Anode Compatibility of Halide Solid‐State Electrolytes

ABSTRACT

Recently, lithium‐metal‐halide (Li‐M‐X) superionic conductors have emerged as promising solid‐state electrolyte (SSEs) for all‐solid‐state lithium batteries (ASSLBs), owing to their high Li⁺ conductivity, excellent compatibility with cathodes, and good mechanical deformability. However, many of these conductors exhibit chemical and electrochemical instability when in contact with the reductive anodes, which hinders the direct use of lithium metal as anodes with extraordinarily high specific capacities and limits the overall energy density of halide‐based ASSLBs. Therefore, enhancing the anode compatibility of halide SSEs has become a critical task in ASSLB development and has attracted broad research interest. Herein, the underlying mechanisms responsible for the instability of halide SSEs against the lithium anode are elucidated based on both experimental observations and theoretical calculations. Recent strategies and progress aimed at improving the compatibility of halide SSEs with lithium anodes are summarized. Moreover, the effects of pressure and volume changes on the interfacial compatibility between halide SSE and lithium metal are discussed. In addition, the current challenges and future research directions are analyzed, aiming to provide theoretical insights and guidance to support further advancements in Li‐M‐X type solid‐state electrolytes and all‐solid‐state lithium batteries.

Halide lithium superionic conductors are emerging solid electrolytes for ASSLBs due to their high ionic conductivity and good cathode compatibility. However, their instability against lithium metal limits practical use. This work reveals interfacial degradation mechanisms and summarizes key strategies to improve anode compatibility for next‐generation high‐energy halide‐based ASSLBs.

1. Introduction

The search for solid‐state electrolytes (SSEs) with high ionic conductivity, good environmental and electrochemical stability is crucial to meet the increasing demands for all‐solid‐state lithium batteries (ASSLBs) with enhanced energy density and safety. In recent decades, various types of SSEs have been developed, including oxide‐based, sulfide‐based, and halide‐based SSEs [1, 2, 3, 4, 5, 6]. Oxide electrolytes feature high safety, excellent chemical stability, and environmental benignity. However, they generally suffer from poor mechanical deformability, and their synthesis typically requires elevated reaction temperatures, which substantially increasing production costs [7, 8]. Sulfide electrolytes exhibit high ionic conductivity and favorable mechanical properties at room temperature. However, their relatively narrow electrochemical stability window prevents their direct pairing with high‐voltage cathode materials. In addition, exposure of sulfide electrolytes to ambient air can generate toxic H₂S gas, which severely limits their practical application [6, 9]. Halide electrolytes offer several advantages over oxides and sulfides, including wide electrochemical windows, chemical compatibility with oxide cathode materials, and favorable scalability [1, 4, 5].

Halide SSEs generally refer to ionic conductors with a halogen anion framework. As early as the 1830s, researchers observed that lithium halides (LiX, X = F, Cl, Br, I) exhibit measurable ionic conductivity [10], but their room‐temperature conductivity was limited (∼10⁻⁷ S cm⁻¹) [11], which initially led to halide SSEs being neglected by researchers in the early days. A significant breakthrough occurred in 2018 when Li₃YCl₆ and Li₃YBr₆ were synthesized with room‐temperature conductivities exceeding 10⁻⁴ S cm⁻¹ [12]. These electrolytes can be directly paired with bare lithium‐cobalt‐oxide cathodes, thereby renewing researchers' interest in halide SSEs. Subsequently, numerous high‐performance halide SSEs have been developed, such as Li₃InCl₆ [13], Li₃ScCl₆ [14], Li₂ZrCl₆ [15], SmCl₃·0.5LiCl [16], Li_1.3Fe_1.2Cl₄ [17], Li₃Y_{1− x} In _x Cl₆, [18], Li_{3− x} M_{1− x} Zr _x Cl₆ (M = Y, Er, Ho) [19, 20, 21], Li_{3− x} (Zr/Hf) _x (Y/Er/Yb/In)Cl₆ [19, 22, 23, 24], Li₂(Sc/In)_2/3Cl₄ [25, 26], Li_{3−3 x} M_{1+ x} Cl₆ (M = Tb, Dy, Ho, Y, Er, Tm) [27]. These halide SSEs exhibit conductivities at the level of 10⁻⁴∼10⁻³ S cm⁻¹ and good stability with high‐voltage oxide cathodes. In recent years, dual‐ion SSEs have also been progressively developed [28, 29, 30, 31, 32, 33, 34, 35]. Notably, oxyhalides LiMOCl₄ (M = Nb, Ta) and Li₃Ta₃O₄Cl₁₀, have been developed and exhibited superior ionic conductivities exceeding 10⁻² S cm⁻¹ [36, 37]. Owing to their high Li⁺ conductivity, mechanical deformability, wide electrochemical window, and good compatibility with oxide cathodes, halide superionic conductors have emerged as one of the most promising SSE classes for ASSLBs [38, 39, 40, 41, 42].

Owing to its high theoretical specific capacity (3860 mAh g⁻¹), low negative electrochemical potential (−3.04 V vs. the standard hydrogen electrode), and excellent gravimetric and volumetric energy densities, Li metal is an ideal anode material for high‐energy‐density batteries [43, 44, 45, 46, 47, 48]. In particular, the use of thin lithium designs is increasingly recommended to further enhance the energy density of ASSLBs [49, 50, 51]. Thin lithium metal anodes (∼25 µm) with an NCM811 cathode can achieve an energy density of approximately 400∼500 Wh kg⁻¹ [50]. The anode‐free design can further enhance the energy density of ASSLBs while reducing the manufacturing challenges associated with electrode processing [52, 53]. Anode‐free solid‐state lithium metal batteries (AFSSLBs) consist of a bare current collector, SSE, and a prelithiated cathode. During the first charging process, Li⁺ ions deintercalated from the cathode are electrochemically plated onto the current collector as metallic lithium, and are subsequently stripped and reinserted into the cathode during discharge [54, 55, 56]. This design eliminates excess lithium metal, relying solely on the cathode‐stored lithium throughout the entire cycling process, corresponding to an approximately zero anode to cathode (N/P) capacity ratio, and the energy density is expected to exceed 500 Wh kg⁻¹ [57]. This design does not require metallic lithium as a starting material, can improve safety, and reduce cost by lowering lithium consumption as well as simplifying processing complexity [52, 53].

However, the development of Li metal anodes with high‐rate capability and high reversibility still remains one of the key challenges in current lithium battery research [58]. Owing to the low electronegativity of Li, most halide SSEs with high‐valence cations tend to be reduced once they come into contact with Li [4, 59, 60]. Such contact often induces severe interfacial reactions, ultimately leading to battery failure [61]. Addressing the interfacial incompatibility between Li metal anodes and halide SSEs requires a fundamental understanding of their interface reaction and reduction stability.

This review analyzes the mechanistic origins of halide SSE instability at the Li metal interface using both experimental insights and theoretical modeling. Strategies and recent progress for improving the interfacial compatibility of halide SSEs with Li metal anodes were summarized. Schematic illustration of the interfacial challenges and modification strategies for halide SSE/Li interfaces is shown in Figure 1. Moreover, key challenges and future research directions are discussed, with the aim of offering a deeper understanding of their chemical/electrochemical reactivity, thereby contributing to transformative advances in Li‐M‐X SSEs and ASSLBs.

FIGURE 1.

Schematic illustration of the interfacial challenges and modification strategies for halide SSE|Li interfaces.

2. Interface Between Li Anode and Halide SSEs

The energy convex hull can be used to quantify the thermodynamic stability of materials, and it is constructed based on the energy data of all compounds within the compositional space (Figure 2a) [62]. The grand potential phase diagram (GPPD) is designed to evaluate the variation law of material stability with external conditions. In lithium‐ion batteries, the electrochemical stability of SSEs at a specific potential relative to the lithium metal chemical potential (${\mu }_{Li}^0$) can be assessed by analyzing phase equilibria under controlled lithium chemical potential conditions. The GPPD constructed as a function of µ _Li provides insight into the electrochemical lithiation and delithiation potentials of SSEs, which correspond to their cathodic and anodic stability limits (Figure 2b).

FIGURE 2.

(a) Grand potential phase diagram revealing phase stability between A and B; (b) grand potential phase diagram revealing stability of A at different Li chemical potential, µ _Li . Reproduced with permission [62]. Copyright (2018), Elsevier Inc. (c) Schematic diagrams of Band‐edge approach. (d) Thermodynamically stable interface between Li|SSE; (e) Unstable interface between Li|SSE, with mixed ionic‐electronic interphase (MCI) forms; (f) Metastable interface between Li|SSE, with kinetically stable self‐limiting solid‐electrolyte interphase (SEI) forms. Reproduced with permission [63]. Copyright (2015), Elsevier B.V.

The gap between these cathodic and anodic limits is referred to as the electrochemical stability window of the solid electrolyte. The electrochemical stability window is determined by several factors, including the alignment of the electrolyte's valence and conduction band edges with the Fermi levels of the electrodes, thermodynamic stability in response to variations in mobile ion concentration, and the phase stability relative to the chemical potentials of its constituent elements [64]. In the band‐edge approach to electrochemical stability, an electrolyte is considered stable against oxidation if its valence band maximum lies below the Fermi level of the cathode with respect to the vacuum (ionization potential). Conversely, it is considered stable against reduction if the conduction band minimum, relative to the vacuum level (electron affinity), is above the Fermi level of the anode [65, 66], as shown in Figure 2c. When exceeding the range of the electrochemical stability window, the SSE undergoes reduction or oxidation reactions, accompanied by the transfer of electrons from the anode to the solid electrolyte and from the solid electrolyte to the cathode, respectively [67].

The optimal SSEs are those that are thermodynamically stable in contact with lithium metal, however, such materials have rarely been reported [68]. Alternatively, an SSE may undergo reduction to form an intermediate interphase, and thermodynamically unstable SSEs can still be employed in contact with Li if the reduction reaction is self‐limiting and the resulting interphase exhibits low resistance. Given the complexity of interfacial reactions between the lithium metal anode and solid electrolytes, the interface problem can be examined from the thermodynamic and kinetic points of view. Three distinct types of SSE|Li interfaces can be classified based on their interfacial reaction characteristics (Figure 2d–f) [63]. (1) Non‐reactive interfaces that are thermodynamically stable; (2) Thermodynamically unstable interfaces, once the Li contacts the solid electrolyte, a side reaction occurs, forming a mixed‐conducting interphase (MCI) with both high ionic and electronic conductivity. The MCI continues to react with Li metal, ultimately leading to battery failure [69]; (3) Interfaces that are thermodynamically unstable but kinetically stabilized due to limited electron transport across the interface, with self‐limiting solid‐electrolyte interphase (SEI) forms [70]. To quantify the growth behavior of such kinetically formed interphases, the Wagner model for diffusion‐controlled solid‐state reactions has been proposed [71]. This model predicts that the interfacial resistance grows parabolically over time and is strongly governed by the ionic and electronic conductivities of the interphase components. An ideal interphase should possess high ionic conductivity and low electronic conductivity to effectively suppress further degradation.

Since solid electrolyte interphases are likely to be nanocomposites comprising at least two distinct phases, kinetic studies must consider the following aspects: (1) charge transfer (CT) processes between Li and interphase components; (2) ionic transport within interphase phases; (3) CT kinetics between interphase components and (4) charge transfer between interphase and SSE [72]. The interfacial overpotential serves as an indicator of electrode reaction reversibility. A high overpotential typically results in reduced power density and energy efficiency. Electrode reactions consist of multiple partial reactions and transport steps, and the key objective of kinetic studies is to identify the rate‐determining step. Various methods have been developed to analyze complex electrochemical kinetics [73, 74, 75, 76].

In addition to interfacial reactions between SSEs and Li anodes, lithium dendrite growth within the SSE can also cause battery failure [77]. Lithium dendrite growth is a complex phenomenon governed by multiple coupled factors, including ionic diffusion, lithium‐ion concentration gradients, temperature, and mechanical stress. In contrast to liquid electrolyte systems, which have been extensively studied, dendrite formation in solid‐state systems exhibits fundamentally different characteristics and remains challenging to probe using non‐invasive techniques. Consequently, its underlying mechanisms remain incompletely understood. In SSEs, the formation of lithium dendrites can generally be categorized into two primary mechanisms. The first involves dendrite nucleation and growth at the electrode‐electrolyte interface, typically arising from non‐uniform lithium plating and stripping (Mechanism 1) [78, 79, 80, 81, 82, 83]. The second proposes that dendrites originate within the bulk of the solid electrolyte, particularly along grain boundaries, where localized reduction of Li⁺ to metallic Li⁰ leads to internal dendrite nucleation (Mechanism 2) [84, 85, 86, 87]. Recently, operando three‐dimensional ⁷Li magnetic resonance imaging (MRI) has enabled the simultaneous visualization of dendrites at both the interface and within the bulk of Li₇La₃Zr₂O₁₂ SSE [88]. This study reveals that Mechanism 1 and Mechanism 2 dominate at different stages of the charging process, and their relative contributions to overall dendrite formation are not fixed. Notably, the interaction between these two mechanisms provides further insight into dendrite growth behavior in high‐performance energy storage systems.

However, current understanding of lithium dendrite growth mechanisms mainly comes from studies on oxide and sulfide solid electrolytes. Halide electrolytes remain less explored due to the formation of relatively thick interphases between Li metal, which complicates direct investigation of dendrite behavior. Therefore, it is necessary to develop halide solid electrolytes that are kinetically stable with lithium metal and to systematically explore their dendrite growth mechanisms.

3. Cathodic Limit of Halide SSEs

Halide SSEs are highly valued for ASSLBs due to their high oxidation potentials (anodic limits). However, reactions with Li metal usually occur, as halide electrolytes tend to form thermodynamically unstable interfaces at the Li metal surface. The instability of halide SSEs against Li metal anodes primarily arises from the high reduction potentials (cathodic limits) of their high‐valence cations. This reduction of the central metal (M ^x+ → M⁰) occurs upon formation of a conductive interphase in contact with Li metal. The inherent instability of halide SSEs with Li metal leads to continuous interfacial reactions, resulting in high interfacial impedance, low coulombic efficiency, and eventually battery failure. To date, few experimental studies have investigated the interfacial composition between Li metal and halide SSEs. Advanced characterization techniques should be employed to comprehensively investigate the Li/halide SSE interface. This section provides an overview of the stability of various halide SSEs in contact with Li anodes, from both experimental and theoretical perspectives.

3.1. Theoretical Predictions for Reduction Stability of Halide SSEs

3.1.1. Thermodynamic Stability of Halide SSEs

Theoretical calculations show that the reduction limit potential of most fast lithium‐ion conductors is higher than that of the lithium metal anode, indicating that they are thermodynamically unstable against lithium metal [68, 89, 90, 91, 92, 93, 94]. The trends in the electrochemical stability window (ESW) of various Li‐M‐X (M represents cations and X comprises anions such as I, Br, Cl, and F) ternary compounds and other solid electrolyte systems are summarized in Figure 3a and Table 1. Compared with nitrides and oxides, halides exhibit relatively higher reduction potentials. Although the ESWs of all Li₃MX₆ systems are broad enough to accommodate the operating voltage range of typical cathode materials, their stability against Li metal anodes remains limited. Only binary lithium ionic compounds are stable at 0 V (vs. Li⁺/Li), as their anions are in the most reduced state (Figure 3a). The reduction potentials of Li₃MCl₆ compounds vary from 0.7‐2.6 V (vs. Li⁺/Li), depending on the species of the M cation. For example, when M is Sc, Y, or a lanthanide element, the reduction potentials range from 0.41 to 0.87 V (vs. Li⁺/Li). In contrast, M from group 13 elements (Al, In, Ga) leads to substantially higher reduction potentials ranging from 1.06‐2.55 V (vs. Li⁺/Li), as illustrated in Figure 3a. The variation in reduction stability among halide SSEs with different central metals (M) originates from differences in the HOMO‐LUMO energy gaps of the corresponding [MCl₆] octahedra. The orbital gap fundamentally reflects the electronic excitation barrier associated with the reduction of the central metal cation and therefore serves as a reliable electronic‐structure descriptor for evaluating reduction stability. In the recent work, the orbital gaps of [InCl₆]³⁻, [ZrCl₆]²⁻and [LuCl₆]³⁻ were calculated to be 4.685 eV, 5.864 eV, and 7.810 eV, respectively. The octahedral orbital gap of In³⁺, Zr^4+, and Lu³⁺ are consistent with the trend of reduction stability: the larger the orbital gap, the stronger the reduction stability, and vice versa [94].

FIGURE 3.

(a) Calculated thermodynamics intrinsic electrochemical stability window of Li‐M‐X ternary compounds in fluorides, chlorides, bromides, iodides, and other electrolytes. (b) Electrochemical stability window of halide SSEs with cationic and anionic substitutions [68, 89, 90, 91, 92, 93, 94]. The reference data were summarized in Table 1. (c) Heatmap of the reaction energy of anode/binary materials between lithium halide SSEs. The interface pseudo‐binary reaction energy was calculated from the materials project databases [95].

TABLE 1.

Electrochemical stability window and phase equilibria at the reduction and oxidation potentials of some representative solid electrolytes.

Compound	Reduction Voltage vs Li+/Li (V)	Phase equilibria at reduction onset	Oxidation Voltage vs Li+/Li (V)	Phase equilibria at oxidation onset	Refs.
LiF	0	LiF	6.36	F2	[91]
LiYF4	0.36	LiF, Y	6.56	F2, YF3	[89]
Li3AlF6	1.06	LiF, Al	6.48	F2, AlF3	[89]
Li3GaF6	2.28	LiF, Ga	6.59	F2, GaF3	[89]
Li3InF6	2.55	LiF, In	6.36	F2, InF3	[89]
Li2TiF6	1.9	LiF, TiF3	6.71	F2, LiF4	[89]
LiCl	0	LiCl	4.27	Cl2	[91]
Li2MgCl4	0.82	LiCl, Mg	4.23	Cl2, MgCl2	[89]
Li2ZnCl4	1.91	LiCl, Zn	4.21	Cl2, ZnCl2	[89]
LiAlCl4	1.54	LiCl, Al	4.45	Cl2, AlCl3	[89]
Li3AlCl6	1.59	LiCl, Al	4.26	Cl2, AlCl3	[94]
LiGdCl4	0.69	LiCl, Gd2Cl3	4.24	Cl2, GdCl3	[89]
LiGaCl4	2.11	LiCl, LiGaCl3	4.41	Cl2, GaCl3	[89]
Li2ZrCl6	1.75	LiCl, ZrCl3	4.25	Cl2, ZrCl4	[90]
Li2.5In0.5Zr0.5Cl6	2.28	InCl2, ZrCl4, LiCl	4.25	Li3InCl6, ZrCl4, Cl2	[90]
Li2.5Y0.5Zr0.5Cl6	1.75	YCl3, ZrCl3, LiCl	4.25	YCl3, ZrCl4, Cl2	[90]
Li2.5Er0.5Zr0.5Cl6	1.75	Li3ErCl6, ZrCl3, LiCl	4.25	Li3ErCl6, ZrCl4, Cl2	[90]
Li3InCl6	2.38	LiCl, In2Cl3	4.3	Cl2, InCl3	[89]
Li3In0.5Lu0.5Cl6	2.28	Li3InCl6, LuCl3, LiCl	4.25	Li3InCl6, LuCl3, Cl2	[90]
Li3In0.5Ho0.5Cl6	2.28	InCl2, HoCl3, LiCl	4.25	Li3InCl6, HoCl3, Cl2	[90]
LiTaCl6	1.86	LiCl, Ta	4.25	Cl2, TaCl5	[94]
LiNbCl6	2.17	LiCl, Nb	4.25	Cl2, NbCl5	[94]
LiTaOCl4	2.24	LiTa4O8, TaCl4, LiCl	4.06	Cl2, TaCl3O	[96]
LiNbOCl4	2.90	NbCl2O, LiCl	4.06	Cl2, NbCl3O	[96]
Li3ScCl6	0.87	LiCl, Sc	4.21	Cl2, ScCl3	[89]
Li3Sc0.5Lu0.5Cl6	0.91	Sc5Cl8, LuCl3, LiCl	4.25	ScCl3, LuCl3, Cl2	[90]
Li3Sc0.5Y0.5Cl6	0.91	0Sc5Cl8, YCl3, LiCl	4.25	ScCl3, YCl3, Cl2	[90]
Li3Sc0.5Ho0.5Cl6	0.91	0Sc5Cl8, HoCl3, LiCl	4.25	ScCl3, HoCl3, Cl2	[90]
Li3YCl6	0.62	LiCl, Y	4.21	Cl2, YCl3	[89]
Li3Y0.5Ho0.5Cl6	0.71	YCl3, LiCl, Ho	4.25	HoCl3, YCl3, Cl2	[90]
Li3Y0.5Dy0.5Cl6	0.66	YCl3, LiCl, Dy	4.25	YCl3, DyCl3, Cl2	[90]
Li3Y0.5Sm0.5Cl6	0.73	YCl3, LiCl, Sm	4.25	YCl3, SmCl3, Cl2	[90]
Li3LaCl6	0.41	LiCl, La	4.21	Cl2, LaCl3	[89]
Li3CeCl6	0.67	LiCl, Ce	4.27	Cl2, CeCl3	[94]
Li3PrCl6	0.60	LiCl, Pr	4.25	Cl2, PrCl3	[94]
Li3NdCl6	0.57	LiCl, Nd	4.21	Cl2, NdCl3	[89]
Li3PmCl6	0.72	LiCl, Pm	4.25	Cl2, PmCl3	[94]
Li3SmCl6	0.67	LiCl, Sm	4.23	Cl2, SmCl3	[89]
Li3GdCl6	0.76	LiCl, Gd	4.25	Cl2, GdCl3	[94]
Li3TbCl6	0.60	LiCl, Tb	4.21	Cl2, TbCl3	[89]
Li3DyCl6	0.66	LiCl, Dy	4.25	Cl2, DyCl3	[94]
Li3HoCl6	0.64	LiCl, Ho	4.25	Cl2, HoCl3	[89]
Li3Ho0.5Dy0.5Cl6	0.71	DyCl3, LiCl, Ho	4.25	HoCl3, DyCl3, Cl2	[90]
Li3Ho0.5Sm0.5Cl6	0.73	HoCl3, LiCl, Sm	4.25	HoCl3, SmCl3, Cl2	[90]
Li3Ho0.5Tb0.5Cl6	0.71	TbCl3, LiCl, Ho	4.25	HoCl3, TbCl3, Cl2	[90]
Li3ErCl6	0.71	LiCl, Er	4.25	Cl2, ErCl3	[94]
Li3TmCl6	0.62	LiCl, Tm	4.22	Cl2, TmCl3	[89]
Li3LuCl6	0.73	LiCl, Lu	4.25	Cl2, LuCl3	[89]
Li3Lu0.5Er0.5Cl6	0.72	Li3ErCl6, LiCl, Lu	4.27	Li3ErCl6, Cl2, LuCl3	[90]
Li3Lu0.5Y0.5Cl6	0.73	YCl3, LiCl, Lu	4.25	YCl3, Cl2, LuCl3	[90]
Li3Lu0.5Ho0.5Cl6	0.71	HoCl3, LiCl, Lu	4.26	HoCl3, Cl2, LuCl3	[90]
LiBr	0	LiBr	3.16	Br2	[91]
LiGaBr4	2.04	LiBr, LiGaBr3	3.26	Br2, GaBr3	[89]
Li3InBr6	2.17	LiBr, InBr2	3.15	Br2, InBr3	[89]
Li3InCl3Br3	2.17	LiBr, LiCl, InBr2	3.31	Li3InCl6, InBr3, Br	[90]
Li3LuCl3Br3	0.72	LiCl, LiBr, Lu	3.15	LuCl3, Br	[90]
Li3TmCl3Br3	0.67	LiCl, LiBr, Tm	3.14	TmCl3, Br	[90]
Li3ErCl3Br3	0.70	LiCl, Li3ErBr6, Er	3.17	ErCl3, Br	[90]
Li3YCl3Br3	0.59	Y2Cl3, LiCl, LiBr	3.17	YCl3, Br	[90]
Li2MgBr4	0.88	LiBr, Mg	3.15	Br2, MgBr2	[89]
LiAlBr4	1.8	LiBr, Al	3.9	Br2, AlBr3	[91]
Li2ZnBr4	1.78	LiBr, Zn	3.15	Br2, ZnBr2	[89]
Li3ScBr6	0.92	LiBr, Sc	3.93	Br2, ScBr2	[89]
Li3YBr6	0.59	LiBr, Y	3.15	Br2, YBr3	[89]
Li3LaBr6	0.61	LiBr, La2Br5	3.15	Br2, LaBr3	[89]
Li3NdBr6	0.59	LiBr, Nd	3.15	Br2, NdBr3	[89]
Li3SmBr6	0.61	LiBr, Sm	3.15	Br2, SmBr3	[89]
Li3GdBr6	0.59	LiBr, Gd	3.15	Br2, GdBr3	[89]
Li3TbBr6	0.67	LiBr, Tb	3.15	Br2, TbBr3	[89]
Li3DyBr6	0.67	LiBr, Dy	3.15	Br2, DyBr3	[89]
Li3HoBr6	0.68	LiBr, Ho	3.15	Br2, HoBr3	[89]
Li3TmBr6	0.68	LiBr, Tm	3.16	Br2, TmBr3	[89]
LiI	0	LiI	2.47	I2	[91]
Li2ZnI4	1.55	LiI, Zn	2.47	I2, ZnI2	[89]
LiGaI4	1.85	LiI, LiGaI3	2.53	I2, GaI3	[89]
LiInI4	1.98	LiI, InI2	2.5	I2, InI3	[89]
Li3OCl	0	Li3OCl	2.98	LiCl, LiClO4	[91]
Li2O	0	Li2O	3.11	O2	[91]
LiAlO2	0.17	Li5AlO4, Li3Al2	3.7	LiAl5O8, O2	[89]
Li4Ti5O12	1.99	Li5Ti7O16, Li2TiO3	3.79	TiO2, Ti2O7	[91]
Li2ZrO3	0.27	Li6Zr2O7, Zr	3.38	O2, ZrO2	[91]
Li7La3Zr2O12	0.027	Zr4O, Li2O, La2O3	3.15	Li2O2, La2Zr2O7, La2O3	[90]
Li4GeO4	1.11	Li2O, Ge	3.4	O2, Li2GeO3	[91]
LiNbO3	1.71	LiNbO2, Li3NbO4	5.1	O2, Nb2O5	[91]
Li3PO4	0.76	Li2O, Li3P	4.21	O2, Li4P2O7	[91]
LiGe2(PO4)3	2.93	GeO2, LiPO3, Ge	4.47	GeO2, O2, GeP2O7	[91]
LiTi2(PO4)3	2.71	Li2Ti2(PO4)3	4.77	Ti5P6O25, Ti2O7, TiP2O7	[91]
Li2S	0	Li2S	2.15	S	[91]
Li3PS4	1.72	Li2S, P	2.36	LiS4, P2S7	[89]
Li10GeP2S12	1.71	Li4GeS4, Li2S, P	2.29	Li3PS4, GeS2, LiS4	[90]
Li4SnS4	1.93	Li2S, SnS	2.35	S, Li2SnS3	[91]
Li6PS5Cl	1.71	LiCl, Li2S, P	2.13	Li3PS4, LiS4, LiCl	[90]
Li3N	0	Li3N	0.48	N2	[92]
Li9N2Cl3	0	Li9N2Cl3	0.5	N2, LiCl, NCl3	[93]
Li2.5N0.5S0.5	0	Li2.5N0.5S0.5	0.5	N2, LiCl, S	[68]

The impacts of both cationic and anionic substitutions on the electrochemical stability window of halide solid electrolytes are further analyzed. As a case study, the reduction potentials of Li₃YCl₆ and Li₃ErCl₆ are 0.62 and 0.71 V, respectively (Table 1). However, the substitution of Zr⁴⁺ raises the reduction potential of Li‐M‐Cl SEs to approximately 1.7 V (M is Y, Er) (Figure 3b). Notably, Li₃InCl₆, displays a significantly higher reduction potential (2.6 V vs. Li⁺/Li), far exceeding that of other Li₃MCl₆ solid electrolytes. The substitution of In³⁺ also raises the reduction potential of Li‐M‐Cl (M = Zr, Ho) SEs to approximately 2.6 V. These observations suggest that these materials are thermodynamically unstable in contact with lithium metal, likely due to the reduction of Zr⁴⁺ or In³⁺, thereby compromising their electrochemical stability [97], Substitutions involving Sc, Y, and lanthanide elements were further investigated. In contrast to Zr⁴⁺ and In³⁺ substitution—which markedly increases the calculated reduction potential, mutual substitution among elements such as Lu, Sc, and Y exerts only a minor influence on the calculated reduction potential (Figure 3b). This finding indicates that the reduction potential of Li‐M‐M′‐X SE is predominantly governed by the central metal M. The anionic substitution (e.g., Br doping) has little influence on the calculated reduction potential (Figure 3b). However, the reduction byproducts of the electrolytes contain significant volumetric fractions of LiBr, which are likely to passivate the interface [98].

The calculated reaction energies between Li‐M‐X ternary compounds and lithium metal (Figure 3c) are relatively high for Li₃InCl₆ (532 meV/atom), Li₃AlCl₆ (367 meV/atom), Li₂ZrCl₆ (443 meV/atom), Li₂ZnCl₄ (449 meV/atom), LiNbCl₆ (839 meV/atom), and LiTaCl₆ (787 meV/atom). These results are consistent with their higher reduction potentials and indicate poor thermodynamic stability against Li metal. In contrast, halide electrolytes with central metals of Mg, Sc, Y, or lanthanides show lower reaction energies (<200 meV/atom), suggesting improved compatibility with lithium metal. The reaction energies between these electrolytes and Li alternative anodes, including In, Si, Li‐In alloys, and Li‐Si alloys, were also evaluated. While the reaction energies with In, Si, and graphite anodes are relatively low, interfacial instability remains a concern for Li‐In and Li‐Si alloy systems. The chemical stability between the binary lithium ionic compounds and Li‐M‐X solid electrolytes was also assessed (Figure 3c). The reaction energies between Li‐M‐X solid electrolytes and Li₂O, Li₃P, and Li₃N exceed 200 meV/atom, leading to chemical decomposition. In contrast, Li₂S and Li₂Se exhibit lower reaction energies (approximately 100 meV/atom), while binary halides (LiF, LiCl, LiBr, and LiI) exhibit excellent chemical compatibility with Li‐M‐X electrolytes. Based on the electrochemical stability window and chemical compatibility, binary halides can be identified as effective protecting materials for stabilizing the halide‐anode interface.

3.1.2. Kinetic Stability of Halide SSEs

Electrochemical kinetic effects of Li|halide‐SSE interfaces are more directly captured by finite‐temperature ab initio molecular dynamics (AIMD) simulations. In the Li₃YCl₅Br system, an atomically resolved Li(110)|Li₃YCl₅Br interface was constructed, and its structural evolution was monitored through finite‐temperature AIMD simulations [99]. The results reveal that Li₃YCl₅Br is strongly reducing against lithium metal: Y³⁺ is reduced to metallic Y, while LiCl/LiBr phases form at the interface. Coordination environment analyses and charge partitioning further demonstrate that this decomposition process is highly exothermic. Ion transport parallel to the interface remains comparable to that in the bulk electrolyte, whereas cross‐interface Li⁺ transport toward the cathode is severely hindered, highlighting the pivotal role of reaction‐induced interfaces in regulating charge transfer. This analysis was further extended to the Li₃YCl₄Br₂ system, where contact with Li metal similarly induces LiCl/LiBr formation and Y³⁺ reduction. However, the resulting interphase adopts a layered and structurally flexible morphology that can dynamically reorganize upon cycling. The resulting dynamic solid electrolyte interphase exhibits mixed ionic‐electronic conductivity and mechanical compliance, which collectively account for the excellent long‐term stability observed in symmetric Li|Li₃YCl₄Br₂|Li cells [100]. For the prototypical Li₃YCl₆ electrolyte, AIMD simulations likewise show that direct contact with Li metal drives Y reduction and LiCl formation [101]. Nevertheless, introducing an electronically insulating yet Li⁺‐conducting interlayer can enforce electron tunneling, thereby suppressing deep cation reduction at the interface.

Recently, machine‐learning interatomic potentials (MLIPs) trained on DFT data are increasingly being applied to better capture the long‐range morphological evolution, dendrite nucleation, and the growth of nanocomposite interphase layers in SSE systems, demonstrating both DFT‐level accuracy and scalability to large spatial and temporal scales [102, 103, 104]. MLIPs have shown promise for bulk halide SSE transport simulations, but explicit ML‐based simulations of reactive Li|halide interfaces remain largely absent in the current literature.

3.2. Experimental Reduction Stability of Halide SSEs

In addition to the intrinsic thermodynamic stability of the material, the interfacial stability is also influenced by the electrochemical reduction products of SSEs. For example, the sulfide SSE Li₆PS₅Cl decomposes to Li₂S, Li₃P, and P₂S₅ when in contact with Li metal or under low potential conditions [105, 106]. These products can block electron pathways while maintaining Li⁺ conductivity, thus preventing further decomposition. This passivation mechanism makes Li₆PS₅Cl compatible with Li metal anodes. However, halide SSEs generally do not form favorable decomposition products when exposed to Li metal.

The reaction products between Li₃InCl₆ (Li₃YCl₆) and Li metal were examined using in situ X‐ray photoelectron emission spectroscopy (XPS), as shown in Figure 4a [60]. The results show that In³⁺ is reduced to In⁰, which can also be observed in the In‐MNN Auger spectrum. It was concluded that the Li₃InCl₆/Li interface is thermodynamically unstable, leading to a steady increase in interfacial resistance. Theoretical calculations and experimental characterizations suggest that the reactivity of halide SSEs with lithium metal is primarily governed by the central metal M. To understand this composition‐dependent reactivity, systematic investigations have been conducted on Li₃InCl₆, Li₂ZrCl₆, and Li₃YCl₆ using in situ XPS to track interfacial phase evolution and passivation behavior [107]. Although these halide SSEs exhibit similar ionic conductivities, they display markedly different surface decomposition behaviors when in contact with metallic Li, as illustrated in Figure 4b. Thermodynamically, all halide SSEs are unstable upon contact with Li, forming multiple stable and metastable phases. Among them, Li₃InCl₆ exhibits the highest reactivity, undergoing complete reduction within 10 min of lithium deposition. Li₃YCl₆ follows with 80% reduction, while Li₂ZrCl₆ exhibits the highest stability, with only 47% reduction. In situ XPS also revealed several metastable phases: In₅Cl₉, YCl, and ZrCl_2, which were previously predicted theoretically but not experimentally verified. Notably, the volumetric fraction and electronic conductivity of these metastable species also influence interfacial reaction kinetics and the degree of passivation.

FIGURE 4.

(a) X‐ray photoelectron In‐3d, Li‐1s, and Auger In‐MNN spectra during/after Li deposition on Li₃InCl₆. Reproduced with permission [60]. Copyright (2020), Wiley‐VCH. (b) Phase evolution in interphases with the time of Li deposition for Li₃YCl₆, Li₃InCl₆, and Li₂ZrCl₆. SEI % includes the contribution of LiCl, Li₂O, and Li₂CO₃ in the Li 1s spectrum. (c) The impedance evolution with time at OCV for Li|Li₃InCl₆|Li, Li|Li₂ZrCl₆|Li, and Li|Li₃YCl₆|Li. Reproduced with permission [107]. Copyright (2024), American Chemical Society.

To further investigate the evolution of interphases over time, researchers performed in situ electrochemical impedance spectroscopy (EIS) measurements on Li|SSE|Li symmetric cells [107], as shown in Figure 4c. In situ EIS is a non‐destructive technique that enables the study of the solid electrolyte interphase, interface evolution, and charge transfer processes in Li|SSE|Li symmetric cells over time. Figure 4c presents the impedance evolution during chemical stability tests of the three halide SEs, revealing distinct differences in their reaction kinetics with lithium metal. Li₃InCl₆ exhibits two semicircles in the impedance spectra after one minute of cell assembly, indicating rapid reaction kinetics with lithium metal, consistent with the in situ XPS observations. A rapid increase in impedance suggests the fast degradation of Li₃InCl₆, with XPS confirming the presence of a Li‐In alloy in the interphase, which exhibits high electronic conductivity. In contrast, Li₂ZrCl₆ does not display a new semicircle until one hour after contact with lithium metal, indicating a lower initial reaction activity of Li₂ZrCl₆ with lithium metal, and is consistent with the results of XPS. However, after 5 h, the impedance growth rate of Li₂ZrCl₆ surpasses that of Li₃InCl₆, suggesting that changes in reaction kinetics may be associated with microstructural evolution and the formation of new chemical species. Li₃YCl₆, does not exhibit the emergence of new semicircles even after 24 h of reaction. Instead, it shows an overall decreasing impedance trend. This indicates that the interface between Li metal has reached chemical equilibrium, enabling effective interfacial passivation, making it a more promising candidate for lithium metal anodes. Notably, it has been highlighted that the reaction between Li₂ZrCl₆ and lithium metal leads to the formation of a highly porous interphase layer, which is the primary reason for its increasing impedance. Furthermore, its kinetic behavior is significantly influenced by pressure variations, a factor that will be discussed in detail in the subsequent section on strategies for enhancing reduction stability.

In addition to Li₃InCl₆, Li₂ZrCl_6, and Li₃YCl₆, many other halides have been studied with regard to the stability against lithium metal anode, such as Li₃ScCl₆ [14], Li_2.73Ho_1.09Cl₆ [27], Li_0.388Ta_0.238La_0.475Cl₃ [108], Li₃LuCl₆ [109], and Li₃YbCl₆ [109]. The Li₃ScCl₆ represents an early demonstration of a Li‐M‐Cl halide solid electrolyte directly cycled against lithium metal, although it still exhibits clear polarization. The overpotential of the symmetric cell of Li_2.73Ho_1.09Cl₆ was lower than that of Li₃ScCl₆, suggesting that Ho may exhibit better stability against Li than Sc. The electrochemical stability window obtained from theoretical calculation confirms that Ho sample has a lower reduction potential (0.64 V vs. Li⁺/Li) compared to the Sc sample (0.91 V vs. Li⁺/Li). In recent years, increasing attention has been devoted to investigating the compatibility between halide electrolytes and lithium metal, leading to several breakthroughs. Li_0.388Ta_0.238La_0.475Cl₃ (LiTaLaCl) exhibits good interfacial compatibility with Li metal electrodes due to a gradient interfacial passivation layer that stabilizes the Li metal electrode during long‐term cycling in a Li‐Li symmetric cell (1 mAh cm⁻²) for over 5000 h [108]. Furthermore, Li₃LuCl₆ and Li₃YbCl₆ were reported to be kinetically stable with lithium owing to the formation of an electronically insulating self‐limiting SEI during cycling [109]. The formation of a self‐limiting SEI layer that contributes to kinetic stability with lithium will be discussed in detail in the following section.

It is worth noting that the compatibility between SSEs and lithium metal is commonly evaluated using Li|SSE|Li symmetric cell by monitoring its cyclability. However, insufficient understanding of the “soft breakdown” phenomenon can lead to misleading cycling stability in some symmetric cells [110, 111, 112]. Ideally, the charge carriers in solid‐state symmetric cells should be exclusively ionic. When both ionic and electronic transport occur in a Li|SSE|Li symmetric cell, the system is considered to undergo soft breakdown (Figure 5a). In contrast, hard breakdown refers to the system where only electronic transport dominates [110, 113].

FIGURE 5.

(a) Schematic diagram of all‐solid‐state batteries with different conduction mechanisms. (b) Li|Li₃ScCl₆|Li symmetric cell under the 0.2 mAh.cm⁻² with a current density of 0.2 mA.cm⁻². (c) Qualitative and quantitative analysis of soft breakdown by CV. (d) Li|Li₆PS₅Cl|Li symmetric cell under the 0.2 mAh.cm⁻² with a current density of 0.2 mA.cm⁻² (e) CV profiles of pristine Li|Li₆PS₅Cl|Li symmetric cells. (f) CV profiles of Li|Li₆PS₅Cl|Li symmetric cells with hard breakdown. (g) Temperature‐dependent EIS profiles of Li|Li₆PS₅Cl|Li symmetric cells with hard breakdown. Reproduced with permission [110]. Copyright (2022), Elsevier B.V.

Soft breakdown is prevalent in solid‐state batteries and has been observed across various SSE systems, including oxide electrolytes (Li_6.5La₃Zr_1.5Ta_0.5O₁₂), halide electrolytes (Li₃ScCl₆), and sulfide electrolytes (Li₁₀GeP₂S₁₂, Li₆PS₅Cl) [110]. For Li₃ScCl₆, the symmetric cell shows an initial increase in polarization during early cycling, followed by a gradual decrease and then stable cycling. This behavior has been proposed to originate from internal soft short circuits, which can result in deceptively stable cycling performance (Figure 5b). This phenomenon cannot be clearly identified only from plating and stripping profiles or standard EIS measurements. To address this issue, cyclic voltammetry (CV) has been proposed as an effective diagnostic tool for detecting soft breakdown (Figure 5c) [110]. An equivalent circuit model was constructed to simulate Li|SSE|Li symmetric cells under soft breakdown. In this model, I _peakrepresents the peak current in CV measurements, I _ionand I _eleccorrespond to the ionic and electronic current contributions, respectively. V is the applied bias, R _SSE denotes the bulk resistance of the SSE, R _int represents the resistance of the interphase between the SSE and lithium metal. R _ctis the charge‐transfer resistance, and R _e is the electronic resistance of the SSE, describing electronic transport within the symmetric cell. Under ideal conditions without interfacial reactions or electrolyte degradation, the CV current remains stable during cycling. When interfacial reactions or SSE degradation occur, the CV current decreases as R _int and R _ct increase. Notably, an increase in CV current only occurs when R _e decreases. Therefore, by analyzing the evolution of current in the CV profiles, soft and hard breakdown phenomena can be effectively distinguished.

To validate this approach, CV tests were carried out on Li|Li₆PS₅Cl|Li symmetric cells after prolonged cycling. Under a current density of 0.2 mA cm^{− 2}, soft breakdown occurred after approximately 20 h and then evolved into hard breakdown after 230 h (Figure 5d). The CV results reveal a sudden increase in peak current after the first cycle, followed by stabilization, indicating the coexistence of ionic and electronic transport. This confirms the presence of “deceptive lithium stability” (Figure 5e). After a hard breakdown occurs, the CV response becomes linear, indicating ohmic behavior, the EIS spectra also exhibit a characteristic inductive straight line (Figure 5f,g). These results demonstrate that CV is a simple and direct method to identify soft breakdown in ASSLBs.

In recent years, increasing attention has been paid to “soft short‐circuit” behavior in ASSLBs [111, 113, 114]. Operando transmission electron microscopy (TEM) has been used to track lithium deposition and current evolution, providing insight into the nanoscale origin of the transition from soft to hard short circuits. Under high electric fields, electronic leakage within SSEs can lead to the formation of Li⁰ in voids or along grain boundaries. These Li⁰ specie then accumulate into chains or clusters, gradually forming electronic pathways that result in soft short circuits. With continued growth, these pathways become fully connected, eventually leading to hard short circuits [111]. In addition, nondestructive characterization techniques, such as operando neutron imaging and X‐ray computed tomography (XCT) are important for detecting and visualizing soft short circuits in ASSLBs [114]. Future studies on halide SSE systems should pay close attention to identifying and eliminating soft breakdown behavior.

4. Strategies for Improving Li Compatibility of Halide SSEs

To address the issues associated with the cathodic limit, most halide SSE‐based ASSLBs employ lithium alloys instead of pure Li metal as the anode due to the relatively higher reduction potential. Among them, Li‐In alloy is the most commonly used. In addition, stabilizing the Li|SSE interface through the addition of extra protective layers has proven to be an effective strategy. The formation of self‐limiting SEI layers can also help mitigate interfacial issues, and this behavior is largely determined by the chemical composition of the halide SSEs. Last but not least, pressure regulation plays a critical role in influencing lithium behavior and battery performance, and should be considered a key focus for future in‐depth research.

4.1. Li Alloys for Stabilizing Lithium Metal Anodes

Lithium alloys serve as a promising strategy for achieving a stable electrolyte‐electrode interface, which is essential for ensuring long‐term cycling stability in solid‐state batteries [115]. Li can be mixed with elements such as In, Al, Mg, Zn, Si, Ag, Bi, etc. to form binary Li alloys [116, 117]. These alloy anodes exhibit improved chemical compatibility with halide SSEs owing to their relatively high reduction potentials (e.g., Li‐In alloy 0.62 V vs. Li⁺/Li) and the formation of kinetically stable interphases [118]. In addition, Li alloys often possess higher diffusion coefficient of Li atom (e.g., Li‐Ag alloy 10⁻⁸ cm²s⁻¹) [119] compared to pure Li metal (5.7 × 10⁻¹¹ cm²s⁻¹) [120], facilitating faster lithium transport at the interface and suppressing dendritic growth [121]. Moreover, alloy anodes can improve the interfacial wettability with halide electrolytes, which helps to maintain uniform contact between the halide SSEs and the alloy anodes [119]. Among numerous lithium alloys, Li‐In alloys remain the most commonly used anode in halide‐based ASSLBs [122]. At present, laboratory‐scale Li‐In anodes are primarily prepared by stacking and pressing lithium foil with indium foil, or by mixing the two metal powders to form composite materials. More complex architectures or advanced prelithiation approaches are rarely employed [123]. However, the use of Li‐In alloy anodes reduces the operating voltage and thus compromises the energy density of halide‐based ASSLBs [6, 124, 125]. Furthermore, Li alloy anodes undergo pronounced volume changes and morphological instability during the lithiation/delithiation process. At high current densities, alloy dendrites may also form, a phenomenon that has often been underestimated in earlier studies. Recent work has clarified that the shear modulus and partial molar volume of both alloy anodes and SSEs jointly determine the interfacial electrochemical behavior and dendrite evolution. Based on these insights, a general design principle has been proposed to improve the morphological stability of alloy anodes. This principle follows a “hard‐hard” and “soft‐soft” pairing strategy between alloy anodes and SSEs. This concept applies to halides and other electrolyte systems, providing a fundamental framework for designing alloy anodes with long cycle life in ASSLBs [126].

Therefore, future studies should place greater emphasis on optimizing the composition, mechanical properties, and cost of alloy anodes, which will facilitate their broader application in ASSLBs.

4.2. Interfacial Buffer Layer for Stabilizing Li Anodes

An alternative approach to addressing anode instability involves eliminating direct contact between lithium and halide electrolytes. A protective interlayer at the Li‐halide interface can facilitate Li⁺ transport across the interface and suppress side reactions without compromising cell voltage. Ideal interlayers are supposed to possess high ionic conductivity, electronic insulation, electrochemical stability, and strong chemical compatibility with both halide electrolytes and lithium metal [127]. A variety of electrolytes have been extensively investigated as interfacial buffer layers, such as sulfides [128], nitrides [129], borohydrides [130, 131], antistructure electrolytes [68, 132, 133], and polymers [134]. We will discuss each of these buffer layers in detail below.

4.2.1. Sulfide Buffer Layers

Previous studies have demonstrated that sulfide electrolytes, such as Li₆PS₅Cl, can function effectively as interfacial buffer layers between halide electrolytes and Li anodes [60, 135]. Although Li₆PS₅Cl has a narrow thermodynamic ESW, it can react with Li and generate electronically insulating phases such as Li₂S, Li₃P, and LiCl [136]. This will block electron transfer and suppress further degradation, generating thin, self‐limiting SEI [128], as shown in Figure 6a(i,ii). Without the Li₆PS₅Cl buffer layers, a mixed ionic‐electronic conductive interphase forms, which in turn promotes the growth of a thick decomposition layer at the Li|Li₃YCl₆ interface. EIS of the Li|Li₃YCl₆|Li symmetric cell shows the emergence and growth of a new semicircle at mid‐frequency over time, indicating continuous interfacial reactions (Figure 6a(iii)), consistent with the formation of conductive Li‐Y alloys with high electronic‐conductivity (Figure 6a(vi)). In contrast, in the EIS curves of the Li|Li₆PS₅Cl|Li₃YCl₆|Li₆PS₅Cl|Li symmetric cell, the semicircle at the middle frequency was absent (Figure 6a(iv)), confirming effective suppression of interfacial reactions by Li₆PS₅Cl (Figure 6a(vii)). This enables stable cycling for over 500 h at an overpotential below 100 mV, whereas the Li₃YCl₆ symmetric cell shows a progressive increase in overpotential to 600 mV (Figure 6a(v)) [135]. In addition to Li₆PS₅Cl, other sulfide electrolytes such as Li₇P₃S₁₁, and Li_6.7Si_0.7Sb_0.3S₅I have also been utilized to enhance interfacial stability in halide‐based ASSLBs [22, 98, 137]. Relevant studies reporting sulfides as buffer layers for halide SSEs are summarized in Table 2.

FIGURE 6.

(a) Insertion of sulfide interlayer between halides and anode. Schematic illustration shows the (i) Li|Li₃YCl₆ design and (ii) Li|Li₆PS₅Cl|Li₃YCl₆ design. Reproduced with permission [128]. Copyright (2021), Royal Society of Chemistry. (iii‐iv) Time‐resolved EIS spectra of Li|Li₃YCl₆|Li and Li|Li₆PS₅Cl|Li₃YCl₆|Li₆PS₅Cl|Li symmetric cells, respectively. (v) Symmetric cell performance comparison. (vi) The mixed electronic and ionic interface between Li metal and Li₃YCl₆. (vii) The Li⁺‐conductive interface between Li₃YCl₆ and Li enabled by a thin layer of Li₆PS₅Cl. Reproduced with permission [135]. Copyright (2021), AAAS. (b) Insertion of nitride interlayer between halides and anode (i) Li plating/stripping curves of Li|Li₂ZrCl₆|Li symmetric cells with and without the β‐Li₃N interfacial layer addition. The EIS spectra change with the number of cycles of the Li|Li₂ZrCl₆|Li symmetric battery without (ii) and with (iii) the β‐Li₃N buffer layer. Reproduced with permission [138]. Copyright (2022), American Chemical Society. (iv) Crystal structure of vacancy‐rich β‐Li₃N SSE. (v) Long‐term electrochemical performance of the NCM83|Li₃InCl₆|Li₃YCl₆|vacancy rich β‐Li₃N|Li all‐solid‐state lithium metal batteries. Reproduced with permission [92]. Copyright (2024), Springer Nature. (c) Insertion of antistructure interlayer between halides and anode (i) Schematics of the crystal structure of Li_2.5N_0.5S_0.5. (ii) Pictures of solid electrolyte pellets being exposed to melt Li metal. (iii) Schematic of symmetric batteries with Li₆PS₅Cl and Li_2.5N_0.5S_0.5 solid electrolytes. Reproduced with permission [68]. Copyright (2024), American Chemical Society. (iv) Schematics of the crystal structure of Li₂OHCl SSE. (v) Comparison of different current densities cycling for symmetric cells of Li|Li₃InCl₆|Li and Li|Li₂OHCl|Li₃InCl₆|Li₂OHCl|Li at 80°C. (vi) Comparison of the In 3d XPS spectra based on pristine Li₃InCl₆, and symmetric cells after cycling without and with Li₂OHCl interlayer. Reproduced with permission [139]. Copyright (2022), American Chemical Society.

TABLE 2.

Summary the addition of buffer layers between halide SSEs and Li metal anodes and electrochemical performances of halide‐based ASSLBs.

Halide SSEs	Cathode	Buffer layer	Anode	Cell performance		Operating Conditions		Refs.
				First cycle CE(%)/capacity (mAh g−1)	Capacity (mAh g−1)/current density/cycle	Temperature (°C)	Stack pressure (MPa)
Li3YCl6	NCM811	Li6PS5Cl	Li	87%/181	∼150/–/100	25	5	[128]
Li3YCl6	LiCoO2	Li6PS5Cl	Li	98.1%/139.1	118.5/0.1C/50	25	—	[135]
Li2Sc2/3Cl4	LiCoO2	Li6.7Si0.7Sb0.3S5I	Li‐In	93.7%/142	120/1C/50	—	187.5	[26]
Li2.7Yb0.7Zr0.3Cl6	NCM622	Li6.7Si0.7Sb0.3S5I	Li‐In	–/170	136/0.2C/150	25	—	[22]
Li2.633Er0.633Zr0.367Cl6	LiCoO2	Li3PS4	Li11Sn6	96.4%/110	∼81/0.5C/200	25	—	[19]
Li2ZrCl6	LiCoO2	Li6PS5Cl	Li‐In	97.9%/137	114/0.5C/100	25	187.5	[15]
Li2ZrCl6	LiCoO2	Li6PS5Cl	Li‐In	91.4%/156	∼138/0.5C/100	30	70	[140]
Li2.6Er0.6Zr0.4Cl6	LiCoO2	Li6PS5Cl	Li‐In	97.4%/147.5	∼90/0.5C/500	27	—	[97]
Li3InCl6‐Li3YCl6	LiCoO2	β‐Li3N	Li	96.9%/139.2	95.21/1C/5000	25	0	[92]
Li3HoCl6	NCM83	Li9N2Cl3	Li	89.3%/208.2	103/0.5C/1500	25	—	[93]
Li1.5ZrCl4.75O0.5	LiCoO2	Li2.5N0.5S0.5	Li	96.4%/125.2	71/1.5C/290	25	31.25	[68]
Li3InCl6	LiFePO4	Li2OHCl	Li	97%/148.8	130/0.1C/100	80	—	[139]
Li3InCl6	LiCoO2	PBO SPE	Li	95.4%/133.4	119.6/0.1C/50	30	0	[134]

4.2.2. Nitride Buffer Layers

In addition to sulfide buffer layers, nitrides also offer a promising route to stabilize the interface between lithium metal and halide SSEs, such as the use of β‐Li₃N and Li₉N₂Cl₃‐based electrolytes [92, 93, 138].

β‐Li₃N was firstly used as an interfacial modification layer to improve the interfacial stability between Li₂ZrCl₆ and the Li anode, as shown in Figure 6b(i–iii). β‐Li₃N exhibits high ionic conductivity and excellent electronic insulation, making it highly compatible with Li metal. Acting as an interlayer, β‐Li₃N effectively prevents direct physical contact between the halide SSE and Li metal, thereby effectively mitigating interfacial side reactions. With the introduction of a β‐Li₃N interlayer, the interfacial impedance between Li₂ZrCl₆ and the Li anode decreased significantly, dropping from 1929 to 400 Ω. Furthermore, the overpotential of Li symmetric cells was decreased from 250 to 50 mV at a current density of 0.1 mA cm⁻² without no significant increase for at least 300 h [138].

A recently reported vacancy‐rich β‐Li₃N SSE demonstrates superionic conductivity and exceptional compatibility with lithium metal [92]. This vacancy‐rich β‐Li₃N SSE exhibits an impressive ionic conductivity of 2.14 × 10⁻³ S cm⁻¹ at 25°C, surpassing most previously reported nitride‐based SSEs. Compared with conventional commercial Li₃N (in which both lithium‐vacancy and nitrogen‐vacancy concentrations are extremely low in the α and β phases), vacancy‐rich β‐Li₃N exhibits markedly increased concentrations of lithium and nitrogen vacancies (Figure 6b(iv)). Moreover, it shows intrinsic chemical stability against lithium with a reduction potential of 0 V, SEM images of pristine and lithium‐exposed vacancy‐rich β‐Li₃N exhibit negligible morphological changes, indicating minimal interfacial reactivity with lithium metal. Furthermore, ex situ X‐ray absorption near‐edge structure (XANES) analysis confirms the chemical stability of this SSE in contact with lithium. All‐solid‐state lithium metal batteries incorporating vacancy‐rich β‐Li₃N as a solid electrolyte interlayer, paired with Li₃InCl₆‐Li₃YCl₆ halide electrolytes and NCM83 cathodes, achieve a high‐capacity retention of 92.5% with 142 mAh g⁻¹ over 3500 cycles at 1.0 C (Figure 6b(v)).

4.2.3. Antistructure Buffer Layers

In addition to the sulfide and nitride buffer layers described above, antistructure SSEs (Li_2.5N_0.5S_0.5, Li₃OX or Li₂OHX, X = Cl, Br) have been developed and demonstrated excellent stability with Li [68, 132, 141, 142]. Recently, we proposed a lithium‐rich antifluorite solid electrolyte Li_2.5N_0.5S_0.5 (Figure 6c(i) [68], which exhibits high room‐temperature ionic conductivity and provides three‐dimensional fast Li⁺ transport pathways. This electrolyte exhibits intrinsic thermodynamic stability against lithium metal anodes (Figure 6c(ii)) and shows excellent interfacial stability in Li‐Li symmetric cells (Figure 6c(iii)).

In addition, Li₂OHCl has also been reported with an antiperovskite structure (Figure 6c(iv) and remains stable in contact with lithium metal even at temperatures exceeding the melting point of lithium [142]. When the cell was heated to 195°C, stable cycling was observed with a molten lithium anode, maintaining excellent Li exchange performance over 4500 min and a low cell voltage polarization of only 0.06 V. This stable cycling behavior is attributed to the formation of a robust self‐limiting SEI, produced via interfacial reactions between molten lithium and Li₂OHCl that generate Li₂O and LiCl. A higher concentration of Li₂O near the lithium anode forms the primary structure of the SEI, creating an interconnected network that protects the crystalline Li₂OHCl from further reaction with molten lithium. The high oxygen concentration within the SEI supports the formation of a lithium oxide layer, which stabilizes the interface between the molten Li anode and the solid electrolyte while preserving ionic conductivity.

Consequently, Li₃InCl₆|Li₂OHCl bilayer halides SSE has been developed [139]. Lithium symmetric cells employing the sandwich electrolyte Li₂OHCl|Li₃InCl₆|Li₂OHCl exhibited stable cycling for over 300 h at 80°C and achieved a high critical current density exceeding 0.6 mA cm⁻², as shown in Figure 6c(v). On the contrary, Li|Li₃InCl₆|Li cells exhibited poor Li stability. The Li₂OHCl initially reacted with Li metal and formed a self‐limiting SEI layer consisting of Li₂O and LiCl. This SEI layer provides excellent interfacial stability, intimate contact with Li metal, electronic isolation, and efficient ionic conductivity, thereby preventing further reaction between Li₂OHCl and Li metal. Additionally, XPS analysis confirmed the electrochemical stability of the Li₂OHCl|Li₃InCl₆|Li₂OHCl sandwich SSE (Figure 6c(vi)). After 300 h of constant‐current cycling, no significant shift in the binding energy of elemental indium was observed in the SSE. In contrast, when Li₃InCl₆ was directly contacted with Li metal, In³⁺ was reduced to metallic indium (In⁰).

Bilayer halide interface Li₃YCl₆|Li₃OCl has also been proposed to mitigate the interfacial compatibility of halide SE with Li metal. The interfacial stability and Li‐ion transport properties of the Li₃YCl₆|Li₃OCl|Li system were investigated using density functional theory (DFT) and ab initio molecular dynamics (AIMD) simulations [143]. AIMD simulations reveal that Y─Cl bonds break and Li─Cl bonds form at the Li₃YCl₆/Li interface. The distance between Y and Cl atoms at the interface is approximately ∼2.8 Å, which is significantly longer than the Y─Cl bond length of ∼2.6 Å in Li₃YCl₆ or YCl₃ [144]. Meanwhile, the Li‐Cl distance at the interface is ∼2.4 Å, closely matching the bond length in LiCl. The AIMD results indicate significant interfacial decomposition of Li₃YCl₆ in contact with lithium metal, along with the potential formation of decomposition products such as LiCl and metallic Y. These findings are consistent with previous experimental observations of the instability of Li₃YCl₆ solid electrolyte when in contact with Li metal [128]. Moreover, the interfacial stability of the Li₃YCl₆|Li₃OCl|Li is significantly improved. The systems remained stable even after 20 ps of AIMD simulation, maintaining an interfacial distance of approximately 2.2 Å after structural relaxation, in contrast to the larger deviations observed in the uncoated Li₃YCl₆|Li interface. This further confirms that, unlike the unstable Li₃YCl₆|Li interface, the bilayer halide|Li interface exhibits superior interfacial stability. This finding provides theoretical insight into strategies for improving the compatibility of halide solid electrolytes with lithium metal.

4.2.4. Borohydrides Buffer Layers

Borohydrides have been reported to exhibit excellent stability against the Li anode [131]. For example, a composite hydride lithium superionic conductor, 0.7Li(CB₉H₁₀)‐0.3Li(CB₁₁H₁₂), has been reported with a high conductivity of 6.7 × 10⁻³ S cm⁻¹ at 25°C [131]. This composite hydride demonstrated stable lithium plating/stripping behavior at 0.2 mA cm⁻², with minimal voltage polarization (6.0 mV) and negligible interfacial resistance (<1 Ω cm²). More importantly, after repeated lithium plating/stripping, the Li|0.7Li(CB₉H₁₀)‐0.3Li(CB₁₁H₁₂)|Li cell exhibited excellent voltage retention and highly stable lithium‐ion transfer at the electrolyte|Li interface. The remarkable interfacial compatibility of 0.7Li(CB₉H₁₀)–0.3Li(CB₁₁H₁₂) with the lithium metal anode is ascribed to its high chemical stability and high physical deformability. Although the integration of borohydrides with halide electrolytes has not been reported, their excellent compatibility with lithium metal suggests that they are promising candidates for interfacial buffer layers at the anode side.

4.2.5. Polymer Buffer Layers

In addition to the aforementioned inorganic electrolytes, researchers have investigated organic buffer layers. For example, a cross‐linked poly (butylene oxide) solid polymer electrolyte (xPBO SPE) thin film was fabricated on the surface of Li₃InCl₆ SSE to improve its compatibility with lithium metal [134]. This method produces a smooth xPBO SPE thin film of approximately 2 µm thick on the surface of Li₃InCl₆, referred to as xPBO@Li₃InCl₆. The xPBO@Li₃InCl₆ lithium symmetric cell exhibited highly stable EIS performance during aging, with interfacial resistance remaining around 200 Ω even after 10 h. Furthermore, cycling tests conducted on the xPBO@Li₃InCl₆ lithium symmetric cell under high current densities and capacity loadings demonstrated exceptionally stable operation over 200 h.

4.2.6. Incompatibility Issues of Buffer Layers With Halide Electrolytes and Solutions

Even though adding buffer layers at the Li‐halide interface can suppress side reactions without compromising cell voltage, the compatibility of buffer layers with halides remains a significant challenge. Recent studies have identified limited chemical compatibility between Li₆PS₅Cl and Li₃InCl₆, as revealed by time‐of‐flight secondary ion mass spectrometry (TOF‐SIMS), which showed an indium sulfide‐rich region at their interface (Figure 7a). This observation indicates that Li₃InCl₆ and Li₆PS₅Cl undergo interfacial reactions, rendering the interface thermodynamically unstable, predominantly resulting in the formation of indium sulfide‐like species [145]. To address this incompatibility issue, researchers have investigated the electrochemical compatibility of various halide and sulfide electrolytes by tuning the central metal M in halides [146]. The study demonstrates that the central metal M governs reaction kinetics with sulfides. Specifically, Li₃InCl₆ and Li₂ZrCl₆ exhibit high chemical reactivity toward Li₆PS₅Cl, whereas Li₃YCl₆, Li₃ScCl₆, and Li₃ErCl₆ display enhanced interfacial stability (Figure 7b). The chemical reactivity correlates with the electronegativity of the central metal. Higher electronegativity (i.e., a higher ionicity‐covalency parameter) increases the covalent character of the M─Cl bond. According to hard and soft acids and bases (HSAB) theory, highly electronegative cations polarize low‐electronegativity anions, promoting covalent bonding and increasing reactivity. For example, Zr⁴⁺, being more electronegative, enhances reactivity with sulfide anions, while Y³⁺ shows the lowest reactivity. These insights guide the rational pairing of halide and sulfide electrolytes in bilayer separators for ASSLBs.

FIGURE 7.

(a) TOF‐SIMS imaging of the interfacial area between Li₃InCl₆ (top) and Li₆PS₅Cl (bottom). Reproduced with permission [145]. Copyright (2023), Wiley‐VCH. (b) Schematic presentation of monolayer and bilayer cell configuration of two representative halide SSEs (Li₃InCl₆ and Li₃YCl₆), in bilayer configuration with Li₆PS₅Cl, with their merits and drawbacks. Reproduced with permission [146]. Copyright (2024), American Chemical Society. Heat map of the reaction energy between halides SSEs and representative sulfide SSEs (Li₆PS₅Cl, Li₁₀GP₂S₁₂, Li₃PS₄, Li₇P₃S_11, and Li_6.7SiSb_0.5S₅I) for (c) Li‐M‐Cl and (d) Li‐M‐Cl‐O SSEs. Reproduced with permission [29]. Copyright (2025), Wiley‐VCH. (e) Adding ALD coating between Li₃InCl₆ and Li₆PS₅Cl. Reproduced with permission [147]. Copyright (2022), American Chemical Society.

Recently, we proposed a dual‐anion strategy involving the incorporation of O²⁻ anions into the halide framework to mitigate the interfacial incompatibility between halide and sulfide electrolytes [29]. Our study revealed that a high concentration of oxygen can be incorporated into Zr‐based halide solid electrolytes without degrading ionic conductivity, while also significantly reducing the interfacial reaction energy with the sulfide electrolyte Li₆PS₅Cl. This modification facilitates the formation of a dense and low‐resistance Li₃PO₄‐rich interphase, thereby substantially enhancing the electrochemical performance of full cells. Furthermore, this O/Cl dual‐anion strategy is applicable to a broad range of representative halide/sulfide electrolyte systems, including Li₆PS₅Cl, Li₁₀GeP₂S₁₂, Li₃PS₄, Li₇P₃S₁₁, and Li_6.7SiSb_0.5S₅I (Figure 7c,d). This work provides valuable insights and practical strategies for enhancing interfacial compatibility between halide and sulfide solid electrolytes, thus promoting the advancement of ASSLBs. Similar observations have also been reported by other researchers [147, 148]. Specifically, the electrostatic potential difference between Li₆PS₅Cl and Li₃InCl₆ was found to hinder Li⁺ migration, resulting in the irreversible decomposition of Li₃InCl₆ into LiCl and InCl₃, which leads to a substantial increase in interfacial resistance and eventual battery failure. In contrast, the electrostatic potential gradient at the Li_1.75ZrO_0.5Cl_4.75|Li₆PS₅Cl interface facilitates Li⁺ diffusion and maintains interfacial stability. Even upon partial decomposition of Li₆PS₅Cl, Li_1.75ZrO_0.5Cl_4.75 effectively suppresses further side reactions through the formation of an oxygen‐rich interphase layer [148]. Recent studies have further investigated the effect of oxygen substitution on the chemical stability of halide electrolytes in contact with Li_5.5PS_4.5Cl_1.5 using density functional theory (DFT) calculations [149]. The calculations indicate that Li_3.25InCl_5.75O_0.25 exhibits improved thermodynamic stability against Li_5.5PS_4.5Cl_1.5 compared to Li₃InCl₆. XPS analyses further reveal that oxygen substitution improves the chemical compatibility with Li_5.5PS_4.5Cl_1.5 by suppressing the formation of impurity phases such as In₂S₃, P₂S₅, and Li₂S.

In addition to halide electrolyte modification, applying a nanoscale (1 or 2 nm) Li₃PO₄ coating between Li₃InCl₆ and Li₆PS₅Cl using atomic layer deposition (ALD) has proven effective in enhancing interfacial stability (Figure 7e) [147]. This uniform, pinhole‐free surface engineering technique enables the construction of halide‐based all‐solid‐state batteries with excellent long‐term stability. This technique is widely employed in the modification of oxide, polymer, and other solid electrolyte systems [122, 150, 151, 152, 153, 154].

4.3. Stabilization of Lithium Metal Anodes by an In Situ Formed Self‐Limiting SEI Layers

Adding artificial interlayers between halide SSEs and Li metal anode is a common strategy, however, it requires additional processing during cell preparation and is not conducive to large‐scale applications. Constructing in situ formed self‐limiting SEI layers is a more effective strategy to stabilize the interface between halide SSEs and Li metal anode. The in situ formed self‐limiting SEI layers should satisfy several requirements, including chemical stability with both halide SSEs and Li metal, high ionic conductivity, electronic insulation, and sufficient mechanical strength to prevent lithium dendrite growth [155]. LiF and LiCl compounds are effective in stabilizing the SEI and preventing lithium dendrite growth due to their high interfacial energy and low electronic conductivity.

Some doping strategies have been employed to enhance the stability of Li₃YBr₆ against Li by substituting fluoride ions (F⁻) [156]. Linear sweep voltammetry (LSV) results for Li₃YBr_5.7F_0.3 showed a negative shift in the reduction onset voltage and a decrease in overall reduction current compared to Li₃YBr₆, while the Li|Li₃YBr_5.7F_0.3|Li symmetrical cells exhibited significantly improved cycling performance (1000 h at 0.1 mA cm⁻²), as shown in Figure 8a,b. This improvement is attributed to the formation of a homogeneously distributed and highly concentrated LiF‐rich interphase between the halide SSE and Li metal. The dense, reticular fluorinated interfacial layer formed in situ effectively hinders lithium dendrite growth and interfacial side reactions at the anode. It also provides better interfacial contact and uniformity compared to artificially fluorinated interphases generated by adding LiFSI. Moreover, Li₂ZrCl_{6 ‐x} F _x has also been reported with improved Li stability [157, 158, 159]. When in contact with Li metal, Li₂ZrCl_{6 ‐x} F _x forms a stable interphase primarily composed of LiF, which passivates the interface by blocking electron transport while allowing lithium‐ion conduction. This effectively prevents further decomposition of the SSE and improves the performance of symmetric cells compared to Li₂ZrCl₆. Consequently, the study of lithium‐metal‐compatible fluoride‐based SSEs represents a promising area of research, as fluorides offer enhanced electrochemical stability at both cathodes and anode [160]. The synthesis of fluorine‐doped halide electrolytes provides a promising route for generating in situ self‐limiting SEI layers, and further studies are needed to optimize electrolyte compositions and cell configurations to stabilize the halide SSE|Li metal interface. However, F⁻ substitution typically compromises ionic conductivity, and the trade‐off between ionic transport and interfacial stability needs further investigation. Recently, a polyanion‐coordination strategy has been proposed for the design of fluorinated solid electrolytes, exemplified by Li_xTi(PO₄)_x/3F₄ [161]. Owing to weaker interactions between Li⁺ and its coordination environment, along with a more flexible framework, the optimized Li_1.3Ti(PO₄)_1.3/3F₄ achieves an ionic conductivity of 1.6 × 10⁻⁵ S cm⁻¹. This value is two orders of magnitude higher than that of Li₂TiF₆. This finding indicates a promising pathway for improving ionic conductivity in fluorinated solid electrolytes.

FIGURE 8.

(a) LSV curve of the Li₃YBr₆ and Li₃YBr_5.7F_0.3 (b) Li plating and stripping in Li|Li₃YBr_5.7F_0.3|Li symmetric cells at 0.1 mA cm⁻² with the capacity of 0.1 mAh cm⁻². Reproduced with permission [156]. Copyright (2021), Wiley‐VCH. (c) Schematic of the gradient structural interphase layer generated at the Li|Li_0.388Ta_0.238La_0.475Cl₃ interface. (d) Cycling stability of the Li|Li_0.388Ta_0.238La_0.475Cl₃|Li symmetric cell. Reproduced with permission [108]. Copyright (2023), Springer Nature. (e) Molecular orbital theory diagram of an [LuCl₆]³⁻ octahedron and a [ZrCl₆]²⁻ octahedron. (f) Schematic of Li₂ZrCl₆|Li MCI interphase. (g) Schematic of Li_2.15Zr_0.85Lu_0.15Cl₆|Li self‐limiting SEI layer. Reproduced with permission [94]. Copyright (2025), Wiley‐VCH. (h) Reduction potentials of representative halide solid electrolytes (SE = Li₃YCl₆, Li₃LuCl₆, and Li₃YbCl₆) vs. Li metal. High‐resolution (i) Yb 4d XPS spectra of the Li₃YbCl₆ surface and (j) Lu 4f XPS spectra of the Li₃LuCl₆ surface after contact with Li metal. Reproduced with permission [162]. Copyright (2025), Wiley‐VCH.

Developing halide electrolytes with satisfying Li compatibility beyond F doping still remains a significant challenge. Recently, a LaCl₃‐based lithium superionic conductor Li_0.388Ta_0.238La_0.475Cl₃ (LiTaLaCl) was reported to possess excellent compatibility with lithium metal [108]. This conductor could form a gradient interfacial passivation layer that stabilized the lithium metal electrode, enabling long‐term cycling of a Li‐Li symmetric cell (1 mAh cm⁻²) for over 5000 h (Figure 8d). This electrolyte possesses a defect‐rich nanocrystalline structure, enabling the formation of dense electrolyte pellets with minimal grain boundaries. This structural feature is beneficial for enhancing interfacial stability with lithium metal. XPS analysis of the LiTaLaCl surface revealed a depth‐dependent change in the chemical state of Ta after 50 h of cycling. The reduction product of Ta in surface area is only 13.4%, which is significantly lower than that of In for Li₃InCl₆ (over 60%) [60]. Additionally, in the interphase layer of SSE, the electrochemically reduced Ta⁰ gradually decreases from 13.4% at the surface to 2.2% at a depth of 3 nm, and La shows a similar change. The gradient reduction of metal species at the interface was attributed to the formation of a LiCl interphase layer that acts as an electrically insulating passivation layer. This layer was also suggested to be effective in reducing ground interfacial strains during stripping/lithium plating and protecting the solid electrolyte from Li metal [105, 163]. Additionally, the dense nanocrystalline structure of LiTaLaCl was crucial for enhancing the stability of the Li/SSE interface (Figure 8c); intimate contact between lithium and the SSE was preserved even after 50 cycles, and the SSE surface remained compact.

Furthermore, we have proposed and validated an “orbital gap modulation” strategy to enhance the reduction stability of halide electrolytes [94]. By incorporating lanthanide elements (e.g., Ho and Lu) into the low‐cost Li₂ZrCl₆ electrolyte, we successfully regulated the local electronic distribution, effectively widened the orbital gap of the metal center, and significantly enhanced the reduction stability of the electrolyte (Figure 8e–g). We also revealed that the introduction of p‐block elements (e.g., Al and In) into the electrolyte leads to a narrowed band gap and reduces the reduction stability. This provides a new approach to overcoming the bottleneck of anode compatibility in halide solid electrolytes.

Recent studies have also evaluated the interfacial stability of Li₃YCl₆, Li₃YbCl_6, and Li₃LuCl₆ in contact with lithium metal [162]. These three materials exhibit similar low reduction potentials (approximately 0.7 V vs. Li⁺/Li, Figure 8h). Further XPS analysis revealed the presence of metallic Y⁰ at the interface, which is highly electronically conductive and facilitates the continuous degradation of Li₃YCl₆. In sharp contrast, Li₃YbCl₆ and Li₃LuCl₆ exhibit distinct interfacial reaction mechanisms, which render them kinetically stable when in contact with lithium metal. As shown in Figure 8i,j, optical images confirmed that both electrolytes retained their original white appearance upon contact with lithium metal, indicating effective passivation. Further XPS characterization revealed the formation of electronically insulating YbCl₂ (bandgap = 4.96 eV) at the Li₃YbCl₆ surface (Figure 8i), which plays a critical role in achieving a self‐limiting surface reduction reaction. Similarly, Li₃LuCl₆ forms a passivating LuCl_x layer at the surface upon contact with lithium (Figure 8j). Notably, our recent work has also calculated the kinetic reaction behavior between the Li₄YbCl₆ electrolyte and Li metal, and revealed the kinetically stable reaction mechanism of Yb²⁺ with Li metal, which is consistent with the results reported in experiments [164].

Moreover, the low electronic conductivity (1.74 × 10⁻¹⁰ S cm⁻¹) of LiTaLaCl is also beneficial for suppressing lithium dendrite growth within the SSE particles [86]. An ideal SSE should exhibit high ionic conductivity and low electronic conductivity to ensure the exclusive transport of lithium ions between the anode and cathode. However, studies focusing on the influence of electronic transport in SSEs remain scarce. A recent comprehensive investigation evaluated the electronic conductivities of several representative SSEs [165]. The study examined Li₃PS₄, Li₇La₃Zr₂O₁₂, and Li₃YCl₆, and reported that Li₃PS₄ exhibited the highest electronic conductivity (3.0 × 10⁻⁹ S cm⁻¹), followed by Li₃YCl₆ (4.5 × 10⁻¹⁰ S cm⁻¹) and Li₇La₃Zr₂O₁₂ (6.2 × 10⁻¹⁰ S cm⁻¹). It was also demonstrated that conventional DC polarization methods using ion‐blocking electrodes severely overestimate the electronic conductivity of SSEs. By employing improved techniques, the authors were able to measure lower electronic conductivities and emphasized the need for more fundamental studies to identify effective strategies for further minimizing the intrinsic electronic conductivity of SSEs.

4.4. Volume Change of Electrolytes and Pressure Optimization

Notably, morphological changes in lithium metal during current loading represent a significant challenge due to the substantial volume fluctuations of the Li electrode. Changes are likely to occur at the Li/SSE interface when the stripping current density exceeds a critical value. The formation and growth of interfacial pores and their correlation with interfacial impedance have been investigated using operando impedance spectroscopy and ex situ SEM [80]. In situ x‐ray computed tomography revealed that pore formation is the primary cause of structural degradation at both Li/SSE and Na/SSE interfaces [166]. Atomic transport at the interface is expected to result in the formation of a disordered surface layer with restricted structural relaxation.

Volumetric changes in halide electrolytes during their reaction with lithium metal attracted little attention. A recent study investigated these effects using Li₃InCl₆, Li₂ZrCl₆, and Li₃YCl₆ as representative examples [107]. The molar volume changes were calculated to be: Li₃InCl₆ (141.13 cc/g) > Li₂ZrCl₆ (42.78 cc/g) > Li₃YCl₆ (32.93 cc/g). The study revealed that the reduction of halide electrolytes proceeds through multiple intermediate stages, each associated with distinct volumetric changes. The observed minimal pressure change in Li₃InCl₆ is due to different mechanical properties of the evolving Li‐In alloy and InCl. Li₂ZrCl₆ exhibited a larger molar volume change than Li₃YCl₆. Moreover, experimental results demonstrated that the rapid reaction kinetics between Li metal and Li₂ZrCl₆ led to the formation of a porous solid electrolyte interphase. This porous interphase plays a critical role in regulating both electronic and ionic transport at the interface and influences the subsequent degradation kinetics of Li₂ZrCl₆. To isolate the impedance contributions of interphase microstructure and porosity from constriction resistance effects, symmetrical cells were tested under incremental pressures (1 MPa steps). The impedance changes of solid electrolytes under varying pressures were recorded, as shown in Figure 9a. For Li₂ZrCl₆, impedance decreased with increasing pressure, whereas for Li₃InCl₆ and Li₃YCl₆, it remained nearly unchanged. This suggests that the interphase formed on Li₂ZrCl₆ possesses higher porosity and is the primary contributor to its overall impedance. The interphase microstructure and porosity at the Li|SSE interface were quantitatively analyzed using focused ion beam scanning electron microscopy (FIB‐SEM) (Figure 9b). The results indicated that the interphase of Li|Li₃YCl₆ exhibited the lowest porosity (2.12%), followed by Li|Li₃InCl₆ (6.99%), while Li|Li₂ZrCl₆ exhibited the highest porosity at 17.64%. These findings were consistent with the impedance spectroscopy results, confirming that the elevated impedance of Li|Li₂ZrCl₆| symmetrical cell primarily arises from the porous microstructure of its interphase.

FIGURE 9.

(a) Pressure and impedance evolution with time for Li|Li₃InCl₆|Li, Li|Li₂ZrCl₆|Li, and Li|Li₃YCl₆|Li symmetric cells. (b) FIB‐SEM images of interphase formed during the chemical reaction between the Li metal and Li₃InCl₆, Li₂ZrCl₆, Li₃YCl₆ SSE, respectively. (c) Impedance and overpotential growth in Li|Li₂ZrCl₆|Li cell at variable pressure and constant stack pressure, respectively. Reproduced with permission [107]. Copyright (2024), American Chemical Society. (d) The electroplating/stripping process of lithium metal symmetric cells under different operating pressures, reproduced with permission [167]. Copyright (2019), Wiley‐VCH. (e) The influence of applied stack pressure on the generation and propagation of lithium dendrites, reproduced with permission [77]. Copyright (2023), Springer Nature.

The microstructure, porosity, ionic and electronic conductivities of the solid electrolyte interphase govern the subsequent reactions between the halide solid electrolyte and lithium metal. Additionally, the applied pressure during electrochemical cycling influences the microstructure and porosity of interphase. Figure 9c presents the variations in overpotential and impedance under both variable and constant stack pressures. A significant discrepancy in the increase of overpotential and impedance was observed between these two conditions. Specifically, the increase in overpotential and impedance in cells operated under variable pressure conditions was nearly ten times higher than that under constant pressure. This finding suggests that the porosity of the interphase is a primary contributor to the increased interfacial impedance and reduced interfacial capacitance at the Li|Li₂ZrCl₆ interface. Analysis of chemical, electrochemical, and mechanical pressure data reveals that chemo‐mechanical interactions play a pivotal role in determining the ultimate interphase structure and porosity, thus impacting the stability of the SSE|Li metal interface.

Pressure not only affects the morphology of the interphase but also plays a critical role in the behavior of lithium metal. Lithium metal exhibits significant ductility at room temperature, making it highly susceptible to deformation under applied pressure. Consequently, its electrochemical performance is significantly influenced by the applied stacking pressure [168]. Previous studies have investigated the lithium plating/stripping behavior in symmetric cells under various applied pressures [167]. As shown in Figure 9d, under an applied pressure of 5 MPa, lithium metal maintains intimate contact with the SSEs throughout cycling, without mechanical creep or dendrite formation. The symmetric cell demonstrates stable cycling exceeding 1000 h without short‐circuiting. When the applied pressure is increased to 25 MPa, lithium metal begins to infiltrate the electrolyte through surface pores without complete penetration. The intrusion reduces the interelectrode distance, thereby lowering the initial overpotential and creating preferential sites for lithium deposition during subsequent cycling. After 48 h of plating/stripping cycling, uncontrolled dendritic growth ultimately results in cell failure. At a higher applied pressure of 75 MPa, lithium metal mechanically penetrates the entire SSE via creep, bypassing the plating/stripping process, which leads to catastrophic mechanical failure of the solid‐state electrolyte. This behavior can be attributed to the fact that 75 MPa exceeds the yield strength of lithium metal by more than two orders of magnitude. Additionally, the solid‐state electrolyte with a porosity of 18% and inadequate densification offers internal pathways for lithium creep.

A recent study further explored the effect of applied pressure on lithium dendrite formation and propagation [77]. Simulations of lithium dendrite growth during plating/stripping at a constant capacity reveal that the dendrite length varies significantly with the applied pressure (Figure 9e). When the applied pressure approaches zero, lithium dendrite formation is effectively suppressed. In contrast, at elevated pressures (e.g., 7 MPa), lithium is mechanically driven into the SSE, thereby accelerating dendrite propagation and crack formation. To validate this result experimentally, a three‐electrode Li|Li₆PS₅Cl cell was constructed. Under a pressure of 0.1 MPa, lithium plating at 4.0 mA cm⁻² causes a gradual voltage drop during the initial cycles, indicating lithium infiltration and a subsequent decline in impedance. However, the cell eventually short‐circuits after 170 cycles. In comparison, at 7 MPa, the cell fails after only 35 cycles, confirming that higher pressure significantly accelerates short‐circuiting. The high ductility of lithium metal underscores the importance of effective stress relaxation in enabling dendrite‐free operation of ASSLBs.

Overall, the optimization of anode‐halide SSE interfaces relies on the synergistic integration of mechanical compliance, chemical stability, ionic conductivity, and electronic insulation. Among the available strategies, anode engineering and interfacial buffer layer construction are the most widely adopted approaches and exhibit broad compatibility across different systems. Alloy anodes can regulate lithium deposition behavior while improving chemical stability and interfacial wettability, thereby suppressing dendrite growth and interfacial failure [118, 119, 120, 121, 169]. However, it inevitably compromises energy density [126, 170]. Accordingly, alloy anodes are more suitable for scenarios where ultra‐high energy density is not the primary requirement. In contrast, interfacial buffer layer construction offers immediate and pronounced benefits in stress accommodation, wettability enhancement, and electronic blocking, owing to its tunable mechanical and chemical characteristics. They also show strong adaptability across a wide range of electrolyte systems [128, 171]. Nevertheless, long‐term cycling may lead to structural changes and compatibility issues with different electrolytes. This places strict demands on thickness control and composition. A balance is needed among ionic transport, mechanical compliance, and interfacial compatibility. In addition, continuous manufacturing of buffer layers needs to be considered for future large‐scale production. At the same time, precise control of interfacial uniformity and thickness remains challenging. Another strategy is to form self‐limiting SEI layers by tuning electrolyte composition or electronic structure. This approach can improve interfacial stability and regulate ionic transport behavior. However, electrolyte‐engineering strategies are inherently system‐dependent and highly sensitive to formation conditions [108, 162]. Although it offers a simple cell design and favorable manufacturability, but requires extensive early‐stage development. Pressure regulation can effectively suppress lithium dendrite growth, enhance cycling stability, improve energy density, and promote interfacial contact. Nevertheless, excessive pressure may induce short‐circuits, and precise pressure control requires extra equipment, which increases complexity and cost.

Synergistic progress in materials innovation, interface engineering, and multiscale process integration is expected to enable systematic optimization from material selection and structural design to manufacturing processes. This will allow unified control over mechanical, chemical, and electrochemical properties. Such advances will provide a solid foundation for the development of high‐performance, long‐lifespan ASSLBs.

5. Advanced Characterization Techniques for ASSLBs

Advanced characterization methodologies could benefit a deeper understanding of the dynamic interfacial evolution, lithium dendrite growth, and ion transport mechanisms. In this section, we highlight several of the state‐of‐the‐art techniques for this investigation, including x‐ray absorption spectroscopy (XAS), x‐ray computed tomography (XCT), X‐ray photoelectron spectroscopy (XPS), solid‐state nuclear magnetic resonance (SS‐NMR), cryo‐transmission electron microscopy (Cryo‐TEM), and neutron depth profiling (NDP).

5.1. X‐Ray‐Based Techniques

XAS is an element‐specific and nondestructive technique used to probe interfacial chemical states and degradation mechanisms in ASSLBs. Recently, operando X‐ray absorption near‐edge structure (XANES) measurements during lithium plating and stripping have confirmed the chemical stability of vacancy‐rich Li₉N₂Cl₃ SSE against lithium metal (Figure 10a). The operando XANES experiments were conducted within a single vacuum X‐ray absorption spectroscopy chamber using an identical electrochemical cell configuration (Figure 10a(i)). Throughout the entire lithium plating/stripping process at different current densities, XANES spectra were collected in real time, and no discernible spectral evolution was observed over the entire 8 h cycling period (Figure 10a(ii–iv)), confirming a stable interface between the Li₉N₂Cl₃ SSE and lithium metal [93].

FIGURE 10.

(a) Operando XANES studies of chemical stability toward lithium metal. (i) Schematic illustration of the operando XANES studies of chemical stability of vacancy‐rich Li₉N₂Cl₃ during the lithium‐plating/stripping process and the configuration of the operando cell for XANES. (ii) Operando XANES spectra at Cl K‐edge with (iii) the corresponding discharge/charge voltage profiles of the Li cycling in the operando cell at several current densities and (iv) with first derivative mapping of vacancy‐rich Li₉N₂Cl₃ during the lithium‐plating/stripping process. Reproduced with permission [93]. Copyright (2023), AAAS. (b) Operando XCT of the Li|Li₆PS₅Cl|Li₃ScCl₆|Li₆PS₅Cl|Li cell before plating and after plating at different current densities. Reproduced with permission [171]. Copyright (2024), Elsevier Inc. (c) (i) Schematic illustration of operando X‐ray photoelectron spectroscopy. (ii) Zr‐3d XPS spectrum of the Li₂ZrCl₆ before and after Li deposition. (iii) Zr‐3d XPS spectrum of the Li_2.15Zr_0.85Lu_0.15Cl₆ before and after Li deposition. Reproduced with permission [94]. Copyright (2026), Wiley‐VCH. (d) Volume view of 3D ⁷Li MRI images of cycled without shorting and shorted Li₇La₃Zr₂O₁₂. Reproduced with permission [88]. Copyright (2025), Springer Nature. (e) Schematic of operando NMR spectroscopy for ASSLBs. Reproduced with permission [172]. Copyright (2026), American Chemical Society. (f) Cryo‐TEM images of the formed interphase layer between individual lithium dendrites and Li₆PS₅Cl electrolyte and their corresponding FFT patterns at (i) 25°C and (ii) 60°C. Reproduced with permission [173]. Copyright (2022), American Chemical Society. (g) Setup and cell configurations for the in situ NDP measurement. Reproduced with permission [174]. Copyright (2017), American Chemical Society. (h) Time‐resolved lithium concentration profiles of Li|Li₂ZrCl₆|Li and Li|Li_2.15Zr_0.85Lu_0.15Cl₆|Li symmetric cells. Reproduced with permission [94]. Copyright (2026), Wiley‐VCH.

XCT is a nondestructive technique that uses x‐ray attenuation to reconstruct three‐dimensional structures. It allows direct observation of internal features and buried interfaces that are not accessible with conventional microscopy. In ASSLBs, this method is widely applied to study interfacial degradation, dendrite growth, and mechanical failure in solid electrolytes during electrochemical cycling. For instance, XCT was employed to investigate crack evolution in multilayer Li₆PS₅Cl|Li₃ScCl₆|Li₆PS₅Cl SSE (Figure 10b). The results reveal that cracks preferentially deflect along mechanically weak interfaces between dissimilar electrolytes, and such crack‐deflection mechanisms effectively suppress dendrite propagation [171].

XPS is a surface‐sensitive technique that can provide quantitative information and chemical‐state analysis of elements at solid‐solid interfaces. Recent progress in operando XPS has enabled real‐time monitoring of potential distribution and species evolution during electrochemical cycling. For example, operando XPS has been applied to investigate interfacial reaction products between the electrolyte and the lithium anode (Figure 10c(i)) [94]. Upon contact with lithium, Li₂ZrCl₆ forms a series of reduced by‐products, including metallic Zr⁰, which leads to continuous growth of the interphase layer (Figure 10c(ii)). In contrast, no metallic Zr⁰ species were detected in the XPS spectra of Li_2.15Zr_0.85Lu_0.15Cl₆ during the initial stage of lithium deposition to 1 h, indicating that the reduction of Zr⁴⁺ can be effectively suppressed (Figure 10c(iii)).

5.2. Solid‐State Nuclear Magnetic Resonance

SS‐NMR spectroscopy is a powerful technique for probing the chemical environments of a wide range of magnetically active nuclei and is particularly well suited for investigating subtle variations in local structures. NMR resonances are highly sensitive to the dynamics of mobile ions (such as Li⁺ and protons), while being largely insensitive to grain boundary effects [133].

Three‐dimensional ⁷Li magnetic resonance imaging (MRI) was performed to investigate the spatial distribution of lithium dendrites in Li|Li₇La₃Zr₂O₁₂|Li symmetric cell [88]. In cells at an intermediate cycling stage with a stable voltage profile, clustered lithium microstructures are observed on the surface of Li₇La₃Zr₂O₁₂. In the short‐circuited cell, 3D MRI results reveal the formation of a dense dendritic network within the Li₇La₃Zr₂O₁₂ electrolyte (Figure 10d). In addition, researchers have developed a quantitative operando NMR technique compatible with a high stack pressure of 95 MPa (Figure 10e) [172]. This method enables the resolution of stepwise lithiation and delithiation mechanisms of silicon anodes in all‐solid‐state batteries. It further reveals asymmetric reaction pathways and temperature‐dependent behavior, providing atomic‐scale insight for the performance optimization of silicon‐based all‐solid‐state batteries.

Other studies have also demonstrated the use of NMR to track dendrite growth and probe lithium‐ion transport across interfaces [175, 176]. Collectively, these findings establish solid‐state NMR as a distinctive, nondestructive characterization technique capable of simultaneously quantifying lithium species, resolving transport pathways, and elucidating the dynamic mechanisms underlying interfacial failure.

5.3. Cryo–Transmission Electron Microscopy

Cryo‐transmission electron microscopy (cryo‐TEM) enables nanoscale characterization of air‐ and moisture‐sensitive lithium dendrites and electrode‐electrolyte interfaces under near‐native conditions, thereby providing critical insights into the structure, composition, and evolution mechanisms of lithium dendrites and interfacial layers. A stable electrode‐electrolyte interface is essential for realizing high‐energy‐density ASSLBs; however, its nanoscale behavior remains poorly understood. Using cryo‐TEM, researchers systematically investigated the nanostructure of interfacial layers formed between individual lithium dendrites and Li₆PS₅Cl electrolyte at different temperatures [173]. At 25°C, a well‐defined crystalline Li₂S interphase with a thickness of approximately 12 nm was observed at the interface, exhibiting effective passivation behavior (Figure 10f(i)). In contrast, at 60°C, the interphase evolves into a polycrystalline Li₂S structure accompanied by an order‐disorder phase transition. The elevated temperature leads to a significant increase in interphase thickness and grain boundary density, resulting in a pronounced rise in interfacial resistance (Figure 10f(ii)).

5.4. Neutron Depth Profiling (NDP)

Neutron depth profiling (NDP) is a non‐destructive analytical technique based on neutron‐nucleus interactions, enabling depth‐resolved quantification of lithium concentration from the surface to the bulk of the electrolyte. This method provides direct experimental insights into lithium diffusion, interfacial deposition behavior, and ion transport mechanisms. Researchers have employed in situ NDP to monitor lithium plating/stripping processes in garnet‐type solid electrolytes, allowing real‐time tracking of interfacial evolution (Figure 10g) [174, 177]. Furthermore, operando NDP has been employed to monitor the dynamic evolution of lithium concentration profiles during lithium plating in three representative solid electrolytes—LiPON, Li₇La₃Zr₂O₁₂, and amorphous Li₃PS₄. The results reveal that the relatively high electronic conductivity is the origin of lithium dendrite formation within solid electrolytes [86]. In our recent work, ex situ NDP was employed to investigate the evolution of lithium concentration distributions at the interfaces of Li|Li₂ZrCl₆ and Li|Li_2.15Zr_0.85Lu_0.15Cl₆ cells during cycling [94]. For the Li_2.15Zr_0.85Lu_0.15Cl₆‐based symmetric cells, only a slight increase in lithium concentration was observed within the probed depth range over cycling, indicating a kinetically stable interface. In contrast, the Li₂ZrCl₆‐based symmetric cells exhibit a continuous increase in lithium concentration during cycling, suggesting persistent accumulation of electrochemically inactive lithium at the interface (Figure 10h).

6. Summary and Perspective

In general, rare‐earth elements, including Sc³⁺, Y^3+, and lanthanides, exhibit considerable reduction stability. The reduction potential of typical chloride‐based Li₃MCl₆ follows the trend Y³⁺ < Sc³⁺< In³⁺. SSEs containing group 3 elements may offer a wider electrochemical stability window. This review summarizes the development and current understanding of halide SSEs with respect to their stability against lithium, and introduces the fundamental concepts of interfacial behavior from both thermodynamic and kinetic perspectives. Most halide SSEs are unstable in contact with Li metal, severely limiting their applicability in all‐solid‐state batteries. Designing a stable interface between halide SSEs and Li anodes is essential. Due to limited insight into the electrochemical stability of halide SSEs, additional characterization techniques and analyses are required to improve the understanding of their interfacial behavior with lithium metal anodes. Despite their relatively low ionic conductivity, fluorine‐based halide SSEs have shown promise in forming stable interfaces with Li metal. Enhancing the ionic conductivity of fluorine‐based halide SSEs through partial substitution and polyanion‐coordination strategy may help mitigate the reductive instability issues of halide SSE.

Additionally, improvement strategies are summarized. One effective strategy involves stabilizing the Li|SSE interface through external protective layers. However, compatibility between the protective layer and the halide electrolyte must be addressed. Secondly, alloying strategies with lithium anodes are also effective, but they reduce the operating voltage and energy density and lead to a certain degradation of the battery energy density. Thirdly, the formation of self‐limiting SEI layers via doping or electronic structure modification offers a highly promising approach. Pressure regulation significantly influences lithium metal behavior and overall battery performance, and should be a key focus in future research.

From a future perspective (Figure 11), the use of lithium metal anodes in halide SSE‐based ASSLBs can be enhanced by promoting self‐limiting interfacial reactions and suppressing lithium filament formation. At the cell level, optimizing the balance between performance and applied stack pressure is essential for halide‐based battery systems, while the integration of advanced characterization techniques provides insights into interfacial behavior.

FIGURE 11.

Perspectives on addressing interfacial instability in halide‐based batteries.

Conflicts of Interest

The authors declare no conflicts of interest.

Data Availability Statement

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