English

• 1. Introduction

  Na-ion batteries (SIBs) are regarded as a promising next generation electrochemical energy storage technology due to the abundance, wide distribution, low cost of sodium resources, as well as similar chemical properties of Na and Li elements, which have been widely investigated in the past decades [1-10]. Kummer et al. first reported a fast Na-ion conduction in Na-β″-Al₂O₃ in 1967 [11]. Later, Ford et al. designed a high-temperature Na-S battery based on this material in 1968. Since then, researchers have made a lot of efforts in sodium ion battery. Up to data, more than 20 companies around the world are working on the development of SIBs, and SIBs are gradually achieving practical application. Traditional SIBs using liquid organic electrolytes are prone to volatility and leakage, resulting in safety risks [12-14]. Differently, all-solid-state SIBs are significantly safer because of the use of solid-state electrolytes (SSEs) with non-volatile and non-leakage features [15-18]. More importantly, solid-state SIBs with sodium metal as anode can provide high energy density comparing to hard carbon anode.

  Solid-state SIBs have become one of hot topics in the future energy storage field [19,20]. The ionic conductivity and stability of SSEs as well as their compatibility with electrode materials in solid-state SIBs are the important factors affecting the performance of SIBs [21-23]. Therefore, it is imperative to synthesize and optimize new Na-ion SSE materials in addition to optimizing and improving electrode materials. Na-ion SSEs can be mainly divided into two catalogs: inorganic and organic polymer SSEs. Although organic polymer SSEs have good flexibility and low interfacial resistance with electrodes [24-29], they face some intrinsic natures, such low ionic conductivity at r.t., narrow electrochemical window, and poor mechanical strength, which hinders their application [30-32]. Adding inorganic fillers and synthesizing inorganic-organic polymer composites are two widely adopted strategies to improve ionic conductivity of organic polymer SSEs. Compared with polymer SSEs, inorganic ones have higher ionic conductivity at r.t. (> 10⁻³ S/cm), wider electrochemical window, better thermal stability, and higher mechanical strength can effectively inhibit the growth of sodium dendrites, and thus have been attracted more attention [23,33-36].

  Over the past decades, considerable progress about the Na-ion SSEs have been achieved. As early as 1968, the fast conduction of sodium ions in oxide Na-β″-Al₂O₃ was reported by Kummer et al. firstly [11]. After that, the Nasicon material with the general formula Na_1+xZr₂Si_xP_3-xO₁₂ (0 ≤ x ≤ 3) was proposed by Goodenough and Hong et al. in 1976 [37,38]. The highest ionic conductivity was obtained up to 6.7 × 10⁻⁴ S/cm and 0.2 S/cm at 25 ℃ and 300 ℃, respectively, for Na₃Zr₂Si₂PO₁₂ (i.e., x = 2 of Na_1+xZr₂Si_xP_3-xO₁₂). In 2011, Evstygneeva et al. reported a new type of oxide Na-ion SSE with layered structure [39]. In addition to oxides mentioned above, anti-perovskites, complex hydride, and sulfide SSEs are also being studied extensively [40-42]. The highest Na-ion conductivity at r.t. is reached as high as 7 × 10⁻² S/cm for Na₂(CB₉H₁₀)(CB₁₁H₁₂). A systematic understanding of the structure, synthesis, modification, and application of Na-ion SSEs is very important, but urgent for the development of all-solid-state SIBs with high performance.

  In this review, we focus on inorganic Na-ion SSEs, including Na-β/β″-Al₂O₃, Nasicon, layered oxides, anti-perovskite, complex hydrides and halides, etc. We first discuss the mechanism of Na-ion transport and the structures of these typical SSEs. We then specify the synthesis methods and enumerate the strategies for enhancing the conductivity and stability of Na-ion SSEs. After summarizing the application of SSEs in solid-state SIBs, and the facial designing and advanced characterizations for interfaces of all-solid-state SIBs, we put forward some view on the future development of inorganic SSEs.

  2. Fundamentals of Na-ion SSEs

  2.1 Mechanism of Na⁺ transport in inorganic SSEs

  Classical models describing Na-ion transport in SSEs treats ion leaps as uncorrelated and independent, which can be described as the transport of a single ion. Ion transportation can be achieved by hopping the mobile ions through various point defects, such as vacancies or interstitials [43,44]. The typical ion diffusion mechanism can be classified into three types: (1) Direct hopping mechanisms between vacancies; (2) direct hopping mechanisms between interstitials; (3) "knock-off" like mechanism between vacancies and interstitials (Figs. 1a-c). Within the crystal structure framework, individual ions are able to hop from one lattice site to another through connected diffusion channels, which is what is known as ion migration. The energy barrier for this ion transport process is determined by the framework. The highest energy along the diffusion pathway determines the energy barrier E_a of the mobile ion. However, the classical diffusion model contradicts the results of existing scientific research. He et al. reported that the ion transport mechanism is a multi-ion concerted migration mechanism for multiple adjacent ions hopping into their nearest sites simultaneously through AIMD simulations [43]. Increasing the Na concentration in SSEs yields high ionic conductivity. The low-energy sites in the structure of Na-ion conductors are preferentially occupied by Na-ions, while the excess Na-ions occupy the high-energy sites as the low-energy sites are fully occupied and cannot accommodate more Na-ions. In a multi-ion concerted migration process, Na-ions located at high-energy sites can effectively migrate downhill and cancel out some of the energy barrier that is felt by Na-ions occupying low-energy sites and undergoing upward migration. This results in lowered energy barriers for Na-ion diffusions (Fig. 1d).

  Figure 1

  

  Figure 1.  Schematic diagram of four typical ion diffusion mechanisms of NaF, purple represents F ions, yellow represents occupied sodium ion sites, gray represents sodium ion vacancies, and green represents interstitial sites. (a) Direct hopping mechanisms between vacancies; (b) Direct hopping mechanisms between interstitials; (c) "Knock-off" like mechanism between vacancies and interstitials; (d) Conceptual diagram of the energy barriers required for single ion diffusion and multiple ions to diffuse in concert.

  DownLoad: Full-Size Img PowerPoint

  2.2 Crystal structures of various inorganic Na-ion SSEs

  Inorganic Na-ion SSEs are a class of solid materials whose ionic conductivity are close to or exceeds that of the molten electrolyte. Two types of sublattices exist in the structure of SSEs. The first is a skeleton of non-moving ions that only undergo thermal vibrations with no ion transport, which cannot contribute to ion conduction in the structure. The second type consisting of the sublattice of mobile ions, which occupy the original lattice sites and undergo position shifts between other vacancies with similar activation energies, may provide a higher ion conductivity. Recently, inorganic Na-ion conductors can be classified into several types based on their compositions, such as oxides, sulfides, complex hydrides, and halides. The detailed crystal structure information of these different SSEs will be carefully summarized in the following parts of this section.

  2.2.1   Na-β/β″-Al₂O₃

  To date, Na-β/β″-Al₂O₃ stands as the sole commercially employed Na-ion SSE, finding primary applications in high-temperature sodium-sulfur batteries and solid-state SIBs [45-47].

  Na-β/β″-Al₂O₃ exhibits a layered structure characterized by two-dimensional Na-ion conduction occurring between the layers. Fig. 2 illustrates the presence of two distinct crystal structures in Na-β/β″-Al₂O₃, denoted as Na-β-Al₂O₃ and Na-β″-Al₂O₃, respectively. The chemical composition of Na-β-Al₂O₃ is typically represented as Na₂O(8–11) Al₂O₃. It adopts a hexagonal crystal system (P6₃/mmc) with lattice parameters a = b = 5.58 Å, c = 22.45 Å. Na-β-Al₂O₃ exhibits a stacking structure composed of two spinels [48-51]. The spinel layers adjacent to the Na-conducting layer are linked via O²⁻ within the Na-conducting layer, wherein the electrostatic attraction of O²⁻ to Na⁺ is different in the two crystal structures.

  Figure 2

  

  Figure 2.  Crystal structure of Na-β/β″-Al₂O₃. The yellow ball represents the sodium ion, the red represents the oxygen ion, and the gray represents the aluminum ion.

  DownLoad: Full-Size Img PowerPoint

  Among them, the large electrostatic attraction of O²⁻ and Na⁺ in Na-β-Al₂O₃ leads to a smaller unit cell volume, and the Na-conducting layer can accommodate less sodium ions, so it has a lower ionic conductivity. In contrast, the reduced electrostatic attraction between O²⁻ and Na⁺ in the Na-β″-Al₂O₃ structure results in a larger unit cell volume compared to the Na-β-Al₂O₃ structure. Moreover, the Na-β″-Al₂O₃ structure exhibits a higher sodium ion content within the two-dimensional ion transport plane. These factors contribute to enhanced facilitation of sodium ion migration and consequently lead to higher ionic conductivity in the Na-β″-Al₂O₃ structure. However, the thermodynamic stability of Na-β″-Al₂O₃ is relatively poor, posing challenges in obtaining a pure phase of this material. At high temperatures, Na-β″-Al₂O₃ is prone to decomposition into Al₂O₃ and Na-β-Al₂O₃. Moreover, the ionic conductivity of polycrystalline Na-β″-Al₂O₃ is influenced by the ratio between Na-β″-Al₂O₃ and Na-β-Al₂O₃ phases. Furthermore, Na-β″-Al₂O₃ exhibits greater sensitivity to moisture in the air and possesses inferior mechanical strength, thereby imposing additional constraints on its practical application.

  2.2.2   Nasicon

  Nasicon-type (Na_1+xZr₂Si_xP_3-xO₁₂ (0 ≤ x ≤ 3)) SSE is another prominent class of Na-ion conductors that has received significant attention for its potential application in Na-ion SSEs. It was firstly proposed and studied by Goodenough and Hong in 1976 [37,38]. When x = 2, the composition Na₃Zr₂Si₂PO₁₂ exhibited remarkable ionic conductivity of 6.7 × 10⁻⁴ S/cm at r.t., 1.2 × 10⁻³ S/cm at 60 ℃ and 0.2 S/cm at 300 ℃. In Nasicon-structured SSEs, two distinct crystal structures can be observed: a monoclinic structure (C2/c) within the compositional range of 1.6 ≤ x ≤ 2.2, and a rhombohedral structure (R-3c). These crystal structures are illustrated in Figs. 3a and b, respectively. The monoclinic phase can transform into the rhombohedral phase upon heating to a temperature range of 420–450 K. The rhombohedral Nasicon structure exhibits a three-dimensional framework composed of corner-shared ZrO₆ and (P/Si)O₄ polyhedral units. Within this framework, Na-ions are situated and can be conducted isotropically within the rhombohedral Nasicon structure, there exist two crystallographically distinct Na-ion sites, denoted as Na1 and three Na2 sites. The monoclinic phase can be considered a distortion of the rhombohedral phase, whereby the structural units ZrO₆ and (P/Si)O₄ are remained during the distortion process but the overall symmetry of the structure is diminished. Furthermore, the structural distortion induces a modification in the occupancy of Na-ions, resulting in the division of the three Na2 sites into newly formed Na sites, namely one Na2 site and two Na3 sites. These additional Na-ion sites, created as a result of the distortion, serve as exchange sites for facilitating Na-ion transport. This phenomenon enhances the overall Na-ion mobility and promotes ion conductivity (Figs. 3c-e). In the monoclinic phase, the diffusion pathways for Na-ions include the Na1-Na2 channel and the Na1-Na3 transport channel. In contrast, the diffusion channel in the rhombohedral phase is primarily through the Na1-Na2 channel.

  Figure 3

  

  Figure 3.  Structures of Nasicon, different yellow balls represent different sites of sodium ions, red represents oxygen ion, blue octahedron represents [ZrO₆] octahedron, pink tetrahedron represents [Si/PO₄] tetrahedron: (a) Rhombohedral phase; (b) Monoclinic phase; (c) Na1-Na2 transport channel of rhombohedral phase; (d) Na1-Na2 transport channel and (e) Na1-Na3 transport channel of monoclinic phase.

  DownLoad: Full-Size Img PowerPoint

  2.2.3   P2-type layered oxides

  P2-type layered oxides are a new type of Na-ion SSE, which was first reported by Evstygneeva et al. in 2011 [39]. Four layered compounds, namely Na₂M₂TeO₆ (M = Ni, Co, Zn, Mg), were successfully synthesized and characterized, revealing their hexagonal crystal structures (Figs. 4a and b). Na₂M₂TeO₆ exhibits a honeycomb-like architecture, where each TeO₆ octahedron is surrounded by six MO₆ octahedra in the layers via edge-sharing interactions. Na-ions are located between layers occupying different positions, showing partial occupancy. The anisotropic conduction of Na-ions on the ab plane can be attributed to the limited permeability of Na-ions through the oxide layer. In this compound, three crystal graphically distinct Na-ion sites, denoted as Na1, Na2, and Na3, are present in each unit cell, with respective multiplicities of three, two, and one. And they have similar parameters of the hexagonal P2-type layered compounds: a = 5.20–5.28 Å, c = 11.14–11.31 Å, but different stacking sequences along c. In which Na₂Mg₂TeO₆ and Na₂Co₂TeO₆ were not obtain pure phases. And Na₂Ni₂TeO₆ exhibits a highest ionic conductivity of 10.1–10.8 S/m at 300 ℃. However, its ionic conductivity can only reach 0.0008–0.0034 S/cm at r.t., which may be related to the generation of glass during the sintering process. The formation of glass phase increased its density, but the poor conductivity of the glass phase results in a high activation energy of up to 0.55 eV. Additionally, it is worth noting that the D-shells of the Ni and Co cations in the Na₂Ni₂TeO₆ and Na₂Co₂TeO₆ samples exhibit partial filling, and the oxidation state may undergo changes to enhance electronic conductivity. The measured electronic conductivity contribution of Na₂Co₂TeO₆, determined through the DC polarization method, is found to be only between 0.0031 S/cm and 0.0044 S/cm at 300 ℃. Similarly, the electronic contribution of Na₂Ni₂TeO₆ is even lower. However, it should be noted that these materials may not be the optimal choices for SSEs due to the presence of variable valence elements such as Ni and Mn. During the charging and discharging process of the battery, redox reactions involving these elements can occur, resulting in electrolyte instability and self-discharge behavior of the battery.

  Figure 4

  

  Figure 4.  Crystal structures of (a) Na₂Ni₂TeO₆ and (b) Na₂M₂TeO₆ (M = Zn, Co, Mg). Different yellow balls represent different sites of sodium ions, the green octahedron represents [TeO₆], the purple octahedron represents [NiO₆], and the pink octahedron represents [ZnO₆].

  DownLoad: Full-Size Img PowerPoint

  2.2.4   Anti-perovskites

  Perovskites represent a class of structural types characterized by the connectivity of octahedra through shared vertices. From the perspective of structural chemistry, perovskites have a rich elemental composition, as well as a symmetrical diversity and a large number of symbiotic structures due to octahedral distortion. Anti-perovskites, being isostructural to perovskites, can be described by the general formula ABX₃, where the positive and negative ions occupy positions opposite to those in the perovskite structure [52,53]. Precisely, in perovskite structures, the A-site and B-site ions are cations, whereas the X-site is occupied by an anion, as exemplified by CaTiO₃. In contrast, anti-perovskites feature an arrangement where the A-site and B-site ions are anions, while the X-site is occupied by a cation, as observed in Na₃OCl (Fig. 5). On the one hand, in terms of crystal structure, anti-perovskites exhibit a comparable structural tolerance and diversity to perovskites. Additionally, in anti-perovskites, the oxygen ions are positioned at the center, while the larger A-site ions are situated at the face center of the cube, enabling jump-type ionic conduction. In recent years, Na-rich anti-perovskites demonstrated an exciting solution as a new class of ionic conductors [54-56]. These materials primarily consist of Na, O, and either Cl or Br, rendering them lightweight in nature. They exhibit favorable ionic conductivity and demonstrate excellent electrochemical stability towards Na. The anti-perovskite structure offers inherent flexibility and adjustability, facilitating chemical modifications aimed at enhancing the ionic conductivity.

  Figure 5

  

  Figure 5.  Plan view of Na₃OCl along the [111] direction. Blue ball represents oxygen ion, green represents chloride ion, yellow represents sodium ion.

  DownLoad: Full-Size Img PowerPoint

  2.2.5   Sulfide

  Sulfide-based SSEs exhibit higher ionic conductivity and lower grain boundary impedance compared to oxide SSEs, owing to the larger atomic radius of sulfur (S) relative to oxygen (O). The introduction of S atoms in place of O atoms leads to lattice expansion and the formation of pathways conducive to Na-ion transport. Furthermore, the lower electronegativity of S compared to O results in a reduced binding capacity for Na-ions, facilitating their movement and thereby enhancing the overall ionic conductivity. Moreover, sulfide SSEs possess soft characteristics and can be synthesized at lower temperatures, eliminating the need for high-temperature ceramic sheet sintering required for oxide SSEs [57]. The powdered sulfide material can be cold-pressed to establish favorable electrode-electrolyte interfaces, simplifying the preparation process of solid-state batteries. However, it should be noted that sulfide SSEs exhibit poor stability and are susceptible to reaction with moisture in the air, leading to the release of toxic H₂S gas [58-60]. This limitation significantly hinders their application in practical settings.

  Na₃PnS₄ (Pn = S, Sb): Na₃PS₄ exhibits two distinct crystal structures: tetragonal phase (P-42₁c; a = b = 6.9520 Å, c = 7.0699 Å) and cubic phase (I-43 m; a = b = c = 7.0699 Å). These two phases differ slightly in their structural arrangement. In the tetragonal phase, Na-ions distributed in one tetrahedral site and one octahedral site, while in the cubic phase, Na-ion is distributed in two distorted interstitial sites. The structure of Na₃PS₄ is shown in Figs. 6a and b. Among them, the conductivity of the cubic Na₃PS₄ is much higher than that of tetragonal Na₃PS₄, primarily attributed to their structural disparities. However, recent investigations have confirmed that the ionic conductivity of the tetragonal phase can also be higher than that of the cubic phase. In 2018, Takeuchi et al. synthesized tetragonal Na₃PS₄ with Na-rich vacancies by an improved preparation process [35], which had an ionic conductivity of 3.39 × 10⁻³ S/cm at 25 ℃. This significant finding challenged the previously held notion regarding the ionic conductivity of the tetragonal and cubic phases. Subsequent theoretical investigations have further confirmed that the local structure of the two crystal forms of Na₃PS₄ does not exhibit noticeable differences. As a result, the disparity in crystal structure may not necessarily be a determining factor influencing its ionic conductivity. Instead, the variation in ionic conductivity could be attributed to the concentration of vacancies within the system.

  Figure 6

  

  Figure 6.  Crystal structures of Na₃PS₄. (a) Tetragonal and cubic phase, (b) crystal structure of Na₁₁Sn₂PS₁₂. The blue ball represents tin ions, pink represents phosphorus ions, green represents sulfur ions, and yellow represents sodium ions.

  DownLoad: Full-Size Img PowerPoint

  Na₁₁Sn₂PnSn₁₂ (Pn = P, Sb; Sn = S, Se): In recent years, Na₁₁Sn₂PS₁₂ has emerged as a promising Na-ion conductor, garnering significant interest among researchers [61]. Zhang et al. first reported a Na₁₁Sn₂PS₁₂ electrolyte with a three-dimensional structure in 2017 [62]. Remarkably, this electrolyte demonstrated has an ionic conductivity of 1.4 × 10⁻³ S/cm at r.t. It is proved to be a new structure with space group I4₁/acd: 2, a = b = 13.6148(3) Å, c = 27.2244(7) Å. The structural resemblance of Na₁₁Sn₂PS₁₂ to LGPS (Li₁₀GeP₂S₁₂) is evident (Fig. 6b). Within this framework, Na⁺ are situated within octahedral sites, facilitating their transport along an interconnected three-dimensional pathway comprised of isoenergetic Na-S octahedra, which is further augmented by partial vacancy intersections. Notably, the ion transport pathway observed in Na₁₁Sn₂PS₁₂ represents the first instance of a fully three-dimensional faceted octahedral site pathway documented in sulfide SSEs.

  2.2.6   Halide

  The general formula for a sodium halide solid electrolyte is Na_3-xM_1-xM'_xX₆ (M can be a transition metal element, a lanthanide, a boron element, or a nitrogen element; M' can partially or completely replace Zr; X can be one or more halogens), and the radii of M and X determine the crystal structure of the material to some extent [44]. Common crystal structures include triangle P31c, triangle R-3, monoclinic crystal P21/n and monoclinic crystal C2/m. The crystal structure of the halide solid electrolyte is composed of stable NaX₆ and MX₆ octahedrons. The different NaX₆ connection modes determine the migration path of Na⁺ and affect the conductivity of Na⁺. For example, the NaCl₆ octahedra in Na₃YCl₆ are stacked face to face along the c-axis or side to side along the and axis. The former provides a fast migration path along the c-axis by directly migrating Na⁺ from adjacent octahedral locations (Fig. 7).

  Figure 7

  

  Figure 7.  Crystal structure of Na₃YCl₆. Different yellow balls represent different sites of sodium ions, the gray color represents yttrium ions and the green color represents chloride ions.

  DownLoad: Full-Size Img PowerPoint

  2.2.7   Complex hydrides

  Complex hydrides are a group of compounds composed of metal cations (e.g., Na⁺, Li⁺) and complex anions ([BH₄]⁻, [B₁₀H₁₀]²⁻, [B₁₂H₁₂]²⁻), which have received extensive attention due to their lighter weight, better electrochemical stability and excellent deformation properties. In 2012, Orimo et al. first presented the conduction of Na-ion in complex hydrides [63]. The ionic conductivity of NaAlH₄ was determined to be 2.1 × 10⁻¹⁰ S/cm, while that of Na₃AlH₆ was measured at 6.4 × 10⁻⁷ S/cm at r.t. Notably, the ionic conductivity of Na₃AlH₆ significantly increased to 4.2 × 10⁻⁴ S/cm at 433 K. The crystal structures of NaAlH₄ and Na₃AlH₆ are depicted in Figs. 8a and b, respectively. Although high ionic conductivity was not achieved in this study, it served as a significant milestone in exploring complexed hydrides as SSEs for Na-ions. Subsequently, they studied the Na(BH₄)-Na(NH₂)-NaI system [64]. The sample with the composition Na(BH₄)_0.5(NH₂)_0.5 exhibited the best ionic conductivity, reaching 2 × 10⁻⁶ S/cm. The excellent performance is attributed to its special anti-calcite structure, which has Na-ion vacancies and facilitates the transport of Na-ions. They believed that this material is promising for use in SIBs.

  Figure 8

  

  Figure 8.  Crystal structures of (a) NaAlH₄, (b) Na₃AlH₄, (c) low-temperature ordered monoclinic phase and (d) high-temperature disordered cubic phase. Reprinted with permission [65]. Copyright 2014, Royal Society Chemistry.

  DownLoad: Full-Size Img PowerPoint

  In 2014, Udovic et al. reported a novel complexed hydride SSE denoted as Na₂B₁₂H₁₂, this material exhibits an intriguing ordered-disordered phase transition, where the disordered phase driven by entropy becomes thermodynamically stable at elevated temperatures, enabling enhanced ion conduction kinetics. This unique phase behavior contributes to the fast transport of ions inside the material [65]. The two distinct structures of Na₂B₁₂H₁₂ are shown in Figs. 8c and d. The material is an ordered monoclinic phase at low temperature, while at higher temperatures, it undergoes a phase transition to a disordered body-centered cubic structure with a significant presence of cation vacancies. Notably, the disordered cubic Na₂B₁₂H₁₂ exhibits a remarkable enhancement in ionic conductivity, reaching up to 0.1 S/cm within the temperature range of 540 K to 573 K. However, the ionic conductivity of Na₂B₁₂H₁₂ at r.t. is not optimal, leading to a limited temperature range for its effective utilization as SSE.

  3. Synthesis methods for different sodium solid electrolytes

  The processing of SSEs typically involves two main steps. The first step involves synthesizing substances that have a specific chemical composition, physical state, crystal structure, and desired properties. The second step involves fabrication of the SSE materials, such as shaping and fixing them through firing, sintering, filming, and other methods. However, with inorganic oxide SSEs, it can be difficult to give them desired shapes, forms, and sizes due to their inherent brittleness, which limits the types of fabrication methods that can be used. In contrast, sulfide materials like Na₃PS₄ are generally more deformable and can be shaped under local stresses. However, the specific synthesis method used can also affect the resulting ionic conductivity of the SSE.

  For oxide SSE ceramics, the classical two-step processing method by solid-state reaction usually requires high temperature and long-time sintering, even higher than 1200 ℃ for several days sometimes. And high temperature is needed to facilitate the diffusion of the ions through the product layer. Therefore, increasing the surface area of reactants, preparing a homogeneous mixture of reactants, adding of impurities that can cause distortion in the crystal lattice and weaken the chemical bonds, etc. can promote the solid-state reaction. Additionally, metastable structure with higher formation energy comparing to stable one, such as Na-β″-Al₂O₃ comparing to Na-β-Al₂O₃, could exhibit high ionic conductivity. Introducing foreign ions or impurities stabilize such metastable structure is important for deigning SSEs with excellent Na-ion conduction. This part will describe in the following part.

  Liquid-phase method is another commonly employed technique for synthesizing SSEs. Compared with the solid phase method, the power produced by the liquid-phase method is more uniform. This approach involves dissolving the raw materials in a suitable solvent, followed by a stirring-induced liquid-phase reaction and subsequent heat treatment. It is particularly useful for coating electrode materials with electrolytes. However, when working with sulfide materials, caution must be exercised as they can undergo undesired side reactions with polar solvents. In Fig. 9, schematic diagrams depicting the synthesis processes using solid-state and liquid-phase methods are presented.

  Figure 9

  

  Figure 9.  Schematic diagram of synthesis of (a) solid state reaction and (b) liquid phase methods.

  DownLoad: Full-Size Img PowerPoint

  Na-β/β″-Al₂O₃ is generally synthesized using a conventional solid-phase method, which involves calcination at temperatures ranging from 1200 ℃ to 1250 ℃, followed by sintering at 1600 ℃. This method has been extensively studied due to its high yield and straightforward preparation process. However, this high-temperature sintering tends to lead to the loss of Na and the abnormal grain growth, which will weaken the materials and adversely affect its ionic conductivity. Furthermore, this method can produce NaAlO₂ at grain boundaries, which is hygroscopic and make Na-β/β″-Al₂O₃ susceptible to degradation. Therefore, Na-β/β″-Al₂O₃ prepared by solid-phase method is usually unstable. As a result, several methods including sol-gel synthesis, coprecipitation method, chemical vapor deposition and vapor-phase synthesis have proven to be effective alternatives to the traditional solid-state reaction. Shan et al. synthesized TiO₂ doped Na-β/β″-Al₂O₃ via a sol-gel method with tetrabutyl titanate (TBT) as the precursor for TiO₂ [66], the relative density reached 99.8%, this result showed that this method can accelerate grain growth and improve the densification. Butee et al. [67] synthesized Na-β/β″-Al₂O₃ by citrate sol-gel rote using glycerine as fuel, instead of ethylene glycol, and the synthesized Na-β/β″-Al₂O₃ has high density and high ionic conductivity. Virkar et al. [68] used Y-ZrO₂ and α-Al₂O₃ as raw materials and obtained dense ceramic sheets by a vapor phase process after high temperature sintering at 1450 ℃ for 17 h. Using this method can effectively enhance the chemical stability and mechanical strength of Na-β/β″-Al₂O₃.

  The traditional solid-phase method has been the most commonly used approach for synthesizing Nasicon-type electrolytes, owing to its high yield and relatively simple preparation process. While the traditional solid-phase method is effective, it has some limitations, including the need for high temperatures (over 1200 ℃) and long duration (over 20 h), which can result in abnormal grain growth and loss of Na and P during sintering. Additionally, poorly conducting ZrO₂ phases may be formed at grain boundaries due to the high sintering temperatures, leading to a decrease in ionic conductivity. Therefore, the density and the purity of the materials are the key to affect the ionic conductivity of the grain boundaries. A lot of works has been devoted to the synthesis of pure phase Na₃Zr₂Si₂PO₁₂. Naqash et al. synthesized Nasicon compositions with sodium excess, i.e., Na₃Zr₂Si₂PO₁₂ + xNa₂O (0 ≤ x ≤ 0.2), demonstrated that the excess Na effectively compensated for the volatilized Na element at high temperature and effectively suppressed the formation of impurity of impurity phase [69]. The sample with x = 0.2 showed the least sodium deficiency and the highest ionic conductivity.

  In addition, sol-gel method, hydrothermal method, solution-assisted solid-phase reaction (SASSR), liquid feed-flame spray pyrolysis (LF-FSP), spray freeze/freeze drying and low-temperature rapid microwave sintering method are also widely used to synthesize Nasicon materials with high densities. Compared with the traditional solid-phase method, sol-gel method can reduce the sintering temperature required for densification effectively, thus avoiding the formation of impurity phases. And the samples prepared by sol-gel method is more uniform than solid-phase method, therefore, this method is widely used for the synthesis of Nasicon electrolytes. Zhang et al. demonstrated that the sintering temperature could be reduced by the sol-gel method [70]. The results showed that the main phase of Nasicon prepared by sol-gel method was formed when sintered at 850 ℃, and the sample obtained by sintering at 950 ℃ was well crystalline.

  Solid phase reactions and liquid phase method are the main method for sulfide SSEs. Solid phase reaction involves high temperature sintering and annealing to obtain SSEs with a high degree of crystallinity. For sulfide SSEs, the sintering temperature is usually 200–800 ℃. The liquid phase method involves dissolving various raw materials in a solvent to achieve homogeneous mixing of the reactants. The reduced particle size of the precursor powder shortens the mass transfer process, reduces the temperature of the reaction and densification and avoids the formation of tramp phases. A few typical examples of the preparation of sulfide electrolytes using different methods are given next.

  In 2015, a cubic Na₃PS₄ SSE was firstly synthesized by liquid-phase reaction with NMF [71]. And its ionic conductivity was 2.6 × 10⁻⁶ S/cm. The presence of impurity Na₃POS₃ in the electrolyte made by this method is one of the reasons for the low ionic conductivity, the conductivity of the impurity phase is very low (about 10⁻⁸ order of order of magnitude), and the second reason may be due to the increase in R_gb leading to the decrease of conductivity. And changing the selecting compositions and solvents may improve ionic conductivity. Kim et al. synthesized Na₃SbS₄ using Na₂S, Sb₂S₃ and S as raw materials using the solution method, which used water as a solvent to avoid the reaction between organic solvents and sulfides [72].

  4. Strategies for enhancing the conductivity and stability of Na-ion SSEs

  The suitable SSE is crucial in the design of all-solid-state SIBs. An ideal SSE ought to have high ionic conductivity at r.t., electronic insulation, excellent electrochemical and interfacial stability. Among these characteristics, ionic conductivity stands out as the most significant performance parameter for SSEs. Consequently, extensive research works was devoted to enhancing the ionic conductivity of various electrolyte materials. Notably, sulfide SSEs are particularly sensitive to moisture and can undergo hydrolysis, releasing H₂S gas. Therefore, improving both the ionic conductivity and stability of sulfide SSEs is of utmost importance. Fig. 10 illustrates the reported conductivities of different inorganic sodium ion conductors, which are discussed in the subsequent sections.

  Figure 10

  

  Figure 10.  Reported conductivities of Na-ion SSEs.

  DownLoad: Full-Size Img PowerPoint

  4.1 Modification of Na-β/β″-Al₂O₃

  The pure phase of Na-β″-Al₂O₃ is thermodynamically metastable, and upon heating to 1500 ℃, it decomposes into Al₂O₃ and Na-β-Al₂O₃. Meanwhile Na-β″-Al₂O₃ is moisture sensitive with poor mechanical properties [73]. Ionic doping is commonly employed for stabilizing the Na-β″-Al₂O₃ phase. Zhu et al. discovered that after the introduction of Mg²⁺ and Li⁺, the β"-Al₂O₃ content in the obtained samples was above 90% [74]. While the β″-Al₂O₃ phase content in the undoped samples was only 5%, which indicates that Mg and Li can stabilize the Na-β″-Al₂O₃ effectively, and it was also found that Mg²⁺ was more beneficial to improve the symmetry of Al(IV) in the Na-β″-Al₂O₃ phase than Li⁺. Wei et al. examined the effects on the microstructure and electrical properties of Na-β″-Al₂O₃ by doping with TiO₂ [75]. The mechanical properties of the obtained samples were enhanced significantly with the increase of the doping amount, and the mechanical properties increased to more than 280 MPa when the TiO₂ content was higher than 1%. Xu et al. studied the influences of the Nb₂O₅ level to microstructure, mechanical performance and ionic conductivity of the Na-β″-Al₂O₃ [76]. They found that when the doping amount was 1 wt%, the Na-β″-Al₂O₃ phase content was 96.82%, the relative density could reach 98.93%. At the same time, the bending strength was also improved to 295 MPa. In addition to single ion doping, co-doping of multiple ions is also a common modification method. Yang et al. added a certain amount of TiO₂ and ZrO₂ when synthesizing Na-β″-Al₂O₃, in which ZrO₂-doped sample usually has high mechanical properties but poor electrical conductivity, while the TiO₂-doped sample usually high ionic conductivity [77]. Introducing Ti⁴⁺ and Zr⁴⁺ into Na-β″-Al₂O₃ could enhance ionic conductivity and mechanical performance. The synthesized sample with the addition of TiO₂ and ZrO₂ has improved bending properties up to 196.3 MPa with ionic conductivity of 0.2 S/cm at 350 ℃. Moreover, dopants such as Li₂O [74], MnO₂ [78], MgO [79], NiO [80], CaO [49], ZrO₂ [81] and Y₂O₃ [68] have been shown to be effective in increasing the Na-β″-Al₂O₃ phase content as well as improving the mechanical properties.

  4.2 Modification of Nasicon

  The conduction mechanism of Na⁺ within the grains and at the boundaries of Nasicon SSEs is different, therefore, efforts should be directed towards improving both the grain conduction and the grain boundary conduction to achieve overall enhancement in ionic conductivity.

  4.2.1   Improve bulk conductivity

  According to the Arrhenius equation, the ionic conductivity of the SSEs is proportional to n·exp(-E_a/K_BT), where n_c is the concentration of Na-ions and E_a is the activation energy. It is evident that achieving higher ionic conductivity for a given temperature requires a low E_a and a high content (n_c) of Na-ions. The activation energy E_a is associated with the size of bottleneck through which the Na-ions must pass during transport. The size of this bottleneck determines the ease of Na-ion transportation. Consequently, the regulation of Na-ion concentration within the cell and the adjustment of the size of the three-dimensional transport channels for Na-ions represent the two primary approaches for enhancing ionic conductivity.

  Nasicon type SSEs possess an open structure, and their ionic conductivity can be improved by optimizing their structure through element doping or chemical substitution. Early work mainly focused on the NaZr₂(PO₄)₃ system, after the replacement, the ionic conductivity was improved, but it was lower than Na₃Zr₂Si₂PO₁₂ [82-85]. Therefore, the current research is mainly to replace Na₃Zr₂Si₂PO₁₂. Based on the charge balance principle, the introduction of low-valence substitution in Nasicon-type SSEs leads to an increase in the concentration of Na⁺ within the unit cell, which helps compensate for the positive charges resulting from the substitution. This can potentially enhance the ionic conductivity of the electrolyte. However, it is observed that the ionic conductivity does not always increase linearly with increasing doping amounts of low-valence ions. Experimental finds show that here exists an optimal doping level beyond which further increase in the dopant concentration actually leads to a decrease in the ionic conductivity. Theoretical calculations offer insights into this phenomenon by considering the maximum number of Na-ion occupancies within the unit cell, which is typically 4. The presence of sodium vacancies in the crystal structure is crucial for facilitating the movement of Na-ions. Therefore, when all the sodium vacancies are completely occupied, the available sites for Na-ions to jump between become limited. This saturation of sodium vacancies hinders the mobility of Na-ions and consequently results in a decrease in the overall ionic conductivity, even with higher dopant concentrations. After studying the composition of many electrolytes, the optimal concentration of Na per formula unit is around 3.3 mol [86]. And this conclusion was verified in subsequent reports. Ma et al. prepared Na_3+xSc_xZr_2-xSi₂PO₁₂ by a solution-assisted solid-phase method, which provides a wide choice of starting materials and requires simple equipment suitable for large-scale preparation, and the relative densities of the samples in this work prepared by this method can reach 95% [87]. Results showed that the substitution will not only increase the concentration of Na-ions and decreases the phase transition temperature. The optimal ionic conductivity of the samples achieved 4.0 × 10⁻³ S/cm. Consequently, the introduction of low-valence ions such as Mg²⁺, Zn²⁺, Al³⁺, and others has been demonstrated to be effective in increasing the ionic conductivity of these materials [88,89]. Nevertheless, it should be noted that certain elements with variable valence, such as Co, Fe, Cr, among others, are not suitable for doping in this particular system. The introduction of these elements may result in electrolyte instability and self-discharge of the battery [90].

  In recent years, co-doping strategy have been widely investigated. By utilizing the synergistic effects of two different dopant ions, the co-doping approach can effectively further enhance conductivity. Typically, the two ions involved in co-doping contribute to the enhancement of ionic conductivity through different mechanisms. One of the ions usually adopts a lower valence to increase the Na-ion concentration within the unit cell. The other ion, with a larger radius, can widen the Na-ion transport channel or facilitate sintering to reduce the resistance at the boundaries. In addition, ions that decrease the temperature of phase transition can be selected as co-doping ions. He et al. a study on the Mg²⁺/F⁻ co-doped Na₃Zr₂Si₂PO₁₂ strategy, aimed at enhancing the ionic conductivity of Na₃Zr₂Si₂PO₁₂ SSEs. In this strategy, Mg²⁺ and Zr⁴⁺ ions were chosen as co-dopants due to their identical radii (0.72 Å) [91]. As a result, Mg²⁺ doping increases the Na-ion concentration within the lattice without causing structural distortion. The introduction of F⁻ ions facilitates the enhancement of sintering activity in the ceramic material, leading to a gradual reduction of grain boundaries as well as a significant growth in material densities. Co-doping of Mg²⁺ and F⁻ ions resulted in an optimal ionic conductivity of 2.21 × 10⁻³ S/cm for the synthesized samples. Another ion commonly studied for co-doping is Sc³⁺, which possesses a similar ionic radius to Zr⁴⁺ (r(Zr⁴⁺) = 0.72 Å, r(Sc³⁺) = 0.745 Å) but a lower valence state. The incorporation of Sc³⁺ has been shown to effectively lower the phase transition temperature, making it a favorable co-doping ion. Omar et al. synthesized Sc/Yb co-doped Na₃Zr₂Si₂PO₁₂ SSEs [92]. The introduction of Sc³⁺ increases the concentration of Na-ions within the cell, while the introduction of Yb ions cause lattice expansion, which can broaden the size of Na-ion transport channels. In addition to co-substitution at the Zr site, simultaneous substitution at the Zr site and P is also effective in improving the ionic conductivity. Yang et al. simultaneously replaced part of the P with Si and part of the Zr with Zn to obtain Na_3.4Zr_1.9Zn_0.1Si_2.2P_0.8O₁₂ with an ionic conductivity of 5.27 × 10⁻³ S/cm [93]. The substitution of Zr with Zn in the SSEs expands the lattice due to the slightly larger radius of Zn compared to Zr. This expansion results in an increased bottleneck size of the Na1-Na-Na1 transport channel, which facilitates the transportation of Na-ions. Furthermore, the lower valence state of Zn ions contributes to an increased carrier concentration, leading to a ratio of Na-ion concentration to vacancy concentration of 3.4:0.6 within the system. This higher carrier concentration contributes to a higher ionic conductivity. The effectiveness of simultaneous Zr and P substitution has also been confirmed in other studies, such as the co-doping of P and Mg ions.

  4.2.2   Decrease grain boundary resistance

  The transport of ions at grain boundaries poses challenges due to the interruption of ion transport channels, leading to difficulties in ion diffusion. Consequently, the energy barrier for ion transport at grain boundaries is higher compared to the bulk material, resulting in lower ionic conductivity at the grain boundaries. The density and purity of the materials play crucial roles in influencing the ionic conductivity at grain boundaries. A lot of works has been devoted to the synthesis of pure phase Na₃Zr₂Si₂PO₁₂. Naqash et al. synthesized Nasicon compositions with sodium excess, i.e., Na₃Zr₂Si₂PO₁₂ + xNa₂O (0 ≤ x ≤ 0.2), demonstrated that the excess Na effectively compensated for the volatilized Na element at high temperature and effectively suppressed the formation of impurity of impurity phase [69]. The sample with x = 0.2 showed the least sodium deficiency and the highest ionic conductivity (1.6 × 10⁻³ S/cm at 25 ℃).

  In addition, sol-gel method, hydrothermal method, solution-assisted solid-phase reaction (SASSR), spray freeze/freeze drying and low-temperature rapid microwave sintering method are also widely used to synthesize Nasicon materials with high densities [94-99]. Sol-gel method can reduce the sintering temperature required for densification effectively, thus avoiding the formation of impurity phases. And the samples prepared by sol-gel method is more uniform than solid-phase method, therefore, this method is widely used for the synthesis of Nasicon electrolytes. Zhang et al. demonstrated that the temperature could be reduced by sol-gel method, results showed that the main phase of Nasicon prepared by sol-gel method was formed when sintered at 850 ℃, and the sample obtained by sintering at 950 ℃ was well crystalline [70]. In addition, many novel preparation methods have also been used to prepare Nasicon materials. Ma et al. prepared Na_3+xSc_xZr_2-xSi₂PO₁₂ by a solution-assisted solid-phase method, which provides a wide choice of starting materials and requires simple equipment suitable for large-scale preparation, and the relative densities of the samples in this work prepared by this method can reach 95% [87]. The optimal ionic conductivity of the samples achieved 4.0 × 10⁻³ S/cm.

  The liquid phase sintering method is widely recognized as an effective approach to improve the density of materials. Oh et al. sintered NZSP in the Na₂SiO₃ with a melting point of 1088 ℃, the formation of the melting pool of Na₂SiO₃ can improve the grain and grain boundary conductivity, and the highest ionic conductivity is 1.45 × 10⁻³ S/cm (Figs. 11a and b) [100]. Shao et al. employed NaF as a co-sintering agent in the sintering process to increase the density of Nasicon materials [101]. SEM images revealed that the ceramic sheet consisted of angular grains and "binder-like" amorphous glassy materials after NaF was added. This amorphous phase altered the grain boundaries, leading to an improvement in density and enhancing the ionic conductivity of Na₃Zr₂Si₂PO₁₂ from 4.5 × 10⁻⁴ S/cm to 1.7 × 10⁻³ S/cm (Fig. 11c).

  Figure 11

  

  Figure 11.  (a) Schematic diagram of using Na₂Si₂O₃ as NZSP sintering additive and (b) ionic conductivity and relative density of these samples. Reprinted with permission [100]. Copyright 2019, American Chemical Society. (c) Ionic conductivity and E_a of samples with different levels of NaF. Reprinted with permission [101]. Copyright 2019, Elsevier.

  DownLoad: Full-Size Img PowerPoint

  4.3 Modification of layered oxides

  As well as other types of SSEs, elemental doping is the main strategy commonly used to improve the ionic conductivity of layered oxide SSEs. Li et al. investigated Ga³⁺ doping on the structure and properties of Na₂Zn₂TeO₆ (NZTO) [102]. They found that a small amount of Ga³⁺ substitutes for Zn²⁺ would introduce more Na⁺ vacancies in the interlayer gaps (Fig. 12a), which greatly reduce strong Na⁺-Na⁺ coulomb interactions and exhibit a superionic conductivity of 1.1 × 10⁻³ S/cm at r.t. Ga³⁺ doping improves the ionic conductivity significantly for the following reasons: (1) Ga³⁺ doping with smaller radius induces longer Na-O bonds and larger ion migration channels, which also reduces the attraction between Na⁺ and O²⁻; (2) Sufficient Na-site vacancies in NZTO-G0.1 not only increase the concentration of current carriers, but also reduce the migration energy of Na-ions. And the similar radii of Ga³⁺ and Zn²⁺ will not lead to distortion of the crystal structure. In 2018, Wu et al. analyzed the conduction of Na⁺ in Ga³⁺-doped Na₂Zn₂TeO₆ from the point of view defect chemistry for the first time [103]. The results indicate that the introduction of Ga³⁺ increases the concentration and mobility of mobile sodium ions, ultimately leading to an increase in the grain bulk conductivity. Furthermore, the existence of ZnGa defects with effective positive charges decreases the charge density in the space-charge layer, which reduces the Schottky barrier height, ultimately resulting in an increase in grain boundary conductivity (Figs. 12b and c). Ion doping with larger radius can also effectively improve the ionic conductivity. Deng et al. explored the effect of Ca²⁺ doping on the structure and properties of Na₂Zn₂TeO₆ [104]. The refinement results show that the unit cell parameters a and c and the distance between the two layers increase with the increase of Ca²⁺ doping content. At the same time, the ionic conductivity of Na₂Zn₂TeO₆ at r.t. is increased to 7.5 × 10⁻⁴ S/cm with the activation energy of 0.225 eV.

  Figure 12

  

  Figure 12.  (a) Transformation of partial crystal structure before and after Ga substitution in NZTO. Reprinted with permission [102]. Copyright 2018, Elsevier. (b) Total ionic conductivities of Na_1.95Zn_1.95Ga_0.05TeO₆ and Na₂Zn₂TeO₆ and (c) the σ_bulk and σ_gb of two samples at different temperatures Reprinted with permission [103]. Copyright 2018, Elsevier.

  DownLoad: Full-Size Img PowerPoint

  In addition to Na₂M₂TeO₆ (M = Ni, Co, Zn, Mg), Smaha et al. reported another layered Na ion conductor, Na_3-xSn_2-xSb_xNaO₆ (with a space group of C2/c) in 2015 [105]. This finding expanded the range of layered materials capable of facilitating sodium ion conduction. When x = 0.8, the maximum conductivity of Na_2.2Sn_1.2Sb_0.8NaO₆ at 500 ℃ is 1.43 × 10⁻³ S/cm and the activation energy is 0.63 eV. The authors believed that replacing Sn⁴⁺ with Sb⁵⁺ would create Na vacancies and thus promote higher Na-ion mobility.

  4.4 Modification of anti-perovskites

  Modulation of local structural features by elemental doping or creating vacancies is a common way to improve ionic conductivity. In 2015, Wang et al. synthesized Na₃OCl and Na₃OBr with anti-perovskite structure, and found that the ionic conductivity of Na-rich anti-perovskite can be improved by changing the element species at the halogen site and creating Na-ion vacancies (Fig. 13a) [106]. The Na₃OBr_0.6I_0.4 obtained by introducing I⁻ into Na₃OBr possesses the highest ionic conductivity of 0.43 mS/cm at 200 ℃. The sodium ion conductivity increases sequentially from Na₃OCl to Na₃OBr to Na₃OBr_0.6I_0.4, which proves that the mismatch effect caused by the incorporation of larger halide ions at the A site can improve the ionic conductivity. Then high-valence Sr³⁺ are doped into Na sites to introduce Na vacancies, and the obtained Na_2.9Sr_0.05OBr_0.6I_0.4 has an ionic conductivity of 19 mS/cm at 200 ℃. This work confirms that unequal alkaline earth mental ion doping can enhance the ionic conductivity of anti-perovskite SSEs effectively. On the other hand, the free transport volume in the anti-perovskite structure can also be changed by modifying the anions on the A-site. And the free space in the anti-perovskite structure depends on the size of the anion once the rigid backbones are fixed. Larger anions will leave a smaller free space in the framework. Therefore, the A site preferentially selects anions with smaller radii. But if the anion at A-site is too small to fill the cavity, it will cause the anti-perovskite structure to collapse or distorted, leaving little free space for ion migration. This work showed the benefit of cation mixing in anti-perovskite electrolytes with the Na₃OBr and Na₃OBr_0.6I_0.4 samples, substituting the smaller A anion with a larger anion may maintain the general formula of free space to provide a stable structure, thus, the activation energy of the mixed cation sample is much lower than that of a single cation (0.63 eV for Na₃OBr_0.6I_0.4 and 0.76 eV for Na₃OBr), as shown in Fig. 13b. To further improve the conductivity, researchers are focusing their attention to anion groups. Substituting X with a cluster (i.e., Na₃S(BCl₄) and Na₃O(BF₄)) was predicted to be an effective way to achieve higher ionic conductivity [107]. In 2019, Sun et al. synthesized Na₃OBH₄ with an anti-perovskite structure by solid-phase method (Fig. 13c) [108]. Its ionic conductivity is 4.4 × 10⁻³ S/cm, which is four orders of magnitude higher than the existing Na₃OX (X = Cl, Br, I) (Fig. 13d). The synthesis of Na₃OBH₄ with higher ionic conductivity may be an important advance in the development of Na-rich anti-perovskite electrolytes. Gao et al. synthesized Na₃ONO₂ by a low temperature solid-state reaction, ESI measurements showed that the highest ionic conductivity in Na₃ONO₂ (0.37 mS/cm, E_a: 0.385 eV) around 485 K [109]. And they used Neutron powder diffraction refinements and DFT calculations reveal the mechanism of the conductivity enhancement. Above 485 K, the NO₂⁻ rotation is more intensified than that at lower temperature, which can significantly facilitate Na⁺ ion migration via Na–O interactions.

  Figure 13

  

  Figure 13.  (a) Crystal structure of Na₃OX (X = Cl, Br, I) and (b) Arrhenius plots for Na₃OCl, Na₃OBr, Na₃OBr_0.6I_0.4 and Na_2.9Sr_0.05OBr_0.6I_0.4 samples. Reprinted with permission [106]. Copyright 2015, Elsevier. (c) Crystal structure of Na₃OBD₄ and (d) Nyquist plots of hot-pressed Na₃OBH₄ pellet. Reprinted with permission [108]. Copyright 2019, American Chemical Society.

  DownLoad: Full-Size Img PowerPoint

  In addition, a large amount of theoretical computational work has driven the development of Na-rich anti-perovskite materials. Wan et al. studied the effect of substitutional defects on the Na migration energy through NEB computations and AIMD simulations [110]. They investigated the effects of divalent and trivalent dopants (Mg²⁺, Ca²⁺, Sr²⁺, Ba²⁺, Al³⁺, and Ga³⁺) on the ionic transport and conductivity in Na₃OCl, the results showed that Ba is the most stable substitutional defect among the ones studied. It was also found that Ca²⁺ may be the most potential doped ion because it leads to the lowest dopant-vacancy binding energy. Goldmann et al. investigated the cation doping mechanisms and ionic conductivity in the anti-perovskite Na₃OCl in an atomic-scale perspective [111]. They found that the most favorable aliovalent cation dopants are Mg²⁺, Ca²⁺, Al³⁺ and Ga³⁺ with Na-vacancy charge compensation; an increase in Na-vacancy concentration would promote Na-ion conductivity. The highest conductivity is found for the Mg-doped system of the order of 10⁻⁵ S/cm at 500 K, but lower conductivities are predicted for the trivalent Al and Ga dopants.

  Currently, there is a limited variety of Na-rich anti-perovskite SSEs available, and their ionic conductivity at r.t. does not meet the requirements for practical applications. Correspondingly, the development of all-solid-state SIBs utilizing these electrolytes is still limited. Therefore, it is crucial to actively explore new types of anti-perovskite materials in order to expand the range of Na-ion SSEs options. Additionally, there is a need to investigate alternative preparation methods for anti-perovskite electrolytes. For instance, in the case of Na₃OBH₄ mentioned earlier, the author demonstrated that hot-pressing sintering method yielded significantly higher ionic conductivity (4.4 × 10⁻³ S/cm) at r.t. compared to samples prepared using the traditional cold-pressed sintering method [108]. This also shows that the preparation process of the sample has a great influence on its performance. Furthermore, anti-perovskite SSEs should be studied from multiple perspectives. Previous studies on F ion conductors and O ion conductors have shown that the ability of ion conductivity can be regulated by controlling the pressure, while the effect of pressure on the ionic conductivity of Na-rich anti-perovskite materials has not been investigated [112,113]. In 2016, Wang et al. invested the structure stability of anti-perovskite Na₃OBr and Na₄OI₂ SSEs under high pressure by in-situ SR-XRD [114]. The results demonstrate that both the cubic Na₃OBr structure and tetragonal Na₄OI₂ structure remain stable under high pressure conditions up to 23 GPa. However, it should be noted that the variation of ionic conductivity of anti-perovskites as a function of pressure has not been reported.

  4.5 Modification of sulfide SSEs

  4.5.1   Improve ionic conductivity

  Na₃PnS₄ (Pn = P, Sb): Element doping is a widely employed approach to enhance the ionic conductivity of Na₃PS₄ materials. The effect of doping ions on the ionic conductivity can be attributed to several factors, including following aspects. On the one hand, low-valence ions such as Sn⁴⁺, Ge⁴⁺, Ti⁴⁺, replace P⁵⁺ ions, in order to maintain the overall electrical neutrality, excess Na-ions are often introduced, which is beneficial to reduce the activation energy and improve the conductivity of Na-ions. Replacing S²⁻ with low-valent anions such as F⁻, Cl⁻, Br⁻ can introduce more Na-ion vacancies, increase the effective hopping probability of Na-ions between adjacent sites, and thus improve the ionic conductivity. On the other hand, ionic doping with larger radius (such as As⁵⁺) expands the lattice and elongates the Na-S bonds, thereby increasing the ionic conductivity. We summarize the modification methods of Na₃PnS₄ (Pn = P, Sb) solid electrolyte according to the different doping sites.

  In 2014, Tanibata et al. reported for the first time (100-x) Na₃PS₄·xNa₄SiS₄ glass-ceramic, in which the glass-ceramic containing 6% Na₄SiS₄ had an ionic conductivity of 7.4 × 10⁻⁴ S/cm at r.t. [115]. And it was found that the introduction of Si ions caused excess Na ions to occupy Na2 sites, which is also the reason for the improvement of its ionic conductivity. In 2019, Hayashi et al. reported the W-doped Na₃SbS₄ solid electrolyte Na_2.88Sb_0.88W_0.12S₄, the ionic conductivity can reach 3.2 × 10⁻² S/cm at r.t. [116]. Because of the ionic radius of Sb is larger than that of P, and replacing P with Sb can effectively expand the ion transport channel and achieve higher ionic conductivity (Figs. 14a and b) [33,117,118]. Yu et al. replaced a part of P ions in Na₃PS₄ with As³⁻ ions, and synthesized a solid electrolyte composed of Na₃P_0.62As_0.38S₄, with a high ionic conduction of 1.46 × 10⁻³ S/cm at r.t. [119]. Weng et al. synthesized WS₂ -doped Na_3-xSb_1-xW_xS₄ and WO₂-doped Na_3-xSb_1-xWxS_4–2xO_2x (x = 0.025, 0.05, 0.075, 0.1) solid electrolytes, The room temperature ionic conductivity of the sample composed of Na_2.95Sb_0.95W_0.05S₄ and Na_2.95Sb_0.95W_0.05S_3.9O_0.1 can achieve 10.37 and 8.49 mS/cm [120].

  Figure 14

  

  Figure 14.  (a) Raman spectra and (b) XRD patterns of Pristine Na₃SbS₄·9H₂O, as-prepared Na₃SbS₄, air-exposed Na₃SbS₄ and reheated air-exposed Na₃SbS₄. Reprinted with permission [33]. Copyright 2016, Wiley Online Library.

  DownLoad: Full-Size Img PowerPoint

  Equivalent doping has been shown to be effective in enhancing ionic conductivity. In 2015, Zhang et al. replaced S with Se lead to Na₃PSe₄ with an ionic conductivity of 1.16 × 10⁻³ S/cm, and reduced the Ea of the electrolyte to 0.21 eV [121]. The increase of ionic conductivity could result from two factors, firstly, the substitution of Se^2- with larger ionic radius for S²⁻ can expand the lattice; Second, Se²⁻ with greater polarizability will weaken the binding energy between Na ions and S^2- framework, thereby improving the mobility of Na ions. Ceder et al. combined theoretical calculations and experiments to study the Na⁺ conduction mechanism in cubic Na₃PSe₄ [122]. The study found that when Na₃PSe₄ has defects, it is conducive to the rapid conduction of Na⁺, indicating that the mechanism of ion transport in Na₃PSe₄ is vacancy transport mechanism.

  Guided by first-principles calculations, Chu et al. synthesized a series of novel Cl-doped tetragonal Na₃PS₄ (t-Na_3-xPS_4-xCl_x) samples by solid phase reactions [123]. The various reaction materials were ground, calcined, reground and re-pelletized. These pellets were then treated by spark plasma sintering. The sample with the composition Na_2.9375PS_3.9375Cl_0.062 exhibited the highest ionic conductivity of 1.0 × 10⁻³ S/cm at r.t., and the assembled full cell displayed reversible charging and discharging at a rate of 0.1 C. The researchers demonstrated, using density functional theory calculations, that the exceptional performance of the Cl-doped Na₃PS₄ electrolyte in this battery can be attributed not only to the high conductivity of Na-ions in the SSEs but also to the formation of an electronically insulating and ionically conductive interface layer at the electrode-electrolyte interface due to the presence of Cl⁻.

  In 2018, Moon et al. synthesized Ca-doped Na₃PS₄ by the solid-phase method, found that Na⁺ in t-Na₃PS₄ was replaced by Ca²⁺, forming a cubic phase of Na_3–2xCa_xPS₄, and at the same time generating Na⁺ vacancies, which greatly improved the ionic conductivity. When x = 0.135, it can reach 1 × 10⁻³ S/cm at 25 ℃ [124].

  Na₁₁Sn₂PnSn₁₂ (Pn = P, Sb; Sn = S, Se): In 2018, Duchardt et al. replaced the S atom in Na₁₁Sn₂PS₁₂ with Se to obtain an electrolyte with a composition of Na_11.1Sn_2.1P_0.9Se₁₂ [125]. Its performance test found that the ionic conductivity was almost unchanged before and after Se substitution, but the activation energy decreased significantly after substitution (Na₁₁Sn₂PS₁₂: 0.39 eV and Na_11.1Sn_2.1P_0.9Se₁₂: 0.30 eV). Structural analysis shows that the substitution of Se atoms with larger radii for S atoms can broaden the ion transport channel, and the substitution of Se for S makes the lattice softer, lowering the migration barrier for ion transport in the lattice, thereby lowering the activation energy. In 2020, Jia et al. reported a halogen ion-doped Na_3.67[Sn_0.67Si_0.33]_0.67P_0.33S₄ electrolyte, in which Na_3.57[Sn_0.67Si_0.33]_0.67P_0.33S_3.9I_0.1 was obtained by I⁻ doping [126]. The ionic conductivity of the electrolyte can reach 1.08 × 10⁻³ S/cm at r.t. Structural studies show that doping I with a larger radius than S can expand the lattice and accelerate the transport of Na ions. In 2021, Liu et al. synthesized a new type of solid electrolyte Na₁₀SnSb₂S₁₂, which showed a room temperature ionic conductivity of 0.52 mS/cm and a low activation energy of 0.23 eV [14].

  4.5.2   Chemical and electrochemical stability

  Despite the achievement of practical RT ionic conductivity in sulfide SSEs, their sensitivity to moisture and air remains a significant challenge that hinders their widespread application. The instability of sulfide SSEs in the presence of moisture and air can be explained by the "theory of hard and soft acids and bases". According to this theory, there is a preference for the interaction between hard acids and hard bases, as well as soft acids and soft bases. In the case of sulfide SSEs, S²⁻ acts as a soft acid and tends to bind with soft bases to form stable structures. On the other hand, P⁵⁺ acts as a hard acid and tends to bind with hard bases such as O²⁻. Consequently, most P-based sulfide SSEs exhibit lower stability in the presence of air. However, Na₃SbS₄, which contains Sb⁵⁺ as a soft acid, exhibits improved chemical stability. Wang et al. investigated the air stability of Na₃SbS₄ and observed the formation of Na₃SbS₄·9H₂O when exposed to air, as indicated by Raman and XRD results (Figs. 14c and d). Nevertheless, after sintering at 150 ℃ for 1 h, Na₃SbS₄ returned to its original form, demonstrating excellent chemical stability in air [33]. In addition to this, partial substitution of P in Na₃SbS₄ with As and Sn also had a significant effect on its chemical stability [119,127].

  The electrochemical stability of sulfide SSEs is typically assessed using cyclic voltammetry, where inert electrodes are used as blocking electrodes and Na metal is used as the counter electrode. The electrochemical window, which indicates the range of stable electrochemical potentials, is an important parameter for evaluating the stability of solid electrolytes. For instance, Na₃PS₄ has been reported to have an electrochemical window of stability up to 5 V (vs. Na⁺/Na) as measured by cyclic voltammetry. However, experimental studies have shown that Na₃PS₄ is not stable when in contact with the Na negative electrode. Yue et al. conducted tests on the Na-Sn-C|Na₃PS₄|Na₃PS₄Na₂S-C system and observed the oxidation of Na₃PS₄ at approximately 2.0 V during the first charge. This finding indicates that the actual electrochemical stability window for sulfide electrolytes may be less favorable than the results obtained from cyclic voltammetry tests. Therefore, it is crucial to consider the actual electrochemical behavior and stability of sulfide electrolytes under specific operational conditions, particularly in contact with the negative electrode, to ensure their reliable performance in practical applications.

  Sodium metal possesses a high energy density, but its reactivity poses challenges for achieving stable interfaces with electrolyte materials. The stability of the electrolyte in relation to sodium metal plays a crucial role in determining the formation of a stable interface. When the electrolyte is thermodynamically stable relative to sodium metal, a stable interface is formed upon contact. Conversely, if the electrolyte is thermodynamically unstable, rapid chemical reactions occur, leading to an increase in interfacial impedance, a decrease in ion diffusion rate, and ultimately, the failure of the cell. To enhance the electrochemical stability of sulfide SSEs, it is essential to design electrolyte materials and interfaces in a rational manner. One approach involves the doping of Cl⁻ ions into Na₃PS₄, as demonstrated by Chu et al. This Cl⁻ doping can passivate or stabilize the electrode-electrolyte interface, thereby improving the electrochemical stability of the electrolyte [123]. By carefully modifying the composition and structure of the electrolyte, it is possible to achieve improved stability and performance in sulfide solid electrolyte-based systems. Such strategies for enhancing electrochemical stability are crucial for enabling the practical application of sulfide solid electrolytes, ensuring their long-term performance and reliability in advanced energy storage devices.

  4.6 Modification of halide SSEs

  In 1995, Mathias et al. reported Na₃MX₆ (X = Cl, Br) SSEs for the first time [128]. Na₃SmBr₆ and Na₃GdCl₆ showed an ionic conductivity of only 10⁻⁵ S/cm, while Na₃YbBr₆ and Na₃YbCl₆ showed an even lower ionic conductivity of 10⁻⁶ S/cm. Such low ionic conductivity renders them unsuitable for applications in all-solid-state SIBs. In recent years, Li-ion halides (Li₃YCl₆, Li₃YBr₆, Li₂ZrCl₆) have been reported as promising Li-ion conductors [129-133]. The outstanding properties of high ionic conductivity at r.t. and excellent chemical and electrochemical stability have garnered significant interest in the study of lithium-ion conductors. While Li₃YCl₆ and Li₃YBr₆ have been extensively investigated, sodium halides such as Na₃YCl₆ and Na₃YBr₆ have received comparatively less attention. This is attributed to the fully occupied 2d and 4e sites in Na₃YCl₆ (NYC) (space group: P21/n), which restrict ion hopping and consequently result in lower ionic conductivity. In 2021, Wu et al. addressed this limitation by introducing high-valence Zr⁴⁺ doping to create sodium vacancies, thereby improving the ionic conductivity of these materials [134]. They synthesized a series of Zr⁴⁺-doped Na_3-xY_1-xZr_xCl₆ samples, remarkably, the introduction of Zr⁴⁺ doping led to a substantial improvement in r.t. ionic conductivity. When x = 0.75, the doped sample exhibited an impressive ionic conductivity of 6.6 × 10⁻⁵ S/cm, which is three orders of magnitude higher compared to the undoped sample.

  Recently, Zhang et al. synthesized a Na₃YbCl₆-based halide electrolyte using high-energy ball milling, and the optimized NaYbZrCl_0.75 sample has a high ionic conductivity of 6.6 × 10⁻⁵ S/cm and a wide electrochemical window of 4.1 V [135]. Xu et al. synthesized a series of F-doped NaAlCl_4−xF_x (x = 0, 0.1, 0.3, 0.5, and 0.7) and borohydride (BH₄)-doped NaAlCl_4−x(BH₄)_x (x = 0.1 and 0.3) halide solid electrolytes using a straightforward ball-milling method, and these samples maintained the same orthorhombic structure as NaAlCl₄ [136]. The Na⁺ conductivity of the appropriately doped NaAlCl_4−xF_x SEs was found to be higher than that of NaAlCl₄. Specifically, NaAlCl_3.3F_0.7 exhibited significantly enhanced Na⁺ conductivity (3.56 × 10⁻⁵ S/cm at 30 ℃) compared to NaAlCl₄ (4.14 × 10⁻⁶ S/cm at 30 ℃). Fu et al. developed a new class of halide heterogeneous structure electrolytes by exploiting the structural differences between high and low coordination halide frameworks [137]. Halide heterogeneous structure electrolytes containing UCl₃-type high coordination frameworks and amorphous low coordination frameworks achieved a Na⁺ conductivity of 2.7 mS/cm at room temperature. Recently, Hu et al. successfully designed and synthesized a novel amorphous NaTaCl₆ halide solid-state electrolyte that exhibits an ultra-high ionic conductivity of 4 × 10⁻³ S/cm at room temperature [138]. The excellent ionic conductivity achieved can be attributed to the formation of a reconfigured amorphous poly(TaCl₆) octahedral network triggered by high-energy mechano-chemical reactions, which efficiently drives sodium ions into the open amorphous halide network, resulting in weaker Na-Cl interactions.

  4.7 Modification of complex hydrides

  Among the reported composite hydride solid-state electrolytes, Na₂B₁₂H₁₂ in its disordered phase has been found to exhibit the highest ionic conductivity at elevated temperatures. Consequently, recent investigations in complex hydrides have aimed to reduce the phase transition temperature of Na₂B₁₂H₁₂. In 2015, Tang et al. reported that the ionic modifications can effectively lower the phase transition temperature [139]. They replaced B in Na₂B₁₂H₁₂ with C atom to form NaCB₁₁H₁₂ (Figs. 15a and b), and found that the material had a phase transition of 308 K (600 K for NaCB₁₁H₁₂) and an ionic conductivity of 0.12 S/cm at 383 K. Battaglia et al. mixed two precursors of Na₂B₁₂H₁₂ and Na₂B₁₀H₁₀ in a ratio of 1:1 to obtain Na₂(B₁₂H₁₂)_0.5(B₁₀H₁₀)_0.5 electrolyte, and the ionic conductivity of the material reached 0.9 mS/cm at 20 ℃ [140]. XRD and DSC characterizations showed that this compound has a face-centered cubic structure at r.t. (the same as that of Na₂B₁₂H₁₂ at 180 ℃), which also confirmed that anion doping can effectively reduce the phase transition temperature.

  Figure 15

  

  Figure 15.  (a) Structure diagram of B₁₂H₁₂²⁻ and CB₁₁H₁₂⁻. (b) The ionic conductivities of Li⁺ and Na⁺ species in LiCB₁₁H₁₂ and NaCB₁₁H₁₂ were measured as a function of inverse temperature. Reprinted with permission [139]. Copyright 2015, RSC publishing. (c) The ionic conductivity of pristine and ball-milled Na₂B₁₂H₁₂ was compared after QENS measurements. Reprinted with permission [141]. Copyright 2016, Elsevier. (d) A description of the synthetic route of NaBH₄@Na₂B₁₂H₁₂. Reprinted with permission [142]. Copyright 2022, American Chemical Society.

  DownLoad: Full-Size Img PowerPoint

  The ball-milling method has demonstrated effectiveness in reducing the phase transition temperature of Na₂NH₂B₁₂H₁₂. Tang et al. observed that this method, through the reduction of crystallite size and induction of disordering effects, leads to the stabilization of a high-temperature-like superionic-conducting phase at r.t. [141]. The effectiveness of the high-temperature disordered phase was further substantiated through XRD and other characterization techniques. Performance testing also revealed a significant improvement in the ionic conductivity of the ball-milled samples (Fig. 15c). Furthermore, the synthesis of mixed anionic hydrides has demonstrated its efficacy in reducing the phase transition temperature.

  In 2022, Luo et al. successfully synthesized a novel solid-state electrolyte, NaBH₄@Na₂B₁₂H₁₂, featuring a distinct core-shell structure (Fig. 15d) [142]. The sample exhibited a high ionic conductivity of 10⁻⁴ S/cm at 115 ℃. The results showed that the core-shell structure may facilitate NaBH₄/Na₂B₁₂H₁₂ interfacial sites for sodium ion hopping. This, in turn, improves the ionic conductivity, providing a novel approach for enhancing the ionic conductivity of complex hydride SSEs.

  The primary objective of developing SSEs is to successfully apply them to solid-state batteries, achieving superior safety, high energy/power density, and extended cycling performance. In this section, we will provide a comprehensive overview and analysis of the current state of solid sodium batteries based on the aforementioned SSE.

  5. Solid-state batteries applications for Na-ion SSEs

  The primary objective of developing SSEs is to successfully apply them to SIBs, achieving superior safety, high energy/power density, and extended cycling performance. In this section, we will provide a comprehensive overview and analysis of the current state of SIBs based on the aforementioned SSEs.

  5.1 Solid-state SIBs based on oxide SSEs

  Na-β″-Al₂O₃, renowned for its remarkable ionic conductivity, has emerged as a pioneering SSE for sodium batteries [143]. ZEBRA batteries and high-temperature Na-S batteries were discovered based on it. Its operating temperature is generally around 300 ℃, which not only increases the operational and maintenance costs of the battery but also increases safety hazards [144-146]. Due to the low room temperature ionic conductivity of Na-β-Al₂O₃, large grain boundary impedance, and poor sodium affinity, it cannot work at lower temperatures [147]. To reduce the operating temperature of Na-β-Al₂O₃-based batteries, it is necessary to enhance the room temperature ionic conductivity of Na-β-Al₂O₃ electrolytes, and to improve the solid-solid contact interface between the electrolyte and electrode materials. Recently, there has been emerging research on solid-state SIBs utilizing Na-β″-Al₂O₃ electrolytes, which exhibit promising performance even at ambient temperature. Zhao et al. used a strip casting and sintering process to prepare Na-β″-Al₂O₃ thin films with a thickness of about 100 µm [148]. The battery used gel NaTi₂(PO₄)₃ as the cathode and Na metal the anode. The reversible discharge capacity remained at 100 mAh/g after 50 cycles (Fig. 16a). Na-β″-Al₂O₃ is highly stable with Na and has a wide electrochemical window. Nevertheless, the inadequate interface contact between the Na metal anode and electrolyte leads to elevated interfacial resistance and non-uniform Na⁺ deposition [13,149]. To address these issues, Wen et al. proposed a dual-functional layer composed of three-dimensional cross-linked carbon fibers and Sn particles, which can be applied to the surface of Na-β″-Al₂O₃. This modification enhances interface wettability and facilitates the uniform deposition of Na⁺ [150]. A battery with Na₃V₂(PO₄)₃ as the cathode showed a capacity retention of 99.7% after 100 cycles at 5 C, indicating that the modified anode interface can adapt to high current density (Fig. 16b). Research has shown that, similar to Li₇La₃Zr₂O₁₂, eliminating hydroxyl groups and carbon contamination on the surface of Na-β″-Al₂O₃ through heat treatment is also crucial for improving the critical current density. Therefore, Battaglia et al. performed heat treatment on Na-β″-Al₂O₃ in an argon atmosphere and found that the interfacial resistance can be less than 10 Ω cm², and the current density can reach 12 mA/cm² (Fig. 17a) [151]. In addition, Liu et al. introduced a minor quantity of ionic liquid to the composite cathode, which can adhere to the surface of the SSE and maintain good interface contact with the SSE (Fig. 17b) [152]. The Na_0.66Ni_0.33Mn_0.67O₂/Na-β″-Al₂O₃/Na battery designed in this way has excellent rate performance and an exceptional cycle life, exceeding 10,000 cycles. Kim et al. even added liquid electrolyte to wet the interface at both sides, resulting battery showed significant advantages in capacity performance and cycling stability compared to liquid-state batteries (Fig. 17c) [145]. In addition, employing advanced techniques such as pulsed laser deposition for in-situ cathode deposition onto the electrolyte surface is a favorable approach to optimize the contact between the cathode and electrolyte [153].

  Figure 16

  

  Figure 16.  (a) Cycling performance of the SIBs based on the Na-β″-Al₂O₃ thin film electrolyte. Reprinted with permission [148]. Copyright 2016, Elsevier. (b) Schematic illustration of Na₃V₂(PO₄)₃/SC-treated-β″-Al₂O₃/Na solid-state batteries and corresponding electrochemical performance. Reprinted with permission [150]. Copyright 2022, Elsevier.

  DownLoad: Full-Size Img PowerPoint

  Figure 17

  

  Figure 17.  (a) XPS spectra and surface composition analysis results of Na-β″-Al₂O₃ surfaces before and after heat treatment. Reprinted with permission [151]. Copyright 2019, Wiley Online Library. (b) Schematic illustration of Na_0.66Ni_0.33Mn_0.67O₂/β″-Al₂O₃/Na solid-state batteries with ionic liquid added in the cathode side and corresponding electrochemical performance. Reprinted with permission [152]. Copyright 2016, American Chemical Society. (c) Schematic illustration of solid-state batteries with liquid electrolyte added in the both sides of the interface between the cathode and anode electrodes and corresponding electrochemical performance. Reprinted with permission [145]. Copyright 2016, Elsevier.

  DownLoad: Full-Size Img PowerPoint

  Analogous to Na-β″-Al₂O₃, Nasicon-type SSEs also have excellent ionic conductivity, but the significant interface resistance arising from solid-solid contact in the all-solid-state SIBs greatly limits its performance [154]. Given their inherent brittleness, ceramic electrolytes often necessitate hot pressing to achieve optimal interface contact with electrode materials. Zhang et al. designed a Na/Na_3.3Zr_1.7La_0.3Si₂PO₁₂/Na₃V₂(PO₄)₃ battery [155]. Nonetheless, significant polarization was observed at room temperature, resulting in a modest reversible capacity of only 85 mAh/g during the initial cycle. Even when the cell was operated at 80 ℃, the capacity dropped sharply after 40 cycles. The unsatisfactory electrochemical performance can be attributed to inadequate contact between the electrode material and SSEs. Despite the advantageous properties of high ionic conductivity and high elastic modulus in Nasicon, the interface contact deteriorated over long-term cycling due to volume changes of the active material, making the cell unable to function properly even at high temperatures [156]. To improve interface compatibility, Zhang et al. also assembled two other types of cells, using a liquid electrolyte (NaPF₆/EC-DMC) and an ionic liquid (PP13FSI) as wetting agents, respectively (Fig. 18a) [155]. The incorporation of a liquid electrolyte partially enhanced the electrochemical performance. However, the specific capacity experienced a substantial decline after 250 cycles. This can be ascribed to the evaporation and/or decomposition of the liquid electrolyte. Additionally, the NVP/IL/SE/Na cell demonstrated the best cycling stability. This can be attributed to the enhanced interface contact between the electrode material and the solid electrolyte, forming a buffer layer. This buffer layer effectively mitigates interface resistance, minimizes cathode material volume change, and achieves excellent electrochemical performance. In addition to in-situ generation of intermediate layers by adding ionic liquids, some researchers directly add buffer layers between the cathode and electrolyte. Flexible solid materials are introduced between the cathode and electrolyte. Goodenough's team introduced plastic crystal electrolytes, which reduced interface impedance, increased cycle life, and allowed for high-rate performance, cycling more than 100 times at 5 C (Fig. 18b) [157]. Similarly, this interface layer is also introduced to the anode/electrolyte interface. Zhou and his colleagues added CPMEA between Na metal and Nasicon electrolyte (Fig. 19a) [158]. After 70 cycles at 0.2 C and 65 ℃, a stable capacity of about 102 mAh/g was maintained, indicating that the Na/CPMEA/Nasicon interface has good stability and dendrite suppression ability. Ran et al. infused antioxidant polyacrylonitrile (PAN) and antioxidant polyethylene oxide (PEO) into a layered Nasicon framework comprising a compact core layer and a porous outer layer, and the polymer formed a tight and stable interface with the electrode, greatly reducing the interface resistance (Fig. 19b) [159]. The Na/SCE/NVPF battery demonstrates impressive cycling performance. After 460 cycles, the battery exhibits a capacity of 94 mAh/g with a capacity retention of 81%. In addition to soft organic interfacial layers, inorganic interfacial layers also find many applications. Yin et al. first introduced an AlF₃ coating between the Na metal and SSE, which can react with Na during the initial cycles to generate a Na⁺ conductive buffer layer in-situ, increasing the interfacial area and suppressing Na dendrite growth by increasing the critical dendrite length (Fig. 20a) [160]. Wang et al. used a thermal decomposition method to construct an ultra-thin coating of SnS₂ between Nasicon and Na anode, and designed a Na₃V₂(PO₄)₃/Na₃Zr₂Si₂PO₁₂-SnS₂/Na battery, which exhibited remarkable rate performance (Fig. 20b) [161]. Furthermore, composite anodes are also an effective way to promote interfacial contact. Luo et al. added amorphous SiO₂ to the molten Na metal, which reduced surface tension and enabled close contact between the Na-SiO₂ composite material and Nasicon electrolyte, significantly reducing the interfacial resistance by 16 times and improving CCD performance by 5 times (Fig. 20c) [162]. Na-Na₁₅Sn₄ composite alloy is also frequently used to increase the diffusion coefficient of Na⁺ in the anode.

  Figure 18

  

  Figure 18.  (a) Schematic diagram of Na₃V₂(PO₄)₃/IL/Na_3.3Zr_1.7La_0.3Si₂PO₁₂/Na solid-state batteries and corresponding electrochemical performance. Reprinted with permission [155]. Copyright 2017, Wiley Online Library. (b) Schematic illustration of solid-state batteries with plastic-crystal electrolyte applied in the cathode and corresponding electrochemical performance. Reprinted with permission [157]. Copyright 2017, Wiley Online Library.

  DownLoad: Full-Size Img PowerPoint

  Figure 19

  

  Figure 19.  (a) Schematic illustration of solid-state batteries with CMPEA added in the anode side and corresponding electrochemical performance. Reprinted with permission [158]. Copyright 2017, American Chemical Society. (b) Schematic illustration of solid-state sodium ion batteries with sandwich composite electrolyte and corresponding electrochemical performance. Reprinted with permission [159]. Copyright 2021, Elsevier.

  DownLoad: Full-Size Img PowerPoint

  Figure 20

  

  Figure 20.  (a) Schematic illustration of the AlF₃-modified Nasicon based cells and corresponding electrochemical performance. Reprinted with permission [160]. Copyright 2020, Elsevier. (b) Schematic illustration of the SnF₂-modified Nasicon based cells and corresponding electrochemical performance. Reprinted with permission [161]. Copyright 2021, Royal Society Chemistry. (c) Schematic illustration of the SiO₂-modified Nasicon based cells and corresponding electrochemical performance. Reprinted with permission [162]. Copyright 2020, American Chemical Society.

  DownLoad: Full-Size Img PowerPoint

  In summary, both Na-β-β″-Al₂O₃ and Nasicon-type SSEs have similar strategies for optimizing the performance of batteries. Firstly, the grain boundary impedance of the oxide electrolyte is high, and high-temperature treatment is needed to improve material density and reduce grain boundary impedance. Secondly, common strategies to improve the electrode-electrolyte interface contact include thin film formation, liquid wetting or polymer addition, in-situ deposition, and porous structure design. Table 1 summarizes the performance of all-solid-state SIBs based on oxide Ni-ion SSEs [150,152,155,160,161,163-169].

  Table 1

  

  Table 1.  Summary of solid-state SIBs using oxide-based SSEs.

  DownLoad: CSV

  5.2 Solid-state SIBs based on sulfide SSEs

  Sulfide-based SSEs offer distinct advantages over oxide-based SSEs. They have higher ionic conductivity and are typically softer in texture, allowing for better contact with electrode materials through simple cold pressing and resulting in lower interfacial impedance. Therefore, the advancement of SIBs utilizing sulfide-based SSEs has attracted widespread attention. Nevertheless, several challenges remain to be addressed. Firstly, it pertains to the inherent stability of sulfide-based SSEs as numerous studies have highlighted their susceptibility to instability when exposed to air, leading to potential side reactions with water. Secondly, it is associated with ensuring the interface stability between sulfide-based SSEs and electrode materials [170-173].

  Among the Various Na-ion SSEs investigated, the cubic Na₃PS₄ stands out for its exceptional Na-ion conductivity, surpassing 10⁻⁴ S/cm at room temperature [174]. Therefore, Wang et al. suggested the utilization of Na₃PS₄ as both a solid electrolyte and an active material, thereby establishing an inherent interface contact between the electrolyte and electrode [175]. The battery Na-Sn-C/Na₃PS₄/Na₃PS₄Na₂S-C delivers a high reversible capacity of 869.2 mAh/g at 60 ℃ (Fig. 21a). In addition, Adams et al. proposed partially replacing P⁵⁺ with Sn⁴⁺, leading to the highest ionic conductivity of the Na_3.1Sn_0.1P_0.9S₄ electrolyte composition (0.52 mS/cm) [176]. The Na_2+2δFe_2-δ(SO₄)₃/Na_3.1Sn_0.1P_0.9S₄/Na₂Ti₃O₇ battery was also prepared. After 100 cycles at 2 C at 80 ℃, 82% of the initial capacity was retained (Fig. 21b). To address the air stability of the sulfide itself, Yao et al. produced a quaternary nano-solid electrolyte Na₁₀SnSb₂S₁₂ by a liquid-phase method, which exhibited excellent air stability [14]. The nano-sized electrolyte particles enhanced the interface contact between the electrolyte and electrode, resulting in a reduction in interface resistance. Therefore, the assembled TiS₂/Na₁₀SnSb₂S₁₂/Na SIBs exhibited good cycling performance (Fig. 21c). Furthermore, Na₃P_0.62As_0.38S₄ has good water stability and was used to prepare the TiS₂/Na₃P_0.62As_0.38S₄/Na-Sn solid-state battery [177]. The charge-discharge curve of the battery exhibited a significant irreversible capacity.

  Figure 21

  

  Figure 21.  (a) TEM images and corresponding elemental mappings of the NPS-nano-Na₂S-C nanocomposite cathode and electrochemical performance of the Na-Sn-C/Na₃PS₄/Na₃PS_4—Na₂S-C solid-state batteries. Reprinted with permission [175]. Copyright 2017, American Chemical Society. (b) Cycling performance of the Na_2+2δFe_2-δ(SO₄)₃/Na_3.1Sn_0.1P_0.9S₄/Na₂Ti₃O₇ solid-state batteries at different temperatures. Reprinted with permission [176]. Copyright 2017, Royal Society Chemistry. (c) The quantity of H₂S gas produced from Na₁₀SnSb₂S₁₂ exposed to humid air for different time durations and electrochemical performance of the TiS₂/Na₁₀SnSb₂S₁₂/Na solid-state batteries. Reprinted with permission [14]. Copyright 2021, Elsevier.

  DownLoad: Full-Size Img PowerPoint

  The narrow electrochemical window of sulfide electrolytes makes them prone to oxidation and reduction decomposition during battery cycling when matched with positive and anodes, leading to rapid performance degradation. Taking the well-established and widely studied Na₃PS₄ as an example, theoretical calculations show that its electrochemical window is about 1.55–2.25 V, which is poorly compatible with the electrochemical stability of high-voltage oxide cathodes, limiting its long cycle performance and practicality [178]. Therefore, Wu et al. constructed a battery, comprising a composite cathode composed of NaCrO₂ and NYZC, Na₃PS₄ electrolyte, and Na-Sn anode (Fig. 22a) [134]. At 20 ℃ and 0.1 C, the battery's first-cycle Coulombic efficiency increased sharply (from 71.9% to 97.6%). At 40 ℃ and 1 C conditions, it could cycle more than 1000 times with a capacity retention of 89.3%. The NYZC_0.74—NaCrO₂-VGCF composite cathode material exhibited stable electrochemical performance. However, in traditional cathode composite materials obtained by mechanical mixing, the electrolyte and cathode exhibit limited point-to-point contact, leading to reduce the utilization of the active material. Zhang et al. used a water-based solution to coat the surface of Se-doped SPAN active material with a layer of Na₃SbS₄, achieving close contact and uniform distribution between the electrolyte and electrode (Fig. 22b) [179]. The incorporation of the coating on the cathode resulted in a notable enhancement in Na⁺ diffusion kinetics and a reduction in cathode diffusion impedance compared to the corresponding mixed cathode. In addition, the NSS coating also reduced the volume expansion effect of the cathode. The all-solid-state SIBs using the SeSPAN-Ar@NSS (1.23 mg/cm² area loading) cathode could provide good reversible capacity for over 150 cycles. Besides employing coating techniques to improving the cathode/electrolyte interface, finding cathode materials that are more compatible with sulfide solid electrolytes is also a direction for improving the cathode interface. Nagata et al. synthesized a series of amorphous Na_0.7CoO₂—Na_xMO_y (M = N, S, P, B, or C) composite electrode materials and found that this layered transition metal oxide and sodium-containing oxyacid salt composite amorphous material can effectively improve the discharge specific capacity of electrode material [180]. At the same time, the Yao's group also successively reported high-capacity quinone-based organic cathode materials Na₄C₆O₆ and pyrene-4, 5, 9, 10-tetrone (PTO) (Fig. 22c) [183]. These materials exhibited good chemical and electrochemical compatibility with Na₃PS₄ electrolyte and resulted in very high specific capacity in all-solid-state SIBs, with outstanding stability, cycling stably for 400 and 500 cycles at 60 ℃, respectively.

  Figure 22

  

  Figure 22.  (a) Schematic illustration of NaCrO₂+NYZC/Na₃PS₄/Na₂Sn solid-state batteries and corresponding electrochemical performance. Reprinted with permission [134]. Copyright 2021, Springer Nature. (b) Schematic illustration of SeSPAN-Ar@NSS/Na₃SbS₄/Na₁₅Sn₄ solid-state batteries and corresponding electrochemical performance. Reprinted with permission [179]. Copyright 2020, Elsevier. (c) Schematic illustration of SIBs with organic cathode and corresponding electrochemical performance. Reprinted with permission [180,181]. Copyright 2018, American Chemical Society. Copyright 2019, Cell Press.

  DownLoad: Full-Size Img PowerPoint

  There are two primary challenges that need to be addressed in all-solid-state SIBs based on sulfide electrolytes with Na metal: interface compatibility and suppression of Na dendrite growth. Studies have shown that sulfide electrolytes such as Na₃PS₄ can decompose with metal Na to generate Na₂S and Na₃P, where the electronic conductivity of Na₃P is high, which can lead to continuous decomposition of the interface during cycling, increasing the total resistance. In recent times, numerous researchers have employed intermediate buffer layers, including ZrO₂, Sc₂O₃, and HfO₂, to enhance the interface compatibility between Na metal and sulfide electrolytes [182]. Li designed a phase-transition copolymer (PPP-NaTFSI) as the interlayer between the Na anode and Na₃SbS₄ (Fig. 23a) [183]. This work not only prevents the decomposition of sulfide but also facilitates the uniform deposition of sodium. Tian et al. discovered a hydrated phase Na₃SbS₄·8H₂O (Fig. 23b) [184]. The symmetric Na cell with untreated Na₃SbS₄ showed a continuously increasing voltage, while the cell with surface-hydrated Na₃SbS₄ exhibited only a small increase. The study showed that the hydrate reacted partially with metal Na upon contact, forming passive products NaH and Na₂O, which limited further decomposition of the Na₃SbS₄ solid electrolyte, reduced the growth of interface impedance during cycling, and improved interface stability with metal Na. Ionic liquids have low vapor pressure, good thermal stability, and good flow/wetting properties. Additionally, they also have good electrochemical compatibility with metal Na and Na alloy anodes. However, controlling the amount of ionic liquid additive is a key factor in improving battery performance. Therefore, Wang and colleagues presented a straightforward approach to optimize the utilization of ionic liquids (Fig. 23c) [185]. They used a PVDF membrane featuring sub-micron vertical channels to restrict the flow of ionic liquid, which enabled a uniform coating of BMTFSI with a thickness less than 500 nm on Na₃SbS₄ particles. The application of the coating offered several advantages. Firstly, it enhanced the physical contact between NSS particles, facilitating improved ion transport within the system. Secondly, it played a crucial role in effectively stabilizing the interface between NSS and metallic Na. The NVP/NSS/Na SIBs with an ionic coating structure exhibited excellent cycling and rate performance with over 400 cycles.

  Figure 23

  

  Figure 23.  (a) Schematic illustration of solid-state batteries with a phase-transition copolymer (PPP-NaTFSI) as the interlayer in the anode side and corresponding electrochemical performance. Reprinted with permission [183]. Copyright 2021, Wiley Online Library. (b) Schematic illustration of Na₃SbS₄ electrolyte with surface hydrate coating and corresponding SXRD profiles of Na/Na₃SbS₄/Na symmetric cell. Reprinted with permission [184]. Copyright 2019, Elsevier. (c) Schematic illustration of solid-state SIBs with PVDF membrane with sub-micron vertical channels and corresponding electrochemical performance. Reprinted with permission [185]. Copyright 2022, Elsevier.

  DownLoad: Full-Size Img PowerPoint

  Surprisingly, Li et al. started from the electrolyte and proposed using Na₄P₂S_4.5O_2.5 oxysulfide as an additive to prepare a stable electrolyte-electrode interface by constructing a Na₃SbS₄ composite electrolyte (Fig. 24) [186]. For the first time, the behavior of different sub-lattices during cycling was revealed, and the migration of O ions was observed in the oxysulfide electrolyte. The Na//Na symmetrical cell showed high critical current density and long cycling stability of sodium plating/stripping, far exceeding that of the Na₃SbS₄ single-phase electrolyte. The assembled Se_0.8S_7.2@PAN/NaPSO+NSS/Na₁₅Sn₄ all-solid-state battery cycled 1000 times under room temperature and 0.3 C, demonstrating excellent electrochemical performance.

  Figure 24

  

  Figure 24.  (a) Mechanism schematic illustration of the Se_0.8S_7.2@PAN/NaPSO+NSS/Na₁₅Sn₄ solid-state batteries. (b) Raman spectra, SEM image and Ionic conductivity of NaPSO+NSS. (c) electrochemical performance of the Se_0.8S_7.2@PAN/NaPSO+NSS/Na₁₅Sn₄ solid-state batteries. (d) Cross-sectional SEM images and EDS mapping for the Se_0.8S_7.2@PAN/NaPSO+NSS/Na₁₅Sn₄ and corresponding electrochemical performance. Reprinted with permission [186]. Copyright 2022, Wiley Online Library.

  DownLoad: Full-Size Img PowerPoint

  In general, in contrast to the rigid and fragile characteristics of oxide solid electrolytes, sulfide-based materials are relatively soft and can have good contact with electrode materials through simple cold pressing. Therefore, the strategy of using liquid wetting agents at the interface is less commonly employed in sulfide-based all-solid-state SIBs. The main focus is on matching compatible positive and anode materials or introducing stable interfacial layers on both sides of the electrodes to prevent electrolyte oxidation and decomposition during battery cycling, which can lead to rapid performance degradation. Table 2 provides a summary of SIBs based on sulfide Na-ion SSEs [134,179,181,182,187-194].

  Table 2

  

  Table 2.  Summary of solid-state SIBs using sulfide-based SSEs.

  DownLoad: CSV

  5.3 Solid-state SIBs based on anti-perovskites SSEs

  In recent years, researchers have started investigating Anti-perovskites Na-ion SSEs, which can exhibit extremely high ionic conductivity. However, these electrolytes exhibit high moisture susceptibility and have a tendency to absorb humidity [55,108]. Additionally, similar to oxide-type Na-ion SSEs, these electrolytes have relatively high grain boundary impedance, which limits their application in full cells. Braga et al. utilized precursor materials similar to Na₃OCl and conducted a series of heat treatments after mixing the raw materials in a deionized water environment. This process resulted in the formation of an amorphous glassy electrolyte, Na_3-xH_xOCl (0 < x < 1) [195]. Na/ferrocene cells assembled with Na_3-xH_xOCl as the solid electrolyte demonstrate good cycling stability at room temperature (Fig. 25a). Unlike traditional oxide SSEs, perovskite-type Na-ion SSEs can be synthesized at lower temperatures, making them easier to process. They also exhibit higher ionic conductivity. If the solid-solid interface contact issue is properly addressed, this material holds great potential for application in all-solid-state SIBs. Fan et al. synthesized an innovative anti-perovskite SSEs, Na_2.98Mg_0.01SO₄F_0.95Cl_0.05, using low-cost precursors through a solid-state reaction at 500 ℃ (Fig. 25b) [54]. The incorporation of Mg²⁺ and Cl⁻ led to a three-order-of-magnitude increase in ionic conductivity, reaching approximately 10⁻⁴ S/cm at 60 ℃. Based on this electrolyte, a Fe[Fe(CN)₆]₃/Na_2.98Mg_0.01SO₄F_0.95Cl_0.05/Na-Sn solid-state battery was assembled, demonstrating reversible operation. The battery exhibited an initial discharge capacity of 91.0 mAh/g, which was followed by a sustained reversible capacity of 77.0 mAh/g. Jiang et al. investigated a vacancy-deficient cluster-ion based anti-perovskite solid-state electrolyte called Na₂BH₄NH₂ (Fig. 25c) [196]. A NaSn|Na₂BH₄NH₂|NaSn symmetric cell was cycled for 500 h at a current density of 0.1 mA/cm². Furthermore, the compatibility of Na₂BH₄NH₂ electrolyte with electrode materials was demonstrated using TiS₂ cathode, indicating good compatibility between Na₂BH₄NH₂ and electrode materials.

  Figure 25

  

  Figure 25.  (a) Cycling performance of the Na/ferrocene cells assembled with Na_3-xH_xOCl. Reprinted with permission [195]. Copyright 2017, Royal Society Chemistry. (b) Schematic illustration of Fe[Fe(CN)₆]₃/Na_2.98Mg_0.01SO₄F_0.95Cl_0.05/Na-Sn and corresponding electrochemical performance. Reprinted with permission [54]. Copyright 2020, Elsevier. (c) Electrochemical performance of NaSn|Na₂BH₄NH₂|NaSn symmetric cell and TiS₂|Na₂BH₄NH₂|NaSn solid-state batteries. Reprinted with permission [196]. Copyright 2023, Wiley Online Library.

  DownLoad: Full-Size Img PowerPoint

  5.4 Solid-solid SIBs based on complex hydrides SSEs

  Orimo's group was among the first to develop sodium-based complex hydrides SSEs and investigate their electrochemical performance in solid-state batteries [63]. These complex hydrides typically exhibit a characteristic first-order phase transition at a certain temperature, with a several-orders-of-magnitude increase in conductivity from the low-temperature phase to the high-temperature phase [197]. However, the relatively high phase transition temperature has limited their application in all-solid-state SIBs. Over the past few years many researchers have successfully lowered the phase transition temperature of some materials to room temperature or even eliminated the phase transition through chemical modifications, size control, partial anion substitution, or mixing different anions [198-201]. These advancements have greatly facilitated their application in solid-state batteries. Although the optimized sodium ion complex hydrides electrolytes have significantly improved conductivity, the reported all-solid-state SIBs using complex hydrides electrolytes are still limited due to factors such as large interfacial resistance and poor cycling stability. So far, most all-solid-state SIBs using complex hydrides electrolytes have employed TiS₂ or NaCrO₂ cathodes with a voltage range of around 3 V. In recent years, some researchers have utilized mixed borohydrides as solid electrolytes in all-solid-state SIBs. Yoshida et al. obtained pseudo-binary complex hydrides as Na-ion SSEs by mechanically ball-milling pure Na₂B₁₀H₁₀ and Na₂B₁₂H₁₂ [202]. The ion conductivity of the mixture showed a direct correlation with the fraction of the body-centered cubic phase. It reached the maximum value when the molar ratio was 1:3. The initial specific discharge capacity of the cell assembled with Na metal anode and TiS₂ cathode is close to the theoretical capacity of TiS₂ (239 mAh/g). Duchêne et al. developed a NaCrO₂|Na₂(B₁₂H₁₂)_0.5(B₁₀H₁₀)_0.5|Na battery [203]. By impregnating the electrolyte solution onto the cathode particles to improve the electrode-electrolyte contact, a stable solid-solid interface was established between the electrolyte and the cathode, enhancing the battery performance (Fig. 26a). This resulted in reversible and stable cycling. After 20 cycles at 0.05 C, the capacity retention exceeded 90%, and after 250 cycles at 0.2 C, the capacity retention reached 85%. The Duchêne research group also infiltrated the soluble-processed complex hydrides Na₄(B₁₂H₁₂)(B₁₀H₁₀) into a porous electrode made through slurry casting [204]. They reported that an all-solid-state SIBs with a NaCrO₂ electrode infiltrated with Na₄(B₁₂H₁₂)(B₁₀H₁₀) maintained 95.6% of its initial capacity after 100 cycles at 0.5 C (Fig. 26b).

  Figure 26

  

  Figure 26.  (a) Schematic illustration of NaCrO₂|Na₂(B₁₂H₁₂)_0.5(B₁₀H₁₀)_0.5|Na solid-state batteries and corresponding electrochemical performance. Reprinted with permission [203]. Copyright 2013, Royal Society Chemistry. (b) Schematic illustration of solid-state batteries with a NaCrO₂ electrode infiltrated with Na₄(B₁₂H₁₂)(B₁₀H₁₀) and corresponding electrochemical performance. Reprinted with permission [204]. Copyright 2020, Elsevier.

  DownLoad: Full-Size Img PowerPoint

  Further investigations have focused on a subgroup of complex hydrides, borohydrides, and their C-derived compounds known as carba‑closo-borates, which exhibit excellent conductivity values even at room temperature. Murgia et al. recorded a fast Na⁺ conductor, Na₄(CB₁₁H₁₂)₂(B₁₂H₁₂) [205]. Integrating this electrolyte into a Na/NaCrO₂ battery demonstrated excellent electrochemical performance (Fig. 27a). The specific capacity of the battery was approximately 100 mAh/g at an output current density of 24 mA/g. In-situ EIS testing during cycling revealed a stable Na-Na symmetric cell, with the conductor enabling stable Na plating/stripping for over 500 h, exhibiting limited polarization and low interfacial resistance. Brighi et al. synthesized a novel compound, Na_2-x(CB₁₁H₁₂)_x(B₁₂H₁₂)_1-x (Fig. 27b) [206]. Symmetric cells with Na-Sn alloy electrodes exhibited reversible Na⁺ shuttling and limited polarization over a working period exceeding 700 h, demonstrating their electrochemical stability. In addition, He and colleagues proposed an innovative Na⁺ conductor Na₃NH₂B₁₂H₁₂ (Fig. 27c) [207]. It has an impressive electrochemical window of 10 V vs. Na⁺/Na. By utilizing this electrolyte, an all-solid-state SIBs was constructed with a TiS₂ cathode and Na foil anode. Operating at 80 ℃ and 0.1 C, the battery demonstrated over 200 reversible cycles with a capacity retention rate exceeding 50%.

  Figure 27

  

  Figure 27.  (a) Electrochemical performance of NaCrO₂/Na₄(CB₁₁H₁₂)₂(B₁₂H₁₂)/Na solid-state batteries. Reprinted with permission [205]. Copyright 2019, Elsevier. (b) Electrochemical performance of NaSn/Na_4/3(CB₁₁H₁₂)_2/3(B₁₂H₁₂)_1/3/NaSn symmetric cell. Reprinted with permission [206]. Copyright 2018, Elsevier. (c) CV test results for Na/Na₃NH₂B₁₂H₁₂/Pt and electrochemical performance of TiS₂/Na₃NH₂B₁₂H₁₂/Na solid-state batteries. Reprinted with permission [207]. Copyright 2018, Elsevier.

  DownLoad: Full-Size Img PowerPoint

  The electrochemical window of most complex hydrides is narrow, with the electrochemical stability primarily limited to 4 V. This incompatibility with 4 V cathode materials results in lower energy density. To address this issue, Remhof's research group showcased the effectiveness of a hydroborate-based electrolyte, Na₄(CB₁₁H₁₂)₂(B₁₂H₁₂), which comprises two distinct cage-like anions with varying oxidation stability. This electrolyte demonstrated successful interface passivation with high-voltage cathodes operating at 4 V, effectively mitigating impedance growth during cycling [208]. They achieved the first demonstration of a 4 V-level complex hydrides electrolyte-based all-solid-state SIBs. The battery employed Na₃(VOPO₄)₂F as the cathode and Na metal as the anode, without any artificial coating or protection. The configuration of the battery was Na₃(VOPO₄)₂F/Na₄(CB₁₁H₁₂)₂(B₁₂H₁₂)/Na. When cycled to 4.15 V vs. Na⁺/Na at 0.1 C, the battery exhibited a discharge capacity of 104 mAh/g, and at 0.2 C, the capacity was 99 mAh/g. After 800 cycles at room temperature, the capacity retention rates were 78% and 76%, respectively (Fig. 28a). Furthermore, Guo and collaborators achieved, for the first time, an electrochemical stability window using the novel complex hydrides Na₃B₂₄H₂₃ (6.7 V vs. Na⁺/Na) as the electrolyte [209]. A high-voltage all-solid-state sodium battery, Na[Ni_1/3Fe_1/3Mn_1/3]O₂/Na₃B₂₄H₂₃–5Na₂B₁₂H₁₂/Na, was assembled using Na₃B₂₄H₂₃–5Na₂B₁₂H₁₂ in conjunction with a 4 V-level cathode. After 100 cycles at 60 ℃ and 0.1 C, the capacity retention rate reached 84% (Fig. 28b). This discovery provides an opportunity for the development of novel complex hydrides electrolyte-based all-solid-state SIBs with high voltage and high energy density.

  Figure 28

  

  Figure 28.  (a) Schematic illustration of Na₃(VOPO₄)₂F/Na₄(CB₁₁H₁₂)₂(B₁₂H₁₂)/Na solid-state batteries and corresponding electrochemical performance. Reprinted with permission [208]. Copyright 2020, Royal Society Chemistry. (b) Schematic illustration of Na[Ni_1/3Fe_1/3Mn_1/3]O₂/Na₃B₂₄H₂₃–5Na₂B₁₂H₁₂/Na solid-state sodium batteries and corresponding electrochemical performance. Reprinted with permission [209]. Copyright 2023, Elsevier.

  DownLoad: Full-Size Img PowerPoint

  In general, complex hydrides electrolytes have shown superior compatibility with metallic Na, high chemical/stability, good mechanical properties, low weight density, excellent solution processability, and low toxicity. Particularly, their high deformability and high solubility are advantageous for processing into solid-state batteries. However, the complex synthesis process and electrochemical instability limit their application in solid-state batteries. So far, only a few complex hydride SSEs have been developed and applied in all-solid-state SIBs. Table 3 summarizes the solid-state batteries based on complex hydrides Na-ion electrolytes [202,203,205,207-211]. Furthermore, although complex hydrides have a notable impact on stabilizing the deposition/stripping of sodium, the properties of their electrochemical decomposition products are still unclear and require further investigation.

  Table 3

  

  Table 3.  Summary of solid-state SIBs using complex hydrides SSEs.

  DownLoad: CSV

  5.5 Solid-state SIBs based on halide SSEs

  In recent years, lithium halide superionic conductors have been widely studied. Generally, lithium halides can maintain oxidation limits exceeding 4 V, demonstrating good (electro)chemical stability with active materials such as LiCoO₂ and NCM811. Although sodium halides are not as good as lithium halides in terms of high-pressure stability and ion conductivity, they can still be used as ion conductor additives in high-pressure oxide-based cathodes instead of sulfide electrolytes. For example, Na₂ZrCl₆, Na_3-xY_1-xZr_xCl₆, and NaAlCl₄ (synthesized by ball milling/mechanochemical methods) have shown good compatibility with NaCrO₂. Li et al. found that NaYbZrCl_0.75 additive exhibits good compatibility with NaCrO₂, inducing a positive effect on the volume change during cathode cycling. It induces the formation of a special core-shell structure on the surface of NaCrO₂ grains, with uniform NaYbZrCl_0.75 encapsulation, promoting uniform ion transport and alleviating (electro)chemical mechanical stress. The all solid-state SIBs with NaYbZrCl_0.75 additive exhibits good cycling stability, maintaining a capacity retention of 74.1% after 1000 cycles at 25 ± 3 ℃ and 83.0% after 1900 cycles at 60 ℃ (Fig. 29a) [135]. Fu et al. developed a novel class of halide heterostructure electrolytes with good deformability and high-pressure stability, enabling stable cathode/electrolyte interfaces with Na_0.85Mn_0.5Ni_0.4Fe_0.1O₂ cathode through simple cold pressing, achieving a capacity retention of 91.0% after 100 cycles at 0.2 C in all solid-state SIBs [212]. In addition, Hu et al. successfully combined a Na-ion halide electrolyte (NaTaCl₆) with a poly-anion-type Na₃V₂(PO₄)₃ cathode for the first time, demonstrating stable long-term cycling performance after 4000/600/1500 cycles at 3/1/0.5 C rates, with capacity retentions of 81%/95%/98%, respectively. The comprehensive performance of the reported all solid-state SIBs surpasses all previously reported results (Fig. 29b) [138].

  Figure 29

  

  Figure 29.  (a) Schematic diagram illustrating the structural changes in the composite cathode of NaCrO_2—NYbZrCl_0.75/Na₃PS₄/Na₂Sn solid-state batteries and corresponding electrochemical performance. Reprinted with permission [135]. Copyright 2023, Elsevier. (b) Schematic illustration of NVP/NaTaCl₆/Na₃PS₄/Na₁₅Sn₄ solid-state sodium batteries and corresponding electrochemical performance [138]. Copyright 2023, Cell Press.

  DownLoad: Full-Size Img PowerPoint

  After optimization, Na-ion halides electrolyte generally exhibit good ion conductivity, wide electrochemical oxidation windows, and a certain degree of toughness, making them excellent positive electrode ion conductor additives for use in all-solid-state batteries. In conclusion, halide electrolytes hold promising prospects in the field of batteries, especially suitable for high voltage and high-performance requirements, thus offering more efficient and sustainable energy storage solutions for future renewable energy applications.

  6. Advanced characterizations for solid-state sodium batteries

  Although many researchers have made significant efforts to address interface issues in electrode/electrolyte interfaces, there are still many unanswered questions, such as the exact composition and structure of the interfaces, their evolution processes, their correlation with battery performance, and the formation mechanisms. To convincingly answer these questions, more advanced characterization techniques need to be developed to uncover clues. Currently, some literature has used certain characterization techniques to reveal the structural features and functional mechanisms of the constructed electrode/electrolyte interfaces accurately and intuitively. Furthermore, different characterization techniques can be classified as in-situ or ex-situ. In-situ characterization, compared to ex-situ characterization, directly reveals the internal electrochemical processes of redox reactions, allowing real-time monitoring of the evolution of interface composition and structure. While in-situ characterization provides direct evidence of the internal redox reactions in batteries, such tests are often operationally challenging and have specific requirements for the testing equipment. Therefore, the available types of techniques for in-situ characterization are still quite limited.

  Similar to lithium, dendritic growth of sodium also occurs in the form of branching, causing short circuits in sodium metal batteries. Considerable investigation has been undertaken to explore the mechanisms governing the growth and deposition of lithium dendrites. However, there is limited research on sodium dendrite growth and sodium deposition. The dynamics of sodium dendrite deposition/stripping are still not well understood, which hampers the development of strategies for uniform sodium deposition and stripping. Therefore, real-time observation of metal sodium deposition and stripping is crucial to understanding their growth mechanisms and inhibiting the formation of sodium dendrites. Huang et al. conducted real-time observation studies on sodium dendrite growth and stripping in a carbon dioxide atmosphere using advanced aberration-corrected environmental transmission electron microscopy (ETEM) [213]. It is worth noting that the deposition rate can be adjusted by controlling the applied voltage. Therefore, the morphology of deposited sodium can be controlled by adjusting the applied voltage, which may provide important clues for mitigating dendritic growth in all-solid-state SIBs. In a recent investigation, Huang et al. performed in-situ optical microscopy (OM) examinations to observe the growth of Na dendrites in a symmetric cell Na/β″-Al₂O₃/Na [214]. Complementing this analysis, they employed in-situ SEM to examine the dynamic interplay between Na deposition and the propagation of cracks. The outcomes of their research demonstrated that crack formation preceded the deposition of Na, with subsequent Na deposition inducing the generation of additional cracks. Notably, Na deposition persisted alongside crack propagation until a short circuit event was triggered. This study substantially contributes to our comprehension of the failure mechanisms inherent in all-solid-state SIBs employing Na-β″-Al₂O₃ SSEs.

  So far, there has been relatively limited research on the mechanical properties related to sodium dendrites due to the technical difficulties in preparing samples suitable for nanomechanical testing. Liu et al. conducted real-time observation of the growth characteristics of sodium dendrites and simultaneous measurement of their mechanical properties using an environmental transmission electron microscope-atomic force microscope (ETEM-AFM) platform [215]. The measured maximum strength of the sodium dendrites reached up to 203 MPa, over 300 times higher than bulk Na. The results also show that the Na dendrites creep through the cracks and pores in SSEs, resulting in the failure of the battery. Therefore, reducing the defect size in the SSEs is crucial for mitigating battery failure induced by sodium dendrites. To sum up, the findings of this work provide important insights into mitigating dendritic short circuits in solid-state sodium metal batteries.

  Furthermore, these characterization techniques also contribute to understanding cathode/electrolyte interface issues, such as interface kinetics and cathode/electrolyte reactions. Zhang's team used in-situ TEM to reveal the fundamental mechanism of Se catalytic conversion reactions in solid-state Na-SeS₂ nanobatteries [192]. The formation of amorphous Na-Se_xS_y, which significantly reduces the energy barrier of the redox reaction, was observed for the first time during the conversion process. Furthermore, the discoveries made by in-situ TEM guide the design of cathode components. Optimized composite cathodes in all-solid-state Na-SeS₂ batteries demonstrate excellent cycling capability based on the insights gained from in-situ TEM.

  Before addressing interface issues, it is important to recognize that the bare Na metal's contact with ambient gases during production and transportation sets the initial conditions for SEI formation, thus establishing the initial conditions for subsequent anode and electrolyte interface reactions. Therefore, understanding the stability of metallic Na in dry air is crucial. Li et al. conducted a more in-depth study on the chemical reactions between alkali metals and ambient gases and improved the electrochemical performance of Na by controlling interface stability [216]. They used in-situ ETEM to observe the in-situ formation of Na metal particles in specifically designed solid-state cells under high vacuum. They found that to maintain the stability of Na in dry air, in-situ surface treatment (exposure to a CO₂ atmosphere) prior to the oxidation process, introducing a protective layer (Na₂CO₃), can address the issue of spontaneously forming low-quality oxide passivation layers.

  Undoubtedly, the advanced characterization techniques and analysis methods mentioned above provide valuable insights into the structure, composition, and electrochemical properties of electrode/electrolyte interfaces. Nevertheless, it is essential to acknowledge that every technique or method possesses inherent limitations and disadvantages. Thus, it is imperative to employ a combination of multiple techniques to achieve comprehensive characterization and analysis, enabling us to derive dependable and precise conclusions. Bruce et al. from the University of Oxford employed T2 relaxation-diffusion magnetic resonance imaging (direct T2 contrast MRI) to investigate the interface microstructure of the Na/Na-β″-Al₂O₃/Na symmetric battery before and after cycling of the metallic Na electrode [217]. Additionally, to further visualize and validate the microstructure of the dendrites, Bruce combined other non-in-situ methods such as X-ray computed tomography and SEM to track and observe morphological changes inside the battery. Dendritic growth with a fractal morphology was observed in short-circuited cells, consistent with the ²³Na MRI results. The synergistic use of multiple methods mentioned above provides a deeper understanding of crack formation, microstructure growth, and the morphology of Na dendrites.

  In summary, the intuitive and accurate clues obtained through various advanced characterization techniques provide insights for the development of new approaches to address electrode/electrolyte interface issues, thereby promoting research and development of high-performance all-solid-state SIBs. The continuous advancement and application of these techniques will further drive the progress of sodium battery technology, providing better solutions for sustainable energy storage.

  7. Summary and outlook

  All-solid-state SIBs are considered an ideal next-generation energy storage device for its significant cost and safety advantages. As the core part of all-solid-state SIBs, SSEs have captured the interest of researchers as they show great advantages in the cycling performance and safety performance of batteries. In this paper, we review the research progress of different types of inorganic SSEs, including the structure of electrolytes, synthesis methods, modification strategies, applications in SIBs and so on, as well as the surface design and advanced characterization of solid-state sodium battery interfaces.

  Inorganic Na-ion SSEs have received a lot of attention due to their good electrochemical and safety properties, but there are some problems, such as alumina trioxide, the only commercially available Na-ion SSE, which has a low room temperature ionic conductivity. The modified Nasicon electrolyte has improved r.t. ionic conductivity and also has good chemical and electrochemical stability. However, there are interface problems with the electrodes. Sulfide electrolytes are soft in texture and have high r.t. ionic conductivity, but sulfides are poorly air-stable, etc. Based on these problems, scientists have made great efforts to improve the ionic conductivity and chemical/electrochemical stability of inorganic Na-ion SSEs.

  In order to achieve practical applications of all-solid-state SIBs, we believe that future research will focus mainly on the following parts: (1) The role of Na-ion SSEs in cathode mixture of all-solid-state SIBs needs to be unraveled. Clarifying the relationship between redox activity and electrochemical stability in Na-ion SSEs is crucial. This aids in guiding the interface and material design for all-solid-state SIBs. (2) Na-ion halide electrolytes offer advantages over Na-ion sulfide electrolytes in terms of lower interfacial impedance and higher voltage stability. However, their application is currently limited by their lower ionic conductivity at room temperature. Significant progress has been made in developing novel electrolytes with high room temperature sodium ion conductivity, with an increasing number of sodium halide solid electrolytes exhibiting superior conductivity being reported. Therefore, the development of Na-ion halide electrolytes with excellent comprehensive performance can facilitate their application in the field of all-solid-state SIBs. (3) The thickness of the electrolyte is crucial for the energy density of the battery. Therefore, the design of electrolyte membranes based on current Na-ion SSEs is significant for the practical application of all-solid-state SIBs.

  Thus, the development of all-solid-state SIBs still face various challenges that require further exploration and investigation.

  Declaration of competing interest

  The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.

  Acknowledgments

  This work was supported by the National Natural Science Foundation of China (Nos. 22175070, 22293041). This work was also supported by the National Key Research and Development Program (Nos. 2021YFB2500200, 2021YFB2400300), the National Natural Science Foundation of China (No. 52177214), and China Fujian Energy Devices Science and Technology Innovation Laboratory Open Fund (No. 21C-OP202211).

  
  
• 1. [1]

     K. Song, C. Liu, L. Mi, et al., Small 17 (2021) e1903194.

  2. [2]

     C. Yang, S. Xin, L. Mai, et al., Adv. Energy Mater. 11 (2020) 2000974.

  3. [3]

     J. Hwang, S. Myung, Y. Sun, Chem. Soc. Rev. 46 (2017) 3529–3614.

  4. [4]

     Z. Jiang, C. Yu, S. Chen, et al., Scripta Mater. 227 (2023) 115303.

  5. [5]

     Z. Wang, C. Tang, Z. Wang, et al., Energy Mater. Adv. 4 (2023) 1–13.

  6. [6]

     Q. Zhang, X. Shen, Q. Zhou, et al., Energy Mater. Adv. 2022 (2022) 1–11.

  7. [7]

     L. Shen, S. Deng, R. Jiang, et al., Energy Storage Mater. 46 (2022) 175–181.

  8. [8]

     H. Wan, J. Mwizerwa, F. Han, et al., Nano Energy 66 (2019) 104109.

  9. [9]

     H. Wan, W. Weng, F. Han, et al., Nano Today 33 (2020) 100860.

  10. [10]

      H. Wan, J. Mwizerwa, X. Qi, et al., ACS Appl. Mater. Interfaces 10 (2018) 12300–12304. doi: 10.1021/acsami.8b01805

  11. [11]

      Y. Yuno fano, J. Kummer, J. Inorg, Nuclear Chem. 29 (1967) 2453–2466.

  12. [12]

      F. Colò, F. Bella, J.R. Nair, et al., J. Power Sources 365 (2017) 293–302.

  13. [13]

      J. Wu, R. Zhang, Q. Fu, et al., Adv. Funct. Mater. 31 (2020) 2008165.

  14. [14]

      G. Liu, X. Sun, X. Yu, et al., Chem. Engin. J. 420 (2021) 127692.

  15. [15]

      R. Zhao, Y. Wu, Z. Liang, et al., Energy Environ. Sci. 13 (2020) 2386–2403. doi: 10.1039/d0ee00153h

  16. [16]

      W. Richards, L. Miara, Y. Wang, et al., Chem. Mater. 28 (2015) 266–273.

  17. [17]

      Z. Zhang, S. Wenzel, Y. Zhu, et al., ACS Appl. Energy Mater. 3 (2020) 7427–7437. doi: 10.1021/acsaem.0c00820

  18. [18]

      L. Shen, J. Yang, G. Liu, et al., Mater. Today Energy 20 (2021) 100691.

  19. [19]

      Y. Wang, S. Song, C. Xu, et al., Nano Mater. Sci. 1 (2019) 91–100.

  20. [20]

      G. Huang, P. Zheng, W. Li, et al., Funct. Mater. Lett. 14 (2021) 2130005. doi: 10.1142/s179360472130005x

  21. [21]

      X. Zhu, K. Wang, Y. Xu, et al., Energy Storage Mater. 36 (2021) 291–308.

  22. [22]

      E. Quartarone, P. Mustarelli, Chem. Soc. Rev. 40 (2011) 2525–2540. doi: 10.1039/c0cs00081g

  23. [23]

      L. Ran, A. Baktash, M. Li, et al., Energy Storage Mater. 40 (2021) 282–291.

  24. [24]

      P. Ding, Z. Lin, X. Guo, et al., Mater. Today 51 (2021) 449–474.

  25. [25]

      L. Gao, B. Tang, H. Jiang, et al., Adv. Sustain. Syst. 6 (2021) 2100389.

  26. [26]

      A. Li, X. Liao, H. Zhang, et al., Adv. Mater. 32 (2020) e1905517.

  27. [27]

      L. Tian, Y. Liu, Z. Su, et al., J. Mater. Chem. A 9 (2021) 23882–23890. doi: 10.1039/d1ta06269g

  28. [28]

      L. Yue, J. Ma, J. Zhang, et al., Energy Storage Mater. 5 (2016) 139–164.

  29. [29]

      Q. Zhao, P. Chen, S. Li, et al., J. Mater. Chem. A 7 (2019) 7823–7830. doi: 10.1039/c8ta12008k

  30. [30]

      G. Wang, X. Zhu, A. Rashid, et al., J. Mater. Chem. A 8 (2020) 13351–13363. doi: 10.1039/d0ta00335b

  31. [31]

      A. Manthiram, X. Yu, S. Wang, Nat. Rev. Mater. 2 (2017) 16103.

  32. [32]

      J. Zhang, J. Zhao, L. Yue, et al., Adv. Energy Mater. 5 (2015) 1501082.

  33. [33]

      H. Wang, Y. Chen, Z.D. Hood, et al., Angew. Chem. Int. Ed. 55 (2016) 8551–8555. doi: 10.1002/anie.201601546

  34. [34]

      S. He, Y. Xu, Y. Chen, et al., J. Mater. Chem. A 8 (2020) 12594–12602. doi: 10.1039/c9ta12213c

  35. [35]

      S. Takeuchi, K. Suzuki, M. Hirayama, et al., J. Solid State Chem. 265 (2018) 353–358.

  36. [36]

      Z. Wu, S. Chen, C. Yu, et al., Chem. Engin. J. 442 (2022) 136346.

  37. [37]

      H. Hong, Mat. Res. Bull. 11 (1976) 173–182.

  38. [38]

      J. Goodenough, H. Hong, A. Kafalas, Mat. Res. Bull. 11 (1976) 203–220.

  39. [39]

      M. Evstigneeva, V. Nalbandyan, A. Petrenko, et al., Chem. Mater. 23 (2011) 1174–1181. doi: 10.1021/cm102629g

  40. [40]

      W. Xia, Y. Zhao, F. Zhao, et al., Chem. Rev. 122 (2022) 3763–3819. doi: 10.1021/acs.chemrev.1c00594

  41. [41]

      Y. Sadikin, M. Brighi, P. Schouwink, et al., Adv. Energy Mater. 5 (2015) 1501016.

  42. [42]

      Y. Qie, S. Wang, S. Fu, et al., J. Phys. Chem. Lett. 11 (2020) 3376–3383. doi: 10.1021/acs.jpclett.0c00010

  43. [43]

      X. He, Y. Zhu, Y. Mo, Nat. Commun. 8 (2017) 15893.

  44. [44]

      G. Liu, J. Yang, J. Wu, et al., Adv. Mater. 36 (2024) 2311475. doi: 10.1002/adma.202311475

  45. [45]

      C. Delmas, Adv. Energy Mater. 8 (2018) 1703137.

  46. [46]

      K. Hueso, M. Armand, T. Rojo, Energy Environ. Sci. 6 (2013) 734–749. doi: 10.1039/c3ee24086j

  47. [47]

      H. Li, H. Fan, Z. Liu, et al., Sens. Actuator. B: Chem. 255 (2018) 1445–1454.

  48. [48]

      X. Lu, G. Xia, J.P. Lemmon, et al., J. Power Sources 195 (2010) 2431–2442.

  49. [49]

      S. Lee, D. Lee, S. Lee, et al., Bull. Mater. Sci. 39 (2016) 729–735. doi: 10.1007/s12034-016-1199-6

  50. [50]

      G. Yamaguchi, Bull. Chem. Soc. Jpn. 41 (1968) 93–99. doi: 10.1246/bcsj.41.93

  51. [51]

      M. Peters, J. Phys. Chem. 73 (1969) 1774–1780.

  52. [52]

      L. Boyer, P. Edwardson, Ferroelectrics 104 (1990) 417–422. doi: 10.1080/00150199008223849

  53. [53]

      R. Niewa, Z. Anorg. Allg. Chem. 639 (2013) 1699–1715. doi: 10.1002/zaac.201300063

  54. [54]

      S. Fan, M. Lei, H. Wu, et al., Energy Storage Mater. 31 (2020) 87–94.

  55. [55]

      H. Nguyen, S. Hy, E. Wu, et al., J. Electrochem. Soc. 163 (2016) A2165–A2171. doi: 10.1149/2.0091610jes

  56. [56]

      Y. Yu, Z. Wang, G. Shao, J. Mater. Chem. A 6 (2018) 19843–19852. doi: 10.1039/c8ta08412b

  57. [57]

      C. Wei, R. Wang, Z. Wu, et al., Chem. Engin. J. 476 (2023) 146531.

  58. [58]

      C. Wei, C. Yu, R. Wang, et al., J. Power Sources 559 (2023) 232659.

  59. [59]

      C. Wei, S. Chen, C. Yu, et al., Appl. Mater. Today 31 (2023) 101770.

  60. [60]

      C. Wei, C. Liu, Y. Xiao, et al., Adv. Funct. Mater. 34 (2024) 2314306.

  61. [61]

      Q. Luo, C. Yu, C. Wei, et al., Ceram. Int. 49 (2023) 11485–11493.

  62. [62]

      Z. Zhang, E. Ramos, F. Lalère, et al., Energy Environ. Sci. 11 (2018) 87–93. doi: 10.1039/c7ee03083e

  63. [63]

      H. Oguchi, M. Matsuo, S. Kuromoto, et al., J. Appl. Phys. 111 (2012) 036102.

  64. [64]

      M. Matsuo, S. Kuromoto, T. Sato, et al., Appl. Phys. Lett. 100 (2012) 203904.

  65. [65]

      T. Udovic, M. Matsuo, A. Unemoto, et al., Chem. Commun. 50 (2014) 3750–3752.

  66. [66]

      S. Shan, L. Yang, X. Liu, et al., J. Alloys Compd. 563 (2013) 176–179.

  67. [67]

      S. Butee, K. Kambale, M. Firodiya, Process. Appl. Ceram. 10 (2016) 67–72.

  68. [68]

      L. Ghadbeigi, A. Szendrei, P. Moreno, et al., Solid State Ionics 290 (2016) 77–82.

  69. [69]

      S. Naqash, F. Tietz, E. Yazhenskikh, et al., Solid State Ionics 336 (2019) 57–66.

  70. [70]

      Z. Zhang, S. Shi, Y. Hu, et al., J. Inorg. Mater. 28 (2013) 1255–1260.

  71. [71]

      S. Yubuchi, A. Hayashi, M. Tatsumisago, Chem. Lett. 44 (2015) 884–886. doi: 10.1246/cl.150195

  72. [72]

      T. Kim, K. Park, Y. Choi, et al., J. Mater. Chem. A 6 (2018) 840–844. doi: 10.1039/c7ta09242c

  73. [73]

      A. Virkar, R. Gordon, J. Am. Ceram. Soc. 60 (1977) 58–61. doi: 10.1111/j.1151-2916.1977.tb16094.x

  74. [74]

      C. Zhu, J. Xue, J. Alloys Compd. 517 (2012) 182–185.

  75. [75]

      X. Wei, Y. Cao, L. Lu, et al., J. Alloys Compd. 509 (2011) 6222–6226.

  76. [76]

      D. Xu, H. Jiang, Y. Li, et al., Eur. Phys. J. Appl. Phys. 74 (2016) 10901. doi: 10.1051/epjap/2016150466

  77. [77]

      L. Yang, S. Shan, X. Wei, et al., Ceram. Int. 40 (2014) 9055–9060.

  78. [78]

      H. Erkalfa, Z. Baykara, Ceram. Int. 24 (1998) 81–90.

  79. [79]

      G. Chen, J. Lu, X. Zhou, et al., Ceram. Int. 42 (2016) 16055–16062.

  80. [80]

      C. Zhu, Y. Hong, P. Huang, J. Alloys Compd. 688 (2016) 746–751.

  81. [81]

      Y. Viswanathan, A. Virkar, J. Mater. Sci. 18 (1983) 109–113.

  82. [82]

      K. Yuria Saito, T. Asai, H. Kageyama, et al., Solid State Ionics 58 (1992) 327–331.

  83. [83]

      A. Winand, P. Tarte, J. Mater. Sci. 25 (1990) 4008–4013.

  84. [84]

      T. Miyajim, J. Tamaki, M. Matsuoka, et al., Solid State Ionics 124 (1999) 201–211.

  85. [85]

      K. Koji Kawada, T. Okura, Funct. Mater. Lett. 14 (2021) 2141001.

  86. [86]

      M. Guin, F. Tietz, J. Power Sources 273 (2015) 1056–1064.

  87. [87]

      Q. Ma, M. Guin, S. Naqash, et al., Chem. Mater. 28 (2016) 4821–4828. doi: 10.1021/acs.chemmater.6b02059

  88. [88]

      M. Samiee, B. Radhakrishnan, Z. Rice, et al., J. Power Sources 347 (2017) 229–237.

  89. [89]

      A. Jolley, G. Cohn, G. Hitz, et al., Ionics 21 (2015) 3031–3038. doi: 10.1007/s11581-015-1498-8

  90. [90]

      L. Liu, X. Qi, Y. Shao, et al., Energy Storage Sci. Techonol. 6 (2017) 961–980.

  91. [91]

      S. He, Y. Xu, X. Ma, et al., ChemElectroChem 7 (2020) 2087–2094. doi: 10.1002/celc.201902052

  92. [92]

      S. Pal, R. Saha, G. Kumar, et al., J. Phys. Chem. C 124 (2020) 9161–9169. doi: 10.1021/acs.jpcc.0c00543

  93. [93]

      Y. Jing, G. Liu, M. Avdeev, et al., ACS Energy Lett. 5 (2020) 2835–2841.

  94. [94]

      R. Fuentes, F. Figueiredo, M. Soares, et al., J. Eur. Ceram. Soc. 25 (2005) 455–462.

  95. [95]

      A. Ignaszak, P. Pasierb, R. Gajerski, et al., Thermochim. Acta 426 (2005) 7–14.

  96. [96]

      B. Yan, L. Kang, M. Kotobuki, et al., Mater. Technol. 34 (2018) 356–360.

  97. [97]

      F. Ejehi, S. Marashi, M. Ghaani, et al., Ceram. Int. 38 (2012) 6857–6863.

  98. [98]

      H. Leng, J. Huang, J. Nie, et al., J. Power Sources 391 (2018) 170–179.

  99. [99]

      S. Liu, C. Zhou, Y. Wang, et al., ACS Appl. Mater. Interfaces 12 (2020) 3502–3509. doi: 10.1021/acsami.9b11995

  100. [100]

       J. Oh, L. He, A. Plewa, et al., ACS Appl. Mater. Interfaces 11 (2019) 40125–40133. doi: 10.1021/acsami.9b14986

  101. [101]

       Y. Shao, G. Zhong, Y. Lu, et al., Energy Storage Mater. 23 (2019) 514–521.

  102. [102]

       Y. Li, Z. Deng, J. Peng, et al., Chem. Eur. J. 24 (2018) 1057–1061. doi: 10.1002/chem.201705466

  103. [103]

       J. Wu, Q. Wang, X. Guo, J. Power Sources 402 (2018) 513–518.

  104. [104]

       Z. Deng, J. Gu, Y. Li, et al., Electrochim. Acta 298 (2019) 121–126.

  105. [105]

       R. Smaha, J. Roudebush, J. Herb, et al., Inorg. Chem. 54 (2015) 7985–7991. doi: 10.1021/acs.inorgchem.5b01186

  106. [106]

       Y. Wang, Q. Wang, Z. Liu, et al., J. Power Sources 293 (2015) 735–740.

  107. [107]

       P. Hong Fang, ACS Appl. Mater. Interfaces 11 (2019) 963–972.

  108. [108]

       Y. Sun, Y. Wang, X. Liang, et al., J. Am. Chem. Soc. 141 (2019) 5640–5644. doi: 10.1021/jacs.9b01746

  109. [109]

       H. Zhang. Lei Gao, Y. Wang, et al., J. Mater. Chem. A 8 (2020) 21265–21272.

  110. [110]

       T. Wan, Z. Lu, Ciucci F, J. Power Sources 390 (2018) 61–70.

  111. [111]

       M. Clarke. B. Goldmann, J. Dawson, et al., J. Mater. Chem. A 10 (2022) 2249–2255.

  112. [112]

       C. Cazorla, D. Errandonea, Phys. Rev. Lett. 113 (2014) 235902. doi: 10.1103/PhysRevLett.113.235902

  113. [113]

       A. Barriocanal, M. Varela, Z. Sefrioui, et al., Science 321 (2008) 676–680.

  114. [114]

       Y. Wang, T. Wen, C. Park, et al., J. Appl. Phys. 119 (2016) 025901.

  115. [115]

       N. Tanibata, K. Noi, A. Hayashi, et al., ChemElectroChem 1 (2014) 1130–1132. doi: 10.1002/celc.201402016

  116. [116]

       A. Hayashi, N. Masuzawa, S. Yubuchi, et al., Nat. Commun. 10 (2019) 5266.

  117. [117]

       A. Banerjee, K.H. Park, J.W. Heo, et al., Angew. Chem. Int. Ed. 55 (2016) 9634–9638. doi: 10.1002/anie.201604158

  118. [118]

       L. Zhang, D. Zhang, K. Yang, et al., Adv. Sci. 3 (2016) 1600089.

  119. [119]

       Z. Yu, S.L. Shang, J.H. Seo, et al., Adv. Mater. 29 (2017) 1605561.

  120. [120]

       W. Weng, G. Liu, Y. Li, et al., Appl. Mater. Today 27 (2022) 101448.

  121. [121]

       L. Zhang, K. Yang, J. Mi, et al., Adv. Energy Mater. 5 (2015) 1501294.

  122. [122]

       S. Bo, Y. Wang, J.C. Kim, et al., Chem. Mater. 28 (2015) 252–258.

  123. [123]

       I. Chu, C. Kompella, H. Nguyen, et al., Sci. Rep. 6 (2016) 33733.

  124. [124]

       C. Moon, H. Lee, K. Park, et al., ACS Energy Lett. 3 (2018) 2504–2512. doi: 10.1021/acsenergylett.8b01479

  125. [125]

       M. Duchardt, S. Neuberger, U. Ruschewitz, et al., Chem. Mater. 30 (2018) 4134–4139. doi: 10.1021/acs.chemmater.8b01656

  126. [126]

       H. Jia, X. Liang, T. An, et al., Chem. Mater. 32 (2020) 4065–4071. doi: 10.1021/acs.chemmater.0c00872

  127. [127]

       S. Shang, Z. Yu, Y. Wang, et al., ACS Appl. Mater. Interfaces 9 (2017) 16261–16269. doi: 10.1021/acsami.7b03606

  128. [128]

       M. Meyer, Z. Anorg, Allg. Chem. 621 (1995) 457–463.

  129. [129]

       T. Asano, A. Sakai, S. Ouchi, et al., Adv. Mater. 30 (2018) e1803075.

  130. [130]

       S. Wang, Q. Bai, A.M. Nolan, et al., Angew. Chem. Int. Ed. 58 (2019) 8039–8043. doi: 10.1002/anie.201901938

  131. [131]

       M. Gombotz, H.M.R. Wilkening, ACS Sustain. Chem. Eng. 9 (2020) 743–755.

  132. [132]

       R. Schlem, A. Banik, S. Ohno, et al., Chem. Mater. 33 (2021) 327–337. doi: 10.1021/acs.chemmater.0c04352

  133. [133]

       S. Chen, C. Yu, C. Wei, et al., Energy Mater. Adv. 4 (2023) 1–10.

  134. [134]

       E. Wu, S. Banerjee, H. Tang, et al., Nat. Commun. 12 (2021) 1256.

  135. [135]

       L. Li, J. Yao, R. Xu, et al., Energy Storage Mater. 63 (2023) 103016.

  136. [136]

       X. Xu, Y. Li, X. Wang, et al., J. Solid State Electrochem. 28 (2024) 3501–3507. doi: 10.1007/s10008-024-05838-1

  137. [137]

       J. Fu, S. Wang, D. Wu, et al., Adv. Mater. 36 (2024) 2308012.

  138. [138]

       Y. Hu, J. Fu, J. Xu, et al., Matter 7 (2024) 1018–1034.

  139. [139]

       W. Tang, A. Unemoto, W. Zhou, et al., Energy Environ. Sci. 8 (2015) 3637–3645.

  140. [140]

       L. Duchene, R.S. Kuhnel, D. Rentsch, et al., Chem. Commun. 53 (2017) 4195–4198.

  141. [141]

       W. Tang, M. Matsuo, H. Wu, et al., Energy Storage Mater. 4 (2016) 79–83.

  142. [142]

       X. Luo, A. Rawal, M.S. Salman, et al., ACS Appl. Nano Mater. 5 (2022) 373–379. doi: 10.1021/acsanm.1c03187

  143. [143]

       J. Oh, L. He, B. Chua, et al., Energy Storage Mater. 34 (2021) 28–44.

  144. [144]

       D. Reed, G. Coffey, E. Mast, et al., J. Power Sources 227 (2013) 94–100.

  145. [145]

       I. Kim, J. Park, C. Kim, et al., J. Power Sources 301 (2016) 332–337.

  146. [146]

       G. Zhang, Z. Wen, X. Wu, et al., J. Alloys Compd. 613 (2014) 80–86.

  147. [147]

       C. Zhao, L. Liu, X. Qi, et al., Adv. Energy Mater. 8 (2018) 1703012.

  148. [148]

       K. Zhao, Y. Liu, S. Zhang, et al., Electrochem. Commun. 69 (2016) 59–63.

  149. [149]

       H. Yang, B. Zhang, K. Konstantinov, et al., Adv. Energy Sustain. Res. 2 (2021) 2000057.

  150. [150]

       H. Lai, J. Wang, M. Cai, et al., Chem. Engin. J. 433 (2022) 133545.

  151. [151]

       M. Bay, M. Wang, R. Grissa, et al., Adv. Energy Mater. 10 (2019) 1902899.

  152. [152]

       L. Liu, X. Qi, Q. Ma, et al., ACS Appl. Mater. Interfaces 8 (2016) 32631–32636. doi: 10.1021/acsami.6b11773

  153. [153]

       P. Kehne, C. Guhl, L. Alff, et al., Solid State Ionics 341 (2019) 115041.

  154. [154]

       C. Li, R. Li, K. Liu, et al., Interdisciplinary Mater. 1 (2022) 396–416. doi: 10.1002/idm2.12044

  155. [155]

       Z. Zhang, Q. Zhang, J. Shi, et al., Adv. Energy Mater. 7 (2017) 1601196.

  156. [156]

       Y. Li, M. Li, Z. Sun, et al., Energy Storage Mater. 56 (2023) 582–599.

  157. [157]

       H. Gao, L. Xue, S. Xin, et al., Angew. Chem. Int. Ed. 56 (2017) 5541–5545. doi: 10.1002/anie.201702003

  158. [158]

       W. Zhou, Y. Li, S. Xin, et al., ACS Cent. Sci. 3 (2017) 52–57. doi: 10.1021/acscentsci.6b00321

  159. [159]

       L. Ran, M. Li, E. Cooper, et al., Energy Storage Mater. 41 (2021) 8–13.

  160. [160]

       X. Miao, H. Di, X. Ge, et al., Energy Storage Mater. 30 (2020) 170–178.

  161. [161]

       X. Wang, J. Chen, Z. Mao, et al., J. Mater. Chem. A 9 (2021) 16039–16045. doi: 10.1039/d1ta04869d

  162. [162]

       H. Fu, Q. Yin, Y. Huang, et al., ACS Mater. Lett. 2 (2019) 127–132. doi: 10.1109/icme.2019.00030

  163. [163]

       X. Yu, Y. Yao, X. Wang, et al., Energy Storage Mater. 54 (2023) 221–226.

  164. [164]

       Q. Ni, Y. Xiong, Z. Sun, et al., Adv. Energy Mater. 13 (2023) 2300271.

  165. [165]

       Y. Lu, J.A. Alonso, Q. Yi, et al., Adv. Energy Mater. 9 (2019) 1901205.

  166. [166]

       C. Wang, Z. Sun, Y. Zhao, et al., Small 17 (2021) e2103819.

  167. [167]

       Y. Zhao, C. Wang, Y. Dai, et al., Nano Energy 88 (2021) 106293.

  168. [168]

       D. Li, C. Sun, C. Wang, et al., Energy Storage Mater. 54 (2023) 403–409.

  169. [169]

       J. Yang, G. Liu, M. Avdeev, et al., ACS Energy Lett. 5 (2020) 2835–2841. doi: 10.1021/acsenergylett.0c01432

  170. [170]

       Z. Jiang, S. Chen, C. Wei, et al., Chin. Chem. Lett. 35 (2024) 108561.

  171. [171]

       J. Liang, X. Li, C. Wang, et al., Energy Mater. Adv. 4 (2023) 1–14.

  172. [172]

       Q. Luo, L. Ming, D. Zhang, et al., Energy Mater. Adv. 4 (2023) 1–12. doi: 10.53819/81018102t4181

  173. [173]

       L. Ming, D. Liu, Q. Luo, et al., Chin. Chem. Lett. (2023) 109087.

  174. [174]

       B. Tang, P. Jaschin, X. Li, et al., Mater. Today 41 (2020) 200–218.

  175. [175]

       J. Yue, F. Han, X. Fan, et al., ACS Nano 11 (2017) 4885–4891. doi: 10.1021/acsnano.7b01445

  176. [176]

       R. Rao, H. Chen, L. Wong, et al., J. Mater. Chem. A 5 (2017) 3377–3388.

  177. [177]

       Z. Yu, S.L. Shang, J. Seo, et al., Adv. Mater. 29 (2017) 1605561.

  178. [178]

       Q. Liu, X. Zhao, Q. Yang, et al., Adv. Mater. Technol. 8 (2023) 2200822.

  179. [179]

       Z. Zhang, H. Cao, M. Yang, et al., J. Energy Chem. 48 (2020) 250–258.

  180. [180]

       Y. Nagata, K. Nagao, M. Deguchi, et al., Chem. Mater. 30 (2018) 6998–7004. doi: 10.1021/acs.chemmater.8b01872

  181. [181]

       F. Hao, X. Chi, Y. Liang, et al., Joule 3 (2019) 1349–1359.

  182. [182]

       H. Tang, Z. Deng, Z. Lin, et al., Chem. Mater. 30 (2017) 163–173.

  183. [183]

       Y. Li, W. Arnold, S. Halacoglu, et al., Adv. Funct. Mater. (2021) 31.

  184. [184]

       Y. Tian, Y. Sun, D.C. Hannah, et al., Joule 3 (2019) 1037–1050.

  185. [185]

       Z. Wang, L. Zhang, X. Shang, et al., Chem. Eng. J. 428 (2022) 123094.

  186. [186]

       L. Li, R. Xu, L. Zhang, et al., Adv. Funct. Mater. 32 (2022) 2203095.

  187. [187]

       H. Wan, J.P. Mwizerwa, X. Qi, et al., ACS Nano 12 (2018) 2809–2817. doi: 10.1021/acsnano.8b00073

  188. [188]

       H. Wan, J.P. Mwizerwa, X. Qi, et al., ACS Appl. Mater Interfaces 10 (2018) 12300–12304. doi: 10.1021/acsami.8b01805

  189. [189]

       J. Yue, X. Zhu, F. Han, et al., ACS Appl. Mater Interfaces 10 (2018) 39645–39650. doi: 10.1021/acsami.8b12610

  190. [190]

       X. Chi, Y. Liang, F. Hao, et al., Angew. Chem. Int. Ed. 57 (2018) 2630–2634. doi: 10.1002/anie.201712895

  191. [191]

       T. An, H. Jia, L. Peng, et al., ACS Appl. Mater. Interfaces 12 (2020) 20563–20569. doi: 10.1021/acsami.0c03899

  192. [192]

       Z. Zhang, Z. Wang, L. Zhang, et al., Adv. Sci. 9 (2022) e2200744.

  193. [193]

       J. Park, J.P. Son, W. Ko, et al., ACS Energy Lett. 7 (2022) 3293–3301. doi: 10.1021/acsenergylett.2c01514

  194. [194]

       X. Chi, Y. Zhang, F. Hao, et al., Nat. Commun. 13 (2022) 2854.

  195. [195]

       M. Braga, N. Grundish, A. Murchison, et al., Energy Environ. Sci. 10 (2017) 331–336.

  196. [196]

       R. Jiang, C. Song, J. Yang, et al., Adv. Funct. Mater. 33 (2023) 2301635.

  197. [197]

       Y. Pang, Y. Liu, J. Yang, et al., Mater. Today Nano 18 (2022) 100194.

  198. [198]

       L. de Kort, O. Brandt Corstius, V. Gulino, et al., Adv. Funct. Mater. 33 (2023) 2209122.

  199. [199]

       J. Cuan, Y. Zhou, T. Zhou, et al., Adv. Mater. 31 (2019) e1803533.

  200. [200]

       D. Souza, A. D'Angelo, T. Humphries, et al., Dalton Trans. 51 (2022) 13848–13857. doi: 10.1039/d2dt01943d

  201. [201]

       W. Tang, K. Yoshida, A. Soloninin, et al., ACS Energy Lett. 1 (2016) 659– 664. doi: 10.1021/acsenergylett.6b00310

  202. [202]

       T.K. Yoshida, A. Unemoto, M. Matsuo, et al., Appl. Phys. Lett. 110 (2017) 103901.

  203. [203]

       L. Duchêne, R. Kühnel, E. Stilp, et al., Energy Environ. Sci. 10 (2017) 2609–2615.

  204. [204]

       L. Duchene, D.H. Kim, Y.B. Song, et al., Energy Storage Mater. 26 (2020) 543–549.

  205. [205]

       F. Murgia, M. Brighi, R. Černý, Electrochem. Commun. 106 (2019) 106534.

  206. [206]

       M. Brighi, F. Murgia, Z. Łodziana, et al., J. Power Sources 404 (2018) 7–12.

  207. [207]

       L. He, H. Lin, H. Li, et al., J. Power Sources 396 (2018) 574–579.

  208. [208]

       R. Asakura, D. Reber, L. Duchêne, et al., Energy Environ. Sci. 13 (2020) 5048–5058. doi: 10.1039/d0ee01569e

  209. [209]

       M. Jin, S. Cheng, Z. Yang, et al., Chem. Engin. J. 455 (2023) 140904.

  210. [210]

       L. Duchêne, D.H. Kim, Y.B. Song, et al., Energy Storage Mater. 26 (2020) 543–549.

  211. [211]

       K. Niitani, S. Ushiroda, H. Kuwata, et al., ACS Energy Lett. 7 (2021) 145–149.

  212. [212]

       J. Fu, S. Wang, D. Wu, et al., Adv. Mater. 36 (2024) e2308012.

  213. [213]

       L. Geng, C. Zhao, J. Yan, et al., J. Mater. Chem. A 10 (2022) 14875–14883. doi: 10.1039/d2ta02513b

  214. [214]

       L. Geng, D. Xue, J. Yao, et al., Energy Environ. Sci. 16 (2023) 2658–2668. doi: 10.1039/d3ee00237c

  215. [215]

       Q. Liu, L. Zhang, H. Sun, et al., ACS Energy Lett. 5 (2020) 2546–2559. doi: 10.1021/acsenergylett.0c01214

  216. [216]

       Y. Li, Q. Liu, S. Wu, et al., J. Am. Chem. Soc. 19 (2023) 10576–10583. doi: 10.1021/jacs.2c13589

  217. [217]

       G. Rees, D. Spencer Jolly, Z. Ning, et al., Angew. Chem. Int. Ed. 60 (2021) 2110–2115. doi: 10.1002/anie.202013066

• 

  Figure 1  Schematic diagram of four typical ion diffusion mechanisms of NaF, purple represents F ions, yellow represents occupied sodium ion sites, gray represents sodium ion vacancies, and green represents interstitial sites. (a) Direct hopping mechanisms between vacancies; (b) Direct hopping mechanisms between interstitials; (c) "Knock-off" like mechanism between vacancies and interstitials; (d) Conceptual diagram of the energy barriers required for single ion diffusion and multiple ions to diffuse in concert.

  下载: 全尺寸图片 幻灯片
  

  Figure 2  Crystal structure of Na-β/β″-Al₂O₃. The yellow ball represents the sodium ion, the red represents the oxygen ion, and the gray represents the aluminum ion.

  下载: 全尺寸图片 幻灯片
  

  Figure 3  Structures of Nasicon, different yellow balls represent different sites of sodium ions, red represents oxygen ion, blue octahedron represents [ZrO₆] octahedron, pink tetrahedron represents [Si/PO₄] tetrahedron: (a) Rhombohedral phase; (b) Monoclinic phase; (c) Na1-Na2 transport channel of rhombohedral phase; (d) Na1-Na2 transport channel and (e) Na1-Na3 transport channel of monoclinic phase.

  下载: 全尺寸图片 幻灯片
  

  Figure 4  Crystal structures of (a) Na₂Ni₂TeO₆ and (b) Na₂M₂TeO₆ (M = Zn, Co, Mg). Different yellow balls represent different sites of sodium ions, the green octahedron represents [TeO₆], the purple octahedron represents [NiO₆], and the pink octahedron represents [ZnO₆].

  下载: 全尺寸图片 幻灯片
  

  Figure 5  Plan view of Na₃OCl along the [111] direction. Blue ball represents oxygen ion, green represents chloride ion, yellow represents sodium ion.

  下载: 全尺寸图片 幻灯片
  

  Figure 6  Crystal structures of Na₃PS₄. (a) Tetragonal and cubic phase, (b) crystal structure of Na₁₁Sn₂PS₁₂. The blue ball represents tin ions, pink represents phosphorus ions, green represents sulfur ions, and yellow represents sodium ions.

  下载: 全尺寸图片 幻灯片
  

  Figure 7  Crystal structure of Na₃YCl₆. Different yellow balls represent different sites of sodium ions, the gray color represents yttrium ions and the green color represents chloride ions.

  下载: 全尺寸图片 幻灯片
  

  Figure 8  Crystal structures of (a) NaAlH₄, (b) Na₃AlH₄, (c) low-temperature ordered monoclinic phase and (d) high-temperature disordered cubic phase. Reprinted with permission [65]. Copyright 2014, Royal Society Chemistry.

  下载: 全尺寸图片 幻灯片
  

  Figure 9  Schematic diagram of synthesis of (a) solid state reaction and (b) liquid phase methods.

  下载: 全尺寸图片 幻灯片
  

  Figure 10  Reported conductivities of Na-ion SSEs.

  下载: 全尺寸图片 幻灯片
  

  Figure 11  (a) Schematic diagram of using Na₂Si₂O₃ as NZSP sintering additive and (b) ionic conductivity and relative density of these samples. Reprinted with permission [100]. Copyright 2019, American Chemical Society. (c) Ionic conductivity and E_a of samples with different levels of NaF. Reprinted with permission [101]. Copyright 2019, Elsevier.

  下载: 全尺寸图片 幻灯片
  

  Figure 12  (a) Transformation of partial crystal structure before and after Ga substitution in NZTO. Reprinted with permission [102]. Copyright 2018, Elsevier. (b) Total ionic conductivities of Na_1.95Zn_1.95Ga_0.05TeO₆ and Na₂Zn₂TeO₆ and (c) the σ_bulk and σ_gb of two samples at different temperatures Reprinted with permission [103]. Copyright 2018, Elsevier.

  下载: 全尺寸图片 幻灯片
  

  Figure 13  (a) Crystal structure of Na₃OX (X = Cl, Br, I) and (b) Arrhenius plots for Na₃OCl, Na₃OBr, Na₃OBr_0.6I_0.4 and Na_2.9Sr_0.05OBr_0.6I_0.4 samples. Reprinted with permission [106]. Copyright 2015, Elsevier. (c) Crystal structure of Na₃OBD₄ and (d) Nyquist plots of hot-pressed Na₃OBH₄ pellet. Reprinted with permission [108]. Copyright 2019, American Chemical Society.

  下载: 全尺寸图片 幻灯片
  

  Figure 14  (a) Raman spectra and (b) XRD patterns of Pristine Na₃SbS₄·9H₂O, as-prepared Na₃SbS₄, air-exposed Na₃SbS₄ and reheated air-exposed Na₃SbS₄. Reprinted with permission [33]. Copyright 2016, Wiley Online Library.

  下载: 全尺寸图片 幻灯片
  

  Figure 15  (a) Structure diagram of B₁₂H₁₂²⁻ and CB₁₁H₁₂⁻. (b) The ionic conductivities of Li⁺ and Na⁺ species in LiCB₁₁H₁₂ and NaCB₁₁H₁₂ were measured as a function of inverse temperature. Reprinted with permission [139]. Copyright 2015, RSC publishing. (c) The ionic conductivity of pristine and ball-milled Na₂B₁₂H₁₂ was compared after QENS measurements. Reprinted with permission [141]. Copyright 2016, Elsevier. (d) A description of the synthetic route of NaBH₄@Na₂B₁₂H₁₂. Reprinted with permission [142]. Copyright 2022, American Chemical Society.

  下载: 全尺寸图片 幻灯片
  

  Figure 16  (a) Cycling performance of the SIBs based on the Na-β″-Al₂O₃ thin film electrolyte. Reprinted with permission [148]. Copyright 2016, Elsevier. (b) Schematic illustration of Na₃V₂(PO₄)₃/SC-treated-β″-Al₂O₃/Na solid-state batteries and corresponding electrochemical performance. Reprinted with permission [150]. Copyright 2022, Elsevier.

  下载: 全尺寸图片 幻灯片
  

  Figure 17  (a) XPS spectra and surface composition analysis results of Na-β″-Al₂O₃ surfaces before and after heat treatment. Reprinted with permission [151]. Copyright 2019, Wiley Online Library. (b) Schematic illustration of Na_0.66Ni_0.33Mn_0.67O₂/β″-Al₂O₃/Na solid-state batteries with ionic liquid added in the cathode side and corresponding electrochemical performance. Reprinted with permission [152]. Copyright 2016, American Chemical Society. (c) Schematic illustration of solid-state batteries with liquid electrolyte added in the both sides of the interface between the cathode and anode electrodes and corresponding electrochemical performance. Reprinted with permission [145]. Copyright 2016, Elsevier.

  下载: 全尺寸图片 幻灯片
  

  Figure 18  (a) Schematic diagram of Na₃V₂(PO₄)₃/IL/Na_3.3Zr_1.7La_0.3Si₂PO₁₂/Na solid-state batteries and corresponding electrochemical performance. Reprinted with permission [155]. Copyright 2017, Wiley Online Library. (b) Schematic illustration of solid-state batteries with plastic-crystal electrolyte applied in the cathode and corresponding electrochemical performance. Reprinted with permission [157]. Copyright 2017, Wiley Online Library.

  下载: 全尺寸图片 幻灯片
  

  Figure 19  (a) Schematic illustration of solid-state batteries with CMPEA added in the anode side and corresponding electrochemical performance. Reprinted with permission [158]. Copyright 2017, American Chemical Society. (b) Schematic illustration of solid-state sodium ion batteries with sandwich composite electrolyte and corresponding electrochemical performance. Reprinted with permission [159]. Copyright 2021, Elsevier.

  下载: 全尺寸图片 幻灯片
  

  Figure 20  (a) Schematic illustration of the AlF₃-modified Nasicon based cells and corresponding electrochemical performance. Reprinted with permission [160]. Copyright 2020, Elsevier. (b) Schematic illustration of the SnF₂-modified Nasicon based cells and corresponding electrochemical performance. Reprinted with permission [161]. Copyright 2021, Royal Society Chemistry. (c) Schematic illustration of the SiO₂-modified Nasicon based cells and corresponding electrochemical performance. Reprinted with permission [162]. Copyright 2020, American Chemical Society.

  下载: 全尺寸图片 幻灯片
  

  Figure 21  (a) TEM images and corresponding elemental mappings of the NPS-nano-Na₂S-C nanocomposite cathode and electrochemical performance of the Na-Sn-C/Na₃PS₄/Na₃PS_4—Na₂S-C solid-state batteries. Reprinted with permission [175]. Copyright 2017, American Chemical Society. (b) Cycling performance of the Na_2+2δFe_2-δ(SO₄)₃/Na_3.1Sn_0.1P_0.9S₄/Na₂Ti₃O₇ solid-state batteries at different temperatures. Reprinted with permission [176]. Copyright 2017, Royal Society Chemistry. (c) The quantity of H₂S gas produced from Na₁₀SnSb₂S₁₂ exposed to humid air for different time durations and electrochemical performance of the TiS₂/Na₁₀SnSb₂S₁₂/Na solid-state batteries. Reprinted with permission [14]. Copyright 2021, Elsevier.

  下载: 全尺寸图片 幻灯片
  

  Figure 22  (a) Schematic illustration of NaCrO₂+NYZC/Na₃PS₄/Na₂Sn solid-state batteries and corresponding electrochemical performance. Reprinted with permission [134]. Copyright 2021, Springer Nature. (b) Schematic illustration of SeSPAN-Ar@NSS/Na₃SbS₄/Na₁₅Sn₄ solid-state batteries and corresponding electrochemical performance. Reprinted with permission [179]. Copyright 2020, Elsevier. (c) Schematic illustration of SIBs with organic cathode and corresponding electrochemical performance. Reprinted with permission [180,181]. Copyright 2018, American Chemical Society. Copyright 2019, Cell Press.

  下载: 全尺寸图片 幻灯片
  

  Figure 23  (a) Schematic illustration of solid-state batteries with a phase-transition copolymer (PPP-NaTFSI) as the interlayer in the anode side and corresponding electrochemical performance. Reprinted with permission [183]. Copyright 2021, Wiley Online Library. (b) Schematic illustration of Na₃SbS₄ electrolyte with surface hydrate coating and corresponding SXRD profiles of Na/Na₃SbS₄/Na symmetric cell. Reprinted with permission [184]. Copyright 2019, Elsevier. (c) Schematic illustration of solid-state SIBs with PVDF membrane with sub-micron vertical channels and corresponding electrochemical performance. Reprinted with permission [185]. Copyright 2022, Elsevier.

  下载: 全尺寸图片 幻灯片
  

  Figure 24  (a) Mechanism schematic illustration of the Se_0.8S_7.2@PAN/NaPSO+NSS/Na₁₅Sn₄ solid-state batteries. (b) Raman spectra, SEM image and Ionic conductivity of NaPSO+NSS. (c) electrochemical performance of the Se_0.8S_7.2@PAN/NaPSO+NSS/Na₁₅Sn₄ solid-state batteries. (d) Cross-sectional SEM images and EDS mapping for the Se_0.8S_7.2@PAN/NaPSO+NSS/Na₁₅Sn₄ and corresponding electrochemical performance. Reprinted with permission [186]. Copyright 2022, Wiley Online Library.

  下载: 全尺寸图片 幻灯片
  

  Figure 25  (a) Cycling performance of the Na/ferrocene cells assembled with Na_3-xH_xOCl. Reprinted with permission [195]. Copyright 2017, Royal Society Chemistry. (b) Schematic illustration of Fe[Fe(CN)₆]₃/Na_2.98Mg_0.01SO₄F_0.95Cl_0.05/Na-Sn and corresponding electrochemical performance. Reprinted with permission [54]. Copyright 2020, Elsevier. (c) Electrochemical performance of NaSn|Na₂BH₄NH₂|NaSn symmetric cell and TiS₂|Na₂BH₄NH₂|NaSn solid-state batteries. Reprinted with permission [196]. Copyright 2023, Wiley Online Library.

  下载: 全尺寸图片 幻灯片
  

  Figure 26  (a) Schematic illustration of NaCrO₂|Na₂(B₁₂H₁₂)_0.5(B₁₀H₁₀)_0.5|Na solid-state batteries and corresponding electrochemical performance. Reprinted with permission [203]. Copyright 2013, Royal Society Chemistry. (b) Schematic illustration of solid-state batteries with a NaCrO₂ electrode infiltrated with Na₄(B₁₂H₁₂)(B₁₀H₁₀) and corresponding electrochemical performance. Reprinted with permission [204]. Copyright 2020, Elsevier.

  下载: 全尺寸图片 幻灯片
  

  Figure 27  (a) Electrochemical performance of NaCrO₂/Na₄(CB₁₁H₁₂)₂(B₁₂H₁₂)/Na solid-state batteries. Reprinted with permission [205]. Copyright 2019, Elsevier. (b) Electrochemical performance of NaSn/Na_4/3(CB₁₁H₁₂)_2/3(B₁₂H₁₂)_1/3/NaSn symmetric cell. Reprinted with permission [206]. Copyright 2018, Elsevier. (c) CV test results for Na/Na₃NH₂B₁₂H₁₂/Pt and electrochemical performance of TiS₂/Na₃NH₂B₁₂H₁₂/Na solid-state batteries. Reprinted with permission [207]. Copyright 2018, Elsevier.

  下载: 全尺寸图片 幻灯片
  

  Figure 28  (a) Schematic illustration of Na₃(VOPO₄)₂F/Na₄(CB₁₁H₁₂)₂(B₁₂H₁₂)/Na solid-state batteries and corresponding electrochemical performance. Reprinted with permission [208]. Copyright 2020, Royal Society Chemistry. (b) Schematic illustration of Na[Ni_1/3Fe_1/3Mn_1/3]O₂/Na₃B₂₄H₂₃–5Na₂B₁₂H₁₂/Na solid-state sodium batteries and corresponding electrochemical performance. Reprinted with permission [209]. Copyright 2023, Elsevier.

  下载: 全尺寸图片 幻灯片
  

  Figure 29  (a) Schematic diagram illustrating the structural changes in the composite cathode of NaCrO_2—NYbZrCl_0.75/Na₃PS₄/Na₂Sn solid-state batteries and corresponding electrochemical performance. Reprinted with permission [135]. Copyright 2023, Elsevier. (b) Schematic illustration of NVP/NaTaCl₆/Na₃PS₄/Na₁₅Sn₄ solid-state sodium batteries and corresponding electrochemical performance [138]. Copyright 2023, Cell Press.

  下载: 全尺寸图片 幻灯片

  Table 1.  Summary of solid-state SIBs using oxide-based SSEs.

  下载: 导出CSV

  Table 2.  Summary of solid-state SIBs using sulfide-based SSEs.

  下载: 导出CSV

  Table 3.  Summary of solid-state SIBs using complex hydrides SSEs.

  下载: 导出CSV
•