Unveiling a Stable Polysulfide Transport Framework in a Fluorine‐Free Li‐S Batteries

TL;DR

A fluorine-free electrolyte for lithium-sulfur batteries improves performance by creating a stable framework that reduces polysulfide shuttling and enhances cycle life.

Contribution

A fluorine-free phosphate-based electrolyte enables a new redox mechanism in Li-S batteries, shifting to a Li3PS4-Li2S pathway with improved stability.

Findings

The new electrolyte achieves a stable capacity of 765 mAh g−1 over 200 cycles.

The Li-P-S framework suppresses polysulfide diffusion and promotes solid-state-like behavior.

Dynamic phase transitions in Li-P-S complexes drive reversible conversion and sustained performance.

Abstract

Lithium‐sulfur batteries are limited by severe polysulfide shuttling, unstable lithium interfaces, and parasitic redox reactions, which degrade cycle life and efficiency. These issues stem from the soluble nature of polysulfide intermediates and the poor interfacial control in conventional electrolytes. Herein, we report a fluorine‐free phosphate‐based electrolyte that fundamentally shifts the redox chemistry from the typical S8‐Li2S pathway to a reversible Li3PS4‐Li2S mechanism. This in situ transformation produces Li‐rich thiophosphate phases that suppress polysulfide diffusion, alleviate cathode expansion, and promote solid‐state‐like behavior. Spectroscopic and microscopic analyses revealed dynamic phase transitions within the Li‐P‐S complexes that drive reversible conversion and sustained performance. As a result, the cells delivered a stable capacity of 765 mAh g−1 over 200 cycles. This work presents a fluorine‐free electrolyte design that establishes a new class of solid‐state‐like Li‐S redox chemistry, offering a promising route toward durable, high‐performance lithium‐sulfur batteries.

This work unveils a lithium thiophosphate reaction pathway activated by trigger P2S5 units, which initiate fast self‐assembly and shift the system away from conventional Li‐S reactions. The reversible Li‐P‐S framework delivers stability, safety, and high energy density, while enabling a distinct solid‐liquid hybrid electrochemical mode that provides a viable route for next‐generation sulfur‐based energy storage.

Figures

Electrochemical performance of Li‐S cells with conventional LiTFSI and Li‐P‐S electrolytes. (a) Cycling performance of Li‐S cells with conventional LiTFSI electrolyte and Li‐P‐S electrolyte at 0.2C. (b) Voltage profiles of cells with LiTFSI and (c) Li‐P‐S electrolytes.

Interfacial kinetics and reaction pathway evolution in Li‐S cells using LiTFSI and Li‐P‐S electrolytes. (a) Reference cycling voltage profiles of LiTFSI and Li‐P‐S cells. (b) DRT maps during the 30th cycle charge process for LiTFSI and (c) Li‐P‐S cells. (d) DRT maps during the 30th cycle discharge process for LiTFSI and (e) Li‐P‐S cells.

XPS analysis of the sulfur cathode surface in Li‐P‐S cell. (a) S 2p and (b) P 2p spectra of the pristine sulfur cathode before cycling. (c) S 2p and (d) P 2p spectra of the sulfur cathode after 100 cycles.

Postcycling structural and compositional analysis of the sulfur cathode from the Li‐P‐S cell. (a) Cross‐sectional TEM image showing the dense, layered morphology of the cycled cathode on an Al current collector. (b) High‐resolution TEM image revealing lattice fringes with interplanar spacings of 5.12 and 3.95 Å, corresponding to the (210) and (211) planes of β‐Li3PS4, respectively. (c) SAED pattern indexed to crystalline β‐Li3PS4. (d) XRD pattern of the cathode after cycling, confirming the phase transition from elemental sulfur to β‐Li3PS4. (e) Raman spectrum of the activated cathode showing the presence of thiophosphate and polysulfide species. (f) Deconvoluted Raman peaks at ∼406 and ∼465 cm−1, assigned to the vibrational modes of P2S7 4− units and polysulfide chains, respectively.

Comparative electrochemical analysis of Li‐S cells using LiTFSI and Li‐P‐S electrolytes. (a) GITT profiles after 30 activation cycles for LiTFSI and (b) Li‐P‐S cells. (c) CV curves of both cells recorded after 30 cycles at a scan rate of 0.05 mV s−1. (d) Calculated Li+ diffusion coefficients derived from the GITT data in (a) and (b) for LiTFSI and (e) Li‐P‐S. (f) Tafel plots fitted from CV‐derived reduction peaks C1/C1# and (g) C2/C2#, and oxidation peaks (h) A1 and (i) A1#.

SEM images of the Li metal (left) and S cathode (right) surfaces after 100 cycles in the LiTFSI and Li‐P‐S cells. (a), (c), (e), and (g) correspond to the LiTFSI cell; (b), (d), (f), and (h) correspond to the Li‐P‐S cell.

Proposed interfacial reaction mechanism in the Li‐P‐S cell during initial and subsequent cycling. Schematic illustration of the in situ formation and evolution of Li3PS4/Li7P3S11 at the cathode‐electrolyte interface during the initial and repeated charge‐discharge cycles.

Full text

Introduction

1

Li metal batteries are considered to be the most promising electrochemical storage devices for high‐energy‐density storage because of the extremely high specific capacity (3840 mAh g^−1^) and the lowest reduction potential (−3.04 V vs. standard hydrogen electrode) [1]. When paired with sulfur‐based cathodes, the resulting lithium‐sulfur batteries (LSB) offer a high theoretical capacity (1672 mAh g^−1^), lightweight components, and low material cost. However, practical deployment is hindered by several intrinsic issues, including poor electrical and ionic conductivity, large volume changes during cycling, and dissolution of polysulfide intermediates (PS), which leads to active material loss and interface degradation. To achieve suitable ionic conductivity, highly conductive bis(trifluoromethanesulfonyl)imide (LiTFSI) salts are widely used in liquid‐electrolyte (LE) Li‐S batteries. While effective, LiTFSI is known to corrode aluminum current collectors [2] and raises environmental concerns [3, 4], particularly in large‐scale applications. More critically, LiTFSI‐based LEs promote the diffusion of polysulfides, which deposit as insulating Li_2_S on Li and significantly decrease cycling stability [5, 6, 7]. Because of the strong solubility of long‐chain polysulfides in such systems, current strategies cannot completely prevent sulfur transfer, especially under high sulfur loading, which is required for practical applications [8, 9].

Building on these long‐standing challenges, recent studies have shown that the solvation environment of LiTFSI‐based ethers fundamentally governs polysulfide dissolution, shuttle intensity, and interfacial stability, reinforcing that conventional LE systems intrinsically favor uncontrolled PS mobility [10]. In parallel, radical‐stabilization approaches demonstrate that electrolyte‐species coordination can dramatically alter the sulfur redox sequence, highlighting the critical role of electrolyte chemistry in directing reaction pathways [11]. These insights collectively suggest that modifying polysulfide species represents an effective lever for suppressing shuttle and reshaping the conversion mechanism. One practical implementation is the use of pre‐dissolved lithium polysulfides (Li_2_S_x_), which levels concentration gradients at the cathode‐electrolyte interface and simultaneously supplies Li^+^ to the system. Xu et al. [12] demonstrated that such configurations could deliver up to 1450 mAh g^−1^ for 50 cycles. However, dissolved species easily passivate the lithium surface and trigger parasitic reactions, resulting in long‐term instability [2]. Another related strategy uses P_2_S_5_‐Li_2_S additives to generate a protective layer on the lithium anode, thereby preventing undesirable interfacial reactions [13]. These Li‐P‐S complexes react with lithium in situ to form a sulfide‐based solid interface with a low Young's modulus (∼20 GPa), which passivates the anode and improves interfacial stability [13, 14, 15]. These thiophosphate additives appear to effectively moderate the reactivity of the lithium anode in LE systems. However, most studies do not mention the reaction of thiophosphate additives with the sulfur cathode. In particular, sulfide‐based solid‐state electrolytes usually have a relatively narrow electrochemical window [14, 16], causing them to decompose simultaneously at both the cathode and anode sides. Their redox products, which overlap with Li‐S potential, remain largely uncharacterized. Existing density functional theory (DFT) studies focus mainly on bulk solid frameworks (e.g., Li_10_GeP_2_S_12_, Li_7_P_3_S_11_, Li_3_PS_4_) [17, 18], offering limited insight into dynamic interfacial processes in liquid systems.

Several recent studies have described the reversible lithium thiophosphate decomposition process in all‐solid‐state LSBs. Since Li‐S and sulfide solid‐state electrolyte redox (S/S^2−^) reactions occur almost concurrently within the same potential window, many studies point out that part of the overall capacity originates from the contribution of the reversible sulfide electrolyte decomposition process after the initial cycle [19, 20, 21]. More recently, Wang et al., demonstrated that integrating mixed ionic‐electronic conductors (MIECs) into high‐sulfur‐content cathodes can extend electrochemically active sites beyond conventional three‐phase boundaries, thereby improving sulfur utilization [22]. While this conduction‐pathway broadening approach has shown considerable promise, it still relies on substantial conductive additive content and does not directly address the intrinsic sluggishness of the sulfur conversion reaction. In this context, the present study introduces a complementary strategy, where an in situ formed Li‐P‐S interphase participates directly in reversible redox and phase transition processes. This chemically driven activation not only promotes interfacial charge transfer and accelerates Li_2_S formation but also mitigates the accumulation of electrochemically inactive sulfur, offering an alternative and synergistic route to overcoming conversion‐reaction limitations in high‐loading Li‐S systems.

Herein, we report a novel Li‐P‐S‐based LE design wherein the cathode and anode are simple S/carbon‐black/binder and Li metal, respectively. Unlike conventional designs, this system enables the in situ transformation of elemental sulfur into electrochemically active thiophosphate phases through cycling, thereby reducing PS dissolution and alleviating interfacial stress at the solid‐liquid interface. Our comprehensive characterizations revealed the formation of a homogeneous interfacial layer composed of Li_3_PS_4_ and Li_7_P_3_S_11_ across both electrodes, improving the physical contact and lowering the resistance. Lithium symmetric cells assembled with the in situ generated Li‐P‐S layer maintained a low overpotential of 5 mV for over 1000 h. In coin cells with 4 mg cm^−2^ sulfur loading, the capacity stabilized at ∼720 mAh g^−1^ over the first 100 cycles, follow with a slight increase in the later stages, reaching an average capacity of around 740 mAh g^−1^ in 200 cycles period. The extended cycle results in Figure S1 show that after 400 cycles, the cell capacity has a slight decline again and eventually stabilizes at about 660 mAh g^−1^. These results demonstrate a unique self‐reorganization mechanism for sulfide species and provide a promising strategy for enhancing the long‐term stability of Li‐S batteries.

Results and Discussion

2

Figure 1 compares the cycling performances and voltage profiles of Li‐S batteries with our designed Li‐P‐S electrolyte and the conventional electrolyte (LiTFSI in DME‐DOL) at 0.2C in the potential range of 1.6–2.8 V. As shown in Figure 1a, both cells exhibited an activation behavior at the initial stage of cycling. The charging and discharging capacity gradually increased from 255 mAh g^−1^ at the first cycle to 485 mAh g^−1^ for LiTFSI cell and from 789 to 816 mAh g^−1^ for the Li‐P‐S cell, respectively. A significant capacity drop was observed in the LiTFSI cell starting from the 15th cycle, whereas the Li‐P‐S cell demonstrated stable capacity retention over 200 cycles. The voltage profiles in Figure 1b,c show that both cells exhibit two distinct plateaus in charging and discharging. In LiTFSI cells, as previous reports suggest [23], formation and dissolution of high‐order PS (Li_2_S_x_, where x = 4–8) at the higher voltage plateau (2.3–2.4 V), while at the lower voltage plateau (∼2.1 V), further reduction of long‐chain PS to Li_2_S_2_ and/or Li_2_S occurs. For the case of the LiTFSI cell, the shuttle effect can be inhibited in the early stage by the addition of LiNO_3_, which improves the coulombic efficiency at the initial cycles. However, a rapid capacity fade was observed once LiNO_3_ was depleted, leading to a pronounced capacity drop beginning at the 15th cycle (Figure 1a). This behavior indicates an internal short circuit. On the other hand, excellent capacity retention and a high sulfur utilization rate were achieved in Li‐P‐S cells. Figure 1b,c show the constant current (CC) charging/discharging mode for LiTFSI and Li‐P‐S cell, respectively. Under CC mode, the LiTFSI cell exhibits a gentle slope near the end of the discharge curve, which represents the sluggish Li_2_S_2_/Li_2_S solid‐state conversion. For Li‐P‐S cell, a sharp drop is observed at the end of discharge, which suggests the presence of a secondary phase transition. Moreover, Li‐P‐S cell has a higher reaction plateau in the entire potential range, indicating that the polarization barrier during the electrochemical reaction has been significantly improved. This observation indirectly suggests that the Li‐P‐S electrolyte exists not simply as a surface covering layer.

Electrochemical performance of Li‐S cells with conventional LiTFSI and Li‐P‐S electrolytes. (a) Cycling performance of Li‐S cells with conventional LiTFSI electrolyte and Li‐P‐S electrolyte at 0.2C. (b) Voltage profiles of cells with LiTFSI and (c) Li‐P‐S electrolytes.

To gain insight into Li‐ion transport behavior, electrochemical impedance spectroscopy (EIS) was analyzed using the distribution of relaxation times (DRT) method at the 30th cycle (Figure S2 and Note S1) for both cells. In Figure 2, the 2D mapping DRT spectra of the two cells differ significantly. The discontinuous low‐frequency response at 2.3–2.7 V during the charge (Figure 2b,c) is likely attributed to the solid‐liquid transition, where dissolved PS oxidized and deposited onto the electrode [24]. As shown in Figure  2c, the Li‐P‐S cell exhibits a distinct mid‐frequency response (∼100 Hz) during charge, absent in LiTFSI cells, indicating the emergence of an intermediate phase with unique ionic transport properties. This feature appears only above 2.5 V and intensifies with cycling (Figure S3), suggesting progressive formation of a self‐constructed Li‐conductive phase. Importantly, the 2D DRT mappings collected at the 60th, 90th, and 120th cycles (Figure S3) preserve the same characteristic mid‐frequency relaxation mode without the appearance of additional spectral branches. The persistence of this DRT signature suggests that the interfacial ion‐transport pathway reaches a stable configuration shortly after activation and remains mechanistically unaltered throughout extended cycling.

Interfacial kinetics and reaction pathway evolution in Li‐S cells using LiTFSI and Li‐P‐S electrolytes. (a) Reference cycling voltage profiles of LiTFSI and Li‐P‐S cells. (b) DRT maps during the 30th cycle charge process for LiTFSI and (c) Li‐P‐S cells. (d) DRT maps during the 30th cycle discharge process for LiTFSI and (e) Li‐P‐S cells.

By contrast, the discharge profiles of the two cells exhibit fundamentally different transport behaviors. In LiTFSI cells, once the discharge voltage drops below 2 V, the mid‐frequency response (100 Hz) quickly shifts to lower frequencies (1 Hz), followed by a steady, solid diffusion response at 0.1 Hz. This discrete shift is typically associated with the charge transfer process of the sulfur cathode and indicates the presence of two distinct Li storage stages, each characterized by different charge transfer kinetics [25, 26]. The subsequent formation of electrically insulating Li_2_S/Li_2_S_2_ toward the end of discharge further impedes charge transfer and alters the ion transport pathway [27]. Conversely, the Li‐P‐S cell exhibits a continuous propagation of the interfacial relaxation mode (∼10 Hz), reflecting a fundamentally different ion‐transport mechanism in which intermediate species remain confined rather than engaging in parasitic shuttling. Unlike the discrete phase transitions characteristic of conventional Li‐S reactions, the DRT analysis indicates that the Li‐P‐S configuration redirects the reaction pathway and stabilizes the final products, markedly accelerating the evolution toward steady‐state conduction [28, 29].

X‐ray Photoemission Spectroscopy analysis was performed on the cycled cathode of Li‐P‐S cell, focusing on the P 2p and S 2p regions. In Figure 3a, two doublets at 164.1 and 168.2 eV in the pristine cathode correspond to elemental S and SO_x_ [30]. After charging to 2.8 V, these signals disappeared, replaced by strong peaks at 161.2 and 163.5 eV assigned to PS_4_ ^3−^ and P_2_S_7_ ^4−^ species [31]. The P 2p spectrum (Figure  3d) confirms these species with corresponding doublets at 133.4 and 134.6 eV [32]. These results suggest that elemental S is converted into Li_3_PS_4_ or Li_7_P_3_S_11_ phase during the first charge, as lithium thiophosphates refill vacancies left by consumed sulfur. It is noted that the signal of S never reappeared in the XPS spectrum, even for cells fully recharged to over 2.8 V. Contrary to previous studies [33], the self‐constructed Li_3_PS_4_/Li_7_P_3_S_11_ does not decompose back to S and P_2_S_5_ in subsequent cycling. We attribute the contrasting behavior of lithium thiophosphates in solid vs. liquid systems to the high polysulfide concentration in Li‐P‐S cells. As reported by Rauh et al. [34], the elemental S spontaneously reduces to S^2−^ and integrates into PS clusters in solution. In such an environment, elemental sulfur cannot remain stable, explaining why Li_3_PS_4_/Li_7_P_3_S_11_ in Li‐P‐S cells do not oxidize back to elemental sulfur.

XPS analysis of the sulfur cathode surface in Li‐P‐S cell. (a) S 2p and (b) P 2p spectra of the pristine sulfur cathode before cycling. (c) S 2p and (d) P 2p spectra of the sulfur cathode after 100 cycles.

For comparison, XPS spectra of P 2p and S 2p for the Li counter electrode after activating cycle (Figure S4) showed features similar to those of the sulfur cathode, suggesting the formation of Li_3_PS_4_/Li_7_P_3_S_11_ on the Li surface. This suggests an additional function of our electrolyte in suppressing dendrite growth through the formation of structurally flexible and stable thiophosphate during cycling [35]. To further validate the stability of this interphase, long‐term plating/stripping tests were conducted using symmetric Li/Li cells with Li‐P‐S electrolyte. As shown in Figure S5a, the cell maintained a stable square‐wave voltage profile for over 1000 h under 1 mA cm^−2^ and 1 mA h cm^−2^. Only a slight increase in voltage hysteresis from 4 to 8 mV after 400 h (Figure S5b,c) was observed, attributed to the densification of the Li_3_PS_4_/Li_7_P_3_S_11_ interphase and reduced electrolyte exposure. On the other hand, the absence of runaway polarization or voltage drift over extended cycling suggests that parasitic polysulfide redox processes are effectively mitigated, further supporting that the thiophosphate interphase suppresses shuttle‐induced instability. In contrast, LiTFSI‐based symmetric cells exhibited a sharp overpotential growth, likely due to dendritic growth and unstable SEI formation, consistent with prior reports [36]. Nyquist plots before and after cycling (Figure S5d) further confirm the interfacial evolution and improved impedance stability induced by thiophosphate formation.

Ex situ XRD and high‐resolution TEM analyses were conducted for activated Li‐P‐S cells (Figure 4). TEM cross‐sections (Figure 4a) show that the cycled cathode develops a glassy matrix, and the macropores present in the pristine sulfur‐carbon layer disappear after activation. To further examine the structural evolution, cross‐sectional SEM of the cathode at different cycling stages (Figure S6) reveals a nonmonotonic thickness change. The pristine electrode exhibits a thickness of ∼50–60 µm with a highly porous morphology. After the 10th cycle, the composite layer swells to nearly 100 µm, which can be attributed to electrolyte uptake and the transient formation of Li_2_S_x_ and Li_3_PS_4_/Li_7_P_3_S_11_ species within the porous framework. With continued cycling to the 100th cycle, the overall thickness decreases to below 50 µm, and the structure becomes markedly denser, indicating the collapse of internal voids and consolidation of the thiophosphate‐rich network. The XRD patterns (Figure 4d) clearly show that the sulfur cathode has gone through a process of phase reorganization. The pristine sulfur cathode exhibited pure α‐sulfur peaks, but no sulfur‐related diffraction was found after activation, indicating that the elemental sulfur vanished after repeated cycling. To identify the short‐range order, a selected area electron diffraction (SAED) test was performed. The activated cathode presents an amorphous region mixed with crystalline micrograins (Figure 4b), and the SAED pattern (Figure 4c) confirmed to be β‐Li_3_PS_4_ phase. The difference between XRD and TEM can be reasonably attributed to the ultra‐fine granularity of Li_3_PS_4_ (d_grain_∼10 nm). The crystallographic evidence supplements the spectroscopic analysis mentioned earlier.

Postcycling structural and compositional analysis of the sulfur cathode from the Li‐P‐S cell. (a) Cross‐sectional TEM image showing the dense, layered morphology of the cycled cathode on an Al current collector. (b) High‐resolution TEM image revealing lattice fringes with interplanar spacings of 5.12 and 3.95 Å, corresponding to the (210) and (211) planes of β‐Li3PS4, respectively. (c) SAED pattern indexed to crystalline β‐Li3PS4. (d) XRD pattern of the cathode after cycling, confirming the phase transition from elemental sulfur to β‐Li3PS4. (e) Raman spectrum of the activated cathode showing the presence of thiophosphate and polysulfide species. (f) Deconvoluted Raman peaks at ∼406 and ∼465 cm−1, assigned to the vibrational modes of P2S7 4− units and polysulfide chains, respectively.

To further characterize the chemical structure, Raman scattering was performed on pristine and fully activated cathodes (Figure 4e,f). For the pristine cathode, the Raman spectra were in good agreement with XRD pattern. The intense peak at 470 cm^−1^ with a minor shoulder at 465 cm^−1^ represents the S‐S bond stretching of the S_8_ ring structure, and the broader 439 cm^−1^ band corresponds to S‐S stretching modes [37]. After activation, the top curve in Figure 4e displays two distinct peak sets, one corresponding to thiophosphate polyhedral species centered at 408 cm^−1^ [38] and the other to PS species distributed across 450–500 cm^−1^ region [39]. Deconvolution (Figure  4f) reveals PS_4_ ^3−^ and P_2_S_7_ ^4−^ at 420 and 407 cm^−1^ [40], and overlapping short‐chain, long‐chain, and radical PS signals at 453, 465, and 495 cm^−1^ [41]. These findings are consistent with the XPS results, confirming that elemental sulfur is replaced by Li_3_PS_4_/Li_7_P_3_S_11_ in high PS concentration cells. Importantly, Raman spectra provide evidence of the coexistence of PS and Li_3_PS_4_/Li_7_P_3_S_11_, which confirms the hypothesis that lithium thiophosphate can be stable in high‐concentration PS cells.

Up to this point, our observations demonstrate that during the initial charge, Li_2_S is electrochemically oxidized and simultaneously reacts with the surrounding mP_2_S_5_‐nLi_2_S_x_ glass‐ceramic electrolyte to form Li_3_PS_4_ and Li_7_P_3_S_11_ like environments at the Li_2_S/electrolyte interface. Due to the limited mass transfer distance within the cell, it is presumed that this interfacial reaction only involves a narrow region around the electrode‐electrolyte interface. Furthermore, because the reaction is highly reversible, the overall electrolyte concentration of the system does not have a significant change within a few cycles. Structurally, this interphase formation can be viewed as a local reorganization of the thiophosphate network. The original glass‐ceramic matrix contains mixed coordination environments, including cross‐linked units such as P_2_S_7_ ^4−^ or P_2_S_6_ ^4−^, together with PS_4_ ^3−^ tetrahedra. After the first charge, S consumption and the interfacial reaction increase the fraction of PS_4_ ^3−^ units and Li‐rich thiophosphate species characteristic of Li_3_PS_4_ and Li_7_P_3_S_11_. This tendency is consistent with the stronger spectral features assigned to PS_4_ ^3−^ in the Raman/XPS results after activation (Figure 3 and Figure 4f), as well as the weak reflections corresponding to Li_3_PS_4_/Li_7_P_3_S_11_ observed in the ex situ TEM and XRD patterns (Figure 4c,d). In other words, the rate of this phase transition reaction may be discontinuous. In the early stages of the cycle, the initially present elemental sulfur is rapidly replaced by Li‐P‐S substances and enters the LiPS‐Li_2_S reaction cycle. While after the early cycles, this phase transition reaction will tend to reach a stable dynamic equilibrium. The ionic transport behavior observed in the in situ DRT further corroborates this interfacial reconstruction. Once the Li‐rich thiophosphate layer is formed, the 2D DRT mappings collected at the 30th, 60th, 90th, and 120th cycles preserve the same characteristic mid‐frequency relaxation mode without the appearance of additional features. The persistence of this spectral signature indicates that the interfacial ion‐transport pathway rapidly reaches a stable configuration and remains mechanistically unaltered during extended cycling. Such cycle‐invariant relaxation behavior is consistent with a self‐limited evolution of the Li_3_PS_4_/Li_7_P_3_S_11_ interphase, rather than a continuously evolving interface. Beyond kinetic and structural stabilization, this interfacial reconstruction also has profound implications for cell safety [42]. Prior studies have shown that soluble high‐order polysulfides (e.g., Li_2_S_6_/Li_2_S_8_) serve as major thermal initiators in Li‐S systems, driving parasitic redox at the Li surface and accelerating cathode‐originated runaway reactions [43, 44]. In the Li‐P‐S configuration, sulfur species are not cycled through a soluble reservoir but are immobilized within inorganic thiophosphate domains, eliminating the shuttle‐mediated pathways required to trigger these exothermic reactions. Mechanistically, the glassy Li_3_PS_4_/Li_7_P_3_S_11_ interphase functions analogously to thermoresponsive protective layers proposed for safe lithium‐metal operation, enabling uniform Li deposition and preventing current‐density amplification that typically leads to dendrite growth [45, 46, 47]. This localized compositional change is beneficial for both ion transport and cycling stability. Li_3_PS_4_ and Li_7_P_3_S_11_ are known to exhibit high Li^+^ conductivity; therefore, the formation of a Li‐rich thiophosphate interphase lowers the interfacial resistance and facilitates Li^+^ exchange between Li_2_S and the bulk electrolyte, which further improved rate performance of the Li‐P‐S cell (in comparison with the LiTFSI cell), as shown in Figure S7. Consistently, the interfacial resistance extracted from EIS decreases after the first few cycles and then stabilizes, indicating that the reconstructed interphase rapidly reaches a new equilibrium and enables a reversible Li_2_S conversion process.

To gain more insight into the kinetic behaviors of both cells, the galvanostatic intermittent titration technique (GITT) analysis was employed. As shown in Figure 5a,b, both LiTFSI and Li‐P‐S cells were adopted for the GITT test after 30 activation cycles (0.25 C rate for 5 min, followed by 60 min rest). The LiTFSI cell displayed classic Li‐S redox features but retained only ∼120 mAh g^−1^ in GITT cycles, which is about one‐half of its capacity in galvanostatic cycles ∼250 mAh g^−1^ (Figure 1c). This clearly demonstrates the occurrence of polysulfide shuttling during the long relaxation period, which ultimately leads to severe self‐discharge of common Li‐S cells. As a comparison, the Li‐P‐S cell exhibited well‐defined voltage plateaus, higher overall capacity, and significantly reduced relaxation‐induced polarization. These features indicate improved reaction kinetics and more efficient sulfur utilization, as reflected in smoother plateau evolution and lower Li‐ion diffusion polarization. The absence of severe relaxation decay further suggests suppressed intermediate shuttling and a more direct conversion pathway during charge‐discharge. To clarify the relationship between redox reactions at steady state, cyclic voltammetry (CV) tests were performed (also after 30 cycles) on the LiTFSI and Li‐P‐S cell and compared in Figure 5c. For the LiTFSI cell, starting from the open‐circuit voltage (OCV), a positive scan yielded one oxidative peak (A1) accompanied by a high‐voltage shoulder (A2) at 2.39 and 2.42 V, while the negative scanning showed two reductive peaks (C1, C2) at 2.32 and 2.01 V, respectively.

Comparative electrochemical analysis of Li‐S cells using LiTFSI and Li‐P‐S electrolytes. (a) GITT profiles after 30 activation cycles for LiTFSI and (b) Li‐P‐S cells. (c) CV curves of both cells recorded after 30 cycles at a scan rate of 0.05 mV s−1. (d) Calculated Li+ diffusion coefficients derived from the GITT data in (a) and (b) for LiTFSI and (e) Li‐P‐S. (f) Tafel plots fitted from CV‐derived reduction peaks C1/C1# and (g) C2/C2#, and oxidation peaks (h) A1 and (i) A1#.

The two‐step reduction pair corresponds to the double plateau in the GITT discharge curve, as shown in Figure 5a. The smaller C1 peak indicates the S to soluble PS [Li_2_S_n_ (4 ≤ n ≤ 8)] reduction, corresponding to the discharge plateau at 2.2–2.4 V and steep slope in the 2.1–2.2 V range in the GITT test. The C2 cathodic peak represents a process of short‐chain PS to lithium sulfide (either Li_2_S_2_ or Li_2_S) conversion, contributing two‐thirds of overall cell capacity (Figure S8) [33]. When it comes to Li‐P‐S cell, results show the cathodic peaks in CV plot significantly shift to lower potential (C1^#^: 2.26 V, C2^#^: 1.91 V) and anodic peaks (A1^#^: 2.40 V, A2^#^: 2.59 V) to higher potentials. Except for the peak location. The GITT and CV plots present an almost equal capacity distribution for each reduction step (Figure S8). This behavior is similar to profiles associated with sulfide electrolyte decomposition [21, 41, 48]. According to the computational predictions, complete reduction of Li_7_P_3_S_11_ yields Li_3_P and Li_2_S_2_ [16, 17]. This redox pathway produces a multi‐step profile closely matching the CV and GITT features observed in our Li‐P‐S cell. Tafel plots (Figure S9) further reveal higher exchange current densities in the Li‐P‐S system (8.64 and 3.82 mA cm^−2^ for cathodic and anodic processes, respectively), indicative of faster charge‐transfer kinetics and reduced interfacial resistance. Although the apparent Li⁺ diffusion coefficient (*D_Li+_ *) obtained from GITT (Figure 5g,h; Note S2) is lower in the Li‐P‐S cell, this reflects its distinct reaction mechanism rather than slower kinetics. In LiTFSI electrolyte, sulfur conversion proceeds through dissolved Li_2_S_x_ intermediates and is governed by bulk diffusion. By contrast, in the Li‐P‐S system, sulfur is immobilized within a Li_3_PS_4_/Li_7_P_3_S_11_‐rich interphase, causing the reaction to occur primarily at a solid–solid interface and suppressing long‐range polysulfide transport. Under such interfacial‐dominated conditions, GITT captures only local Li⁺ redistribution and therefore yields a lower apparent *D_Li+_

• [49, 50, 51]. This behavior aligns with the self‐healing concept reported by Xu et al. [12], where stabilization of polysulfides near the cathode surface diminishes concentration gradients and reduces the apparent diffusion coefficient despite faster overall kinetics. A similar stabilization arises in the Li‐P‐S electrolyte: the thiophosphate‐rich interphase regulates polysulfide speciation and minimizes bulk mass transport. Consequently, the lower *D_Li+_
• values are consistent with surface‐controlled conversion kinetics and correlate well with the smoother discharge plateaus, reduced relaxation polarization, and cycle‐invariant DRT features. Nevertheless, the significantly enhanced redox reversibility and charge‐transfer kinetics, supported by Tafel and cycling data, confirm that ionic diffusion is not rate‐limiting. Instead, the system operates via a surface‐controlled conversion mechanism governed by interfacial processes. Based on the galvanostatic and the CV results, it is reasonable to assume that the original Li‐S nature is replaced by the Li_7_P_3_S_11_ redox reaction in Li‐P‐S cell.

To further quantify the kinetic characteristics of each redox step, a separate set of Tafel analyses was performed to extract apparent activation energies from the linear regions of log(i)‐V plots (Figure 5f–i; see Note S3 and Figure S10 for details). These values reflect the energy barriers for interfacial charge‐transfer reactions. Notably, the lower activation energies observed in the Li‐P‐S cell, particularly during the late‐discharge and early‐charge stages, are consistent with the persistent mid‐frequency features observed in the DRT spectra. This correlation suggests that the redox behavior is governed by a modified interfacial reaction pathway facilitated by the Li‐P‐S electrolyte, enabling more accessible and reversible charge transfer.

As shown in Figure 6, scanning electron microscopy (SEM) observations reveal distinct electrode surface morphologies. For the Li‐P‐S cell (Figure 6b,d,f,h), the electrode surface is uniformly coated with a thin film, and the interfacial boundary is observed to advance progressively with continued galvanostatic cycling. As discussed in previous sections, this self‐constructing interface plays a dominant role in shaping the electrochemical behavior of the system. Specifically, surface defects and electrochemically active sites appear to be passivated by the formation of Li_3_PS_4_/Li_7_P_3_S_11_, thereby shielding the electrode from direct attack by reactive polysulfide species. In contrast, the electrode from the LiTFSI‐based cell (Figure 6a,c,e,g) exhibits rough, fragmented surfaces, indicating structural degradation and correlating with voltage fluctuations, possibly due to internal short circuits (Figure S11).

SEM images of the Li metal (left) and S cathode (right) surfaces after 100 cycles in the LiTFSI and Li‐P‐S cells. (a), (c), (e), and (g) correspond to the LiTFSI cell; (b), (d), (f), and (h) correspond to the Li‐P‐S cell.

Conclusions

3

Combining all the evidence, the results of spectroscopy, crystallography, and real‐time electrochemistry, all point to a conclusion different from most of the Li‐S cells reported, which actually inspired the idea explored in this work. Based on our observations, herein, we propose a process of an in situ conversion from pristine sulfur to glassy/or polycrystalline Li_3_PS_4_/Li_7_P_3_S_11_ phase, supported by the crystallographic and Raman analysis. As illustrated in Figure 7, by assembling an S‐C composite cathode into a Li‐P‐S electrolyte, sulfur is spontaneously reduced by the Lewis bases of PS and Li‐P‐S clusters, generating Li_2_S or lithium thiophosphate on the electrode surface. The subsequent first charging oxidizes Li_2_S and lithium thiophosphate into metastable Li_3_PS_4_/Li_7_P_3_S_11_, enabling the cell into the subsequent cycling route.

Proposed interfacial reaction mechanism in the Li‐P‐S cell during initial and subsequent cycling. Schematic illustration of the in situ formation and evolution of Li3PS4/Li7P3S11 at the cathode‐electrolyte interface during the initial and repeated charge‐discharge cycles.

In summary, this study developed a fluorine‐free Li‐P‐S LE system and explored how the coordination of phosphorus pentasulfide and PS replaces the well‐known Li‐S nature. More than just a functional interface, the Li‐P‐S plating layer is to provide a stable capacity source for the entire electrochemical system and even replace the original Li‐S redox reaction after multiple cycles. A novel modified redox pathway shifts the traditional S_8_‐Li_2_S reaction toward reversible Li_3_PS_4_‐Li_2_S conversion, enabling controllable thiophosphate phase transformation and providing a powerful solution to the persistent challenges of polysulfide shuttling and lithium dendrite formation. Ultimately, this study not only paves the way for the development of high‐energy and long‐lasting Li‐S systems but also providing a new sight into the idea for the reversibility of sulfide solid electrolytes.

Experimental Section/Methods

4

Preparation of Lithium Polysulfides

4.1

Lithium polysulfides, Li_2_S_x_ (5 ≤ x ≤ 8), were prepared by stirring stoichiometric Li_2_S and S powders (>99% trace metals basis, Sigma‐Aldrich) in the mixtures of DME (≥ 98%, Sigma‐Aldrich), DOL (≥ 99%, Sigma‐Aldrich) solvent in 1:1 (v/v) ratio at 70°C for 12 h in an Argon‐filled glovebox (O_2_< 0.2 ppm, H_2_O < 0.1 ppm, vigor USA)

Preparation of Lithium Thiophosphate Complexes

4.2

To prepare lithium thiophosphate (mP_2_S_5_‐nLi_2_S_x_) complexes in DME and DOL mixed solution, a controlled amount of P_2_S_5_ (≥99%, Sigma–Aldrich) was added to the as‐prepared Li_2_S_x_ solutions, such that the mixture fell within the phase‐stable molar ratio window of 1:1 to 1:3. The solution was then stirred at 70°C for 5 h. The characteristic light yellow to dark red appearance of the complexes was observed after stirring.

Preparation of LiPS‐Based Electrolyte

4.3

mP_2_S_5_‐nLi_2_S_x_ complexes, and/or, Li_2_S_x_ solution were added into DME, DOL solvent (DME:DOL = 1:1 (v, v)) as electrolytes. The P_2_S_5_ concentrations were kept at 0.75 M in all electrolyte solutions as the baseline. Lithium thiophosphate complexes in the mixture of DOL with other glyme‐based solvents can be prepared in a similar method.

Preparation of Blank Electrolyte Solutions

4.4

1.0 M LiTFSI (99.95% trace metals basis, Sigma‐Aldrich), and 3 wt.% LiNO_3_ (99.99% trace metals basis, Sigma‐Aldrich) powder were dissolved into the as‐prepared DME and DOL mixture solution (1:1 (v/v)) and stirred at room temperature in a glove box for 3 h.

Sulfur Cathode Preparation

4.5

To fabricate the sulfur cathode, 70 wt.% of Sulfur powder (≥99.0%, Sigma‐Aldrich) was dry mixed with 20 wt.% of super P carbon black (Denka, 50% compressed) and 10 wt.% of pectin (pectin from apple, Sigma‐Aldrich) powder, followed by dispersing in DI water of 1700 rpm at ambient environment for slurry fabricating. After 3 h stirring, and process by casting, drying, and rolling, the cathode pieces were punched into 14 mm discs and ready for coin cell assembling.

Coin Cell Assembly

4.6

CR2032‐type coin cells (Ubiq Co., Taiwan) were assembled in an argon‐filled glove box (Vigor, USA, H_2_O/O_2_< 0.1 ppm). The sulfur cathode discs were pre‐dried at 120°C under vacuum for 12 h prior to cell construction. Lithium‐metal foils (Ø15.8 mm) and separators (Asahi, Japan) were pre‐wetted with electrolyte (∼25 µL absorbed) to ensure uniform ionic contact. Cells using the LiTFSI‐based electrolyte are referred to as LiTFSI cells, whereas cells using the polysulfide‐rich electrolyte are denoted as Li‐P‐S cells. The electrolyte dosage was standardized based on the electrolyte‐to‐sulfur ratio (E/S). Both LiTFSI and Li‐P‐S cells were assembled using 8.5 µL of electrolyte per 1 mg of sulfur (E/S = 8.5 µL mg^−1^), ensuring a direct comparison without electrolyte starvation or flooding. The areal sulfur loading of the cathode was maintained between 4.2 and 4.5 mg cm^−2^.

Material Characterization

4.7

The structure of the cathode was characterized by powder X‐ray diffraction (XRD, Rigaku, Japan) with Cu Kα radiation in the 2θ range of 10°–80°. The morphology of the samples was observed by a field emission analysis electron microscope (FEI Inspect‐F SEM). The surface state of the lithium anode and sulfur cathode, which after cycling was determined by hard X‐ray photoelectron spectroscopy XPS using ULVAC‐PHI (Quantes) with Al Kα radiation, and data are averaged from ten cells. Microstructure and phase determination were performed using a high‐resolution transmission electron microscope (TEM, JEOL, JEM‐2100F). To protect the surface layer from damage during sample preparation and later observation, the ex‐situ specimens were sputtered with a 10 nm platinum layer by electron beam and subsequently cut by SEIKO SMI3050SE Dual‐Beam focused ion‐beam milling before the TEM test.

Author Contributions

F.Y.Wu, P.W. Chi, P.M. Wu, and M.K. Wu conceptualized the study, developed the research framework, and conducted the final review of the manuscript. F.Y. Wu and P.W. Chi performed the experiments, and all authors analyzed the data. The manuscript and figures were prepared collaboratively by F.Y. Wu, P.W. Chi, P.M. Wu, and M.K. Wu. All authors have reviewed and revised the manuscript thoroughly.

Conflicts of Interest

The authors declare no conflicts of interest.

Supporting information

Supporting File: advs73584‐sup‐0001‐SuppMat.docx.