in\nour selection. In summary, while the high DN value of TFA- is a significant advantage, it is\nnot the sole reason for our choice. The combination of its strong interaction with Li⁺, the\nbeneficial inorganic derivatives formed during reduction, and its favorable solubility\ncharacteristics collectively make TFA- an ideal candidate for our research.\nRegarding the reviewer's statement: “the ability of these anions to coordinate with Li+ ions\nis similar”. We appreciate the reviewer’s insights regarding the DN value, binding energy\nand Li+-anion interaction. While there are indeed reports indicating that Li+-TFA- has a\nlower binding energy compared to Li+-NO - (Angew. Chem. Int. Ed. 63, e202314876\n3\n(2024)). Besides, the literature also provides direct evidence that TFA-, with its higher DN\nvalue (34), preferentially coordinates with Li+ over NO - (DN value of 22.2) (J. Phys. Chem.\nAdditionally, several studies have highlighted that the strong interaction between the\ncarbonyl group (C=O, in TFA-) and Li+ promotes the dissociation of LiNO within the solvent\n3\n(Angew. Chem. Int. Ed. 63, e202403668 (2024); Nano Energy 104, 107881(2022); Adv.\nEnergy Mater. 10, 1903843 (2022)). Therefore, we argue that the binding of TFA- to Li+ is\nstronger than that of NO - to Li+.\nconcept. Indeed, the reduction priority of anions in an actual battery environment is\ninfluenced by multiple factors, including the electric field and concentration. Density\nfunctional theory (DFT) calculations can only serve as a supplementary tool to illustrate the\nreduction tendencies of anions. Referring to the literature noted by the reviewer (Energy\nEnviron. Sci., 11, 2306 (2018)), it’s mentioned that: “it is more correct to speak of potential\nof electrolyte reduction at negative potentials”. Therefore, to further support our arguments,\nwe have conducted linear sweep voltammetry (LSV) analysis on the studied electrolytes\n(Supplementary Fig. 14). The results show that the BE+Cs electrolyte containing TFA-\nexhibits a reduction peak at 1.2-1.0 V, which corresponds to the reduction of TFA- (Angew.\nChem. Int. Ed. 135, e202306236 (2023); Small 20, 2408164 (2024)). This finding aligns\nwith the theoretical calculation results, confirming that TFA- has a higher reduction priority.\nWe thank the reviewers again for their careful reading and valuable feedback. We hope\nthat our response clarifies the concerns and demonstrates the validity of our approach.\nIn response to your comments, we would like to highlight how our study contributes to\nexisting knowledge and advances the field in several key ways below:\nFirstly, we have proposed a novel mechanism for the formation of a stable SEI based on\nthe de-solvation of heteroalkali cations, rather than relying solely on the conventional\nelectrostatic shielding effect. By Employing electrochemical in-situ FT-IR characterization,\nwe directly visualize the dynamic involvement of more anions and fewer free solvent during\nthe plating process of Li metal electrode. This provides a deeper understanding of the SEI\nformation mechanism at the molecular level.\nSecondly, both the cation and anion of the CsTFA additive we proposed play a significant\nrole in enhancing the performance of the Li anode. Attributed to the synergistic effect, our\nLi|Cu half cells achieve remarkable CE (99.17% of CE and 99.77% of CE ) in\nCycle Aurbach\ncommercial ether electrolyte of normal concentration. This level of Li utilization efficiency\nis exceptional compared to other electrolyte systems reported in the literature, including\nboth conventional normal-concentration and high-concentration/locally high-concentration\nelectrolyte systems.\nThirdly, our Li|LFP full cell (with an N/P ratio of 1.36) and lean-Li|LFP full cell (with an N/P\nratio of 0.1-0.4) demonstrate impressive cycling stability. Notably, achieving high stability\nfor Li|LFP full cells at an N/P ratio of 0.1-0.4 has not been realized in the reported literature.\nFurthermore, our proposed linear correlation and fit of Li-excess to cycle life provide a more\naccurate indication of electrolyte performance and enable better prediction of battery\nlifetime.\nIn summary, our study not only offers a novel perspective on SEI formation but also delivers\nsignificant practical improvements in Li battery performance. We believe these\ncontributions will advance the field and inspire further research.\nThank you for your valuable time and consideration, and we remain open to any further\nquestions or feedback.\nThe corresponding content has been incorporated into the revised manuscript.\n----------------------------------------------------------------------------------------------------------------------\nSupplementary Fig. 14 Linear sweep voltammetry (LSV) of Li|Cu half cells in BE+Cs\nand BE electrolyte within the voltage range from 3.0 V to 0 V at 5 mV s-1.\n(See Supplementary Fig. 14)\n----------------------------------------------------------------------------------------------------------------------\nReviewer #2 (Remarks to the Author):\nIn this work, the authors propose the use of a low ion-pairing salt - CsTFA to foster a LiF-\nrich solid electrolyte interphase (SEI) that enables improved stripping and plating of the Li-\nmetal anode. The authors show that the Coulombic efficiency (CE) of Li stripping/plating is\nimproved as the alkali metal ion size increases. The authors demonstrate that the large\nsize of Cs+-ion facilitates faster de-solvation and plates with a compact inorganic-rich SEI\nthrough various characterization techniques. The cell performance is under different\namounts of Li excess (varying N/P) highlights the drastic benefits of introducing the CsTFA\nadditive. The correlation and linear fit of Li-excess to cycle life is interesting as well. This\nmanuscript can be accepted after some major revisions based on the following comments:\nResponse: We sincerely appreciate your positive feedback on our manuscript. Based on\nyour detailed comments, we have made improvements to our work. Our responses to your\ncomments are provided below. We hope these revisions meet your expectations.\n1. Previous literature reports have shown the benefits of adding Cs+ cation. Why did the\nauthors choose TFA- as the preferred anion? Why not other high DN/ weakly pairing/\nF-rich anions such as bromide, Iodide, tetrafluoroborate, nitrate? It is important to\nshow experimentally the effect of the anion across donor numbers as this is something\nthat the authors claim using theory. Aurbach CE tests with various anions is very\ncritical to highlight the importance of TFA- anion.\nResponse: Thank you very much for highlighting this important issue. We sincerely hope\nour responses can dispel your doubts and meet your satisfaction.\nCs+-containing salts have indeed been widely reported as electrolyte additives in numerous\nstudies before (J. Am. Chem. Soc. 135, 4450-4456 (2013); Nano Energy 134, 110582\n(2025); Chem. Eng. J. 451, 138580 (2023); Adv. Energy Mater. 12, 2102834 (2022); Nat\nCommun 14, 8414 (2023)). These studies have primarily focused on the conventional\nelectrostatic shielding effect to inhibit dendritic growth and the newly discovered non-LiF-\nSEI inorganic compounds (CsFSI). In contrast, our work aims to explore the intrinsic effects\nof Cs+ salt additives on the de-solvation of the electrolyte and the SEI generation process\nthrough in-situ spectroscopy techniques combined with direct physical and chemical\ncharacterizations.\nOur rationale for selecting TFA- as the preferred anion is based on several key\nconsiderations:\n1. High polarity and electron-donating capacity: TFA- is a highly polar, all-fluorinated anion\nwith an exceptionally high electron-donating capacity and the highest DN value among\navailable Li salts (PNAS 121 (5), e2316914121 (2024); Small 20, 2408164 (2024)). This\nenables TFA- to interact strongly with Li+, allowing it to preferentially enter the\ninner solvation shell.\n2. Lower LUMO energy level: theoretical calculations and previous literature (Adv. Energy\nMater. 10, 1903843 (2020); Angew. Chem. Int. Ed. 63, e202403668 (2024); Angew. Chem.\nInt. Ed. 63, e202314876 (2024)) indicate that TFA- has a relatively lower LUMO energy\nlevel compared to solvents. This property endows TFA- with a distinct voltage response\neffect and preferential reductive decomposition behavior compared to solvent molecules\nand other Li salts.\n3. Formation of beneficial inorganic derivatives: during the initial reduction process, TFA-\ncan form inorganic derivatives such as LiF and Li O. These compounds are known to\n2\nenhance the compatibility of the Li anode and improve the stability of the SEI properties\n(Nat Energy 9, 817-827 (2024); Nat Commun 12, 186 (2021)).\n4. Solubility in aprotic solvents: as reported in the literature (J. Phys. Chem. B 116, 36,\n11323-11331 (2012); J. Comput. Chem. 22, 827-834 (2001)), TFA- salts exhibit low lattice\nenthalpy, which ensures good solubility in aprotic solvents. This is crucial for their\ncompatibility with our electrolyte system.\nIn summary, our decision to use CsTFA was driven by its high polarity, reduction sensitivity,\nperfluorinated structure, ease of availability, lower cost, and moderate solubility in ether\nelectrolytes. These factors collectively make CsTFA an ideal choice for our research.\nWe appreciate your suggestion to experimentally validate the effect of the anion across\ndifferent donor numbers. In response, we have supplemented the Aurbach CE tests with\nvarious anions in our studied BE and SE electrolytes (Supplementary Fig. 1 and 2). The\nresults indicate that, arise from the anions of commonly used Li salts (PF -, TFSI- and FSI-),\n6\nother anions (BF -, CLO -, OTF-, NO - and TFA-) generally exhibit higher Aurbach CE\n4 4 3\nvalues as their DN values increase. Notably, TFA-, which has the highest DN value among\nthese anions, achieves the highest Aurbach CE. This observation underscores the\nsignificant role of the anion’s electron-donating capacity in enhancing Li plating/stripping\nefficiency.\nThe corresponding content has been incorporated into the revised manuscript.\n----------------------------------------------------------------------------------------------------------------------\nSupplementary Fig. 1 Aurbach CE of Li|Cu half cells in SE electrolyte with 0.05 M\nLiPF , 0.05 M LiTFSI, 0.05 M LiBF , 0.05 M LiClO , 0.05 M LiFSI, 0.05 M LiOTF,\n6 4 4\n0.05 M LiNO and 0.05 M LiTFA at 0.5 mA cm-2 with a fixed capacity of 1.0 mAh cm-\n3\n2.\nSupplementary Fig. 2 Aurbach CE of Li|Cu half cells in BE electrolyte with 0.05 M\nLiPF , 0.05 M LiTFSI, 0.05 M LiBF , 0.05 M LiClO , 0.05 M LiFSI, 0.05 M LiOTF,\n6 4 4\n0.05 M LiNO and 0.05 M LiTFA at 0.5 mA cm-2 with a fixed capacity of 1.0 mAh cm-\n3\n2.\n(See Supplementary Fig. 1 and 2)\n----------------------------------------------------------------------------------------------------------------------\n2. Part of the reason for using Cs+ is to also form a good seeding layer for Li deposition.\nIt would be useful to see cross-section SEMs of the anodes with increasing amount\nof Li-deposition that highlight the critical role played by Cs+.\nResponse: Thank you for your comments regarding the need for cross-sectional SEM\nimages to highlight the critical role of Cs⁺ in Li deposition. In response, we have conducted\nadditional experiments to obtain cross-sectional SEM images of the anodes with varying\namounts of Li deposition (1, 2, 3 and 5 mAh cm-2) (Supplementary Fig. 26). The results\nreveal that, with the addition of CsTFA, the deposited Li evolves into a much larger and\nless voided morphology along the cross-sectional direction. Moreover, the denser Li\ndeposits also result in thinner overall thicknesses. These findings highlight the significant\nimpact of CsTFA on the morphology of Li deposits, which is crucial for enhancing the\nstability and performance of the Li anode.\nBeyond the effect of the seeding layer, our work primarily investigates the de-solvation of\nheteroalkali cations (Cs+) and its influence on the chemical and physical composition of the\nThe corresponding content has been incorporated into the revised manuscript.\n----------------------------------------------------------------------------------------------------------------------\nSupplementary Fig. 26 Cross-sectional SEM images of the deposited Li obtained from\nLi|Cu half cells in (a) BE+Cs, (b) BE, (c) SE+Cs and (d) SE electrolyte at 0.5 mA cm-2\nwith a fixed capacity of 1.0 mAh cm-2, 2.0 mAh cm-2, 3.0 mAh cm-2 and 5.0 mAh cm-\n2.\n(See Supplementary Fig. 26)\n----------------------------------------------------------------------------------------------------------------------\n3. It is great that the authors have used TOF-SIMS to highlight the depth-profile of\nvarious SEI components with BE+Cs electrolyte. however, it would be useful and\nimportant to contrast the SEI difference between BE and BE+Cs electrolytes.\nResponse: Thank you very much for your valuable comments. We appreciate your\nsuggestion to contrast the difference in the SEI between BE and BE+Cs electrolytes. We\nhave supplemented the TOF-SIMS results of the SEI on deposited Li in BE electrolyte\n(Supplementary Fig. 31, 34 and 35). During negative mode sputtering, detected secondary\nion fragments such as LiF-, LiO-, Li N-, CO -, C HO- are characteristic of LiF, Li O, Li N,\n3 3 2 2 3\nLi CO and organic species, respectively. By comparing the differential intensity depth\n2 3\nprofiles of TOF-SIMS results between BE+Cs and BE electrolytes, we observed that the\nSEI in BE+Cs electrolyte contains a higher concentration of inorganic components (LiF,\nLi O, Li N) and exhibits more active Li beneath the SEI layer. In contrast, the SEI in BE\n2 3\nelectrolyte is enriched with organic components and a higher amount of Li CO , which is\n2 3\ngenerally considered detrimental to SEI stability. These supplemented results are\nconsistent with the Cryo-TEM images and depth-profiling XPS spectra results.\nThe corresponding content has been incorporated into the revised manuscript.\n----------------------------------------------------------------------------------------------------------------------\nSupplementary Fig. 31 (a,b) The depth profiles of TOF-SIMS for several interested\nsecondary ion fragments obtained by sputtering SEI on deposited Li in (a) BE+Cs and\n(b) BE electrolyte. (c-h) The differential intensity depth profiles of TOF-SIMS for (c)\nLi-, (d) LiF-, (e) Li N-, (f) LiO-, (g)C HO- and (h) CO - fragments of SEI in BE+Cs and\n3 2 3\nBE electrolyte.\nSupplementary Fig. 34 (a-c,g-i) The depth profile of Li-, LiF-, LiO-, Li N- , C HO- and\n3 2\nLiF - secondary ion fragments of SEI formed in BE electrolyte. (d-f,j-l) Corresponding\n2\n3D distributions.\nSupplementary Fig. 35 (a-c) The depth profile of LiO -, LiN- and CO - secondary ion\n2 3\nfragments of SEI formed in BE electrolyte. (d-f) Corresponding 3D distributions.\n(See Supplementary Fig. 31, 34 and 35)\n----------------------------------------------------------------------------------------------------------------------\n4. While the authors calculate the N/P using only the excess lithium available at the\nanode, this is not the right definition. The accepted definition is the total lithium\ncapacity in the cell (N) versus the practical positive electrode capacity (P)\n[https://advanced.onlinelibrary.wiley.com/doi/full/10.1002/aesr.202000110]. Therefore,\nan anode-free cell would have an N/P of 1. This way of calculating also clearly\nrepresents the percent of Li-excess in the cell. It is highly recommended that the\nauthors use this method of calculating N/P.\nResponse: Thank you for your thoughtful comment regarding the calculation of the N/P\nratio. We truly appreciate your input and the reference provided. In fact, the definition of\nN/P ratio in the view supported by the reviewer and the one used in our previous\nmanuscript both describe the same fundamental concept: the capacity ratio between the\nanode and the cathode. The difference lies in the starting point of the calculation. While\nthe reviewer suggests that the N should refer to the total Li inventory of the battery, our\napproach focuses on the capacity of the negative side of the battery.\nHowever, after careful consideration and a thorough review of the relevant literature, we\nrespectfully maintain our original approach to calculating the N/P ratio based on the\ncapacity of the negative electrode relative to that of the positive electrode. We believe this\napproach provides a clearer and more direct indication of the capacity balance between\nthe anode and cathode, which is crucial for evaluating the performance and lifetime of the\nfull cell. The N/P ratio we advocate is calculated by the following equation:\nN/P ratio = 𝐶𝑎𝑝𝑎𝑐𝑖𝑡𝑦 / 𝐶𝑎𝑝𝑎𝑐𝑖𝑡𝑦\n𝑁𝑒𝑔𝑎𝑡𝑖𝑣𝑒 𝑃𝑜𝑠𝑖𝑡𝑖𝑣𝑒\nThank you again for your valuable feedback! We remain open-minded regarding the\ncalculation of N/P ratio and welcome further insights from reviewer.\nOur references are listed below.\n----------------------------------------------------------------------------------------------------------------------\nReference:\nNat Energy 19, 15008 (2016) [High-capacity battery cathode prelithiation to offset initial\nlithium loss | Nature Energy];\nNat Energy 7, 1031-1041 (2022) [Quantification of reversible and irreversible lithium in\npractical lithium-metal batteries | Nature Energy];\nNat Energy 3, 889-898 (2018) [Langmuir–Blodgett artificial solid-electrolyte interphases for\npractical lithium metal batteries | Nature Energy];\nNat Energy 10, 191-204 (2025) [Breaking the molecular symmetricity of sulfonimide anions\nfor high-performance lithium metal batteries under extreme cycling conditions | Nature\nEnergy];\nNature 627, 101-107 (2024) [Ligand-channel-enabled ultrafast Li-ion conduction | Nature];\nNature 614, 694-700 (2023) [Electrolyte design for Li-ion batteries under extreme operating\nconditions | Nature];\nScience 375, 739-745 (2022) [Self-assembled monolayers direct a LiF-rich interphase\ntoward long-life lithium metal batteries | Science];\nScience 357, 279-283 (2017) [Highly elastic binders integrating polyrotaxanes for silicon\nmicroparticle anodes in lithium ion batteries | Science];\nJ. Am. Chem. Soc. 146, 31, 21721-21728 (2022) [Simple Framework for Simultaneous\nAnalysis of Both Electrodes in Stoichiometric Lithium–Sulfur Batteries | Journal of the\nAmerican Chemical Society];\nJ. Am. Chem. Soc. 144, 1, 212-218 (2022) [A Successive Conversion–Deintercalation\nDelithiation Mechanism for Practical Composite Lithium Anodes | Journal of the American\nChemical Society];\nJoule 4, 2, 285-291 (2020) [Lithium-Sulfur Batteries: Attaining the Critical Metrics -\nScienceDirect];\nJoule 5, 5, 1011-1015 (2021) [Achieving highly reproducible results in graphite-based Li-\nion full coin cells - ScienceDirect];\n----------------------------------------------------------------------------------------------------------------------\n5. If the Li-excess is calculated as described above, it'll also be very interesting to\ncalculate the Li inventory loss rate. This is a great way to showcase the high CE of\nthe electrolytes.\nResponse: Thank you for your suggestion regarding the calculation of the Li inventory loss\nrate (LILR). We have supplemented the calculated LILR of our lean-Li full cells in the\nrevised manuscript (Fig. 6e and 6f, Supplementary Fig. 41 and 42 and Supplementary\nTable 7.). The results indicate that the LILR values of our lean-Li full cells using the BE+Cs\nelectrolyte range from 0.138% to 0.395%. These relatively low values demonstrate the high\nstability and compatibility of the electrolyte with the Li anode.\nThe formula of LILR we use is shown below:\nLILR = 100×(1-𝑛−√𝑡 𝑄 /𝑄 )\n𝑛 𝑡\nwhich is referenced from the literature: Energy Environ. Sci. 13, 2501-2514 (2020) [Lithium\ndegradation in lithium–sulfur batteries: insights into inventory depletion and interphasial\nevolution with cycling - Energy & Environmental Science (RSC Publishing)]\nThe corresponding content has been incorporated into the revised manuscript.\n----------------------------------------------------------------------------------------------------------------------\nFig. 6 | e,f, Long-term cycling performance and calculated LILR of Li|LFP cells with\nultra-low N/P ratios at 0.2 C formation and 0.5 C cycling.\n(See Fig. 6e and 6f)\nSupplementary Fig. 41 Long-term cycling performance and calculated LILR of\nLi|LFP cells with N/P ratios from 0.104 to 0.263 at 0.2 C formation and 0.5 C cycling\nin BE+Cs electrolyte.\nSupplementary Fig. 42 Long-term cycling performance and calculated LILR of\nLi|LFP cells with N/P ratios from 0.280 to 0.415 at 0.2 C formation and 0.5 C cycling\nin BE+Cs electrolyte.\n(See Supplementary Fig. 41 and 42)\nSupplementary Table 7. The range of ultra-low negative/positive electrode capacity\n(N/P) ratios and the calculated average LILR.\nCapacity of Capacity of\nTransition point Average\nN/P ratios negative positive\ncycle number LILR (%)\n(mA h cm-2) (mA h cm-2)\n0.104 0.4 3.823 24 0.395\n0.107 0.4 3.723 28 0.280\n0.122 0.4 3.281 35 0.276\n0.128 0.4 3.111 38 0.239\n0.169 0.4 2.363 48 0.298\n0.20 0.4 1.989 53 0.169\n0.221 0.8 3.616 58 0.200\n0.263 0.8 3.043 62 0.186\n0.28 0.8 2.835 75 0.167\n0.309 0.8 2.584 94 0.138\n0.32 0.8 2.499 102 0.158\n0.34 0.8 2.346 112 0.138\n0.407 1.5 3.684 131 0.190\n0.415 1.5 3.616 147 0.197\n(See Supplementary Table 7.)\n----------------------------------------------------------------------------------------------------------------------\n6. It is interesting that there is capacity fade observed in anode-free cells without the loss\nin CE. How can the authors explain this?\nResponse: Thank you for your insightful question regarding the capacity fade observed in\nanode-free cells without a loss in CE. In fact, this phenomenon has indeed been noted in\nseveral previous studies, particularly in anode-free Cu|LFP cells using ether-based\nelectrolytes (Angew. Chem. Int. Ed. 62, 1-9, e202304978 (2023); Adv. Mater. 35, 1-12,\n2210130 (2023); Adv. Mater. 36, 1-10, 2401114 (2024); Adv. Mater. 36, 1-9, 2407648\n(2024)). However, these studies all lack detailed explanations for this phenomenon. To\naddress this, we have supplemented our analysis with calculated CE, capacity retention,\nand Li inventory retention rate (LIRR) for the Cu|LFP cell (Supplementary Fig. 45).\nAdditionally, an explanatory diagram is provided below to aid in understanding.\nThe phenomenon of capacity decay without a loss in CE is indeed intriguing. From our\nperspective, the CE and LIRR of the Cu|LFP cell are not consistently high throughout the\nentire cycling lifetime. Here is a detailed breakdown of the observed behavior:\n1. Initial cycle stage (0-25 cycles): In the initial cycles, the cell contains a partial excess of\nLi inventory compared to the cathode inventory. The discharge capacity, combined with the\nexcess Li on the Cu collector, constitutes the total active Li in the cell. Due to the fragile\nand unstable nature of the freshly formed SEI, the continued growth of the SEI and the\ngeneration of \"dead\" Li gradually deplete the active Li through intense ether electrolyte\ncorrosion. During this stage, both CE and LIRR are relatively low, and capacity decay is\nsignificant.\n2. Middle cycle stage (25-100 cycles): As the cycling progresses, the partial excess of Li\ninventory is gradually consumed. However, the ongoing process of SEI enhancement and\nthe continuous generation of \"dead\" Li continues to consume active Li. The active Li of\ninitial-turn discharge capacity is gradually consumed. As a result, CE remains low but\ngradually increases over time. Capacity decay persists, but at a reduced rate compared to\nthe initial cycle stage.\n3. Later cycle stage (100-200 cycles): In the later stages of cycling, the SEI becomes more\nreinforced and stable. The depletion of active Li is now primarily due to the generation of\n\"dead\" Li. Consequently, the capacity decay of the Cu|LFP cell continues but at a much\nmilder rate. Both CE and LIRR reach relatively high levels, with CE approaching 99.5%.\nIn summary, the capacity fade observed in anode-free cells without a significant loss in CE\ncan be attributed to the dynamic changes in SEI stability and the consumption of active Li\nduring different stages of cycling.\nWe hope this detailed explanation provides a comprehensive understanding of this\nphenomenon and addresses your concerns. We are open to further discussion and look\nforward to your continued feedback.\nThe formula of LIRR we use is shown below:\nLIRR = 100×(𝑄 /𝑄 )1/𝑛\n𝑛 𝑖𝑛𝑡\nwhich is referenced from the literature:\nAdv. Energy Mater. 11, 1-18 , 2000804, (2021) [Anode-Free Full Cells: A Pathway to High-\nEnergy Density Lithium-Metal Batteries - Nanda - 2021 - Advanced Energy Materials -\nWiley Online Library]\nThe corresponding content has been incorporated into the revised manuscript.\n----------------------------------------------------------------------------------------------------------------------\nSupplementary Fig. 45 (a) Cycling charge-discharge curves, (b) calculated CE, (c)\ncalculated capacity retention and (d) calculated Li inventory retention rate (LIRR) of\nanode-free Cu|LFP cells at 0.3 C charge and 0.5 C discharge.\n(See Supplementary Fig. 45)\n----------------------------------------------------------------------------------------------------------------------\n7. There are several typos and grammatical errors in the manuscript. It is recommended\nthat they are addressed prior to publication.\nResponse: Thank you for your careful review of our manuscript. We have conducted a\ncomprehensive language edit and made necessary formatting adjustments to ensure that\nthe manuscript meets the journal's standards. We believe these revisions have improved\nthe clarity and readability of the manuscript. We are open to further discussion and look\nforward to your continued feedback.\nReviewer #3 (Remarks to the Author):\nIn this work, the authors have proposed and validated an innovative cation-de-solvation\nmechanism involved in SEI formation in lithium-metal batteries. The proposed\nmethodology demonstrates excellent electrochemical performance at low N/P ratios, which\nholds particular promise for practical energy density improvements. The mechanistic\ninsights into how heteroalkali cations modulate solvation structures provide a valuable\nframework for advancing electrolyte additive design. Overall, I think this work would\nbroaden the perspective on interfacial processes and the de-solvation of electrodes. The\npresented results are interesting and significant. The manuscript is well-written and\ninformative; however, I would like to offer some suggestions for minor revisions to further\nimprove its clarity and impact.\nResponse: We sincerely appreciate your positive feedback on our manuscript. Based on\nyour detailed comments, we have made improvements to our work. Our responses to your\ncomments are provided below. We hope these revisions meet your expectations.\n1. The authors have emphasized the relationship between the cation de-solvation and\nSEI generation, how to define the de-solvation and SEI generation? And what's the\nsubordination between them?\nResponse: Thank you very much for your comments. We have provide detailed definitions\nand explanations regarding the subordination between these two processes below.\nCation de-solvation refers to the process where solvated cations (e.g., Li⁺ ions) lose their\nsolvent molecules as they approach the electrode surface or the SEI layer. This step is\ncrucial for the cations to be able to move through the SEI and participate in the charge\ntransfer reactions at the electrode surface (Nat Energy 10, 502-512 (2025); PNAS 121 (4)\n(2024); Nat Commun 14, 3536 (2023)).\nSEI generation refers to the reduction reactions of both organic and inorganic compounds\non the anode surface of the electrode during the initial cycles. The SEI plays a vital role in\nstabilizing the electrode and regulating ion transport (Nat Energy 6, 303-313 (2021); Nat\nCommun 16, 3813 (2025)).\nThe subordination relationship between cation de-solvation and SEI generation could be\nunderstood as follows: (1). during the initial charging process, the electrolyte undergoes\nreduction reactions at the electrode surface, leading to the formation of the SEI. (2). once\nthe SEI layer is formed, it influences the de-solvation process of cations. The chemical\ncomposition and structure of SEI affect how easily the cations lose their solvent molecules.\n(3). the ease of cation de-solvation, in turn, affects the stability and further evolution of SEI.\nEfficient de-solvation can lead to smoother ion transport, which could help in maintaining\nthe integrity of SEI. Conversely, if de-solvation is hindered, it may result in the formation of\nan unstable or incomplete SEI, impacting the overall performance and longevity of the\nbattery.\nIn summary, cation de-solvation and SEI generation are closely interconnected processes.\nThe formation of the SEI layer provides a critical environment for cation de-solvation, while\nthe efficiency of de-solvation influences the stability and further development of the SEI.\nWe have emphasized this relationship in our study to highlight the importance of\nunderstanding and optimizing these processes for better battery performance.\nWe hope this clarification addresses your concerns. We are open to further discussion and\nlook forward to your continued feedback.\n2. How to understand the statement in the second paragraph of Introduction: “those with\nlarger radii display more pronounced solvent-phobic characteristics and allow them to\napproach the electrode surface more closely”? Indeed, why cations of larger radii\nfacilitate closer proximity to the electrode surface? The statement here is vague, and\na brief explanation of the relevant reason would be beneficial for the readers.\nResponse: Thank you for your valuable comments. We understand that the statement may\nseem vague, the explanations are provided below.\nFor the explanation of statement: “those with larger radii display more pronounced solvent-\nphobic characteristics”: larger cations generally have less concentrated charges, that is,\nlower charge density compared to smaller cations (Chem catalysis 2, 1267-1276 (2022);\nNat Commun 14, 612 (2024)). As a result, the electrostatic attraction between the larger\ncations and the solvent molecules is weaker. This weaker interaction makes the larger\ncations to attract relatively fewer solvent molecules, exhibiting solvent-phobic\ncharacteristics.\nFor the explanation of statement: “cations of larger radii facilitate closer proximity to the\nelectrode surface”: because of the more concentrated charges and the weaker cation-\nsolvent interaction, solvated large cations contain less solvent molecules and smaller size\nof solvated structure (Chem catalysis 2, 1267-1276 (2022); J. Am. Chem. Soc. 144, 1589-\n1602 (2022); Phys. Chem. Chem. Phys. 22 16301-16313 (2020)). As a result, the smaller\nsolvated large cations could approach the electrode surface more easily under electric field.\n3. Suggestions for rephrasing: in the first paragraph of Introduction, the sentence\n“ultimately leading to dendritic growth and low coulombic efficiency (CE) during Li\nplating and stripping” might be revised for better precision. Specifically, “dendritic\ngrowth” and “low coulombic efficiency” are not parallel concepts but rather sequential\nconsequences. I would recommend rephrasing this sentence to more accurately reflect\nthe progression of these phenomena.\nResponse: We sincerely thank the reviewer for careful reading. As suggested, we have\nrevised the whole statement into: “The initially formed SEI is fragile and prone to fracture\nwhen undergoing displacements during cycling. This fresh fracture increases the contact\nsurface area of the electrolyte and electrode, leading to the reactive consumption of active\nLi, which in turn promotes the growth of dendritic Li, ultimately resulting in low coulombic\nefficiency (CE) during Li plating and stripping7.” (See page 2, introduction)\nWe hope these revisions meet your expectations.\n4. In Fig. 1: the legend mentions the role of the “electric field”, but this term is not clearly\ndefined in Fig. 1. Adding a brief description or annotation of the “electric field” in the\nfigure would help the readers better understand its significance in the context of the\nillustration.\nResponse: Thank you for the good suggestion. We have marked the electric field in the\nrevised schematic diagram. We hope these revisions meet your expectations and look\nforward to your continued feedback.\n----------------------------------------------------------------------------------------------------------------------\nFig. 1 | Schematic diagram of the SEI formation and Li depositon. a, Stable SEI\nexhibits intact and LiF-enriched, which is mainly derived from the de-solvation of\nsolvated Cs+. Black dotted line represents the binding of Cs+ with solvent and anion.\nRed dotted line represents the de-solvating network formed by adsorbed Cs+ under\nelectric field. b, Unstable SEI exhibits chipped with a messy rippled appearance, which\nis derived from the de-solvation of solvated Li+.\n(See Fig. 1)\n----------------------------------------------------------------------------------------------------------------------\n5. Suggestions for revision of the legends: in Fig. 2b, the legend mentions “binding\nenergies to Na+”, which should be revised to “binding energies to Li+” to match the\ncontext of this work. Additionally, the legends (a-h) of Fig. 6 appear to be incorrect,\nlikely due to a rearrangement of the figure without corresponding updates to the\nlegends. I would suggest to double-check and revise these legends to ensure\nconsistency and accuracy.\nResponse: We feel sorry for our carelessness. In our resubmitted manuscript, the typo is\nrevised. Thanks for your correction and we look forward to your continued feedback.\n6. In Fig. 5: in diagrams (e) and (f), the labels for the thickness of the SEI layer are\ninconsistent. Specifically, “SEI layer ~25nm” in Fig. 5f and “SEI layer ~ 17nm” in Fig.\n5e. Please revise them into an identical format.\nResponse: We sincerely thank the reviewer for careful reading. As suggested, we have\nharmonized the above inconsistencies into “SEI layer ~ 25nm” in Fig. 5f and “SEI layer ~\n17nm”.\nThe corresponding content has been incorporated into the revised manuscript and we look\nforward to your continued feedback.\n----------------------------------------------------------------------------------------------------------------------\nFig. 5 | e-h, Cryo-TEM images and the morphology of Li deposited on a Cu grid in\nBE+Cs (e,g) and BE (f,h) electrolytes.\n(See Fig. 5e-g)\n----------------------------------------------------------------------------------------------------------------------\n7. Electrochemical impedance data of full cell: this work presents the excellent full-cell\nperformance for the extreme anode-poor case, but it lacks data on the electrochemical\nimpedance variation of full cells. This information would be highly valuable in\nsubstantiating the claimed improvement in the interface stability by the additives. I\nwould suggest the authors consider supplementing additional work with impedance\nanalyses to provide a more comprehensive understanding of the electrochemical\nbehavior and the interface stability of their system.\nResponse: Thank you for your insightful comments and suggestions. We have\nsupplemented electrochemical impedance spectroscopy of Li|LFP full cells after different\ncycle number. As shown, the Li|LFP full cells in BE+Cs and SE+Cs electrolyte display\nsmaller electrochemical impedance than those in BE and SE electrolyte. The results\ndemonstrate that the added CsTFA could reduce the internal resistance and polarization\nduring battery cycling process, as evidenced by the reduced R and R .\nct SEI\nThe corresponding content has been incorporated into the revised manuscript.\n----------------------------------------------------------------------------------------------------------------------\nSupplementary Fig. 39 Nyquist plots of Li|LFP full cells tested (a) before cycling,\n(b) after 2 cycles’ formation, (c) after 10 cycles and (d) after 50 cycles.\n(See Supplementary Fig. 39)\n----------------------------------------------------------------------------------------------------------------------"}, {"round": 2, "reviewer_comments": {"reviewer_1": "In this round of revision, the authors have made great efforts to defend their viewpoints. However, I still have several technical concerns:\n(1) In the supplementary LSV tests, a small reduction peak is observed at approximately 1.7 V (vs. Li⁺/Li) for the BE electrolyte. Please provide an explanation for this feature.\n(2) It is recommended that the authors refer to the article “Inorganic-rich Solid Electrolyte Interphase for Advanced Lithium Metal Batteries in Carbonate Electrolytes” for calculating the reduction potentials of various species, instead of relying solely on HOMO/LUMO energy level estimations.\n(3) The binding energies calculated using the DMOL3 module may lack sufficient accuracy. For instance, the binding energy of Li⁺–TFA⁻ is reported to be –6.98 eV (672 kJ/mol), a value comparable to the lattice energy of ionic crystals. Considering that these additives are dissolved in the electrolyte, it is more appropriate to compute the ion-pair binding energies using the Gaussian software under implicit solvent conditions to better reflect the solution-phase environment.\n(4) The role of Cs⁺ in modulating the solvation structure is currently inferred indirectly through IR, 7Li NMR, and Raman spectroscopy. Could the authors consider using Cs NMR to directly probe the interaction between Cs⁺ and the solvation environment?\n(5) As I pointed out in my previous comments, the functional performance of NO3⁻ and TFA⁻ appears quite similar (see Fig. S1–2). A key contribution of this work is the proposal of a selection criterion for effective anions. Based on this criterion, could the authors provide a list of additional anion candidates with greater potential?\n(6) In the authors’ response, they state that the choice of TFA⁻ was partially based on its chemical composition and its influence on SEI formation. However, this justification remains unconvincing. For example, LiNO3 does not contain F, yet the CEs of electrolytes with LiNO3 and LiTFA are comparable (Fig. S1–2). Furthermore, the CE trends appear to correlate more consistently with the DN of the anions. Among the tested anions, some contain F or O, while others do not. Please clarify this.\n(7) Authors claim that LiTFA has high solubility. To support this viewpoint, please provide the solubility data of lithium salts with different anions in the electrolyte solvent.\n(8) The angstrom (Å) symbols in the RDF plots are inconsistent in font style compared to the rest of the text. Please revise accordingly.\n------\nComments on the responses to Reviewer #2:\nI think most of the points raised by Reviewer2 have been effectively addressed. However, I am not sure about whether the responses to Q4 and Q6 can meet R2's standard.\n1. As for Q4, the authors did not use the N/P ratio calculation methods suggested by R2, but personally I think it is not such significant.\n2. As for Q6, I cannot understand the term \"partial excess of Li inventory\" since the cell is anode free?", "reviewer_3": "The authors addressed my comments, I would recommend the acceptance."}, "author_response": "Response to referees’ comments\nDear Referees:\nThank you very much for your precious time in reviewing our manuscript (Manuscript\nnumber: NCOMMS-25-05148A) and providing quite valuable comments. We have read\nyour comments carefully and tried our best to revise manuscript accordingly. We think all\ncomments are quite helpful and important to raise the quality of the revised manuscript. All\ncomments have been point-by-point replied as below and the revised parts have been\nhighlighted in the revised manuscript with YELLOW background. We sincerely look forward\nto the reviewer's approval and expect this revised manuscript can reach the high standard\nof“Nature communication”.\nReviewer comments\nReviewer #1 (Remarks to the Author):\nIn this round of revision, the authors have made great efforts to defend their viewpoints.\nResponse: We thank the reviewer sincerely for the meticulous review that have helped us\nstrengthen our manuscript. Based on the detailed comments, we have made\nimprovements to our work. Our responses to your concerns are provided below, we hope\nthese revisions meet the reviewer’s expectations.\n(1) In the supplementary LSV tests, a small reduction peak is observed at approximately\n1.7 V (vs. Li /Li) for the BE electrolyte. Please provide an explanation for this feature.\nResponse: We thank the reviewer for careful observation. We consider the diminutive\n⁺\nreduction peak to be negligible; it likely originates from an inadvertent activation of the coin\ncell—triggered by an uncontrollable perturbation such as a sudden temperature fluctuation,\na positional shift, or minor movement during measurement on the electrochemical\nworkstation. To eliminate any potential ambiguity, we have adjusted the vertical scale of\nSupplementary Fig. 19.\nThe corresponding content has been incorporated into the revised manuscript.\n----------------------------------------------------------------------------------------------------------------------\nSupplementary Fig. 19 Linear sweep voltammetry (LSV) of Li|Cu half cells in BE+Cs\nand BE electrolyte within the voltage range from 3.0 V to 0 V at 5 mV s-1.\n(See Supplementary Fig. 19)\n----------------------------------------------------------------------------------------------------------------------\n(2) It is recommended that the authors refer to the article “Inorganic-rich Solid Electrolyte\nInterphase for Advanced Lithium Metal Batteries in Carbonate Electrolytes” for calculating\nthe reduction potentials of various species, instead of relying solely on HOMO/LUMO\nenergy level estimations.\nResponse: We thank the reviewer for this valuable suggestion about the calculation of\nreduction potentials. Guide by the literature (Angew. Chem. Int. Ed. 2021, 60, 3705-3715,\n202012005, DOI: 10.1002/anie.202012005; Nat Energy 2025, DOI: 10.1038/s41560-025-\n01842-5; Nanotechnology 2015, 26, 354003, DOI: 10.1088/0957-4484/26/35/354003), we\nhave now computed the one-electron reduction potentials of all relevant species. The\ncalculated values are 1.227 V for LiNO , 1.134 V for LiTFA, 1.075 V for LiTFSI, 0.974 V for\nDME, and 0.682 V for DOL. These results, together with the HOMO/LUMO analysis,\n₃\nconfirm that LiTFA is reduced immediately after LiNO , providing a firm thermodynamic\nrationale for the observed interfacial chemistry.\n₃\nThe corresponding data and discussion have been added to the revised manuscript.\n----------------------------------------------------------------------------------------------------------------------\nSupplementary Fig. 18 Calculated reduction potentials of optimized Li+-solvent,\nLiTFSI, LiNO and LiTFA complexes from M052X calculations within the SMD\n3\nimplicit solvation model (acetone, ε=20.49).\n(See Supplementary Fig. 18)\nIt is noted that:\nReduction potentials are performed by Gaussian 09 software package from M05-2X\ncalculations within the SMD implicit solvation model (acetone, ε = 20.49). The Li+-complex\ngeometry optimizations are performed with the pre-equilibrium structure at the B3LYP\ndensity functional and 6-31+G** basis set. The reduction potentials (E ) of complex within\nred\nthe SMD implicit solvation are calculated by the following formula:\nE = - ∆G ( ) - ∆G ( ) + ∆G ( ) - ∆G ( ) /F - 1.4\nred gas gas sol sol\nwhere ∆G (M-), ∆G (M0) are− the free en0ergies of red−uced and n0eutral complexes in gas-\ngas [gas 𝑀𝑀 𝑀𝑀 𝑀𝑀 𝑀𝑀] phase at 298.15 K. ∆G (M-), ∆G (M0) are the free energies of reduced and neutral\nsol sol\ncomplexes within the SMD implicit solvation model. F is the Faraday constant (9.6485 ×\n----------------------------------------------------------------------------------------------------------------------\n(3) The binding energies calculated using the DMOL3 module may lack sufficient accuracy.\nFor instance, the binding energy of Li –TFA is reported to be –6.98 eV (672 kJ/mol), a\nvalue comparable to the lattice energy of ionic crystals. Considering that these additives\n⁺ ⁻\nare dissolved in the electrolyte, it is more appropriate to compute the ion-pair binding\nenergies using the Gaussian software under implicit solvent conditions to better reflect the\nsolution-phase environment.\nResponse: We thank the reviewer for the meticulous query regarding the accuracy of the\nDMol³ calculations. Following the precedent set in the literature (Angew. Chem. Int. Ed.\n2021, 60, 3705–3715, DOI: 10.1002/anie.202012005; Nat. Energy 2025, DOI:\n10.1038/s41560-025-01842-5; Nanotechnology 2015, 26, 354003, DOI: 10.1088/0957-\n4484/26/35/354003), we recomputed all Li –anion binding energies with Gaussian 09\nusing the SMD continuum solvation model (acetone, ε = 20.49). The results reaffirm that\n⁺\nbinding strength increases monotonically with the donor number (DN) of the anion,\nmirroring the trend reported in Ref.(PNAS 2024, 121, e2316914121, DOI:\n10.1073/pnas.2316914121). Because this solvent-corrected benchmark reproduces the\nrelative order obtained with our original DMol³ approach, we retain the DMol³ values while\nnoting the excellent agreement. We thank the reviewer again for their careful\nconsiderations.\n----------------------------------------------------------------------------------------------------------------------\nFig. Donor number of commonly used anions and their binding energies to Li+.\nIt is noted that:\nThe calculations of binding energies are performed by Gaussian 09 software package from\nwithin the SMD implicit solvation model (acetone, ε = 20.49). The Li+-complex geometry\noptimizations are performed with the pre-equilibrium structure at the B3LYP density\nfunctional and 6-31+G** basis set. The binding energies (E ) of complex within the\nbinding\nSMD implicit solvation are calculated by the following formula:\nE = E - - E\nbinding total\nwhere E\nLi\n+ is the free energy of isolated Li+ in th\n𝐸𝐸\ne\n𝐿𝐿 𝐿𝐿\ng+asn phase, E\nn\nis the free energy of an\nisolated anion immersed in implicit solvent.\n----------------------------------------------------------------------------------------------------------------------\n(4) The role of Cs in modulating the solvation structure is currently inferred indirectly\nthrough IR, 7Li NMR, and Raman spectroscopy. Could the authors consider using Cs NMR\n⁺\nto directly probe the interaction between Cs and the solvation environment?\nResponse: We thank the reviewer for their valuable suggestion. Recognizing the need for\n⁺\ndirect evidence of CsTFA solvation and its influence on the solvation structure, we have\nrecorded ¹³³Cs NMR spectra of the electrolytes; the resulting data are presented in\nSupplementary Fig. 12.\nThe corresponding content has been incorporated into the revised manuscript.\n----------------------------------------------------------------------------------------------------------------------\nSupplementary Fig. 12 133Cs NMR spectra of (a) BE and (b) BE+Cs electrolyte.\n(See Supplementary Fig. 12)\n----------------------------------------------------------------------------------------------------------------------\n(5) As I pointed out in my previous comments, the functional performance of NO3 and\nTFA appears quite similar (see Fig. S1–2). A key contribution of this work is the proposal\n⁻\n⁻\nof a selection criterion for effective anions. Based on this criterion, could the authors\nprovide a list of additional anion candidates with greater potential?\nResponse: We thank the reviewer for this pertinent observation. While LiNO is widely used\nas an electrolyte additive, its effectiveness in today’s commercial ether- and ester-based\n₃\nformulations has reached a plateau. The persistently high N/P ratio and the unavoidable\nextra excess Li remain the primary bottlenecks for LMBs. Consequently, even state-of-the-\nart electrolytes containing LiNO still leave substantial room for performance gains. In this\nstudy, we have focused on establishing rational criteria for anion selection of electrolyte\n₃\nadditives and would like to further improve the performance of commercial electrolytes.\nEncouraged by the reviewer’s interest, we outline below a set of next-step strategies and\na shortlist of anionic candidates that merit further investigation.\n1) Based on the advantages of hetero-halide (LiCl/LiBr/LiI) interfacial layers (ACS Energy\nLett. 2024, 9, 2546–2553,DOI: 10.1021/acsenergylett.4c00548; Energy Environ. Sci.\n2022, 15, 3236–3245, DOI: 10.1039/d2ee01358d; Nat. Commun. 2023, DOI:\n10.1038/s41467-023-36934-x) and research on non-perfluorinated structure (Sci. Adv.\n2025, 11, eads4014, DOI: 10.1126/sciadv.ads4014) reported in the literature, after\ninvestigating the relevant availability, We consider to employ the heterohalogenation\nand partial defluorination of our TFA- with the shortest carbon chain. Some candidate\nanions are shown below:\n2) Guided by recent demonstrations that the effect of carbon chain length and fluorine\ncontent percentage dictate interfacial stability and ion-transport kinetics, (Nat.\nCommun. 2023, DOI: 10.1038/s41467-023-38724-x; Nat. Commun. 2024, DOI:\n10.1038/s41467-024-53807-z), some anions species with longer carbon chains and\nperfluorinated/non-perfluorinated structures are provided as follows:\n3) By substituting the carboxylate with a sulfinate motif, which was recently reported to\nendow interfaces with superior mechanical resilience and ionic conductivity (Nature\n2025, DOI: 10.1038/s41586-024-08465-y), we have designed a suite of fluorosulfinic-\nacid anions that merit further evaluation:\n4) In addition we consider the organicisation and nitridation/fluorination of the cationic\nparts:\n(6) In the authors’ response, they state that the choice of TFA was partially based on its\nchemical composition and its influence on SEI formation. However, this justification\n⁻\nremains unconvincing. For example, LiNO3 does not contain F, yet the CEs of electrolytes\nwith LiNO3 and LiTFA are comparable (Fig. S1–2). Furthermore, the CE trends appear to\ncorrelate more consistently with the DN of the anions. Among the tested anions, some\ncontain F or O, while others do not. Please clarify this.\nResponse: We thank the reviewer for their incisive remarks. Our selection of the TFA\nanion was guided by several criteria, with chemical composition and its ability to engineer\n⁻\nan optimized SEI layer being the most decisive. Figures S4,5 report the Aurbach CE tests\nperformed, while the Aurbach’s method is a short-term protocol offering a rapid snapshot\nof partial Li deposition/stripping, which could be referenced but is not sufficient to fully\ncapture long-term electrolyte performance. In fact, the cyclic Li-Cu/cathode to anode half\ncells/full cells deserves more attention. The anions compared were chosen on the basis of\ntheir DN values, not on the presence or absence of F or O. Many fluorine/oxygen atoms in\nthe structures serve primarily as stabilizing substituents rather than active SEI layer\nprecursors. Although NO is widely reported to generate Li N/Li O-rich SEI layer (Energy\nEnviron. Sci., 2024, 17, 4519–4530, DOI: 10.1039/d4ee00399c; Nat. Commun., 2024, 15,\n₃⁻ ₃ ₂\n9356, DOI: 10.1038/s41467-024-53807-z), such layers often suffer from poor interfacial\nstability and accelerated corrosion (J. Am. Chem. Soc., 2025, 147, 9365–9377, DOI:\n10.1021/jacs.4c15759). Thus, the benefit of NO is intrinsically limited by this trade-off. In\ncontrast, fluorinated additives are known to deliver LiF-rich SEI layer that markedly\n₃⁻\nenhance interfacial stability and enable prolonged cycling (Nat. Commun., 2025, 16, 3344,\nDOI: 10.1038/s41467-025-58627-3; Nat. Energy, 2024, 9, 817–827, DOI: 10.1038/s41560-\n024-01519-5). Besides, the chosen perfluorinated anionic TFA could achieve an effective\nfluorination of SEI layer in the first cycle, which facilitates: (1) a fast pull-up of cycling CE\nup to 99.36% after a fast-activated first cycle and (2) a long-term stable cycling lifetime with\na relatively high level of CE.\nWe have made a simple comparison and highlight of the long-term cycling of Li-Cu cell in\nthe following figure.\nFig. A comparison and highlight of the long-term cycling of Li-Cu cell.\n(7) Authors claim that LiTFA has high solubility. To support this viewpoint, please provide\nthe solubility data of lithium salts with different anions in the electrolyte solvent.\nResponse: We thank the reviewer for their valuable suggestion.\nThe corresponding content has been incorporated into the revised manuscript.\n----------------------------------------------------------------------------------------------------------------------\nSupplementary Table 1. Solubility tests of LiPF , LiTFSI, LiBF , LiClO , LiFSI, LiOTF,\n6 4 4\nLiNO and LiTFA in DME, DOL and BE electrolyte.\nLi salts Solubility in DME Solubility in DOL\nelectrolyte\nLiPF 0.2 M < c < 0.4 M 0.5 M < c < 1 M 0.8 M < c\nLiBF 4 M < c < 6 M 0.5 M < c < 2 M 0.8 M < c\nLiOTF 2 M < c < 4 M 2 M < c < 4 M 0.8 M < c\nLiNO 1 M < c < 2 M c < 0.2 M 0.4 M < c < 0.8 M\nSupplementary Fig. 1 Solubility tests of LiPF , LiTFSI, LiBF , LiClO , LiFSI, LiOTF,\n6 4 4\nLiNO and LiTFA in DME solvent, the undissolved Li salts are depicted by red arrow\n3\nand dashed circle.\nSupplementary Fig. 2 Solubility tests of LiPF , LiTFSI, LiBF , LiClO , LiFSI, LiOTF,\n6 4 4\nLiNO and LiTFA in DOL solvent, the undissolved Li salts are depicted by red arrow\n3\nand dashed circle.\nSupplementary Fig. 3 Solubility tests of LiPF , LiTFSI, LiBF , LiClO , LiFSI, LiOTF,\n6 4 4\nLiNO and LiTFA in BE electrolyte, the undissolved Li salts are depicted by red arrow\n3\nand dashed circle.\n(See Supplementary Fig. 1-3)\n----------------------------------------------------------------------------------------------------------------------\n(8) The angstrom (Å) symbols in the RDF plots are inconsistent in font style compared to\nthe rest of the text. Please revise accordingly.\nResponse: We thank the reviewer for careful observation. We have revised the\ncorresponding symbols in Fig. 2 and Supplementary Fig. 11.\nThe corresponding content has been incorporated into the revised manuscript.\nFig. 2 | Theoretical calculations and solvation chemistry. g,h, Radial distribution\nfunctions and coordination numbers of Li+-ODOL, Li+-ODME, Li+-OTFSI-, Li+-\nONO3- and Li+-OTFA- pairs in in BE+Cs (g) and BE (h) electrolytes.\nSupplementary Fig. 11 Linear sweep voltammetry (LSV) of Li|Cu half cells in BE+Cs\nand BE electrolyte within the voltage range from 3.0 V to 0 V at 5 mV s-1.\n(See Fig. 2 and Supplementary Fig. 11)\n----------------------------------------------------------------------------------------------------------------------\nOverall, we are deeply grateful to Reviewer #1 for the thorough review, incisive scientific\nreasoning, and constructive recommendations. We welcome continued dialogue and look\nforward to any further feedback.\n------\nComments on the responses to Reviewer #2:\nI think most of the points raised by Reviewer2 have been effectively addressed. However,\nI am not sure about whether the responses to Q4 and Q6 can meet R2's standard.\nResponse: We sincerely appreciate your positive feedback on our manuscript. Based on\nyour comments, our responses are provided below.\n1. As for Q4, the authors did not use the N/P ratio calculation methods suggested by R2,\nbut personally I think it is not such significant.\nResponse: Thank you for your comment. Actually the calculation methods of the N/P ratio\nwe used here are generic, so we have not modified it.\n2. As for Q6, I cannot understand the term \"partial excess of Li inventory\" since the cell is\nanode free?\nResponse: We apologize for the lack of detailed introduction of the term we used, actually\nthe \"partial excess of Li inventory\" is defined as the residual Li after first cycle, which could\nbe included in the initial irreversible capacity from the charging/discharging curve in the\nfirst turn.\nWe hope these explanations meet your expectations!"}]

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