Polarizability-enhanced ionic transport in rare-earth-free halide–sulfide electrolytes: Li₂ZrSCl_4−xBr_x

Abstract

Commercially viable all-solid-state batteries (ASSBs) rely on solid electrolytes (SEs) that combine high ionic conductivity, electrochemical stability, low cost, and scalable production. Here, we report a series of rare-earth-free solid electrolytes, Li₂ZrSCl_4−xBr_x (0 ≤ x ≤ 4), synthesized via a rapid, energy-efficient mechanochemical route. The optimized composition, Li₂ZrSCl_1.3Br_2.7, exhibits an ionic conductivity of ∼1.03 mS cm⁻¹, one order of magnitude higher than Li₂ZrCl₆. Structural and morphological analyses using XRD, SEM/EDS, and ⁶Li MAS NMR reveal that progressive Br⁻ substitution drives structural disorder, enhances anion polarizability, and yields dynamically disordered Li⁺ environments conducive to rapid ion migration. Compared to the semi-crystalline Li₂ZrSCl₄ and fully brominated Li₂ZrSBr₄ end members, Li₂ZrSCl_1.3Br_2.7 achieves an optimal structural disorder and Li⁺ ion mobility, resulting in enhanced ionic conductivity. When used as a catholyte in an ASSB with TiS₂ as the cathode active material, Li₂ZrSCl_1.3Br_2.7 exhibits good rate capability and stable long-term cycling performance. This work highlights the viability of Li₂ZrSCl_4−xBr_x as a high-performance and inexpensive solid electrolyte, combining fast Li⁺-ion transport, electrochemical stability, and scalable synthesis, making it a promising candidate for commercial ASSBs.

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Introduction

The advancement of solid-state batteries (SSBs) hinges on developing solid electrolytes (SEs) that combine high ionic conductivity, broad electrochemical stability, and cost-effectiveness.^1,2 Traditional liquid electrolytes face inherent drawbacks such as flammability, leakage, and limited electrochemical stability, which hinder the advancement of safer and more efficient batteries.^3,4 Halides are a promising class of SEs for the next-generation ASSBs because of their compatibility with high-voltage cathode materials.^5,6 Despite the advantages of halide SEs, achieving high Li⁺-ion conductivities-up to ∼10⁻³ S cm⁻¹ – has largely relied on the use of expensive transition and rare-earth trivalent elements, such as in materials like Li₃YCl₆, Li₃ErCl₆, Li₃HoCl₆, and Li₃InCl_6.^7–11 More recently, pentavalent cations such as Ta⁵⁺ and Nb⁵⁺ have been introduced in compounds like LiTaOCl₄, LiNbOCl₄, NaTaOCl₄, further expanding the compositional landscape of high-performing but costly halide SEs.^12–14 While these materials often exhibit favorable electrochemical properties, the use of rare-earth metals introduces challenges related to high cost, limited global supply, and supply-chain concentration, which can hinder large-scale manufacturing and long-term commercialization of solid-state batteries. Consequently, the development of rare-earth-free halide electrolytes based on earth-abundant elements is an important strategy for improving material sustainability and economic viability.¹⁵ In contrast, Li₂ZrCl₆, a rare-earth-free alternative, has garnered increasing attention owing to its lower cost and encouraging electrochemical properties.^16–18 However, its room temperature ionic conductivity still falls short for practical solid-state batteries,^19,20 highlighting the need for further enhancements.

Recent efforts to enhance the performance of Zr-based halide SEs have focused on compositional and structural modifications. These approaches include cationic substitutions at the Zr⁴⁺ site in Li₂ZrCl₆ – for example, Li_2+xZr_1−xIn_xCl₆,²¹ and Li_2.4Zr_0.8Zn_0.2Cl₆,²² – which have been shown to improve moisture stability. Similarly, anionic substitutions, such as replacing Cl⁻ with F⁻ in Li_2.3ZrCl_6.1F_0.2,²³ have been employed to widen the electrochemical stability window. More recently, researchers have explored dual-anion systems that retain the beneficial features of halides, such as high polarizability and structural softness, while integrating anions like oxide or sulfide to tailor the local coordination environment and facilitate Li⁺ transport.^24–26 Notably, superionic conductivities around ∼10⁻³ S cm⁻¹ have been observed for Li₃ZrCl₄O_1.5,²⁷ Li_3.1ZrCl_4.9O_1.1,²⁸ Li₃Zr_0.75OCl₄,²⁹ and ZrO₂–2LiCl–Li₂ZrCl₆.³⁰ Additionally, Xu and co-workers reported a Li⁺ conductivity of 0.86 mS cm⁻¹ for Li_2.2ZrCl_5.8S_0.2.³¹ However, these materials require long milling durations, LiCl precursor, and/or secondary heat treatments, posing challenges for industrial scalability. Motivated by these advancements and limitations, we sought to further expand the compositional space by developing a new series of rare-earth-free solid electrolytes, Li₂ZrSCl_4−xBr_x (0 ≤ x ≤ 4). This system leverages a diversified anion framework and mixed halide composition, leading to improvements in ionic conductivity from scalable synthesis.

In this work, we report the synthesis of the Li₂ZrSCl_4−xBr_x composition by directly combining Li₂S and ZrX₄ (X = Cl or Br), avoiding the commonly used LiCl precursor. This approach shortens the milling time, as pre-formed LiCl is highly crystalline and less reactive. The resulting Li₂ZrSCl₄ (LZSC), similar to recently reported S²⁻ doped Li₂ZrCl₆ (LZC),³² exhibits crystalline characteristics, with observed P3̄m1 and C2/m phases, showing a modest conductivity improvement from ∼0.1 in LZC to ∼0.45 mS cm⁻¹ in LZSC, accompanied by a substantial reduction in activation energy from ∼0.41 to ∼0.37 eV, highlighting sulfur's role in softening the anion framework and lowering the Li⁺ migration barriers. Further Br⁻ substitution promotes amorphization and increases lattice softness through enhanced anion polarizability, yielding the optimized composition Li₂ZrSCl_1.3Br_2.7 (LZSCB), which reaches ∼1.03 mS cm⁻¹ at room temperature. While the activation energy remains nearly constant across the series (∼0.36–0.37 eV), the conductivity enhancement arises from an increased pre-exponential factor associated with fast Li⁺ ion dynamics, whereas the fully Br-substituted end member, Li₂ZrSBr₄ (LZSB), shows reduced conductivity. Importantly, the amorphous Li₂ZrSCl_1.3Br_2.7 material can be synthesized through a single-step 2-hour mechanochemical process, offering a practical and scalable route for industrial production. To gain deeper insight into the structure, composition, and surface morphology, a suite of characterization techniques – including NMR, XRD, SEM, EDS, and Raman spectroscopy – was utilized. As a catholyte in an ASSB comprising TiS₂ as cathode active material, Li₆PS₅Cl separator, and Li/In alloy anode, LZSCB demonstrates a good room-temprature cycling performance with 94% capacity retention (248 mAh g⁻¹) over 70 cycles at 0.2C. Our findings highlight LZSCB with remarkable conductivity, stability, and ease of production with low cost, positioning it for applications in commercially viable ASSB technologies.

Materials & methods

Synthesis

Lithium sulfide (99.5%, Alfa Aesar), ZrCl₄ (99.9%, Alfa Aesar), and ZrBr₄ (99.9%, Alfa Aesar) were stored in an argon-filled glovebox (VTI) with moisture and O₂ levels below 0.1 ppm. Subsequently, a stoichiometric amount of the starting precursors was manually ground for 15 min to achieve a homogeneous powder. The ground powder was subsequently loaded into a 20 mL zirconia milling jar with three 10 mm-sized zirconia balls, and vacuum sealed within the argon-filled glovebox. The sealed jar was placed in a High-Energy Ball Mill (8000 M Mixer/Mill®) and milled for 2 h. The processed sample was thereafter stored in the argon-filled glovebox for further characterization.

Scanning electron microscopy (SEM) and energy dispersive X-ray spectroscopy (EDS)

The samples were mounted on aluminum stubs using carbon tape to ensure electrical conductivity during SEM analysis. SEM imaging was conducted using a JEOL JSM-7000F microscope operated at an accelerating voltage of 15 kV. Prior to imaging, samples were sputter-coated with a thin layer of iridium to minimize charging effects. EDS analysis was performed in conjunction with SEM to determine the elemental composition of the samples. EDS spectra were collected and analyzed using an Oxford Instruments X-Max detector and Aztec software to quantify the elemental distribution across different regions of the samples. The working distance was maintained at 10 mm to optimize both imaging resolution and EDS signal detection.

Solid-state NMR

⁶Li and ⁷Li solid-state magic-angle-spinning (MAS) NMR spectra were acquired using a Bruker Avance III system equipped with a 500 MHz (11.74 T) ultrashield wide-bore magnet (89 mm). The Larmor frequencies were 73.6 MHz for ⁶Li and 194.4 MHz for ⁷Li. Samples were sealed in 2.5 mm zirconia rotors under an argon atmosphere and measured under an MAS rate of 24 kHz. Single-pulse ⁶Li and ⁷Li NMR experiments were performed with a π/2 pulse length of 3.3 µs and recycle delays of 200 s and 20 s, respectively. To study the ion dynamics in Li₂ZrSCl_4−xBr_x, the inversion-recovery approach was employed to measure ⁷Li T₁ NMR relaxation times at variable temperatures. The T₁ experiments were performed on a 300 MHz spectrometer operating at the ⁷Li Larmor frequency of ∼116 MHz, with chemical shift calibration against LiCl_(s) at −1.1 ppm. ³⁵Cl and ⁷⁹Br MAS NMR experiments were performed using an 800-MHz Bruker spectrometer with a field strength of 19.6T. The Larmor frequencies for ³⁵Cl and ⁷⁹Br were 208.3 MHz and 40 MHz, respectively. A single-pulse sequence was used for both ³⁵Cl and ⁷⁹Br experiments with a π/2 pulse length of 5.6 µs and 2.2 µs each and a recycle delay of 10 s and 5 s, respectively. Sample powders were packed under argon into 3.2 mm ZrO₂ pencil rotors and spun at 16 or 17 kHz. ³⁵Cl experiments were calibrated to a 0.1 M NaCl in D₂O solution at 0 ppm. ⁷⁹Br NMR experiments were calibrated to solid KBr at 54.5 ppm.

Raman spectroscopy

Raman spectra were collected on a Horiba JY LabRam HR Evolution spectrometer using a 532 nm excitation laser and a 600 g per mm grating. Spectra were acquired over the 100–1000 cm⁻¹ range with a laser power of 20 mW and an integration time of 4 seconds. The resulting data were processed using LabSpec 6 software.

Electrochemical impedance spectroscopy (EIS)

The synthesized sample was hand-ground and pressed at 300 MPa in an 8 mm diameter PEEK mold to make a ∼0.8 mm thick pellet for AC impedance measurements. The pellet was assembled in a split cell with steel as the blocking electrode. The electrochemical impedance measurements were carried out using a Gamry electrochemical analyzer, spanning frequencies from 5 MHz to 1 Hz, with an applied voltage of 10 mV. Ionic conductivities were calculated from the impedance values extracted from the Nyquist plots and pellet dimensions. A CSZ Microclimate chamber was used to perform impedance measurements across a range of temperatures from −20 to 60 °C. Electronic conductivity was evaluated using DC polarization in a symmetric steel|Li₂ZrSCl_4−xBr_x|steel configuration. Steady-state currents were recorded at applied voltages of 100, 200, 300, and 400 mV.

Lab X-ray diffraction

The sample was sealed on a zero-background holder using Kapton film and analyzed using a Rigaku SmartLab X-ray diffractometer equipped with a Cu Kα radiation source (λ = 0.154 nm) in Bragg–Brentano geometry. Diffraction data were collected over a 2θ range of 10° to 70°, with a step size of 0.03° and a scan rate of 1.5° min⁻¹.

Linear sweep voltammetry (LSV) measurements

Approximately 150 mg of Li₂ZrSCl_4−xBr_x was weighed into a 10 mm diameter polyether ether ketone (PEEK) cylinder and pressed at 250 MPa for 1 min. A composite with a mass ratio of 9 : 1 of Li₂ZrSCl_4−xBr_x to carbon nanofiber was hand-ground in an agate mortar for 15 minutes. An 8-mg portion of this composite was spread over one side of the Li₂ZrSCl_4−xBr_x pellet, contained in the PEEK cylinder, to serve as the working electrode, and pressed at 350 MPa for 1 minute. Indium foil (∼7.9 mm diameter) and lithium foil (∼4.8 mm diameter), both with a thickness of 0.1 mm, were hand-pressed onto the opposite side of the pellet. The cell was then placed into a stainless-steel casing with a constant pressure of 250 MPa. The LSV measurement was performed with a scan rate of 0.1 mV s⁻¹ using a BioLogic SP-300.

Assembly of the all-solid-state battery

All-solid-state battery (ASSB) half-cells were assembled in custom-made 10 mm diameter stainless-steel cells, using stainless-steel plungers as current collectors. All cell preparation and handling were performed in an Ar-filled glovebox with H₂O and O₂ levels below 0.1 ppm. The composite catholyte was prepared by ball milling TiS₂ powder (Sigma-Aldrich, 99.9%) with the synthesized Li₂ZrSCl_4−xBr_x electrolyte in a 1 : 2 weight ratio (TiS₂ : SE) for 30 min. Approximately 12 mg of the catholyte mixture was uniformly distributed onto one side of the separator pellet, corresponding to an areal capacity of approximately 1.22 mAh cm⁻², and compacted at 300 MPa for 10 s to ensure intimate interfacial contact. The separator consisted of a ∼100 mg Li₆PS₅Cl pellet, prepared following previously reported procedures,³³ and pressed at 300 MPa for 10 s. On the opposite side of the separator, a 7.9-mm-diameter indium foil was placed directly in contact with the pellet, followed by a 4.8-mm-diameter lithium foil (∼1 mg) to form the Li–In alloy anode in situ. The assembled cell was sealed with vacuum grease to prevent air exposure. Electrochemical testing was conducted at room temperature under a constant stack pressure of ∼35 MPa. Galvanostatic charge–discharge measurements were performed over a voltage window of 1.0–2.5 V vs. Li–In at various C-rates, where 1C was defined based on the theoretical capacity of TiS₂ (239 mAh g⁻¹).

Results & discussions

Average structure and morphology

The X-ray diffraction patterns for the as-milled Li₂ZrSCl_4−xBr_x (x = 0, 2.0, 2.7, and 4.0) and Li₂ZrCl₆ reveal a clear structural evolution with increasing Br content, shown in Fig. 1a. For x = 0 (LZSC), the pattern indexes to a two-phase mixture comprising a trigonal P3̄m1 polymorph and a monoclinic C2/m polymorph. A transition from Li₂ZrCl_6, which typically crystallizes in the trigonal phase.^15,30 The corresponding Rietveld refinement plot, in Fig. 1b, Tables S1 and S2, confirms the coexistence of these two phases, with good agreement between the observed and calculated patterns.

Fig. 1 Structural evolution and morphology of Li₂ZrSCl_4−xBr_x (0 ≤ x ≤ 4) (a) powder X-ray diffraction of Li₂ZrSCl_4−xBr_x showing progressive peak shifts (dashed line), peak broadening, and intensity loss with increasing Br⁻ content. (b) Rietveld refinement of Li₂ZrSCl₄ (x = 0), confirming a biphasic structure comprising trigonal P3̄m1 and monoclinic C2/m phases. (c) Crystal structure models of the P3̄m1 and C2/m polymorphs of Li₂ZrSCl₄ and schematic representation of the structural evolution upon Br⁻ substitution, leading to an amorphous Li₂ZrSCl_1.3Br_2.7 network (d) SEM micrograph and corresponding EDS elemental maps showing uniform distribution of Zr, S, Cl, and Br in Li₂ZrSCl_1.3Br_2.7.

The refined structural models for both polymorphs are illustrated in Fig. 1c, featuring the layered trigonal and monoclinic frameworks of LZSC and the transition toward an amorphous configuration upon partial substitution of Cl⁻ by the larger Br⁻ anion. This transition is evident in the diffraction patterns in Fig. 1a, where the Bragg peaks broaden, diminish in intensity, and shift toward lower 2θ (∼35°, 50°), consistent with lattice expansion induced by Br⁻ addition. At higher substitution levels (x = 2.0, 2.7, 4.0), as in Li₂ZrSCl_1.3Br_2.7, the diffraction pattern becomes increasingly featureless, with broad humps replacing sharp Bragg peaks. This behavior indicates extensive amorphization and collapse of the long-range crystalline order, Fig. 1c, while short-range structural motifs are retained. The low-intensity peak at ∼27° is likely associated with the Li₂S secondary phases.

Fig. 1d shows a scanning electron microscopy (SEM) image of the LZSCB sample (see Fig. S1 for LZSC), revealing uniform particle agglomeration and a smooth surface texture, features characteristic of an amorphous material. No evidence of crystalline domains or phase separation was observed, indicating structural homogeneity. Energy dispersive X-ray spectroscopy (EDS) mapping (also in Fig. 1d) confirmed the even distribution of constituent elements – Zr, S, Cl, and Br – throughout the sample. As expected, lithium was not detected due to EDS limitations in detecting low-atomic-number elements. To address this, nuclear magnetic resonance (NMR) was employed, offering better sensitivity for detecting and quantifying light elements such as lithium.

Local structural analysis

A comparison of the ⁶Li MAS NMR spectra of LZSC and LZSCB is shown in Fig. 2a. The ⁶Li MAS NMR spectrum of LZSC shows two resonances at −0.68 and −0.54 ppm, consistent with the coexistence of trigonal and monoclinic phases identified by XRD (Fig. 1). The signal at −0.68 ppm (blue) likely corresponds to the more ordered trigonal phase, while the downfield resonance, −0.54 ppm, arises from Li environments in the less symmetric monoclinic lattice.

Fig. 2 Local structural environment analysis of Li⁺ ions (a) ⁶Li MAS NMR spectra of Li₂ZrSCl₄ and Li₂ZrSCl_4−xBr_x. (b) Select ⁷Li MAS NMR spectrum obtained from inversion recovery experiments showing the different relaxation behavior of the Li sites; the full series of inversion-recovery spectra is shown in the SI, Fig. S2. (c) Raman Spectra of Li₂ZrSCl₄ and Li₂ZrSCl_1.3Br_2.7 compared to related structures.

For LZSCB, only one broad isotropic resonance is observed, centered around −0.54 ppm. The line width increases from ∼19.7 Hz in LZSC to 26.9 Hz in LZSCB. The increased chemical shift dispersion with Br⁻ incorporation indicates increased structural disorder introduced by Br⁻ substitution, which broadens the range of Li-anion coordination environments and results in the wider ⁶Li peak observed for LZSCB. This is also consistent with the extensive amorphization observed in the XRD patterns. Li₂S impurity phases are observed in both samples at ∼2.4 ppm, these could result from unreacted precursors. This residual Li₂S is also detected as a secondary impurity by XRD, evidenced by weak diffraction features. Owing to the highly disordered or partially amorphous nature of the Li₂ZrSCl_4−xBr_x framework, quantitative phase analysis based on XRD is inherently uncertain, as peak broadening and overlap can influence refined weight fractions. ⁶Li MAS NMR directly probes Li speciation and provides a more reliable estimate of the Li₂S content in this system. Using a long recycle delay (200 s), the integrated ⁶Li NMR signal indicates that ∼17% of Li environments in LZSCB correspond to Li₂S, which converts to an estimated ∼2.3 wt% Li₂S¹⁴ indicating that Li₂S is present as only a minor impurity. Similarly, a ∼3.6 wt% Li₂S is observed for LZSC. A summary of the NMR-based phase quantifications is provided in the SI (Table S3).

To better resolve the different Li sites observed in LZSC, the T₁-edited ⁷Li MAS NMR spectrum is shown in Fig. 2b and S2, obtained from spin-lattice inversion recovery experiments at a recovery time of 0.9 s after inversion. Li⁺ sites with varied ion dynamics will exhibit different T₁ relaxation behavior and thus show different recovery rates to their equilibrium states. At an intermediate relaxation time shorter than 20 s, e.g., 0.9 s, the enlarged difference in intensity can help better resolve different components in the ⁷Li NMR spectrum. Fig. 2b reveals three Li⁺ sites in LZSC with slightly different T₁ relaxation times, indicating a heterogeneous local environment of Li⁺ and slow exchange among these lithium sites. Whereas LZSCB shows a symmetric NMR peak, well described by a single-component fit, suggesting a more homogeneous distribution of Li sites or fast chemical exchange among different Li sites, characteristic of fast ion conductors.³⁴

The Raman spectra of the Li₂ZrSCl_4−xBr_x series, presented in Fig. 2c, confirm the presence of distorted ZrX₆ (X = Cl, Br, S) octahedra as the base structural motifs,³⁵ with broad, overlapping bands indicating an amorphous, disordered lattice. Compared to sharp peaks in reference compounds (ZrCl₄, Li₂S, Li₂ZrCl₆), the LZSC and LZSCB spectra show more peaks from the Zr–Cl/Br/S bending and stretching modes and significant broadening of these peaks due to variations in bond lengths and angles contributing to complex vibrational dynamics. The broad peaks at >500 cm⁻¹ are likely to originate from Zr–S stretching vibrations in distorted or mixed-ligand octahedra such as ZrSCl_4−xBr_x. Additionally, the absence of the Li₂S peak at ∼370 cm⁻¹ confirms sulfur incorporation into the framework.³¹

Anion substructure

Anionic local structures in Li₂ZrSCl_4−xBr_x were directly probed using ³⁵Cl (spin-3/2) and ⁷⁹Br (spin-3/2) MAS NMR spectroscopy (Fig. 3). ³⁵Cl is a spin-3/2 quadrupolar nucleus, which makes its NMR behavior highly sensitive to local symmetry and electric field gradients (EFGs) around Cl⁻ ions. The ³⁵Cl MAS NMR of LZSC shows a sharp resonance at 9.6 ppm, suggesting a relatively symmetric local environment. The slight asymmetry is from 2nd-order quadrupolar effects. The ³⁵Cl NMR spectrum of LZSCB (Fig. 3a) exhibits an upfield shift (isotropic peak at 2.5 ppm) along with a broad (610 Hz), asymmetric lineshape, indicating increased site asymmetry and disorder from Br⁻ substitution in the LZSCB structure.

Fig. 3 Halide local environments in Li₂ZrSCl_4−xBr_x probed with NMR (a) ³⁵Cl MAS NMR spectra of Li₂ZrSCl₄ and Li₂ZrSCl_1.3B_2.7 (b) ⁷⁹Br MAS NMR spectra of Li₂ZrSBr₄ and Li₂ZrSCl_1.3B_2.7.

⁷⁹Br is a spin-3/2 quadrupolar nucleus, much like ³⁵Cl, but with a larger quadrupolar moment, thus more sensitive to local structural symmetry. The ⁷⁹Br MAS NMR spectra of Li₂ZrSCl_4−xBr_x further reveal the impact of halide substitution on the local structure (Fig. 3b). Both LZSB and LZSCB display broad resonances on the order of several kHz, characteristic of moderate quadrupolar interactions. In LZSB, a relatively broad (4.2 kHz) isotropic peak is observed at 116 ppm, while in LZSCB, the resonance becomes even broader (8.4 kHz), suggesting a higher degree of local structural disorder and the presence of mixed halide environments.

Ion dynamics probed with T₁ NMR relaxation time measurements

To probe Li⁺ dynamics in Li₂ZrSCl_4−xBr_x, ⁷Li T₁ NMR relaxation times were measured across a temperature range of 25 to 85 °C using the inversion recovery method under static conditions, as shown in Fig. 4a. LZSC shows a relatively large T₁ across the temperature range, 1.35 s at 25 °C, indicating slower Li⁺ dynamics. LZSCB, on the other hand, shows shorter T₁ values (starting at ∼0.37 s at 25 °C, decreasing to ∼0.15 s at 85 °C), suggesting faster Li⁺ motion in the Br⁻ substituted material.

Fig. 4 (a) Variable-temperature ⁷Li static T₁ relaxation times for Li₂ZrSCl₄ and Li₂ZrSCl_1.3Br_2.7 (b) linewidth narrowing of the ⁷Li static NMR spectra with increasing temperature, due to increased motion.

According to the Bloembergen–Purcell–Pound (BPP) theory³⁶ presented in eqn (1), T₁ is dependent on the correlation time τ_c of ionic motion.

where τ_c, denotes the average residence time between consecutive Li⁺ ion hops, ω_o is the Larmor frequency, γ is the magnetogyric ratio, µ_o is the magnetic permeability of free space, ℏ is the reduced Planck constant, and r_o is the interatomic distance.

In the slow-motion regime, wherein ω_oτ_c ≫ 1, increasing temperature leads to faster motion and shorter T₁ time due to the correlation time (τ_c) of the molecular motions approaching the Larmor frequency (ω_o). Both systems operate in the slow or intermediate motion regime, but LZSCB exhibits faster ion dynamics than LZSC, as indicated by the consistently lower T₁ values across the temperature window. Linewidth decreases with temperature in both samples (Fig. 4b), indicating dynamic averaging of dipolar interactions as lithium mobility increases. LZSCB generally shows broader lines (∼620 Hz at RT) compared to LZSC (∼400 Hz), in agreement with increased structural disorder.

Ion transport properties determined with EIS

Fig. 5a displays the Nyquist plot and the corresponding equivalent circuit fit (inset) for the SEs at room temperature. The data were modeled using an equivalent circuit composed of two resistor–constant phase element (R‖CPE) pairs in series with an additional CPE, accounting for total ionic transport and electrode interface contributions. Attempts to deconvolute bulk and grain boundary resistances, even at low temperatures (−20 °C; Fig. S3), were unsuccessful; thus, the extracted values represent total ionic conductivity.³⁷ Among the compositions studied, the LZSCB exhibits the highest ionic conductivity of ∼1.03 mS cm⁻¹, a nearly tenfold increase compared to LZC, which shows a conductivity of 0.10 mS cm⁻¹. The intermediate composition Li₂ZrSCl₂Br₂ (LZSCB2) shows an intermediate ionic conductivity of ∼0.57 mS cm⁻¹, while the end members LZSC and LZSB exhibit lower conductivities of 0.45 and 0.43 mS cm⁻¹, respectively (Fig. 5c and Table S4). To further investigate the transport behavior, variable-temperature (VT) impedance spectroscopy was performed (Fig. S3). The resulting Arrhenius plots (Fig. 5b and S4) show a linear dependence of conductivity on inverse temperature across the entire range, indicating the absence of phase transitions or thermal degradation. Activation energies extracted from the slopes of the Arrhenius plots were 0.41 eV (LZC), 0.37 eV (LZSC), 0.36 eV (LZSCB), 0.36 eV (LZSCB2), and 0.37 eV (LZSB).

Fig. 5 Charge transport properties of Li₂ZrSCl_4−xBr_x (0 ≤ x ≤ 4) determined with AC electrochemical impedance spectroscopy and DC polarizations. (a) Representative Nyquist plot and corresponding equivalent circuit fitting at 25 °C for Li₂ZrSCl_1.3Br_2.7 (b) Arrhenius-type plot for the series (c) chart showing the comparative ionic conductivity and activation energy for Li₂ZrCl₆ and Li₂ZrSCl_4−xBr_x (d) plots from DC polarization measurements to determine the electronic conductivity.

LZC exhibits a relatively high activation energy (∼0.41 eV) and low ionic conductivity (∼0.1 mS cm⁻¹), indicating strongly hindered Li⁺ migration. Upon sulfur incorporation, the activation energy decreases to ∼0.36–0.37 eV across the Li₂ZrSCl_4−xBr_x series, confirming that sulfur introduces a softer, more polarized anion lattice that lowers the intrinsic migration barrier. In comparison, Br⁻ substitution has little effect on the activation energy but strongly influences ionic conductivity through changes in the pre-exponential factor, which increases systematically and reaches a maximum for LZSCB (Table S4).³⁸ This trend indicates that the conductivity enhancement is driven primarily by increased Li⁺ ion jumping rates. This correlates with the enhanced Li⁺ dynamics observed in the ⁷Li T₁ NMR measurements.

The superior performance of LZSCB relative to LZSC and LZSB arises from an optimal balance between structural disorder and anion polarizability.³⁹ LZSCB is largely amorphous, Li⁺ transport is governed by short-range interactions rather than long-range order. In the more crystalline LZSC, Li⁺ migration is partially hindered by the residual crystalline domains and stronger Li–Cl interactions, which create deeper potential wells and limited site connectivity. Partial substitution of Cl⁻ with the larger, more polarizable Br⁻ in LZSCB further softens the anion framework, introduces intrinsic structural frustration within the disordered matrix, and increases the local free volume, as evidenced by lattice expansion and increased disorder in XRD (Fig. 1a). These effects enhance Li⁺ site accessibility, weaken electrostatic interactions, and likely couple to low-frequency anion motions that promote dynamic Li⁺ hopping.^24,40–44

In contrast, the fully Br-substituted LZSB becomes more ordered, compared with partially Br-substituted compositions, leading to decreased Li⁺ ion dynamics. While partial Br substitution enhances ionic transport relative to the Cl-rich end member, the x = 2 composition (LZSCB2) does not reach the maximum conductivity, demonstrating that compositional tuning beyond simple Br addition is required to achieve optimal transport behavior. Accordingly, LZSCB (x = 2.7) achieves superior Li⁺ transport by inducing optimal structural disorder.^39,45,46

The pressure dependence of ionic conductivity was evaluated under varying stack pressures, as shown in Fig. S5. Ionic conductivity increases with applied pressure due to improved particle–particle contact and reduced interfacial resistance,⁴⁷ and then plateaus at ∼1.02 mS cm⁻¹ between 35 and 40 MPa, with no further enhancement at higher pressures. Furthermore, electronic conductivity measurements (Fig. 5d, S6 and S7) show very low values (∼10⁻⁹–10⁻¹⁰ S cm⁻¹) across the Li₂ZrSCl_4−xBr_x series, confirming that the total measured conductivity is dominated by ionic transport. Notably, Li₂S is a poor Li⁺ conductor at room temperature⁴⁸ and, given its low volume fraction, has a negligible influence on ionic conductivity, which is dominated by the Li₂ZrSCl_4−xBr_x framework.

Electrochemical stability and half-cell cycling

The electrochemical stability window (ESW) of the compounds was assessed using linear sweep voltammetry (LSV), shown in Fig. S8. In the Li–In|SE|SE–C setup used for this measurement, the solid electrolyte (Li₂ZrSCl_4−xBr_x) was positioned between a Li–In alloy anode and a composite cathode consisting of carbon and the SE. Li–In alloy was used as the working electrode in the LSV measurements to minimize interfacial instability and enable a reliable assessment of the electrolyte's intrinsic electrochemical stability. Compared to pure Li metal, Li–In forms a less reactive interface with solid electrolytes, suppressing parasitic reactions and interphase growth that can otherwise dominate the LSV response.⁴⁹ The LSV results reveal that Li₂ZrSCl₄ exhibits the widest electrochemical stability window (∼0.7–2.6 V), due to the high oxidative stability of Cl⁻. Partial Br substitution in Li₂ZrSCl_1.3Br_2.7 slightly narrows the window to ∼1–2.6 V, with oxidative decomposition onset occurring near 2.6 V. Fully substituted Li₂ZrSBr₄ displays the narrowest range (∼1.3–2.0 V), with both earlier oxidation and reduction, indicating it is the least electrochemically stable. Fig. S8d summarizes this trend, where the broader voltage window of the all-Cl compound aligns with the fact that Br⁻ is more readily oxidized than Cl⁻.⁵⁰

When integrated into composite cathodes to facilitate ion transport, Li₂ZrSCl_4−xBr_x SEs demonstrated excellent cycling performance in ASSB half-cells. These cells were assembled using a TiS₂:2Li₂ZrSCl_4−xBr_x composite cathode, a Li–In alloy anode, and a Li₆PS₅Cl separator. Prior to full-cell testing, the compatibility between LZSCB and a Li metal anode was evaluated using time-resolved impedance spectroscopy on a Li|LZSCB|Li symmetric cell (see Fig. S9). The continuous increase in interfacial impedance observed in the Li|LZSCB|Li symmetric cell indicates the non-self-limiting formation of a resistive interphase at the Li/electrolyte interface, consistent with reductive decomposition of the halide framework under direct contact with Li metal. Similar behavior has been reported for a broad range of halide solid electrolytes and reflects their limited intrinsic stability against Li.⁵¹ In this work, Li₆PS₅Cl was intentionally used as a separator to decouple cathode performance from known interfacial challenges arising from direct contact between halide electrolytes and Li metal. This cell configuration allows LZSCB to be evaluated specifically as a catholyte, focusing on its ionic transport and compatibility within composite cathodes. Also, preliminary cycling experiments with NMC and LiFePO₄ (LFP) composite cathodes resulted in rapid cell failure (see Fig. S10) due to electrolyte instability at higher voltages, further supporting the selection of TiS₂ as an appropriate cathode active material for evaluating the LZSCB electrolyte system.

As shown in the voltage profiles (Fig. 6a, b and S11), at a current rate of 0.1C, initial discharge capacities reached 415 mAh g⁻¹ for LZSC, 382 mAh g⁻¹ for LZSCB, and 330 mAh g⁻¹ for LZSB. These discharge capacities at low current densities exceed the theoretical capacity of TiS₂ (239 mAh g⁻¹), a behavior commonly reported for TiS₂–sulfide–electrolyte composite cathodes when capacities are normalized to the TiS₂ mass. This excess capacity arises from interfacial and electrolyte-derived electrochemical contributions, enabled by intimate contact between TiS₂ and the solid electrolyte.^52–55 In the present work, using ball-milled TiS₂ together with a ball-milled halide catholyte (Li₂ZrSCl_4−xBr_x) increases the interfacial area and thus the capacity contributions from interfacial reactions. These contributions, often kinetically limited, are most pronounced at low C-rates and early cycles and are progressively suppressed at higher rates, where the electrochemical response becomes increasingly dominated by TiS₂ intercalation. The capacity stabilized after the second cycle, suggesting the formation of a passivating interphase that supports reversible cycling. Rate performance tests further showcased the superior behavior of LZSCB, which delivered 322, 289, 248, 200, and 133 mAh g⁻¹ at 0.1, 0.2, 0.5, 1, and 2C, respectively. In comparison, LZSC delivered 323, 263, 200, 141, and 62 mAh g⁻¹, while LZSB realized 330, 304, 235, 160, 109, and 48 mAh g⁻¹, at corresponding C-rates. The higher capacities observed for the LZSCB cell are linked to its enhanced ionic conductivity and the improved utilization of the active cathode CAM material. Notably, LZSCB retained 81% of its initial capacity after 80 cycles at 0.2C (Fig. 6c), underscoring its long-term cycling stability. These improvements are likely a result of its high ionic conductivity and enhanced compatibility at the cathode–electrolyte interface. Li₂S, present as a minor impurity, is electrochemically stable within the relevant voltage window and does not adversely affect electrolyte stability.⁵⁶ Even at a high current rate of 2C, LZSCB maintained a significantly higher specific capacity (133 mAh g⁻¹) compared to LZSC (62 mAh g⁻¹) and LZSB (48 mAh g⁻¹), highlighting its potential for application in high-performance ASSBs.

Fig. 6 Electrochemical profiles of half cells using (a) Li₂ZrSCl₄ and (b) Li₂ZrSCl_1.3Br_2.7 in the composite cathodes at 0.1C, 0.2C, 0.5C, 1C, and 2C, respectively. (c) The corresponding rate performance and long-term cycling stability cycled at 0.2C, where C = 239 mAh g⁻¹ referenced to the theoretical capacity of TiS₂. Li₂ZrSBr₄ is included for comparison.

To directly probe interfacial stability during long-term cycling, we measured the EIS before and after cycling (see Fig. S12 and Table S5). EIS data reveal increased bulk and interfacial resistances after cycling (R_bulk: 107.8 to 146.6 Ω; Interface 1: 14.9 to 24.7 Ω; Interface 2: 35.5 to 685.5 Ω). Owing to depressed arcs and overlapping time constants, assignments are made based on frequency position and literature precedent.^57–59 The mid-frequency response (Interface 1) is attributed to the TiS₂/catholyte|Li₆PS₅Cl interface and the low-frequency response (Interface 2) to the Li–In|Li₆PS₅Cl interface. The pronounced growth of the low-frequency resistance is consistent with progressive anode-side interphase formation, where sulfide electrolytes are known to form Li₂S-, Li₃P-, and LiCl-rich reduction products,⁶⁰ whereas cathode-side interphases typically involve sulfur- and phosphorus-rich species like S, polysulfides,^55,61 or in our case Li–Ti–P–S–Br–Cl phases. Despite this interfacial evolution, the cell maintains high coulombic efficiency and good capacity retention, indicating that the formed interphases are largely passivating rather than performance-limiting.

Conclusion

In this study, we developed a family of rare-earth-free halide–sulfide solid electrolytes, Li₂ZrSCl_4−xBr_x (0 ≤ x ≤ 4), synthesized via a rapid, energy-efficient, two-hour mechanochemical method using a mixed-anion design strategy. Structural analysis from XRD and NMR data reveals that partial Br⁻ substitution in Li₂ZrSCl₄ induces progressive amorphization, enhanced anion polarizability, and dynamic local disorder, leading to significantly improved Li⁺ transport. The optimized composition, amorphous Li₂ZrSCl_1.3Br_2.7, exhibited a high ionic conductivity of 1.03 mS cm⁻¹, outperforming its end members and an order of magnitude higher than that of crystalline Li₂ZrCl₆ compound. Sulfur incorporation introduces a softer, more polarizable anion lattice that lowers the intrinsic Li⁺ migration barrier from ∼0.41 eV in Li₂ZrCl₆ to ∼0.36–0.37 eV in Li₂ZrSCl_4−xBr_x. Whereas Br⁻ substitution further enhances lattice softness and promotes increased amorphization and dynamic disorder, the resulting conductivity enhancement is governed primarily by an increased pre-exponential factor, which peaks for Li₂ZrSCl_1.3Br_2.7, indicating optimal structural disorder that promotes fast Li⁺ ion dynamics. When applied as a catholyte in all-solid-state battery half-cells, LZSCB enabled high capacities, excellent rate capability, and stable long-term cycling, retaining 81% capacity after 80 cycles at 0.2C. The good performance of the solid-state battery cells employing Li₂ZrSCl_1.3Br_2.7, especially at high rates, is attributed to improved ionic conductivity and favorable interfacial compatibility between cathode and solid electrolyte. These findings demonstrate that compositional tuning via mixed halides is an effective strategy to develop scalable, high-performance chalcohalide electrolytes for next-generation solid-state batteries.

Author contributions

Thilina N. D. D. Gamaralalage: data curation, formal analysis, investigation, validation, visualization, writing – review & editing. Pawan K. Ojha: data curation, formal analysis, investigation, validation, visualization. Bright O. Ogbolu: data curation, formal analysis, investigation, validation, visualization, writing – original draft, writing – review & editing. Md Mahinur Islam: investigation, validation. Tehreem Toheed: investigation, validation. Sean C. Wilkerson: investigation. Yan–Yan Hu: conceptualization, formal analysis, funding acquisition, project administration, resources, software, supervision, validation, writing – review & editing.

Conflicts of interest

The authors declare no conflict of interest.

Data availability

The data supporting this article have been included as part of the Supplementary information (SI). Supplementary information is available. See DOI: https://doi.org/10.1039/d5sc08802j.

Acknowledgements

The authors gratefully acknowledge support from the National Science Foundation under grant no. DMR-1847038. Solid-state NMR measurements were conducted at the National High Magnetic Field Laboratory, supported by NSF Cooperative Agreement no. DMR-1644779 and DMR-2128556 and the State of Florida. Raman spectroscopy was carried out at the Materials Characterization Laboratory, Department of Chemistry and Biochemistry, Florida State University (FSU075000MAC), with technical assistance from Dr J. S. Raaj Vellore Winfred.

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Footnote

† These authors contributed equally.

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