Balancing vacancy and carrier concentration in halide solid electrolytes for all-solid-state sodium batteries

doi:10.21203/rs.3.rs-4978944/v1

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Balancing vacancy and carrier concentration in halide solid electrolytes for all-solid-state sodium batteries

https://doi.org/10.21203/rs.3.rs-4978944/v1

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Designing halide solid electrolytes (SEs) with high ionic conductivity and good (electro)chemical stability is essential for the advancement of all-solid-state sodium-ion batteries (ASSIBs). Unfortunately, most sodium-based halide SEs experience limited ionic conductivities and ambiguous correlation between their structure features and ion transport properties. Here we report a novel strategy to significantly boost the conductivities of sodium halides by regulating vacancy and charge carrier concentrations through a facile NaCl-poor method. This approach achieves a balanced structure with optimal vacancy and carrier content, rendering several-fold conductivities enhancement of series sodium halides. Furthermore, a fluorination-induced amorphization protocol is employed to enhance (electro)chemical stability and interfacial compatibility without detrimentally influencing conductivities. When paired with an uncoated Na₃V₂(PO₄)₃ cathode and a Na₃PS₄-coated Na₁₅Sn₄ anode, the Na_0.5ZrCl₄F_0.5 catholyte enables the battery to run for 300 cycles, retaining 94.4% of its initial discharge capacity at room temperature. This study provides a versatile pathway for creating inorganic ion conductors with high conductivity and long-term cyclability, advancing the development of ASSIBs.

Physical sciences/Materials science/Materials for energy and catalysis/Batteries

Physical sciences/Energy science and technology/Energy storage/Batteries

All-solid-state sodium-ion batteries (ASSIBs) are poised to revolutionize grid-scale stationary storage due to their potential for higher energy density and enhanced safety over conventional liquid electrolyte systems^1–3. Moreover, the abundance and low cost of sodium raw materials further bolster the appeal of ASSIBs for large-scale applications^4–6. The key to their effective operation lies in the development of solid electrolytes (SEs) with high room-temperature (RT) ionic conductivity (σ), superior interfacial (electro)chemical stability with electrodes, and adequate deformability to ensure interface affinity^7–9. Despite extensive research into various sodium SEs such as polymers, oxides, and sulfides, the practical deployment of ASSIBs has been hindered by the low RT conductivities of polymers, the high sintering temperatures required for oxides, and the narrow electrochemical windows and air sensitivity of sulfides^10–12. Recently, halide-based Li superionic conductors featuring monovalent anions (F⁻, Cl⁻, Br⁻, I⁻) have garnered attention for their high ionic conductivities (> 10^− 3 S cm^− 1) at RT, high-voltage stability, and compressibility¹³. This has sparked interest in the development of halide-based Na-ion solid electrolytes (HNISEs) as potential alternatives. Nonetheless, the design principles of Li-ion halide SEs cannot be imprudently adopted for Na-ion analogs, mainly attributed to the much larger ionic radius of Na⁺ (1.06 Å) than Li⁺ (0.76 Å)¹⁴. Consequently, most reported HNISEs exhibit inadequate conductivities at RT (~ 10^− 5 S cm^− 1). For instance, Na₃YCl₆ showed a conductivity of 9.85 × 10^− 8 S cm^− 1 at RT, considerably lower than its Li-ion counterparts (Li₃YCl₆: 5.1 × 10^− 4 S cm^− 1, Li₃YBr₆: 1.7 × 10^− 3 S cm^− 1, RT) as reported by Asano in 2018^15,16. Additionally, Jung et al. proposed trigonal Na₂ZrCl₆ and orthorhombic NaAlCl₄ with moderate conductivities of 1.8 × 10^− 5 S cm^− 1 and 3.9 × 10^− 6 S cm^− 1 at 30 ℃, respectively^17,18. The insufficient conductivities of HNISEs fundamentally challenge the capacity release and rate performance of ASSIBs.

In typical inorganic SEs, ion conduction primarily relies on vacancies and interstitial sites for ion hopping¹⁹. Increasing vacancy concentration in crystal lattices through hetero-valent doping has proven effective in enhancing conductivity by generating more sites for charge carrier transport³. For example, aliovalent Zr⁴⁺ substitution in Na_2.25Y_0.25Zr_0.75Cl₆¹⁵, Na_2.4Er_0.4Zr_0.6Cl₆²⁰, and Na_2.25Yb_0.25Zr_0.75Cl₆²¹ enhances RT conductivities to 6.6 × 10^− 5 S cm^− 1, 3.5 × 10^− 5 S cm^− 1, and 6.6 × 10^− 5 S cm^− 1, respectively. Recently, Meng et al. designed a class of Na_2.25−xY_0.25Zr_0.75Cl_6−x SEs achieved with improved σ of 4.0 × 10^− 4 S cm^− 1 at RT, attributed to the presence of a mixed nanocrystalline and amorphous phase structure that promotes the formation of additional vacancies and free volumes²². Typically, creating vacancies in structures reduces carrier numbers, while higher carrier content often corresponds to fewer vacancies, indicating a strong negative correlation. However, current research tends to examine the effects of vacancy and carrier concentration separately, overlooking their potential interplay. The intricate relationship between charge carriers, vacancy content, and their combined impact on the bulk conductivity of HNISEs remains poorly understood.

In addition to the intrinsic conductivities of Na-ion SEs, their interface stability with electrodes is crucial for the efficient operation of ASSIBs. Incorporating fluorine into crystal structures is a promising modification tactic to improve this stability. Fluorine-doped ion conductors exhibit outstanding electrochemical stability and interfacial compatibility due to the high electronegativity of F⁻, which facilitates the in-situ formation of an F-rich stable interphase between electrode and electrolyte interfaces^23–25. Moreover, recent research has also shown that fluorinated chloride SEs exhibit superior interface compatibility with sulfides²⁶. However, fluorinating inorganic ceramic SEs can reduce their conductivities owing to the robust ionic bonding between fluorine anions and mobile carriers, which hinders cation diffusion within the structures^27,28. Therefore, balancing ion transport and electrochemical stability in fluoride-doped HNISEs presents a significant challenge.

Herein, we precisely tune the vacancy concentration and Na-ion content within Na_2 + xM_xZr_1−xCl₆-type (M = Yb, Er, x = 0, 0.25, 0.4) HNISEs using a NaCl-poor strategy without adding any doping agents. Moreover, we investigate the effects of vacant-site and Na-ion contents on SEs' ion transport capability. Contrary to previous understandings^3,12,29, achieving a rational equilibrium between vacancy concentrations and mobile charge carrier numbers within HNISE structures is pivotal for ion migration efficiency. Building on this understanding, we developed a series of NaCl-poor HNISEs (Na_xZrCl_4+x, Na_xEr_0.4Zr_0.6Cl_3.6+x, and Na_xYb_0.25Zr_0.75Cl_3.75+x) that achieved the highest conductivities and lowest activation energy (E_a) under an optimized balance of vacant-sites and Na-ion. Additionally, we incorporated an appropriate proportion of F anions into the optimized Na_0.5ZrCl_4.5−yF_y structure. This modification enhanced σ (Na_0.5ZrCl₄F_0.5, 1.1 × 10^− 4 S cm^− 1 at 25 ℃) through the fluorination-induced amorphization effect and provided excellent electrochemical compatibility with both electrodes and asymmetric sulfides SE interfaces. Using Na_0.5ZrCl₄F_0.5 as the catholyte, coated-free Na₃V₂(PO₄)₃-based ASSIBs demonstrated stable cycling over 300 cycles at 0.1 C under RT. Our novel insights offer significant design flexibility and potential for developing high conductivity and (electro)chemically stable Na-ion SE systems.

Ion transport scenarios under various vacancy and Na-ion concentration

The crystal structures of superionic conductors critically determine their ion transport capabilities and migration bottleneck sizes. In crystalline inorganic SEs, ionic hopping predominantly occurs between adjacent sites with the lowest energy barriers, whether through single-ion or multi-ion concerted migration³⁰. Furthermore, charge carriers typically conduct through the grain bulk and along boundaries via interconnected diffusion channels formed by the anion framework. Broadly, the conductivity of inorganic SEs can be succinctly expressed by Eq. (1):

σ = qnµ (1)

Where q, n, and µ are indexed to the elementary charge, charge carrier concentrations, and the mobility of carriers, respectively. Notably, carrier mobility is influenced by various factors, such as defect concentration, interatomic forces, disorder, and channel size³. In halogen anion frameworks, ion conduction mechanisms are shaped by vacancy and mobile charge carrier concentrations³¹, which typically exhibit a negative correlation. Excess charge carriers hinder ionic hopping, while too many vacancies increase migration distances. Thus, theoretical concentration-balanced thresholds for carriers and vacancies likely exist, with SEs achieving maximum conductivity when these thresholds are met. As illustrated in Fig. 1a, we disentangled the impact of carrier and vacant-site contents on Na-ion diffusion through three designed scenarios. In scenario (Ⅰ), where Na-ion content > threshold (\(\:{\text{T}}_{{\text{Na}}^{\text{+}}}\)) and vacancy content < threshold (T_vacancy) in chloride frameworks, Na-ion hopping is hindered due to the limited number of available vacancies. Thus, the ion conduction process is dominated by vacancies. On the contrary, scenario (Ⅱ) shows that when Na-ion content < \(\:{\text{T}}_{{\text{Na}}^{\text{+}}}\) and vacancy content > T_vacancy, ion hopping is limited by the reduced mobility of Na-ion charge carriers. Consequently, the ion conduction process becomes ion-dominated. Scenario (Ⅲ) demonstrates that when both Na-ion and vacancy content are at their respective thresholds, Na-ion conduction is most efficient, indicating optimal ionic transport kinetics. Notably, a strong correlation exists between the concentration thresholds of vacancies and charge carriers in the structure. As vacancy content nears its threshold, charge carrier concentrations also approach their theoretical limits. Conversely, when these concentrations deviate from their thresholds, the variation trends of vacancies and charge carriers become diametrically opposite.

Meanwhile, the experimental results strongly support these viewpoints. As illustrated in Fig. 1b, optimizing vacancy and charge carrier contents using a NaCl-poor method yields compositions such as Na_0.5ZrCl_4.5, Na_0.6Er_0.4Zr_0.6Cl_4.2, and Na_0.4Yb_0.25Zr_0.75Cl_4.15, delivering higher conductivities of 8.1 × 10^− 5 S cm^− 1, 5.7 × 10^− 5 S cm^− 1, and 2.0 × 10^− 4 S cm^− 1 than their original samples at 25 ℃, respectively. As a proof-of-concept study, the representative trigonal phase structure (space group: P—3m1) Na_xZrCl_4+x SEs were selected to research their structure changes induced by the NaCl-poor process. Then, the electrochemical behaviors of ASSIBs based above electrolyte systems will be systematically evaluated.

Structural analysis, micro-morphology, and ion conduction

Figure 2a shows the ternary Na-Zr-Cl composition diagram, highlighting the stoichiometric ratios explored in this study (Na_xZrCl_4+x, where 0.2 ≤ x ≤ 2) with red markers. The Na_xZrCl_4+x SEs were synthesized using a mechanochemical ball-milling method with varying molar ratios of NaCl and ZrCl₄. The powder X-ray diffraction (XRD) patterns (Fig. 2b) revealed that the crystal structures of the Na_xZrCl_4+x series correspond to typical trigonal phases, isostructural with Li₂ZrCl₆ and Li₃YCl₆^16,32. Furthermore, as the NaCl molar ratio decreases, the XRD diffraction peaks of the samples show reduced intensity and a slight rightward shift. This suggests partial volume contraction in the crystal structure and lower product crystallinity. The weak crystallinity of the Na_xZrCl_4+x samples primarily resulting from reduced grain size and strain-broadening effects induced by the intense mechanochemical ball-milling. During the synthesis process, the electrolyte particles would be crushed and creating structural disorder and defects, which contributed to promoting ion hopping^33,34. Figures 2c-d and Figures S1-2 illustrate the XRD Rietveld refinements and corresponding crystal structures of representative Na_0.5ZrCl_4.5 and Na₂ZrCl₆ SEs. Detailed structural parameters are provided in Tables S1-2. The crystallographic atom sites within Na₂ZrCl₆ and Na_0.5ZrCl_4.5 SEs are identical, and both compounds share the same space group, P—3m1. This similarity indicates that the reduction in NaCl composition does not alter the electrolyte's microstructure, facilitating the examination of how vacant-site/Na-ion ratios affect the conductivities of HNISEs. Furthermore, Raman spectroscopy (Figure S3) reveals similar spectral profiles for Na₂ZrCl₆ and Na_0.5ZrCl_4.5 SEs, further supporting that the NaCl-poor process has minimal impact on the electrolyte structures. Distinctly, Na_0.5ZrCl_4.5 SEs exhibit more vacancies (Na and Cl vacant sites) and smaller lattice sizes compared to Na₂ZrCl₆ due to the NaCl-poor effect. The micro-morphologies of Na₂ZrCl₆, Na_0.5ZrCl_4.5, and Na₃PS₄ SEs are compared by scanning electron microscopy (SEM) images (Figure S4). The Na_0.5ZrCl_4.5 products present smaller particle sizes and lower porosity (under 300 MPa) than the others, indicating reduced crystallinity and improved deformability in NaCl-poor SEs.

Figure 2e compares the Arrhenius plots of the Na_xZrCl_4+x (0.2 ≤ x ≤ 2) series based on electrochemical impedance spectra (EIS) at different measuring temperatures. In addition, the correlative σ and E_a, along with their EIS Nyquist plots of all products at 25 ℃ are demonstrated in Fig. 2f, Figure S5, and Table S3. As a result, all NaCl-poor samples exhibit improved conductivities and reduced E_a. The optimized Na_0.5ZrCl_4.5 composition achieves the highest σ of 8.1 × 10^− 5 S cm^− 1 and lowest E_a of 0.35 eV at 25 ℃, representing a more than six-fold enhancement compared to the pristine Na₂ZrCl₆ sample (1.3 × 10^− 5 S cm^− 1 at 25 ℃). Moreover, the electron conductivities of Na_xZrCl_4+x SEs are estimated to be approximately 10^− 9 S cm^− 1 (Figure S6, Table S7), indicating their favorable insulating properties. The significant improvement in conductivity of the NaCl-poor SEs likely arises from an increased number of vacant sites within the structures, which provide more free volume and larger ion transport channels with reduced bottleneck size. To be noted, the conductivity change trend of Na_xZrCl_4+x SEs displays a distinct volcano-shaped curve, suggesting that neither insufficient nor excessive NaCl within the structures (correlated with vacancy and carrier concentrations) is optimal for Na-ion diffusion.

To verify the aforementioned conjecture, we compared the changes in Na-ion content and vacant-site concentrations as mole number x increased (Fig. 2g and Figure S7). The NaCl-poor condition in Na₂ZrCl₆ reduces the Na-ion content in the Cl sublattices while increasing the vacancy concentrations (both Na and Cl vacant sites). For example, when x = 0.2 (Na_xZrCl_4+x), Na-ion in the Na_0.2ZrCl_4.2 structure occupies only 5% of the total ion sites, while the concentrations of Na and Cl vacant sites reach 95% and 30%, respectively (Fig. 2g (Ⅰ)). When x increases to 0.5, the electrolytes achieve their highest conductivity. At this point, both the Na-ion (12.5%) and vacancy contents (Na vacancy: 87.5%, Cl vacancy: 25%) reach their optimal thresholds (Fig. 2g (Ⅱ)). In stark contrast, further increasing the Na/Cl ion content (x = 0.8, 2) leads to decreased Na and Cl vacancy concentrations (Fig. 2g (Ⅲ-Ⅳ)), resulting in a corresponding reduction in conductivity (Figs. 2e-f). Based on the above analysis, we concluded that an optimal balance between mobile charge carrier and vacancy concentrations is crucial for efficient ion conduction in sodium superionic conductors. Building on this inference, we developed a series of Na-ion halide SEs systems with significantly improved conductivities by adjusting carrier and vacancy contents through facile poor-NaCl strategies. As demonstrated in Figures S8-9 and Tables S4-S5, the optimized compositions of Na_0.6Er_0.4Zr_0.6Cl_4.2 and Na_0.4Yb_0.25Zr_0.75Cl_4.15 exhibited superior conductivities of 5.7 × 10^− 5 S cm^− 1 and 2.0 × 10^− 4 S cm^− 1, respectively, compared to their pristine counterparts (Na_2.4Er_0.4Zr_0.6Cl₆: 1.9 × 10^− 5 S cm^− 1; Na_2.25Yb_0.25Zr_0.75Cl₆: 2.9 × 10^− 5 S cm^− 1) at 25 ℃. The above results strongly demonstrate that enhancing the conductivity of inorganic solid electrolytes by balancing vacancy and carrier concentrations is a universal and feasible strategy.

To elucidate the effect of electrolyte structural features on ion conduction, the bond valence site energy (BVSE) method was employed to calculate ionic transport channels and corresponding migration energy barriers. In the Na₂ZrCl₆ structure (Fig. 3a), the [Na1-Na2-Na1] chain along the [001] direction was identified as the most efficient 1D transport channel, with a migration barrier of 0.497 eV (marked in red in Figs. 3c-d and g). Additionally, the Na-ion can also migrate along the [Na2-i1-Na2], [Na2-i2-i3-Na1], and [Na1-i4-Na1] directions (marked in purple, blue/orange, and green, respectively, in Figs. 3c-d and g). These paths connect the [Na1-Na2-Na1] chain to form a 3D transport network, albeit with higher migration barriers of 0.673 eV, 0.696 eV, and 0.722 eV, respectively. In contrast, the Na_0.5ZrCl_4.5 structure (Fig. 3b) exhibits the lowest ion migration barrier of 0.479 eV along the [Na1-i1-Na1] direction (marked in red in Figs. 3e-f and h), effectively interconnecting the 1D [Na1-Na2-Na1] transport channels to form the most efficient 3D ionic percolating networks. Moreover, the other ionic transport channels along the [Na1-Na2-Na1] and [Na1-i2-Na1] directions (marked in green and blue, respectively, in Figs. 3e-f and h) also exhibit relatively low migration barriers of 0.591 eV. These findings further corroborate the more favorable Na-ion transport environments within NaCl-poor phase structures. It should be noted that BVSE analysis primarily characterizes ionic diffusion behaviors in crystal structures with periodic atomic arrangements and does not account for contributions from other non-ideal factors, such as amorphous matrices, defects, and surface structures. These features are typically nonperiodic owing to the robust ball-milling reactions. However, the crystalline component in NaCl-poor samples delivers lower migration energy barriers compared to pristine samples prepared under the same conditions. Thus, the effective transport barriers calculated from BVSE analysis provide powerful evidence of the significantly improved conductivities in NaCl-poor SEs. Furthermore, nonperiodic compositions likely play an important role in ionic migration, warranting systematic research applying advanced characterization techniques with exceptionally fine spatial resolution, such as aberration-corrected transmission electron microscopy³². Additionally, the conductivities of NaCl-poor electrolyte systems could be further improved by optimizing the proportions of nonperiodic components within the structures.

The general design strategy of the HNISEs in enhancing (electro)chemical stability

Ordinarily, the fluorine-ion substitution can effectively improve the (electro)chemical stability of HNISEs but at the cost of sacrificed conductivities due to the lower polarizability and smaller ionic radii of F⁻ (119 pm) compared to Cl⁻ (167 pm) or Br⁻ (182 pm)³⁵. Interestingly, in this study, the reduction in conductivity of fluorinated HNISEs can be well mitigated by the amorphization-induced protocol. Therefore, to further improve the electrochemical compatibility of NaCl-poor SEs with electrodes, varying amounts of F⁻ were incorporated into the Na_0.5ZrCl_4.5−yF_y (0 ≤ y ≤ 1) series through mechanochemical synthesis. XRD results (Fig. 4a) show that the pattern peaks diminish as the F⁻ content increases, indicating that these fluorinated SEs tend to become amorphous during the synthesis procedures. Concurrently, Raman spectrums further suggest the successful fluorination of Na_0.5ZrCl_4.5 without introducing additional impurities (Fig. 4b). Furthermore, the peak positions in the spectra of fluorinated samples emerge slight left shift (higher wavenumber) compared to pristine Na_0.5ZrCl_4.5 (~ 328, 158, and 124 cm^− 1), primarily due to the formation of shorter F-Zr bonds within the structures. The SEM image (Fig. 4c) of Na_0.5ZrCl₄F_0.5 reveals small particle sizes (< 5 µm) consisting of irregular micro- and nano-particles. The reduced average particle and non-crystalline structure result in denser cold-pressed pellets (~ 98.2% relative density, Table S8) with decreased porosity (Figure S10a). The improved compressibility of the powder samples is more favorable for intimate solid-solid interface contact between electrode materials and SEs³⁶. Furthermore, the corresponding energy-dispersive X-ray spectroscopy (EDS) mapping implies the uniform dispersion of Zr, Cl, and F elements within the samples, indicating successful chemical fluorination of Na_0.5ZrCl_4.5 rather than physical mixing of fluorides (Figures S10b-d).

The ion conduction properties of the Na_0.5ZrCl_4.5−yF_y (0 ≤ y ≤ 1) compositional series were assessed via EIS and Arrhenius plots (Figures S11 a-b). The conductivities of Na_0.5ZrCl_4.5−yF_y SEs reveal a pattern characterized by an initial increase followed by a subsequent decrease. At an F doping content near 0.5, the electrolyte achieves the highest σ of 1.12 × 10^− 4 S cm^− 1 at 25 ℃ (Fig. 4d). Furthermore, the Na_0.5ZrCl₄F_0.5 SEs display electron conductivities of approximately 10^− 10 S cm^− 1 (Figure S11c, Table S7), exhibiting excellent electron insulating properties. In brief, when the F doping content (y) is ≤ 0.5, the contribution of the amorphous effect to ion transport exceeds the constrained force between F⁻ anion and Na⁺ cation. Thus, the conductivities of Na_0.5ZrCl_4.5−yF_y (0 < y ≤ 0.5) SEs exhibited a growth trend. Whereas, for y > 0.5, the conductivities dramatically decrease because of the stronger attraction between F⁻ and carriers. Recent studies have demonstrated that the amorphization of SEs can significantly improve their ion conduction capabilities, owing to the unique local chemistry environments and ion migration mechanisms in amorphous domains^11,37–40. Consequently, the increased conductivity observed with partial substitution of F⁻ is mainly attributed to the amorphization effect induced by fluorination. However, excessive F⁻ tends to form robust bonds with mobile charges, impeding the effective diffusion of Na⁺ within crystal structures. Furthermore, the mechanism underlying the conductivity enhancement in amorphous fluorinated-based SEs requires further investigation in future studies. Given the favorable balance between ion conduction and (electro)chemical stability in fluorinated SEs, the optimized Na_0.5ZrCl₄F_0.5 samples will be the primary focus of subsequent research.

Another critical challenge for HNISEs is their sensitivity to moisture due to the deliquescent properties of most halides⁴¹. In this work, the moisture stability of various SEs was assessed using EIS Nyquist plots and XRD patterns. After 10 min of exposure to ambient relative humidity (RH, 35 ± 5%), the XRD pattern of Na_0.5ZrCl₄F_0.5 showed no additional impurity peaks (Figure S12). Furthermore, the conductivity of Na_0.5ZrCl₄F_0.5 remained stable after humidity exposure (Fig. 4e), demonstrating good moisture tolerance. In contrast, the XRD pattern of the Na_0.5ZrCl_4.5 pellet delivered attenuated peak intensity and reduced conductivity following similar moisture treatments (Figure S13). The improved humidity stability of Na_0.5ZrCl₄F_0.5 can be ascribed to the stronger Zr/Na-F bonds formed within the structures through F substitution^24,42,43.

The electrochemical stability of Na_0.5ZrCl₄F_0.5 was evaluated using cyclic voltammetry (CV) tests in an asymmetric configuration. Vapor-grown carbon fibers (VGCF)/SEs composites were used as the working electrode, while a Na₁₅Sn₄ (0.1 V vs Na/Na⁺) alloy served as the reference electrodes. In composite electrodes, the incorporation of VGCF allows for continuous and efficient electron transport, enabling precise monitoring of electrochemical reaction signals⁴⁴. Figure 4f demonstrates the positive scan curves for various SEs at a scan rate of 0.1 mV s^− 1. Notable, Na_0.5ZrCl₄F_0.5 exhibits the lowest integrated current of 1.34 mAV g^− 1, in contrast to Na₂ZrCl₆ (4.35 mAV g^− 1) and Na_0.5ZrCl_4.5 (4.26 mAV g^− 1), manifesting the fluorination process significantly enhances the oxidation stability of HNISEs within high-voltage regimes. X-ray photo-electron spectroscopy (XPS) was employed to assess the high-voltage stability of these HNISEs. When oxidation is up to 5 V, oxidation peaks related to \(\:{\text{Cl}}_{\text{x}}^{\text{-}}\) are evident in the spectra of Na₂ZrCl₆ and Na_0.5ZrCl_4.5, indicating their relative instability at elevated voltages (Figs. 4g-h). In contrast, the spectra for Na_0.5ZrCl₄F_0.5 (Figs. 4i-j) show no obvious oxidation peaks for \(\:{\text{Cl}}_{\text{x}}^{\text{-}}\) and \(\:{\text{F}}_{\text{x}}^{\text{-}}\), highlighting its exceptional anti-oxidation properties attributed to the stronger electronegativity of fluorine. Additionally, Zr 3d spectra exhibit minimal variation, indicating the maintenance of Zr in its highest oxidation state, Zr⁴⁺ (Figure S14).

In addition to impressive oxidation stability, Na_0.5ZrCl₄F_0.5 also demonstrates a decreased absolute reduction current of 0.0266 mA at 0 V (Figure S15), corroborating the improvement of intrinsic reduction stability against Na₁₅Sn₄ anode by fluorination strategy. Further analysis of the interface kinetic stability between Na_0.5ZrCl₄F_0.5 and Na₁₅Sn₄ involved cycling Na₁₅Sn₄/Na₁₅Sn₄ symmetrical cells. As illustrated in Figure S16, the initial plating/stripping potential of the Na_0.5ZrCl₄F_0.5 and Na_0.5ZrCl_4.5-based cells were approximately 0.8 V at a current density of 0.1 mA cm^− 2 and a capacity of 0.1 mAh cm^− 2. Remarkably, the Na_0.5ZrCl₄F_0.5-based symmetrical cell maintained stable plating/stripping for 800 h with minimal overpotential increases (~ 1 V), whereas its Na_0.5ZrCl_4.5-based counterpart delivered a progressive overpotential increased to approximately 2.6 V over the same duration. In contrast, the Na₂ZrCl₆-based symmetrical cell displayed a significant overpotential increase and shorter plating/stripping times (exceeding 5 V within 73 h), primarily due to the lowest conductivity and poor reduction stability of the electrolytes. Additionally, the improved cycling stability in Na₁₅Sn₄/Na_0.5ZrCl₄F_0.5/Na₁₅Sn₄ cells likely originates from the in-situ formation of a kinetically stable interphase layer between the fluorinated SEs and the anode interface (e.g., NaF, see Figure S17), which helps mitigate ongoing parasitic reactions during cycling^25,45–47.

Although Na_0.5ZrCl₄F_0.5-based symmetrical cells exhibit reduced plating/stripping overpotentials, these remain relatively high (~ 0.8 V) to effectively release the full battery's capacity. Therefore, the Na₃PS₄ (NPS, ~ 1.4 × 10^− 4 S cm^− 1 at 25 ℃), known for its exceptional kinetic stability against Na-Sn alloys^15,48, was incorporated as an intermediate electrolyte layer between the HNISEs and the Na₁₅Sn₄ anode. Moreover, recent studies have shown that fluorination of halides is advantageous for promoting interfacial compatibility with sulfides²⁶. We explored this by comparing the electrochemical behaviors of two symmetrical cell configurations: both utilized HNISEs as interlayers and NPS as outer layers. As depicted in Figure S18, both configurations initially exhibited low overpotentials (~ 0.2 V). However, the Na_0.5ZrCl_4.5-based sandwich symmetrical cell presented increasing plating/stripping potentials over time, signaling interfacial incompatibility between Na_0.5ZrCl_4.5 and NPS during cycling. In contrast, the Na_0.5ZrCl₄F_0.5-based sandwich symmetrical cell maintained excellent cycling stability with negligible overpotential increase, underscoring the beneficial impact of fluorine doping on halide/sulfide interfacial compatibility. Based on these findings, a dual-electrolyte layer strategy was adopted for ASSIBs, positioning Na_0.5ZrCl₄F_0.5 adjacent to the cathode and NPS near the Na₁₅Sn₄ alloy anode.

Electrochemical energy storage behaviors of ASSIBs

The electrochemical performance of mechanochemically synthesized Na_0.5ZrCl₄F_0.5, Na_0.5ZrCl_4.5, and Na₂ZrCl₆ HNISEs in ASSIBs was evaluated using Na₃V₂(PO₄)₃ (NVP) cathodes and VGCF conductive additives, as shown in Fig. 5a. Detailed material characteristics are provided in Figures S19-20. For comparative purposes, we constructed a single-layer cell configuration employing only NPS, without HNISEs (Figure S21a). As depicted in Fig. 5b, under a voltage window of 1.9 to 3.7 V at RT and a 0.1 C rate, the initial Coulombic efficiency (ICE) of the Na_0.5ZrCl₄F_0.5 cell was markedly higher at 94.5%, with a discharge capacity of 86.4 mAh g^− 1, in contrast to the Na_0.5ZrCl_4.5 and Na₂ZrCl₆ cells, which achieved ICE of 88.7% and 84.2% with discharge capacities of 70.3 mAh g^− 1 and 57.7 mAh g^− 1, respectively. The enhanced electrochemical stability of the Na_0.5ZrCl₄F_0.5 electrolyte appears to shield the other HNISEs and NPS from oxidation through interaction with NVP. Nonetheless, the cell employing NPS electrolyte exhibited a rapid decline in discharge capacity and increased interfacial resistance following the initial charge, attributed to the poor cathodic stability under high voltages, as further detailed in Figures S21b-c⁴⁹.

The rate and cycle capability of batteries employing Na_0.5ZrCl₄F_0.5, Na_0.5ZrCl_4.5, and Na₂ZrCl₆ as catholytes demonstrate comparable capabilities, as shown in Figs. 5c and d. Specifically, the Na_0.5ZrCl₄F_0.5 cell exhibits a higher specific capacity of 46.5 mAh g^− 1 at 0.3 C, significantly outperforming the capacities of 25.3 and 4.6 mAh g^− 1 observed in the Na_0.5ZrCl_4.5 and Na₂ZrCl₆ cells, respectively. When the current density is reduced back to 0.1 C, a reversible capacity of 82.3 mAh g^− 1 is restored. The improved rate capacity of the Na_0.5ZrCl₄F_0.5 cells primarily stems from their higher conductivity and enhanced compatibility with fluorinated electrolytes. It is noteworthy that the capacity limitations at higher rates are largely due to the suboptimal conductivities of the electrolyte layers, which are less than 1 mS cm^− 1. The Na_0.5ZrCl₄F_0.5-based cell also displays a substantial capacity retention of 94.4% over 300 cycles at 0.1 C, markedly surpassing the performance of Na_0.5ZrCl_4.5 and Na₂ZrCl₆ cells, which retain only 49.5% and 17.7% of their initial capacities after 150 and 100 cycles, respectively. Furthermore, the lower overpotential and enhanced cycling reversibility of the Na_0.5ZrCl₄F_0.5 cell are further demonstrated in the charge/discharge voltage profiles and corresponding dQ/dV integral curves (Figures S22-23).

The EIS tests were conducted to assess the interfacial stabilities between cathodes and HNISEs in both pre- and post-long-term cycling. As illustrated in Figure S24, Nyquist plots present two distinct semicircles at high and mid frequencies, corresponding to the impedance of separating electrolyte layers and composite cathode interfaces, respectively⁵⁰. Following extended cycling, the Na_0.5ZrCl₄F_0.5 cell (~ 711 Ω) demonstrated significantly lower interfacial resistance compared to Na_0.5ZrCl_4.5 (~ 4237 Ω) and Na₂ZrCl₆ (~ 7346 Ω) cells, indicating superior electrochemical stability of Na_0.5ZrCl₄F_0.5 catholytes during cycling. Moreover, the in-situ galvanostatic-EIS (GEIS) combined with the distribution of relaxation time (DRT) method was performed to evaluate the detailed electrochemical evolution of HNISE cells during charging. Figures S25a-c show no distinct changes in resistance during the charging process (2.9 to 3.7 V) for the Na_0.5ZrCl₄F_0.5 cell, yet notable impedance growth is observed in Na_0.5ZrCl_4.5 and Na₂ZrCl₆ analogs. Semi-quantitative DRT results (Figures S25d-f) categorize relaxation time (τ) into four groups—S1, S2, S3, and S4—across these three cells. Peak changes in the S3 (10^− 4 < τ < 10⁰ s) and S4 (10⁰ < τ < 10¹ s) regions are predominantly linked to the cathode/electrolyte interface and diffusion resistance, respectively⁵¹. For the Na_0.5ZrCl₄F_0.5 cell, peak intensities in the S3 and S4 regions show minimal changes with relaxation time, underscoring excellent electrochemical stability within composite cathodes. Conversely, peak intensities in S3 and S4 regions increase notably with relaxation time for Na_0.5ZrCl_4.5 and Na₂ZrCl₆ cells, indicative of irreversible parasitic reactions at cathode/electrolyte interfaces, leading to deteriorating interface stability and heightened cell impedance. In conclusion, based on EIS and DRT analyses, Na_0.5ZrCl₄F_0.5 catholytes exhibit superior high-voltage stability and interface compatibility with polyanionic-type positive electrodes, thereby boosting ASSIBs' performance.

To investigate the influences of chemical-mechanical failure on full cells, we compared the micro-morphological evolution of composite cathode surfaces using SEM images of pristine and cycled samples. Figure 6a illustrates that all three electrode types show similar microcrack areas on their surfaces before cycling, manifesting insufficient connection between HNISEs and cathode materials. After several hundred cycles, the Na_0.5ZrCl₄F_0.5-based composite electrodes delivered a uniform and dense structure with no discernible microcracks, in contrast to the other two types, which displayed more and longer cracks. This conversion stems from the continuous enhancement of NVP/Na_0.5ZrCl₄F_0.5 interfacial mechano-chemical stability during cycling, alongside the inherent high deformability of fluorinated SEs. Additionally, Na-ion diffusion coefficients (\(\:{\text{D}}_{{\text{Na}}^{\text{+}}}\)) for three configurations were calculated through the galvanostatic intermittent titration technique (GITT), a standard method for assessing ion transport kinetics within composite cathode⁵². As demonstrated in Fig. 6b and Figures S26, the Na_0.5ZrCl₄F_0.5-based cathode composites display higher \(\:{\text{D}}_{{\text{Na}}^{\text{+}}}\) and lower overpotential during charging, affirming the favorable kinetic and thermodynamic stability of Na_0.5ZrCl₄F_0.5 catholyte within the electrodes. Furthermore, the calculated interfacial reaction energies between HNISEs and electrodes provide further support for these observations (Fig. 6c and Table S9). The Na_0.5ZrCl₄F_0.5 (~ 84 meV) catholytes and their phase equilibria products (Na₂ZrF₆ ~ 15.5 meV) exhibit a lower interface reaction driving force with NVP cathodes compared to the Na_0.5ZrCl_4.5 (~ 88 meV) counterpart, suggesting a more favorable interfacial compatibility between the fluorinated electrolyte and NVP. It is noteworthy that the mutual reaction energies between HNISEs and anodes (Na metal and Na₁₅Sn₄) are significantly higher than those at the cathode interface, highlighting the inherent thermodynamic instability of the HNISEs/anode interface. However, favorable interfacial decomposed insulating phases (such as NaF and NaCl) can be formed from fluorinated SEs during cycling, which helps inhibit the interfacial parasitic reactions and the continuous decomposition of HNISEs^25,53. Thus, fluorinated SEs demonstrate greater stability at the anode interface, consistent with the stable cycling observed in Na_0.5ZrCl₄F_0.5-based symmetrical cells (Figure S16). In a word, the Na halide SEs could achieve significantly improved RT conductivities using the NaCl-poor strategy owing to an effective balance of vacancy and carrier concentrations within their crystal structures. Moreover, the introduction of fluorine into structures further enhanced their (electro)chemical stability and interface compatibility. Surprisingly, the fluorinated electrolyte also demonstrated boosted conductivity, deformability, and moisture stability, largely attributed to the amorphous effect and stronger bonding within structures (as summarized and illustrated in Figs. 6d and e).

In conclusion, we have developed a universal strategy for enhancing ion transport in Na halide-based SEs by adopting a NaCl-poor method. This approach optimally adjusts vacancy and carrier concentrations within the lattice, establishing critical balancing concentration thresholds that significantly enhance Na⁺ diffusion efficiency. As a result, the conductivities of these NaCl-poor SEs substantially exceed those of their standard counterparts. Furthermore, integrating fluoride into zirconium-based chloride SEs not only augments their electrochemical stability and interface compatibility with electrodes but also increases conductivity and moisture tolerance. This improvement is primarily attributed to the formation of an amorphous phase and the strengthened Zr-F bonds within the electrolyte matrix. Employing Na₃V₂(PO₄)₃ as a cathode and Na_0.5ZrCl₄F_0.5 as a catholyte, ASSIBs exhibit excellent cycling stability, maintaining approximately 94% of their initial capacity after 300 cycles at a discharge rate of 0.1 C. This study lays the groundwork for the design of advanced Na-ion halide SEs, optimizing both conductivity and (electro)chemical stability, thus promoting the durability and safety of ASSIBs.

Materials synthesis

High-purity anhydrous powders, specifically NaCl (Aladdin, 99.99%), ZrCl₄, ZrF₄, YbCl₃, and ErCl₃ (each from Aladdin at 99.9%), served as the precursor materials. The series sodium chloride-based SEs, designated as Na_xZrCl_4.5+x, Na_xEr_0.4Zr_0.6Cl_3.6+x, Na_xYb_0.25Zr_0.75Cl_3.75+x, and the fluorinated variant Na_0.5ZrCl_4.5−yF_y, were synthesized through high-energy ball-milling. This process employed a 20:1 ball-to-precursor mass ratio at 600 rpm over 10 hours, adhering to stoichiometric guidelines. Similarly, Na₃PS₄ (NPS) SEs were produced by ball-milling a stoichiometric mixture of Na₂S (Aladdin, 95%) and P₂S₅ (Macklin, 99%) under the same conditions, followed by annealing at 270 ℃ for 2 hours within vacuum quartz tubes.

Characterizations

X-ray diffraction (XRD) analysis was conducted using a D/MAX-RB Rigaku device with Cu Kα radiation to identify material phases and structures. To minimize air exposure, Kapton tapes were used to cover the sample surfaces. Phase quantification was refined using the Rietveld method with GSAS II software⁵⁴. The morphology and elemental composition of the samples were examined via scanning electron microscopy (SEM, JEOL JSM-7001F) and energy dispersive spectroscopy (EDS, Oxford). Porosity and relative density estimates were derived from SEM images, using the Adjust-Threshold plugin in ImageJ⁵⁵. Raman spectroscopic data were obtained with a JY HR800 confocal Raman microscope. Chemical compositions and states were assessed through X-ray photoelectron spectroscopy (XPS, ESCALAB 250Xi, Thermo Fisher). Ionic conductivities of synthesized SEs at temperatures ranging from 25 to 80 ℃ were measured using alternating current (AC) impedance spectroscopy over a frequency range of 0.1 Hz to 1 MHz, at a 10 mV amplitude, utilizing a Solartron 1260 + 1287 electrochemical station. Electronic conductivities were determined via the direct current (DC) polarization technique, applying a 1 V bias for 1800 s with stainless steel ion-blocking electrodes. Electrochemical stability was assessed through cyclic voltammetry (CV) using a CHI660B workstation. Measurements were conducted in a Na₁₅Sn₄|NPS/HNISEs|HNISEs-VGCF (vapor-grown carbon fibers) cell configuration within a 0 to 5 V range (vs. Na/Na⁺) at a scan rate of 0.1 mV s^− 1. The mass ratio of SEs to VGCF was maintained at 7:3, with an approximate hybrid mass loading of 10 mg. A symmetric cell configuration of Na₁₅Sn₄/SEs/Na₁₅Sn₄ was constructed to evaluate interfacial stability against the Na₁₅Sn₄ anode, testing at a current density of 0.1 mA cm² and a capacity of 0.1 mAh cm². Interfacial and full-cell resistances across different catholytes were characterized by electrochemical impedance spectroscopy (EIS). The model ASSIBs were assembled using Na₃V₂(PO₄)₃ (NVP) as the cathode material, Na₁₅Sn₄ (Na-Sn) as the anode, and a combination of halides SEs and/or NPS as the electrolyte separator layers. The cathode composites were prepared by ball-milling NVP, SEs, and VGCF in a mass ratio of 45:50:5 at 300 rpm for 30 minutes. The Na-Sn alloy anode was prepared through the ball-milling of sodium metal (Aladdin, 99.9%) and tin powder (Aladdin, 99.99%) at 300 rpm for 20 h. Dense electrolyte layers, 10 mm in diameter, were formed by compressing 40 mg of Na_0.5ZrCl₄F_0.5 and 60 mg of NPS at 100 MPa. Subsequently, 10 mg of cathode material and 30 mg of Na₁₅Sn₄ powder were uniformly spread on either side of the electrolyte pellets and compacted at 370 MPa for 2 mins. Galvanostatic charge-discharge cycles for the ASSIBs were conducted between 1.9 and 3.9 V versus Na-Sn/Na⁺ at 25 ℃ using a LAND CT2001A cell test system under a constant pressure of 70 MPa. Galvanostatic intermittent titration technique (GITT) tests were carried out at a rate of 0.1 C, with each charging step lasting 10 minutes followed by a one-hour rest period. Electrochemical impedance spectroscopy (EIS) assessed both full cell and interfacial resistances across various catholytes, employing a frequency range from 0.1 Hz to 10⁶ Hz at a 10 mV amplitude. To enable comparative analysis of interfacial impedance across various halide-based SE cells, we normalized the imaginary components (-Z'') of the impedance. In-situ EIS measurements were performed at various potentials (2.9, 3.1, 3.3, 3.5, and 3.7 V versus Na-Sn/Na⁺) at a rate of 0.1 C.

Theoretical Calculation

Bond valence site energy (BVSE) calculations were carried out using the softBV⁵⁶ program, which utilized structural models derived from XRD Rietveld refinement. The program assessed the energies of various Na sites within the crystal structures on a 3D grid with 0.1 Å resolution, employing the transferable Morse-type softBV force field. Zones of low BVSE, indicative of potential Na-ion conduction pathways, were identified. VESTA⁵⁷ software was used to visualize the crystal structures and the potential energy isosurfaces.

The interface pseudo-binary reaction energies were calculated as:

ΔE_D(SSE, electrodes, x) = E_eq(C_interface(C_SEs, C_electrodes, x)) - 𝐸_interface(SSE, electrodes, 𝑥)

Where C_SEs and C_electrodes represent the compositions of SEs and electrode materials, respectively, normalized to one atom per formula unit. The variable (x) denotes the molar fraction of the SEs. Material energies were obtained from the Materials Project database⁵⁸.

Data availability

The data that support the findings of this study are available from the corresponding author upon reasonable request.

Acknowledgments

Financial support from the National Key R&D Program of China (2023YFB2503902), National Natural Scientific Foundation of China (U21A2080), and Natural Science Foundation of Beijing Municipality (Z200011) are gratefully acknowledged.

Author contributions

L.Z.F. and C.W.N conceived the project. M.W. designed and conducted the experiments, including materials synthesis, characterization and data analysis. X.L. and L.G. contributed to data interpretation via computational studies. H.L., D.L., and X.Q. provided support in electrochemical testing and manuscript revision. All authors engaged actively in discussion and iterative refinement of the manuscript.

Competing Interests

The authors declare no conflict of interest.

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Balancing vacancy and carrier concentration in halide solid electrolytes for all-solid-state sodium batteries

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Abstract

Figures

Introduction

Results and Discussion

Ion transport scenarios under various vacancy and Na-ion concentration

Structural analysis, micro-morphology, and ion conduction

The general design strategy of the HNISEs in enhancing (electro)chemical stability

Electrochemical energy storage behaviors of ASSIBs

Methods

Materials synthesis

Characterizations

Theoretical Calculation

Declarations

References

Additional Declarations

Supplementary Files

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