New advances in solid-state electrolytes: from halides to oxyhalides

Abstract

The emerging area of halide solid electrolytes has garnered significant interest due to its superior ionic conductivity, broad electrochemical stability range, and strong compatibility with positive electrode materials. In recent times, there has been a continuous emergence of diverse halide systems. This study provides a comprehensive review of the synthesis strategies classification schemes, structural attributes, ion transport mechanisms, and modified means of halides. Notably, the paper delves into a detailed analysis of the structural details and ion conduction mechanisms of several representative halides. Furthermore, this review also examines the intrinsic connection between the recently discovered oxyhalides and halides and figures out the primary obstacles and prospective directions for advancement through a thorough review and analysis.

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Qingtao Wang

Qingtao Wang received his BSc in 2005 from Shandong University and received his PhD in 2010 from Lanzhou Institute of Chemical Physics, Chinese Academy of Sciences. He joined Northwest Normal University in July 2010. His current research interests include: (1) solid electrolytes of sulfides, halides and oxyhalides; (2) electrocatalysts for energy conversion.

Zhenyang Shen

Zhenyang Shen, a master's student in the group of Associate Professor Qingtao Wang at Northwest Normal University, is currently researching solid electrolytes of halides and oxyhalides.

Yongmei Zhou

Yongmei Zhou, a master's student in the group of Associate Professor Qingtao Wang at Northwest Normal University, is currently researching solid electrolytes of halides and oxyhalides.

Ying Liu

Ying Liu received her BE in 1996 from Lanzhou university of technology telecommunication department. She joined Northwest Normal University in June 1996. Her current research interests include: (1) chemical automation instrumentation and detection technology, process control; (2) chemical sensors.

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1 Introduction

With the increasingly serious global energy shortage and environmental pollution issues, replacing traditional, non-renewable fossil fuels with clean, sustainable, and renewable energy has become an inevitable trend. Conventional lithium-ion batteries are susceptible to combustion and detonation owing to their limited thermal stability and low ignition threshold. In contrast, solid electrolytes have garnered attention in the realm of battery technology for their enhanced safety, stability, cost-effectiveness, and extended operational lifespan. This has prompted significant research efforts and ongoing advancements within the scientific community in this domain in recent times. As the core component of the new energy storage device, the solid electrolyte needs to meet the ionic conductivity of more than 10⁻³ S cm⁻¹ at room temperature, while requiring high electrochemical stability and excellent compatibility with the electrode material.^1–9 Among the sulfides, oxides, polymers, halides, and oxyhalides electrolytes that have been explored, oxides have high ionic conductivity of 10⁻³ to 10⁻⁴ S cm⁻¹ and have good chemical stability against lithium metal, but the use of solid oxide electrolytes alone cannot provide satisfactory electrochemical performance due to high boundary resistance due to insufficient interfacial contact.¹⁰ The ionic conductivity of solid electrolytes based on sulfides is similar to that of liquid electrolytes. However, the practical utility of these solid electrolytes is constrained by their limited electrochemical stability window, high susceptibility to moisture, and tendency to react readily with water in the atmosphere, leading to the generation of hydrogen sulfide gas.^11–15 Researchers prefer polymer electrolytes due to their wide electrochemical stability and high flexibility. They are also valued for their cost-effectiveness and straightforward preparation methods, which contribute to their widespread adoption and utilization in commercial settings. However, the ionic conductivity of such electrolytes is generally low at room temperature.¹⁶ Although the electrolyte system based on polyethylene oxide is recognized as a promising member of the polymer application, it usually needs to improve its performance by modifying means such as crosslinking and copolymerization, and its excellent electrochemical performance is often only shown at high-temperature conditions.^17,18

Compared to the three types of electrolytes mentioned above, research on halide solid electrolytes (HSEs) is still relatively scarce. The focus is on achieving high ionic conductivity (>10⁻³ S cm⁻¹) at room temperature, a wide electrochemical window, and good compatibility with cathode materials, which has become a popular research direction.^19,20 The development timeline of halide solid electrolytes is shown in Fig. 1. In 1930, the ionic conductive behavior of LiX (X = Cl, F, Br, I) was reported, but the low ionic conductivity did not attract much attention.²¹ In 2018, Asano and colleagues synthesized Li₃YCl₆ and Li₃YBr₆ with exceptional ionic conductivity at room temperature using mechanical ball milling and annealing processes. The conductivity achieved was several orders of magnitude higher than that of previously reported HSEs. Subsequent research on HSEs has entered a more in-depth stage.^21–23 Although HSEs has excellent electrochemical performance, its widespread application in practical use is hindered by economic barriers and practicality concerns stemming from the expensive metal elements it is composed of. In this context, research on cost control of HSEs is particularly crucial.²⁴ The Li₂ZrCl₆ raw material reported by Ma et al.²⁵ is several orders of magnitude cheaper than the most advanced HSEs, while also exhibiting a high ionic conductivity (0.81 mS cm⁻¹) at room temperature and good moisture resistance. As of 2022, there is still a significant gap in the reported ion conductivity and sulfide of HSEs, which is worth noting that Tanaka and others²⁶ have studied a new oxyhalide LiMOCl₄ (M = Ta, Nb) with an ion conductivity as high as 12.4 mS cm⁻¹. This new material can rival or even surpass the liquid electrolyte in lithium-ion batteries at room temperature. It is considered the most promising candidate for practical applications in all-solid-state batteries (ASSBs).

Fig. 1 Development timeline of halide electrolytes and emerging oxyhalides for ASSBs.

This article provides an overview of the development of halide solid electrolytes and the innovative evolution from halides to oxyhalides. It then details three common synthesis methods and five major classification systems of halides, and delves into four types of crystal structures and their corresponding ion transport mechanisms. Finally, the article discusses modification strategies for halides and reveals the inherent connection between halides and oxyhalides.

2 Categories of halide solid electrolytes

Regarding the classification of halides, there are currently two main classifications: (1) dividing them into three categories based on different types of metal elements:¹⁹ (1) elements from IIIB (M = Sc, Y, La–Lu). (2) Elements from IIIA (M = Al, Ga, Ln). (3) Divalent metal elements (M = Ti, V, Cr, Mn, Fe, Co, Ni, Cu, Zn, Cd, Mg, Pb) in the Li_aMX_b. (2) According to the amount of transition metal elements substituted for Li ions,³ they are divided into (1) Li_a-M-Cl₆, (2) Li_a-M-Cl₄, and (3) Li_a-M-Cl₈ halides (M = transition metal element). Tuo et al.²⁷ extended the classification of the IV based on the research of Zr and Hf elements. With the research on LiTaCl₆ and LiMCl₅Xⁿ⁻_1/n (M = Ta, Nb, Xⁿ⁻ = Cl⁻, F⁻, O²⁻) electrolytes^28,29 a halide of VB metal element has gained public attention. When classifying based on the degree of Li-ion substitution by transition metal elements, different main groups or subgroup elements can be grouped into the same category. Here, we will continue to further classify based on the diverse types of metal elements.

2.1 Crystal structures of the IIIB metal elements (Sc, Y, La–Lu)

The halides of the IIIB metal elements include Li₃MX₆ (X = Cl, Br, I) and LiMF₄. Meyer et al.^19,30,31 confirmed through single-crystal X-ray diffraction data of the ternary chlorides Li₃ErCl₆, Li₃YbCl₆, and Li₃ScCl₆ that they crystallize in three different structural types. They classified (M = Y, Tb–Tm) as a trigonal crystal system with the P3̄m1 space group, (M = Y, Yb, Lu) as an orthorhombic crystal system with the pnma space group, and Sc as having a cubic close-packed anion sublattice with a crystal structure being monoclinic (space group C2/m). Part of the crystal structure of Li₃MCl₆ is shown in Fig. 2a.³² YCl₃ is a representative of MX₃ salts. Its physical properties make it equivalent to the end (rare) member of the lanthanide series,³³ which may also be one of the reasons why Y and some lanthanide metal elements have the same space group structure. With the high stability of the YX₆³⁻ (X = Cl or Br) octahedra and the expectation that each Y can provide two cation vacancies, Asano et al.^22,33 synthesized Li₃YCl₆ (LYC) and Li₃YBr₆ (LYB) with ionic conductivities reaching 10⁻³ S cm⁻¹ at room temperature after cold pressing. LYC is composed of hexagonally close-packed (hcp) anions, with a crystal structure of trigonal symmetry (space group P3̄m1) (Fig. 2b), whereas LYB is composed of cubic close-packed (ccp) anions, with all cations located in octahedral positions, exhibiting a monoclinic structure (space group C2/m) (Fig. 2c). Sun et al.³⁴ synthesized Li_xScCl_3+x using various ratios of LiCl and ScCl₃. When x = 3, the room temperature ionic conductivity reached 3.02 mS cm⁻¹. Similar to Y, ScCl₆³⁻ also has an octahedral structure, but it is based on a monoclinic structure with ccp anion arrangement (Fig. 2d). Subsequently, a disordered spinel structure Li₂Sc_2/3Cl₄ with an ionic conductivity of 1.5 mS cm⁻¹ at room temperature was reported.³⁵ Li is randomly distributed on tetrahedral and octahedral sites available in the lattice. Similar to the previously reported spinel Li₂MgCl₄,^36,37 Li₂MgCl₄ is composed of edge-sharing (Mg₁/Li₂)Cl₆ octahedra, while Li₂Sc_2/3Cl₄ contains edge-sharing (Sc₁/Li₄)Cl₆ octahedra.

Fig. 2 (a) Crystal structures outlined with the unit cells for trigonal and monoclinic Li₃MCl₆ (or Li₂MCl₆).³² (b) The triangular structure of Li₃YCl₆ and (c) the monoclinic structure of Li₃YBr₆.²² Reprinted with permission from ref. 22 and 32. Copyright (2018 and 2021) Wiley. (d) Typical crystal structure of Li₃ScCl₆.³⁴ Reprinted with permission from ref. 34. Copyright (2020) American Chemical Society.

Lanthanum (La)^38,39 in the lanthanide series has garnered attention as a research hotspot due to its distinctive UCl₃-type crystal structure (space group P63/m). This structure sets it apart from the typical hexagonal close-packed (hcp) and face-centered cubic close-packed (ccp) anion frameworks. The fundamental unit of this framework is the lanthanum chloride nine-coordinated “three vertex three prism” unit, distinguished by the abundance of channel sites capable of hosting monovalent cations and offering migration pathways for them.

Xu et al.⁴⁰ studied the new spinel structure Li₃MI₆ (M = Sc, Y, La) with octahedral Li occupancy using density functional theory (DFT) calculations. They predicted that Li₃MI₆ is monoclinic (space group C2) with a layered structure stacked O4 (octahedron 4) along the c-direction (Fig. 3a). There are two main types of stable lithium occupation sites: octahedral sites found in transition metal layers and those within pure lithium layers. The LiI₆ octahedra in transition metal layers have a slightly larger volume and longer Li–I bond length than those in pure lithium layers. This setup, with around one-third of the octahedral sites empty, promotes the diffusion of Li⁺ ions. Zeier et al.³¹ successfully synthesized Li₃ErI₆. The Pawley fitting provided information on the lattice and overall symmetry, and only the C2 and C2/c space groups obtained reflections that matched both the observed and predicted data. Although these two space groups are highly similar and have the same polyhedral connectivity, C2/c exhibits superior symmetry. By combining synchrotron X-ray diffraction and neutron diffraction techniques, it was determined that Li₃ErI₆ belongs to a monoclinic structure (space group C2/c) (Fig. 3b). Compared to Li₃ErCl₆, it was found that Li₃ErI₆ has a more flexible lattice and faster ion transport.

Fig. 3 (a) O4 stacking sequence of Li₃MI₆ (M = Y, Sc, La) cells.⁴⁰ (b) Structural frame and full structure of Li₃ErI₆.³¹ Reprinted with permission from ref. 31 and 40. Copyright (2019 and 2020) American Chemical Society.

All ternary halides of Li₃MBr₆ exhibit a monoclinic structure.^19,40,41 Zeier et al.⁴² utilized temperature-dependent neutron diffraction to investigate Li₃YBr₆ with the C2/m space group. This structure is distinct from Li₃YCl₆ and can be described as having a layered configuration with the (002) crystal plane.

2.2 Crystal structures of the IVB metal elements (Zr, Ti, Hf)

Research on the IVB metal elements of HSEs is currently focused on Li₂ZrCl₆,²⁵ which is popular due to its relatively high ionic conductivity at room temperature and considerable cost-effectiveness. The Li₂ZrCl₆ (α-LZC) after planetary ball milling has a P3̄m1 space group that is remarkably similar to Li₃YCl₆. After annealing at 350 °C, in addition to higher crystallinity, the crystal structure undergoes a change from P3̄m1 to a structure similar to Li₃InCl₆ (C2/m) called β-LZC (Fig. 4a). However, the ionic conductivity decreases by two orders of magnitude. Recently reported material Li₃TiCl₆ (LTC)⁴³ with combined electrolyte, positive electrode, and negative electrode functions, after ball milling treatment, exhibits weaker diffraction peaks compared to LTC after annealing at 300 °C. From the X-ray diffraction (XRD) patterns, there is no fundamental difference between the two samples, indicating their consistency in overall crystal structure. Refinement of the XRD of LTC after ball milling and annealing reveals a close resemblance to the crystal structure of Li₃InCl₆ in the C2/m space group (Fig. 4b). Li⁴⁴ and colleagues obtained another ideal layered structure Li₂MCl₆ (M = Zr, Hf) using the Stochastic Surface Walking global optimization combined with global Neural Network potential (SSW-NN) method. The structure consists of two layers formed by the sharing of a common edge between LiCl₆⁵⁻ and Zr/HfCl₆⁵⁻ octahedra, displaying a novel layered structure distinct from previous reports (Fig. 4c).

Fig. 4 (a) The crystal structures of α-LZC (left) and β-LZC(right).²⁵ (b) The structural model of LTC superimposed with the Li-ion potential map.⁴³ Reprinted with permission from ref. 25 and 43. Copyright (2021 and 2023) Springer Nature. (c) The ideal layer structure of Li₂MCl₆ (M = Zr, Hf) crystal.⁴⁴ Reprinted with permission from ref. 44. Copyright (2022) American Chemical Society.

2.3 Crystal structure of the IIIA metal elements (Al, Ga, In)

Research on aluminum halides can be traced back to the 1970s.⁴⁵ Weppner et al.⁴⁶ calculated that NaAlCl₄ has an orthorhombic structure, while LiAlCl₄ and KAlCl₄ have monoclinic structures. Gregory et al.⁴⁷ first synthesized the iodide LiAlI₄ (monoclinic, space group P21/c), and room temperature diffraction data confirmed that LiAlI₄ and LiAlBr₄ are isostructural with chloride analogues. Li⁺ and Al³⁺ ions occupy octahedral and tetrahedral interstices, respectively. The extended structure is composed of distorted LiX₆ octahedra and AlX₄ tetrahedra, with the two octahedra connected at shared edges to form “Li₂X₁₀ dimers” (Fig. 5a–c). Regarding the IIIA elements, the current research focus is still on In element. In 1998, Koji Yamada et al.⁴⁸ first reported Li₃InBr₆, which undergoes a phase transition at high temperatures with high ionic conductivity and maintains phase stability when returning to low temperatures. For the high-temperature phase Li₃InBr₆,⁴⁹ it exhibits a monoclinic structure with space group C2/m, and in the structure, In³⁺ is in a disordered state at octahedral positions. Yamada et al.⁵⁰ found that diffusion of In³⁺ and Li⁺ occurs throughout the crystal when the sample is kept at 420 K for two hours. One-third of the cation sites are occupied by vacancies due to substitution, indicating that In³⁺ is a good candidate for introducing vacancies into lithium-based solid electrolytes. Although the structure of the low-temperature phase is unknown, it is believed to be an ordered crystal based on the signals detected in this phase. In 1986, Wolfgang and his team⁵¹ determined the crystal structure of LiGaX₄ (X = Cl, F, Br) using single crystal X-ray methods. The compound possesses LiX₆ octahedra and GaX₄ tetrahedra, similar to the structure of LiAlCl₄. Interestingly, Jung and others⁵² designed an amorphous 2LiCl–GaF₃ electrolyte, which exhibits a high ionic conductivity of 3.6 mS cm⁻¹ at room temperature and high viscoelasticity. This unique “clay” property is believed to be due to the formation of gallium-containing, mobile, and complex anions Ga-F-X (X representing different halogens).

Fig. 5 (a) Crystal structure of LiAlI₄ (P21/c) projected along the b-axis. (b) Showing the linkage between an LiX₆ octahedron and neighboring AlX₄ tetrahedra. (c) Highlighting the positions of the Li⁺ cations with respect to the isolated haloaluminate anions: Li (4e, green spheres), Al (4e, light blue spheres), and I (4e, purple).⁴⁷ Reprinted with permission from ref. 47. Copyright (2021) American Chemical Society. (d) X-ray diffraction patterns of the as-prepared ball milled and annealed Li₃InCl₆ samples. (e) Structure of annealed-Li₃InCl₆. (f) Schematic unit showing typical Li⁺ location around the InCl₆³⁻ octahedra.⁵⁵ Reprinted with permission from ref. 55. Copyright (2019) Royal Society of Chemistry.

The Li₃InCl₆ (LIC) reported by Lutz et al.⁵³ was the best lithium-ion conductor at that time, with an ionic conductivity of 0.2 mS cm⁻¹ at 300 °C. Schmidt et al.⁵⁴ obtained colorless single crystals of LIC through the Bridgman-type method using a mixture of LiCl and InCl₃ in a 3 : 1 molar ratio, highlighting the potential of indium-based halides. Sun et al.⁵⁵ reported a high stability and high crystallinity halide electrolyte LIC, with a monoclinic crystal structure (space group C2/m). After annealing, it exhibits a high ionic conductivity of 1.49 mS cm⁻¹ at room temperature. The XRD spectrum after annealing shows higher crystallinity compared to before annealing. The structure is essentially a distorted rock-salt LiCl structure, with octahedral interstices occupied by Li⁺, In³⁺, and vacancies in a ratio of 3 : 1 : 2. The anions are arranged in a ccp structure, which has lower symmetry compared to the body-centered cubic sublattice of the anions (Fig. 5d–f). Subsequently, the group prepared LIC with ionic conductivity up to 2.04 mS cm⁻¹ at room temperature in an aqueous solution system.⁵⁶ The crystal structure of this LIC remains the same as the samples synthesized by ball milling and sintering methods. It is worth noting that the ionic conductivity of this LIC can still be restored after removing the adsorbed water.

2.4 The amorphous state of the VB metal elements (Nb, Ta)

The main molecular formula of the VB metal elements is LiMCl₆, and current reports on it are all amorphous. Ishiguro et al.⁵⁷ found that TaCl₅ has the ability to vitrify various lithium salts (LiF, LiCl, LiBr, LiI, Li₂O, LiOH, Li₂O₂, Li₂S), and can produce a glassy phase superionic conductor LiTaCl₅Xⁿ⁻_1/n with an ion conductivity as high as 10.95 mS cm⁻¹ at room temperature. By observing XRD, it was clearly found that there were no crystalline peaks in LiTaCl₆, confirming that the substance is amorphous. In addition, a similar situation was also observed in the XRD of LiNbCl₆. Li et al.²⁹ used ab initio molecular dynamics (AIMD) to simulate annealing and create a glassy model of LiTaCl₆ (Fig. 6). In the model, high-density Cl⁻ aggregate with Li⁺ ions bridged by Cl⁻, with an average of 4.8 Cl⁻ surrounding each Li⁺. The coordination number of Cl⁻ ions around Ta is 5.7, with most Ta atoms existing in the form of TaCl₆⁻ and some in the form of TaCl₅. The detailed parameters of the solid electrolytes of the aforementioned four types of halides can be found in Table 1.

Fig. 6 (a) Structure of glassy LiTaCl₆: snapshot from AIMD. (b) Coordination polyhedra and LiCl_n network in the snapshot.²⁹ Reprinted with permission from ref. 29. Copyright (2024) Wiley.

Table 1 Ionic conductivity and crystal structure of some metal halides

Group	Material	Structure	Conductivity (mS cm^−1)	Ref.
IIIB	Li3ScCl6	Monoclinic, C2/m	3 at 25 °C	35
	Li2Sc2/3Cl4	Cubic	1.5 at 25 °C	36
	Li3YCl6	Trigonal	0.51 at 25 °C	22
	Li3YCl6	Orthorhombic, Pnma	3.07 at 450 °C (calculate)	27
	Li3YBr6	Monoclinic, C2/m	0.72 at 25 °C	22
	Li3LaI6	Monoclinic, C2	1.23 at 25 °C (calculate)	44
	Li3ErCl6	Trigonal, P3̄m1	3 at 25 °C (calculate)	19
	Li3ErI6	Monoclinic, C2/m	0.65 at 25 °C	32
IVB	Li3TiCl6	Monoclinic, C2/m	0.115 at 25 °C	44
	Li3TiCl6	Monoclinic, C2/m	1.04 at 300 °C	44
	Li2ZrCl6	Trigonal, P3̄m1	0.81 at 25 °C	25
	Li2ZrCl6	Monoclinic, C2/m	5.81 × 10^−3 at 350 °C	25
	Li2HfCl6	—	∼0.5 at 25 °C (calculate)	45
VB	LiTaCl6	Amorphous	10.95 at 25 °C	30
	LiTaCl6	Amorphous	7.16 at 25 °C	28
	LiTaCl6	Amorphous	11 at 25 °C (calculate)	61
	LiNbCl6	Amorphous	12.192 at 25 °C (calculate)	29
IIIA	LiAlCl4	Monoclinic	1.2 × 10^−3 at 25 °C	42
	LiAlI4	Monoclinic, P21/c	0.012 at 25 °C	43
	LiAlBr4	Monoclinic, P21/c	0.033 at 25 °C	43
	Li3InCl6	—	0.2 at 300 °C	47
	Li3InCl6	Monoclinic, C2/m	1.49 at 25 °C	48
	Li3InCl6	Monoclinic, C2/m	2.04 at 25 °C	49

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2.5 Crystal structure of divalent metal elements (Fe, Cd, Mg, Co, Zn, V)

In A₂BCl₄, various structures are generated based on the different ionic radius ratios. Electrolytes composed of divalent metal ions are mainly divided into four crystal systems: spinel system (cubic) with the formula Li_aMX₄ (M = Ti, V, Cr, Mn, Fe, Cd, Mg), olivine system (orthorhombic) of Li₂ZnCl₄, distorted Li₂MCl₄ (M = Cr, Fe, Co), and Suzuki structure (Li₆MX₈, M = V, Fe, Co). Please refer to Table 2 for more details. Spinel exists in two types: normal and inverse. In the normal spinel, the lithium ion conduction properties mainly depend on specific positions in the octahedral structure, while the inverse spinel utilizes both tetrahedra and octahedra in the lithium ion conduction process. This is also an important factor leading to the lower ionic conductivity of normal spinel compared to inverse spinel. Additionally, an attractive feature of spinel is that various transition metal cations can be stabilized in different distorted structures, offering possibilities for various conduction mechanisms.⁵⁸

Table 2 Crystal structure and ionic conductivity of some divalent metal halides

Material	Structure	Conductivity (mS cm^−1)	Ref.
Li2MgCl4	Inverse spinel	0.14 at 400 °C	60
Li2CdCl4	Inverse spinel	0.3 at 400 °C	60 and 65
Li1.9Cd1.05Cl4	Inverse spinel	0.35 at 400 °C	61
Li2MnCl4	Inverse spinel	0.14 at 400 °C	60
Li2FeCl4	Inverse spinel	0.63 at 400 °C	60
Li2ZnCl4	Normal spinel	1 × 10^−3 at 200 °C	63
Li6FeCl8	Suzuki	2 at 200 °C	75
Li6CoCl8	Suzuki	0.7 at 200 °C	77
Li2ZnCl4	Olivine	0.2 at 280 °C	64
Li2CrCl4	Distorted spinel	1.5 at 400 °C	71
Li2CoCl4	Distorted spinel	1 at 300 °C	71 and 72

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Research on Li₂MCl₄ can be traced back to 1975 when Van Loon et al.⁵⁹ reported four ternary chlorides Li₂MCl₄ (M = Mg, Mn, Fe, Cd). Calculations showed that they have a spinel structure, with half of the lithium ions existing in tetrahedra surrounded by chlorine ions within the crystal, while the other half of the lithium ions and M ions are distributed in octahedra. This distribution is believed to resemble a NaCl structure with cationic defects. Lutz et al.⁶⁰ found that when Li₂MgCl₄ is heated to 535 °C, Li₂MnCl₄ (460 °C), and Li₂CdCl₄ (385 °C), the spinel structure reversibly transforms into a defective cubic structure. In addition, these researchers^61–64 measured the ion conductivity of Li₂MCl₄ between room temperature and 500 °C. At high temperatures, the ternary chlorides become highly disordered, and the ion conductivity continues to increase. For instance, the ion conductivity of Li₂CdCl₄ increased from 0.12 ± 0.03 at 300 °C to 1.3 ± 0.1 S cm⁻¹ at 500 °C. Li_1.9Cd_1.05Cl₄ exhibits an ion conductivity as high as 0.35 S cm⁻¹ at 400 °C, but its decomposition potential is only 1.1–1.3 V, limiting its application in batteries.

Van Loon, Sassmannshausen and others^65,66 found that Li₂ZnCl₄, after low-temperature modification, crystallizes differently from other known halides with a spinel structure. All Li ions are located at octahedral sites, while Zn ions are at tetrahedral sites (normal spinel structure). After high-temperature modification (215 °C), it transforms into an olivine structure. Lutz and others^66–68 conducted a more detailed study on this structure. They found that a colorless single crystal Li₂ZnCl₄ with a normal spinel structure can be obtained by heating a stoichiometric mixture of LiCl and ZnCl₂ in a sealed glass ampoule at 403 K for two months. Heating at 532 K for a week followed by quenching in ice water results in olivine-type Li₂ZnCl₄ (crystal structure shown in Fig. 7). Additionally, the transformation from normal spinel to olivine also changes the space group from Fd3m to Pnma. In the olivine structure,⁶⁹ Cl ions are densely packed in a hexagonal arrangement around Li ions in the octahedra. The activation energy for conduction through octahedra is slightly lower than in the normal spinel structure. The lower calculated density (d = 1.529) compared to normal spinel (d = 1.601) indicates that the olivine structure is more open.

Fig. 7 (a) Spinel (cubic) and (b) olivine (orthorhombic) structures of the Li₂ZnCl₄ compound.⁶⁷ Reprinted with permission from ref. 67. Copyright (2022) Elsevier.

Kanno and others⁷⁰ discovered that Li₂MCl₄ (M = Cr, Fe, Co) has a distorted spinel structure. The distortion is caused by a 1 : 1 ordering of Li and divalent metal ions on the octahedral sites. This distorted structure is defined as a deficient ordered rock-salt type SnMn₂S₄ (space group Cmmm).^71,72 Importantly, the distorted structure results in a significant decrease in ionic conductivity.

The Suzuki phase was proposed by Suzuki in 1960.⁷³ This structure is a rock salt derivative with an upper structure of LiCl type, where cations (Li⁺ and M²⁺) are arranged in an ordered manner, and there are vacancies at octahedral positions.⁷⁴ Unlike the anti-perovskite structure, the MX₆ octahedra in the Suzuki structure are isolated from each other. According to reports by H. D. Lutz and others,⁷⁴ the formation of the Suzuki phase is only possible when the radius ratio R^II_M/R^I_M < 1, such as in the case of Fe, V, and Co. Unlike LiCl,^74–76 the presence of vacancies in Li₆MCl₈ promotes faster conduction of Li⁺ ions. Experimental results show that when measuring ionic conductivity at 350 °C, the arrangement of cations changes from ordered to disordered, further confirming the critical influence of cation arrangement at octahedral positions on ion migration.

In general, divalent metal electrolytes encompass a wide range of electrolyte types and intricate structures, with overall ion conductivity not being very optimal. The most notable structure is the spinel structure. Unfortunately, research on this particular structure has been exceedingly limited in recent years. Consequently, an urgent challenge currently confronting us is to conduct more thorough and extensive research on these metals.

3 Synthesis method of halide solid electrolytes

Currently, the preparation methods for HSEs typically include solid-phase methods (ball milling or ball milling and sintering) and liquid-phase methods. The two synthesis methods of Li₃InCl₆ are exemplified in Fig. 8. In addition, there are also a few reports on gas-phase methods⁷⁷ and ammonia-assisted wet chemical methods.⁷⁸

Fig. 8 (a) Schematics of the synthesis of the Li₃InCl₆ solid electrolytes, including ball-milling (R₁) followed by an annealing (R₂) step. (b) Illustration of water-mediated synthesis route for Li₃InCl₆ and the reversible interconversion between the hydrated Li₃InCl₆·xH₂O and dehydrated Li₃InCl₆.⁵⁶ Reprinted with permission from ref. 56. Copyright (2019) Wiley.

3.1 Solid-phase reaction synthesis (ball milling, ball milling and sintering)

Halides exhibit strong hygroscopic properties, and the majority of solid-phase synthesis methods necessitate an inert gas atmosphere, such as argon or nitrogen. Early synthesis processes often require a significant amount of time, as seen in the case of Li₂MCl₄ (M = Mg, Mn, Cd),⁶¹ where the raw materials must be heated for a week after grinding. Similarly, Li₃MBr₆ (M = Sm–Lu, Y)⁴¹ was synthesized in a quartz ampoule at 400 °C for two weeks. This preparation method greatly limits industrial production.

High-energy ball milling has now become one of the mainstream methods for solid electrolyte synthesis. It refines the raw materials through the impact between mill ball and the original materials, providing the activation energy required for the reaction, and significantly reducing preparation time. Schlem et al.⁴² optimized the traditional ampoule synthesis method and prepared Li₃YCl₆ and Li₃YBr₆ using ball milling technique. There are mainly two types of popular ball milling methods: (1) Direct ball milling method, which does not require any additional processing of the materials. For example, Li_2.99Y_0.99Zr_0.01Br_1.6Cl_4.4⁷⁹ can be obtained through one-step ball milling, and the non-periodic features induced by intense ball milling lead to high ionic conductivity. (2) In a two-step ball milling method, Li₂ZrCl₆ prepared by Chen et al.⁸⁰ was initially ball milled for 12 hours at 600 rpm with a low ratio of ball and material, and in the second step, ball milling was conducted for 48 hours at speeds ranging from 600 to 900 rpm with higher ratios of ball and material. Through comparison, Li₂ZrCl₆ synthesized using the two-step ball milling method exhibited higher ionic conductivity.

Another common approach is to conduct heat treatment following ball milling. Some halides exhibit poor crystallinity after basic mechanical ball milling, resulting in low ionic conductivity. High-temperature annealing is the most commonly used method to improve crystallinity in materials like Li₃InCl₆ and Li₂Sc_2/3Cl₆.^35,81 Annealing generally increases the crystallinity of the product and enhances ionic conductivity. However, research has found that annealed Li₃InCl₆ not only increases ionic conductivity but also leads to an increase in impurities in the sample, attributed to the formation of InOCl under the influence of humidity.⁸² Nevertheless, the proportion of impurities is very small. After annealing, the ion conductivity of Li₂ZrCl₆ and Li₃ErCl₆^25,31 decreased. The reason for this is that the annealing process suppressed the metastable non-periodic features, which resulted in a decrease in ion conductivity as the number of non-periodic features decreased.

3.2 Liquid phase synthesis

A common method for synthesizing halides is solid-phase synthesis, but the products produced through this process are sensitive to moisture. Sun et al.⁵⁶ reported a method for preparing Li₃InCl₆ in an aqueous solution system, as shown in Fig. 7b. Li₃InCl₆·xH₂O is generated by reacting LiCl and InCl₃ in appropriate molar ratios at room temperature in H₂O solution, and pure phase LIC is obtained by vacuum heating at 200 °C. At this point, it is in a highly crystalline state, with an ion conductivity higher than that of LIC produced by solid-phase methods. Additionally, the reversible conversion between LIC and LIC·2H₂O ensures resistance to air and humidity. Liu et al.⁸² tested the effects of high vacuum (HV; 10⁻³ Torr), low vacuum (LV; 10⁻¹ Torr), Ar, and N₂ (1 atm) drying environments on the microstructure and ion conductivity of LIC. The study found that drying under HV conditions was conducive to forming smaller particles. The removal of crystalline water during heating benefited from shorter diffusion lengths and larger surface areas. Additionally, the ion conductivity was improved (2.70 mS cm⁻¹) under reduced oxygen contamination. This study provides insights for synthesizing high-quality LIC and is also beneficial for the synthesis of other water-mediated solid electrolytes.

4 Conduction mechanism of Li⁺ in halide solid electrolytes

The classic diffusion model describes ion transport as a single ion jumping from one lattice site to another through interconnected diffusion channels within the crystal structure framework. This is currently the basis for understanding ion diffusion in solids. The tetrahedral coordination of Li in sulfides (with a framework of covalent bonds) is geometrically similar to the body-centered cubic (bcc) anion sub-lattice, and the diffusion rate of Li⁺ in the bcc structure is superior to other closed frameworks. The oxides require covalent doping to achieve high Li⁺ conduction through cooperative migration. The halides are based on tight anion stacking structures with ionic bonds for Li⁺ migration, where Cl⁻ and Br⁻ can rapidly transport Li⁺ in different anion sub-lattices. Compared to sulfides and oxides, halides have fewer lattice restrictions, and the fundamental difference between ionic and covalent bonds in terms of gaining and losing electrons, results in different migration phenomena and diffusion mechanisms for Li⁺.^34,83–85 The structure of halides is the main factor determining ionic conductivity. This section will provide a detailed discussion on the diffusion mechanism of Li⁺ in four crystal structures and the influencing factors.

4.1 Conduction mechanism of various crystal structures

4.1.1 Monoclinic structure (space group C2/m). Liang et al.³⁴ discussed the effect of the value of x in Li_xScCl_3+x³⁰ on Li⁺ diffusion. When x = 3, there is an optimal Li⁺ diffusion channel and vacancy concentration. Li⁺ jumps between two adjacent octahedral sites through the tetrahedral interstice, forming a three-directionally isotropic channel (Oct–Tet–Oct) (Fig. 9a). When x > 3, although the structure remains unchanged from Li₃ScCl₆, the decrease in Sc occupancy leads to an increase in Li⁺ concentration and a decrease in diffusion vacancies for Li. When x < 3, the Coulomb forces between Sc³⁺ and Li⁺ repel each other, blocking the tetrahedral interstice sites near Sc³⁺, which leads to a decrease in ion jump rate. In Li₃YBr₆,²² the migration paths of Li⁺ in all directions are connected through tetrahedral gaps (Oct1–[Tet1 or Tet2]–Oct2, Oct1–Tet3–Oct3) to form a three-dimensional channel (Fig. 9b). Jia et al.⁸⁶ discovered through AIMD simulation that the highly symmetric ccp structure in Li₃InCl₆ forms an isotropic three-dimensional channel for the motion trajectory of Li⁺. The adjacent Li⁺ ions are restricted to a large octahedral unit due to the longer bond length between In and Cl, while the larger channels along the migration path make it easier for Li⁺ to diffuse. The highly localized position of In in the unit cell stabilizes the entire structure, enabling Li⁺ migration while inhibiting Cl⁻ diffusion.

Fig. 9 (a) The Li⁺ migration pathways in ccp-anion stacking sublattice of Li₃ScCl₆ structure.³⁴ Reprinted with permission from ref. 34. Copyright (2020) American Chemical Society. (b) Li⁺ migration path in Li₃InBr₆.²² Reprinted with permission from ref. 22. Copyright (2018) Wiley. (c) Different transition areas for possible lithium diffusion in Li₃ErCl₆.⁸⁷ Reprinted with permission from ref. 87. Copyright (2019) Wiley. (d) Ionic conduction pathway of Li₃YCl₆. The gray spheres represent Li vacancies.⁸⁸ Reprinted with permission from ref. 88. Copyright (2021) American Chemical Society. (e) and (f) Li-ion migration pathways of Li₂ZrCl₆.³² Reprinted with permission from ref. 32. Copyright (2021) Springer Nature. (g) and (h) 3D Li⁺ ion diffusion pathway of disordered spine Li_2/3ScCl₄.⁹¹ Reprinted with permission from ref. 91. Copyright (2023) Royal Society of Chemistry. (i) Top view of the LaCl₃ lattice along the c axis to show the unit cell (dashed square) and the abundant, intrinsically pre-existing channels of inner diameter 4.6 Å. (j) Side view of the vacancy-contained LaCl₃ lattice indicating Li⁺ migration along the 1D channel (red spheres) and between adjacent channels (bidirectional arrows; vacancies are represented by the grey tricapped trigonal prisms). (k) Li⁺ probability density, represented by green isosurfaces from AIMD simulations at 900 K in the vacancy-contained LaCl₃ lattice. (l) Isolated Li⁺ probability density isosurfaces by removal of all [LaCl₉] polyhedrons to show the excellent interconnectivity of the 3D Li⁺ migration pathways.⁹² Reprinted with permission from ref. 92. Copyright (2023) Springer Nature.

4.1.2 Trigonal structure (space group P3̄m1). Li₃YCl₆, Li₃ErCl₆, and Li₂ZrCl₆ all have a trigonal structure (space group P3̄m1), and are the most extensively studied in this structure. Zeier et al.⁸⁷ proposed four possible migration pathways for the polyhedra of Li⁺ diffusion region in Li₃ErCl₆: (1) Jumping along the z-direction on the shared face of LiCl₆⁵⁻ octahedra to form the O_h(6h)–O_h(6g) pathway. (2) Jumping on the (002) plane, O_h(6h)–T_d–O_h(6h). (3) Jumping on the (001) plane, O_h(6g)–T_d–O_h(6g). (4) Jumping between two adjacent tetrahedra, T_d–T_d. Complete occupation of the 6g site may lead to faster anisotropic migration of pathways 1 and 2 (Fig. 9c). In Li₃YCl₆, two adjacent octahedra are connected along the a and b axes by sharing edges,⁸⁸ and they are connected along the c axis by sharing faces. On the a and b axes, Li⁺ ions move from octahedral (O) site to the adjacent octahedral (O) site via the tetrahedral (T) site, creating two O–T–O migration paths. On the c axis, Li⁺ jumps between adjacent octahedral (O) sites, forming an O–O path (Fig. 9d). The climbing image nudged elastic band (CI-NEB) method calculates the migration energy barriers of three paths as follows: two O–T–O migration paths (0.27–0.52 eV), and O–O path (0.19 eV). The results show that in LYC, Li⁺ is more inclined to migrate along the c-axis. Ertekin et al.⁸⁹ conducted a more in-depth study on Li⁺ conduction in LYC. When the anion framework is frozen, Li⁺ diffusion is almost completely inhibited, confirming that the migration of Li⁺ is closely related to the vibration of the densely packed anion framework. Li et al.⁸⁶ demonstrated through calculations that the migration trajectory of the c-axis is more continuous than that of the a or b axes. The difference in migration energy barriers in three directions can be explained by the weaker bonding between Li and Cl in the LiCl₆ octahedron, making it easier for Li to leave the host. This anisotropic diffusion pathway leads to an overall decrease in the ionic conductivity of LYC.

Compared to other isomorphic chloride systems, Li₂ZrCl₆²⁵ has smaller lattice parameters and a lower occupancy rate of Zr⁴⁺ sites. Calculations using the bond valence site energy (BVSE) method indicate that the Li1–Li2–Li1 path in the [001], direction is the optimal pathway for Li⁺ migration simultaneously with the Li1–i1–Li1–i2–Li1 and Li2–i3–Li2 paths form an anisotropic three-dimensional network (Fig. 9e). However, this crystal structure does not facilitate Li⁺ migration. The high ionic conductivity of Li₂ZrCl₆ at room temperature is attributed to the non-periodic features induced by intense ball milling.

4.1.3 Orthorhombic structure (space group Pnma). After annealing, Li₃YCl₆ will exhibit an orthorhombic structure with a space group of Pnma, which differs mainly from P3̄m1 in the arrangement of octahedra. Additionally, in the Pnma space group, some Li sites are occupied, and there are no vacant sites for Y.⁹⁰ Although LYC has two different structures, the transport pathways of Li⁺ all form anisotropic three-dimensional networks and preferentially diffuse along the c-axis.

4.1.4 Cubic structure (spinel). The common feature of Li₂Sc_2/3Cl₄ and Li₂MCl₄ (M = Mg, Fe, Cr, Mn) is the disordered distribution of Li⁺ on the octahedral sites. Li₂Sc_2/3Cl₄ is the first spinel-type superionic halide, with an ion conductivity three orders of magnitude higher than other spinel phases. In the crystal lattice of Li₂Sc_2/3Cl₄, there are multiple Li sites within the unit cell. The Li⁺ at other sites can migrate through a rigid framework formed by Li/Sc sharing positions, creating a 3D diffusion pathway with the shared octahedral and tetrahedral Li1,2,3 sites (Fig. 9f). The new diffusion pathways and numerous vacancies contribute to the relatively high ionic conductivity of Li₂Sc_2/3Cl₄.^35,37,60 The Li₂Sc_2/3Cl₄ computational model established by Han et al.⁹¹ suggests that the promotion of fast Li⁺ transport is due to the small electrostatic repulsion between cations. In addition to the fully connected three-dimensional network, there are additional diffusion pathways at other sites to facilitate Li⁺ migration. In Li₂MgCl₄,³⁵ Li⁺ migrates through the octahedral surface shared with the tetrahedral edges. The lower ionic conductivity is attributed to fewer Li⁺ vacancies, and the presence of Mg²⁺ in the main diffusion path of Li⁺ hinders its migration.

4.1.5 UCl₃-type crystal framework. In the UCl₃ lattice,³⁹ chloride ions are arranged in a non-close-packed manner, forming abundant one-dimensional large-sized channels. By inducing lanthanum vacancies, it is possible to achieve continuous conductivity in three-dimensional space, endowing lithium ions with the ability to freely migrate throughout the entire crystal network. Yao et al.⁹² prepared Li_0.388Ta_0.238La_0.475Cl₃ based on the characteristics of the UCl₃ lattice, which exhibits a high ionic conductivity of 3.02 mS cm⁻¹ at room temperature. Although the LaCl₃ lattice is rich in octahedral positions, its inherent single-dimensional transport path is insufficient to support high ionic conductivity (Fig. 9i). By doping with Ta to introduce lanthanum vacancies, these vacancies trigger strong exchange interactions between large-sized high-speed ion channels and adjacent channels, thereby constructing initial three-dimensional penetrating ion flow channels (Fig. 9j–l). In addition, this lattice has strong tolerance and can accommodate various ions from +1 to +6 without disrupting its basic UCl₃-type framework structure.

4.1.6 Amorphous halides. The current report indicates that amorphous halides can be specifically categorized into two groups: glassy and clay-like states. Yao and his team's²⁸ research on LiTaCl₆ through molecular dynamics simulations revealed that the migration of Li ions occurs between the octahedral and tetrahedral channels, involving gap transitions between Li–Ta–Cl. This demonstrates a three-dimensional directional transport mechanism similar to Li₂ZrCl₆. However, the ion migration path in LiTaCl₆ includes more intermediates between Li–Ta–Cl, and all energy barriers during the migration process are lower than in Li₂ZrCl₆. The reason why the clay-like 2LiCl–GaF₃⁹³ exhibits high ionic conductivity is due to the presence of mixed anions adjusting the local structure and energy characteristics. In this material, Ga shows enhanced coordination flexibility, mainly attributed to two points: firstly, Cl⁻ creates a relatively spacious space around Ga due to its smaller ionic radius, available for Cl⁻ and F⁻ coordination; secondly, the electronegativity of Cl (3.16) is close to F (3.98), promoting a more uniform electron density distribution in Ga(Cl, F)_n, beneficial for maintaining stability under a wider range of coordination numbers (n = 4, 5, 6). The environment in which Li⁺ is located is defined by the mixed halide anions around Ga, collectively promoting short-range ordered structures in the amorphous framework of 2LiCl–GaF₃.

Observing and analyzing the typical crystal structures mentioned above reveals that compounds with monoclinic structures featuring cubic close-packing, such as Li₃ScCl₆, Li₃YBr₆, and Li₃InCl₆, exhibit similar or identical fast ion migration pathways, leading to superior ion conductivity compared to several other types. It is worth noting that although a single UCl₃-type lattice does not support fast ion conduction due to its one-dimensional channel characteristics, the introduction of external elements to create lanthanide vacancies and construct a three-dimensional ion migration network enables fast ion conduction. This discovery provides insights for the future development of novel electrolyte materials from the perspective of crystal structure design.

4.2 The impact of defects on Li⁺ conduction

Defects provide additional diffusion space and storage sites for ions, which are considered the most effective solution for fundamentally improving ion diffusion kinetics.⁹⁴ In solid halide electrolytes, the primary factor influencing ion conduction is the vacancies created by defects to facilitate their movement. Vacancy is one of the common crystal defects. Ren et al.⁹⁵ constructed vacancy defect structures of Li₃InCl₆ and Li₃ScCl₆ (Fig. 10a and b). The concentration of Li inside the unit cell decreases with the introduction of Li vacancy defects, leading to a significant increase in Li⁺ diffusion capability. These defects provide Li with transport channels while reducing the migration energy barrier. Research shows that the number of Li site defects within an appropriate range can enhance the migration rate of Li⁺ in Li₃InCl₆ and Li₃ScCl₆. Kim et al.⁹⁶ established four models, including normal Li₃InCl₆, In occupying 4g sites (In_4g), Cl vacancies (V_Cl), and the coexistence of In_4g and V_Cl, to study the effect of vacancy defects on the ionic conductivity (Fig. 10c–f). The presence of vacancy defects reduces the migration energy barrier of Li⁺, promotes partial Li penetration to the vacant octahedral sites in the In layer, increases the concentration of Li layer vacancies, accelerates the diffusion rate of Li⁺, and results in higher ionic conductivity.

Fig. 10 (a) and (b) Li₃ScCl₆ and Li₃InCl₆ defect models constructed according to the vacancy formation energy, where the red part is the distribution of lithium site defects.⁹⁵ Reprinted with permission from ref. 95. Copyright (2023) Elsevier. (c) Supercell with no point defects introduced (LIC), (d) with a 4g site occupancy by In (LIC_In_4g), (e) with a Cl vacancy (LIC_V_Cl), (f) with both a 4g site occupancy by In and a Cl vacancy (LIC_In_4g–V_Cl).⁹⁶ Reprinted with permission from ref. 96. Copyright (2022) Royal Society of Chemistry. (g and h) Crystal structure of LYBC-HP viewed along different orientations. (i and j) Proposed Li⁺ diffusion pathways along ab-plane and c-direction.⁹⁷ Reprinted with permission from ref. 97. Copyright (2021) American Chemical Society.

Stacking faults are a common type of crystal defect. Liu et al.⁹⁷ obtained tetrahedral (T–Li) and octahedral (O–Li) coexisting Li₃Y(Br₃Cl₃) by mixing Br and Cl ions (Fig. 10g–j). Axial stacking defects result in local order and long-range disorder within the crystal. T–Li creates more vacancies on the octahedra, leading to a cooperative diffusion mechanism, which is one of the reasons for the high ionic conductivity of Li₃Y(Br₃Cl₃). Sebti et al.⁹⁸ found that the mechanosynthesis of Li₃YCl₆ (LYC) exhibited a phenomenon of a high concentration of stacking faults and the coexistence of lithium layer defects. The presence of stacking faults and lithium layer defects establishes additional interconnected sites, creates more interlayer channels, and reduces the transport barrier of Li⁺. However, when exposed to environments above 60 °C, the defect concentration and Li⁺ migration rate decreased, indicating a strong correlation between defect concentration and the ionic conductivity of LYC. This also demonstrates the important role of defects in promoting the structural conduction of Li⁺.

The arrangement of cations is a crucial factor that determines the high ionic conductivity of halide superionic conductors. Zhong et al.⁹⁹ simulated the effect of Zr disorder/order on Li₂ZrCl₆ ion migration. In their study of the α-hcp and β-fcc phases of LZC, they found that compared to the ground state order, when Zr is completely disordered, Li⁺ diffusion exhibits Arrhenius behavior at all temperatures. The Li/vacancy disorder caused by Zr disorder promotes the rapid conduction of Li⁺, but spontaneous disorder of Zr does not occur at any reasonable temperature, only non-equilibrium high-energy processes (such as mechanical ball milling) can achieve a disordered state. This also explains the decrease in ionic conductivity of LZC after high-temperature annealing.

In summary, defects within the crystal structure when designing halide solid electrolytes are an important research focus that cannot be ignored. In addition, multiple factors such as grain boundary characteristics, impurity content, and mechanical-chemical interactions during the synthesis process significantly influence the transport behavior of Li⁺ at various levels.

4.3 Modified strategy

4.3.1 Different/same-valence cation doping. Doping with cations of different or same valence is a commonly used method for modifying electrolytes, mainly by enhancing ion conductivity by inducing lattice distortion, increasing vacancies, and altering the crystal structure. The ionic conductivity before and after modification is shown in Table 3.

Table 3 Comparison of ionic conductivity before and after modification

Modified strategy	Original sample	mS cm^−1	Modified	mS cm^−1	Ref.
Cation doping	Li2MgCl4	0.00019	Li1.6Zn0.2MgCl4	0.012	101
	Li3ScCl6	0.8	Li2.375Sc0.375Zr0.625Cl6	2.2	102
	Li2ZrCl6	0.4	Li2.25Zr0.75Fe0.25Cl6	1	33
	Li3ErCl6	0.33	Li2.633Er0.633Zr0.367Cl6	1.1	104
	Li3YCl6	0.51	Li2.5Y0.5Zr0.5Cl6	1.4	104
	Li2ZrCl6	0.32	Li2.25Zr0.75In0.25Cl6	1.08	105
	Li2ZrCl6	0.5	Li2.8Zr0.2In0.8Cl6	1.4	107
Fluoridation	Li3YCl6	0.224	Li3YCl5.97F0.03	0.239	109
	Li3InCl6	0.55	Li3InCl5.8F0.2	1.2	110
High-entropy	Li3InCl6	0.849	Li2.75Y0.16Er0.16Yb0.16In0.25Zr0.25Cl6	1.171	111
	Li3InCl6	1.27	Li2.8In0.2Sc0.2Yb0.2Lu0.2Zr0.2Cl6	2.13	112

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Chen et al.¹⁰⁰ reported Zn-doped Li_1.6Zn_0.2MgCl₄, which showed an increase in ion conductivity by two orders of magnitude compared to Li₂MgCl₄ (1.2 × 10⁻⁵ S cm⁻¹). The Li⁺ is slightly larger than Zn²⁺, and when Zn enters the Li₂MgCl₄ lattice, the unit cell volume decreases by 0.89%. However, due to charge balance, the substitution of Zn²⁺ for Li⁺ creates Li vacancies (Fig. 11a and b). Element substitution not only affects the vacancy concentration but also smoothens the overall migration energy barrier. Vacancies within a certain range contribute to enhancing Li⁺ conductivity, but when there are too many vacancies, the ion conductivity decreases instead. Wu et al.¹⁰¹ introduced Zr/Hf into Li₃ScCl₆ using a strategy of adjusting the vacancy concentration to prepare Li_2.375Sc_0.375Zr_0.625Cl₆ and Li_2.375Sc_0.375Hf_0.375Cl₆. After successfully substituting Zr⁴⁺/Hf⁴⁺, due to lattice contraction in the c-direction, the jump distance of Li⁺ along the axis becomes shorter. Simultaneously, an empty Li site is instantly occupied by another Li ion, resulting in excellent ion conductivity due to the rapid ion exchange rate and larger diffusion bottleneck. Usami and colleagues¹⁰² found that doping Li₃InCl₆ with Nb and Zr resulted in an increase in the volume of the unit cell. The relationship between the sizes of the three ions is as follows: In³⁺ > Zr⁴⁺ > Nb⁵⁺, and the unit cell volume replaced by Nb⁵⁺ is larger than that of Zr⁴⁺. It is not difficult to see that the increase in lattice constant is not directly caused by the substitution of cations. When Nb or Zr enters the unit cell, vacancies are created, causing electrostatic repulsion between internal anions, leading to an increase in unit cell volume as the vacancy space expands. In addition to the impact of vacancies on Li⁺ transport, first-principles molecular dynamics (FPMD) calculations show that at 373 K–673 K, Li⁺ transitions from existing only in In-deficient layers to three-dimensional diffusion within and between layers. At 673 K and above, a phase transition occurs where Li/vacancy arrangements become disordered (Fig. 11c and d), and the increase in disorder also facilitates the migration of Li⁺ to some extent. Replacing Zr with Fe in Li₂ZrCl₆³² not only reduces the Coulomb repulsion between Li⁺ and other metal cations, lowering the migration energy barrier of Li⁺ but also, after Fe enters the lattice, the unit cell volume decreases, and the channels between the octahedra being connected along the c-axis direction, making it easier for Li⁺ to diffuse.

Fig. 11 (a) Crystal structure of Li_2-2xZn_xMgCl₄. (b) The diagram of Li vacancy induced by Zn substitution.¹⁰⁰ Reprinted with permission from ref. 100. Copyright (2023) Elsevier. (c) Population density of Li ions in LIC (Li₄₈In₁₆Cl₉₆) derived from FPMD calculations at 373 and 673 K, respectively. (d) Diffusion pathway of Li ions at 673 K.¹⁰² Reprinted with permission from ref. 102. Copyright (2023) AIP Publishing.

The trigonal structured Li₃YCl₆ is considered a poor ion conductor. Park et al.¹⁰³ reported Zr-doped Li₃MCl₆ (M = Y, Er), where phase transitions occur at higher doping levels. The rearrangement of metal ions and tilting of octahedra transform the trigonal phase into an orthorhombic phase. The additional tetrahedral sites created play a crucial role in altering the overall migration energy barrier. A substantial number of vacancies create three-dimensional channels for Li⁺ migration, resulting in a significant increase in ionic conductivity at room temperature by an order of magnitude. Based on the theory of hard and soft acids and bases, compounds formed by soft acid In are less likely to react with hard base H₂O. Xie et al.¹⁰⁴ prepared Li_2.25Zr_0.75In_0.25Cl₆, where the covalent substitution strategy effectively increased the ionic conductivity. Excessive substitution and annealing, however, can cause phase transitions and decrease the ionic conductivity. Transmission electron microscopy revealed that the samples after ball milling exhibited a disordered structure, which could be a contributing factor to the high ionic conductivity. Ma and Tu^105,106 reported a mixed-phase Li_2.8Zr_0.2In_0.8Cl₆ containing P3̄m1 and C2/m space groups after ball milling, which completely transformed into a monoclinic structure after annealing. When the In substitution amount is between 0.3–0.8, In doping results in lattice expansion in three directions, and the migration barrier for Li⁺ along all migration paths gradually decreases. When the substitution amount exceeds 0.8, the unit cell abnormally shrinks, and the reason for the decrease in ionic conductivity, in addition to the increase in migration energy barrier, may also be attributed to the insufficient Li vacancies available to facilitate Li⁺ migration.

Currently, in the methods of doping halides with cations of different or same valence, a very inspiring idea is to alter the crystal structure, transitioning from trigonal to orthorhombic, and from orthorhombic to a monoclinic phase with isotropic ion conduction. The monoclinic phase, through lattice expansion induced by doping, can enhance ionic conductivity, offering a pathway for future modifications to the structure of HSEs.

4.3.2 Fluoridation strategy. Fluorination is the most commonly used strategy to enhance the stability of electrolytes. In addition, the introduction of fluoride ions into the lattice can lead to local distortion, which can impact the ionic conductivity. Kim et al.¹⁰⁷ found additional vacancies in fluorinated Li₃InCl₆. The migration barrier near F is lower than at other positions, but the Coulomb force of F on Li suppresses the movement range of surrounding Li, affecting Li⁺ diffusion. Therefore, controlling the concentration of F reasonably can achieve the best doping effect. When Li moves to an empty site in the octahedral layer of In, the initial site forms a new vacancy, which promotes the migration of Li⁺. This is not only the main migration channel for Li⁺, but also the key to the entire diffusion process. Zhong et al.¹⁰⁸ reported Li₃YCl_5.97F_0.03, where trace amounts of F do not affect the ionic conductivity, but the LiF formed at the interface of electrolyte and negative electrode can not only enhance the oxidation stability of the electrolyte but also inhibit its decomposition. Wang et al.¹⁰⁹ prepared Li₃InCl_5.8F_0.2, and found, through the Arrhenius equation, that the introduction of F can significantly reduce the activation energy, leading to a notable increase in ionic conductivity. Additionally, the modified Li₃InCl₆ exhibited a significant decrease in electronic conductivity, effectively suppressing the growth of lithium dendrites and extending the cycle life of the battery.

4.3.3 High-entropy modification. High-entropy strategy is an effective approach to promote Li/vacancy disorder and increase configurational entropy. Huang et al.¹¹⁰ reported a new high-entropy halide Li_2.75Y_0.16E_r0.16Yb_0.16In_0.25Zr_0.25Cl₁₆ (HE-LIC). The original Li₃InCl₆ is a single phase with a monoclinic structure. However, when Y and Zr are introduced, the structure becomes a mixed phase. With the further introduction of Er and Yb, the structure transforms into a monoclinic phase with a small amount of rock salt LiCl. This also indicates that the high-entropy-driven structural stability can promote the solid solution of various cations in LIC. The lattice distortion caused by introducing different cations serves a dual function: on one hand, it can effectively regulate the diffusion kinetics of Li⁺ and Cl⁻, on the other hand, this distortion process creates unobstructed three-dimensional migration channels for Li⁺, thereby achieving high ionic conductivity. Similarly, Zhao et al.¹¹¹ improved the configurational entropy by increasing the number of main elements, thereby altering the thermodynamic stability and kinetic properties of the material to achieve fast ion conduction in the preparation of Li_2.8In_0.2Sc_0.2Yb_0.2Lu_0.2Zr_0.2Cl₆. This material exhibits an ionic conductivity of up to 2.13 mS cm⁻¹.

5 Oxyhalides

The topological structure of a specific anion arrangement determines the inherent mobility of Li⁺. However, closely packed anions in halides are not conducive to Li⁺ conduction, which seems to be related to monovalent halide anions. Introducing a mixed double anion strategy in halide matrices can enhance ion conductivity, as demonstrated in Li₃Y(Br₃Cl₃).^26,85,97 Emly and others studied oxyhalides with an anti-perovskite structure. Unlike other electrolytes, Li₃OCl and Li₃OBr exhibit a high number of vacancies for Li⁺ transport, offering a new avenue for modifying halides.^112–115 The emergence of new lithium-ion conductors is largely the result of expanding known superionic compounds into new compositional spaces.⁸⁵ A mixed anion system oxyhalides based on halides has now been reported.

5.1 Crystalline oxyhalides

Tanaka et al.²⁶ prepared LiMOCl₄ (M = Nb, Ta) with a conductivity greater than 10 mS cm⁻¹ at room temperature. The use of high-valence cations to form corner-sharing polyhedra is expected to create pathways for Li⁺ migration. LiNbOCl₄ (LNOC) has an orthorhombic structure (space group Cmc21), consisting of corner-sharing NbO₂Cl₄ octahedral chains and LiCl tetrahedra located between the chains (Fig. 12a–c). Compared to Li₃YCl₆ and Li₃YBr₆, the crystal lattice framework of LNOC is more open, with larger diffusion channels. This is also a significant factor contributing to the high ionic conductivity of LNOC. Further understanding of the diffusion mechanism of LNOC reveals that the moving ions diffuse independently without correlation, following a typical ion hopping mechanism similar to diffusion in liquids. Additionally, its higher oxidation stability enables long cycling and high capacity in batteries, offering new perspectives for solid-state electrolytes.

Fig. 12 (a) Crystal structure of the LNOC. The Li-ion probability density map (yellow isosurface) along the (b) b-axis and (c) a-axis direction calculated by AIMD simulations at 800 K.²⁶ Reprinted with permission from ref. 26. Copyright (2023) Wiley. (d and e) Atomic model for the Li₃Zr_0.75OCl₄ materials.¹¹⁶ (f) Reverse Monte Carlo (RMC) fit (red line) to the experimental G(r) (gray circle) of Li₃ZrCl₄O_1.5. Several possible basic building blocks in Li₃ZrCl₄O_1.5 are shown on the left side. On the right side, the schemes show the connectivity that leads to the edge-sharing and corner-sharing Zr-centered polyhedra in Li₃ZrCl₄O_1.5, respectively.¹²⁰ Reprinted with permission from ref. 116 and 120. Copyright (2023 and 2024) American Chemical Society.

Ma et al.¹¹⁶ reported Li₃Zr_0.75OCl₄ with a monoclinic structure. Currently, besides introducing expensive metals such as In and Er in the modification of Li₂ZrCl₆, it is challenging to achieve an ionic conductivity exceeding 1 mS cm⁻¹ at room temperature. Starting with the monoclinic structure Li₃ScCl₆, which possesses rapid ion transport characteristics, the product obtained by replacing Sc with 0.75Zr is not a pure phase. Considering that vacancies are needed between transition metals to balance the charge difference between Sc³⁺ and Zr⁴⁺, O is used to replace Cl to introduce anionic vacancies. This results in Li₃Zr_0.75OCl₆ with the same structure as Li₃ScCl₆ (Fig. 12d and e). In addition, unlike Sc, the Zr sites are not fully occupied, containing a large number of vacancies, which may provide an additional pathway for Li⁺ migration. BVSE analysis shows that the C2/m structure of Li₃Zr_0.75OCl₆ is more conducive to lithium ion migration compared to the P3̄m1 structure of Li₂ZrCl₆. The migration energy barrier for the three-dimensional network pathway is only 0.306 eV, and ion conduction can occur in non-periodic features and crystal phases, both of which collectively determine the overall ionic conductivity.

Yu et al.¹¹⁷ enhanced the ionic conductivity of Li₂ZrCl₆ by substituting hard alkalis. The ionic conductivity of the monoclinic (metastable phase) Li_3.1ZrCl_4.9O_0.1 reached 1.3 mS cm⁻¹. In addition to the fast ion conduction of the monoclinic structure, the energy barrier on the Li⁺ migration path is lower after introducing O²⁻. Oxygen substitution stabilizes the Li interstitial sites, but this method is only effective for structures where tetrahedra and metal octahedra share corners. Li et al.¹¹⁸ reported a monoclinic Li_2.22Zr_1.11Cl_5.33O_0.67 with a room temperature ionic conductivity exceeding 1 mS cm⁻¹. The lower activation energy is closely related to the changes in crystal structure and the occupation of lithium in different interstitial positions. When Li occupies more tetrahedral sites, the required jump distance for its migration is correspondingly shortened, which is crucial for enhancing high ionic conductivity. Due to the stronger Li⁺–O²⁻ bonds compared to Li⁺–Cl⁻, Li_2.22Zr_1.11Cl_5.33O_0.67 exhibits higher structural thermal stability and solvent stability.

5.2 Amorphous oxyhalides

Sun et al.¹¹⁹ constructed an amorphous xLi₂O-MCly (M = Nb, Ta) with a mixed anion structure. After optimization, the ion conductivity can reach 6.6 mS cm⁻¹. The disordered and irregular [TaCl₅–O] structure enables the formation of Li–Cl sub-lattices in the amorphous region, and the oxygen at bridging sites expands the migration channel radius, facilitating the entry of Li⁺. Furthermore, the unsaturated Ta–Cl⋯Li bonds between Li–Cl have weaker coulombic forces, enabling Li⁺ to transition easily between sites. A similar situation is also observed in Hf. In conclusion, the combination of amorphous terminal Cl⁻ with abundant O²⁻ plays a crucial role in the rapid migration of Li⁺. In addition, the group also controlled the amorphization degree of Li₂ZrCl₆ by introducing O,¹²⁰ conducted an in-depth analysis of the Li₃ZrCl₄O_1.5 structure, where the polyhedra centered on Zr are mainly connected by corner-sharing Cl and O, with Zr–X (X = O, Cl) bond lengths of disordered and distorted Zr–O/Cl polyhedra being averaged. The remaining polyhedra exhibit edge-sharing and a small amount of face-sharing (Fig. 12f). This unique structure enables its ionic conductivity to reach (1.35 ± 0.07) mS cm⁻¹.

Li_1.75ZrCl_4.75O_0.5 reported by Ma et al.¹²¹ exhibits an ionic conductivity of 2.42 mS cm⁻¹ at room temperature, with the structure mainly being amorphous (where the crystalline phase is less than 20%: trigonal and monoclinic). It is important to emphasize that effective ion transport is achieved through a large amount of amorphous material when the crystal structure transitions to an amorphous state by introducing O into Li₂ZrCl₆. Hu et al.¹²² reported a viscoelastic amorphous LiAlCl₄-75%O (LACO75) electrolyte, in which the ionic conductivity reached 1 mS cm⁻¹ at 30 °C by replacing chlorine in tetrachloroaluminate with oxygen. This change resulted in an improvement of 3 orders of magnitude compared to LiAlCl₄. As the oxygen content increased, the oxygen-containing tetrachloroaluminate transitioned a brittle molten salt a ductile glass. AIMD simulations revealed that the LACO75 structure is completely disordered, with oxygen connecting isolated aluminum atoms to form an Al–O–Al network. In molten the strong Al–O bonds are challenging to break, and during the contraction process of atomic rearrangement, they are further constrained by the Al–O–Al network, leading to the formation of a glassy state. The high ionic conductivity is attributed to the movement/rotation of segments, shorter Li⁺ hopping paths, and the presence of Al–O–Al chains.

6 Summary and outlook

This article reviews the synthesis methods, structure, corresponding ion conduction mechanisms, modification strategies of halide solid electrolytes, and finally elaborates on the breakthrough from halides to oxyhalides. Firstly, existing halide electrolytes are classified into five categories based on different types of metal elements, and detailed explanations are provided for the crystal structures and space groups of each category. Secondly, the preparation processes of common synthesis methods such as high-energy ball milling, annealing, and liquid-phase method are summarized. Thirdly, starting from four typical structures, the ion conduction mechanisms are organized and summarized, followed by a detailed explanation of three modification methods. Finally, the article describes the connection between halides and oxyhalides and introduces the structures of existing oxyhalides. At present, research on halides primarily focuses on the following aspects:

(1) Ionic conductivity. In order to effectively reduce the migration energy barrier of Li⁺, researchers use various modification methods: (1) doping with cations of different or same valence to increase Li vacancies. (2) substituting larger radius central metal elements for smaller ones to expand the lattice and enlarge Li⁺ migration channels. (3) adjusting the crystal structure. For instance, Li₃InCl₆ with a monoclinic structure demonstrates excellent ionic conductivity at room temperature, whereas economically advantageous Li₂ZrCl₆ does not exhibit remarkable ionic conductivity at room temperature. By substituting In for Zr to induce a transition from a trigonal to a monoclinic structure to enhance ionic conductivity.

(2) Synthesis process. High-energy ball milling is currently the most widely used synthesis method. The non-periodic characteristics resulting from intense ball milling lead to high ionic conductivity. However, the lengthy processing time and the substantial disparity between theoretical simulations and actual results have yet to be addressed. Recently, the proposed two-step ball milling method has somewhat enhanced the ionic conductivity of Li₂ZrCl₆, but the 60-hour synthesis process is discouraging. Research has shown that Li₃InCl₆ prepared by liquid-phase synthesis exhibits excellent electrochemical performance. However, it remains to be verified whether this synthesis method is applicable to other halide electrolytes.

(3) Oxygen anion intervention. Recently, research on the introduction of O²⁻ into halides has been particularly focused on Li₂ZrCl₆. This modification strategy involves adjusting the structure and chemical properties of materials by introducing oxygen anions, aiming to enhance their ionic conductivity and electrochemical stability, which is particularly crucial for the application of solid-state batteries.

At the current stage, halides exhibit relatively balanced performance compared to sulfides and oxides, indicating promising prospects. However, most central metal atoms are relatively expensive, and there is still a significant gap in ion conductivity of halides and advanced sulfide solid electrolytes at room temperature. Here are some ideas: (1) Develop new low-cost halides. Current research focuses on In, Zr, and Y. Although Zr has low cost, its room temperature ion conductivity is not ideal. In and Y have higher costs, which seriously affects commercialization. Therefore, the exploration of new halide systems is still urgent. (2) Exploration of spinel-type materials. Research on spinel-type electrolytes is still limited, but Li₂Sc_2/3Cl₄ reported by Xueliang Sun and others in 2022, exhibits excellent ion conductivity. We estimate that there is still significant room for improvement in spinel-type halides. (3) Oxygen anion involvement in other systems. Apart from Zr, Nb, Hf, and Ta, there have been no reports on other oxyhalides, but there are numerous existing halide systems. Combining research on the monoclinic structure or coexistence with other phases of Zr-based oxyhalides, can oxygen anions be directly introduced into halides with monoclinic structures to improve electrochemical performance? There has been little exploration of the amorphous state in oxyhalides. It is unknown whether crystalline or amorphous forms will be more advantageous. Further research is needed to understand the unknown properties and issues.

From the research on halide solid electrolytes in 2018 to the excellent performance of oxyhalides shown in 2023, it signifies that this field is in a rapid development stage, indicating that there will be more innovations and breakthroughs in the coming years. The potential applications of halides and oxyhalides in battery technology are particularly noteworthy, as they are expected to address the safety issues and performance limitations brought about by existing liquid electrolytes, and promote the development of battery technology towards higher energy density, longer cycle life, and greater safety.

Data availability

No primary research results, software or code have been included and no new data were generated or analysed as part of this review.

Conflicts of interest

There are no conflicts to declare.

References

1. A. Manthiram, X. Yu and S. Wang, Lithium battery chemistries enabled by solid-state electrolytes, Nat. Rev. Mater., 2017, 2, 16103.
2. K. Zaghib, M. Dontigny, A. Guerfi, P. Charest, I. Rodrigues, A. Mauger and C. M. Julien, Safe and fast-charging Li-ion battery with long shelf life for power applications, J. Power Sources, 2011, 196, 3949–3954.
3. S. Chen, C. Yu, Q. Luo, C. Wei, L. Li, G. Li, S. Cheng and J. Xie, Research Progress of Lithium Metal Halide Solid Electrolytes, Acta Phys.-Chim. Sin., 2022, 39, 2210032.
4. N. Nitta, F. Wu, J. T. Lee and G. Yushin, Li-ion battery materials: present and future, Mater. Today, 2015, 18, 252–264.
5. G. Jia, Z. Deng, D. Ni, Z. Ji, D. Chen, X. Zhang, T. Wang, S. Li and Y. Zhao, Temperature-dependent compatibility study on halide solid-state electrolytes in solid-state batteries, Front. Chem., 2022, 10, 952875.
6. Z.-Y. Gu, X.-T. Wang, Y.-L. Heng, K.-Y. Zhang, H.-J. Liang, J.-L. Yang, E. H. Ang, P.-F. Wang, Y. You, F. Du and X.-L. Wu, Prospects and perspectives on advanced materials for sodium-ion batteries, Sci. Bull., 2023, 68, 2302–2306.
7. J.-Z. Guo, Z.-Y. Gu, M. Du, X.-X. Zhao, X.-T. Wang and X.-L. Wu, Emerging characterization techniques for delving polyanion-type cathode materials of sodium-ion batteries, Mater. Today, 2023, 66, 221–244.
8. S. Ling, B. Deng, R. Zhao, H. Lin, L. Kong, R. Zhang, Z. Lu, J. Bian and Y. Zhao, Revisiting the Role of Hydrogen in Lithium–Rich Antiperovskite Solid Electrolytes: New Insight in Lithium Ion and Hydrogen Dynamics, Adv. Energy Mater., 2022, 13, 2202847.
9. Z. Zeng, X. Shi, M. Sun, H. Zhang, W. Luo, Y. Huang, B. Huang, Y. Du and C.-H. Yan, Stable all-solid-state Li-Te battery with Li3TbBr6 superionic conductor, Nano Res., 2023, 16, 9344–9351.
10. P. Jiang, G. Du, J. Cao, X. Zhang, C. Zou, Y. Liu and X. Lu, Solid–State Li Ion Batteries with Oxide Solid Electrolytes: Progress and Perspective, Energy Technol., 2023, 11, 2201288.
11. M. Shoji, E. J. Cheng, T. Kimura and K. Kanamura, Recent progress for all solid state battery using sulfide and oxide solid electrolytes, J. Phys. D: Appl. Phys., 2019, 52, 103001.
12. R. Rajagopal, Y. Subramanian, Y. J. Jung, S. Kang and K.-S. Ryu, Preparation of Metal-Oxide-Doped Li7P2S8Br0.25I0.75 Solid Electrolytes for All-Solid-State Lithium Batteries, ACS Appl. Mater. Interfaces, 2023, 15, 21016–21026.
13. H. Muramatsu, A. Hayashi, T. Ohtomo, S. Hama and M. Tatsumisago, Structural change of Li2S–P2S5 sulfide solid electrolytes in the atmosphere, Solid State Ionics, 2011, 182, 116–119.
14. J. Wu, S. Liu, F. Han, X. Yao and C. Wang, Lithium/Sulfide All–Solid–State Batteries using Sulfide Electrolytes, Adv. Mater., 2021, 33, 2000751.
15. J. Wu, L. Shen, Z. Zhang, G. Liu, Z. Wang, D. Zhou, H. Wan, X. Xu and X. Yao, All-Solid-State Lithium Batteries with Sulfide Electrolytes and Oxide Cathodes, Electrochem. Energy Rev., 2021, 4, 101–135.
16. C. Tang, Z. Xue, S. Weng, W. Wang, H. Shen, Y. Xiang, L. Liu and X. Peng, A Biodegradable Polyester-Based Polymer Electrolyte for Solid-State Lithium Batteries, Nanomaterials, 2023, 13, 3027.
17. K. Xie and L. Shi, Ultrathin PEO based electrolyte for high voltage lithium metal batteries enabled by polymer host-plasticizer interactions, J. Energy Storage, 2023, 68, 107640.
18. Z. Yao, K. Zhu, X. Li, J. Zhang, J. Chen, J. Wang, K. Yan and J. Liu, 3D poly(vinylidene fluoride–hexafluoropropylen) nanofiber-reinforced PEO-based composite polymer electrolyte for high-voltage lithium metal batteries, Electrochim. Acta, 2022, 404, 139769.
19. X. Li, J. Liang, X. Yang, K. R. Adair, C. Wang, F. Zhao and X. Sun, Progress and perspectives on halide lithium conductors for all-solid-state lithium batteries, Energy Environ. Sci., 2020, 13, 1429–1461.
20. H.-X. Mei, P. Piccardo, A. Cingolani and R. Spotorno, Unconventional solid-state electrolytes for lithium-based batteries: Recent advances and challenges, J. Power Sources, 2023, 553, 232257.
21. J. Kendall, E. D. Crittenden and H. K. Miller, A study of the factors influencing compound formation and solubility in fused salt mixtures, J. Am. Chem. Soc., 1923, 45, 963–996.
22. T. Asano, A. Sakai, S. Ouchi, M. Sakaida, A. Miyazaki and S. Hasegawa, Solid Halide Electrolytes with High Lithium-Ion Conductivity for Application in 4 V Class Bulk-Type All-Solid-State Batteries, Adv. Mater., 2018, 30, 1803075.
23. B. He, F. Zhang, Y. Xin, C. Xu, X. Hu, X. Wu, Y. Yang and H. Tian, Halogen chemistry of solid electrolytes in all-solid-state batteries, Nat. Rev. Chem., 2023, 7, 826–842.
24. C. Wang, J. Liang, J. T. Kim and X. Sun, Prospects of halide-based all-solid-state batteries: From material design to practical application, Sci. Adv., 2022, 8, eadc9516.
25. K. Wang, Q. Ren, Z. Gu, C. Duan, J. Wang, F. Zhu, Y. Fu, J. Hao, J. Zhu, L. He, C.-W. Wang, Y. Lu, J. Ma and C. Ma, A cost-effective and humidity-tolerant chloride solid electrolyte for lithium batteries, Nat. Commun., 2021, 12, 4410.
26. Y. Tanaka, K. Ueno, K. Mizuno, K. Takeuchi, T. Asano and A. Sakai, New Oxyhalide Solid Electrolytes with High Lithium Ionic Conductivity >10 mS cm−1for All–Solid–State Batteries, Angew. Chem., Int. Ed., 2023, 62, e202217581.
27. K. Tuo, C. Sun and S. Liu, Recent Progress in and Perspectives on Emerging Halide Superionic Conductors for All-Solid-State Batteries, Electrochem. Energy Rev., 2023, 6, 17.
28. F. Li, X. Cheng, G. Lu, Y.-C. Yin, Y.-C. Wu, R. Pan, J.-D. Luo, F. Huang, L.-Z. Feng, L.-L. Lu, T. Ma, L. Zheng, S. Jiao, R. Cao, Z.-P. Liu, H. Zhou, X. Tao, C. Shang and H.-B. Yao, Amorphous Chloride Solid Electrolytes with High Li-Ion Conductivity for Stable Cycling of All-Solid-State High-Nickel Cathodes, J. Am. Chem. Soc., 2023, 145, 27774–27787.
29. M. Lei, B. Li, H. Liu and D. E. Jiang, Dynamic Monkey Bar Mechanism of Superionic Li–ion Transport in LiTaCl6, Angew. Chem., Int. Ed., 2023, 63, e202315628.
30. A. Bohnsack, F. Stenzel, A. Zajonc, G. Balzer, M. S. Wickleder and G. Meyer, Ternare chloride der Selten-Erd-Elemente mit lithium, Li₃MCl₆ (M = Tb-Lu, Y, Sc): Synthese, Kristallstrukturen and Ionenbewegung, Z. Anorg. Allg. Chem., 1997, 623, 1067–1073.
31. R. Schlem, T. Bernges, C. Li, M. A. Kraft, N. Minafra and W. G. Zeier, Lattice Dynamical Approach for Finding the Lithium Superionic Conductor Li3ErI6, ACS Appl. Energy Mater., 2020, 3, 3684–3691.
32. H. Kwak, D. Han, J. Lyoo, J. Park, S. H. Jung, Y. Han, G. Kwon, H. Kim, S. T. Hong, K. W. Nam and Y. S. Jung, New Cost-Effective Halide Solid Electrolytes for All-Solid-State Batteries: Mechanochemically Prepared Fe3+-Substituted Li2ZrCl6, Adv. Energy Mater., 2021, 11, 2003190.
33. G. N. Papatheodorou, Raman spectroscopic studies of yttrium(III) chloride–alkali metal chloride melts and of Cs2NaYCl6 and YCl3 solid compounds, J. Chem. Phys., 1977, 66, 2893–2900.
34. J. Liang, X. Li, S. Wang, K. R. Adair, W. Li, Y. Zhao, C. Wang, Y. Hu, L. Zhang, S. Zhao, S. Lu, H. Huang, R. Li, Y. Mo and X. Sun, Site-Occupation-Tuned Superionic LixScCl3 + xHalide Solid Electrolytes for All-Solid-State Batteries, J. Am. Chem. Soc., 2020, 142, 7012–7022.
35. L. Zhou, C. Y. Kwok, A. Shyamsunder, Q. Zhang, X. Wu and L. F. Nazar, A new halospinel superionic conductor for high-voltage all solid state lithium batteries, Energy Environ. Sci., 2020, 13, 2056–2063.
36. C. Cros, L. Hanebali, L. Latié, W. Villeneuve and W. Gang, Structure, ionic motion and conductivity in some solid-solutions of the LiCI-MCI₂ systems (M = Mg, V, Mn), Solid State Ionics, 1983, 9, 139–148.
37. M. Partik, M. Schneider and H. D. Lutz, Kristallstrukturen von MgCr2O4−Typ Li2VCl4 und Spinell–Typ Li2MgCl4 und Li2CdCl4, Z. Anorg. Allg. Chem., 1994, 620, 791–795.
38. X. Li, Y. Xu, C. Zhao, D. Wu, L. Wang, M. Zheng, X. Han, S. Zhang, J. Yue, B. Xiao, W. Xiao, L. Wang, T. Mei, M. Gu, J. Liang and X. Sun, The Universal Super Cation–Conductivity in Multiple–cation Mixed Chloride Solid–State Electrolytes, Angew. Chem., Int. Ed., 2023, 62, e202306433.
39. C. X. Li Feng, L. U. O. Jinda and Y. A. O. Hongbin, Metal chloride solid-state electrolytes and all-solid-state batteries: State-of-the-art developments and perspectives, Energy Storage Sci. Technol., 2024, 13, 193–211.
40. Z. Xu, X. Chen, K. Liu, R. Chen, X. Zeng and H. Zhu, Influence of Anion Charge on Li Ion Diffusion in a New Solid-State Electrolyte, Li3LaI6, Chem. Mater., 2019, 31, 7425–7433.
41. A. Bohnsack, G. Balzer, H. U. Güdel, M. S. Wickleder and G. Meyer, Ternäre Halogenide vom Typ A3MX6. VII [1]. Die Bromide Li3MBr6 (M=Sm, Lu, Y): Synthese, Kristallstruktur, Ionenbeweglichkeit, Z. Anorg. Allg. Chem., 1997, 623, 1352–1356.
42. R. Schlem, A. Banik, S. Ohno, E. Suard and W. G. Zeier, Insights into the Lithium substructure of the superionic conductors Li₃YCl₆ and Li₃YBr₆, Chem. Mater., 2021, 33, 327–337.
43. K. Wang, Z. Gu, Z. Xi, L. Hu and C. Ma, Li3TiCl6 as ionic conductive and compressible positive electrode active material for all-solid-state lithium-based batteries, Nat. Commun., 2023, 14, 1396.
44. F. Li, X. Cheng, L.-L. Lu, Y.-C. Yin, J.-D. Luo, G. Lu, Y.-F. Meng, H. Mo, T. Tian, J.-T. Yang, W. Wen, Z.-P. Liu, G. Zhang, C. Shang and H.-B. Yao, Stable All-Solid-State Lithium Metal Batteries Enabled by Machine Learning Simulation Designed Halide Electrolytes, Nano Lett., 2022, 22, 2461–2469.
45. R. A. H. W. Weppner, Ionic conductivity of alkali metal chloroaluminates, Phys. Lett. A, 1976, 58, 245–248.
46. W. Weppner and R. A. Huggins, Ionic conductivity of solid and liquid LiAlCl₄, J. Electrochem. Soc., 1977, 124, 35–38.
47. N. Flores-González, N. Minafra, G. Dewald, H. Reardon, R. I. Smith, S. Adams, W. G. Zeier and D. H. Gregory, Mechanochemical Synthesis and Structure of Lithium Tetrahaloaluminates, LiAlX4 (X = Cl, Br, I): A Family of Li-Ion Conducting Ternary Halides, ACS Mater. Lett., 2021, 3, 652–657.
48. Y. Tomita, A. Fuji-i, H. Ohki, K. Yamada and T. Okuda, New Lithium Ion Conductor Li₃InBr₆ Studied by ⁷Li NMR, Chem. Lett., 1998, 27, 223–224.
49. K. Yamada, K. Kumano and T. Okuda, Lithium superionic conductors Li3InBr6 and LiInBr4 studied by 7Li, 115In NMR, Solid State Ionics, 2006, 177, 1691–1695.
50. K. I. K. Yamada and T. Okuda, Impedance spectroscopy and structural investigation of Li⁺ conductor Li3InBr6, Solid State Ionics, 2002, 621–628.
51. W. Hönle, B. Hettich and A. Simon, Darstellung und Kristallstrukturen von LiGaCl4 and LiGaI4, Z. Naturforsch B. J. Chem. Sci., 1987, 42, 248–250.
52. S.-K. Jung, H. Gwon, G. Yoon, L. J. Miara, V. Lacivita and J.-S. Kim, Pliable Lithium Superionic Conductor for All-Solid-State Batteries, ACS Energy Lett., 2021, 6, 2006–2015.
53. H. J. Steiner and H. D. Lutz, Neue schnelle Ionenleiter vom Typ MMIIICl6 (MI = Li, Na, Ag; MIII = In, Y), Z. Anorg. Allg. Chem., 1992, 613, 26–30.
54. M. O. Schmidt, M. S. Wickleder and G. Meyer, TernaÈre Halogenide vom Typ A3MX6. VIII Zur Kristallstruktur von Li3InCl6, Z. Anorg. Allg. Chem., 1999, 625, 539–540.
55. X. Li, J. Liang, J. Luo, M. N. Banis, C. Wang, W. Li, S. Deng, C. Yu, F. Zhao, Y. Hu, T.-K. Sham, L. Zhang, S. Zhao, S. Lu, H. Huang, R. Li, K. R. Adair and X. Sun, Air-stable Li3InCl6 electrolyte with high voltage compatibility for all-solid-state batteries, Energy Environ. Sci., 2019, 12, 2665–2671.
56. X. Li, J. Liang, N. Chen, J. Luo, K. R. Adair, C. Wang, M. N. Banis, T. K. Sham, L. Zhang, S. Zhao, S. Lu, H. Huang, R. Li and X. Sun, Water-Mediated Synthesis of a Superionic Halide Solid Electrolyte, Angew. Chem., Int. Ed., 2019, 58, 16427–16432.
57. Y. Ishiguro, K. Ueno, S. Nishimura, G. Iida and Y. Igarashib, TaCl5-glassified Ultrafast Lithium Ion-conductive Halide Electrolytes for High-performance All-solid-state Lithium Batteries, Chem. Lett., 2023, 52, 237–241.
58. R. Kanno, Y. Takeda and O. Yamamoto, Structure, ionic conductivity and phase transforiviation of double chloride spinels, Solid State Ionics, 1988, 28–30, 1276–1281.
59. C. J. J. Van Loon and J. DE Jong, Some chlorides with the inverse spinel structure: Li₂TCl₄ (T=Mg,Mn,Fe,Cd), Acta Crystallogr., Sect. B: Struct. Crystallogr. Cryst. Chem., 1975, 31, 2549–2550.
60. H. D. Lutz, W. Schmidt and H. Haeuseler, Zur Kenntnis der Chlorid–Spinelle Li₂MgCl₄, Li₂MnCl₄, Li₂FeCl₄, Li₂CdCl₄, Z. Anorg. Allg. Chem., 1979, 453, 121–126.
61. R. Kanno, Y. Takada, K. Takada and O. Yamamoto, Ionic Conductivity and Phase Transition of the Spinel System Li_{2 − 2x} M _{1 + x} Cl₄ (M = Mg, Mn, Cd), J. Electrochem. Soc., 1984, 131, 469–474.
62. C. W. R. Nagel, J. Senker and H. D. Lutz, ⁶Li and ⁷Li MAS-NMR studies on fast ionic conducting inverse spinel-type Li_{2 − 2x}Mg_1+x Cl₄ and normal spinel-type Li₂ZnCl₄, Solid State Ionics, 2000, 130, 169–173.
63. H. D. Lutz, W. Schmidt and H. Haeuseler, Chloride spinels: a new group of solid lithium electrolytes, J. Phys. Chem. Solids, 1981, 42, 287–289.
64. Y. T. R. Kanno and O. Yamamoto, Ionic conductivity of solid lithium ion conductors with the spinel structure: Li₂MCl₄ (M = Mg, Mn, Fe, Cd), Mater. Res. Bull., 1981, 16, 999–1005.
65. C. J. J. Van Loon and D. Visser, Chlorides with the chrysoberyl structure: Na₂CoCl₄ and Na₂ZnCl₄, Acta Crystallogr., Sect. B: Struct. Crystallogr. Cryst. Chem., 1977, 33, 188–190.
66. M. Sassmannshausen, I. Solinas and H. D. Lutz, Crystal structure of dilithium zinc tetrachloride, Li₂ZnCl₄ with spinel- and olivine-type structure, Z. Kristallogr. Cryst. Mater., 1996, 211, 819–820.
67. A. Vijay and R. D. Eithiraj, Structural, electronic, optical, and thermoelectric properties of Li2ZnCl4 based on density functional theory computations, J. Phys. Chem. Solids, 2022, 171, 111037.
68. W. S. P. Kuske and H. D. Lutz, Neutron diffraction studies on spinel type Li2ZnCl4, Mater. Res. Bull., 1988, 23, 1805–1808.
69. R. Kanno, Y. Takeda, M. Mort and O. Yamamoto, Ionic conductivity and structure of double chloride Li₂ZnCl₄ in the Licl-ZnCl₂ system, Chem. Lett., 1989, 18, 223–226.
70. R. Kanno, K. Takeda and O. Yamamoto, Structure, ionic conductivity and phase transformation of double chloride spinels, Solid State Ionics, 1988, 28–30.
71. H. D. Lutz, Structural Phase Transition and Nonstoichiometry of Li2FeCl4-Neutron Diffraction Studies, J. Solid State Chem., 1993, 107, 245–249.
72. H. D. L. M. Schneider, Polymorphie und Pseudosymmetrie von Li2CoCl4, Z. Kristallogr. Cryst. Mater., 1993, 203, 183–197.
73. K. Suzuki, X-ray Studies on Precipitation of Metastable Centers in Mixed Crystals NaCl-CdCl2, J. Soc. Jpn., 1961, 16, 67–78.
74. H. D. Lutz, Y. KUSK and K. Wussow, Neue Lithiumchlorid-Suzukiphasen: Li₆MCI₈( M = Fe, Co, N), Z. Anorg. Allg. Chem., 1987, 553, 172–178.
75. G. Villeneuve, L. Latié, C. Christian and H. Paul, NMR of ⁷Li in the paramagnetic “Suzuki-type” phase Li₆VCl₈, Mater. Res. Bull., 1984, 19, 1515–1526.
76. T. M. L. Hanebali, C. Cros and P. Hagenmuller, Sur le systeme LiCl·VCl2, Mater. Res. Bull., 1981, 16, 887–901.
77. T. Oi and K. Miyauchi, Amorphous thin film ionic conductors of mLiF.nAlF₃, Mater. Res. Bull., 1981, 16, 1281–1289.
78. C. H. Wang, J. W. Liang, J. Luo, J. Liu, X. N. Li, F. P. Zhao, R. Y. Li, H. Huang, S. Q. Zhao, L. Zhang, J. T. Wang and X. L. Sun, A universal wet-chemistry synthesis of solid-state halide electrolytes for all-solid-state lithium-metal batteries, Sci. Adv., 2021, 7, eabh1896.
79. X. Wang, E. Zhao and W. Guo, A dual-matrix hybridization strategy for synthesis of cationic-anionic double-doped Li2.99Y0.99Zr0.01Br1.6Cl4.4 halide solid electrolyte, J. Solid State Electrochem., 2023, 27, 3597–3602.
80. S. Chen, C. Yu, C. Wei, L. Peng, S. Cheng and J. Xie, Revealing milling durations and sintering temperatures on conductivity and battery performances of Li_2.25Zr_0.75Fe_0.25Cl₆ electrolyte, Chin. Chem. Lett., 2023, 34, 107544.
81. P. Molaiyan, S. E. Mailhiot, K. Voges, A. M. Kantola, T. Hu, P. Michalowski, A. Kwade, V.-V. Telkki and U. Lassi, Investigation of the structure and ionic conductivity of a Li3InCl6 modified by dry room annealing for solid-state Li-ion battery applications, Mater. Des., 2023, 227, 111690.
82. H.-W. Liu, C.-C. Lin, P.-Y. Chang, S.-C. Haw, H.-S. Sheu, J.-M. Chen, C.-C. Chen, R.-J. Jeng and N.-L. Wu, Reducing oxy-contaminations for enhanced Li-ion conductivity of halide-based solid electrolyte in water-mediated synthesis, J. Solid State Electrochem., 2022, 26, 2089–2096.
83. X. He, Y. Zhu and Y. Mo, Origin of fast ion diffusion in super-ionic conductors, Nat. Commun., 2017, 8, 15893.
84. S. Wang, Q. Bai, A. M. Nolan, Y. Liu, S. Gong, Q. Sun and Y. Mo, Lithium Chlorides and Bromides as Promising Solid-State Chemistries for Fast Ion Conductors with Good Electrochemical Stability, Angew. Chem., Int. Ed., 2019, 58, 8039–8043.
85. Y. Wang, W. D. Richards, S. P. Ong, L. J. Miara, J. C. Kim, Y. Mo and G. Ceder, Design principles for solid-state lithium superionic conductors, Nat. Mater., 2015, 14, 1026–1031.
86. X. Jia, M. Zhou, R. Zhang and G. Li, Li+ transportation mechanisms in the halide solid-state electrolytes Li3YCl6 and Li3InCl6, Emerging Mater. Res., 2022, 11, 220–227.
87. R. Schlem, S. Muy, N. Prinz, A. Banik, Y. Shao-Horn, M. Zobel and W. G. Zeier, Mechanochemical Synthesis: A Tool to Tune Cation Site Disorder and Ionic Transport Properties of Li3MCl6 (M = Y, Er) Superionic Conductors, Adv. Energy Mater., 2019, 10, 1903719.
88. J. Qiu, M. Wu, W. Luo, B. Xu, G. Liu and C. Ouyang, Insights into Bulk Properties and Transport Mechanisms in New Ternary Halide Solid Electrolytes: First-Principles Calculations, J. Phys. Chem. C, 2021, 125, 23510–23520.
89. B. Ahammed and E. Ertekin, Configurational Disorder, Strong Anharmonicity, and Coupled Host Dynamics Lead to Superionic Transport in Li3YCl6 (LYC), Advanced Materials, 2024, 36, 2310537.
90. L. Hu, J. Zhu, C. Duan, J. Zhu, J. Wang, K. Wang, Z. Gu, Z. Xi, J. Hao, Y. Chen, J. Ma, J.-X. Liu and C. Ma, Revealing the Pnma crystal structure and iontransport mechanism of the Li3YCl6 solid electrolyte, Cell Rep. Phys. Sci., 2023, 4, 101428.
91. S. Hyun, H. Chun, M. Hong, J. Kang and B. Han, First-principles study on ultrafast Li-ion diffusion in halospinel Li2Sc2/3Cl4 through multichannels designed by aliovalent doping, J. Mater. Chem. A, 2023, 11, 4272–4279.
92. Y.-C. Yin, J.-T. Yang, J.-D. Luo, G.-X. Lu, Z. Huang, J.-P. Wang, P. Li, F. Li, Y.-C. Wu, T. Tian, Y.-F. Meng, H.-S. Mo, Y.-H. Song, J.-N. Yang, L.-Z. Feng, T. Ma, W. Wen, K. Gong, L.-J. Wang, H.-X. Ju, Y. Xiao, Z. Li, X. Tao and H.-B. Yao, A LaCl3-based lithium superionic conductor compatible with lithium metal, Nature, 2023, 616, 77–83.
93. V. L. Sawankumar, V. Patel, H. Liu, E. Truong, Y. Jin, E. Wang, L. Miara, R. Kim, H. Gwon, R. Zhang, I. Hung, Z. Gan, S.-K. Jung and Y.-Y. Hu, Charge-clustering induced fast ion conduction in 2LiXGaF3: A strategy for electrolyte design, Sci. Adv., 2023, 9, eadj9930.
94. Z. Liu, L. Qin, X. Cao, J. Zhou, A. Pan, G. Fang, S. Wang and S. Liang, Ion migration and defect effect of electrode materials in multivalent-ion batteries, Prog. Mater. Sci., 2022, 125, 100911.
95. C. Zhang, H. Zhuang, Z. Qi, X. Liu and Y. Ren, The effect of defects for the ion transport of Li3ScCl6 and Li3InCl6 with the interface of lithium metal anode: A first-principles study, Mater. Today Commun., 2023, 35, 105764.
96. E. Kim and Y. Kim, nvestigation of the effect of point defects on the Li-ion conductivity of Li3InCl6, Dalton Trans., 2022, 51, 18159–18168.
97. Z. Liu, S. Ma, J. Liu, S. Xiong, Y. Ma and H. Chen, High Ionic Conductivity Achieved in Li3Y(Br3Cl3) Mixed Halide Solid Electrolyte via Promoted Diffusion Pathways and Enhanced Grain Boundary, ACS Energy Letters, 2020, 6, 298–304.
98. E. Sebti, H. A. Evans, H. Chen, P. M. Richardson, K. M. White, R. Giovine, K. P. Koirala, Y. Xu, E. Gonzalez-Correa, C. Wang, C. M. Brown, A. K. Cheetham, P. Canepa and R. J. Clément, Stacking Faults Assist Lithium-Ion Conduction in a Halide-Based Superionic Conductor, J. Am. Chem. Soc., 2022, 144, 5795–5811.
99. P. C. Zhong, S. Gupta, B. W. Deng, K. J. Jun and G. Ceder, The effect of cation-disorder on lithium transport in halide superionic conductors, ACS Energy Letters, 2024, 9, 2775–2781.
100. Z. Gao, S. Wang, F. Sun, Z. Yu, H. Song, Z. Liu and H. Chen, Enhanced electrochemical properties of Li2MgCl4 by Zn substitution for all-solid-state batteries, J. Solid State Chem., 2023, 328, 124361.
101. R. Li, P. Lu, X. Liang, L. Liu, M. Avdeev, Z. Deng, S. Li, K. Xu, J. Feng, R. Si, F. Wu, Z. Zhang and Y.-S. Hu, Superionic Conductivity Invoked by Enhanced Correlation Migration in Lithium Halides Solid Electrolytes, ACS Energy Letters, 2024, 9, 1043–1052.
102. T. Usami, N. Tanibata, H. Takeda and M. Nakayama, Analysis of ion conduction behavior of Nb- and Zr-doped Li₃InCl₆ -based materials via material simulation, APL Mater., 2023, 11, 121107.
103. K.-H. Park, K. Kaup, A. Assoud, Q. Zhang, X. Wu and L. F. Nazar, High-Voltage Superionic Halide Solid Electrolytes for All-Solid-State Li-Ion Batteries, ACS Energy Letters, 2020, 5, 533–539.
104. S. Chen, C. Yu, S. Chen, L. Peng, C. Liao, C. Wei, Z. Wu, S. Cheng and J. Xie, Enabling ultrafast lithium-ion conductivity of Li 2 ZrCl 6 by indium doping, Chin. Chem. Lett., 2022, 33, 4635–4639.
105. X. Luo, X. Hu, Y. Zhong, X. Wang and J. Tu, Degradation Evolution for Li2ZrCl6 Electrolytes in Humid Air and Enhanced Air Stability via Effective Indium Substitution, Small, 2023, 20, 2306736.
106. K. Wang, Z. Gu, H. Liu, L. Hu, Y. Wu, J. Xu and C. Ma, High–Humidity–Tolerant Chloride Solid–State Electrolyte for All–Solid–State Lithium Batteries, Adv. Sci., 2024, 11, 2305394.
107. Y. Kim and S. Choi, Investigation of the effect of F-doping on the solid-electrolyte property of Li3InCl6, J. Power Sources, 2023, 567, 232962.
108. M. Yamagishi, C. Zhong, D. Shibata, M. Morimoto and Y. Orikasa, Effect of Fluorine Substitution in Li3YCl6 Chloride Solid Electrolytes for All-solid-state Battery, Electrochemistry, 2023, 91, 037002–037002.
109. Q. Wang, X. Ma, Q. Liu, D. Sun and X. Zhou, Fluorine-doped Li3InCl6 to enhance ionic conductivity and air stability, J. Alloys Compd., 2023, 969, 172479.
110. Z. Song, T. Wang, H. Yang, W. H. Kan, Y. Chen, Q. Yu, L. Wang, Y. Zhang, Y. Dai, H. Chen, W. Yin, T. Honda, M. Avdeev, H. Xu, J. Ma, Y. Huang and W. Luo, Promoting high-voltage stability through local lattice distortion of halide solid electrolytes, Nat. Commun., 2024, 15, 1481.
111. Q. Wang, Y. Zhou, X. Wang, H. Guo, S. Gong, Z. Yao, F. Wu, J. Wang, S. Ganapathy, X. Bai, B. Li, C. Zhao, J. Janek and M. Wagemaker, Designing lithium halide solid electrolytes, Nat. Commun., 2024, 15, 1050.
112. A. Emly, E. Kioupakis and A. Van der Ven, Phase Stability and Transport Mechanisms in Antiperovskite Li3OCl and Li3OBr Superionic Conductors, Chem. Mater., 2013, 25, 4663–4670.
113. Y. Zhao and L. L. Daemen, Superionic conductivity in lithium-rich anti-perovskites, J. Am. Chem. Soc., 2012, 134, 15042–15047.
114. S. Muy, J. Voss, R. Schlem, R. Koerver, S. J. Sedlmaier, F. Maglia, P. Lamp, W. G. Zeier and Y. Shao-Horn, High-Throughput Screening of Solid-State Li-Ion Conductors Using Lattice-Dynamics Descriptors, iScience, 2019, 16, 270–282.
115. H. Kageyama, K. Hayashi, K. Maeda, J. P. Attfield, Z. Hiroi, J. M. Rondinelli and K. R. Poeppelmeier, Expanding frontiers in materials chemistry and physics with multiple anions, Nat. Commun., 2018, 9, 772.
116. J. Wang, F. Chen, L. Hu and C. Ma, Alternate Crystal Structure Achieving Ionic Conductivity above 1 mS cm−1 in Cost-Effective Zr-Based Chloride Solid Electrolytes, Nano Lett., 2023, 23, 6081–6087.
117. K.-H. Park, S. Y. Kim, M. Jung, S.-B. Lee, M.-J. Kim, I.-J. Yang, J.-H. Hwang, W. Cho, G. Chen, K. Kim and J. Yu, Anion Engineering for Stabilizing Li Interstitial Sites in Halide Solid Electrolytes for All-Solid-State Li Batteries, ACS Appl. Mater. Interfaces, 2023, 15, 58367–58376.
118. B. Li, Y. Li, H.-S. Zhang, T.-T. Wu, S. Guo and A.-M. Cao, Fast Li+-conducting Zr4+-based oxychloride electrolyte with good thermal and solvent stability, Science China Materials, 2023, 66, 3123–3128.
119. S. Zhang, F. Zhao, J. Chen, J. Fu, J. Luo, S. H. Alahakoon, L.-Y. Chang, R. Feng, M. Shakouri, J. Liang, Y. Zhao, X. Li, L. He, Y. Huang, T.-K. Sham and X. Sun, A family of oxychloride amorphous solid electrolytes for long-cycling all-solid-state lithium batteries, Nat. Commun., 2023, 14, 3780.
120. S. Zhang, F. Zhao, L.-Y. Chang, Y.-C. Chuang, Z. Zhang, Y. Zhu, X. Hao, J. Fu, J. Chen, J. Luo, M. Li, Y. Gao, Y. Huang, T.-K. Sham, M. D. Gu, Y. Zhang, G. King and X. Sun, Amorphous Oxyhalide Matters for Achieving Lithium Superionic Conduction, J. Am. Chem. Soc., 2024, 146, 2977–2985.
121. L. Hu, J. Wang, K. Wang, Z. Gu, Z. Xi, H. Li, F. Chen, Y. Wang, Z. Li and C. Ma, A cost-effective, ionically conductive and compressible oxychloride solid-state electrolyte for stable all-solid-state lithium-based batteries, Nat. Commun., 2023, 14, 3807.
122. T. Dai, S. Wu, Y. Lu, Y. Yang, Y. Liu, C. Chang, X. Rong, R. Xiao, J. Zhao, Y. Liu, W. Wang, L. Chen and Y.-S. Hu, Inorganic glass electrolytes with polymer-like viscoelasticity, Nat. Energy, 2023, 8, 1221–1228.

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