Localization of micron-scale lithium polysulfide-rich phases causes anode heterogeneity in lean-electrolyte lithium–sulfur batteries

Localization of micron-scale lithium polysulfide-rich phases causes anode heterogeneity in lean-electrolyte lithium–sulfur batteries | Research Square window.SnipcartSettings = { analytics: { enabled: false } }; (function() { var accessVector = localStorage.getItem('access_vector') || ''; window.dataLayer = window.dataLayer || []; if (accessVector) { window.dataLayer.push({ user: { profile: { profileInfo: { snid: accessVector } } } }); } })(); (function(w,d,s,l,i){w[l]=w[l]||[];w[l].push({'gtm.start':new Date().getTime(),event:'gtm.js'});var f=d.getElementsByTagName(s)[0],j=d.createElement(s),dl=l!='dataLayer'?'&l='+l:'';j.async=true;j.src='https://www.googletagmanager.com/gtm.js?id='+i+dl;f.parentNode.insertBefore(j,f);})(window,document,'script','dataLayer','GTM-K279D39R'); Browse Preprints In Review Journals COVID-19 Preprints AJE Video Bytes Research Tools Research Promotion AJE Professional Editing AJE Rubriq About Preprint Platform In Review Editorial Policies Our Team Advisory Board Help Center Sign In Submit a Preprint Cite Share Download PDF Article Localization of micron-scale lithium polysulfide-rich phases causes anode heterogeneity in lean-electrolyte lithium–sulfur batteries Hee-Tak Kim, Hyeokjin Kwon, Sejin Kim, Ilju Kim, Cinthya Paulina, and 7 more This is a preprint; it has not been peer reviewed by a journal. https://doi.org/ 10.21203/rs.3.rs-7916426/v1 This work is licensed under a CC BY 4.0 License Status: Posted Version 1 posted You are reading this latest preprint version Abstract Lithium–sulfur (Li–S) batteries offer exceptionally high gravimetric energy density, yet their cycling stability remains limited by uneven Li anode reaction. Here, we reveal that the emergence of a micron-scale morphological heterogeneity in the Li anode originates from the localized formation of a micron-scale lithium polysulfide (LiPS)-rich phase in the electrolyte. Notably, the chemical composition of the LiPS-rich phase contains minimal amounts of anions such as FSI− or NO3−, which are typically responsible for forming stable solid-electrolyte interphase (SEI); this deficiency obstructs the formation of Li2O-containing SEI layer, resulting in the localized dendritic Li morphology. We further demonstrate that fragmenting the LiPS-rich phase reduces its locality, suppresses anode heterogeneity, and significantly improves cycling stability. Physical sciences/Energy science and technology/Energy storage/Batteries Physical sciences/Chemistry/Electrochemistry/Batteries Figures Figure 1 Figure 2 Figure 3 Figure 4 Figure 5 Main The growing demand for energy storage systems with higher energy densities is driving the development and adoption of next-generation batteries across various fields. Lithium–sulfur (Li–S) batteries are one of the most promising candidates for next-generation battery systems due to their high theoretical gravimetric density (>2600 Wh kg −1 ), low cost and environmental friendliness 1-4 . Achieving such high energy densities necessitates operating Li-S batteries with minimal electrolyte amounts, as electrolyte accounts for a significant portion of the cell's weight (>40 wt%) 2,5,6 . However, Li-S batteries typically show low Coulombic efficiency (CE) and short lifespan operation (commonly < 40 cycles) under lean electrolyte conditions [Electrolyte / Sulfur (E/S) ratio < 3 μL mg s −1 ] 7-9 . The limited cycling stability is primarily attributed to the dissolution of lithium polysulfides (LiPS, (Li 2 S X , X = 4, 6, 8)) 10-13 , which are generated from the electrochemical conversion of solid-state cathode active materials (i.e., S 8 or Li 2 S) and readily diffuse from the S cathode to the Li anode. On the Li anode, dissolved LiPS is reduced to Li 2 S 2 /Li 2 S, corroding the Li electrode. The issue becomes more pronounced under lean-electrolyte conditions, where LiPS accumulates within the limited volume of liquid electrolyte 14-16 . Consequently, lean electrolyte conditions exacerbate Li anode instability, resulting in markedly reduced cycling performance 16-20 . However, the specific influence of dissolved LiPS in the electrolyte on Li anode behavior has yet to be thoroughly clarified, and several fundamental questions remain unresolved. 1) Does the Li 2 S 2 /Li 2 S-containing solid electrolyte-interphase (SEI), formed via reactions between the Li anode and LiPS, inherently causes adverse effects on Li morphology? Some studies have shown that a moderate amount of LiPS can actually improve Li morphology, suggesting that corrosion of LiPS itself may not be inherently detrimental to anode performance 21,22 . 2) Under lean-electrolyte conditions, is the degradation of the Li anode at elevated LiPS concentrations fully accounted for by an increased corrosion, or are additional mechanisms at play? Attributing Li anode degradation solely to severe corrosion lacks sufficient causal evidence. For examples, certain electrolytes containing reactive species (e.g., LiFSI or LiDFOB, etc.) have been reported to promote dense Li deposition by accelerating the formation of a inorganic-rich SEI, ultimately extending battery cycle life 23,24 . More detailed understanding of the mechanisms by which LiPS influence Li anode is essential for developing strategies to overcome the short cycle life of Li-S batteries under lean-electrolyte conditions. Herein, we propose more fundamental mechanism of LiPS-induced instability of Li anode that accelerates cell degradation under lean electrolyte condition ( Fig. 1a,b ). We discovered that the clustering nature of LiPS 25,26 results in the formation of localized tens of micron-scale LiPS-rich phases in the electrolyte, inducing significant heterogeneity of Li anode. Notably, the chemical composition of the LiPS-rich phase contains minimal amounts of salt anions such as FSI − or NO 3 − , which are typically responsible for forming stable SEI. This deficiency of FSI − or NO 3 − in LiPS-rich phase is found to exert a detrimental influence on Li morphology by obstructing formation of Li 2 O-containing SEI, a species known to be found in electrolytes that induces a homogeneous Li conformation. Based on this understanding, we propose an additive strategy to reduce the localized nature of PS-rich phases by fragmenting them using various salt anions that can control the degree of LiPS-rich phase locality. With the designed electrolyte, our assembled prototype 0.7 Ah stacked pouch cell could deliver higher cycling stability of 134 cycles under lean electrolyte condition of E/S ratio 2.3 μL mg s −1 (E/C ratio = 1.9 g Ah −1 ), compared to 46 cycles at reference electrolyte. This study expands understanding of LiPS-containing catholytes, proposing that degradation of Li anodes is due to local phase formation and subsequent SEI changes by LiPS, rather than the prevailing view that degradation is due to the intrinsic corrosiveness of LiPS. Sharp decrease in lifespan of Li-S batteries under lean electrolyte conditions Given that high LiPS concentrations in the electrolyte are often considered crucial under lean conditions, LiPS are likely to exert a critical influence on the Li. Optical microscopy (OM) images of Li electrodes collected from two Li–S pouch cells ( Supplementary Fig. 1 ) with different E/S ratios reveal the formation of micrometer-scale domains with LiPS residues with a distinct spatial distribution across the electrode surface ( Fig. 2a ). At an E/S ratio of 5.0 μL mg s − 1 , only a few LiPS deposits are observed on the Li electrode. In contrast, at an E/S ratio of 2.3 μL mg s − 1 , LiPS precipitation increases markedly, spreading extensively across the electrode. The characteristic sizes of these LiPS-rich phases were determined to be D 50 = 7 μm and D 90 = 22 μm ( Supplementary Fig. 2 ). These results highlight that LiPS accumulates on the Li electrode at the micrometer scale, forming localized phases. It has been reported that Li–S batteries exhibit a significant decline in cycling stability under lean electrolyte conditions 9 . To investigate this behavior, we assembled a prototype 0.7 Ah pouch cells ( Fig. 2b and Supplementary Fig. 3 ) and evaluated their cycling performance. As shown in Fig. 2c , the pouch cells sustained over 200 cycles before reaching 70% capacity retention at an E/S ratio of 5.0 and 3.5 μL mg s − 1 , which provides a sufficient amount of electrolyte for cell operation. In contrast, under a leaner condition with an E/S ratio of 2.3 μL mg s − 1 , only 46 cycles were achieved, reflecting significant degradation in cycling stability. Voltage profiles during cycling further reveal that the electrolyte volume strongly affects cell performance, as evidenced by a pronounced increase in overpotential over successive cycles at an E/S ratio of 2.3 μL mg s − 1 , in contrast to the more stable profiles observed at E/S ratios of 5.0 and 3.5 μL mg s − 1 ( Supplementary Fig. 4 ). The Li anodes collected from a Li–S pouch cell at an E/S ratio of 2.3 μL mg s − 1 after 46 cycle operation exhibited uneven morphology, implying severe degradation (Fig. 2d). Scanning electron microscopy (SEM) images of the Li electrodes after 20 cycles from the pouch cells show that electrolyte deficiency leads to more irregular Li deposition. ( Fig. 2e ). Notably, the morphological evolution at 2.3 μL mg s − 1 features distinct structural characteristics, where the deposited Li exhibits heterogeneous surface morphologies comprising both dendritic and granular domains. This heterogeneity was observed across tens of micrometers, indicating substantial reaction heterogeneity. Energy dispersive X-ray spectroscopy (EDS) analysis was conducted to examine the spatial distribution of the chemical composition of the SEI layer. At an E/S ratio of 5 μL mg s − 1 —where stable cycling was observed—S atoms, mainly originating from LiPS, showed relatively low locality. At an E/S ratio of 3.5 μL mg s − 1 , where porous Li structures began to appear, S distribution became more localized. Notably, at an E/S ratio of 2.3 μL mg s − 1 —where a sharp decline in cycling performance was observed—the spatial distribution of S within the Li electrode exhibited strong localization. The size of the localized regions progressively increases with decreasing E/S ratio. To sum up, observations indicate that under lean electrolyte conditions, micron-scale LiPS-rich phases accumulate locally on the Li electrode, while heterogeneous Li structures also emerge at comparable length scales, which we identify as a primary cause of Li anode degradation in Li–S full cells. Taken together, these findings lead us to hypothesize that the formation of localized LiPS-rich phases is a key factor driving Li heterogeneity and, ultimately, premature cell failure, as will be discussed in the following section. Micron-scale LiPS-rich phases in the catholyte Li 2 S 4 was selected as the representative LiPS species for catholyte modeling, as it predominates at the maximum LiPS concentration during discharge in Li–S full cells 27-29 . Catholytes containing 0.2, 0.5, 0.8, and 1.0 M Li 2 S 4 were prepared in the reference electrolyte (FN, 0.75 M LiFSI in 1, 2-dimethoxyethane (DME): 2-methylfuran (2MF) 4:1 in volume with 5 wt% LiNO 3 ). The 1.0 M Li 2 S 4 concentration corresponds to the maximum Li 2 S 4 level expected at an E/S ratio of 2.3 μL mg s –1 ( Fig. 3a , Supplementary Note 1 and Supplementary Fig. 5 ), while the lower concentrations represent more flooded-electrolyte conditions. The catholytes are opaque as shown in insets in Fig. 3b , implying their inhomogeneous structure. OM images of the catholytes reveal the formation of micron-scale, localized LiPS-rich phases, which grow in size with increasing the LiPS concentration ( Fig. 3b,c ). While the diameters of LiPS-rich phases in FN + 0.2 M Li 2 S 4 were D 50 =4 μm and D 90 =5 μm, increasing the Li 2 S 4 concentration to 0.5 M, 0.8 M, and 1.0 M resulted in size increases to D 50 = 4 μm, 6 μm, 9 μm, and D 90 = 9 μm, 12 μm, and 21 μm, respectively ( Fig. 3c ). In contrast, no visible features were observed in the LiPS-free FN electrolyte. These results indicate that, under lean electrolyte conditions where LiPS concentrations are high, LiPSs significantly aggregate, leading to the formation of more discrete and localized phases. Importantly, the OM observation implies that the increasingly aggregated state of LiPS with LiPS concentration observed in 7 Li nuclear magnetic resonance (NMR) spectroscopy analysis ( Supplementary Fig. 6 ) and MD simulations ( Supplementary Fig. 7 ) not only reflects changes in coordination structures at the molecular scale, but also leads to the formation of micron-scale inhomogeneous LiPS-rich phases. Given that more uneven Li morphologies were observed at higher LiPS concentrations in Fig. 2e , we hypothesized that the emergence of heterogeneity of Li anode is closely associated with localized LiPS-rich phases at electrolyte. As shown in Fig. 3d, Li|Ni half cells employing catholytes with different Li 2 S 4 concentrations exhibited distinct Li morphologies after cycling. Both the FN electrolyte and the FN + 0.2 M Li 2 S 4 catholyte promoted dense and uniform Li structure, whereas FN electrolyte with 0.5 M Li 2 S 4 led to heterogeneous Li morphology, featuring the coexistence of granular and porous Li domains. Such uneven Li morphology closely resembles that observed at the Li metal anode in the Li–S full cell. In particular, we observed a significant difference in elemental distribution between the two Li regions; as shown in Fig. 3e , the distribution of the S element is noticeably concentrated in the dendritic region compared to the granular region. This suggests the possibility that the formation of dendritic regions is induced by the tens of micron-scale LiPS-rich phases. In a later section, we explore how SEI properties correlate with Li morphology in S-concentrated regions. Dendritic and porous Li structure, with its high surface area, generally exacerbates side reactions with the electrolyte and leads to lower CE. Inclusion of more LiPS in FN electrolyte resulted in lower initial CE ( Supplementary Fig. 8 ) and subsequent average CE measured by modified Aurbach protocol ( Fig. 3f ). While the FN electrolyte achieved an average CE of 99.40%, the addition of 0.2 M, 0.5 M, 0.8 M, and 1.0 M LiPS decreased it to 99.12 %, 98.96 %, 98.37 %, and 97.98 %, reflecting the adverse influence of LiPS-induced morphological heterogeneity in Li. Characterization of SEI layer induced by LiPS-rich phases The reason behind the formation of dendritic Li induced by the LiPS-rich phase remains to be clarified. The SEI characteristics of Li deposits formed in FN + 1.0 M Li 2 S 4 electrolyte were investigated using cryogenic transmission electron microscopy (cryo-TEM) ( Fig. 4a-c ). At the granular Li region, the SEI layer exhibited crystalline features composed of Li 2 O crystallites, as further confirmed by the selected area electron diffraction (SAED) pattern. The formation of Li 2 O component in SEI is mainly attributed to the decomposition of FSI − or NO 3 − anions, which has been reported to contribute to the development of stable SEI layers and granular Li morphology 30,31 . In contrast, the dendritic Li structure exhibited only weak crystalline signals. A similar lack of crystalline features was also observed on Li dendrites formed in the FN + 0.5 M Li 2 S 4 ( Supplementary Fig. 9 ). Notably, clear crystalline signals were detected when using FN electrolyte without LiPS species ( Supplementary Fig.10 ). When comparing the (fluorine + nitrogen)/sulfur ratio from EDS analysis—which reflects the relative contributions of LiFSI and LiNO 3 versus LiPS to SEI formation—a gradual decrease is observed with increasing Li 2 S 4 concentration in the catholyte ( Fig. 4d and Supplementary Fig. 11 ). This suggests that SEI layers formed by LiPS-rich phases contain a lower proportion of LiFSI and LiNO 3 -derived components, which have been associated with granular Li morphology 32,33 . To further clarify the chemical environment of LiPS-rich phase, Raman spectroscopy was conducted on FN and FN + 1.0 M Li 2 S 4 electrolytes ( Fig. 4e ). The spectrum of the LiPS-rich phase shows strong S X 2− and S 4 2− signals at 400–500 cm −1 , along with a weak DME signal at ~850 cm –1 27,34,35 , indicating the presence of only a small amount of solvent molecules. Notably, signals corresponding to FSI − and NO 3 − , which are present in the FN electrolyte, are barely detectable 36,37 . These observations collectively account for the markedly suppressed Li 2 O formation in contact with LiPS-rich phases. Molecular dynamics (MD) simulations also exhibited the localized distribution of LiPS within the electrolyte. At higher concentrations, the S 4 2 − anion exhibited enhanced coordination with Li⁺, consistent with the clustering tendency of LiPS ( Supplementary Fig. 7 and 12–14 ). Notably, these LiPS-containing clusters hardly contain other anions (i.e., FSI⁻ and NO 3 − ). The reduced presence of FSI⁻ and NO 3 − in these LiPS-containing clusters limits their decomposition, thereby suppressing the formation of LiFSI and LiNO 3 -derived SEI layers that are essential for effective passivation of the Li anode. We further investigated the behavior of other soluble LiPS species, specifically Li 2 S 8 and Li 2 S 6 ( Supplementary Fig. 15-16) . The sulfur concentration in LiPS species was adjusted to ~14 wt% in the electrolyte (i.e., 0.5 M Li 2 S 8 and 0.67 M Li 2 S 6 ) to reflect their actual levels in Li–S full cells 38 . Although longer-chain LiPS have higher redox potentials—implying greater corrosivity toward the Li anode—their impact on Li reversibility and morphology was not more detrimental. In fact, more uniform Li deposition and higher CE compared to Li 2 S 4 were observed. MD simulations showed that these species exhibit a weaker tendency to form discrete phase domains due to reduced clustering propensities ( Supplementary Fig. 17-18 ) which was also observed from previous studies 39 . OM observation for bulk catholytes containing Li 2 S 8 and Li 2 S 6 displayed smaller LiPS-phases indicated by smaller D 50 and D 90 values ( Supplementary Fig. 19 ). These findings suggest that the spatial localization of LiPS, rather than their inherent corrosivity, plays a more critical role in determining Li anode stability. Fragmenting LiPS-rich phases to decrease locality Following the identification of the detrimental effects of localized LiPS-rich phases on the Li electrode, LiI was selected to examine its ability to suppress such localization. I⁻, known for its high donor number (DN) (29 kcal mol −1 ) and high Li + binding energy ( Supplementary Fig. 20) 40 , is expected to inhibit LiPS clustering by competitively coordinating with Li⁺, thereby suppressing the formation of large LiPS-rich phases. To verify this effect, 0.1 M LiI was added to FN electrolyte, denoted as FNI electrolyte. For comparison, 0.1M LiFSI (DN of 9.5 kcal mol −1 ) and LiNO₃ (DN of 19 kcal mol −1 ), which are baseline salts, were also introduced, yielding FNF and FNN electrolytes, respectively. These electrolytes were designed to evaluate how anion properties influence the extent of LiPS-rich phase formation. The downfield shifts of 7 Li NMR signal upon incorporation of 1.0 M Li 2 S 4 were compared among the electrolytes ( Fig. 5a, Supplementary Fig. 21 ). These shifts originate from the coordination of PS anions with Li + as LiPS-rich phases evolve. FNF exhibited a shift of 0.18 ppm, nearly identical to FN (0.19 ppm), indicating that the addition of 0.1M LiFSI had little effect on Li + PS interactions. In FNN, the shift was slightly smaller at 0.16 ppm, reflecting a modest reduction in Li + PS coordination due to the 0.1M NO 3 − anion. Notably, FNI showed a substantially smaller shift of 0.09 ppm, demonstrating that 0.1M LiI effectively suppresses Li + PS coordination, thereby mitigating the development of LiPS-rich phases. Such changes in the local environment of Li⁺ are, as previously investigated, associated with the micron-scale clustering of LiPS. To investigate the regulation of LiPS-rich phases by additional anions, the OM images of the 1.0 M Li 2 S 4 -containing electrolytes were captured, and the size distributions of LiPS-rich phases were determined ( Fig. 5b,c ). FNF and FNN had limited impact on regulating LiPS-rich phases, consistent with the NMR findings. Compared to FN + 1.0 M Li 2 S 4 (D 50 = 9 μm and D 90 = 21 μm, Fig. 3c), the D 50 /D 90 values slightly decreased to D 50 = 7 μm and D 90 = 19 μm for FNF + 1.0 M Li 2 S 4 , and D 50 = 8 μm and D 90 = 17 μm for FNN + 1.0 M Li 2 S 4 . In contrast, FNI further reduced the size of LiPS-rich phases, lowering the D 50 and D 90 values to 7 and 13 μm, respectively—approximately 62% of those measured in FN electrolyte. This result indicates that the disruption of LiPS clustering with I − leads to fragmentation of LiPS-rich phases, thereby decreasing locality. The correlation between the locality of LiPS-rich phases and the uniformity in Li morphology was confirmed through the SEM images of the cycled Li electrodes ( Fig. 5b ). FNF and FNN showed concurrent evolution of porous and granular Li morphologies, with the porous regions corresponding to areas with high S content in the EDS images, indicating the formation of localized S-rich SEI. In contrast, FNI exhibited smooth and uniform Li morphology and diminished localization of S content, indicating a more homogeneous SEI layer. These findings demonstrate that mitigating the localization of LiPS-rich phases within the electrolytes, particularly via the addition of I - , promotes uniform Li deposition. For the 1.0 M Li 2 S 4 -containing electrolytes, the CE of the Li electrode was measured using the modified Aurbach method. Despite a decrease in ionic conductivity ( Supplementary Fig. 22 ), the CE increased with the DN of the added salt anion ( Fig. 5d and Supplementary Fig. 23 ); FNF (98.07%), FNN (98.30%), and FNI (98.58%) exhibited an increase in CE compared to FN (97.98%). The higher Li reversibility in the LiPS-containing FNI emphasizes the importance of suppressing the localization. Under LiPS-free conditions, the average CE of FNI (99.37%) was nearly identical with that of FN (99.40%, Fig. 3f ), indicating that the improved Li reversibility is associated with the reduced locality of LiPS-rich phases ( Supplementary Fig. 24 ). We assembled Li-S full cells to demonstrate the efficacy of the approaches in enhancing full cell cycling stability ( Fig. 5e,f ). The FN cell showed rapid capacity fading due to Li instability caused by LiPS-rich phases, resulting in only 46 cycles of operation. FNF and FNN, which moderately improved Li reversibility and Li deposition morphology, showed extended cycling performances (58 and 84 cycles for FNF and FNN, respectively), with FNN outperforming FNF due to the higher DN of NO 3 − compared to FSI − . In contrast, FNI, which markedly reduced the locality of LiPS-rich phases, delivered superior cycling stability, retaining 70% of their capacity after more than 134 cycles. Notably, the observed improvement is not attributable to the iodine redox reaction, commonly used to enhance redox kinetics, as the operating voltage window lies outside the redox potential of LiI (oxidation potential of ~3.0 V vs. Li/Li + ) 41,42 . We analyzed the Li electrodes after 20 cycles using SEM and EDS ( Supplementary Fig. 25) . Pronounced heterogeneity in Li morphology, indicative of the locality of LiPS-rich phases, was observed for FNF and FNN, whereas FNI exhibited a much more uniform Li surface. The spatial distribution of S signals overlapped well with the porous Li regions, indicating that S-rich SEI preferentially formed on the locally porous Li structures. Suppressing the locality of LiPS-rich phases thus facilitates more uniform Li morphology, leading to improved cycling stability. A prototype pouch cell was assembled using the FNI electrolyte at a lower E/S ratio of 1.5 μL mg s −1 (E/C ratio = 1.3 g Ah −1 ) to validate the cycling stability improvements achieved by mitigating the locality of LiPS-rich phases. ( Supplementary Fig. 26-27 and Table 1 ). Due to very low dosage of electrolyte, our prototype pouch cell achieved an energy density of 524 Wh kg −1 (including tabs and packages, 324 Wh kg −1 ). Although Li-S full cells typically operate for only a few cycles under such a low E/S ratio due to instability of Li electrode 7 , our cell demonstrated stable operation for 45 cycles, underscoring the importance of controlling locality of LiPS-rich phases to enhance cycling stability under lean electrolyte conditions. Conclusions In this study, we investigated the effects of LiPS on the Li anode, which are determined by micron-scale locality of LiPS-rich phases. Due to the clustering nature of LiPS, LiPS-rich phases form under lean electrolyte conditions. Sporadically distributed micron scale LiPS-rich phases, characterized by low FSI − and NO 3 − concentrations, undermine anode stability by hindering the formation of LiFSI or LiNO 3 -derived SEI layer and facilitating non-uniform Li growth, thereby contributing to unstable cell performance. We demonstrated that decreasing locality of LiPS-rich phases can significantly reduce anode heterogeneity, as evidenced by the improved cycling stability. We conclude that the locality of LiPS-rich phases is the root cause of instability in lean-electrolyte Li-S batteries. The insights gained from this study will provide a foundation for developing engineering strategies to address the issue and pave the way for the implementation of high energy density, yet stable, Li-S battery systems. Methods Materials preparation. 1,2-Dimethoxyethane (DME) (anhydrous, 99.5%), 2-methylfuran (2MF) (99%), lithium nitrate (LiNO 3 ) (99.9%), and lithium iodide (LiI) (99.9%) were purchased from Sigma Aldrich. Lithium bis(fluorosulfonyl)imide (LiFSI) (99.9%) was obtained from Nippon shokubai. Lithium foils with thicknesses of 40 µm, 150 µm, and 450 µm, as well as double-sided Li-Cu-Li (50 µm), were sourced from Honjo Metal. Preparation of electrolytes. The reference electrolyte (FN) was formulated by dissolving 0.75 M LiFSI in a solvent mixture of 1,2-dimethoxyethane (DME) and 2-methylfuran (2-MF) (volume ratio 4:1), with the addition of 5 wt% LiNO 3 as an additive. To prepare specific electrolytes, 0.1 M of additional salts were introduced into the FN electrolyte. Various lithium polysulfide-containing catholytes were synthesized by stoichiometrically combining Li 2 S and S 8 (Sigma Aldrich). All electrolyte preparations were conducted in an argon-filled glovebox, maintaining H 2 O and O 2 levels below 1 ppm. Characterizations of electrodes and electrolytes. 7 Li NMR measurements were performed using an AS400 spectrometer. To eliminate the influence of DMSO-d6 on the measurements, coaxial NMR tubes were employed. DMSO-d6 was contained in the inner tube, while the specific electrolytes were loaded into the outer tube. SEM imaging and EDS mapping were conducted using a field-emission scanning electron microscope (JEOL, JSM-IT800). Li electrodes from Li|Ni half cells were collected for characterization under conditions analogous to Li stripping and plating in Li–S full cells. Cryo-TEM analysis was carried out using a Glacios microscope (Thermo Fisher) to minimize atmospheric reactions and electron beam damage. Lithium was electrochemically deposited onto a TEM grid [200-mesh Nickel (Ni) with a lacey carbon film] at a current density of 1.0 mA cm − 2 and a capacity of 1.0 mAh cm − 2 . Ni grids were chosen instead of copper to prevent reactions between lithium polysulfides and the current collector 43 . After preparation, the sample was immediately placed in liquid nitrogen following transfer into a microtube. Raman spectroscopy was performed using a PL spectrometer (Renishaw, inVia Qontor) with a 324 nm laser. Electrolyte samples were deposited onto a glass slide and covered with a quartz cover slip. The apertures were then sealed with epoxy resin to prevent unwanted reactions between the atmosphere and lithium polysulfides. Optical microscopy. Optical microscopy was conducted using optic microscope (Novel optics). To get images of various bulk electrolyte and catholytes, 10 µL bulk electrolyte was dropped on slide glass, then covered with cover-glass. The apertures were tightly sealed with epoxy resins. Computational calculations. DFT and MD calculations were performed using the Materials Studio package. The DMol³ module was employed for DFT calculations, utilizing the Generalized Gradient Approximation (GGA) with the Perdew-Burke-Ernzerhof (PBE) gradient-corrected exchange-correlation functional. Geometric optimizations were conducted on all molecules (DME, 2MF, FSI − , NO 3 − , S 4 2 − , S 6 2 − , and S 8 2 − ) prior to calculations, with the energy convergence criterion set to below 10-6 eV. Molecular dynamics (MD) simulations were carried out using the Forcite module within the Materials Studio package. Periodic simulation cells were equilibrated through 1 ns of NPT (constant number of particles, pressure, and temperature) followed by NVT (constant number of particles, volume, and temperature) ensemble simulations. The Nosé algorithm, with a Q ratio of 0.1, was used for thermostatting. After the 1 ns NVT equilibration, an additional 5 ns of NVT simulation was conducted to achieve full equilibrium. Snapshots of polysulfides (PS) and the overall coordination numbers were extracted from the final equilibrated systems. Electrochemical performance measurements and characterizations. Electrochemical measurements were performed using 2032-type coin cells and pouch cells. For coin cells, 150 µm thick Li metal was used as the anode, with 20 µL of various electrolytes for general testing. Polypropylene (PP) separators (Celgard) were used for all cell configurations. To measure CE , the modified Aurbach method was applied following established literature protocols: 1) Pre-formation cycles involved plating and stripping 5 mAh cm − 2 of Li at 0.5 mA cm − 2 . 2) Identical areal capacities of Li were deposited on Ni current collectors. 3) Ten cycles of Li stripping/plating were conducted at 1 mAh cm − 2 and 0.5 mA cm − 2 . 4) A final full Li stripping was performed to 1.0 V to remove the remaining Li reservoir. CE from pre-formation cycles was excluded to eliminate the effects of galvanic corrosion and Li consumption during SEI formation. Two PP separators were used to avoid shot-circuit during lithiation and de-lithiation. For observing Li 2 S deposition at Carbon paper (Wizmac), E/S ratio = 3 μL mg s − 1 was targeted. Due to high dead-volume of coin cell configurations, slightly higher E/S ratio was intentionally used. The loading of the S on cathode was 3mg s cm -2 . Its fabrication was referred to our previous study 44 . Pouch cells were assembled in a dry room (dew point < −40 °C). The cathode, anode, and separator dimensions were 3.3 × 5 cm 2 , 3.5 × 5.2 cm 2 , and 3.7 × 5.3 cm 2 , respectively. The electrolyte was dosed in accordance with the specified E/S ratio. Three double side-coated cathodes (=3-layer stacked pouch cells, 0.27Ah pouch cell) were used for various characterization and assessing rate performances, while 8-layer stacked pouch cells (8 Al current collectors with double-sided cathodes, 0.7Ah pouch cell) were employed for cycling tests. The areal sulfur loading of the cathodes was 2.17 – 2.25 mg s cm -2 . The ratio of sulfur : carbon : binder was 72 : 27 : 1. To minimize contact resistance, pouch cells were compressed between two metal plates at a pressure of 1.5–2.0 bar and operated within a voltage range of 1.80–2.50 V. All cells were rested for 24 hours to achieve full electrolyte wetting. Before the targeted cycling rate, pre-conditioning cycles were conducted at 0.05C (1C = 1675 mAh g s −1 ) and 0.1C two times respectively. For the pouch cell fabricated at E/S ratio = 1.5 µL mg −1 , pre-cycle was performed at 0.025 C and the actual cycling test was conducted at 0.05 C. Galvanostatic cycling tests were performed at 25 °C using a WBCS3000L battery tester. Declarations Data availability All data supporting the findings of this study are available within the paper and its Supplementary Information. Acknowledgements This work was supported by the National Foundation of Korea (NRF) grant funded by the Korea government (MSIT) (RS-2023-00261543) and (RS-2005-25441256) and LG Energy Solution-KAIST Frontier Research Laboratory (2024). We acknowledge the support of LG Energy Solution in the development of the sulfur cathode electrode. Author contributions S. Kim proposed the research and conceived the idea. H. Kwon and H.-T. K. supervised the project. S. Kim performed the electrochemical measurements. S. Kim, I. Kim, and C. Paulina carried out and analyzed the cryo-TEM and SEM experiments. S. Kim, Y. Shin, and D. Kim conducted the DFT and MD simulations. S. Chung and C. Park assisted with pouch cell fabrication. J. Han, J. Kim and B. Cho contributed to optical microscopy analysis. S. Kim, H. Kwon, and H.-T. K. co-wrote and revised the manuscript. Competing interests The authors declare no competing interests. References Manthiram, A., Fu, Y., Chung, S.-H., Zu, C. & Su, Y.-S. Rechargeable lithium–sulfur batteries. Chemical reviews 114 , 11751-11787 (2014). Zhou, G., Chen, H. & Cui, Y. Formulating energy density for designing practical lithium–sulfur batteries. Nature Energy 7 , 312-319 (2022). Manthiram, A., Chung, S. H. & Zu, C. Lithium–sulfur batteries: progress and prospects. Advanced materials 27 , 1980-2006 (2015). Li, Z. et al. Lithiated metallic molybdenum disulfide nanosheets for high-performance lithium–sulfur batteries. 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Recent Progress for Concurrent Realization of Shuttle‐Inhibition and Dendrite‐Free Lithium–Sulfur Batteries. Advanced Materials 35 , 2212116 (2023). Ren, W., Ma, W., Zhang, S. & Tang, B. Recent advances in shuttle effect inhibition for lithium sulfur batteries. Energy Storage Materials 23 , 707-732 (2019). Wild, M. et al. Lithium sulfur batteries, a mechanistic review. Energy & Environmental Science 8 , 3477-3494 (2015). Shi, H. et al. Challenges and solutions for lithium–sulfur batteries with lean electrolyte. Advanced Functional Materials 33 , 2306933 (2023). Guo, J. et al. Shelf life of lithium–sulfur batteries under lean electrolytes: status and challenges. Energy & Environmental Science 17 , 1695-1724 (2024). Liu, Y. et al. Electrolyte solutions design for lithium-sulfur batteries. Joule 5 , 2323-2364 (2021). Chen, Z.-X. et al. Failure analysis of high-energy-density lithium‒sulfur pouch cells. Energy Storage Materials 53 , 315-321 (2022). Bi, C.-X. et al. Galvanic Corrosion of Lithium Metal Anodes in Lithium–Sulfur Batteries. Journal of the American Chemical Society (2025). Kim, S. C. et al. Solvation-property relationship of lithium-sulphur battery electrolytes. Nature Communications 15 , 1268 (2024). Gao, X. et al. Electrolytes with moderate lithium polysulfide solubility for high-performance long-calendar-life lithium–sulfur batteries. Proceedings of the National Academy of Sciences 120 , e2301260120 (2023). Li, W. et al. The synergetic effect of lithium polysulfide and lithium nitrate to prevent lithium dendrite growth. Nature communications 6 , 7436 (2015). Yan, C. et al. Lithium metal protection through in-situ formed solid electrolyte interphase in lithium-sulfur batteries: The role of polysulfides on lithium anode. Journal of Power Sources 327 , 212-220 (2016). Jie, Y. et al. Towards long-life 500 Wh kg− 1 lithium metal pouch cells via compact ion-pair aggregate electrolytes. Nature Energy 9 , 987-998 (2024). Boyle, D. T. et al. Corrosion of lithium metal anodes during calendar ageing and its microscopic origins. Nature Energy 6 , 487-494 (2021). Gupta, A., Bhargav, A., Jones, J.-P., Bugga, R. V. & Manthiram, A. Influence of lithium polysulfide clustering on the kinetics of electrochemical conversion in lithium–sulfur batteries. Chemistry of Materials 32 , 2070-2077 (2020). Gupta, A. & Manthiram, A. Unifying the clustering kinetics of lithium polysulfides with the nucleation behavior of Li 2 S in lithium–sulfur batteries. Journal of Materials Chemistry A 9 , 13242-13251 (2021). Lang, S., Yu, S.-H., Feng, X., Krumov, M. R. & Abruña, H. D. Understanding the lithium–sulfur battery redox reactions via operando confocal Raman microscopy. Nature communications 13 , 4811 (2022). Yu, X. & Manthiram, A. A class of polysulfide catholytes for lithium–sulfur batteries: energy density, cyclability, and voltage enhancement. Physical Chemistry Chemical Physics 17 , 2127-2136 (2015). Xing, C. et al. Regulating liquid and solid-state electrolytes for solid-phase conversion in Li–S batteries. Chem 8 , 1201-1230 (2022). Zhao, Y. et al. Electrolyte engineering for highly inorganic solid electrolyte interphase in high-performance lithium metal batteries. Chem 9 , 682-697 (2023). Zhen, C. et al. Revealing Lithium Nitrate-Mediated Solid-Electrolyte Interphase of Lithium Metal Anode via Cryogenic Transmission Electron Microscopy. Nano Letters 24 , 6714-6721 (2024). Ko, S. et al. Electrode potential influences the reversibility of lithium-metal anodes. Nature Energy 7 , 1217-1224 (2022). Zhang, W. et al. Single-phase local-high-concentration solid polymer electrolytes for lithium-metal batteries. Nature Energy 9 , 386-400 (2024). Chu, H. et al. Unraveling the dual functionality of high‐donor‐number anion in lean‐electrolyte lithium‐sulfur batteries. Advanced Energy Materials 10 , 2000493 (2020). McBrayer, J. D., Beechem, T. E., Perdue, B. R., Apblett, C. A. & Garzon, F. H. Polysulfide speciation in the bulk electrolyte of a lithium sulfur battery. Journal of The Electrochemical Society 165 , A876 (2018). Kwon, H. et al. Weakly coordinated Li ion in single-ion-conductor-based composite enabling low electrolyte content Li-metal batteries. Nature Communications 14 , 4047 (2023). Amirov, A. et al. Effect of lithium perchlorate addition on LiNO3–KNO3 nitrate eutectic. Ionics 30 , 6089-6096 (2024). Song, Y.-W. et al. Phase equilibrium thermodynamics of lithium–sulfur batteries. Nature Chemical Engineering 1 , 588-596 (2024). Andersen, A. et al. Structure and dynamics of polysulfide clusters in a nonaqueous solvent mixture of 1, 3-dioxolane and 1, 2-dimethoxyethane. Chemistry of Materials 31 , 2308-2319 (2019). Linert, W., Jameson, R. F. & Taha, A. Donor numbers of anions in solution: the use of solvatochromic Lewis acid–base indicators. Journal of the Chemical Society, Dalton Transactions , 3181-3186 (1993). Ren, Y., Zhao, T., Liu, M., Zeng, Y. & Jiang, H. A self-cleaning Li-S battery enabled by a bifunctional redox mediator. Journal of Power Sources 361 , 203-210 (2017). Wu, F. et al. Lithium iodide as a promising electrolyte additive for lithium-sulfur batteries: mechanisms of performance enhancement. Advanced Materials (Deerfield Beach, Fla.) 27 , 101-108 (2014). Nanda, S. & Manthiram, A. Delineating the lithium–electrolyte interfacial chemistry and the dynamics of lithium deposition in lithium–sulfur batteries. Advanced Energy Materials 11 , 2003293 (2021). Chu, H. et al. Achieving three-dimensional lithium sulfide growth in lithium-sulfur batteries using high-donor-number anions. Nature communications 10 , 188 (2019). Additional Declarations There is NO Competing Interest. Supplementary Files Supplementarymaterial.docx Article File Cite Share Download PDF Status: Posted Version 1 posted You are reading this latest preprint version Research Square lets you share your work early, gain feedback from the community, and start making changes to your manuscript prior to peer review in a journal. As a division of Research Square Company, we’re committed to making research communication faster, fairer, and more useful. We do this by developing innovative software and high quality services for the global research community. Our growing team is made up of researchers and industry professionals working together to solve the most critical problems facing scientific publishing. Also discoverable on Platform About Our Team In Review Editorial Policies Advisory Board Help Center Resources Author Services Accessibility API Access RSS feed Manage Cookie Preferences © Research Square 2026 | ISSN 2693-5015 (online) Privacy Policy Terms of Service Do Not Sell My Personal Information {"props":{"pageProps":{"initialData":{"identity":"rs-7916426","acceptedTermsAndConditions":true,"allowDirectSubmit":true,"archivedVersions":[],"articleType":"Article","associatedPublications":[],"authors":[{"id":538242589,"identity":"7748efc7-7820-4c5e-b33c-74e8b398eaac","order_by":0,"name":"Hee-Tak Kim","email":"data:image/png;base64,iVBORw0KGgoAAAANSUhEUgAAAZAAAAAyAQMAAABI0h/eAAAABlBMVEX///8AAABVwtN+AAAACXBIWXMAAA7EAAAOxAGVKw4bAAAA9UlEQVRIie3SMQrCMBSA4RcKcQl2k4ge4pWAIngZKXSq4AmKIjj1AIqDVyi4OlQCuvQA3awITgUVF4cOtkVxMgouDvmHktfy0TcEQKf7w7B8hoAUDPkcgQzVJP9eEmoXI/+eADBRvPhM2pXxNrlFIKq16Iq3gedBRSZktnpPOv6mN/JjaNFGf9nzUXJgDpLgqFgsdq0hO0O3ICHDMF/MBZKECrI7XUZZSdzDOkOPg5l+IDEjY1Yu5ho2Q4MDz/8SqEjkWPNmxAVtOEI0UdYn/IjrqYpsZXJJN11rMbcP9TTzTNO093tfQR7x15FCcRt0Op1O91N3w0pQPjCslnwAAAAASUVORK5CYII=","orcid":"https://orcid.org/0000-0003-4578-5422","institution":"Korea Advanced Institute of Science and Technology","correspondingAuthor":true,"prefix":"","firstName":"Hee-Tak","middleName":"","lastName":"Kim","suffix":""},{"id":538242590,"identity":"9fd492b4-e1f0-4de4-908f-4ae048612479","order_by":1,"name":"Hyeokjin Kwon","email":"","orcid":"https://orcid.org/0000-0001-9468-034X","institution":"Korea Advanced Institute of Science and Technology","correspondingAuthor":false,"prefix":"","firstName":"Hyeokjin","middleName":"","lastName":"Kwon","suffix":""},{"id":538242591,"identity":"afcdf25f-ea6f-4c2f-921d-27a82393938d","order_by":2,"name":"Sejin Kim","email":"","orcid":"","institution":"Korea Advanced Institute of Science and Technology","correspondingAuthor":false,"prefix":"","firstName":"Sejin","middleName":"","lastName":"Kim","suffix":""},{"id":538242592,"identity":"bf60d7fe-2c54-4bac-8e1a-1eeff0e82b90","order_by":3,"name":"Ilju Kim","email":"","orcid":"","institution":"Korea Advanced Institute of Science and Technology","correspondingAuthor":false,"prefix":"","firstName":"Ilju","middleName":"","lastName":"Kim","suffix":""},{"id":538242593,"identity":"25992380-6955-4bfd-82e2-dcfe87a184c0","order_by":4,"name":"Cinthya Paulina","email":"","orcid":"","institution":"Korea Advanced Institute of Science and Technology","correspondingAuthor":false,"prefix":"","firstName":"Cinthya","middleName":"","lastName":"Paulina","suffix":""},{"id":538242594,"identity":"db910b04-7ab4-4443-91a1-0b34c84bcfa3","order_by":5,"name":"Yewon Shin","email":"","orcid":"","institution":"Korea Advanced Institute of Science and Technology","correspondingAuthor":false,"prefix":"","firstName":"Yewon","middleName":"","lastName":"Shin","suffix":""},{"id":538242595,"identity":"afbaf342-6050-485c-8cdd-78d36f1314e7","order_by":6,"name":"Dongwoo Kim","email":"","orcid":"","institution":"Korea Advanced Institute of Science and Technology","correspondingAuthor":false,"prefix":"","firstName":"Dongwoo","middleName":"","lastName":"Kim","suffix":""},{"id":538242596,"identity":"30a69ecd-a0fa-4661-bbdd-34d51dd05bf9","order_by":7,"name":"Sungwoong Chung","email":"","orcid":"","institution":"Korea Advanced Institute of Science and Technology","correspondingAuthor":false,"prefix":"","firstName":"Sungwoong","middleName":"","lastName":"Chung","suffix":""},{"id":538242597,"identity":"dfc20158-e47b-45f7-8413-feef21d67363","order_by":8,"name":"Changhoon Park","email":"","orcid":"","institution":"Korea Advanced Institute of Science and Technology","correspondingAuthor":false,"prefix":"","firstName":"Changhoon","middleName":"","lastName":"Park","suffix":""},{"id":538242598,"identity":"6f5e60a9-aa2b-428d-9cc0-d3ac56943544","order_by":9,"name":"Jaewoong Han","email":"","orcid":"","institution":"Korea Advanced Institute of Science and Technology","correspondingAuthor":false,"prefix":"","firstName":"Jaewoong","middleName":"","lastName":"Han","suffix":""},{"id":538242599,"identity":"d86b6f71-5838-47f4-923b-9d143ef90f5b","order_by":10,"name":"Ji Hun Kim","email":"","orcid":"","institution":"Korea Advanced Institute of Science and Technology","correspondingAuthor":false,"prefix":"","firstName":"Ji","middleName":"Hun","lastName":"Kim","suffix":""},{"id":538242600,"identity":"b2a0b675-6874-4f6f-ab29-b264e3a92efd","order_by":11,"name":"Byung-Kwan Cho","email":"","orcid":"","institution":"Korea Advanced Institute of Science and Technology","correspondingAuthor":false,"prefix":"","firstName":"Byung-Kwan","middleName":"","lastName":"Cho","suffix":""}],"badges":[],"createdAt":"2025-10-22 05:31:16","currentVersionCode":1,"declarations":"","doi":"10.21203/rs.3.rs-7916426/v1","doiUrl":"https://doi.org/10.21203/rs.3.rs-7916426/v1","draftVersion":[],"editorialEvents":[],"editorialNote":"","failedWorkflow":false,"files":[{"id":95099566,"identity":"af0d5034-205f-48bf-912b-1668472d7352","added_by":"auto","created_at":"2025-11-04 09:51:06","extension":"png","order_by":1,"title":"Figure 1","display":"","copyAsset":false,"role":"figure","size":109341,"visible":true,"origin":"","legend":"\u003cp\u003e\u003cstrong\u003eMechanistic insights into LiPS-induced Li instability.\u003c/strong\u003e \u003cstrong\u003ea,\u003c/strong\u003e Schematic illustration of the conventional understanding of LiPS effects on the Li electrode. Corrosive reactions between LiPS and Li lead to poor reversibility. \u003cstrong\u003eb,\u003c/strong\u003e Schematic of the mechanism proposed in this work. LiPS forms microscale, localized phases within the electrolyte, resulting in spatially heterogeneous interactions with the Li electrode. This results in lower reversibility and uneven Li deposition due to limited formation of Li\u003csub\u003e2\u003c/sub\u003eO-containing SEI layers.\u003c/p\u003e","description":"","filename":"1.png","url":"https://assets-eu.researchsquare.com/files/rs-7916426/v1/908c472812435c370494e64b.png"},{"id":95099567,"identity":"22e52c46-1896-4f23-8cd7-ace7f6a70ba9","added_by":"auto","created_at":"2025-11-04 09:51:06","extension":"png","order_by":2,"title":"Figure 2","display":"","copyAsset":false,"role":"figure","size":453547,"visible":true,"origin":"","legend":"\u003cp\u003e\u003cstrong\u003eEffect of electrolyte amount on cycling stability.\u003c/strong\u003e \u003cstrong\u003ea\u003c/strong\u003e, Optical microscopy (OM) images of Li electrodes collected immediately after discharging Li-S pouch cells to 2.1 V at 0.1 C under different E/S ratios of 5.0 and 2.3 μL mg\u003csub\u003es\u003c/sub\u003e⁻¹. \u003cstrong\u003eb,\u003c/strong\u003e Schematic configuration of the assembled Li–S pouch cell. \u003cstrong\u003ec,\u003c/strong\u003e Cycling performance of 0.7 Ah Li–S pouch cells at various E/S ratios. The cell was cycled at a rate of 0.2 C (1 C= 1675 mA g\u003csub\u003es\u003c/sub\u003e\u003csup\u003e−1\u003c/sup\u003e). Cycling was continued until the cell retained 70% of its original capacity. Three parallel pouch cells were also tested at an E/S ratio of 2.3 μL mg\u003csub\u003es\u003c/sub\u003e\u003csup\u003e−1\u003c/sup\u003e. \u003cstrong\u003ed\u003c/strong\u003e, S cathodes and Li anodes collected from Li–S pouch cell after 46 cycles operation at 0.2C with an E/S ratio of 2.3 μL mg\u003csub\u003es\u003c/sub\u003e\u003csup\u003e−1\u003c/sup\u003e. \u003cstrong\u003ee,\u003c/strong\u003e Scanning electron microscopy (SEM) and energy-dispersive X-ray spectroscopy (EDS) images of S on deposited Li after 20 cycles at 0.2 C.\u003c/p\u003e","description":"","filename":"2.png","url":"https://assets-eu.researchsquare.com/files/rs-7916426/v1/698eaf1f9025bc02677cce7e.png"},{"id":95099569,"identity":"c3152104-60a6-43e4-81fd-8744b616e118","added_by":"auto","created_at":"2025-11-04 09:51:06","extension":"png","order_by":3,"title":"Figure 3","display":"","copyAsset":false,"role":"figure","size":272215,"visible":true,"origin":"","legend":"\u003cp\u003e\u003cstrong\u003eFormation of LiPS-rich phases in the electrolyte and their correlation with Li morphology.\u003c/strong\u003e \u003cstrong\u003ea,\u003c/strong\u003e Ultraviolet-visible spectra of catholytes (0.5 M and 1.0 M Li\u003csub\u003e2\u003c/sub\u003eS\u003csub\u003e4\u003c/sub\u003e in FN) and the FN electrolyte extracted from a pouch cell (E/S = 2.3 μL mg\u003csub\u003es\u003c/sub\u003e\u003csup\u003e–1\u003c/sup\u003e) after partial discharge. The cell operation was halted immediately before the voltage dip at the onset of the second voltage plateau, as shown in the inset voltage profile. All samples were diluted 10\u003csup\u003e4\u003c/sup\u003e times prior to measurement. \u003cstrong\u003eb, \u003c/strong\u003eOM images of bulk electrolytes with varying Li\u003csub\u003e2\u003c/sub\u003eS\u003csub\u003e4\u003c/sub\u003e concentrations showing the distribution of LiPS-rich phases. \u003cstrong\u003ec\u003c/strong\u003e, Corresponding size distribution histograms of the LiPS-rich domains. \u003cstrong\u003ed, e,\u003c/strong\u003e SEM images of Li electrodes operated for 15 cycles at 1 mA cm\u003csup\u003e–2\u003c/sup\u003e / 1 mAh cm\u003csup\u003e–2\u003c/sup\u003e in the catholytes with different Li\u003csub\u003e2\u003c/sub\u003eS\u003csub\u003e4\u003c/sub\u003e concentrations (d) and corresponding energy-dispersive X-ray spectroscopy (EDS) elemental maps of S atom (e). \u003cstrong\u003ef,\u003c/strong\u003e CE measurements of the Li|Ni cells with catholytes of varying Li\u003csub\u003e2\u003c/sub\u003eS\u003csub\u003e4\u003c/sub\u003e concentrations, measured using the modified Aurbach method.\u003c/p\u003e","description":"","filename":"3.png","url":"https://assets-eu.researchsquare.com/files/rs-7916426/v1/4b2fa67a80a99c43f53074e9.png"},{"id":95099571,"identity":"11cf9788-3d4e-44ac-af71-5aeabef4ef60","added_by":"auto","created_at":"2025-11-04 09:51:06","extension":"png","order_by":4,"title":"Figure 4","display":"","copyAsset":false,"role":"figure","size":266574,"visible":true,"origin":"","legend":"\u003cp\u003e\u003cstrong\u003eCharacterization of the SEI layer induced by the LiPS-rich phase.\u003c/strong\u003e a, Low-magnification cryo-transmission electron microscopy (cryo-TEM) image of Li deposited on a Ni grid in FN + 1.0 M Li\u003csub\u003e2\u003c/sub\u003eS\u003csub\u003e4\u003c/sub\u003e at a current density of 1 mA cm\u003csup\u003e–2\u003c/sup\u003e and an areal capacity of 1 mAh cm\u003csup\u003e–2\u003c/sup\u003e. b, c, Cryo-TEM images and corresponding selected area electron diffraction (SAED) patterns of the granular (b) and dendritic (c) Li morphologies. d, (Fluorine + nitrogen)/sulfur ratio obtained from EDS analysis of the dendritic Li region for the catholytes with different Li\u003csub\u003e2\u003c/sub\u003eS\u003csub\u003e4\u003c/sub\u003e concentrations. Inset: (F+N)/S ratio at the dendritic Li region and granular Li regions plated in FN +1.0 M Li\u003csub\u003e2\u003c/sub\u003eS\u003csub\u003e4\u003c/sub\u003e electrolyte. e, Raman spectra of FN and the LiPS-rich electrolyte phase in FN +1.0 M Li\u003csub\u003e2\u003c/sub\u003eS\u003csub\u003e4\u003c/sub\u003e. Inset: OM image of FN +1.0 M Li\u003csub\u003e2\u003c/sub\u003eS\u003csub\u003e4\u003c/sub\u003e showing the positions\u003csub\u003e \u003c/sub\u003e(#1, #2) for\u003csub\u003e \u003c/sub\u003eRaman analysis.\u003c/p\u003e","description":"","filename":"4.png","url":"https://assets-eu.researchsquare.com/files/rs-7916426/v1/31c2706a60873ef8938dfcc0.png"},{"id":95099568,"identity":"d0cfcc71-857d-4071-bfed-6e8d7aa63911","added_by":"auto","created_at":"2025-11-04 09:51:06","extension":"png","order_by":5,"title":"Figure 5","display":"","copyAsset":false,"role":"figure","size":244930,"visible":true,"origin":"","legend":"\u003cp\u003e\u003cstrong\u003eBeneficial effect of fragmenting LiPS-rich phases. a\u003c/strong\u003e, Changes in ⁷Li NMR chemical shift with the incorporation of 1.0 M Li\u003csub\u003e2\u003c/sub\u003eS\u003csub\u003e4\u003c/sub\u003e for FN, FNF, FNN, and FNI electrolytes. \u003cstrong\u003eb\u003c/strong\u003e, Optical microscopy images of LiPS-rich phases formed in the catholytes (top), and the corresponding SEM (middle) and EDS S mapping images (bottom) of the Li deposits collected from Li–Ni half-cells cycled with the respective catholytes for 15 cycles at 1 mA cm\u003csup\u003e−2\u003c/sup\u003e / 1 mAh cm\u003csup\u003e−2\u003c/sup\u003e. \u003cstrong\u003ec\u003c/strong\u003e, Histogram of LiPS-rich phase size distributions in the catholytes. \u003cstrong\u003ed,\u003c/strong\u003e CE measurements of the Li|Ni cells with catholytes of varying Li\u003csub\u003e2\u003c/sub\u003eS\u003csub\u003e4\u003c/sub\u003e concentrations, measured using the modified Aurbach method. \u003cstrong\u003ee\u003c/strong\u003e, Cycling performance of 0.7 Ah Li–S pouch cells tested at E/S ratio of 2.3 μL mg\u003csub\u003es\u003c/sub\u003e\u003csup\u003e−1\u003c/sup\u003e. The cells were cycled at a rate of 0.2 C (1 C= 1675 mA g\u003csub\u003es\u003c/sub\u003e\u003csup\u003e−1\u003c/sup\u003e). and cycling was terminated when 70% of the initial capacity was retained. \u003cstrong\u003ef,\u003c/strong\u003e Full-cell voltage profiles of Li-S pouch cells suing FNF, FNN, and FNI electrolytes.\u003c/p\u003e","description":"","filename":"5.png","url":"https://assets-eu.researchsquare.com/files/rs-7916426/v1/2f5088587fbbf312e0eb1f41.png"},{"id":96605057,"identity":"cffeecf2-8d7d-4c65-915a-c0f69c0f7b50","added_by":"auto","created_at":"2025-11-24 09:17:54","extension":"pdf","order_by":0,"title":"","display":"","copyAsset":false,"role":"manuscript-pdf","size":2694227,"visible":true,"origin":"","legend":"","description":"","filename":"manuscript.pdf","url":"https://assets-eu.researchsquare.com/files/rs-7916426/v1/3d4fd72e-4fec-4600-9f62-2ccfc582ab29.pdf"},{"id":95099570,"identity":"0cf2d506-cb55-45d5-8540-e920f2b135ab","added_by":"auto","created_at":"2025-11-04 09:51:06","extension":"docx","order_by":1,"title":"","display":"","copyAsset":false,"role":"supplement","size":12601228,"visible":true,"origin":"","legend":"Article File","description":"","filename":"Supplementarymaterial.docx","url":"https://assets-eu.researchsquare.com/files/rs-7916426/v1/c13d1aea42c7b0c514810b67.docx"}],"financialInterests":"There is \u003cb\u003eNO\u003c/b\u003e Competing Interest.","formattedTitle":"Localization of micron-scale lithium polysulfide-rich phases causes anode heterogeneity in lean-electrolyte lithium–sulfur batteries","fulltext":[{"header":"Main","content":"\u003cp\u003eThe growing demand for energy storage systems with higher energy densities is driving the development and adoption of next-generation batteries across various fields. Lithium\u0026ndash;sulfur (Li\u0026ndash;S) batteries are one of the most promising candidates for next-generation battery systems due to their high theoretical gravimetric density (\u0026gt;2600 Wh kg\u003csup\u003e\u0026minus;1\u003c/sup\u003e), low cost and environmental friendliness\u003csup\u003e1-4\u003c/sup\u003e. Achieving such high energy densities necessitates operating Li-S batteries with minimal electrolyte amounts, as electrolyte accounts for a significant portion of the cell\u0026apos;s weight (\u0026gt;40 wt%)\u003csup\u003e2,5,6\u003c/sup\u003e. However, Li-S batteries typically show low Coulombic efficiency (CE) and short lifespan operation (commonly \u0026lt; 40 cycles) under lean electrolyte conditions [Electrolyte / Sulfur (E/S) ratio \u0026lt; 3 \u0026mu;L mg\u003csub\u003es\u003c/sub\u003e\u003csup\u003e\u0026minus;1\u003c/sup\u003e]\u003csup\u003e7-9\u003c/sup\u003e. The limited cycling stability is primarily attributed to the dissolution of lithium polysulfides (LiPS, (Li\u003csub\u003e2\u003c/sub\u003eS\u003csub\u003eX\u003c/sub\u003e, X = 4, 6, 8))\u003csup\u003e10-13\u003c/sup\u003e, which are generated from the electrochemical conversion of solid-state cathode active materials (i.e., S\u003csub\u003e8\u003c/sub\u003e or Li\u003csub\u003e2\u003c/sub\u003eS) and readily diffuse from the S cathode to the Li anode. On the Li anode, dissolved LiPS is reduced to Li\u003csub\u003e2\u003c/sub\u003eS\u003csub\u003e2\u003c/sub\u003e/Li\u003csub\u003e2\u003c/sub\u003eS, corroding the Li electrode. The issue becomes more pronounced under lean-electrolyte conditions, where LiPS accumulates within the limited volume of liquid electrolyte\u003csup\u003e14-16\u003c/sup\u003e. Consequently, lean electrolyte conditions exacerbate Li anode instability, resulting in markedly reduced cycling performance\u003csup\u003e16-20\u003c/sup\u003e.\u003c/p\u003e\n\u003cp\u003eHowever, the specific influence of dissolved LiPS in the electrolyte on Li anode behavior has yet to be thoroughly clarified, and several fundamental questions remain unresolved. 1) Does the Li\u003csub\u003e2\u003c/sub\u003eS\u003csub\u003e2\u003c/sub\u003e/Li\u003csub\u003e2\u003c/sub\u003eS-containing solid electrolyte-interphase (SEI), formed via reactions between the Li anode and LiPS, inherently causes adverse effects on Li morphology? Some studies have shown that a moderate amount of LiPS can actually improve Li morphology, suggesting that corrosion of LiPS itself may not be inherently detrimental to anode performance\u003csup\u003e21,22\u003c/sup\u003e. 2) Under lean-electrolyte conditions, is the degradation of the Li anode at elevated LiPS concentrations fully accounted for by an increased corrosion, or are additional mechanisms at play? Attributing Li anode degradation solely to severe corrosion lacks sufficient causal evidence. For examples, certain electrolytes containing reactive species (e.g., LiFSI or LiDFOB, etc.) have been reported to promote dense Li deposition by accelerating the formation of a inorganic-rich SEI, ultimately extending battery cycle life\u003csup\u003e23,24\u003c/sup\u003e. More detailed understanding of the mechanisms by which LiPS influence Li anode is essential for developing strategies to overcome the short cycle life of Li-S batteries under lean-electrolyte conditions.\u003c/p\u003e\n\u003cp\u003eHerein, we propose more fundamental mechanism of LiPS-induced instability of Li anode that accelerates cell degradation under lean electrolyte condition (\u003cstrong\u003eFig. 1a,b\u003c/strong\u003e). We discovered that the clustering nature of LiPS\u003csup\u003e25,26\u003c/sup\u003e results in the formation of localized tens of micron-scale LiPS-rich phases in the electrolyte, inducing significant heterogeneity of Li anode. Notably, the chemical composition of the LiPS-rich phase contains minimal amounts of salt anions such as FSI\u003csup\u003e\u0026minus;\u003c/sup\u003e or NO\u003csub\u003e3\u003c/sub\u003e\u003csup\u003e\u0026minus;\u003c/sup\u003e, which are typically responsible for forming stable SEI. This deficiency of FSI\u003csup\u003e\u0026minus;\u003c/sup\u003e or NO\u003csub\u003e3\u003c/sub\u003e\u003csup\u003e\u0026minus;\u003c/sup\u003e in LiPS-rich phase is found to exert a detrimental influence on Li morphology by obstructing formation of Li\u003csub\u003e2\u003c/sub\u003eO-containing SEI, a species known to be found in electrolytes that induces a homogeneous Li conformation. Based on this understanding, we propose an additive strategy to reduce the localized nature of PS-rich phases by fragmenting them using various salt anions that can control the degree of LiPS-rich phase locality.\u0026nbsp;With the designed electrolyte, our assembled prototype 0.7 Ah stacked pouch cell could deliver higher cycling stability of 134 cycles under lean electrolyte condition of E/S ratio 2.3 \u0026mu;L mg\u003csub\u003es\u003c/sub\u003e\u003csup\u003e\u0026minus;1\u003c/sup\u003e (E/C ratio = 1.9 g Ah\u003csup\u003e\u0026minus;1\u003c/sup\u003e), compared to 46 cycles at reference electrolyte. This study expands understanding of LiPS-containing catholytes, proposing that degradation of Li anodes is due to local phase formation and subsequent SEI changes by LiPS, rather than the prevailing view that degradation is due to the intrinsic corrosiveness of LiPS.\u003cbr\u003e\u0026nbsp;\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003eSharp decrease in lifespan of Li-S batteries under lean electrolyte conditions\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003eGiven that high LiPS concentrations in the electrolyte are often considered crucial under lean conditions, LiPS are likely to exert a critical influence on the Li. Optical microscopy (OM) images of Li electrodes collected from two Li\u0026ndash;S pouch cells (\u003cstrong\u003eSupplementary Fig. 1\u003c/strong\u003e) with different E/S ratios reveal the formation of micrometer-scale domains with LiPS residues with a distinct spatial distribution across the electrode surface (\u003cstrong\u003eFig. 2a\u003c/strong\u003e). At an E/S ratio of 5.0 \u0026mu;L mg\u003csub\u003es\u003c/sub\u003e\u003csup\u003e\u0026minus;\u003c/sup\u003e\u003csup\u003e1\u003c/sup\u003e, only a few LiPS deposits are observed on the Li electrode. In contrast, at an E/S ratio of 2.3 \u0026mu;L mg\u003csub\u003es\u003c/sub\u003e\u003csup\u003e\u0026minus;\u003c/sup\u003e\u003csup\u003e1\u003c/sup\u003e, LiPS precipitation increases markedly, spreading extensively across the electrode. The characteristic sizes of these LiPS-rich phases were determined to be D\u003csub\u003e50\u003c/sub\u003e = 7 \u0026mu;m and D\u003csub\u003e90\u003c/sub\u003e = 22 \u0026mu;m (\u003cstrong\u003eSupplementary Fig. 2\u003c/strong\u003e). These results highlight that LiPS accumulates on the Li electrode at the micrometer scale, forming localized phases.\u003c/p\u003e\n\u003cp\u003eIt has been reported that Li\u0026ndash;S batteries exhibit a significant decline in cycling stability under lean electrolyte conditions\u003csup\u003e9\u003c/sup\u003e. To investigate this behavior, we assembled a prototype 0.7 Ah pouch cells (\u003cstrong\u003eFig. 2b\u003c/strong\u003e and \u003cstrong\u003eSupplementary Fig. 3\u003c/strong\u003e) and evaluated their cycling performance. As shown in \u003cstrong\u003eFig. 2c\u003c/strong\u003e, the pouch cells sustained over 200 cycles before reaching 70% capacity retention at an E/S ratio of 5.0 and 3.5 \u0026mu;L mg\u003csub\u003es\u003c/sub\u003e\u003csup\u003e\u0026minus;\u003c/sup\u003e\u003csup\u003e1\u003c/sup\u003e, which provides a sufficient amount of electrolyte for cell operation. In contrast, under a leaner condition with an E/S ratio of 2.3 \u0026mu;L mg\u003csub\u003es\u003c/sub\u003e\u003csup\u003e\u0026minus;\u003c/sup\u003e\u003csup\u003e1\u003c/sup\u003e, only 46 cycles were achieved, reflecting significant degradation in cycling stability.\u0026nbsp;Voltage profiles during cycling further reveal that the electrolyte volume strongly affects cell performance, as evidenced by a pronounced increase in overpotential over successive cycles at an E/S ratio of 2.3\u0026nbsp;\u0026mu;L mg\u003csub\u003es\u003c/sub\u003e\u003csup\u003e\u0026minus;\u003c/sup\u003e\u003csup\u003e1\u003c/sup\u003e, in contrast to the more stable profiles observed at E/S ratios of 5.0 and 3.5\u0026nbsp;\u0026mu;L mg\u003csub\u003es\u003c/sub\u003e\u003csup\u003e\u0026minus;\u003c/sup\u003e\u003csup\u003e1\u003c/sup\u003e (\u003cstrong\u003eSupplementary Fig. 4\u003c/strong\u003e). \u003cstrong\u003eThe Li anodes collected from a Li\u0026ndash;S pouch cell at an E/S ratio of 2.3\u0026nbsp;\u003c/strong\u003e\u0026mu;L mg\u003csub\u003es\u003c/sub\u003e\u003csup\u003e\u0026minus;\u003c/sup\u003e\u003csup\u003e1\u003cstrong\u003e\u0026nbsp;\u003c/strong\u003e\u003c/sup\u003e\u003cstrong\u003eafter 46 cycle operation exhibited uneven morphology, implying severe degradation (Fig. 2d).\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003eScanning electron microscopy (SEM) images of the Li electrodes after 20 cycles from the pouch cells show that electrolyte deficiency leads to more irregular Li deposition. (\u003cstrong\u003eFig. 2e\u003c/strong\u003e). Notably, the morphological evolution at 2.3 \u0026mu;L mg\u003csub\u003es\u003c/sub\u003e\u003csup\u003e\u0026minus;\u003c/sup\u003e\u003csup\u003e1\u003c/sup\u003e features distinct structural characteristics, where the deposited Li exhibits heterogeneous surface morphologies comprising both dendritic and granular domains. This heterogeneity was observed across tens of micrometers, indicating substantial reaction heterogeneity. Energy dispersive X-ray spectroscopy (EDS) analysis was conducted to examine the spatial distribution of the chemical composition of the SEI layer. At an E/S ratio of 5 \u0026mu;L mg\u003csub\u003es\u003c/sub\u003e\u003csup\u003e\u0026minus;\u003c/sup\u003e\u003csup\u003e1\u003c/sup\u003e\u0026mdash;where stable cycling was observed\u0026mdash;S atoms, mainly originating from LiPS, showed relatively low locality. At an E/S ratio of 3.5 \u0026mu;L mg\u003csub\u003es\u003c/sub\u003e\u003csup\u003e\u0026minus;\u003c/sup\u003e\u003csup\u003e1\u003c/sup\u003e, where porous Li structures began to appear, S distribution became more localized. Notably, at an E/S ratio of 2.3 \u0026mu;L mg\u003csub\u003es\u003c/sub\u003e\u003csup\u003e\u0026minus;\u003c/sup\u003e\u003csup\u003e1\u003c/sup\u003e\u0026mdash;where a sharp decline in cycling performance was observed\u0026mdash;the spatial distribution of S within the Li electrode exhibited strong localization. The size of the localized regions progressively increases with decreasing E/S ratio.\u003c/p\u003e\n\u003cp\u003eTo sum up, observations indicate that under lean electrolyte conditions, micron-scale LiPS-rich phases accumulate locally on the Li electrode, while heterogeneous Li structures also emerge at comparable length scales, which we identify as a primary cause of Li anode degradation in Li\u0026ndash;S full cells. Taken together, these findings lead us to hypothesize that the formation of localized LiPS-rich phases is a key factor driving Li heterogeneity and, ultimately, premature cell failure, as will be discussed in the following section.\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003eMicron-scale LiPS-rich phases in the catholyte\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003eLi\u003csub\u003e2\u003c/sub\u003eS\u003csub\u003e4\u003c/sub\u003e was selected as the representative LiPS species for catholyte modeling, as it predominates at the maximum LiPS concentration during discharge in Li\u0026ndash;S full cells\u003csup\u003e27-29\u003c/sup\u003e. Catholytes containing 0.2, 0.5, 0.8, and 1.0 M Li\u003csub\u003e2\u003c/sub\u003eS\u003csub\u003e4\u003c/sub\u003e were prepared in the reference electrolyte (FN, 0.75 M LiFSI in 1, 2-dimethoxyethane (DME): 2-methylfuran (2MF) 4:1 in volume with 5 wt% LiNO\u003csub\u003e3\u003c/sub\u003e). The 1.0 M Li\u003csub\u003e2\u003c/sub\u003eS\u003csub\u003e4\u003c/sub\u003e concentration corresponds to the maximum Li\u003csub\u003e2\u003c/sub\u003eS\u003csub\u003e4\u003c/sub\u003e level expected at an E/S ratio of 2.3 \u0026mu;L mg\u003csub\u003es\u003c/sub\u003e\u003csup\u003e\u0026ndash;1\u003c/sup\u003e (\u003cstrong\u003eFig. 3a\u003c/strong\u003e, \u003cstrong\u003eSupplementary Note 1\u003c/strong\u003e and \u003cstrong\u003eSupplementary Fig. 5\u003c/strong\u003e), while the lower concentrations represent more flooded-electrolyte conditions.\u003c/p\u003e\n\u003cp\u003eThe catholytes are opaque as shown in insets in \u003cstrong\u003eFig. 3b\u003c/strong\u003e, implying their inhomogeneous structure. OM images of the catholytes reveal the formation of micron-scale, localized LiPS-rich phases, which grow in size with increasing the LiPS concentration (\u003cstrong\u003eFig. 3b,c\u003c/strong\u003e). While the diameters of LiPS-rich phases in FN + 0.2 M Li\u003csub\u003e2\u003c/sub\u003eS\u003csub\u003e4\u003c/sub\u003e were D\u003csub\u003e50\u003c/sub\u003e=4 \u0026mu;m and D\u003csub\u003e90\u003c/sub\u003e=5 \u0026mu;m, increasing the Li\u003csub\u003e2\u003c/sub\u003eS\u003csub\u003e4\u003c/sub\u003e concentration to 0.5 M, 0.8 M, and 1.0 M resulted in size increases to D\u003csub\u003e50\u003c/sub\u003e= 4\u0026nbsp;\u0026mu;m, 6\u0026nbsp;\u0026mu;m, 9\u0026nbsp;\u0026mu;m, and D\u003csub\u003e90\u003c/sub\u003e= 9 \u0026mu;m, 12 \u0026mu;m, and 21 \u0026mu;m, respectively (\u003cstrong\u003eFig. 3c\u003c/strong\u003e). In contrast,\u0026nbsp;no visible features were observed in the LiPS-free FN electrolyte. These results indicate that, under lean electrolyte conditions where LiPS concentrations are high, LiPSs significantly aggregate, leading to the formation of more discrete and localized phases. Importantly, the OM observation implies that the increasingly aggregated state of LiPS with LiPS concentration observed in \u003csup\u003e7\u003c/sup\u003eLi nuclear magnetic resonance (NMR) spectroscopy analysis (\u003cstrong\u003eSupplementary Fig. 6\u003c/strong\u003e)\u0026nbsp;and MD simulations (\u003cstrong\u003eSupplementary Fig. 7\u003c/strong\u003e) not only reflects changes in coordination structures at the molecular scale, but also leads to the formation of micron-scale inhomogeneous LiPS-rich phases.\u003c/p\u003e\n\u003cp\u003eGiven that more uneven Li morphologies were observed at higher LiPS concentrations in \u003cstrong\u003eFig. 2e\u003c/strong\u003e, we hypothesized that the emergence of heterogeneity of Li anode is closely associated with localized LiPS-rich phases at electrolyte. As shown in \u003cstrong\u003eFig. 3d,\u003c/strong\u003e Li|Ni half cells employing catholytes with different Li\u003csub\u003e2\u003c/sub\u003eS\u003csub\u003e4\u003c/sub\u003e concentrations exhibited distinct Li morphologies after cycling. Both the FN electrolyte and the FN + 0.2 M Li\u003csub\u003e2\u003c/sub\u003eS\u003csub\u003e4\u003c/sub\u003e catholyte promoted dense and uniform Li structure, whereas FN electrolyte with\u0026nbsp;\u0026nbsp;\u0026nbsp;0.5 M\u0026nbsp;Li\u003csub\u003e2\u003c/sub\u003eS\u003csub\u003e4\u003c/sub\u003e led to heterogeneous Li morphology, featuring the coexistence of granular and porous Li domains. Such uneven Li morphology closely resembles that observed at the Li metal anode in the Li\u0026ndash;S full cell. In particular, we observed a significant difference in elemental distribution between the two Li regions; as shown in \u003cstrong\u003eFig. 3e\u003c/strong\u003e, the distribution of the S element is noticeably concentrated in the dendritic region compared to the granular region. This suggests the possibility that the formation of dendritic regions is induced by the tens of micron-scale LiPS-rich phases. In a later section, we explore how SEI properties correlate with Li morphology in S-concentrated regions.\u003c/p\u003e\n\u003cp\u003e\u0026nbsp; \u0026nbsp; \u0026nbsp; \u0026nbsp; \u0026nbsp; \u0026nbsp; Dendritic and porous Li structure, with its high surface area, generally exacerbates side reactions with the electrolyte and leads to lower CE. Inclusion of more LiPS in FN electrolyte resulted in lower initial CE (\u003cstrong\u003eSupplementary Fig. 8\u003c/strong\u003e) and subsequent average CE measured by modified Aurbach protocol (\u003cstrong\u003eFig. 3f\u003c/strong\u003e). While the FN electrolyte achieved an average CE of 99.40%, the addition of 0.2 M, 0.5 M, 0.8 M, and 1.0 M LiPS decreased it to 99.12 %, 98.96 %, 98.37 %, and 97.98 %, reflecting the adverse influence of LiPS-induced morphological heterogeneity in Li.\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003eCharacterization of SEI layer\u003c/strong\u003e \u003cstrong\u003einduced by LiPS-rich phases\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003eThe reason behind the formation of dendritic Li induced by the LiPS-rich phase remains to be clarified. The SEI characteristics of Li deposits formed in FN + 1.0 M Li\u003csub\u003e2\u003c/sub\u003eS\u003csub\u003e4\u003c/sub\u003e electrolyte were investigated using cryogenic transmission electron microscopy (cryo-TEM) (\u003cstrong\u003eFig. 4a-c\u003c/strong\u003e). At the granular Li region, the SEI layer exhibited crystalline features composed of Li\u003csub\u003e2\u003c/sub\u003eO crystallites, as further confirmed by the selected area electron diffraction (SAED) pattern. The formation of Li\u003csub\u003e2\u003c/sub\u003eO component in SEI is mainly attributed to the decomposition of FSI\u003csup\u003e\u0026minus;\u003c/sup\u003e or NO\u003csub\u003e3\u003c/sub\u003e\u003csup\u003e\u0026minus;\u003c/sup\u003e anions, which has been reported to contribute to the development of stable SEI layers and granular Li morphology\u003csup\u003e30,31\u003c/sup\u003e. In contrast, the dendritic Li structure exhibited only weak crystalline signals. A similar lack of crystalline features was also observed on Li dendrites formed in the FN + 0.5 M Li\u003csub\u003e2\u003c/sub\u003eS\u003csub\u003e4\u003c/sub\u003e (\u003cstrong\u003eSupplementary Fig. 9\u003c/strong\u003e). Notably, clear crystalline signals were detected when using FN electrolyte without LiPS species (\u003cstrong\u003eSupplementary Fig.10\u003c/strong\u003e).\u003c/p\u003e\n\u003cp\u003eWhen comparing the (fluorine + nitrogen)/sulfur ratio from EDS analysis\u0026mdash;which reflects the relative contributions of LiFSI and LiNO\u003csub\u003e3\u003c/sub\u003e versus LiPS to SEI formation\u0026mdash;a gradual decrease is observed with increasing Li\u003csub\u003e2\u003c/sub\u003eS\u003csub\u003e4\u003c/sub\u003e concentration in the catholyte (\u003cstrong\u003eFig. 4d\u003c/strong\u003e and\u0026nbsp;\u003cstrong\u003eSupplementary Fig. 11\u003c/strong\u003e). This suggests that SEI layers formed by LiPS-rich phases contain a lower proportion of LiFSI and LiNO\u003csub\u003e3\u003c/sub\u003e-derived components, which have been associated with granular Li morphology\u003csup\u003e32,33\u003c/sup\u003e. To further clarify the chemical environment of LiPS-rich phase, Raman spectroscopy was conducted on FN and FN + 1.0 M Li\u003csub\u003e2\u003c/sub\u003eS\u003csub\u003e4\u003c/sub\u003e electrolytes (\u003cstrong\u003eFig. 4e\u003c/strong\u003e). The spectrum of the\u0026nbsp;LiPS-rich phase shows strong S\u003csub\u003eX\u003c/sub\u003e\u003csup\u003e2\u0026minus;\u003c/sup\u003e and S\u003csub\u003e4\u003c/sub\u003e\u003csup\u003e2\u0026minus;\u003c/sup\u003e signals at 400\u0026ndash;500 cm\u003csup\u003e\u0026minus;1\u003c/sup\u003e, along with a weak DME signal at ~850 cm\u003csup\u003e\u0026ndash;1\u0026nbsp;\u003c/sup\u003e\u003csup\u003e27,34,35\u003c/sup\u003e, indicating the presence of only a small amount of solvent molecules. Notably, signals corresponding to FSI\u003csup\u003e\u0026minus;\u003c/sup\u003eand NO\u003csub\u003e3\u003c/sub\u003e\u003csup\u003e\u0026minus;\u003c/sup\u003e, which are present in the FN electrolyte, are barely detectable\u003csup\u003e36,37\u003c/sup\u003e. These observations collectively account for the markedly suppressed Li\u003csub\u003e2\u003c/sub\u003eO formation in contact with LiPS-rich phases.\u003c/p\u003e\n\u003cp\u003eMolecular dynamics (MD) simulations also exhibited the localized distribution of LiPS within the electrolyte. At higher concentrations, the S\u003csub\u003e4\u003c/sub\u003e\u003csup\u003e2\u003c/sup\u003e\u003csup\u003e\u0026minus;\u003c/sup\u003e anion exhibited enhanced coordination with Li⁺, consistent with the clustering tendency of LiPS (\u003cstrong\u003eSupplementary Fig. 7 and 12\u0026ndash;14\u003c/strong\u003e). Notably, these LiPS-containing clusters hardly contain\u0026nbsp;other anions (i.e.,\u0026nbsp;FSI⁻ and NO\u003csub\u003e3\u003c/sub\u003e\u003csup\u003e\u0026minus;\u003c/sup\u003e). The reduced presence of FSI⁻ and NO\u003csub\u003e3\u003c/sub\u003e\u003csup\u003e\u0026minus;\u003c/sup\u003e in these LiPS-containing clusters limits their decomposition, thereby suppressing the formation of LiFSI and LiNO\u003csub\u003e3\u003c/sub\u003e-derived SEI layers that are essential for effective passivation of the Li anode.\u003c/p\u003e\n\u003cp\u003eWe further investigated the behavior of other soluble LiPS species, specifically Li\u003csub\u003e2\u003c/sub\u003eS\u003csub\u003e8\u003c/sub\u003e and Li\u003csub\u003e2\u003c/sub\u003eS\u003csub\u003e6\u003c/sub\u003e (\u003cstrong\u003eSupplementary Fig. 15-16)\u003c/strong\u003e. The sulfur concentration in LiPS species was adjusted to ~14 wt% in the electrolyte (i.e., 0.5 M Li\u003csub\u003e2\u003c/sub\u003eS\u003csub\u003e8\u003c/sub\u003e and 0.67 M Li\u003csub\u003e2\u003c/sub\u003eS\u003csub\u003e6\u003c/sub\u003e) to reflect their actual levels in Li\u0026ndash;S full cells\u003csup\u003e38\u003c/sup\u003e. Although longer-chain LiPS have higher redox potentials\u0026mdash;implying greater corrosivity toward the Li anode\u0026mdash;their impact on Li reversibility and morphology was not more detrimental. In fact, more uniform Li deposition and higher CE compared to Li\u003csub\u003e2\u003c/sub\u003eS\u003csub\u003e4\u003c/sub\u003e were observed. MD simulations showed that these species exhibit a weaker tendency to form discrete phase domains due to reduced clustering propensities (\u003cstrong\u003eSupplementary Fig. 17-18\u003c/strong\u003e) which was also observed from previous studies\u003csup\u003e39\u003c/sup\u003e. OM observation for bulk catholytes containing Li\u003csub\u003e2\u003c/sub\u003eS\u003csub\u003e8\u003c/sub\u003e and Li\u003csub\u003e2\u003c/sub\u003eS\u003csub\u003e6\u003c/sub\u003e displayed smaller LiPS-phases indicated by smaller D\u003csub\u003e50\u003c/sub\u003e and D\u003csub\u003e90\u003c/sub\u003e values (\u003cstrong\u003eSupplementary Fig. 19\u003c/strong\u003e). These findings suggest that the spatial localization of LiPS, rather than their inherent corrosivity, plays a more critical role in determining Li anode stability.\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003eFragmenting LiPS-rich phases to decrease locality\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003eFollowing the identification of the detrimental effects of localized LiPS-rich phases on the Li electrode, LiI was selected to examine its ability to suppress such localization. I⁻, known for its high donor number (DN) (29 kcal mol\u003csup\u003e\u0026minus;1\u003c/sup\u003e) and high Li\u003csup\u003e+\u003c/sup\u003e binding energy (\u003cstrong\u003eSupplementary Fig. 20)\u003c/strong\u003e\u003csup\u003e40\u003c/sup\u003e, is expected to inhibit LiPS clustering by competitively coordinating with Li⁺, thereby suppressing the formation of large LiPS-rich phases. To verify this effect, 0.1 M LiI was added to FN electrolyte, denoted as FNI electrolyte. For comparison, 0.1M LiFSI (DN of 9.5 kcal mol\u003csup\u003e\u0026minus;1\u003c/sup\u003e) and LiNO₃ (DN of 19 kcal mol\u003csup\u003e\u0026minus;1\u003c/sup\u003e), which are baseline salts, were also introduced, yielding FNF and FNN electrolytes, respectively. These electrolytes were designed to evaluate how anion properties influence the extent of LiPS-rich phase formation.\u003c/p\u003e\n\u003cp\u003eThe downfield shifts of \u003csup\u003e7\u003c/sup\u003eLi NMR signal upon incorporation of 1.0 M Li\u003csub\u003e2\u003c/sub\u003eS\u003csub\u003e4\u003c/sub\u003e were compared among the electrolytes (\u003cstrong\u003eFig. 5a, Supplementary Fig. 21\u003c/strong\u003e). These shifts originate from the coordination of PS anions with Li\u003csup\u003e+\u003c/sup\u003e as LiPS-rich phases evolve. FNF exhibited a shift of 0.18 ppm, nearly identical to FN (0.19 ppm), indicating that the addition of 0.1M LiFSI had little effect on Li\u003csup\u003e+\u003c/sup\u003e PS interactions. In FNN, the shift was slightly smaller at 0.16 ppm, reflecting a modest reduction in Li\u003csup\u003e+\u003c/sup\u003e PS coordination due to the 0.1M NO\u003csub\u003e3\u003c/sub\u003e\u003csup\u003e\u0026minus;\u003c/sup\u003e anion. Notably, FNI showed a substantially smaller shift of 0.09 ppm, demonstrating that 0.1M LiI effectively suppresses Li\u003csup\u003e+\u003c/sup\u003e PS coordination, thereby mitigating the development of LiPS-rich phases.\u003c/p\u003e\n\u003cp\u003eSuch changes in the local environment of Li⁺ are, as previously investigated, associated with the micron-scale clustering of LiPS. To investigate the regulation of LiPS-rich phases by additional anions, the OM images of the 1.0 M Li\u003csub\u003e2\u003c/sub\u003eS\u003csub\u003e4\u003c/sub\u003e-containing electrolytes were captured, and the size distributions of LiPS-rich phases were determined (\u003cstrong\u003eFig. 5b,c\u003c/strong\u003e). FNF and FNN had limited impact on regulating LiPS-rich phases, consistent with the NMR findings. Compared to FN + 1.0 M Li\u003csub\u003e2\u003c/sub\u003eS\u003csub\u003e4\u0026nbsp;\u003c/sub\u003e(D\u003csub\u003e50\u003c/sub\u003e= 9 \u0026mu;m and D\u003csub\u003e90\u003c/sub\u003e= 21 \u0026mu;m, Fig. 3c), the D\u003csub\u003e50\u0026nbsp;\u003c/sub\u003e/D\u003csub\u003e90\u003c/sub\u003e values slightly decreased to D\u003csub\u003e50\u003c/sub\u003e= 7 \u0026mu;m and D\u003csub\u003e90\u003c/sub\u003e= 19 \u0026mu;m \u0026nbsp;for FNF + 1.0 M Li\u003csub\u003e2\u003c/sub\u003eS\u003csub\u003e4\u003c/sub\u003e, and D\u003csub\u003e50\u003c/sub\u003e= 8 \u0026mu;m and D\u003csub\u003e90\u003c/sub\u003e= 17 \u0026mu;m for FNN + 1.0 M Li\u003csub\u003e2\u003c/sub\u003eS\u003csub\u003e4\u003c/sub\u003e. In contrast, FNI further reduced the size of LiPS-rich phases, lowering the D\u003csub\u003e50\u003c/sub\u003e and D\u003csub\u003e90\u003c/sub\u003e values to 7 and 13 \u0026mu;m, respectively\u0026mdash;approximately 62% of those measured in FN electrolyte. This result indicates that the disruption of LiPS clustering with I\u003csup\u003e\u0026minus;\u003c/sup\u003e leads to fragmentation of LiPS-rich phases, thereby decreasing locality. The correlation between the locality of LiPS-rich phases and the uniformity in Li morphology was confirmed through the SEM images of the cycled Li electrodes (\u003cstrong\u003eFig. 5b\u003c/strong\u003e). FNF and FNN showed concurrent evolution of porous and granular Li morphologies, with the porous regions corresponding to areas with high S content in the EDS images, indicating the formation of localized S-rich SEI. In contrast, FNI exhibited smooth and uniform Li morphology and diminished localization of S content, indicating a more homogeneous SEI layer. These findings demonstrate that mitigating the localization of LiPS-rich phases within the electrolytes, particularly via the addition of I\u003csup\u003e-\u003c/sup\u003e, promotes uniform Li deposition.\u003c/p\u003e\n\u003cp\u003e\u0026nbsp; \u0026nbsp; \u0026nbsp; \u0026nbsp; \u0026nbsp; \u0026nbsp;\u0026nbsp;For the 1.0 M Li\u003csub\u003e2\u003c/sub\u003eS\u003csub\u003e4\u003c/sub\u003e-containing electrolytes, the CE of the Li electrode was measured using the modified Aurbach method. Despite a decrease in ionic conductivity (\u003cstrong\u003eSupplementary Fig. 22\u003c/strong\u003e), the CE increased with the DN of the added salt anion (\u003cstrong\u003eFig. 5d\u0026nbsp;\u003c/strong\u003eand\u003cstrong\u003e\u0026nbsp;Supplementary Fig. 23\u003c/strong\u003e); FNF (98.07%), FNN (98.30%), and FNI (98.58%) exhibited an increase in CE compared to FN (97.98%). The higher Li reversibility in the LiPS-containing FNI emphasizes the importance of suppressing the localization. Under LiPS-free conditions, the average CE of FNI (99.37%) was nearly identical with that of FN (99.40%, \u003cstrong\u003eFig. 3f\u003c/strong\u003e), indicating that the improved Li reversibility is associated with the reduced locality of LiPS-rich phases (\u003cstrong\u003eSupplementary Fig. 24\u003c/strong\u003e).\u003c/p\u003e\n\u003cp\u003eWe assembled Li-S full cells to demonstrate the efficacy of the approaches in enhancing full cell cycling stability (\u003cstrong\u003eFig. 5e,f\u003c/strong\u003e). The FN cell\u0026nbsp;showed rapid capacity fading due to Li instability caused by LiPS-rich phases, resulting in only 46 cycles of operation. FNF and FNN, which moderately improved Li reversibility and Li deposition morphology, showed extended cycling performances (58 and 84 cycles for FNF and FNN, respectively), with FNN outperforming FNF due to the higher DN of NO\u003csub\u003e3\u003c/sub\u003e\u003csup\u003e\u0026minus;\u003c/sup\u003e compared to FSI\u003csup\u003e\u0026minus;\u003c/sup\u003e. In contrast, FNI, which markedly reduced the locality of LiPS-rich phases, delivered superior cycling stability, retaining 70% of their capacity after more than 134 cycles. Notably, the observed improvement is not attributable to the iodine redox reaction, commonly used to enhance redox kinetics, as the operating voltage window lies outside the redox potential of LiI (oxidation potential of ~3.0 V vs. Li/Li\u003csup\u003e+\u003c/sup\u003e)\u003csup\u003e41,42\u003c/sup\u003e.\u003c/p\u003e\n\u003cp\u003e\u0026nbsp; \u0026nbsp; \u0026nbsp; \u0026nbsp; \u0026nbsp; \u0026nbsp;\u0026nbsp;We analyzed the Li electrodes after 20 cycles using SEM and EDS (\u003cstrong\u003eSupplementary Fig. 25)\u003c/strong\u003e. Pronounced heterogeneity in Li morphology, indicative of the locality of LiPS-rich phases, was observed for FNF and FNN, whereas FNI exhibited a much more uniform Li surface. The spatial distribution of S signals overlapped well with the porous Li regions, indicating that S-rich SEI preferentially formed on the locally porous Li structures. Suppressing the locality of LiPS-rich phases thus facilitates more uniform Li morphology, leading to improved cycling stability.\u003c/p\u003e\n\u003cp\u003eA prototype pouch cell was assembled using the FNI electrolyte at a lower E/S ratio of 1.5\u0026nbsp;\u0026mu;L mg\u003csub\u003es\u003c/sub\u003e\u003csup\u003e\u0026minus;1\u003c/sup\u003e (E/C ratio = 1.3 g Ah\u003csup\u003e\u0026minus;1\u003c/sup\u003e) to validate the cycling stability improvements achieved by mitigating the locality of LiPS-rich phases. (\u003cstrong\u003eSupplementary Fig. 26-27\u0026nbsp;\u003c/strong\u003eand \u003cstrong\u003eTable 1\u003c/strong\u003e). Due to very low dosage of electrolyte, our prototype pouch cell achieved an energy density of 524 Wh kg\u003csup\u003e\u0026minus;1\u003c/sup\u003e (including tabs and packages, 324 Wh kg\u003csup\u003e\u0026minus;1\u003c/sup\u003e).\u0026nbsp;Although Li-S full cells typically operate for only a few cycles under such a low E/S ratio due to instability of Li electrode\u003csup\u003e7\u003c/sup\u003e, our cell demonstrated stable operation for 45 cycles, underscoring the importance of controlling locality of LiPS-rich phases to enhance cycling stability under lean electrolyte conditions.\u003c/p\u003e"},{"header":"Conclusions","content":"\u003cp\u003eIn this study, we investigated the effects of LiPS on the Li anode, which are determined by micron-scale locality of LiPS-rich phases. Due to the clustering nature of LiPS, LiPS-rich phases form under lean electrolyte conditions. Sporadically distributed micron scale LiPS-rich phases, characterized by low FSI\u003csup\u003e−\u003c/sup\u003e and NO\u003csub\u003e3\u003c/sub\u003e\u003csup\u003e−\u003c/sup\u003e concentrations, undermine anode stability by hindering the formation of LiFSI or LiNO\u003csub\u003e3\u003c/sub\u003e-derived SEI layer and facilitating non-uniform Li growth, thereby contributing to unstable cell performance. We demonstrated that decreasing locality of LiPS-rich phases can significantly reduce anode heterogeneity, as evidenced by the improved cycling stability. We conclude that the locality of LiPS-rich phases is the root cause of instability in lean-electrolyte Li-S batteries. The insights gained from this study will provide a foundation for developing engineering strategies to address the issue and pave the way for the implementation of high energy density, yet stable, Li-S battery systems.\u003c/p\u003e"},{"header":"Methods","content":"\u003cp\u003e\u003cstrong\u003eMaterials preparation.\u003c/strong\u003e 1,2-Dimethoxyethane (DME) (anhydrous, 99.5%), 2-methylfuran (2MF) (99%), lithium nitrate (LiNO\u003csub\u003e3\u003c/sub\u003e) (99.9%), and lithium iodide (LiI) (99.9%) were purchased from Sigma Aldrich. Lithium bis(fluorosulfonyl)imide (LiFSI) (99.9%) was obtained from Nippon shokubai. Lithium foils with thicknesses of 40 \u0026micro;m, 150 \u0026micro;m, and 450 \u0026micro;m, as well as double-sided Li-Cu-Li (50 \u0026micro;m), were sourced from Honjo Metal.\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003ePreparation of electrolytes.\u003c/strong\u003e The reference electrolyte (FN) was formulated by dissolving 0.75 M LiFSI in a solvent mixture of 1,2-dimethoxyethane (DME) and 2-methylfuran (2-MF) (volume ratio 4:1), with the addition of 5 wt% LiNO\u003csub\u003e3\u003c/sub\u003e as an additive. To prepare specific electrolytes, 0.1 M of additional salts were introduced into the FN electrolyte. Various lithium polysulfide-containing catholytes were synthesized by stoichiometrically combining Li\u003csub\u003e2\u003c/sub\u003eS and S\u003csub\u003e8\u003c/sub\u003e (Sigma Aldrich). All electrolyte preparations were conducted in an argon-filled glovebox, maintaining H\u003csub\u003e2\u003c/sub\u003eO and O\u003csub\u003e2\u003c/sub\u003e levels below 1 ppm.\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003eCharacterizations of electrodes and electrolytes.\u003c/strong\u003e\u003csup\u003e\u0026nbsp;7\u003c/sup\u003e\u003cstrong\u003eLi NMR measurements\u003c/strong\u003e were performed using an AS400 spectrometer. To eliminate the influence of DMSO-d6 on the measurements, coaxial NMR tubes were employed. DMSO-d6 was contained in the inner tube, while the specific electrolytes were loaded into the outer tube.\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003eSEM imaging\u003c/strong\u003e and EDS mapping were conducted using a field-emission scanning electron microscope (JEOL, JSM-IT800). Li electrodes from Li|Ni half cells were collected for characterization under conditions analogous to Li stripping and plating in Li\u0026ndash;S full cells.\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003eCryo-TEM analysis\u003c/strong\u003e was carried out using a Glacios microscope (Thermo Fisher) to minimize atmospheric reactions and electron beam damage. Lithium was electrochemically deposited onto a TEM grid [200-mesh Nickel (Ni) with a lacey carbon film] at a current density of 1.0 mA cm\u003csup\u003e\u0026minus;\u003c/sup\u003e\u003csup\u003e2\u003c/sup\u003e and a capacity of 1.0 mAh cm\u003csup\u003e\u0026minus;\u003c/sup\u003e\u003csup\u003e2\u003c/sup\u003e. Ni grids were chosen instead of copper to prevent reactions between lithium polysulfides and the current collector\u003csup\u003e43\u003c/sup\u003e. After preparation, the sample was immediately placed in liquid nitrogen following transfer into a microtube.\u003c/p\u003e\n\u003cp\u003eRaman spectroscopy was performed using a PL spectrometer (Renishaw, inVia Qontor) with a 324 nm laser. Electrolyte samples were deposited onto a glass slide and covered with a quartz cover slip. The apertures were then sealed with epoxy resin to prevent unwanted reactions between the atmosphere and lithium polysulfides.\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003eOptical microscopy.\u0026nbsp;\u003c/strong\u003eOptical microscopy was conducted using optic microscope (Novel optics). To get images of various bulk electrolyte and catholytes, 10 \u0026micro;L bulk electrolyte was dropped on slide glass, then covered with cover-glass. The apertures were tightly sealed with epoxy resins.\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003eComputational calculations.\u003cstrong\u003e\u0026nbsp;DFT and MD calculations\u003c/strong\u003e\u003c/strong\u003e were performed using the Materials Studio package. The DMol\u0026sup3; module was employed for DFT calculations, utilizing the Generalized Gradient Approximation (GGA) with the Perdew-Burke-Ernzerhof (PBE) gradient-corrected exchange-correlation functional. Geometric optimizations were conducted on all molecules (DME, 2MF, FSI\u003csup\u003e\u0026minus;\u003c/sup\u003e, NO\u003csub\u003e3\u003c/sub\u003e\u003csup\u003e\u0026minus;\u003c/sup\u003e, S\u003csub\u003e4\u003c/sub\u003e\u003csup\u003e2\u003c/sup\u003e\u003csup\u003e\u0026minus;\u003c/sup\u003e, S\u003csub\u003e6\u003c/sub\u003e\u003csup\u003e2\u003c/sup\u003e\u003csup\u003e\u0026minus;\u003c/sup\u003e, and S\u003csub\u003e8\u003c/sub\u003e\u003csup\u003e2\u003c/sup\u003e\u003csup\u003e\u0026minus;\u003c/sup\u003e) prior to calculations, with the energy convergence criterion set to below 10-6\u0026nbsp;eV.\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003eMolecular dynamics (MD) simulations\u003c/strong\u003e were carried out using the Forcite module within the Materials Studio package. Periodic simulation cells were equilibrated through 1 ns of NPT (constant number of particles, pressure, and temperature) followed by NVT (constant number of particles, volume, and temperature) ensemble simulations. The Nos\u0026eacute; algorithm, with a Q ratio of 0.1, was used for thermostatting. After the 1 ns NVT equilibration, an additional 5 ns of NVT simulation was conducted to achieve full equilibrium. Snapshots of polysulfides (PS) and the overall coordination numbers were extracted from the final equilibrated systems.\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003eElectrochemical performance measurements and characterizations.\u003cstrong\u003e\u0026nbsp;Electrochemical measurements\u003c/strong\u003e\u003c/strong\u003e were performed using 2032-type coin cells and pouch cells. For coin cells, 150 \u0026micro;m thick Li metal was used as the anode, with 20 \u0026micro;L of various electrolytes for general testing. Polypropylene (PP) separators (Celgard) were used for all cell configurations. To measure \u003cstrong\u003eCE\u003c/strong\u003e, the modified Aurbach method was applied following established literature protocols: 1) Pre-formation cycles involved plating and stripping 5 mAh cm\u003csup\u003e\u0026minus;\u003c/sup\u003e\u003csup\u003e2\u003c/sup\u003e of Li at 0.5 mA cm\u003csup\u003e\u0026minus;\u003c/sup\u003e\u003csup\u003e2\u003c/sup\u003e. 2) Identical areal capacities of Li were deposited on Ni current collectors. 3) Ten cycles of Li stripping/plating were conducted at 1 mAh cm\u003csup\u003e\u0026minus;\u003c/sup\u003e\u003csup\u003e2\u003c/sup\u003e and 0.5 mA cm\u003csup\u003e\u0026minus;\u003c/sup\u003e\u003csup\u003e2\u003c/sup\u003e. 4) A final full Li stripping was performed to 1.0 V to remove the remaining Li reservoir. CE from pre-formation cycles was excluded to eliminate the effects of galvanic corrosion and Li consumption during SEI formation. Two PP separators were used to avoid shot-circuit during lithiation and de-lithiation.\u003c/p\u003e\n\u003cp\u003eFor observing Li\u003csub\u003e2\u003c/sub\u003eS deposition at Carbon paper (Wizmac), E/S ratio = 3 \u0026mu;L mg\u003csub\u003es\u003c/sub\u003e\u003csup\u003e\u0026minus;\u003c/sup\u003e\u003csup\u003e1\u003c/sup\u003e was targeted. Due to high dead-volume of coin cell configurations, slightly higher E/S ratio was intentionally used. The loading of the S on cathode was 3mg\u003csub\u003es\u003c/sub\u003e cm\u003csup\u003e-2\u003c/sup\u003e. Its fabrication was referred to our previous study\u003csup\u003e44\u003c/sup\u003e.\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003ePouch cells\u003c/strong\u003e were assembled in a dry room (dew point \u0026lt; \u0026minus;40 \u0026deg;C). The cathode, anode, and separator dimensions were 3.3 \u0026times; 5 cm\u003csup\u003e2\u003c/sup\u003e, 3.5 \u0026times; 5.2 cm\u003csup\u003e2\u003c/sup\u003e, and 3.7 \u0026times; 5.3 cm\u003csup\u003e2\u003c/sup\u003e, respectively. The electrolyte was dosed in accordance with the specified E/S ratio. Three double side-coated cathodes (=3-layer stacked pouch cells, 0.27Ah pouch cell) were used for various characterization and assessing rate performances, while 8-layer stacked pouch cells (8 Al current collectors with double-sided cathodes, 0.7Ah pouch cell) were employed for cycling tests. The areal sulfur loading of the cathodes was 2.17 \u0026ndash; 2.25 mg\u003csub\u003es\u003c/sub\u003e cm\u003csup\u003e-2\u003c/sup\u003e. The ratio of sulfur : carbon : binder was 72 : 27 : 1. To minimize contact resistance, pouch cells were compressed between two metal plates at a pressure of 1.5\u0026ndash;2.0 bar and operated within a voltage range of 1.80\u0026ndash;2.50 V. All cells were rested for 24 hours to achieve full electrolyte wetting. Before the targeted cycling rate, pre-conditioning cycles were conducted at 0.05C (1C = 1675 mAh g\u003csub\u003es\u003c/sub\u003e\u003csup\u003e\u0026minus;1\u003c/sup\u003e) and 0.1C two times respectively. For the pouch cell fabricated at E/S ratio = 1.5 \u0026micro;L mg\u003csup\u003e\u0026minus;1\u003c/sup\u003e, pre-cycle was performed at 0.025 C and the actual cycling test was conducted at 0.05 C. Galvanostatic cycling tests were performed at 25 \u0026deg;C using a WBCS3000L battery tester.\u003c/p\u003e"},{"header":"Declarations","content":"\u003cp\u003e\u003cstrong\u003eData availability\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003eAll data supporting the findings of this study are available within the paper and its Supplementary Information.\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003eAcknowledgements\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003eThis work was supported by the National Foundation of Korea (NRF) grant funded by the Korea government (MSIT) (RS-2023-00261543) and (RS-2005-25441256) and LG Energy Solution-KAIST Frontier Research Laboratory (2024). We acknowledge the support of LG Energy Solution in the development of the sulfur cathode electrode.\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003eAuthor contributions\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003eS. Kim proposed the research and conceived the idea. H. Kwon and H.-T. K. supervised the project. S. Kim performed the electrochemical measurements. S. Kim, I. Kim, and C. Paulina carried out and analyzed the cryo-TEM and SEM experiments. S. Kim, Y. Shin, and D. Kim conducted the DFT and MD simulations. S. Chung and C. Park assisted with pouch cell fabrication. J. Han, J. Kim and B. Cho contributed to optical microscopy analysis. S. Kim, H. Kwon, and H.-T. K. co-wrote and revised the manuscript.\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003eCompeting interests\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003eThe authors declare no competing interests.\u003c/p\u003e"},{"header":"References","content":"\u003col\u003e\n\u003cli\u003eManthiram, A., Fu, Y., Chung, S.-H., Zu, C. \u0026amp; Su, Y.-S. Rechargeable lithium\u0026ndash;sulfur batteries. \u003cem\u003eChemical reviews\u003c/em\u003e \u003cstrong\u003e114\u003c/strong\u003e, 11751-11787 (2014).\u003c/li\u003e\n\u003cli\u003eZhou, G., Chen, H. \u0026amp; Cui, Y. 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Here, we reveal that the emergence of a micron-scale morphological heterogeneity in the Li anode originates from the localized formation of a micron-scale lithium polysulfide (LiPS)-rich phase in the electrolyte. Notably, the chemical composition of the LiPS-rich phase contains minimal amounts of anions such as FSI− or NO3−, which are typically responsible for forming stable solid-electrolyte interphase (SEI); this deficiency obstructs the formation of Li2O-containing SEI layer, resulting in the localized dendritic Li morphology. We further demonstrate that fragmenting the LiPS-rich phase reduces its locality, suppresses anode heterogeneity, and significantly improves cycling stability. ","manuscriptTitle":"Localization of micron-scale lithium polysulfide-rich phases causes anode heterogeneity in lean-electrolyte lithium–sulfur batteries","msid":"","msnumber":"","nonDraftVersions":[{"code":1,"date":"2025-11-04 09:51:01","doi":"10.21203/rs.3.rs-7916426/v1","editorialEvents":[{"type":"communityComments","content":0}],"status":"published","journal":{"display":true,"email":"[email protected]","identity":"researchsquare","isNatureJournal":false,"hasQc":true,"allowDirectSubmit":true,"externalIdentity":"","sideBox":"","snPcode":"","submissionUrl":"/submission","title":"Research Square","twitterHandle":"researchsquare","acdcEnabled":true,"dfaEnabled":false,"editorialSystem":"","reportingPortfolio":"","inReviewEnabled":false,"inReviewRevisionsEnabled":true}}],"origin":"","ownerIdentity":"d615d6c6-8b92-4e0e-b02e-0974eb73ca0a","owner":[],"postedDate":"November 4th, 2025","published":true,"recentEditorialEvents":[],"rejectedJournal":[],"revision":"","amendment":"","status":"posted","subjectAreas":[{"id":57248724,"name":"Physical sciences/Energy science and technology/Energy storage/Batteries"},{"id":57248725,"name":"Physical sciences/Chemistry/Electrochemistry/Batteries"}],"tags":[],"updatedAt":"2025-11-24T05:05:48+00:00","versionOfRecord":[],"versionCreatedAt":"2025-11-04 09:51:01","video":"","vorDoi":"","vorDoiUrl":"","workflowStages":[]},"version":"v1","identity":"rs-7916426","journalConfig":"researchsquare"},"__N_SSP":true},"page":"/article/[identity]/[[...version]]","query":{"redirect":"/article/rs-7916426","identity":"rs-7916426","version":["v1"]},"buildId":"XKTyCvWXoU3ODBz1xrDgd","isFallback":false,"isExperimentalCompile":false,"dynamicIds":[84888],"gssp":true,"scriptLoader":[]}