English

• 1. Introduction

  The demand for energy storage and conversion technologies continues to increase as the ever-growing intelligent technologies in modern society [1]. Rechargeable batteries, an efficient energy storage and conversion device, have become an indispensable energy storage and supply equipment. At present, lithium-ion batteries (LIBs) are considered as promising rechargeable energy storage technologies [2]. LIBs have greatly improved the level of energy utilization in humans and driven the evolution of information technology and electronics, so three scientists who made pioneering contributions to the development of LIBs were awarded the Nobel Prize for Chemistry in 2019 [3].

  Over the past three decades, LIBs have rapidly become the most successful rechargeable batteries. To develop energy storage systems based on high energy density, more electrode materials with high capacity are constantly explored and investigated, such as nickel-rich cathodes, lithium-rich layered oxides, organic electrodes, silicon anodes, and phosphorus anodes [4–9]. However, the growing demand in diverse application scenarios and limited lithium resources are calling for more novel rechargeable batteries. According to the requirements of energy density, safety and resource richness in different application environments, next-generation rechargeable batteries such as lithium metal batteries, sodium-ion batteries, solid-state batteries, aqueous batteries and flow batteries are experiencing rapid development [10–14]. Lithium metal batteries are the attractive energy storage devices due to the low redox potential (−3.04 V vs. SHE) and high specific capacity (3860 mAh/g) of lithium metal anodes [15–17]. Lithium-sulfur (Li-S) batteries, lithium-air (Li-O₂ or Li-CO₂) batteries and solid-state batteries using lithium metal anodes are expected to have a much higher specific capacity than traditional LIBs [18–20]. With the goal of sustainable energy storage devices, rechargeable batteries based on alkali metal (e.g., Na⁺, K⁺) and polyvalent metal (e.g., Mg²⁺, Zn²⁺, Al³⁺) cations have attracted attention due to their abundant resources, among which sodium-ion batteries have made the most progress [21–25]. In addition, organic electrode materials and low-cost materials receive considerable attention due to the urgent need for sustainable electrode materials and energy storage systems with high energy density [26]. Aqueous batteries with organic electrode materials are more attractive in view of their inherent safety and potential sustainable production capability [27]. Besides these aqueous rechargeable batteries with solid-state electrodes, redox flow batteries using liquid anolytes and catholytes as active species are also one of effective candidates for energy storage, which rely on the redox reaction of soluble molecules or ions [28]. In general, while LIBs dominate the energy market, the development of beyond traditional electrochemical energy based on different application environments has become a hot spot in the research of rechargeable batteries (Fig. 1).

  Figure 1

  

  Figure 1.  Research advances in the field of rechargeable batteries.

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  This review provides a summary of the current important research advances in the field of rechargeable batteries, covering various types of rechargeable batteries and recycling spent power batteries, as well as the state-of-the-art characterization techniques for battery research. By organizing this review, we hope to provide original insights into current energy storage research and lead the future of rechargeable batteries.

  2. Lithium-ion batteries

  2.1 Cathode materials

  2.1.1   LiCoO₂ cathodes

  In 1980, Goodenough et al. first reported the application of layered transition metal oxide LiCoO₂ (LCO) as a cathode material for LIBs, paving the way for the development of new cathode materials [29]. In 1991, LCO played a key role in enabling the first commercial application of LIBs by SONY [30]. During the subsequent two decades, the advancement of LCO was comparatively sluggish. The charge cutoff voltage of the earliest commercial LCO cathode was around 4.2 V, with a specific capacity of approximately 146 mAh/g, which is significantly lower than the theoretical specific capacity (274 mAh/g) [31]. In 2013, LCO with a charge cutoff voltage of 4.35 V came into the market, achieving a specific capacity of 165 mAh/g, thus marking the beginning of a rapid development phase [32]. It is found that when the charge cutoff voltage varies within the range of 4.2–4.5 V, each 0.1 V increase leads to a 10% rise in volume energy density. Significantly, when the charge cutoff voltage ascends from 4.5 V to 4.6 V, the volume energy density of LCO experiences a remarkable increase from 3081 Wh/L to 3721 Wh/L, representing a gain of over 20% [33]. Therefore, the development of 4.6 V (vs. Li⁺/Li) high-voltage LCO has become a major research focus. However, LCO faces more significant challenges under high-voltage conditions, leading to rapid structural deterioration and degraded cycle performance, which hinder the commercialization of 4.6 V high-voltage LCO.

  Firstly, when the charge cutoff voltage exceeds 4.55 V, an adverse phase transition from the O3 phase to the H1–3 phase occurs, resulting in severe lattice contraction and microcracking within the LCO particles. Additionally, the lattice oxgen will be oxidized during deep delithiation due to the overlap between the Co 3d and O 2p orbitals, which is particularly pronounced at the surface, resulting in significant O release, Co dissolution, and irreversible structural degradations [34,35]. The release of O and dissolution of Co in the surface structure result in the formation of a disordered phase with poor ionic conductivity. Consequently, the transport of Li⁺ ions in the surface structure is severely hindered, leading to rapid capacity attenuation. In recent years, there has been increasing research into the involvement of lattice oxygen in redox reactions in LCO [36,37]. At high voltages, the highly oxidative lattice Co⁴⁺ and O^α- (α < 2) on LCO surface trigger more pronounced side reactions, leading to the continuous thickening of the cathode electrolyte interface (CEI) layer with poor ion transmission, as well as the depletion of electrolyte [38]. The structural damage, surface degradation and electrolyte loss deteriorate the electrochemical performances of high-voltage LCO.

  To improve the cycling stability of LCO at high voltages, researchers have proposed various modification strategies, such as element doping, surface coating or both. Element doping, is a simple, effective, and widely used method, enables the modification of the electronic structure, lattice parameters, and other fundamental physical properties of LCO through the introduction of external elements into the lattice. By carefully selecting doping elements, controlling doping concentrations, and adjusting doping sites (such as Li, Co, or O sites), the structural stability and electrochemical performance of LCO can be enhanced. Common doping elements include Mg, Al, Ti, La, Gd, and F, among others, are included [39–41]. Surface coating has been proven to be an effective method for enhancing the surface and interface stability. The coating layer acts as a physical barrier, effective preventing direct contact between LCO and the electrolyte during deep delithiation, thus stabilizing the LCO/electrolyte interface and suppressing the surface degradation. Oxides with high thermodynamic and structural stability, such as Al₂O₃ and ZrO₂, are usually applied [42,43]. Additionally, electronic or ionic conductive materials are employed as coating agents to improve charge transfer and ion transport at the LCO/electrolyte interface. Common coating materials include LiMn_0.75Ni_0.25O₂, Li_1.5Al_0.5Ti_1.5(PO₄)₃, LiCoPO₄, and LiFe_0.4Mn_0.6PO₄, etc. [44,45].

  Combining element doping with surface coating and utilizing the synergistic effect of the two methods is an important strategy to simultaneously improve the LCO structure and interface stability. In some cases, element doping process at high temperature can also produce a coating layer on the surface [46–48]. Since the distribution of dopants is influenced by factors such as the solubility of the dopant in the LCO lattice, the diffusion rate, as well as the doping process. Consequently, under conditions of low solubility or difficult diffusion, the doping elements may accumulate on the surface of LCO, forming a second phase that effectively functions as a surface coating layer. On the other hand, high-temperature treatment is often required during the coating surface to facilitate the solid reaction between the precursor and LCO, which can also induce some elements diffuse into the LCO lattice. Zhang et al. [49] utilized a chemical reconstruction strategy to engineer a collective LCO surface with a three-layer configuration comprising an amorphous Li₃PO₄ outermost layer, a spinel-like intermediate layer, and a Mg-diffusion layer. This innovative design significantly enhances the cycle life of LCO with capacity retention of 82% after 1000 cycles at 4.6 V. Cai et al. [50] constructed a 3 nm thick La_1-xCa_xCoO_3-δ perovskite structure on the surface of LCO through high-temperature heat treatment, along with a La/Ca gradient-doped buffer layer, effectively suppressing interfacial side reactions and surface structural degradation.

  Extensive research has driven the development of LCO to voltages of 4.6 V and beyond, but the ultimate charge cutoff voltage for LCO remains unclear. Approaching the delithiation limit of LCO, along with the search for alternative materials to replace it, represents promising directions for future research.

  2.1.2   LiNi_xCo_yMn_1-x-yO₂ and LiNi_xCo_yAl_1-x-yO₂ cathodes

  The high cost and limited capacity of layered LiCoO₂ cathodes trigger the development of LiNi_xCo_yMn_1-x-yO₂ (NCM) and LiNi_xCo_yAl_1-x-yO₂ (NCA) ternary cathodes for LIBs [51]. To date, increasing Ni content is the main research direction in the field such as high-nickel (x ≥ 0.8) or ultrahigh-nickel (x ≥ 0.9) cathodes, since it can further reduce cost while increasing capacity. However, such high-nickel ternary cathodes face several challenges, such as cycling instability (Li⁺/Ni²⁺ cation mixing), thermal instability, and air instability (surface lithium residual) [52]. To address above problems, several strategies have been reported, such as novel synthesis, single-crystallization, bulk doping, surface coating, gradient structure design.

  High-nickel cathodes are generally synthesized by high-temperature solid sintering method, which resulting in severe Li⁺/Ni²⁺ cation mixing (>5%) that severely hinders the Li⁺ diffusion and deteriorates the capacity and cycle life [53,54]. To regulate Li⁺/Ni²⁺ cation mixing, Luo et al. [55,56] developed a novel ion-exchange method to synthesize NCM from Na-NCM precursor. Benefiting from the mild reaction temperature (<200 ℃), the NCM cathode material inherit a low cation mixing of <1% from Na-NCM precursor. Single-crystallization is another strategy to improve the performance of high-nickel cathodes, due to their advantages of fewer grain boundaries, higher mechanical strength, and lower surface area, when compared with the traditional polycrystalline cathode materials [57]. Nevertheless, single-crystalline high-nickel cathode materials is difficult to be synthesized by a facile method [58,59]. Yoon et al. [60] developed a method for the synthesis of single-crystalline cathodes at mild temperature, in which a mechanochemical activation pretreatment deagglomerates polycrystalline precursors, followed by eutectic salt-assisted calcination.

  In addition, bulk doping as well as surface coating can mitigate irreversible phase transition, crack formation, lattice oxygen release, interfacial side reactions, etc., which helps to improve the performance of high-nickel cathodes [61–63]. For example, high-valance dopants (Ta⁵⁺, Nb⁵⁺, Mo⁶⁺, etc.) reinforce segregation at the grain boundary and thus lead to aligned microstructure as well as high crystallinity [64,65], while low-valence dopants (Al³⁺, etc.) incorporate into the bulk structure and thus stabilize lattice structure [66,67], leading to enhanced electrochemical performance. Besides doping, Fan et al. [68] coated NCM with a NASICON-type ion conductor of Li_1.8Sc_0.8Ti_1.2(PO₄)₃ (LSTP), which significantly mitigated the irreversible phase transition from layered structure to spinel/rock salt and suppressed crack formation. As a result, the 1% LSTP-modified NCM maintained an excellent retention rate of 90.27% after 500 cycles even at a high operating voltage of 4.6 V. Huang et al. [69] modified the NCM with a hydrophobic coating to inhibit the formation of surface lithium residual and thus improve the air stability. In addition, crack formation can be effectively mitigated by gradient structure design [70]. Recently, Li et al. [71] introduced gradient pores in Ni-rich NCM secondary particles, which provide additional void space to buffer the anisotropic volume change of primary particles, and therefore effectively mitigate microcrack formation. With the enhanced chemical-mechanical stability, the ultrahigh-nickel NCM (LiNi_0.96Co_0.02Mn_0.02O₂) exhibited great cycling stability of 80.5% capacity retention after 800 cycles at 1 C in practical pouch-type cells.

  2.1.3   LiNi_0.5Mn_1.5O₄ cathodes

  LiNi_0.5Mn_1.5O₄ (LNMO) is a cathode material for LIBs with a spinel structure, which has two forms. One is the face-centered cubic phase with the space group Fd3m, where Mn and Ni are randomly distributed at the 16d sites, referred to as the disordered LNMO; the other is the simple cubic phase with the space group P4332, where Mn and Ni are orderly distributed at the 4a and 12d sites, known as the ordered LNMO [72].

  LNMO possesses a theoretical specific capacity of 147 mAh/g with a high voltage platform of 4.7 V, resulting in a theoretical specific energy as high as ~690 Wh/kg (only the cathode is considered), which is significantly higher than that of LiFePO₄ (LFP). Additionally, LNMO does not contain the costly Co element, thus resulting in its low costs. Therefore, LNMO holds potential for large-scale application as a replacement for LFP in electric vehicles and electrochemical energy storage systems [73]. However, the poor cycle stability of LNMO resulted from LNMO include lattice distortion induced by the Jahn-Teller effect of Mn³⁺, dissolution of transition metals, interface side reactions and decomposition of the electrolyte hinders its commercialization [74]. To address these issues, researchers have employed many strategies, such as element doping, surface coating, and the electrolyte modification, to improve the electrochemical performance of LNMO.

  Elements doping can effectively reduce the content of Mn³⁺ in the LNMO lattice which alleviating the lattice distortions caused by the Jahn-Teller effect, then significantly enhance the lattice stability of LNMO [75]. Liang et al. [76] employed Sb replacing Mn to reduce the content of Mn³⁺ then significantly inhibited the lattice distortion of LNMO. As a result, Sb doped LNMO material achieved a specific capacity of 99% of its theoretical specific capacity at 1 C, and maintained 72.4% of its initial capacity after 3000 cycles, demonstrating an excellent cycling stability.

  Surface coating can effectively isolate LNMO from direct contact with the electrolyte, inhibiting the occurrence of interfacial side reactions and the dissolution of TM [77]. Kuenzel et al. [78] constructed a TMPO_x (TM = Mn, Ni) coating layer on the surface of LNMO, which not only isolated the electrolyte from direct contact with the material, but also increased the ionic conductivity of the material’s surface. After optimization with TMPO_x coating layer, LNMO exhibited a capacity of 100 mAh/g after 1000 cycles with a capacity retention rate as high as 95%, which tested at a high cut-off potential of 4.95 V with a fast discharge/charge rate of 10 C and at a high temperature of 40 ℃.

  Electrolyte modification can effectively reduce the HOMO value of the electrolyte, enhancing the antioxidant properties and broadening the electrochemical window, which improves the stability of the electrolyte at high voltages, and hinders the occurrence of harmful side reactions [79]. Vald et al. [80] employed dimethyl-2,5-dioxahexanedioate as a solvent of liquid electrolyte, which is stable with LNMO at 5.2 V, significantly enhancing the cycling stability of Li||LNMO batteries. After adopting this solvent, the Li||LNMO batteries displayed a capacity of ~135 mAh/g after 250 cycles at a 0.2 C rate, with an excellent capacity retention as high as 97%.

  2.1.4   Lithium-rich layered oxides

  Lithium-rich manganese-based materials (LRMOs) are comprised of two distinct structural components [81]: LiMO₂ with a layered α-NaFeO₂ structure (space group: R-3m) and Li₂MnO₃ with a rock-salt α-NaFeO₂ structure (space group: C2/m). The overall crystal structure of LRMOs, however, remains a subject of ongoing debate, with two prevailing viewpoints: A single-phase solid solution or a two-phase nanocomposite [82–84]. In Li₂MnO₃, one-third of the transition metal sites in the transition metal layer are occupied by lithium atoms, forming a dumbbell-like atomic arrangement and exhibiting a characteristic honeycomb structure on specific crystallographic planes [85]. During the initial charge of a typical LRMO, xLi₂MnO₃·(1-x)LiMO₂, lithium ions are primarily extracted from the LiMO₂ component at voltages below 4.4 V (Figs. 2a and b) [86,87]. These ions migrate through the electrolyte to the anode, accompanied by the oxidation of transition metal ions (Ni²⁺→Ni⁴⁺, Co³⁺→Co⁴⁺) [88], ultimately transforming LiMO₂ into MO₂. Lithium ions from the octahedral sites within the manganese layer of Li₂MnO₃ migrate to tetrahedral sites in the lithium layer of LiMO₂, replenishing the consumed lithium ions and stabilizing the crystal structure [82,89]. When the charging voltage exceeds 4.4 V, a long voltage plateau appears, corresponding to the activation of Li₂MnO₃. Two primary mechanisms are proposed to explain the voltage plateau observed in Li₂MnO₃ (Fig. 2c) [90,91]: One attributes it to a high-valent manganese redox reaction [92], where half of the Mn⁴⁺ ions are oxidized to Mn⁷⁺. The other ascribes the plateau to a reversible lattice oxygen redox reaction [93,94], involving the reversible redox couple O²⁻/O⁻ during charge/discharge. During the subsequent discharge, lithium ions are re-inserted into both the MO₂ and the activated MnO₂ components. This mechanism enables LRMOs materials to exhibit reversible capacities exceeding 300 mAh/g and energy densities above 900 Wh/kg [95]. Furthermore, their high manganese content, low toxicity, and cost-effectiveness position them as promising candidates for next-generation low-cost, high-energy-density lithium-ion battery cathodes. However, LRMOs face challenges related to capacity and voltage fade, hindering their commercialization. These challenges stem from several factors, including lattice oxygen loss, electrode/electrolyte interfacial side reactions, irreversible structural rearrangements, and irreversible phase transformations induced by transition metal ion migration, which collectively diminish the energy density of the battery system [96–98]. Current mitigation strategies include surface coating, defect engineering, elemental doping, and electrolyte modification [99]. While these approaches can enable LRMOs to achieve capacities exceeding 300 mAh/g at a current density of 20 mA/g [100–102], cycle life and voltage retention often remain suboptimal. Conversely, structural designs that prioritize high-capacity retention often struggle to achieve high specific capacities [103–105]. Therefore, future research should focus on synergistic strategies that combine one or more of these approaches to simultaneously achieve high capacity, superior capacity and voltage retention in LRMOs.

  Figure 2

  

  Figure 2.  (a) Crystal structure of lithium-rich manganese-based materials. (b) Typical first-cycle charge-discharge curves of lithium-rich manganese-based cathode materials. (c) Compositional phase diagram showing the electrochemical reaction pathways for a xLi₂MnO₃·(1-x)LiMO₂ electrode. (c) Reproduced with permission [87]. Copyright 2021, Wiley-VCH.

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  2.1.5   Disordered rock-salt cathodes

  With the rapid advancement of rechargeable LIBs, lithium-rich cathode materials with excellent performance have garnered increasing attention. Currently, lithium-rich cathode materials are primarily classified into lithium-rich layered oxides (LLOs) that have been extensively studied and lithium-rich disordered rock salt oxides (DRXs) that have emerged in recent years [106]. Notably, conventional LLO electrodes are characterized by high specific capacity, high energy density, and relatively low cost. However, problems such as low first-cycle Coulombic efficiency, poor rate capability, and limited cycling stability have hindered their further commercialization. In recent years, DRXs have emerged as a promising alternative. These materials exhibit similar lithium-rich characteristics and charge compensation mechanisms and have been shown to offer not only high reversible specific capacities (>300 mAh/g) but also higher theoretical energy densities (>1000 Wh/kg) and lower irreversible capacity loss during the first cycle. Interestingly, while similarities exist between DRXs and LLOs in terms of chemical composition, structure, Li⁺ migration, and electrochemical cycling behavior, there are also notable differences. Specifically, the chemical compositions of DRXs and LLOs are similar, with their chemical formula generally expressed as Li_1+xTM_1-xO₂ (where TM represents transition metals). However, LLOs typically rely on Ni and Co to maintain structural stability during cycling. In contrast, DRXs can incorporate a variety of transition metals (V, Fe, Cr, Mo, Ti, Mn, etc.). While both structures share a common oxygen lattice arrangement, the key difference lies in their atomic arrangement: in LLOs, lithium and transition metals alternate in a layered configuration, whereas in DRXs, lithium, and transition metals are randomly mixed in the disordered rock salt structure. This disordered arrangement of cations in DRXs creates a unique three-dimensional network that facilitates enhanced Li⁺ diffusion, insertion, and extraction, thereby improving the migration rate and diffusion flux of Li⁺ during cycling. However, despite their promising potential, DRXs still face significant challenges in commercial applications, including poor cycling stability and limited rate capability, which require urgent attention.

  To address the drawbacks of DRXs, researchers have employed various strategies, including surface engineering [107,108], elemental doping [109,110], particle engineering [111], fluorination [112,113], and high-entropy approaches [114], achieving remarkable progress. The most recent modification approaches primarily involve structural design and anionic site engineering. Cai et al. [115] investigated the electrochemical behavior of Li_1+xMn_1–3xTi_2xO₂ (where x = 0.05, 0.1, and 0.15), which transforms electrochemical cycling into partially disordered spinel-like phases with short coherence lengths, referred to as the δ-phase (Fig. 3a). Their findings indicate that Mn-rich DRX compounds tend to form the δ-phase, and the extent of this transformation during cycling significantly impacts long-term electrochemical performance. The presence of Ti⁴⁺ ions inhibit the formation of a fully ordered spinel structure, with a Ti content at or slightly above the minimum required to suppress the two-phase reaction. Notably, the transition from DRX to the δ-phase is more pronounced in Mn-rich, Ti-poor compositions. Hau et al. [116] chemically delighted and heat-treated Li_1.2Mn_0.65Ti_0.15O_1.9F_0.1 at low temperatures to obtain Li_0.7Mn_0.65Ti_0.15O_1.9F_0.1, with an average particle size of approximately 10 µm (Fig. 3b). At a current density of 20 mA/g, the first discharge capacity reached 201 mAh/g, and 110 mAh/g at 500 mA/g. These values represent unprecedented high specific energy and rate capability for micron-sized particles, suggesting that the δ-phase can be formed through a non-in situ chemical decomposition process coupled with low-temperature thermal treatment, which significantly shortens the time needed for the δ-phase transition. Huang et al. [117] successfully synthesized a series of disordered rock salt polyanionic spinel (DRXPS) cathodes with the general formula Li_2+u-vM_2-u[XO₄]_XO_4(1-x), where M denotes a transition metal, XO₄ represents a polyanionic group, and u, v, and x are stoichiometric parameters (Fig. 3c). The rock salt-polyanion integrated structure was found to be an effective strategy for enhancing the stability of DRX cathodes. The DRXPS cathodes exhibited excellent cyclic stability, retaining 70% of their energy density after 100 cycles, alongside impressive rate performance and a highly tunable compositional space.

  Figure 3

  

  Figure 3.  Design of DRX and DRXPS cathodes. (a) The partially disordered spinel phase as an intermediate between the DRX and fully ordered spinel structures. (b) Schematic of the structural evolution of Li_1.2Mn_0.65Ti_0.15O_1.9F_0.1 with each synthesis step. (c) The structure of M₂O₄, M_2−uO₄ and M_2−u[XO₄]_xO_4(1−x). (a-c) Reproduced with permission [115–117]. Copyright 2024, Springer Nature.

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  2.1.6   Polyanionic cathodes

  Polyanionic cathode materials are typically represented by the chemical formula Li_xM_y(XO_n)_z, where M represents transition metal elements such as Fe, Ni, Mn, and X represents elements like S, P, Si, B, etc. These cathode materials have the following advantages: (1) The structural framework is relatively stable during lithiation/delithiation process, possess better thermal stability and safety compared to oxide cathodes; (2) The battery voltage is higher than that of oxide counterparts with the same redox pairs; (3) The ionic character of the M–O bond can be adjusted by varying M and X, which allows for the design of discharge potentials that meet application requirements (inductive effect); (4) They possess a rich variety of crystal structures, providing possibilities for cation and anion substitution in a given structure. However, polyanionic cathodes also have significant drawbacks: (1) Lower electrical conductivity; (2) Limited theoretical specific capacity because of their large molecular weight of the polyanion groups. To further reduce costs and develop iron-based cathodes, in the late 1980s, Goodenough et al. [118] researched a series of iron-based polyanionic oxides Fe₂(XO₄)₃ (X = Mo, W, S). Compared with Fe₂O₃, the P–O covalent bond endows LiFePO₄ with a higher voltage. Based on this, they proposed the inductive effect of cation Xⁿ⁺ to enhance the cathode potential [119].

  Compared to oxides, polyanionic cathodes have a diverse chemical composition and crystal structure, including sulfates, phosphates, silicates, etc. LiFePO₄(LFP) with an olivine structure, has been commercialized because of their stable charge-discharge platform, economic and environmental benefits, and higher safety. To enhance their conductivity and energy density, modifications are primarily made through ion doping, surface coating, morphology control, and addition of lithium-complementary material. LiMnPO₄ with LFP structural approximation has a high voltage of 4.1 V, however, their intrinsic conductivity is low. Martha et al. [120] synthesized carbon-coated LiMn_0.8Fe_0.2PO₄ material through high-temperature solid phase reaction, which exhibited stable cycling with a capacity of approximately 165 mAh/g. Companies in China, including BYD company have begun commercializing it. The synthesis process of pure phase LiVPO₄F is quite complex, this material has a higher discharge voltage of 4.2 V and excellent cycling stability, and it also has potential for application in power batteries. Kim et al. [121] reported a one-step, large scale solid-phase synthesis method for LiVPO₄F. Zhang et al. [122] suppressed the formation of impurity phase Li₃V₂(PO₄)₃ through adding polyvinylidene fluoride (PVDF) as a supplemental fluorine source and conducted B–doping on the thermally decomposed carbon from PVDF to synthesize LiVPO₄F material. The synthesized LiVPO₄F material demonstrated a capacity of 132.9 mAh/g at a 5 C rate. Lander et al. [123] have synthesized a variety of sulfates, including fluorosulfates, hydroxysulfates, and oxosulfates as cathode materials. Among these, the triplite-structured LiFeSO₄F possesses a specific capacity comparable to LFP (~150 mAh/g), and a working voltage that is 0.5 V higher than LFP (3.9 V), with better electrical conductivity than LFP. LiFeSO₄F needs to overcome the issues of slow kinetic performance and the hygroscopic nature of sulfates [124]. Sun et al. [125] prepared the Fe₂O(SO₄)₂ material, which exhibits the best water resistance among sulfates. However, this material is lithium-deficient and has a relatively low voltage (3 V vs. Li⁺/Li⁰). Currently, there is less exploration of modification processes such as sulfate coating and carbon coating compared to phosphates, and there is potential for future performance improvements and commercialization. Additionally, low-cost mixed polyanionic cathode materials are currently less researched, it is also a direction worth exploring in future research [126].

  2.1.7   Organic cathodes

  Organic electrode materials (OEMs) are composed carbon, hydrogen, oxygen, nitrogen, and other non-metallic abundant elements [127,128]. The application of organic compound as cathode materials enables access to metal-free, low-cost, and environmentally friendly LIBs systems, which is of significant research value from the perspective of sustainable development [129]. In the past decade, intensive and innovative efforts have been developed for LIBs, including conducting polymers, organosulfur, conjugated carbonyl compounds, amines compounds, nitroxide radicals-based polymer and conjugated sulfonamides [130,131]. Based on the charge states variation of functional group and the compensation of counter ions, OEMs can be grouped into three types: p-type, n-type, or bipolar-type [128]. Most of n-type organics succeed in its high capacity but with low operating potential, typically display a redox potential lower than 3.0 V vs. Li⁺/Li. It is urgently to explore organic cathode materials with a high-working potential. Among them, conjugated sulfonamides have achieved high redox potentials above 3.0 V vs. Li⁺/Li, which is expected to achieve the 4 V-class organic cathodes [132]. With the continuous increase of concern to resources and environmental issues, a variety of organic structures and redox chemistries have been investigated, including both organic polymers and crystallized organic compounds [133]. Despite receiving increasing attention, research on OEMs is still at its early stage, facing great challenges to reach the goal of practical application [134]. The intrinsic low electrical conductivity and low tap density pose a great challenge to the fast charge/discharge capability and high volumetric energy density. Hence, a large amount of conductive carbon additive is required to obtain sufficient electronic conductivity in the preparation of electrode, which inevitable leads to the reduce of energy density of full cells. The dissolution of solid electrode into electrolyte is the main cause of stability deterioration. To meet these challenges, composite engineering and electrolyte engineering have been employed to avoid adverse reactions for the battery operation and stability. To shorten the distance of organic cathode materials to practical applications, there is an urgent need to make breakthroughs in terms of molecular structure innovation, electrode preparation process optimization, and Ah-level full cells performance verification.

  2.1.8   Halide cathodes

  Halide cathode active materials (CAMs) have gained considerable attention in the advancement of all-solid-state batteries (ASSBs) due to their exceptional deformability and enhanced ionic conductivity. These attributes facilitate the formation of intimate interfaces and reduce the reliance on non-active catholyte content. In 2023, VCl₃ emerged as a promising CAM, exhibiting reversible intercalation-deintercalation electrochemistry within a voltage range of 2.45–3.25 V and an impressive rate capability exceeding 80 mAh/g at 6 C [135]. However, the ionic-insulating nature of VCl₃ necessitated a substantial amount of catholyte to establish Li⁺ transport pathways, resulting in a low CAM content of 45 wt% and, consequently, limiting energy density improvements.

  To address this challenge, Li-containing halide-based cathode materials, such as Li₃VCl₆ and Li₃TiCl₆, were developed, exhibiting high room-temperature ionic conductivities of 0.075 mS/cm and 1.04 mS/cm, respectively [136,137]. These high ionic conductivities enabled the design of catholyte-free cathodes, increasing the CAM content to 95 wt%, representing a 10%-20% improvement compared to traditional oxide CAMs [138,139]. This advancement holds promise for achieving higher energy densities. For example, leveraging a two-electron reaction from Ti⁴⁺ to Ti²⁺, ASSLBs assembled with 95 wt% Li₃TiCl₆ CAM achieved a high capacity of 184.5 mAh/g at 19 mA/g [136]. However, the energy density of these materials remains constrained by the low potentials of the V³⁺/V²⁺ (~3.0 V) and Ti⁴⁺/Ti³⁺ (3.13 V) redox couples and the limited capacity associated with the intercalation-deintercalation mechanism. Additionally, the practical application of these materials faces economic challenges due to the high cost of raw materials, such as TiCl₃ and VCl₃.

  Against this backdrop, Fe-based halide CAMs have emerged as highly promising alternatives. Iron’s higher electronegativity suggests the potential for higher redox voltages, and its abundance and affordability make it a more economical choice compared to titanium and vanadium [140–142]. A particularly striking feature of Fe-based CAMs is their multi-electron transfer capability. Zhou et al. [142] reported that Li_xFeX_x+2 (X = Cl, Br) families demonstrated a three-electron transfer reaction between Fe³⁺ and Fe, relying on an intercalation-conversion coupling mechanism. Notably, catholyte-free ASSLBs employing 95 wt% LiFeCl₃ as the active material delivered an outstanding capacity of 446 mAh/g and an energy density of 912 Wh/kg, outperforming typical oxide and halide CAMs by 1.5 times.

  In summary, halide CAMs exhibit high deformability, superior ionic conductivity, and multi-electron transfer capabilities, making them a promising avenue for advancing ASSLB technology. However, similar to halide solid-state electrolytes, halide CAMs are moisture-sensitive [143]. Thus, developing moisture-tolerant materials remains a critical priority for future research.

  2.2 Anode materials

  2.2.1   Graphite and carbon anodes

  (1) Graphite anodes

  Graphite, a typical layered carbon-based material, can incorporate other heterogeneous particles such as atoms, molecules, ions, or even atomic clusters into its interlayer space through physical or chemical methods, forming a new type of layered compound known as graphite intercalation compounds (GICs). Alkali metal ions (AMIs) can easily form ionic graphite intercalation compounds (AM-GICs) when intercalated into graphite. This is due to the high reactivity of alkali metals, which readily lose their outermost electrons to bond with the nonmetallic carbon atoms. The introduction of electrochemical methods for the preparation of AM-GICs marked a significant turning point in the application of graphite as an anode material for batteries. This also laid the foundation for the concept of rocking-chair battery design, avoiding the direct use of reactive alkali metals as the anode [144].

  Well-defined layered structure and appropriate interlayer spacing of graphite facilitate the intercalation and deintercalation of lithium ions. The lithium intercalation/deintercalation reaction in graphite occurs at potentials between 0 and 0.25 V (vs. Li⁺/Li), and the interlayer spacing of graphite reaches 0.370 nm after full lithiation. The evolution of graphite "stage" structure has long been a focus of scientific research. In 1991, J. R. Dahn and colleagues first proposed the D-H model to describe the changes in carbon structure during lithium intercalation in graphite within organic electrolyte systems [145]. According to this model, the staging transitions proceed from higher stages to "stage 4" then to "stage 3", gradually evolving into a "diluted stage 2" as lithium ions intercalate, before fully transitioning to "stage 2" and eventually reaching "stage 1" forming LiC₆ and achieving a theoretical capacity of 372 mAh/g. With advancements in characterization techniques and deeper research, it has been discovered that the lithium intercalation process in graphite involves more than just simple stage transitions. As research has progressed, the coexistence of the Rudörff and D-H mechanisms, as illustrated in Fig. 4, has emerged as the widely accepted explanation.

  Figure 4

  

  Figure 4.  Schematic diagram of Rudörff model and Daumas-Herold model of lithium storage process in graphite.

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  During the charge and discharge processes, the interlayer spacing of graphite repeatedly expands and contracts. This can lead to co-intercalation of solvents such as lithium and propylene carbonate, which may cause irreversible structural damage to graphite. Such damage can result in layer delamination and shedding, ultimately leading to poor cycling stability. To mitigate these side reactions, graphite used in practical applications is typically coated with a layer of amorphous carbon, forming a "core-shell" structure. This coating prevents direct contact between graphite and the solvent, thereby blocking solvent co-intercalation, improving compatibility with solvents, and reducing side reactions, which enhances the stability of the graphite anode. However, the amorphous carbon coating increases the composite material’s specific surface area and introduces more defects, reducing the initial Coulombic efficiency (ICE). Therefore, the selection of coating materials, coating methods, coating thickness, and post-coating processing must be designed and optimized based on practical requirements. In addition to amorphous carbon coating, surface deposition of a layer of metal or metal oxide is another effective approach. Graphite surfaces contain oxygen-containing functional groups and adsorbed impurities, which can affect the formation of the solid electrolyte interphase (SEI) during electrochemical processes, leading to increased irreversible capacity loss, especially in natural graphite. Surface modification of graphite to enhance the reversible capacity and cycling stability of natural graphite is a commonly used strategy [146]. Among these, surface oxidation is the most widely adopted method, primarily divided into gas-phase oxidation and liquid-phase oxidation. Gas-phase oxidation typically involves high-temperature treatment of graphite with oxidizing agents such as air, oxygen, water, ozone, or carbon dioxide. However, due to the poor uniformity of gas-phase oxidation, it is not well-suited for large-scale industrial production. In contrast, liquid-phase oxidation can effectively address this issue, making it more practical for industrial applications. Strong oxidizing agents such as nitric acid, sulfuric acid, and hydrogen peroxide are generally used. After liquid-phase oxidation, the electrochemical properties of natural graphite, such as cycling stability and reversible capacity, are significantly improved.

  The rapid development of electric vehicles has placed higher demands on the energy density, power density, and cycle life of next-generation LIBs. As a result, specialized graphite structures, such as microporous, nanoporous, and polyhedral graphite, have garnered considerable attention from researchers to address the current bottlenecks and limitations of graphite anodes [147].

  MCMBs are a specialized graphite material formed by processing pitch-based organic substances under an inert atmosphere at high temperatures. This process produces thermodynamically stable condensed polycyclic aromatic hydrocarbons with high molecular weight and planar structures. After further treatment, spherical particles with diameters ranging from 5 µm to 100 µm are obtained. MCMBs possess a dense spherical structure, offering advantages such as high tap density and volumetric capacity compared to natural graphite and other carbon-based anodes [148]. Initial discovery of MCMBs dates back to the 1960s, when researchers identified optically anisotropic small spheres in coal tar pitch, representing the nascent form of mesophase carbon microspheres. In 1973, Japanese researchers synthesized these spherical carbon materials from mesophase pitch and named them MCMBs. Subsequently, extensive research was conducted on this carbon material, and many enterprises have also successfully industrialized MCMB production.

  (2) Amorphous carbon materials

  Amorphous carbon materials (hard carbon and soft carbon) are also important anode materials, especially in sodium-ion battery systems, where they represent the mainstream commercial anode materials. Amorphous carbon materials have low crystallinity and disordered layered structures. Compared to soft carbon, hard carbon is more difficult to graphitize, even at temperatures above 2500 ℃. Soft carbon exhibits a low and stable charge/discharge voltage plateau, large capacity, high efficiency, and good cycling performance, making it suitable for battery applications. Hard carbon, on the other hand, features a stable structure, long charge/discharge cycling life, and a carbon-lithium potential above 0.2 V, providing better safety performance.

  Due to the complex and irregular structure of amorphous carbon materials, their energy storage mechanism is also complex. the mechanism mainly includes intercalation between carbon layers, adsorption in micropores and defects within graphite microcrystals, and filling of voids and clusters formed by stacked layers. Researchers collectively refer to this as the "House of Cards" model [149]. Different storage mechanisms correspond to voltage plateaus in different regions. The crystallinity, crystal size, density, porosity, specific surface area, and heteroatom content of amorphous carbon are closely related to the carbonization conditions and precursor selection, making the study of its energy storage mechanism a highly complex process [150]. The working voltage range of hard carbon anodes is between 0 and 1.5 V (vs. Li⁺/Li), with capacities ranging from 200 mAh/g to 600 mAh/g. The voltage profile consists of two main components: A sloping region in the 0.1–1.0 V range and a plateau region at lower voltages [151,152]. In 1991, Sony Corporation first used hard carbon derived from high-temperature pyrolysis of polyfurfuryl alcohol as an anode material. With the continuous development of hard carbon materials, they have gradually been commercialized as battery anode materials.

  (3) Other carbon materials

  Graphene is highly favored as a battery anode material. With its extremely high specific surface area and excellent electrical conductivity, graphene can significantly enhance the energy density of energy storage batteries. However, the high cost of graphene production poses significant challenges for large-scale manufacturing. Currently, the lithium storage capacity of graphene remains relatively low, requiring further improvements. Moreover, graphene’s high chemical activity may lead to potential battery safety issues. Nevertheless, as a novel material, graphene shows tremendous potential for application in the energy storage field. Its use in energy storage batteries is expected to achieve revolutionary breakthroughs in energy storage technology [153].

  Carbon nanotubes (CNTs) are nanoscale tubular structures composed of carbon atoms. They exhibit exceptional electrical conductivity, excellent thermal stability, and high mechanical strength, making them highly promising for applications in various fields. In battery technology, these properties make carbon nanotubes ideal candidates for electrode materials. In recent years, with the continuous advancement of nanotechnology and battery technology, significant progress has been made in the study of carbon nanotubes as anode materials. Scientists have optimized the electrochemical performance of CNTs by controlling their morphology, structure, and doping elements. These research achievements provide a solid foundation for the application of carbon nanotubes in next-generation high-performance batteries [154,155].

  In recent years, significant research progress has been made in the theory, synthesis, and application of graphdiyne. One notable area of exploration is its application in high-energy-density lithium/sodium-ion batteries. graphdiyne can be synthesized at low temperatures, and its fundamental structural units can be precisely designed. This precise control enables the introduction of more electrochemical active sites and the adjustment of its electronic structure, paving the way for the development of novel, high-performance lithium-ion battery anode materials-advantages that conventional carbon materials cannot achieve. The acetylene-rich structural network of graphdiyne provides abundant storage sites for lithium ions, with a theoretical lithium storage capacity exceedingly twice that of conventional graphite (372 mAh/g), offering significant potential for enhancing the energy density of LIBs. Moreover, graphdiyne possesses three-dimensional ion transport channels, providing a structural foundation for its rapid ion transport properties. As a result, graphdiyne is gradually emerging as a promising anode material for batteries [156].

  2.2.2   Silicon anodes

  LIBs are electrochemical batteries widely used in electronic devices, electric vehicles, and energy storage systems. The high energy density, long calendar life, and wide applicable scenarios make them an important part of modern energy storage technology. The key components of LIB consist of cathode (usually a lithium compound, such as lithium cobalt oxide), anode (usually a carbon-based material, such as graphite), electrolyte and separator. During charging and discharging, Li⁺ ions transfer between the two electrodes. The traditional graphite anode has reached its theoretical capacity up limit and is hard to meet with the demand for higher energy density. Researchers continue to explore next-generation anode materials as a candidate. Silicon anodes, in particular, have attracted attention because of their extremely high theoretical capacity (3579/4200 mAh/g), much higher than the 372 mAh/g of traditional graphite anodes [157,158]. The advantages of silicon anodes include high specific capacity, abundance, low cost, and environmentally friendly [159]. However, it also faces significant challenges, firstly, silicon particle undergoes volume expansion up to 300% during lithiation process. The drastic volume change can lead to pulverization and delamination of electrode, which leads to rapid degradation of the battery capacity. Secondly, the intrinsic poor conductivity of silicon limits the rapid electron transfer through the electrode and affects the rate performance. Thirdly, the issue of low coulombic efficiency (ICE) is also significant with regard to silicon-based anodes. During the initial lithiation process, a significant amount of Li⁺ is consumed to form the solid electrolyte interphase (SEI) film, which results in irreversible loss of active lithium. In summary, silicon anodes tend to have difficulty maintaining structural stability during charge/discharge cycling, which raises a great challenge to the long cycling life of batteries. Silicon anodes can be classified into three main categories: pure silicon anodes, silicon carbon anodes and silicon oxygen anodes [160]. In recent years, researchers have made significant progress in silicon anode research. The development of nanoscale silicon materials to alleviate the problem of volume expansion [161,162], the application of silicon/carbon composites to enhance conductivity [163–165], pre-lithiation [166,167] and improvement of binder [168–170] can improve the ICE of the electrodes and enhance the mechanical strength and stability of the electrodes, and improve the cycling performance of the batteries. Studies on silicon anodes have also revealed the formation of SEI [171,172] and the role of electrolyte additives [173,174]. The stability and composition of SEI layer are crucial to influence the battery performance, a stable SEI layer can effectively prevent further decomposition of the electrolyte, protect the electrode and extend battery life. In massive research, the role of inorganic SEI components in enhancing the SEI stability and ionic conductivity are highlighted, including LiF [175], Li₂CO₃ [176,177], and Li₂O [178,179], etc. The electrolyte additives, on the other hand, are widely used in silicon anode with liquid electrolyte to assist the SEI formation, and thus improve the cycling life and overall battery performance. The development of solid-state electrolytes for silicon anodes in LIBs has received significant development, aiming to address safety concerns and performance limitations associated with traditional liquid electrolytes. Recent research has focused on inorganic solid electrolytes, which offer higher thermal stability and a wider electrochemical stability window compared to organic counterparts, the most widely used solid-state electrolyte including polymer electrolytes of interfacial compatibility [180], sulfide type with good mechanical property [181], perovskite with high ionic conductivity [182]. These developments are crucial for creating safer, more efficient batteries with high energy density and improved cycle life, especially for applications requiring high power and safety, such as electric vehicles and large-scale energy storage systems.

  For silicon anodes, future research directions including two main parts: exploring new research methods and developing innovative technologies to overcome existing challenges. The main development target is to promote the development of silicon anode with high active material content in practical application. For examples, the development of more efficient and stable silicon-based anodes through molecular design and nanoengineering techniques, the enhancement of the conductivity and mechanical properties of electrodes through the optimization of electrode structures and formulations, and the development of novel electrolytes and additives to enhance the stability and functionality of the SEI layer. The successful application of silicon anode is expected to significantly improve the energy density and performance of LIBs, which is of great significance in promoting the development and application of battery technology.

  2.2.3   Phosphorus anodes

  With the burgeoning demand for renewable energy and electric vehicles, enhancing the low-temperature and fast-charging performance of LIBs has emerged as a pivotal area of research [183]. Phosphorus anode materials have garnered considerable attention due to their exceptionally high theoretical specific capacity (2596 mAh/g) and reasonable lithiation potential [184]. The investigation of phosphorus anodes commenced in 2007, primarily centering on their application within LIBs [185]. Compared to conventional graphite anodes, phosphorus not only demonstrates superior capacity during lithiation but also encounters formidable challenges, including significant volume expansion (~300%) and inadequate cycling stability [186,187]. To surmount these obstacles, researchers have explored various synthesis methodologies and the design of composite materials, such as compositing phosphorus nanoparticles with conductive carbon matrices to form P-C or P-O-C bonds, thereby enhancing electronic and ionic conductivity while mitigating volume expansion [188–190].

  The unique physicochemical properties of phosphorus render it an exemplary candidate for low-temperature and fast-charging anodes, attributed to its relatively safe lithiation potential (~0.75 V vs. Li⁺/Li), low ionic diffusion barrier, and remarkable storage capacity. Recent investigations have indicated that phosphorus anodes exhibit reduced electrochemical polarization, which complicates the likelihood of the potential dipping below 0 V during fast charging conditions, thus suppressing lithium plating [191]. To further diminish electrochemical polarization at low temperatures and fast-charging condition, researchers have employed heteroatom doping strategies to catalyze solid-state alloying reactions in phosphorus-based anodes. This accelerates P-P bond cleavage during lithium storage, facilitating an impressive 80% charge in merely 9 min [192]. Given these electrochemical characteristics, phosphorus anodes showcase exceptional potential for applications in low-temperature battery systems [193].

  Regarding electrolyte compatibility for phosphorus anodes, traditional ethylene carbonate (EC)-based electrolytes possess a decomposition potential that is lower than the lithiation potential of phosphorus, complicating the formation of a stable solid electrolyte interphase (SEI) during the initialization stage and resulting in capacity loss. Consequently, it is imperative to utilize solvents or additives with higher reduction potentials to establish a stable SEI before lithiation can commence. For instance, the incorporation of high-reduction-potential additives can lead to the formation of a stable LiF/Li₃PO₄-rich SEI on phosphorus anodes, substantially minimizing the adverse effects of interfacial reactions [184].

  Looking ahead, research on phosphorus anode materials should concentrate on overcoming challenges associated with fast-charging, low-temperature performance and extended cycle life. By judiciously designing composite structures, investigating novel electrolytes, and leveraging emerging technologies such as machine learning, it is anticipated that superior phosphorus anode material systems will be developed. In addition to ongoing chemical optimization and engineering design, addressing the recycling and sustainability of phosphorus anode materials is crucial to fostering a more environmentally friendly battery industry [194].

  2.2.4   Metal oxide anodes

  Metal oxide anodes, owing to their superior reversible capacity, high energy density, and enhanced safety compared to conventional graphite-based material, are regarded as promising candidates for advanced LIBs [195]. According to the electrochemical reaction between anode and Li, the reaction mechanisms can be categorized into three types: (1) Alloying, (2) conversion, and (3) intercalation. The alloying reaction occurs at low potentials (≤1 V vs. Li⁺/Li) in SnO₂, SnO, and ZnO. For instance, SnO₂ has a theoretical capacity of 780 mAh/g, while ZnO has a theoretical capacity of 978 mAh/g. However, the lithium storage capacity of these metal oxides is significantly limited by their large volume expansion (≈240%) during cycling, which leads to pulverization of the anodes. The conversion reaction mechanism involves the formation and decomposition of Li₂O, along with the reduction and oxidation of metals. These metal oxides are of significant interest as potential high-performance anode materials due to their ultrahigh theoretical specific capacity, low operating voltage, reduced toxicity, and low cost. However, as anode materials, they also exhibit significant drawbacks, such as low electronic and ionic conductivity, as well as poor cycling and rate performance.

  Metal oxides for LIBs based on intercalation reaction include TiO₂, Li₄Ti₅O₁₂, Nb₂O₅, and TiNb₂O₇, which enable reversible lithium-ion intercalation while preserving the integrity of the crystal structure. Spinel Li₄Ti₅O₁₂, a typical high-power anode, features excellent structural stability and safety along with low-temperature performance compared with graphite, but its poor electronic conductivity and low theoretical capacity (175 mAh/g) still fails to meet commercial applicability.

  Nb-based anodes, such as Nb₂O₅ and TiNb₂O₇, possess large lattice spacing and open channels that enable fast ion diffusion and reversible intercalation pseudocapacitance, while operating at a high working voltage of 1.65 V vs. lithium anodes, effectively preventing the formation of lithium dendrites and an unstable solid electrolyte interface (SEI) film. Also, the multiple redox couples (Nb³⁺/Nb⁴⁺, Nb⁴⁺/Nb⁵⁺) contribute to their high specific capacity/capacitance (Nb₂O₅ ≈ 200 mAh/g; TiNb₂O₇ ≈ 387 mAh/g), higher than conventional Li₄Ti₅O₁₂ anode. However, the poor electronic conductivity of Nb-based oxides leads to a slower charge transfer rate, thereby increasing the ohmic resistance. Therefore, several strategies have been proposed to improve electron transfer and Li⁺ diffusion rates, thereby enhancing pseudocapacitance kinetics, such as the design of nanostructure, porous materials, oxygen-deficient materials, and multiphase systems. Lou’s group conducted the majority of the work for Nb₂O₅ and TiNb₂O₇, including the investigation of the Li⁺ intercalation mechanism and modification strategy. In 2017, the pseudocapacitive behavior of Nb-based oxides was identified and quantified. Three-dimensionally ordered macroporous (3DOM) Nb₂O₅ and TiNb₂O₇ composed of interconnected single-crystalline nanoparticles were synthesized [196,197]. This architecture facilitates Li⁺ intercalation and fast electron transfer, enabling high-performance lithium-ion pseudocapacitive behavior and excellent electrochemical performance. Additionally, various doping strategies have been developed to enhance the lithium intercalation kinetics of TiNb₂O₇ materials. These include Nb self-doping to introduce defects [198], Tb doping to widen crystal surface distances [199], Ce doping to induce crystal distortion [200], and Zn-Nb co-doping to improve longer cycling stability [201]. To enable the operation of Nb-based materials at low temperatures, a heterostructure-induced built-in electric field was created by selectively nitriding TiNb₂O₇, enabling fast-charging of TiNb₂O₇-based lithium ions batteries at −50 ℃ [202]. Additionally, the active electronic states of TiNb₂O₇ are modulated by dopants and oxygen vacancies, which enhance its low-temperature dynamics under high-mass loading [203]. Furthermore, Lou’s group is actively pursuing the industrialization of TiNb₂O₇. Their results demonstrate promising cycling life, low-temperature performance, and excellent safety in TiNb₂O₇-based LIBs.

  2.3 Electrolyte chemistry

  2.3.1   Carbonate-based electrolytes

  (1) Historical development of carbonates as lithium battery solvents

  Carbonate electrolytes have been the cornerstone of lithium-ion battery systems since their commercialization, primarily due to their outstanding electrochemical stability [204]. Initially, carbonate solvents such as propylene carbonate (PC) and EC were employed. However, the poor compatibility of PC with graphite anodes led to its replacement by binary and multi-component solvent systems [205,206]. Due to the high dielectric constant, EC was retained and blended with low-viscosity carbonate-based solvents to enhance ionic conductivity [207,208].

  To address the demand for cycling stability of Li-metal anodes, fluorinated carbonate solvents have gained significant attention. Fluoroethylene carbonate (FEC), known for its excellent stability with lithium metal and silicon-based anodes, facilitates the formation of a LiF-rich solid electrolyte interphase (SEI) and has been extensively studied and applied [209,210]. In recent years, other fluorinated derivatives, such as difluoroethylene carbonate (DFEC) and methyl 2,2,2-trifluoroethylcarbonate (FEMC) have garnered attention for their superior performance under high voltage conditions, pointing to new directions in next-generation electrolyte development [211–213].

  (2) Classification and characteristics of carbonates

  Carbonates are broadly classified into cyclic and linear types. Cyclic carbonates, including EC, PC, and FEC, exhibit high dielectric constants, enabling efficient lithium salt dissolution, but their high viscosity limits ionic conductivity [214]. In contrast, linear carbonates such as dimethyl carbonate (DMC), diethyl carbonate (DEC), and ethyl methyl carbonate (EMC) are characterized by low viscosity and excellent flow properties [215]. These are typically combined with cyclic carbonates to achieve a balance between ionic conductivity and electrolyte stability, forming binary or multi-component solvent systems.

  (3) Recent advances in carbonate-based electrolytes

  Recent efforts have focused on enhancing carbonate-based electrolytes to meet the dual demands of high energy density and long cycle life. Strategies include fluorinated modifications and synergistic additive effects to improve high-voltage performance and stabilize the electrode/electrolyte interface [216].

  Recent advances in the fluorinated carbonate solvents have significantly improved the stabilization of the electrode interface. For instance, density functional theory (DFT) calculations of the HOMO/LUMO energy levels of fluorinated solvents were employed to identify high-energy fluorinated carbonate solvents (Fig. 5a). An all-fluorinated electrolyte containing ETFEC and FEC was found to promote rapid Li⁺ desolvation and facilitate the formation of a stable, LiF-rich interface while effectively suppressing lithium dendrite growth. This enabled Li||NCM811 batteries to achieve 225 cycles with an average Coulombic efficiency of 99.8% [217]. The synergistic effects of fluorinated carbonates and other functional additives show promise in enriching the inorganic composition of the SEI and cathode–electrolyte interface (CEI), further enhancing the stability of carbonate-based lithium metal batteries. For example, a carbonate electrolyte enhanced with a dual-salt additive (LiTFA-LiNO₃) was developed to stabilize high-voltage lithium metal batteries (Fig. 5b). In this system, TFA⁻ preferentially interacted with trace water rather than PF₆⁻, thereby improving the moisture tolerance of the electrolyte. Simultaneously, NO₃⁻ became significantly enriched at the cathode interface during charging, forming a Li⁺-rich, solvent-coordinated, thermodynamically favorable double layer [218]. Building on these insights, researchers formulated a fluorinated carbonate-based "cocktail" electrolyte by leveraging the synergistic effects of multiple additives (Fig. 5c). This electrolyte promoted the enrichment of LiF and Li₃PO₄ at the electrode/electrolyte interface, effectively suppressing interfacial side reactions and accelerating interfacial reaction kinetics. As a result, it enabled ultra-fast charging under high voltage conditions of 4.6 V [219]. These findings underscore the critical role of carbonate-based electrolytes in future high-energy-density batteries, highlighting their potential to meet the demands of next-generation energy storage systems. These advances underscore the continued importance of carbonate-based electrolytes in future high-energy-density battery systems.

  Figure 5

  

  Figure 5.  (a) The structures and the calculated HOMO/LUMO energies of the solvents. Reproduced with permission [217]. Copyright 2023, Elsevier. (b) Schematic illustration of the reinforced mechanism of the LiTFA−LiNO₃ dual-salt additive on conventional carbonate electrolyte. Reproduced with permission [218]. Copyright 2024, Wiley-VCH. (c) Schematic illustration of the mechanism on the stabilized 4.6 V LCO batteries. Reproduced with permission [219]. Copyright 2024, Royal Society of Chemistry.

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  2.3.2   Ether-based electrolytes

  Ether-based electrolytes are one of the main organic electrolytes for LIBs. When the first Li intercalation cathode was reported by Whittingham in 1972, the ether-based electrolyte-tetrahydrofuran (THF) and dimethoxyethane (DME) mixture was used [220]. In 1985, as the graphite anode was developed for LIBs, electrolyte decomposition, and severe "solvent co-intercalation" led to graphite exfoliation which hindered the development of ether-based electrolytes [221].

  Though the compatibility of some ether-based electrolytes with graphite anode is usually poor, compared with other organic electrolytes, ether-based electrolytes have low melting point and viscosity, and the ether-based solvents have stronger lithium-ion solvation ability, which is conducive to the design of LIBs with high conductivity and good low-temperature performance [222]. Meanwhile, the reduction stability of ether solvents is also excellent than other organic solvents. While matching with the lithium metal anode, the ether electrolyte possesses good stability [223]. Undoubtedly, the development of the next-generation LIBs requires advanced ether-based electrolytes. In recent years, many achievements have been made in the development of advanced ether-based electrolytes.

  The interaction between solute and solvent in electrolytes leads to the recombination of solvent and solute molecules into different solvation complexes, commonly referred to as solvation structures. The different solvation structures will affect the properties of the bulk electrolyte and the electrode/electrolyte interface chemistry, thereby affecting the electrochemical performance of LIBs [224]. Zhang et al. [225] proposed that the localized high-concentration electrolytes formed by adding bis(2,2,2-trifluoroethyl)ether (BTFE) diluent to HCE (4.5 mol/L LiFSI DME) can maintain the anion-derived solvation structure of HCE. The solvation structure promotes the rapid migration of anions to the graphite surface for decomposition, forming an SEI layer containing inorganic components such as LiF, Li₂O, Li₂S, Li₃N, and Li₂S₂O₄ which effectively inhibits the co-intercalation of solvent molecules into the graphite electrode (Fig. 6a). As a result, the graphite electrode cycled 200 cycles at 4 C in this LHCE electrolyte with a capacity retention rate of 85.5%. Zhao et al. [226] used 1,2-dibutoxyethane (EGDE) as a solvent to dissolve LiPF₆ salt. Due to the large steric hindrance of end-on butyl groups, the coordination ability of EDGE solvent with Li⁺ is significantly weakened, resulting in more contact ion pairs (CIPs) and ionic aggregates (AGGs) in the electrolyte. Compared with other linear ether-based electrolytes, electrolytes using EGDE have better compatibility with micro silicon (mSi) anodes which have a capacity of 1901 mAh/g after 500 cycles, and the Coulombic efficiency is as high as 99.92%. Dong et al. [227] introduced 1,3-dioxolane (DOL) and lithium bis(trifluoromethanesulfonyl)-imide (LiTFSI) as functional additives into a weak solvent electrolyte (1 mol/kg LiFSI DX) (Figs. 6b and c). The optimized dual salt electrolyte has an ionic conductivity of 2.26 mS/cm and a high Li⁺ transference number of 0.77 (Fig. 6d). The graphite anode exhibits excellent fast charging performance (210 mAh/g at 5 C) and long cycling performance (600 cycles with a capacity retention rate of 82.0% at room temperature) in this electrolyte.

  Figure 6

  

  Figure 6.  (a) Schematic of the contrast of Li⁺ insertion process between LHCE (1.5 mol/L LiFSI DME-BTFE) and HCE (4.5 mol/L LiFSI DME) in graphite layers. (b) The electrostatic potential maps distributions of DX and DOL. (c) Density functional theory calculation of the binding energy of Li⁺-DX, Li⁺-DOL, Li⁺-FSI⁻ and Li⁺-TFSI⁻. (d) Ionic conductivity and t_Li⁺ of LD (1.0 mol/kg LiFSI in DX), LDD (1.0 mol/kg LiFSI in the mixture of DX and DOL), and LTDD (0.75 mol/kg LiFSI and 0.25 mol/kg LiTFSI in the mixture of DX and DOL) electrolytes. (a) Reproduced with permission [225]. Copyright 2020, Wiley-VCH GmbH. (b-d) Reproduced with permission [227]. Copyright 2024, Elsevier.

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  In summary, the properties of ether-based electrolytes can be effectively optimized by regulating the solvation structure of electrolytes. However, some strategies will sacrifice the advantages of ether-based solvents, such as high ionic conductivity and high reduction stability. The mechanism by which optimized ether electrolytes can enhance the electrochemical performance of graphite and silicon anodes still needs to be further explored in conjunction with advanced in-situ characterization. In the future, it is necessary to combine theoretical calculations and machine learning to obtain more information on the solvation structure of ether-based electrolytes, which is of great significance for the development of electrolytes. Therefore, to develop more practical high-performance ether-based electrolytes, many problems still need to be solved.

  2.3.3   Phosphate-based electrolytes

  Different of carbonate and ether-based solvents, phosphate is generally not used as primary or sole solvent in lithium-ion or lithium metal batteries. Instead, they are typically employed as additives or low-proportion co-solvents [228]. Because Li⁺ of the low concentrations in phosphate-based electrolytes tend to form strong coordination solvation structures with the phosphate, which increases the difficulty of Li⁺ desolations [229]. It ultimately leads to the co-intercalation of the phosphate solvent into graphite, causing graphite exfoliation and continuous electrolyte decomposition at the interface, or it may result in uneven lithium deposition to form lithium dendrites. In the solvation sheaths, the polarized solvent molecules induced by the electrostatic field of Li⁺ are more likely to accept electrons and decompose on the anode surface as their electrophilicity increases. To address the incompatibility between phosphate and electrodes, the general approaches are increasing the salt concentration or constructing the localized high concentration electrolyte (LHCE) to promote coordination between Li⁺ and the anions or to introduce co-solvents or additives to form a stable solid electrolyte interphase (SEI). Fan et al. [230] used 1,1,2,2,3,3,4-heptafluorocyclopentane (HFCP) as diluent to construct LHCE with 1.6 mol/L LiFSI in trimethyl phosphate (TMP) and HFCP with a molar ratio of 1:3. The mesocarbon microbeads (MCMB)||NMC811 full cells based on LHCE not only achieved stable electrochemical cycle, and it had the high capacity retention of 90% over 1300 cycles at 0.5 C. Similar to the fluorinated carbonates, phosphate has flame-retardant properties because most of them have intrinsic non-flammability characteristics [231]. Generally, when the weight fraction of phosphate exceeds 40%, the electrolyte can achieve the zero self-extinguishing time. Additionally, some phosphates also have the ability to scavenge free radicals, thereby blocking the combustion reaction. Typically, volatile carbonate solvents decompose easily and produce highly reactive H· and HO· free radicals when burned, while some phosphate such as ethoxy(pentafluoro)cyclotriphosphazene (PFPN) can decompose into phosphorus radicals in high-temperature environments. These phosphorus radicals can capture the highly active H^· and HO^· radicals via van der Waals forces, releasing reactive free radicals that effectively suppress subsequent free radical chain reactions, preventing the continuation of combustion. Therefore, PFPN is often used as an additive in non-flammable electrolytes (Figs. 7a and b).

  Figure 7

  

  Figure 7.  Phosphate-based electrolyte. Flame-retardant test in the (a) blank electrolyte without PFPN (flammable), (b) PFPN-based electrolyte (non-flammable). (c) Cycling performance of Na||Na₃V₂(PO₄)₂O₂F cells in PFPN-based electrolyte. (d) Schematic illustration of interfacial chemistry and corresponding solvation sheath of the Tri-anion regulated phosphate-based electrolyte. (e) Long-term cycling of NCM||Li cell in CFTN electrolyte at 0.1 C. (f) Schematic of TDNE electrolyte solvation structure. (g) Cycling performance of Li||NCM 622 batteries using TDNE electrolytes at 1 C. (h) Long-term cycling of the HC||Na₄Fe_2.91(PO₄)₂(P₂O₇) pouch cell at 1 C. (a-c) Reproduced with permission [232]. Copyright 2024, Wiley-VCH. (d) Reproduced with permission [233]. Copyright 2024, Elsevier. (e) Reproduced with permission [234]. Copyright 2022, Wiley-VCH. (g) Reproduced with permission [235]. Copyright 2023, American Chemical Society. (h) Reproduced with permission [236]. Copyright 2024, American Chemical Society.

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  More importantly, PFPN has electron-withdrawing F atoms and electron-donating -P=N- group, which can interact with solvents containing lone electrons pairs or electron-deficient atoms to regulate the solvation structure of Li⁺ and form SEI dominated by fluorides and nitrides, improving electrochemical performance. As showed in Fig. 7c, Ma et al. [232] incorporated PFPN as the additive into carbonate-based electrolytes, which not only imparted non-flammable properties but also significantly improved the cycling performance of the 4.5 V high-voltage Na||Na₃V₂(PO₄)₂O₂F battery. After 2000 cycles, the capacity retention rate was 92.4%, with the coulombic efficiency (CE) of 99.71%. In addition to the non-flammable advantages, phosphate is also frequently used to enhance the solubility of nitrates because of their strong coordination structure. Feng et al. [233] dissolved 0.3 mol/L lithium nitrate (LiNO₃) in triethyl phosphate (TEP), and added 1 mol/L LiTFSI and 0.3 mol/L LiDFOB to form a tri-anion strategy to optimize the solvation structure (Fig. 7d). This approach simultaneously stabilized the anode, cathode, and phosphate-based electrolyte, achieving stable electrochemical performance. Similarly, Guo et al. [234] utilized the strong solvating ability of TMP to dissolve 3 mol/L LiNO₃, combining it with the carbonate system and denoted as CFTN. As a result, the NCM||Li cell retained the capacity retention rate of over 80% after 100 cycles at 0.1 C based CFTN electrolyte (Fig. 7e). Besides being used as co-solvents and additives, phosphate can also serve as the main solvent in electrolytes, provided that certain additives are included to adjust the solvation structure and SEI layer. Zheng et al. [235] dissolved 1.5 mol/L LiTFSI in pure TEP and added 0.47 wt% LiNO₃ and 4.53 wt% DME as additives, referred to as TDNE. As showed in Fig. 7f, both LiNO₃ and DME entered the solvation sheath of TEP, synergistically adjusting the solvation structure and preferentially decomposed to form a dense electrode/electrolyte interphase on the lithium anode and NCM cathode. This led to reversible lithium plating/stripping kinetics while suppressing the structural degradation of the NCM cathode. This SEI layer also effectively suppressed side reactions between the electrode and TEP solvent. Based on TDNE, Li||NCM 622 cells exhibited a capacity retention rate of 98.41% after 500 cycles at 1 C when charging to 4.6 V (Fig. 7g). Due to the fact that anisotropic hard carbon (HC) anodes are not affected by the co-intercalation of phosphate, phosphate are also widely used in sodium-ion battery systems. Cao et al. [236] developed a fully phosphate-based electrolyte named MF132 with a molar ratio of 1:3:2 of NaClO₄: TMP: Tris(2,2,2-trifluoroethyl)phosphate (TFEP), which was completely non-flammable. The pouch cells based MF132 exhibited an average CE greater than 99.9% after 2000 cycles at 1 C, with a capacity retention rate of 84.5% (Fig. 7h).

  2.3.4   Sulfone-based electrolytes

  Compared to the high volatility and flammability of conventional carbonate electrolytes, sulfone-based electrolytes show high flash point and low flammability [237,238]. Additionally, due to their low HOMO levels, sulfone solvents demonstrate superior oxidative stability and a broader stable electrochemical window [239,240], which can enhance the high-voltage stability of electrolyte [241,242]. For example, Feng et al. [243] designed a dispersive aggregated electrolyte, in which the sulfolane (SL) solvent enables the electrolyte a decomposition voltage of 5.5 V. Dong et al. [244] designed a sulfone-based electrolyte (1 mol/L LiTFSI in tetramethylene sulfone (TMS) and FEC) with affinity for double electrodes. The oxidation-resistant TMS preferentially adsorbs on the surface of the cathode, preventing the oxidation of the FEC and forming a LiSO₂F-rich interfacial layer, which contributes to the homogenized Li⁺ transportation. The 4.4 V Li||NCM811 cell exhibits 86.1% capacity retention after 500 cycles. SL combining with LiDFOB triggers the formation of CEI composed by Li_xSO_y, Li_xBO_y, realizing LiNi_0.5Mn_1.5O₄ (LNMO) cathode operated at 60 ℃ [245]. As additive or artificial interface layer, sulfone compounds can also benefit to the thermal stability and interfacial stability of electrolytes [246–248].

  Unfortunately, most of the sulfone compounds usually show high viscosity and poor wettability [249], and the high electron density of oxygen atom increases the reactivity of sulfone with the lithium metal, resulting in forming an unstable SEI [250,251]. To further enhance the utility of sulfone solvents, different strategies have been proposed. Tang et al. [252] enhanced the stability of ethyl vinyl sulfone (EVS) at electrode-electrolyte interface by subsequent decomposition of vinylene carbonate (VC) additives, which promoted long-term stable cycling of LMBs. The usage of fluorinated sulfone-based electrolytes is another effective approach. β-Fluorinated sulfones 1,1,1-trifluoro-2-(methylsulfonyl)ethane (TFEMS), in which the sulfonyl group is one carbon atom away from the strong electronegative fluorine group, shows good antioxidant property and stability against graphite anode due to its low reduction potential. Such electrolyte enables the Gr||NCM622 battery with 71% capacity retention after 400 cycles at high cut-off voltage of 4.5 V [253]. Ren et al. [254] introduced the diluent 1,1,2,2-tetrafluoroethyl-2,2,3,3-tetrafluoropropyl ether (TTE) to form a localized high-concentration TMS electrolyte. Except reducing the viscosity, the diluent inhibits the corrosion of the Al collector, realizing the 4.9 V level battery operating at −10 ℃. Lithium salt also affects the properties of sulfone-based electrolyte, the introduction of LiNO₃ changes the solvated structure of highly concentrated SL and promotes more anions to coordinate with Li⁺. As the result, the LiN_xO_y and LiF dominated CEI increase the stability of NMC811, showing negligible capacity decay over 200 cycles [255].

  2.3.5   Nitrile-based electrolytes

  Nitrile compounds commonly show excellent resistance to oxidation under high voltage conditions due to the extremely low highest occupied molecular orbital (HOMO), which have been widely used as solvent or cosolvent in lithium batteries [256–258]. The oxidative stability of nitrile-containing solutions can even reach up to 7.8 V vs. Li⁺/Li [259]. In addition, it is demonstrated that the -CN group could preferentially adsorb onto transition metals of cathode, forming a -CN-TM complex. This interaction prevents other substances contacting with the cathode, thereby suppressing the dissolution of transition metals and enhancing the cyclability and stability of the cathode [260,261]. Due to the strong nucleophilic character of -CN, nitriles exhibit high coordination interaction with Li⁺, thus facilitate to a high solubility of lithium salt and advanced ionic conductivity of electrolyte [256,262]. Compared to commercial carbonate electrolyte (ethylene carbonate/ethyl methyl carbonate, EC/EMC), the modified carbonate electrolyte with 60 vol% isobutyronitrile (IBN) cosolvent not only reduces the viscosity but also facilitates the dissolution of lithium salt, releasing more free lithium ions, which delivers an impressive ionic conductivity of 1.152 mS/cm at −70 ℃ [263]. Furthermore, acetonitrile (AN), acting as a free solvent molecule, induces rapid lithium ion hopping between solvation shells in the FEC-dominated lithium-ion solvation structure, doubling the ionic conductivity to 12 mS/cm [264].

  However, the strong electron-withdrawing effect of the nitrile group makes the adjacent α-H acidic, and the severe decomposition of -CN group is prominent at low potential, which is hard to form a stable SEI film, limiting the application in lithium metal batteries [265,266]. The most common solution is to add film-forming agents which preferentially deposit on the lithium anode and form stable SEI, inhibiting side reactions between nitrile compound and the lithium metal [267]. For instance, adding LiDFOB in succinonitrile (SN) based deep eutectic electrolyte facilitates the formation of anion-derived SEI and improves the stability of the electrolyte, realizing 4.5 V Li||LiCoO₂ cell a 77% retention after 750 cycles [268]. The reduction activity of SN can also be restrained by introducing 1,3,5-trioxane (TXE). This is because that the unstable α-H of SN strongly coordinates with the ether oxygen functional groups [269]. Zhao et al. [270] used dimethylmalononitrile (DMMN) whose two α-H were replaced by methyl groups to avoid the intrinsic instability, the as-obtained electrolyte using LiTFSI as lithium salt achieved the long lifespan for Li||LiFePO₄ cell with almost no capacity decay over 750 cycles. Thanks to the high solubleness of lithium salt in nitrile electrolyte, forming highly concentrated electrolytes (HCEs) and local highly-concentrated electrolytes (LHCEs) is an efficient strategy [271]. In LHCEs, the diluent, such as 1,1,2,2-tetrafluoroethyl-2,2,3,3-tetrafluoropropyl ether (TTE), can be used to adjust the number of free solvent molecules and hence restricts the solvent-based side reactions [272].

  2.3.6   Ionic liquids electrolytes

  Ionic liquid is a new kind of room-temperature molten salt with a low molten point below 100 ℃, which is composed of organic cations and inorganic anions [273,274]. The typical organic cations include imidazolium cations, pyrrolidinium cations, piperidinium cations, tetraalkylammonium cations and tetraalkylphosphonium cations. Based on the type of cations, ionic liquid can be classified into several types including imidazolium-based ionic liquid, pyrrolidinium-based ionic liquid, piperidinium-based ionic liquid, tetraalkylammonium-based ionic liquid and tetraalkylphosphonium-based ionic liquid.

  Compared with conventional organic solvents, ionic liquid has higher thermal stability, low vapor pressure and most importantly the non-flammability. These unique characteristics are completely different from the organic solvents and inorganic salts. Owing to the unique physicochemical properties, ionic liquid showcases great promise as the electrolyte compositions for Li-based batteries such as Li-S batteries or Li-air batteries. They can be either directly used as the electrolyte solvent to replace carbonate solvents, which effectively improve the safety of high-voltage Li-based batteries because of their high thermal stability and non-flammable characteristics. Unfortunately, the high viscosity and high cost of ionic liquids hinder their applications as solvent for electrolyte. Therefore, ionic liquids usually serve as the additives in conventional electrolytes with low viscosity to obtain organic electrolytes with both excellent ionic conductivity and improved safety.

  Typically, the energy gaps between highest occupied molecular orbital (HOMO) and lowest unoccupied molecular orbital (LUMO) of ionic liquid is narrower than the organic solvents. Therefore, ionic liquid is able to be decomposed earlier on both cathodes and anodes to form solid electrolyte interphase (SEI). Especially, the HOMO energy level of most ionic liquids is higher than some Li salts, indicating that the ionic liquid is more effective to be decomposed on the cathode surface to form the anion-rich cathode electrolyte interphase (Fig. 8a) [275]. Hence various reports have demonstrated the applications of ionic liquids in stabilizing the high-voltage cathode of Li-ion batteries. The solvation structures of ionic liquid electrolyte are summarized in Fig. 8b [276]. As shown in Fig. 8b, compared with conventional diluted electrolyte, the anions such as FSI⁻ in the first solvation sheath of ionic liquid electrolyte largely increases. In addition, even when fluorinated diluent is added into the electrolyte, the anions in the first solvation sheath does not decrease. Hence, anion-derived inorganic rich solid electrolyte interphase can be formed on both cathode and the anode with ionic liquid electrolyte, which can stabilize the structure of LiCoO₂ cathode [276]. As a result, improved cycle stability of Li||LiCoO₂ full cells can be achieved with the ionic liquid modified electrolyte [276]. In addition to add co-solvent to decrease the viscosity of the electrolyte, designing new ionic liquid with low viscosity is also a strategy for the practical application of ionic liquid electrolyte. For example, Dai et al. [277] designed a new kind of ionic liquid electrolyte containing 5 mol/L LiFSI and 0.16 mol/L NaTFSI in EMIMFSI. Compared with other ionic liquid, the EMIM cations based ionic liquid has relatively low viscosity and higher ionic conductivity. Hence, it can be directly used as the solvent for Li metal batteries (Fig. 8c). In this ionic liquid electrolyte, the both the imidazole cations and the anions can be decomposed on both LiCoO₂ cathode and Li metal anode to form inorganic rich SEI layer and CEI layer, thus not only stabilizing the structure of LiCoO₂ cathode but also suppress the Li dendrite growth. Consequently, the Li||LiCoO₂ full cell displays excellent cycle life over 1200 times with high-capacity retention of 81%. More importantly, the novel ionic liquid electrolyte is intrinsically non-flammable, which successfully address the safety issue of Li metal batteries [277].

  Figure 8

  

  Figure 8.  Ionic liquid-based electrolyte for Li based batteries: (a) The frontier molecular orbital energy level of different kinds of ionic liquids. Reproduced with permission [275]. Copyright 2021, Elsevier. (b) Solvation structure of ionic liquid-based electrolyte. Reproduced with permission [276]. Copyright 2020, Wiley-VCH. (c) Schematic illustration of the interphase of Li||LiCoO₂ full cell with novel ionic liquid electrolyte. Reproduced with permission [277]. Copyright 2020, Wiley-VCH. (d) Schematic illustration of the interphase of NCM811 cathode with ionic liquid electrolyte. Reproduced with permission [278]. Copyright 2024, Wiley-VCH. (e) Schematic illustration of Li||NCM811 full cell in the ionic liquid modified electrolyte. Reproduced with permission [279]. Copyright 2023, Wiley-VCH. (f) Synthesis procedures of difluorinated ionic liquid for Li metal batteries. Reproduced with permission [280]. Copyright 2024, Wiley-VCH.

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  In addition to LiCoO₂ cathode, ionic liquid can be also decomposed on LiNi_0.8Co_0.1Mn_0.1O₂ (NCM811) cathode to form robust CEI layer to stabilize the structure of NCM811 cathode. For example, Tang et al. [278] reported a type of anion-reinforced solvation structure realized by Py14FSI based ionic liquid additives (Fig. 8d), which achieves a combination of the high ionic conductivity of conventional electrolytes with a solvent-dominant solvation structure and the advantageous interfacial stability characteristic of ionic liquid with an anion-dominant solvation structure. The ionic liquid in the electrolyte can form an exceptionally stable CEI layer on the NCM811 cathode surface, thus not only preventing the detrimental interactions between NCM811 and the electrolyte, but also effectively inhibiting irreversible phase changes and mitigating the occurrence of microcracks in the secondary particle structure of NCM811. As a result, improved cycle life of Li||NCM811 full cell are successfully achieved with ionic liquid modified electrolyte. Owing to the high ionic conductivity of ionic liquid modified electrolyte as well as the ionic liquid derived inorganic-rich CEI and SEI layer on both NCM811 cathode and Li metal anode, the Li||NCM811 full cells exhibit improved power density. Recently, Yuan et al. [279] reported an ionic liquid modified electrolyte for high-voltage Li||NCM811 full cells (Fig. 8e). Owing to the anion enriched solvation structure with high ionic conductivity and the inorganic CEI and SEI layers, the Li||NCM811 full cells showcase excellent fast charging capability. Therefore, even with limited Li metal in the anode, the Li||NCM811 full cells still exhibit long-term cycle stability with high current density [279].

  As the ionic liquids are composed of organic cations and inorganic anions, the structure of organic cations can be designed to satisfy the practical application. Wang et al. [280] designed a novel middle-concentrated ionic liquid electrolyte with an anion-rich solvation structure tuned by difluorinated cations for Li||NCM811 batteries (Fig. 8f). Compared with conventional ionic liquid, the difluorinated cations in the designed ionic liquid can be decomposed on Li metal anode and NCM811 cathode to form LiF-rich solid electrolyte interphase. Therefore, the ionic liquid modified electrolyte not only improves the cycle life of NCM811 full cell but also enhances the safety of Li||NCM811 batteries. In recent year, a series of fluorinated cations based ionic liquid [281,282] have been explored for Li||NCM811 batteries, which effectively prolong the cycle life of the Li||NCM811 batteries owing to the highly fluorinated interphase derived from the cations in the ionic liquids.

  Above all, owing to the unique properties, ionic liquids have successfully applied either as electrolyte solvents or the electrolyte additives for high-voltage Li based batteries. However, it should be noted that ionic liquids are not compatible with the graphite anode in Li-ion batteries due to the co-insertion of large organic cations and Li-ion into the layer of graphite, which could deform the layered structure of graphite anode. Hence, ionic liquids can be only applied for Li metal batteries or Li-ion batteries with silicon anode.

  2.3.7   Polymer-based solid-state electrolytes

  The practical applications of solid polymer-based electrolytes (SPEs) in lithium-metal batteries (LMBs) face two key problems: (1) Unsatisfying physical contact and chemical stabilities of the interfaces between electrodes and SPEs; (2) Restrictive relationship between ionic conductivities and mechanical properties of SPEs. The former causes contact losing and continuous irreversible reactions at the interfaces, thus resulting into the rapid deterioration of rate capability and cycling performance for the high-voltage LMBs [283]. The latter causes the poor inhibition effects of lithium dendrite problems, which increase the possibility of battery short circuit and preparation difficulties of the ultra-thin SPEs, limiting the further improvement of the energy density for the LMBs [284,285]. In this context, researchers have recently proposed new strategies to obtain the SPEs with both high mechanical strength and ionic conductivity, and effectively solve the interface issues at the same time, realizing the applications of SPEs in high-performance LMBs.

  Peng et al. [286] prepared a new lithium superionic conductor (denoted Li-HA-F). Different from the previously reported superionic conductor with bulk structures, Li-HA-F had an ultralong nanofiber structure and ultrahigh room-temperature ionic conductivity (12.6 mS/cm) with good ambient stability (Figs. 9a-c). When this Li-HA-F nanofiber membrane was coupled with poly(ethylene oxide) (PEO)-based solid electrolyte, it enabled the resulting SPE (denoted Li-HA-F CSE) to show high ionic conductivity (4.0 × 10^–4 S/cm at 30 ℃), large Li⁺ transference number (0.66), high breaking strength (9.66 MPa) and wide voltage window (5.2 V) (Figs. 9d and e). It was revealed that the Li-HA-F supplies continuous dual-conductive pathways and causes stable LiF-rich interfaces, leading to excellent performance (Fig. 9f). As a result, the Li/Li half cells with the Li-HA-F SPE showed a high critical current density of 1.4 mA/cm². The LiFePO₄/Li-HA-F CSE/Li and LiNi_0.8Co_0.1Mn_0.1O₂/Li-HA-F CSE/Li solid-state full cells also exhibited good performance in a wide temperature range.

  Figure 9

  

  Figure 9.  (a) Schematic illustrations of the synthetic process for the Li-HA-F nanofibers and Li-HA-F CSE with lithium conductive mechanism and Radar plots which compare the comprehensive performance of nongel SPEs with different fillers. (b) Surface SEM image of the Li-HA-F membrane. (c) Comparison of typical lithium superionic conductors with the Li-HA-F nanofiber. (d) Mechanical strength of the Li-HA-F CSE and PEO-LiTFSI (inset: the optical image of the Li-HA-F CSE at bending state). (e) LSV curves for the Li-HA-F CSE and HA CSE. (f) Binding energy calculation of the (002) surface of HA with and without F doping. (g) Schematic diagram of the design principle for the EDTA-pH PSE. (h) Cross-section SEM image of the EDTA-pH electrolyte. (i) Stress-strain curves of the EDTA-pH and pH electrolyte membranes. (j) Stress and single-edge crack simulations of the EDTA-pH electrolyte. (k) The calculated binding energy of the Li⁺ or LiTFSI^- with the nearest solvent molecules (for the pH electrolyte) or polymer chain (for the EDTA-pH electrolyte). (l) Schematic illustration of the vapor-phase fluorinated approach. (m) The cross-section scanning electron microscope (SEM) image of the F-NBR-g-VEC matrix and corresponding elemental mapping images of C, F, and O (inset: 3D rendering of C-, F-, and O- secondary ion distributions of F-NBR-g-VEC electrolyte). (n) The 3D rendering of LiF^-, Li⁺, and O^- secondary ion distributions on the cycled lithium anode with the F-NBR-g-VEC electrolyte. (o) Schematic illustration of ultraconformal interface between lithium anode and the F-NBR-g-VEC electrolyte. (a-f) Reproduced with permission [286]. Copyright 2024, American Chemical Society. (g-k) Reproduced with permission [287]. Copyright 2024, Wiley-VCH. (l-o) Reproduced with permission [290]. Copyright 2023, Wiley-VCH.

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  Without using inorganic fillers, Liu et al. [287] proposed an efficient strategy for the synthesis of an ultrathin (~5 µm) SPEs (denoted EDTA-PH) with ultrahigh mechanical strength (Figs. 9g and h). As an additive, ethylene diamine tetraacetic acid (EDTA) with electron-donating property induced the conformation transformation of polyvinylidene fluoride-hexafluoropropylene (PVDF-HFP), realizing an utrahigh Young’s modulus (10.6 GPa) by a fine-grain strengthening mechanism (Figs. 9i and j). Moreover, the cis-conformation of PVDF-HFP shortened the Li⁺ pathway and promoted the Li⁺ dissociation and anion immobilization (Fig. 9k), leading to high ionic conductivity (2.47 × 10⁻⁴ S/cm) and transfer number (0.59) of the EDTA-PH electrolyte. The EDTA additives also resulted into the LiF-enriched interfaces between the electrodes and electrolyte with a wide voltage window of 4.7 V. Both half cells and full cells with the electrolytes exhibited good cycling and rate performance. It should be noted that a pouch cell using the ultra-thin electrolyte delivered high energy densities of 516 Wh/kg and 1520 Wh/L (excluding packages).

  The matrixes for traditional SPEs (e.g., PEO and PVDF) have insufficient mechanical resilience to accommodate the volume change of lithium anodes during charge–discharge processes [288,289]. Therefore, the interfacial contact issues cannot be essentially solved. Yang et al. [290] proposed an ultra-conformal SPE (denoted F-NBR-g-VEC) using rubber-derived elastomer as new matrix. To further maintain the interface contact between electrodes and electrolyte, the surface of the rubber-derived electrolyte was fluoridated by a chemical vapor-phase method, which was transformed into a resilient, ultrathin, and mechanically integral LiF-rich layer after cycling in cell (Figs. 9l-n). The ultra-conformal layer chemically bonded the electrolyte and lithium anode and kept their dynamic contact, thus maintaining long-term rapid and stable Li⁺ interfacial transport. Furthermore, uniform lithium deposition and suppressed side reactions were also realized (Fig. 9o). Consequently, the LMBs with the F-NBR-g-VEC electrolyte showed a good cycling performance of 2500 h in half cells and exhibited a good stability in full cells.

  2.3.8   Inorganic solid-state electrolytes

  Organic solid-state electrolytes (SSEs) have been attracting more and more interests, since they have unique advantages such as high chemical and thermal stability, high safety, wide electrochemical window, high conductivity and Li⁺ transference number. Among the various types of inorganic SSE, oxide electrolytes, sulfide electrolytes and novel halide electrolytes are the main families toward the application in solid-state batteries [291,292].

  Oxide SSEs mainly includes Perovskite type Li_3xLa_2/3-xTiO₃, Garnet type Li₇La₃Zr₂O₁₂, NASICON type Li_1+xAl_xM_2–3(PO₄)₃, LISICON type Li₁₄ZnGe₄O₁₆, etc. Normally, oxide SSEs exhibit a wide electrochemical window and good chemical/electrochemical stability. However, the lower ionic conductivity contributed by huge grain boundary resistance, poor mechanical properties, and the high-temperature and high-pressure fabricating process during battery fabricating and packaging pose challenges to the practical application of oxide SSEs [293,294]. Some strategies are usually adopted to promote the development of solid-state lithium batteries based on oxide SSEs, such as composite electrolyte with oxide and polymer, in-situ solidification [295,296]

  Sulfide SSEs show great potential for practical application in solid state batteries [297,298]. Firstly, sulfide SSEs have super-high ionic conductivity, usually 1 mS/cm at room temperature, and even higher [299,300]. Additionally, sulfide SSEs have lower Young’s modulus, which enable simple fabrication processing such cold pressing to achieve grain boundary contacting both in electrolytes and between electrolytes and electrode. However, sulfide SSEs have a narrow electrochemical stability window, influencing the interface stability with electrodes. Especially at the interface between sulfide SSEs and cathodes, high voltage and delithiated cathode cause electrolyte decomposing [301], space charge effects conduct overpotential [302,303], which result in battery performances decay. Sulfide SSEs also exhibit poor thermal stability [304], solvent stability [305] and air stability [306]. They will undergo structural degradation with toxic H₂S gas releasing when exposed to ambient air or solvent. It requires the sulfide SEs to be synthesized and fabricated under particularly exacting circumstances for safety reason.

  Halide SSEs have garnered widespread attention in recent years as a novel family of superionic conductors and become a promising candidate for solid-state lithium batteries [139,307]. Thy aroused the interest of researchers since 2018 that Li₃YCl₆ and Li₃YBr₆ were reported with elevated ionic conductivities (1.7 mS/cm) [308]. After that, a systematic exploration of halide materials Li₃MX₆ (M = Sc, Y, La–Lu; X = Cl, Br) were conducted and they demonstrated high ionic conductivity (>1 mS/cm at room temperature) and good ductility [139,307]. Following up, O atoms were introduced to the lattice to form halide-oxide SSEs, accompanied with excellent properties, such as super high ionic conductivity of 12 mS/cm at room temperature [309], now crystal structures [310], viscoelasticity [311], and cost-effective synthesis [312]. Meanwhile, halide and halide-oxide SSEs have higher oxidation stability voltage, and hence better interface compatibility with cathode. However, the halide and halide-oxide SSEs show poor reduction stability, which could cause the interface degradation with anode. Currently, halide and halide-oxide SSEs are usually studied as solid electrolytes and as ionic conductive additives for solid state composite cathode [313–315]. Further efforts are crucial important to explore more halide and halide-oxide materials with higher ionic conductivity and wider electrochemical window.

  3. Alkali metal rechargeable batteries

  3.1 Lithium-sulfur batteries

  3.1.1   The challenges of lithium-sulfur batteries

  According to the working mechanism of lithium sulfur batteries, the charging and discharging process involves redox reactions of eight electrons. Therefore, it is expected to replace LIBs as the next generation of high-energy density and high-capacity lithium metal battery systems. But lithium sulfur batteries are still far from commercial production due to the complex sulfur phase transition process during their charging and discharging processes. In some urgent and significant challenges that need to be addressed [316,317]. The specific summary is as follows (Fig. 10):

  Figure 10

  

  Figure 10.  Schematic of the intertwined phenomena inducing the impaired reaction activity and the fast degradation of the sulfur cathode and Li metal anode during cycling. Reproduced with permission [317]. Copyright 2023, Wiley-VCH.

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  (1) Insulation properties of sulfur and Li₂S

  The reaction mechanism of lithium sulfur batteries is relatively complex, but the insulation properties of sulfur and Li₂S on electrons and lithium ions directly limit the electrochemical reaction and affect the utilization rate of active materials [318]. Although the solid-liquid reaction proceeds relatively rapidly in the first discharge platform, achieving a theoretical capacity of 418 mAh/g remains highly challenging. In addition, even after multiple cycles, there is still unreacted sulfur. To achieve uniform dispersion and good electrical contact of sulfur, and to promote the reduction reaction by adding conductive additives, sulfur can be reduced to long-chain polysulfides. However, in lower discharge platforms, the reaction kinetics are slower because the conversion reaction involves a solid-solid process, and the polysulfides formed in the upper discharge platform are difficult to completely reduce to Li₂S. In fact, the theoretical capacity ratio of the second discharge platform to the first discharge platform is expected to be 3:1, but the actual capacity ratio is about 2.5 because sulfur often needs to be mixed into the conductive matrix, which leads to a decrease in specific capacity. In addition, due to the slow kinetics of electrons and lithium ions, typical lithium sulfur batteries exhibit higher polarization at lower discharge plateau voltages, which further leads to a decrease in actual energy density [319–321].

  (2) The shuttle effect of polysulfides

  Lithium sulfur batteries generate polysulfides that are soluble in the electrolyte during the reaction process [322]. The polysulfides shuttle to the anode can cause capacity decay and decrease coulombic efficiency, which is known as shuttle effect [323]. Although the dissolution of polysulfides can cause many problems, it is also crucial in achieving high sulfur utilization efficiency. Sulfur reduction can only occur on the surface of conductive materials because sulfur itself is not conductive. Due to the continuous dissolution of polysulfides, the remaining sulfur is exposed to the conductive matrix, allowing for further reduction. Therefore, it is necessary to balance and control the dissolution of polysulfides to achieve high sulfur utilization and high cycling stability [324]. To address these issues, researchers have proposed many solutions, such as adding covering agents, preparing composite materials, and changing electrolyte composition. The covering agent can encapsulate sulfur particles and prevent their dissolution, thereby improving the stability and performance of the battery. The preparation of composite materials can improve the mechanical strength and electrochemical performance of batteries. Changing the electrolyte composition can alleviate the shuttle effect [325,326].

  (3) Growth of lithium dendrites

  Lithium metal is an ideal negative electrode for high-energy density batteries, with a high specific capacity (3860 mAh/g) and low reduction potential (−3.04 V relative to standard hydrogen electrodes). However, the development of its commercial application is constrained by the following issues. Firstly, due to the uneven distribution of surface current density and the concentration gradient of lithium ions at the electrode/electrolyte interface, metal lithium anode can form lithium dendrites, leading to safety issues such as short circuits. Secondly, during the cycling process, repeated lithium deposition/stripping leads to volume changes and cracking of the lithium anode, forming inactive "dead lithium" and reducing the utilization rate of lithium. Thirdly, metallic lithium has high reactivity and spontaneously reacts with organic electrolytes to form a SEI layer. Due to the large volume changes during lithium deposition/stripping, the instability of its SEI structure can lead to sustained irreversible consumption of active lithium. Finally, due to the existence of the "shuttle effect", a passivation layer Li₂S is formed on the surface of the lithium anode, which significantly reduces its efficiency and limits the cycling stability of the battery. To address these issues, further research and exploration are needed on alternative anodes or measures to protect lithium anodes. For example, using high-capacity silicon anodes and composite anodes instead of metal lithium anodes may become a solution. In addition, changing the electrolyte composition and adding additives to the electrolyte have been widely studied to protect the lithium anode and improve battery performance [327–329].

  3.1.2   Progress of research on sulfur cathodes

  (1) Carbon materials

  Carbon materials have become the first conductive sulfur carrier materials developed and applied to the positive electrode of lithium sulfur batteries due to their excellent conductivity, diverse types, and stable structure [330,331]. At present, carbon nanomaterials used as cathodes for lithium sulfur batteries can be classified into four categories based on their dimensions: Zero dimensional, one-dimensional, two-dimensional, and three-dimensional. According to the pore size, it can be divided into three categories: micropores, mesopores, and macropores. According to their structure, they can be divided into carbon nanotubes, carbon nanospheres, carbon nanofibers, graphene, etc. (Fig. 11) [332].

  Figure 11

  

  Figure 11.  Classification diagram of carbon material as cathode of lithium-sulfur batteries. Reproduced with permission [332]. Copyright 2023, Elsevier.

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  Yang et al. [333] prepared vertically arranged sulfur graphene (S-G) nanowalls on conductive substrates and used them as positive electrodes for lithium sulfur batteries. Sulfur nanoparticles are uniformly fixed between graphene layers, and the vertically ordered graphene nanowall array facilitates the diffusion of electrons and lithium ions. The layered porous structure is conducive to regulating volume expansion. S-G has a high reversible capacity of 1261 mAh/g at a current density of 0.25 C, with a reversible capacity exceeding 1210 mAh/g after 120 cycles. However, due to the uneven pore size of graphene, its adsorption capacity for polysulfides is limited, making it unsuitable for commercial production. Fan et al. [334] prepared sulfur porous carbon nanotube (S-PCNT) composite material as a binder free positive electrode for lithium sulfur batteries. Porous carbon nanotubes can provide better conductivity, rich mesoporous structure can provide certain mechanical strength, and have ultra-high specific surface area and abundant adsorption sites, which can efficiently capture sulfur during the preparation process. S-PCNT has an initial specific capacity of 866 mAh/g at a current density of 0.1 C. After 100 cycles, it can still maintain a discharge capacity of 526 mAh/g. Therefore, S-PCNT can provide excellent rate performance, but in general, carbon nanotubes have a small specific surface area and cannot inject more active substances, which is not conducive to improving the energy density of lithium sulfur batteries. Shao et al. [335] prepared porous carbon spheres (PNCS) with large pores and high specific surface area, combined with nitrogen doped graphene (NC), as an efficient positive electrode sulfur carrier (S@PNCS/NC) for lithium sulfur batteries. The high specific surface area and unique porous structure of PNCS facilitate the injection of active substances and the penetration of electrolyte solutions, while the interconnected open pores can significantly shorten the ion/electron transfer pathway. A highly conductive porous carbon sphere composite nitrogen doped graphene cathode can achieve an ultra-high initial discharge capacity of 1445 mAh/g at 0.2 C.

  (2) Metal compound/sulfur cathode

  Metal oxides are widely used as positive electrode materials for lithium sulfur batteries. The unique O²⁻ anion in metal oxides has a strong polar surface and can be connected to lithium polysulfides through strong polarity chemical interactions. Therefore, applying metal oxides to the positive electrode of lithium sulfur batteries can effectively limit polysulfides to the positive terminal. TiO₂ has attracted widespread attention from researchers due to its diverse morphology, excellent conductivity, and polarity characteristics. Moreover, the binding energy between polysulfides and the α-TiO₂ (101) crystal plane (3.59 eV) is higher than that between polysulfides and carbon materials (<1.0 eV) [336]. Li et al. [337] prepared mesoporous TiO₂ nanocrystal cathodes in situ on reduced graphene oxide (S/TiO₂@rGO) (Fig. 12a). Research has shown that TiO₂ nanocrystals have strong chemical adsorption properties for polysulfides and high electron mobility for oxidized graphene. At 0.2 C, the cathode can maintain a capacity of 1116 mAh/g after 100 cycles. When the current density increases to 4.0 C, it can still maintain a capacity of over 60%. Cai et al. [338] designed a TiO₂-VO_x heterostructure (CTVHs) cathode (Fig. 12b). Thanks to the heterojunction structure, the high adsorption energy heterojunction interface serves as a capture center to capture lithium polysulfides, ensuring rapid diffusion of lithium polysulfides to VO_x. In addition, VO_x rich in defects has high catalytic activity and rapid lithium-ion migration rate for polysulfides, effectively achieving adsorption, diffusion, and conversion. It has an initial specific capacity of 1427.6 mAh/g at 0.5 C and can be cycled over 1400 cycles, with a capacity decay rate of only 0.029% per cycle.

  Figure 12

  

  Figure 12.  (a) The synthesis flowchart, SEM diagram and cycle performance of S/TiO₂@rGO. Reproduced with permission [338]. Copyright 2023, Elsevier. (b) The synthesis flowchart, SEM diagram and cycle performance of TiO₂-VO_x cathode. Reproduced with permission [337]. Copyright 2016, American chemical Society.

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  Polar metal sulfides are often combined with carbon materials as the main material for the cathodes of lithium sulfur batteries [339]. Through the polar metal sulfur (M-S) bond, polysulfides are maximally adsorbed, improving the overall support and conductivity of the electrode, thereby alleviating the problem of rapid capacity degradation in lithium sulfur batteries during long cycles to the greatest extent possible. Li et al. [340] decorated Fe₇S₈-MoS₂ heterostructures on layered MoS₂ embedded with N/P-doped carbon nanocapsules as efficient sulfur hosts for Li-S batteries (S)/Fe₇S₈-MoS₂@MoS₂-NPC. The Fe₇S₈-MoS₂ heterostructure combines the advantages of two coupling components, greatly improving its adsorption capacity for lithium polysulfides and accelerating its conversion kinetics. In addition, doped N/P can provide sufficient space and abundant adsorption sites, effectively suppressing their dissolution and shuttle effects. The electrode can achieve an initial capacity of 913.0 mAh/g at 1.0 C and still maintain 751.6 mAh/g after 400 cycles.

  3.1.3   Lithium metal anodes

  Lithium metal anode is considered pivotal for realizing the Li-S battery due to the high theoretical capacity (3860 mAh/g) and low redox potential (−3.040 V vs. SHE) [341,342]. However, its practical application is severely hindered by the Li dendrite, which can lead to short-circuit, explosion and excessive electrolyte decomposition. Generally, there are three main aspects that need to be addressed:

  (1) Suppressing dendrite growth

  Based on the thermodynamic model of Laplace equation [343], the generation of Li dendrites results from uncontrolled stress release, so the stress generated by dendrites can be compressed by a coating film with high Young’s modulus. Natural or artificial SEI can form a passivated and high ion-conductive coating layer on Li metal, which can not only isolate Li metal from the electrolyte, but also provide a certain modulus to inhibit dendrite [344]. Nowadays, the routes of forming coating layer on the Li liquid metal has become an effective issue for harvesting controllable and stable artificial SEI [345,346].

  (2) Inducing dendrite-free deposition

  Based on Sand’s equation, the generation of Li dendrites is attributed to the Li ion deficient layer formed on the electrode surface under the influence of current. The introduction of framework materials with high specific surface areas can reduce the local current density of Li metal deposition and achieve the purpose of inhibiting dendrites. Additionally, based on nucleation thermodynamics (surface energy model and lithiophilic model) [343], researchers paid their attention to the modification of the current collector, and a lot of work effectively proved that the increase of the lateral diffusion rate of Li ions is conducive to make Li deposition inclined to the dendrite-free [347,348]. As a matter of fact, these routes are commonly originated from more homogeneous distribution of ion, electric filed, surface heat and so on [349–351].

  (3) Healing dendrite

  There are two routes for the method. One way involves healing dendrite by Li ion deposition without further dendrite deterioration. For instance, the addition of low-concentration cesium ions inhibits the formation of Li dendrites through the redistribution of potential on the basis of the Nernst equation [352,353]. Another example is iodine redox reactions, which have been shown to repair dendrites [354]. Another way lies in the slight dendrite prior dissolution following the Li⁰_(dendrite) − e⁻ = Li⁺, then a charging protocol of asymmetrical bidirectional current can harvest a healable Li metal anode in a synchronous feedback cycle [355].

  3.1.4   Solid-solution metal anodes

  Lithium metal, operating on a dissolution-deposition mechanism, is the highest specific capacity anode material [356]. Its electrochemical reactions rely on a solid-liquid phase transition, involving the interconversion between lithium metal and lithium ions. Unlike conventional lithium-storage materials that maintain a stable electrode structure during cycling, electrodes utilizing the dissolution-deposition mechanism experience significant volume changes, leading to reaction non-uniformities. At the macroscopic level, this manifests as the redistribution of active material, such as lithium dendrite formation and the accumulation of electronically isolated "dead lithium". These phenomena not only compromise structural integrity and result in active material loss but can also cause internal short circuits, potentially triggering thermal runaway, fires, or even explosions. Therefore, enhancing the uniformity of lithium metal anode reactions is critical for improving overall battery performance.

  At the microscopic level, the uniformity of electrode reactions is governed by two key factors: the electrochemical reaction rate and the diffusion rate. For lithium metal anodes, the kinetics of lithium-ion reduction at the electrode surface are rapid, characterized by a high exchange current density and a reaction rate constant exceeding 0.1 cm/s, leading to an accumulation of lithium atoms on the electrode surface. However, due to lithium metal’s high surface energy, the diffusion energy barrier for lithium atoms on its surface is significant, resulting in a slow self-diffusion rate. Consequently, lithium atoms accumulate on the surface without timely inward diffusion, which promotes lithium dendrite formation. Therefore, suppressing lithium dendrite growth hinges on balancing the reduction rate of lithium ions with the diffusion rate of lithium atoms within the electrode.

  The exchange current density represents the bidirectional (net-zero) current that maintains the equilibrium between two species during a reaction. It serves as a measure of the balanced rates of oxidation and reduction at the electrode under equilibrium conditions [357]. As shown below:

  i0=FKaOαaRβ

  (1)

  where F represents the Faraday constant, K is the rate constant of the electrode reaction, and aOα and aRβ are the symmetry coefficients for the anodic and cathodic reactions, respectively, with α+β=1. aOα and aRβ denote the activities of the oxidized and reduced species, respectively. From Eq. 1, it can be observed that the exchange current density is directly influenced by the activities of the reactants and products. Therefore, from the perspective of electrode design, reducing the activity of solid metallic lithium is an effective strategy to decrease the exchange current density. For pure solid metals, the activity is typically 1, whereas for multi-component solids, the activity can be expressed as follows:

  ai=Xiγi

  (2)

  where Xi represents the molar fraction of the component, and γi denotes the activity coefficient. From Eq. 2, it is evident that the activity of a specific component in a solid phase is related to its molar fraction. Consequently, the exchange current density of lithium metal electrodes can be modulated effectively through alloying strategies.

  On the other hand, the chemical potential gradient and diffusion coefficient are the two key factors determining the atomic diffusion flux in the solid phase [16], as expressed by the following equation [358]:

  Ji=−D∂μi∂x

  (3)

  where J_i represents the diffusion flux of species i, x denotes the distance from the bulk phase to the solid-phase surface, ∂μi∂x is the chemical potential gradient, and D is the diffusion coefficient. Additionally, the diffusion coefficient D is inversely proportional to the atomic migration energy barrier, as expressed by the following equation:

  D∝exp(−Ea/kT)

  (4)

  where E_a is the atomic migration energy barrier, k is the Boltzmann constant, and T is the temperature. Thermodynamic analysis indicates that the chemical potential gradient ∂μi∂x determines the direction of atomic diffusion, while the diffusion coefficient D governs the diffusion rate. Reducing the accumulation of lithium atoms on the electrode surface requires downhill diffusion of lithium atoms and a rapid lithium atomic diffusion rate. This facilitates the inward diffusion of lithium atoms formed at the alloy/electrolyte interface into the bulk phase. Therefore, enhancing the diffusion flux of lithium atoms within the solid phase hinges on creating a chemical potential gradient in the bulk and reducing the atomic migration energy barrier.

  Based on this, we propose optimizing electrode performance through the design of solid-solution alloy phases [359]. Leveraging the solid-solution reaction between alloys enables lithium atoms to grow tens of micrometers deep into the bulk phase, avoiding surface deposition of lithium metal. During lithiation, lithium atoms generated at the alloy/electrolyte interface diffuse into the bulk phase to form alloys. During delithiation, lithium atoms can be extracted from the bulk phase through dealloying. This unique inward growth deposition method prevents dendrite formation and achieves stable cycling with high Coulombic efficiency. The alloy electrode demonstrates an average Coulombic efficiency of 99.8% ± 0.1% over 400 cycles and delivers a high specific capacity of 1660 mAh/g. This approach enables stable cycling of full cells under lean electrolyte conditions and low cathode-to-anode capacity ratios.

  For alloying reactions, the electrochemical alloying of lithium can be categorized into two types: (1) Reconstitution reactions and (2) solid-solution reactions [360]. The alloy phase plays a crucial role in guiding the lithiation/delithiation processes. Understanding the thermodynamic and kinetic properties of alloy phases during these cycles is essential. This includes elucidating (1) the reversibility of alloy phase evolution under thermodynamic equilibrium and (2) the diffusion kinetics of metal atoms in alloy phases and their impact on the uniformity of the alloying/dealloying reactions. We investigated three representative lithium-based alloy systems based on binary phase diagrams: The Li–Ag alloy with a lithium-rich solid-solution phase extending to metallic lithium, the Li–Zn alloy with a narrow solid-solution phase (49–50 at% Li), and the Li–Al alloy containing only intermetallic compound phases [361]. Using thermodynamic data, we calculated the concentration and chemical potential of Li, Ag, Zn, and Al through Gibbs free energy profiles and the Gibbs–Duhem equation. Our findings reveal that in the Li–Ag solid-solution alloy, the presence of a certain degree of lithium solubility results in a continuous variation in chemical potential within the bulk phase. This continuous chemical potential gradient within each solid-solution phase provides a driving force for diffusion. In contrast, Li–Al and Li–Zn intermetallic compounds, which have fixed stoichiometric ratios, exhibit equal chemical potential within a single phase, lacking a gradient for atomic diffusion. The only possible diffusion occurs near phase boundaries due to the chemical potential differences between phases, which results in relatively low diffusion efficiency. Additionally, the atomic interactions in intermetallic compounds are typically stronger than those in solid solutions, as they involve a greater covalent bonding component. This is reflected in the lower mixing enthalpy of intermetallic compounds compared to solid-solution alloys, indicating stronger atomic interactions. Consequently, the higher atomic migration energy barriers in intermetallic compounds hinder diffusion. While past studies have often focused on the tight coupling between Li-alloying reactions and alloy phase transitions, our findings offer valuable insights into the intelligent design of components for advanced secondary metal batteries.

  3.1.5   Electrolytes

  In lithium–sulfur (Li–S) batteries, the electrolyte plays a critical role in achieving high performances, acting not only as an ion-conducting medium but also directly influencing the cathode conversion reaction kinetics and the stability of the anode interface [362]. During cycling, lithium polysulfides (LiPSs) are generated as intermediates, which dissolve in electrolyte to facilitate the charge–discharge process at the cathode [363,364]. However, LiPSs also diffuse to the anode, where they react with metallic lithium and result in the notorious shuttle effect [365,366]. Thus, electrolyte governs the solvation structure of LiPSs, influencing their dissolution, diffusion, cathodic kinetics, as well as the stability of the SEI on Li anode [367,368]. An optimized electrolyte is expected to balance the kinetics of polysulfide conversion at the cathode with anode stability, improving the discharge capacity and prolonging the cycling lifespan of Li–S batteries.

  Tetrahydrofuran (THF) was first reported as a suitable solvent for Li–S batteries in 1988 [369], followed by the study of a 1,3-dioxolane (DOL)-based electrolyte in 1989 [370]. In 1993 and 1994, dimethoxyethane (DME) and DOL were explored as co-solvents for ester- and sulfone-based electrolytes [371]. By 2003, DOL/DME mixtures became the primary solvent system for Li–S battery electrolyte [372]. Fluorinated salts, such as lithium bis(trifluoromethanesulfonyl)imide (LiTFSI), were chosen as the lithium salt due to their inertness toward polysulfides [373]. In 2008, the introduction of LiNO₃ effectively mitigated the corrosion of the lithium anode by LiPSs [374], establishing the benchmark electrolyte for Li–S batteries, composed of 1.0 mol/L lithium salt and 2% LiNO₃ in DOL/DME.

  In the benchmark electrolyte, the presence of free solvents allows the polysulfides to be fully dissolved and presents solvent-separated ion pairs (SSIPs), facilitating the cathodic solid–liquid–solid electrochemical conversions (Fig. 13). This electrolyte meets the initial requirements for both cathode conversion kinetics and anode interface stability. However, under practical lean electrolyte conditions, the significant increase in LiPS concentration severely deteriorates the standard electrolyte [375]. Concentrated LiPSs increase the viscosity and reduce the ionic conductivity. Besides, LiPS oversaturation slows the cathode reactions by solid–solid conversion and exacerbates the side reactions at the anode, limiting the practical application of the benchmark electrolyte [376]. Adjusting the lithium salt concentration or regulating solvents has been shown to mitigate these issues [377]. Enhancing the polysulfide solubility can maintain a liquid phase conversion pathway under lean electrolyte conditions, improving the cathode kinetics and sulfur utilization [378,379]. Such a highly solvating electrolyte (HSE) can be achieved by replacing DOL/DME with solvents strongly solvating LiPSs, such as tetramethylurea (TMU) [380] and dimethyl sulfoxide (DMSO) [381]. However, the high reactivity of these strongly solvating solvents accelerates Li anode corrosion, necessitating further innovations to improve the cycle lifespan in practical high-energy-density Li–S batteries [382].

  Figure 13

  

  Figure 13.  Schematic diagram of the solvation structures of lithium ions and lithium polysulfides in different electrolyte systems for Li–S batteries.

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  An alternative strategy is to reduce LiPS dissolution in electrolyte, which can decouple the cathode conversion process from electrolyte dependence and mitigate anode corrosion [383]. Increasing the lithium salt concentration is a direct approach to control LiPS dissolution. When the salt concentration in electrolyte exceeds a critical threshold, free solvents and SSIPs decrease while contact ion pairs (CIPs) and aggregates (AGGs) increase, resulting in a high-concentration electrolyte (HCE) [384,385]. Research on HCE in Li–S batteries began in 2013, with Oh and coworkers demonstrating that adding 5 mol/L LiTFSI in tetraethylene glycol dimethyl ether (G4) achieved over 99% Coulombic efficiency (CE) [386]. Further, ether-based electrolytes containing 7–12 mol/L fluorinated salts improved the CE by minimizing LiPS dissolution [387,388]. Increased lithium salt concentration also induces an anion-derived SEI on the lithium anode, enhancing the anode stability. However, HCEs present challenges, including poor wettability, high viscosity, and reduced ionic conductivity, which hinder the cathodic kinetics and increase costs.

  Introducing co-solvents with weak solvating power into HCE to form localized high-concentration electrolytes (LHCE) can mitigate the issues associated with high viscosity while preserving the microstructure of HCE [389]. In 2015, Watanabe and coworkers introduced 1,1,2,2-tetrafluoroethyl-2,2,3,3-tetrafluoropropyl ether (TTE) into G4-based ionic liquid electrolytes and examined its effect on Li–S battery performances [390]. This work boosted the development of various LHCEs using new diluents, including 1H, 1H, 5H-octafluoropentyl-1,1,2,2-tetrafluoroethyl ether (OFE) [391], 2,2-dimethyl-3,6,9-trioxaza-2-dodecane [392], hexafluoroisopropyl methyl ether (HFME) [393], and 1,1,2,2-tetrafluoroethyl methyl ether (TME) [394]. These electrolytes effectively reduce LiPS dissolution and suppress lithium dendrite growth, leading to high CE and stable cycling performances. While the LHCE concept is still emerging, its unique electrolyte structure offers considerable design flexibility, potentially meeting the diverse performance requirements.

  The total salt concentration and the weakly solvating solvent content in LHCE can be further reduced to constitute a more practical electrolyte. Zhang’s group developed a series of encapsulating LiPS electrolytes (EPSE) with 1.0 mol/L salt concentration, utilizing weakly solvating solvents besides hydrofluoroether (HFE) such as isopropyl ether (DIPE) [395], diisopropyl sulfide (DIPS) [396], and hexyl methyl ether (HME) [397,398]. When combined with moderately solvating solvents, these electrolytes form a dual-layer solvation structure for LiPSs. The inner layer ensures LiPS dissolution and supports the cathode conversion kinetics, while the weakly solvating outer solvents prevent the parasitic reactions between the lithium anode and LiPSs. The optimal ratio of weakly to relatively strongly solvating solvents is crucial to balance the viscosity, ionic transport, and reaction kinetics. Zhang’s group optimized the outer solvent content, demonstrating that moderately overall solvating EPSEs balance the cathode and anode kinetics, leading to enhanced cycling performances in Li–S batteries [399].

  The LiPSs generated during cycling can dissociate and supplement the lithium salt concentration, enabling the construction of low-concentration electrolytes (LCEs) with reduced lithium salt content. These LCEs lower the viscosity, optimize the interface reactions, and improve the battery performances while significantly reducing the costs. Research on LCEs for Li–S batteries began in 2018. Sun’s group demonstrated that reducing the LiTFSI concentration to 0.5 mol/L decreased electrolyte viscosity from 1.54 mPa s to 0.72 mPa s, improving the wetting properties of cathodes in pouch cells [400]. Wu et al. [401] identified that LCEs at low temperatures can promote Li₂S nucleate and three-dimensional growth, suppress LiPSs aggregate, and induce a dense SEI synergistically with LiNO₃. Yan et al. [402] further found that the desolvation energy barriers for lithium ions were positively correlated with the lithium salt concentrations below 1.0 mol/L. Additionally, Yushin et al. [403] explored the addition of HFE to 0.10 mol/L lithium salt-based electrolyte, achieving stable anode interface formation and improved cycling stability.

  Compared to adjusting the solvent and salt concentrations, introducing additives is a cost-effective and convenient way to enhance the battery performances [404]. Small quantities of additives (typically <10 wt%) can significantly alter the electrolyte properties [405]. Additives, such as metal nitrates [406,407], nitrogen compounds (both inorganic [408,409] and organic [410]), sulfur compounds (e.g., lithium polysulfides [411,412], organic sulfides like linoleic acid [413,414]), and halide-based substances [415,416], can form functional SEI on Li surface, thereby mitigating the side reactions between LiPSs and Li and preventing un-even Li deposition. Compared to the anode, functional additives targeting the sulfur cathode have been more extensively studied. Additives affect both the thermodynamics and kinetics of the cathode through the interactions with the sulfur-containing species [417]. Huang’s research team has pioneered a series of redox mediators [418–420] and redox comediators [421–423] to enhance sulfur cathode utilization and discharge capacity. Based on these additives, they achieved a high actual energy density of 550 Wh/kg in 5 Ah Li–S pouch cells [424,425].

  In conclusion, electrolytes are critical to optimizing Li–S battery performances, with significant advancements made in the iteration of the ether-based benchmark electrolyte. Adjusting salt concentration has led to the development of HCE and LHCE system, while solvent property regulation has resulted in HSE and EPSE systems. Significant progress has also been made with additives targeting both the anode interface and the cathode conversion kinetics. The behavior of high-concentration LiPSs, including their conversion mechanisms and impact on electrolyte microstructure, needs further clarification. Additionally, the synergistic effects of additives in various solvent systems must be explored. During electrolyte development, it is essential to establish quantitative models linking the solvent structure, properties, and battery performances. Solid-state electrolyte exhibits significant potential to improve Li–S battery safety and cycling stability, where future research should focus on optimizing the interface between the solid-state electrolyte and electrodes, developing scalable fabrication methods, and enhancing the ionic conductivity while maintaining mechanical stability.

  3.2 Lithium-air batteries

  3.2.1   Lithium-O₂ batteries

  3.2.1.1   Cathode

  (1) Carbon materials

  Typical Li-O₂ batteries consist of a lithium metal anode paired with an oxygen-breathing cathode, connected via a suitable electrolyte. These batteries lie conceptually between fuel cells and traditional batteries. Their energy output is based on the reversible electrochemical reactions of lithium metal oxidation and oxygen reduction (Li₂ + O₂ ↔ Li₂O₂). Due to lithium metal’s extremely high specific capacity (3860 mAh/g) and lowest theoretical potential (−3.04 V), as well as the utilization of cathode active materials sourced from ambient air, Li-O₂ batteries exhibit an impressive theoretical energy density of up to 3500 Wh/kg [426,427]. This makes them a subject of significant global research interest and positions them as one of the most promising candidates for powering smart electronic devices and electric vehicles.

  However, the development of Li-O₂ batteries remains in its early stages, and several challenges must be addressed. One critical issue is the uncontrollable growth of lithium dendrites, which raises safety concerns for practical applications [428,429]. Similarly, side reactions in liquid electrolytes under high voltage, along with lithium metal corrosion caused by air crossover, negatively impact the electrochemical performance of these batteries [430]. Replacing the liquid electrolyte with a solid-state electrolyte is a promising strategy to mitigate some of these challenges [431,432]. Another significant issue is that the discharge product (Li₂O₂) is both insulating and insoluble, leading to its continuous accumulation in the porous air electrodes during discharge. This accumulation directly affects the battery’s performance, as it depends on the electrode’s ability to store and decompose the discharge product efficiently [433]. Developing high-efficiency cathodes that promote controlled deposition and accelerate the decomposition of LiO₂ is, therefore, a key approach to enhancing the electrochemical performance of Li-O₂ batteries.

  Various materials have been explored as cathodes for Li-O₂ batteries, including noble metals, transition metal oxides, and carbon materials [434–436]. Among these, carbon materials stand out due to their high conductivity, low density, large specific surface area, and porous structure, making them excellent candidates as cathode materials or framework materials for Li-O₂ batteries [437].

  Graphene, known for its ultrathin 2D structure and exceptional physicochemical properties, has been extensively studied [438]. Sun et al. [439] firstly investigated graphene as a cathode material for Li-O₂ batteries, demonstrating its remarkably high discharge capacity. Since then, various graphene-based materials, such as structured graphene, doped graphene, and hybrids with noble metals or transition metal oxides, have been developed. Studies have shown that doping graphene with elements like nitrogen (N) and sulfur (S) enhances its catalytic activity and accelerates the decomposition of discharge products [438,440]. The porous structure of graphene, particularly hierarchical and interconnected pores, improves oxygen diffusion and facilitates the deposition of Li₂O₂, thereby enhancing the battery’s capacity and rate performance (Fig. 14a) [441]. Additionally, graphene serves as an excellent framework material for supporting and dispersing catalysts.

  Figure 14

  

  Figure 14.  (a) SEM images of graphene air electrode with hierarchical and interconnected structure. Reproduced with permission [441]. Copyright 2015, Wiley-VCH. (b) Schematic illustration for the fabrication of free-stand CNTs air electrode. Reproduced with permission [442]. Copyright 2017, Elsevier. (c) Schematic illustration of the cable-type Li-O₂ batteries. Reproduced with permission [446]. Copyright 2018, Wiley-VCH. (d) In situ observation of the morphological evolution of the discharge product with CNTs-based electrodes. Reproduced with permission [449]. Copyright 2021, Wiley-VCH.

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  Carbon nanotubes (CNTs), with their unique 1D structure, have also been utilized to construct electrodes with porous architectures. These structures provide efficient oxygen diffusion pathways and ample space for discharge products (Fig. 14b) [442]. Moreover, self-supporting films and flexible electrodes have been developed using CNTs, enabling the design of Li-O₂ batteries without traditional current collectors and even flexible or cable-type configurations (Fig. 14c) [443–446]. The 1D structure of CNT-based electrodes is particularly valuable for investigating electrochemical mechanisms (Fig. 14d) [447–449].

  Other carbon materials, such as commercial conductive carbon and porous carbon, have been employed as standard electrodes for evaluating catalysts and developing large-scale cells [450]. Functional porous carbons, including mesoporous carbon, macroporous carbon, MOF-derived carbon, biomass-derived carbon, porous carbon nanowires, and porous nanosheets, exhibit unique functional properties that further advance the performance of Li-O₂ batteries [451–453].

  (2) Solid state catalysts

  Among various candidates for air cathodes, carbon-free catalysts have attracted significant attention due to their potential for designing tunable active sites (Fig. 15a) [454,455]. Solid state catalysts, including metals (Au, Ru, Pt, etc.) and metal oxides (Co₃O₄, MnO₂, Fe₃O₄ etc.), have been extensively investigated as carbon-free cathodes for Li-O₂ batteries, demonstrating promising catalytic activity. These findings significantly unveil the feasibility and effectiveness of the carbon-free approach in improving the electrochemical performance of Li−O₂ batteries.

  Figure 15

  

  Figure 15.  (a) Schematic illustration of the essential difference in electrochemical reactions on carbon-based and carbon free cathode surface. (b) Schematic illustration of the change in local electronic structure at Li_xO₂ molecular orbitals upon reactions on the surface of metal alloy cathode. (c) First charge−discharge curves of six types of metal-based cathodes at a current density of 1.0 A/g and a specific capacity limit of 3000 mAh/g. (d) Predicted catalytic activities of typical catalyst including metal oxide based on the established correlation of O₂ desorption and charging voltage with surface acidity. (a, b) Reproduced with permission [455]. Copyright 2023, Wiley-VCH. (c) Reproduced with permission [458]. Copyright 2017, American Chemical Society. (d) Reproduced with permission [460]. Copyright 2015, American Chemical Society.

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  Metal catalysts exhibit considerable catalytic activity in Li-O₂ batteries, benefiting from their tunable electronic structures [456,457]. For instance, Ru-based catalysts as air cathode show high selectivity for boosting the generation of Li_xO₂, a key intermediate for the reversible operation of Li-O₂ battery. Additionally, the incorporation of metal atoms into bimetallic systems, such as Au-Ni or Ir-Ni alloys, has been shown to synergistically enhance ORR and OER performance by optimizing the adsorption energies of reaction intermediates (Fig. 15b), thereby facilitating battery kinetics, and reducing side reactions [458,459]. These advances highlight the potential of metal catalysts working as air cathode in achieving practical Li-O₂ batteries with improved cyclability and energy efficiency (Fig. 15c). However, despite their promising catalytic activity, metal catalysts face challenges such as high priced, unsatisfied stability, and susceptibility to degradation during repeated charge-discharge cycles, which impede their practical application.

  Price is another critical factor influencing the choice of cathode catalysts. Metal oxides have gained significant attention due to their cost-effectiveness, excellent catalytic properties, and relatively high stability (Fig. 15d) [460,461]. Among various candidates, Co_xO_y, Mn_xO_y, and Fe_xO_y stand out as leading oxide catalysts for facilitating the ORR and OER processes in Li-O₂ batteries. Such as Co₃O₄, with its spinel structure, provides a favorable electronic configuration that enhances oxygen adsorption and electron transfer during ORR and OER. Numerous studies have demonstrated that metal oxide catalyst offers high discharge capacities and long cycle life, making them a promising choice for Li-O₂ batteries. However, the relatively low conductivity of metal oxide catalyst limits its electrochemical performance, necessitating the development of hybrid catalysts or conductive supports to enhance its catalytic activity and electron transfer efficiency. Surface modifications, including doping with other metals or introducing defects, have been explored to improve its electrochemical performance.

  (3) Liquid phase catalysts

  In the past decade, liquid-phase catalysts have been used as electron mediators to transfer electron between the electrodes and the active materials due to their dissolution in the electrolyte. This allows the generation and oxidation of Li₂O₂ to take place in solution rather than on the surface of the electrodes, which effectively reduces the polarization voltage and improves the reaction rate, while alleviating the capacity limitations caused by passivation on the surface of the electrodes and pore plugging, and is expected to improve the multiplier performance of the battery. The results are expected to improve the performance of battery multiplication.

  In 2013, combining nanoporous gold electrodes and dimethyl sulfoxide, Bruce et al. [462] successfully stabilized the charging voltage of a lithium-oxygen battery below 3.5 V by adding tetrathiafulvalene (TTF), which demonstrated good reversibility over 100 cycles. The following year, Janek et al. [463] reported another important class of liquid-phase catalysts for the oxygen precipitation reaction, 2,2,6,6-tetramethylpiperidine oxide (TEMPO). The introduction of TEMPO is also effective in reducing the charge polarization, resulting in a reduction of 500 mV in the charging potential of the battery. Since then, a variety of liquid-phase catalysts for oxygen precipitation reactions have been reported and their catalytic mechanisms have been explored, such as lithium halides [464,465], 10-methylphenothiazine (MPT) [466], tris-s-aniline [467,468], dimethylphenazine (DMPZ) [469,470], metal-organic compounds [471,472]. The catalysts have also been reported and their catalytic mechanisms have been explored in depth. Using the same electron transfer function, homogeneous catalysts also play a significant role in catalyzing the oxygen reduction reaction. Owen et al. [473,474] found that ethyl viologen (EV) can shorten the lifetime of lithium superoxide intermediates while increasing the oxygen reduction rate, which is expected to inhibit the side effects of superoxide radicals, and offers the possibility of improving the cycling stability of secondary lithium-oxygen batteries. In 2016, Bruce et al. [475] first proposed that 2,5-di-tert-butyl-1,4-benzoquinone (DBBQ), can form LiDBBQO₂ while transferring electrons, which stabilizes superoxide intermediates and facilitates the generation of liquid-phase Li₂O₂, thus increasing the capacity of the battery up to 80–100 times. Based on this principle, using density-functional theory calculations [476,477], in situ infrared spectroscopy [478], nuclear magnetic resonance (NMR) hydrogen/lithium spectroscopy [479,480], UV–visible spectroscopy [480,481], and electron paramagnetic coexisting vibration [477] methods, a series of quinones have been shown to form intermediates of RM-LiO₂ and thus stabilize the superoxide intermediates, thereby significantly increasing battery capacity while reducing polarization [482–484].

  3.2.1.2   Anodes

  The advantage of ultra-high energy density in Li-O₂ batteries is due to the intrinsically high energy density (~3680 mAh/g) and low potential (−3.04 V vs. SHE) of Li-Metals. The anode study of Li-O₂ batteries is similar to that of Li-ion batteries, so Li-O₂ batteries also face problems such as lithium dendrites, dead lithium, and side reactions at the solid-liquid interface. In addition, the "shuttle effect" of active oxygen and other impurities (CO₃, N₂, H₂O, liquid phase catalysts, etc.) also causes serious side reactions, leading to the reduction of active lithium and impairing the cycle life of the battery. In addition to the preferred lithium metal, other anode materials include Si, Sn, Al, C. However, these materials do not have sufficient lithium embedded capacity, and are prone to deactivation when side reactions occur, which makes their reversibility unsatisfactory. To date, researchers have attempted to address the challenge of using lithium as an anode in the following four ways: (1) Artificially constructing a stable SEI layer to isolate the side reactions [485,486]. (2) Reducing electrolyte activity (ultra concentrated electrolytes, ionic liquids, solid electrolytes, etc.), reducing reactants for side reactions [487–489]. (3) Preparation of selectively permeable diaphragms to eliminate the "shuttle effect" [490,491]. (4) Modulation of lithium metal/collector structure to avoid shape and volume changes [492,493]. However, as one of the core components of lithium oxygen batteries, the research on lithium metal is very challenging. How to improve its stability without compromising other properties such as rate performance and energy density is the aim of future research.

  3.2.1.3   Electrolytes

  A crucial component of Li-O₂ battery is the electrolyte, known as the blood in batteries, which serves as the carrier of Li⁺ transport between the anode and cathode during discharge/charge cycles [494]. Generally, there are four types of Li-O₂ battery, categorized based on the electrolyte and named aprotic, aqueous, hybrid and solid-state Li-O₂ batteries [495]. Compared to the other three types, the aprotic (non-aqueous) Li-O₂ battery has more advantages in energy density, operability and system reversibility. And it has also attracted extensive attention from researchers around the world [496].

  Early research on aprotic Li-O₂ batteries utilized organic carbonates (esters), such as EC, PC and DMC solvent as the component of electrolyte, following their dominance in Li-ion batteries [497]. However, these carbonate components were later confirmed to be unstable in the presence of oxygen-reducing species. A series of spectroscopic characterizations showed that the discharge products of oxygen reduction reaction (ORR) process of aprotic Li-O₂ batteries were mainly carbonates, such as lithium carbonate (Li₂CO₃) and alkyl lithium carbonate (RO-(C_xO-OLi)) rather than the ideal discharge product of Li₂O₂ when carbonate electrolytes were used. Simultaneously, the gaseous CO₂ rather than O₂, is released during charging as the main oxidation product [498,499]. Due to relatively higher oxidation stability than ester solvents, scholars have focused much attention on ether solvents in organic electrolyte of aprotic Li-O₂ battery. Read et al. [500] first investigated ether-based electrolyte and reported good stability as well as excellent rate capability in Li-O₂ battery with dimethoxyethane (DME) and 1,3-dioxolane (DOL) solvents. A few years later, McClosky’s group [501] confirmed that Li₂O₂ is the dominant discharge product in DME-based electrolyte of aprotic Li-O₂ battery during ORR process. Recently, though longer chain ether solvents, especially TEGDME, have been considered as another choice for the electrolytes in aprotic Li-O₂ batteries [502,503], Sharon et al. [504] and Freunberger et al. [505] further confirmed the degradation of ether-based solvents under nucleophilic attack in aprotic Li-O₂ battery (Fig. 16), which was detrimental to the long-term stable cycling of aprotic Li-O₂ system. Then other types of electrolytes have also been investigated in aprotic Li-O₂ batteries, such as acetonitrile (ACN) [506], dimethyl sulfoxide (DMSO) [507,508], dimethylformamide (DMF) [509], ionic liquids [510,511]. However, their compatibility with Li-metal anode in aprotic Li-O₂ batteries needs to be carefully considered. Due to the flammability and volatility defects of traditional organic electrolytes, solid-state electrolytes have also been reported and have shown some promise in aprotic Li-O₂ batteries [512,513]. The first reported all solid-state Li-air battery by Kumar et al. [514] consisted of solid-state electrolyte composite film stacked by the polymer intermediate layer and the glass ceramic Li⁺ conductor, which can achieve real operation under an open-air atmosphere. Yu’s group [515] presented an ultrathin, high-ion-conductive lithium-ion-exchanged zeolite X (LiX) membrane as the sole solid electrolyte to effectively suppress the electrolyte degradation and promote long cycle life of Li-O₂ batteries. Metal-organic frameworks derived solid-state electrolytes along with high chemical/electrochemical robustness can also exhibit superior rate capability and stable cycle life by building multiple wetting interfaces for integrated solid-state Li-O₂ battery [516].

  Figure 16

  

  Figure 16.  Proposed mechanism for the decomposition of ether-based solvents during ORR process in non-aqueous Li-O₂ batteries. Reproduced with permission [505]. Copyright 2011, Wiley-VCH.

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  Both organic electrolytes and solid-state electrolytes are considered to be promising in Li-O₂ batteries. However, many compatible and practical issues still need to be solved and addressed before large-scale applications, such as chemical/electrochemical stability, volatility in open air atmosphere, interfacial contact, triple-phase boundaries, which is closely related to whether the long-term and stable cycling performance of practical Li-O₂ battery can be achieved.

  3.2.2   Lithium-CO₂ batteries

  In the background of carbon neutrality, Li-CO₂ batteries have received much attention for their ability to fix CO₂ and their high theoretical energy density of 1876 Wh/kg [517,518]. Takechi et al. [519] reported the Li-O₂/CO₂ batteries and found that the introduction of CO₂ dramatically increased the discharge capacitance for the first time. Subsequently, researchers have conducted numerous studies on battery systems using pure CO₂ as the working gas [520,521]. The electrochemical redox reaction of CO₂ in a Li-CO₂ cell is 4Li⁺ + 3CO₂ ↔ 2Li₂CO₃ + C [522,523]. Despite their great potential, the continuous accumulation of discharge products at the positive electrode and the slow reaction kinetics during practical use constrain the further development of Li-CO₂ batteries [524]. Therefore, exploring efficient cathode catalysts with good catalytic activity and stability to promote battery reversibility and stability is crucial for the practical application of Li-CO₂ batteries.

  3.3 Sodium-sulfur batteries

  Room temperature sodium-sulfur batteries (RT Na-S) have gained widespread attention due to the high theoretical specific capacity of sulfur (1675 mAh/g), cost-effectiveness, and environmental sustainability [525–527]. In contrast to conventional high-temperature Na-S batteries, RT Na-S systems offer notable advantages regarding operational temperature, material choices, and safety. Nonetheless, challenges such as sluggish electrochemical kinetics, notorious polysulfide shuttle, and limited cycling stability hinder their practical applications.

  Typically, RT Na-S batteries exhibit a configuration (Fig. 17) comprising a Na-metal anode, a Na⁺ conductive electrolyte and a composite sulfur cathode. When elemental sulfur (S₈) is employed as the cathode substance, the initiation of the battery reaction usually takes place upon discharge. Throughout the discharge process, sodium metal undergoes oxidation at the anode, yielding sodium ions and releasing electrons. These sodium ions migrate internally towards the cathode through the electrolyte, while electrons traverse towards the positive electrode via the external electrical circuit. Simultaneously, sulfur undergoes reduction, accepting sodium ions and electrons at the cathode, thereby forming sodium polysulfides.

  Figure 17

  

  Figure 17.  Work principles of RT Na-S batteries.

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  The sulfur conversion behavior in RT Na-S batteries, akin to that of Li-S systems, is critically influenced by the composition and characteristics of electrolytes. Ether-based electrolytes allow for the substantial solubility of long-chain sodium polysulfides, considerably complicating the battery chemistry. A TEGDME-based solution (1.5 mol/L NaClO₄ and 0.3 mol/L NaNO₃) has been adopted as the model electrolyte for the early research of RT Na-S batteries [528], inherited from Li-S analogs [529,530]. Discharge in ether-based Na-S batteries typically involves following four distinct stages with complex solid-liquid-solid transitions. (1) Solid-liquid transition (a plateau at ~2.2 V): Sulfur (S₈) converts to soluble Na₂S₈ through a fast reaction. (2) Liquid-liquid reaction (a slope between 2.2 V and 1.65 V): Soluble Na₂S₂ transitions to Na₂S₄. This stage, although kinetically rapid, is chemically intricate, governed by equilibrium among polysulfide species. (3) Liquid-solid conversion (a plateau at ~1.65 V): Na₂S₄ converts to insoluble Na₂S₃, Na₂S₂, or Na₂S. Variations in capacity and discharge voltage reflect differences in the dominant reactions. (4) Solid-solid conversion (a slope < 1.65 V): Na₂S₃ or Na₂S₂ further converts to Na₂S. This stage is kinetically limited, with high polarization stemming from the poor conductivity of Na₂S₂ and Na₂S.

  Ester-based electrolytes have recently garnered significant attention for RT Na-S batteries. These electrolytes typically comprise EC, propylene carbonate (PC), diethyl carbonate (DEC), and dimethyl carbonate (DMC), often combined with sodium perchlorate (NaClO₄) as the electrolyte salt [531,532]. Additionally, film-forming additives such as fluoroethylene carbonate (FEC) are commonly employed to develop stable cathode electrolyte interphases (CEI) and solid electrolyte interphases (SEI). For example, in a standard ester-based electrolyte formulation (1 mol/L NaClO₄ in EC/PC at a volume ratio of 1:1 with 5% FEC), RT Na-S cells exhibit discharge/charge profiles distinct from those of ether-based systems, characterized by a single slope-like feature [533,534]. This quasi-single-slope behavior is widely attributed to a quasi-solid-solid transformation mechanism, suggesting limited formation of soluble polysulfides. However, it is important to note that trace amounts of dissolvable polysulfides can still be detected under such conditions.

  To summarize, studies on RT Na-S batteries reveal that ether-based solvents offer distinct advantages, including rapid conversion reactions driven by dissolved polysulfides and stable electrode/electrolyte interfaces due to reduced side reactions with sodium metal. However, these benefits are offset by challenges such as polysulfide dissolution and shuttling, as well as safety concerns inherent to ether-based solvents. By contrast, ester-based electrolytes significantly limit the dissolution of sulfur and polysulfides. Yet, the dissolved species, rich in electrons, react with the electrolyte, causing active material losses and capacity decay. Addressing this requires the formation of a durable and stable CEI during the first discharge process to inhibit further dissolution of polysulfides. This goal demands meticulous sulfur cathode and electrolyte design. Considering the chemical and electrochemical stability of ester-based electrolytes, along with their established experiences in lithium-ion battery technology, they represent a promising avenue for commercialization. Future research should prioritize enhancing the performance of ester-based RT Na-S systems, addressing challenges such as sodium anode compatibility and sensitivity to impurities or moisture.

  RT Na-S batteries, with their high energy density and low cost, present a compelling choice for next-generation energy storage solutions. Nonetheless, their electrochemical performance necessitates ongoing enhancement through the synergy of material innovation, battery structural refinement, and system-level engineering. With further technological advancements, RT Na-S batteries are expected to play a crucial role in the future of green energy storage, supporting global energy transformation and sustainable development goals.

  3.4 Sodium-selenium batteries

  Selenium (Se), as an element in the same group as sulfur, has similar molecular structures and physicochemical properties with sulfur, but is more chemically stable [535]. Its theoretical capacity is 675 mAh/g, which is lower than that of sulfur (1675 mAh/g). However, due to selenium’s high density, its theoretical volumetric capacity reaches up to 3253 mAh/cm³. Moreover, selenium has an electronic conductivity as high as 1 × 10⁻³ S/m, which allows selenium as the cathode active material to achieve high utilization rates and fast electrochemical reaction kinetics [536]. In sodium-selenium batteries, selenium undergoes multi-step redox reactions at the cathode, similar to sodium-sulfur batteries. However, compared to sodium-sulfur batteries, sodium-selenium batteries do not exhibit significant shuttle effects [537]. Therefore, sodium-selenium batteries have tremendous potential for development in the field of high-power, high volumetric energy density for small-sized electronic products.

  3.5 Sodium-air batteries

  In 2013, Adelhelm et al. [538] reported for the first time a rechargeable sodium-air battery at room temperature, which uses metal sodium as the negative electrode, carbon fiber gas diffusion layer as the air cathode, 0.5 mol/L NaSO₃CF₃-DEGDME is the electrolyte. During the discharge process, the metal sodium is oxidized to sodium ions, which then migrate to the positive electrode through the separator and electrolyte, where it reacts with air to form sodium superoxide (NaO₂), which is deposited on the surface of the positive electrode. In the subsequent charging process, the reverse reaction of the above reaction is carried out, in which sodium ions are reduced to metal sodium and deposited on the negative electrode, and NaO₂ is oxidized and decomposed to release air [539].

  The above-mentioned sodium-air batteries based on one-electron transfer (O₂/O₂⁻) exhibit extremely low charge overpotential (<200 mV) [540], which is much smaller than that of lithium-air batteries based on two-electron transfer (O₂/O₂²⁻), and the theoretical energy density can reach 1108 Wh/kg, which is much higher than that of existing commercial LIBs. Therefore, it has attracted extensive attention from researchers. However, in subsequent studies, it was found that NaO₂ is not the only discharge product of sodium-air batteries, and the formation of sodium peroxide (Na₂O₂) is often accompanied by the discharge process.

  The theoretical energy density of sodium-air batteries with Na₂O₂ as the discharge product can reach 1605 Wh/kg, but with the obvious increase of charge overpotential, the charge-discharge curve is no different from that of lithium-air batteries, and it is difficult to give full play to the maximum advantages of sodium-air batteries [541]. Thermodynamically, the free enthalpy of Na₂O₂ formation (−449.7 kJ/mol) is slightly greater than that of NaO₂ (−437.5 kJ/mol) at 298 K and standard atmospheric pressure, and the slight difference in free enthalpy is only −12.2 kJ/mol. The similar thermodynamic stability of the two cells leads to the formation of mixed products of the two during discharge [542]. Theoretical calculations also predict the discharge products of sodium-air batteries. First-principles DFT calculations show that NaO₂ is a stable discharge product of Na-O₂ cells under standard test conditions [543]. Subsequently, Ceder et al. [544] of the Massachusetts Institute of Technology calculated the surface energy of sodium oxide crystals and found that the grain size of sodium oxide determines its thermodynamic stability. Unlike the above calculations, they found that Na₂O₂ is stable in the standard state, but nanoscale NaO₂ with lower surface energy is more stable than Na₂O₂ of the same size. In addition to this, factors such as the cathode, electrolyte, will also affect the composition of the discharge product.

  However, until now, it is still unclear which factors play a decisive role in whether the discharge process follows the single-electron reaction (SOX) or the peroxide-forming two-electron reaction pathway. Therefore, priority should be given to solving this problem in order to develop high-performance sodium-air batteries. In the past decade, sodium-air batteries have developed rapidly in terms of charge and discharge mechanism, anode protection, and electrolyte selection. However, its research is still in its infancy, and rechargeable sodium-air batteries are still far from being used in any practical equipment. Before this new type of battery can be utilized, greater progress must be made in the cell reaction mechanism, battery material design, and battery components. Therefore, there is an urgent need for innovation in basic research, material design, and technology development.

  3.6 Potassium-sulfur (selenium) batteries

  Based on the innovative electrochemical conversion reaction, the development of novel electrochemical energy storage devices utilizing sulfur-based elements (sulfur and selenium) as cathode materials for metal-sulfur (selenium) batteries holds significant promise. This approach aims to transcend the traditional capacity limitations of LIBs, which are constrained by the insertion reaction of their cathode materials. The pursuit of high-energy and high-degree energy storage devices is of paramount importance, with lithium-sulfur batteries garnering considerable attention. Given the limited reserves and high costs associated with lithium, potassium-sulfur and potassium-selenium batteries emerge as viable alternatives, abundant in lithium-like elements.

  The electrochemical reaction principle of potassium-sulfur batteries shares similarities with that of lithium-sulfur batteries. However, the larger radius of potassium ions may contribute to a lower overall discharge potential, significant voltage polarization, and an indistinct voltage platform. Unlike lithium-sulfur batteries, the intermediate discharge product of potassium-sulfur batteries, K₂Sn, remains unclear, and the final discharge product at room temperature, whether K₂S₃ or K₂S, remains controversial. The conversion of K₂S₃ to K₂S is challenging, necessitating a higher overpotential for reduction and leading to incomplete sulfur reduction, which in turn decreases the specific capacity of potassium-sulfur batteries.

  As an alkali metal element following lithium and sodium, potassium, serving as a cathode material, exhibits a low reduction potential (−2.93 V vs. the standard hydrogen electrode, second only to lithium) [545,546] and a high specific capacity (particularly in terms of volumetric specific capacity). Potassium-sulfur batteries boast a high theoretical energy density (914 Wh/kg), far exceeding the upper energy density limit of traditional LIBs (approximately 350 Wh/kg) [547]. Compared to lithium and sodium ions, potassium ions exhibit a smaller Stokes radius in organic electrolytes (K⁺: 3.6 Å; Li⁺: 4.8 Å; Na⁺: 4.6 Å) [548], faster ion transmission rates, and a lower system reduction potential. Consequently, potassium-sulfur batteries offer unique advantages over lithium-sulfur and sodium-sulfur batteries.

  Compared to its group sibling sulfur, selenium (Se) has higher electronic conductivity, theoretical specific capacity (675 mAh/g), and volumetric capacity (3253 mAh/cm³) [549]. During the charge and discharge process of potassium-selenium batteries, the selenium cathode only involves a one-step conversion reaction, resulting in less formation of polyselenides, thus the stability of the selenium cathode is higher. However, potassium-selenium batteries still face issues such as rapid capacity decay, short cycle life, and safety concerns. In 2017, Guo et al. [550] established the first model of a potassium-selenium battery.

  3.7 Potassium-air batteries

  In 2013, Wu et al. [551] reported on potassium-oxygen batteries for the first time. This battery utilizes potassium metal as the anode, 0.5 mol/L KPF₆/DME as the electrolyte, and porous carbon as the cathode. The authors demonstrated that the potassium-oxygen battery exhibits a discharge potential of 2.47 V, accompanied by a charging overpotential of only 20–40 mV, which is significantly lower than the charging overpotential of lithium-oxygen batteries (~1.0 V). Following extensive testing and characterization, it was established that the discharge product of the potassium-oxygen battery was potassium superoxide (KO₂), formed through a one-electron transfer process involving O₂/O₂⁻. The low overpotential observed during charging and discharging can be attributed to the high conductivity of KO₂, which exceeds 10 S/cm at room temperature. Although the theoretical energy density of potassium-oxygen battery based on potassium superoxide structure is 934.9 Wh/kg, which is much lower than that of lithium-oxygen battery (3450 Wh/kg), the potassium-oxygen battery is still of great research significance due to the high abundance and low cost of potassium metal. Notably, potassium metal exhibits a reactivity significantly higher than that of lithium and sodium metals. Consequently, the potassium metal anode is prone to substantial corrosion during the operation of potassium-oxygen batteries, ultimately compromising the battery’s cycling performance.

  Despite the rapid advancements in potassium-oxygen battery research over the past decade, which have included positive strides in electrolyte selection, anode and cathode optimization, and reaction mechanism understanding, culminating in the successful operation of potassium-oxygen batteries in an air environment, numerous challenges remain to be addressed before practical application. These include anode protection, practical electrolyte design, cathode optimization, and cost evaluation.

  3.8 Multivalent metal-sulfur batteries

  In addition to alkali metal anodes, sulfur cathodes can also be combined with multivalent metal anodes such as magnesium, zinc, calcium, and aluminum to form multivalent metal-sulfur batteries. Compared with alkali metal anodes, multivalent metal anodes have advantages in terms of abundance, cost, and volumetric energy density [552]. Similar to alkali metal-sulfur batteries, the discharge process of multivalent metal-sulfur batteries is also a multi-step solid-liquid-solid transformation reaction process [553]. In the high-voltage platform region during discharge, elemental sulfur is reduced to high-order polysulfides (Snn2−, with n ≥ 4); in the low-voltage high-platform region, high-order polysulfides are further reduced to low-order polysulfides (Snn2−, with n < 4) and insoluble multivalent metal sulfides (such as MgS, ZnS, CaS, and Al₂S₃).

  Although multivalent metal-sulfur batteries share many similarities with alkali metal-sulfur batteries, the reaction pathways of multivalent metal-sulfur batteries differ significantly due to the differences in metal ions. This is primarily because different metal cations in the electrolyte affect the stability of soluble polysulfide intermediates (Snn2−, n ≥ 4), leading to different reaction pathways for metal-sulfur batteries [554]. Furthermore, the insoluble solid polysulfide products formed by different metal cations also have different electrochemical properties [555].

  Currently, the cycling performance of reported multivalent metal-sulfur batteries is significantly lower than that of alkali metal-sulfur batteries [556]. In addition to the common issues from the sulfur cathode, the passivation of multivalent metal anodes in the presence of soluble polysulfides is more severe, and their reversibility is worse compared to alkali metal anodes. Moreover, the higher valence and complex solvation structure of multivalent metal cations also reduce the kinetic reaction rate at the sulfur cathode.

  3.9 Zinc-air batteries

  The theoretical specific capacity of the zinc anode can reach 820 mAh/g, and the theoretical energy density of Zn-O₂ battery is 1086 Wh/g. Although its energy density is lower than that of Li-O₂ batteries, it still has obvious advantages over LIBs. Moreover, zinc is relatively abundant in the Earth’s crust, is inexpensive, safe, non-toxic, and environmentally friendly. Zinc-based oxygen batteries can undergo multiple electrochemical charge and discharge cycles [557]. Since the Zn-O₂ battery was first proposed in 1868, it has been widely studied. By the 1930s, primary Zn-O₂ batteries achieved commercialization and were applied in fields such as hearing aids [558]. To further expand the application range of Zn-O₂ batteries, the development of electrochemically rechargeable secondary Zn-O₂ batteries is very important. However, there are still many issues that limit the development of secondary Zn-O₂ batteries, mainly including: (1) Serious zinc dendrite growth [559]; (2) decomposition of aqueous electrolyte; (3) passivation layer of oxidized zinc on the anode surface [560]; (4) sluggish kinetics of ORR/OER reactions [561]

  In response to the challenges faced by secondary Zn-O₂ batteries, researchers both domestically and internationally have proposed numerous effective strategies. Regarding the zinc anode, modifications have been made, including alloying, designing and synthesizing three-dimensional porous conductive zinc storage matrices, and surface coating of the zinc surface [562,563]. The even deposition of metallic zinc during the battery cycling process can be promoted, thereby suppressing the growth of dendrites and inhibiting the formation of passivation layers during discharge. For instance, research has found that coating the surface of the zinc anode with a layer of carbon nitride protective film, the nitrogen element in which has a high electronegativity, tends to capture Zn²⁺ easily during the charging process, thereby guiding the uniform nucleation of zinc on the anode surface during charging and suppressing the formation of zinc dendrites [564].

  The electrolyte directly affects the anode and cathode reaction processes of Zn-O₂ batteries, and regulating the composition of the electrolyte can also solve the problems faced by Zn-O₂ batteries and improve battery performance. By adjusting the pH value of the electrolyte, the concentration of salts, and the use of additives (such as BTA, DMSO), it is possible to effectively suppress hydrogen evolution side reactions and reduce the corrosion of zinc [565,566]. In addition to traditional alkaline electrolytes, new types of electrolytes are also used in secondary Zn-O₂ batteries. The use of electrolytes containing large-sized hydrophobic anionic zinc salts [567], or quinone-mediated rechargeable zinc-air flow batteries [568], has been reported to differ from the 4-electron reaction path of the traditional alkaline secondary Zn-O₂ batteries, achieving a more reversible 2-electron transfer for ORR/OER reactions, which improves the utilization rate of zinc.

  On the cathode side, the air electrode can effectively improve battery performance through structural optimization and the design of highly efficient bifunctional catalysts. Designing more efficient air electrode structures to facilitate the mass transfer process of gas-liquid-solid three-phase electrochemical reactions is an important way to reduce electrochemical polarization during charging/discharging processes [569]. The performance requirements for the oxygen cathode include: (1) Good gas diffusion channels; (2) Water molecule diffusion channels to complete the mass transfer process of liquid reactants and products; (3) Providing pathways for electron transfer; (4) Providing active sites for three-phase reactions; (5) Easy to produce and prepare, with reasonable cost. Therefore, it is necessary to design the hydrophilic-hydrophobic component ratio and structure of the cathode reasonably, and place the catalyst layer in a rational position to achieve efficient three-phase reactions [570]. In terms of catalysts, the electrocatalysts currently researched and developed are mainly divided into three categories: noble metal catalysts, carbon-based catalysts, and transition metal catalysts [571]. In recent years, bifunctional catalysts with both ORR and OER catalytic sites have become a research hotspot. The general design approach is to physically mix catalysts with different catalytic activities or to introduce additional active sites in situ on a single catalytically active material. Developing non-noble metal catalysts that are cost-effective and have excellent catalytic performance can effectively reduce the overpotential during battery charging and discharging, thereby improving the energy density and energy efficiency of the battery [572]. For instance, Jing et al. [573] introduced a Cu-ONCs catalyst with atomically dispersed Cu-N_x sites anchored on defective carbon nanofibers, achieving a peak power density of 48.28 mW/cm² in neutral Zn-air batteries and exceptional stability over 3000 cycles. The hierarchical porous structure and Cu-N_x coordination synergistically enhance ORR kinetics and active-site accessibility. Dual-atom Fe, Mn/N-C catalysts regulate Fe³⁺ spin states (intermediate spin t_2g⁴e_g¹) via Mn-N coordination, achieving ORR half-wave potentials of 0.928 V (alkaline) and 0.804 V (acidic) and a peak power density of 160.8 mW/cm² in Zn-air batteries [574]. The spin-state engineering enhances O₂ activation and mitigates Fenton reaction-induced degradation. Recently, a pyrolysis-free Co-PTAPP molecular catalyst with precisely engineered Co-N₂ active centers demonstrates superior bifunctional ORR/OER activity, enabling flexible Zn-air batteries with a high-power density (190.3 mW/cm²) and extended discharge time (500 min) [575]. The Co-N₂ configuration optimizes intermediate adsorption energy, validated by DFT calculations.

  Currently, research on secondary Zn-O₂ batteries in the experimental stage has become increasingly mature. However, most of the reported results focus only on one of the issues, and there is a lack of uniform testing standards. How to integrate existing research findings to develop secondary Zn-O₂ battery prototypes with long shelf life, excellent cycle life, high anode utilization rate, and low production costs has become crucial. In addition, establishing a set of unified testing standards to facilitate data integration and comparison will have a profound impact on accelerating the commercialization process of secondary Zn-O₂ batteries.

  4. Sodium-ion batteries

  4.1 Cathode materials

  4.1.1   Na₃V₂(PO₄)₃ and Na₃V₂(PO₄)₂F₃ cathodes

  Na₃V₂(PO₄)₃ (NVP), as a cathode material for sodium-ion batteries (SIBs), exhibits numerous outstanding characteristics [576–581]. For example, its operating voltage (3.4 V) and theoretical specific capacity (117 mAh/g) are both at a relatively high level, and during the charging and discharging process, it shows excellent stability and reversibility. Notably, the NASICON-type NVP possesses remarkable ionic conductivity (10⁻⁹–10⁻¹¹ cm²/s), allowing sodium ions to be quickly inserted and extracted even under high current densities, perfectly aligning with the demands for fast-charging and high-power output. Additionally, NVP demonstrates excellent thermal stability, which significantly enhances the safety of SIBs. It is noted that in high-temperature environments, the internal chemical reactions of SIBs can become intense, potentially leading to safety hazards. However, the good thermal stability of NVP can effectively mitigate such risks, ensuring safe operation of the SIBs in various operating environments. Finally, NVP can directly utilize the production process of LiFePO₄ (LFP), greatly reducing the difficulty and cost of production. Given that LFP has been widely used in the lithium-ion battery field and its production process is relatively mature, applying it to the production of NVP can fully leverage existing production equipment and technology, injecting strong momentum into the industrialization process and accelerating its advancement. Although NVP exhibits the aforementioned advantages, there are still some issues, such as the fact that while NVP shows a high ionic conductivity, its electronic conductivity is relatively poor. This limits the further improvement of the rate performance and energy density of SIBs. Additionally, the high cost of V in NVP contributes to the overall high cost of the material, which is also a significant factor restricting its industrialization. To address this issue, methods such as doping, surface modification and phase regulation were reported [582–588]. For instance, an effective carbon dot incorporation strategy was utilized by Bai’s group [589], which can significantly improve the electrochemical performances of NVP. The optimization mechanism is summarized as follows. First, by incorporating carbon dots with a ζ potential of 4.54 mV, the nucleation and growth process of NVP can be regulated, leading to a reduction in the particle size of NVP, which effectively shortens the ion diffusion path and accelerates the diffusion kinetics. Next, compared to traditional high-carbon-content compositions, the electrical conductivity of NVP is significantly enhanced through the addition of a small amount of carbon (0.76 wt%), which remarkably balances the competition between the electrical conductivity and electrochemical activity of NVP. Finally, the incorporated carbon dots, rich in hydroxyl (-OH) and amino (-NH) functional groups, promote the rapid decomposition of NaPF₆ and accelerate the deposition kinetics of NaF on the surface of NVP, thereby forming a thin and stable cathode electrolyte interphase (CEI) layer rich in NaF, which greatly enhances the interfacial sodium ion diffusion kinetics. As a consequence, a highly reversible capacity of 112.3 mAh/g at 0.5 C, outstanding rate performance of up to 200 C and stable cycling performance of 98.4% retention after 20,000 cycles were achieved in NVP cathode. Besides, Fe with low cost was employed to substitute the V in NVP by Liang’s group [590], and the Fe substitution can not only realize a highly reversible Fe²⁺/Fe³⁺ redox around 2.5 V, but also unlock the V⁴⁺/V⁵⁺ redox around 4 V. As a result, a greatly improved reversible capacity of 148.2 mAh/g at 0.5 C with an outstanding energy density of 501 Wh/kg was achieved. Recently, a single-phase Na_xV₂(PO₄)₃ (1.5 ≤ x ≤ 2.5) was reported by Masquelier’s group through a straightforward synthetic route [591]. Unlike the traditional NCP, the Na_xV₂(PO₄)₃ demonstrated an unusual single-phase sodium ion storage mechanism, characterized by continuous voltage changes during (de)sodiation. As a result, an average equilibrium working voltage of 3.70 V was observed from single-phase Na₂V₂(PO₄)₃ (benefiting from the activation of the V⁴⁺/V⁵⁺ redox couple), which was much higher than that of the traditional NVP (3.37 V), leading to the enhanced theoretical energy density from 396.3 Wh/kg to 458.1 Wh/kg.

  Na₃V₂(PO₄)₂F (NVPF) exhibits a similar three-dimensional framework structure to NVP, which facilitates abundant channels for sodium ion transport. Due to the high electronegativity of F, it can significantly enhance the operating voltage of NVPF (≈3.9 V), resulting in a higher energy density compared to NVP [592]. Furthermore, the incorporation of F somewhat bolsters the thermal stability of NVPF, thereby improving its safety in high-temperature operating environments. Additionally, by controlling the content and distribution of F, the performance of NVPF can be fine-tuned, providing greater feasibility for the optimized design of NVPF and allowing for customization of material properties based on varying application needs. Despite these advantages, the synthesis process of NVPF is relatively complex and necessitates precise control of reaction conditions and raw material ratios. Failure to achieve this may adversely affect the structure and performance of NVPF, undoubtedly increasing production costs and challenges, which is not conducive to large-scale industrial applications. Moreover, although NVPF exhibits a higher operating voltage compared to NVP, its electrical conductivity is low and the interfacial side reactions become more pronounced at the elevated working voltages, leading to the poor and cycling performances. To overcome the challenges, some effective strategies were conducted [593–599]. For example, a tiny high-entropy doping strategy was utilized to develop the Na₃V_1.94(Cr, Mn, Co, Ni, Cu)_0.06(PO₄)₃O₂F cathode by Yuan’s group [600], which can significantly improve electron/ion diffusion kinetics, greatly enhance the high voltage capacity, and inhibit the lattice expansion during the (de)sodiation. Remarkably, a high energy density of 463 Wh/kg with 93.8% retention after 1000 cycles at 5 C was achieved. Besides, a solvent free mechanochemical approach was reported by Hu’s group for in-situ fabricating NVPF with unique features of nano-crystallization and extra Na-storage site [601]. The effective tactic endowed NVPF with an ultrahigh specific capacity of 142 mAh/g at 0.1 C which was higher than its theoretical specific capacity of 130 mAh/g. Recently, a low-temperature synthesis method for developing kilogram-scale NVPF cathode based on a dynamic lock-and-release mechanism was reported by Bai’s group [602], enabling an efficient development of well crystalline NVPF at 90 ℃ within 3 h via using urea as a dynamic lock-and-release agent. Notably, the as-developed NVPF exhibited superior rate and cycling capabilities within a wide temperature range, including room temperature (80.2% retention after 5000 cycles), low temperature (100% retention after 300 cycles at −20 ℃), and high temperature (84.8% retention after 1600 cycles at 55 ℃).

  4.1.2   Layered oxide cathodes

  Sodium-ion layered oxide cathodes have acquired wide attention owing to their superior theoretical capacities, high working potentials, outstanding tap densities and reducing battery cost, which is regarded as the promising candidate for building large-scale storage system. Particularly, Delmas et al. [603] proposed layered oxides can be divided to O2, O3, P2 and P3 types, in which number refers to repetition periods of oxygen ions, and O and P represents the octahedral and tetrahedral coordination environment of Na ions, respectively. Moreover, the structural formula of layered oxide materials can be written as Na_xTMO₂, where TM represents various transitional metals, which directly influence sodium storage ability and lattice structural stability. For instance, Komaba et al. [604] prepared Ti-substituted P2-Na_2/3Ni_1/3Mn_2/3-xTi_xO₂ by high-temperature calcination, delivering a reversible capacity of 127 mAh/g with relatively high discharge voltage of 3.7 V upon the initial cycle, which indicated Ti modification is beneficial to optimize Na storage performance. Notably, Fe element has been proved as the electrochemical active element within Na-based layered cathode, which sharply contrasts to it in Li-based cathode owing to different ionicity of Na/Li-O bond. Specifically, Okada et al. [605] discovered unique Fe³⁺/Fe⁴⁺ redox reaction occurs based on both O3-NaFeO₂ and NaNi_0.5Fe_0.5O₂ cathode with discharge capacities of 80 and 104 mAh/g, respectively. Moreover, Mu et al. [606] firstly indicated reversible transition of Cu²⁺/Cu³⁺ redox couple can be achieved in O3-Na_0.9[Cu_0.22Mn_0.48Fe_0.30]O₂, which exhibits a superior energy density of 210 Wh/kg with neglect capacity fading over 100 cycles when coupled with hard carbon. Such strategy can significantly reduce the inputs cost of energy storge batteries [607].

  Recently, boosting unique oxygen redox reaction can acted as a burgeoning strategy for Na-based layered cathodes to optimize their output capacities [608]. For example, Yabuuchi et al. [609] designed P2-type Na_5/6[Li_1/4Mn_3/4]O₂ by inducing Li ions into transitional metal layers, which displays an additional voltage plateau at high potential region upon the initial charging process, further resulting in an outstanding specific capacity of about 200 mAh/g with the voltage of 1.5–4.4 V. Besides, mild substitution of transitional metal ions by Mg, Al elements and vacancy defects has been proven to activate oxygen-relative anionic charge mechanism [610,611], which can be attributed to the configurational transition from Na-O-TM and inducing nonbonding O 2p orbitals. For instance, Yabuuchi et al. [612] prepared a Mg-substituted P2-Na_2/3[Mn_0.72Mg_0.28]O₂ cathode, delivering a output capacity of over 200 mAh/g at 10 mA/g with the voltage region of 1.5–4.4 V, which can be assigned to combined redox reaction of both Mn ions and oxygen ions. Moreover, Yamada et al. indicated relatively reversible oxygen redox behavior can be observed in Na_4/7[□_1/7Mn_6/7]O₂ with partial Mn vacancies [613], resulting in a discharge capacity of 75 mAh/g from Na_4/7-x[□_1/7Mn_6/7]O₂ to Na_2/7[□_1/7Mn_6/7]O₂ (Fig. 18). Besides, there are numerous research progresses relating to adjust and optimize oxygen redox reaction and structural stability for Na-based layered cathodes [614,615]. Specifically, Rong et al. [616] reported a P3-Na_0.6[Li_0.2Mn_0.8]O₂ cathode with week Mn-O covalent bonding, which suggests persistent high voltage plateau assigned to highly stable oxygen redox behavior between 3.5–4.5 V. Furthermore, They designed a P2-Na_0.72[Li_0.24Mn_0.76]O₂ material [617], which delivered superior discharge capacity of ~270 mAh/g at 10 mA/g between 1.5 V to 4.5 V. Thus, the above results forcibly confirmed rational utilization of additional oxygen redox reaction is beneficial to achieve high-performance Na-based layered cathode for Na-ion batteries.

  Figure 18

  

  Figure 18.  The charge/discharge profiles of Na_4/7[□_1/7Mn_6/7]O₂ cathode upon the 2nd cycle with the voltage range of 1.5–4.7 V. Reproduced with permission [613]. Copyright 2018, Wiley-VCH.

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  4.1.3   Na_xFe_y(SO₄)_z cathodes

  As a crucial category of polyanionic materials, sodium iron sulfate compounds can be denoted by the chemical formula Na_xFe_y(SO₄)_z (NFS). Benefiting from the stronger inductive effect of SO₄²⁻ compared to other polyanionic groups, along with the widespread availability and cost-effectiveness of iron resources, NFS presents the merits of high operational voltage and cost efficiency [618]. Nevertheless, due to the tendency of sulfate radicals to decompose and release SO₂ gas beyond 450 ℃, coupled with the hygroscopic nature of sulfate radicals rendering them unstable in atmospheric conditions, the widespread application of this material is significantly hindered. Na₂Fe₂(SO₄)₃ is one of the representatives of iron-based sulfate materials and was the first to be studied. It was initially synthesized by Yamada et al. [619] via a low-temperature solid-state method. Unlike most NASICON structured A_xM₂(XO₄)₃-type compounds, Na₂Fe₂(SO₄)₃ forms a unique alluaudite-type framework. As shown in Figs. 19a and b, the Fe ions occupy two sites, Fe(1) and Fe(2), located at the center of the octahedron in the crystal. The two octahedra then form Fe₂O₁₀ dimer units through common edges, which are connected to SO₄ by sharing an O vertex to form a three-dimensional (3D) framework [620]. Na₂Fe₂(SO₄)₃ has a high discharge voltage of ~3.8 V (vs. Na⁺/Na) and exhibits excellent cycling stability and good rate performance (Fig. 19c). In addition, Na₂Fe₂(SO₄)₃ also faces a common issue among polyanionic sodium cathode materials: poor electronic conductivity. To enhance its electrical conductivity, Zhang et al. [621] innovatively combined, except for conventional element doping and carbon coating, Na_2.26Fe_1.87(SO₄)₃ and Na₆Fe(SO₄)₄ to construct a bridging multiscale heterogeneous interface (Figs. 19d and e). The results demonstrated that the Na₆Fe(SO₄)₄ phase with abundant 3D Na⁺ diffusion channels and low barriers could enhance the ionic conductivity of the composite material. Meanwhile, Na_2.26Fe_1.87(SO₄)₃ facilitated the adsorption of the ClO₄⁻ and FEC in the electrolyte, thus forming an inorganic-enriched cathode/electrolyte interphase (CEI) on the cathode surface. Consequently, the designed composite delivered considerable rate performance (73.5 mAh/g at 1200 mA/g) and superior cycle stability (80.69% capacity retention after 1300 cycles at 60 mA/g).

  Figure 19

  

  Figure 19.  (a) The structure of Na₂Fe₂(SO₄)₃ projected along the c axis; and (b) local environment of two independent Fe sites. Green octahedra, yellow tetrahedra and blue spheres show FeO₆, SO₄ and Na, respectively. (c) Galvanostatic charging and discharging profiles of Na_2−xFe₂(SO₄)₃ cathode cycled between 2.0 V and 4.5 V at a rate of C/20 (2Na in 20 h). (d) Schematic illustration of the heterostructure and (e) Na⁺ transfer process in the heterostructure. (f) The structure of Na₂Fe(SO₄)₂ along the c-axis. (g) Cyclic voltammograms for the 1^st to 3^rd cycle for Na₂Fe(SO₄)₂/C at a scan rate of 0.1 mV/s and (h) galvanostatic charge–discharge profiles of Na₂Fe(SO₄)₂ and Na₂Fe(SO₄)₂/C at 0.1 C cycled between 1.5 V and 4.2 V. (i) Na₂Fe(SO₄)₂-HC pouch cell cycling performance. (a-c) Reproduced with permission [619]. Copyright 2014, Springer Nature. (d, e) Reproduced with permission [621]. Copyright 2023, Springer Nature. (g, h) Reproduced with permission [622]. Copyright 2019, Royal Society of Chemistry. (i) Reproduced with permission [623]. Copyright 2025 Elsevier.

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  Compared to Na₂Fe₂(SO₄)₃, an alluaudite-type sulfate Na₂Fe(SO₄)₂ was reported by Jiang’s group in 2019 (Fig. 19f) [622]. Na₂Fe(SO₄)₂ possesses a high working voltage (3.6 V) and a capacity of approximately 82 mAh/g (Figs. 19g and h). Notably, Na₂Fe(SO₄)₂ exhibits high thermal stability (580 ℃) and superior air stability for up to 60 days. Subsequently, Ma et al. [623] employed spherical-shaped Na₂Fe(SO₄)₂ particles as a research model to assess the air stability of this material. Through exposure to varying humidity levels, they observed the formation of a hydrate (Na₂Fe(SO₄)₂·4H₂O) on the surface of the material to protect the particles’ bulk from further hydration, and this hydrate could be easily removed by vacuum heating. Besides, its remarkable cycle performance was also demonstrated after assembling pouch-type full cells (Fig. 19i). At present, extensive efforts have been directed towards optimizing Na₂Fe(SO₄)₂ such as carbon layer modification and morphological structure design [624,625]. These endeavors underscore the expansive potential of Na₂Fe(SO₄)₂ as a promising and sustainable cathode material for sodium-ion batteries, attesting to its capacity for high performance and longevity.

  4.1.4   Na₄Fe₃(PO₄)₂P₂O₇ cathodes

  Polyanionic cathode materials, owing to their high operating voltage and robust structural frameworks that confer excellent cycling stability, have garnered significant attention in the field of sodium-ion batteries in recent years, becoming a focal point of research [626]. Among them, iron-based polyanionic compounds, such as NaFePO₄, Na₂FeP₂O₇, Na₂FePO₄F, Na₃Fe₂(PO₄)₃, Na_4-αFe_2+α/2(P₂O₇)₂, hold promise for significantly reducing energy costs and addressing safety concerns. Notably, the Na₄Fe₃(PO₄)₂P₂O₇/C (NFPP) material, first reported by Kim et al. [627] in 2013, has attracted widespread interest due to its relatively high operating voltage (~3.1 V vs. Na⁺/Na), theoretical capacity (129 mAh/g), and minimal volume change (~4%).

  However, the poor electrical conductivity of NFPP limits its practical electrochemical performance. To address this issue, the primary methods for enhancing the electrochemical performance of materials include carbon coating and cation/anion doping. Carbon coating can significantly improve the electronic conductivity of the material surface, and effectively prevent the direct contact between the electrolyte and the material surface, inhibit the erosion of the electrolyte to the material, thus extending the cycle life of the material. Nevertheless, carbon coating is limited to enhancing the surface conductivity and cannot improve the intrinsic conductivity of the material, leaving challenges in areas such as rate performance unresolved. In comparison, ion doping not only improves the intrinsic conductivity of the material but also introduces foreign atoms that support the structural framework, preventing structural collapse and significantly enhancing cycling stability. These optimization strategies provide promising avenues for the practical application of NFPP materials.

  Given this, Qi et al. [628] proposed substituting Fe with Cu, successfully inducing lattice distortions, enriching the local electronic cloud distribution, narrowing the bandgap, and enhancing electrochemical reaction kinetics. This strategy significantly improved the energy density and cycling stability of NFPP. Recently, they also pioneered a scalable synthesis of hollow core–shelled Na₄Fe_2.4Ni_0.6(PO₄)₂P₂O₇ with tiny-void space (THoCS-0.6Ni) via a one-step spray-drying combined with calcination process [629]. Similarly, Li et al. [630] explored Mn substitution to synthesize Na₄Fe_2.7Mn_0.3(PO₄)₂P₂O₇/rGO, resulting in a marked increase in capacity alongside exceptional low-temperature performance. Xiong et al. [631] introduced Mg doping to NFPP, achieving the synthesis of a pure phase, wherein Mg incorporation notably enhanced the cycling stability of the material. Furthermore, Fei et al. [632] designed a dual-site defect strategy by introducing Fe vacancies to suppress impurity formation and doping heteroallene La at oxygen sites to induce sodium vacancies. This approach mitigated electrostatic interactions between sodium ions and the anionic framework, broadened sodium-ion diffusion channels, reduced the bandgap, and enhanced Na-ion migration pathways, thereby endowing the material with excellent rate capability and stable sodium storage performance.

  4.1.5   Prussian blue analogues cathodes

  Prussian blue (PB) and Prussian blue analogues (PBAs) are a class of transition metal compound materials with a three-dimensional rigid open framework structure [633,634]. PBAs have attracted extensive attention due to their numerous advantages such as stable crystal structure, abundant porous structure that facilitates intercalation/deintercalation, excellent redox activity, high theoretical specific capacity, simple synthesis methods, and low cost, becoming one of the preferred targets for SIB cathode materials [635,636]. However, PBAs still face some challenges in the actual synthesis and application process. For example, during the synthesis of PBAs by traditional methods, a large number of Fe(CN)₆ vacancies and crystallization water are easily produced. The existence of vacancies will reduce the crystallinity of the material, destroy the stability of the lattice structure, decrease the active sites, and disrupt the electron transfer pathways. Excessive water can hinder the migration of alkali metal ions and cause side reactions with the electrolyte, leading to the collapse of the material during the sodium ion intercalation process and resulting in the attenuation of cycle performance. Secondly, during the charge/discharge process, PBA materials encounter issues such as irreversible phase transitions, lattice distortion, and the Jahn-Teller effect. These issues not only lead to alterations in the lattice volume but also induce significant structural stress. In addition, metal ions are bonded to organic ligands that lack redox activity, and these ligands fail to establish a conjugation effect with the metal ions. Consequently, this leads to suboptimal carrier mobility and diminished electronic conductivity. Lastly, there are side reactions at the interface between PBAs and electrolytes, and the CEI film formed on the surface has poor stability, which directly affects the cycle life of SIBs.

  The optimization strategies for PBAs mainly include bulk structure optimization, surface structure optimization, and electrolyte optimization, which can be summarized into the following five parts.

  (1) Improvements in synthesis methods and preparation processes. During co-precipitation or hydrothermal processes, the synthesis methods and production processes of materials can be adjusted by controlling the reaction rate, introducing chelating agents, using vacuum drying to remove water, employing reducing agents or protective gases, etc., to optimize the crystal growth process and obtain high-quality crystal structures [637,638]. Peng et al. [639] innovatively designed an additive-free "ice-assisted" method that suppresses the crystallization rate and the oxidation of Fe²⁺, preparing PBA materials with an extremely stable framework and ample Na⁺ storage sites. When operated at 150 mA/g, the PBA sample demonstrates an elevated specific capacity of 123 mAh/g and maintains 64.9% of its initial capacity following 3000 cycles at a higher current density of 750 mA/g. Of paramount significance is the fact that the elimination of defects substantially contributes to the bolstering of thermal stability, resulting in excellent all-climate performance. Furthermore, to prevent the impact of water solvents on the crystal structure, some anhydrous synthesis methods have been developed. For example, Chou et al. [640] utilized a ball-milling solid-state method to design a "salt-in-water" nano-reactor and successfully prepared high-quality MnHCF-S-170 without any additives. It exhibits an impressive specific capacity of 164 mAh/g at a current density of 10 mA/g, and demonstrates excellent Na⁺ storage capabilities across a wide temperature spectrum from −10 ℃ to 50 ℃.

  (2) Elemental composition optimization. Modifying PBA materials through ion doping is also an effective means of optimizing the stability of the lattice structure. The commonly used doping ions nowadays are mainly transition metal ions and other metal ions such as K⁺ and Ba²⁺ [641,642]. Doping can achieve several functions: It can increase the electronic conductivity and the migration rate of Na⁺; larger-radius cations can alleviate lattice strain during the insertion and extraction of Na⁺, enhancing the cycling stability of PBAs; the introduction of certain electrochemically active ions can increase the specific capacity of PBAs. Yu et al. [643] adopted a nickel-copper co-doping modification strategy to develop a novel manganese-based PBA compound (MnCuNi-PBA) for use as a cathode material in sodium-ion batteries. This material was synthesized by co-precipitation, effectively controlling the Jahn-Teller effect in the lattice structure, resulting in a product material with high reversible specific capacity and cycling stability. Recently, high-entropy doping has been applied to PBA materials. By introducing multiple metal ions into the PB crystal structure, it is possible to increase the configurational entropy of the material, optimize electrical conductivity, and enhance electrochemical stability [644,645]. Tang et al. [646] synthesized hollow high-entropy metal hexacyanoferrates (HEPBAs) with five metals of Ni, Cu, Zn, Co, and Mn, in which Mn, Co, Fe, and Cu as active elements are reaction sites, while Ni and Zn as inert elements contribute to structure stability. Thanks to the "cocktail effect", the specific capacity of HEPBAs maintains a 71.1% retention at 100 C compared to 0.1 C, and shows only 22.2% capacity loss after 25,000 cycles at 100 C.

  (3) Surface modification. Surface coating of PBA materials with inorganic compounds, conductive carbon materials, or conductive polymers can effectively prevent the occurrence of side reactions on the material surface and enhance the conductivity, thereby improving the cycle life of the material. Hu et al. [647] coated a Co_xB skin on the surface of manganese hexacyanoferrate (MnHCF) by room-temperature reaction of Co²⁺ and NaBH₄. The Co_xB shell delivers good electronic and ionic conductivity, superb mechanical flexibility, decreased Mn dissolution, and suppressed internal anisotropic stress. The Co_xB coated sample shows a capacity of 98 mAh/g with a capacity loss of 26% over 2500 cycles at 10 C.

  (4) Special structural design. In recent years, the optimization of material performance through special structural design methods such as constructing concentration gradients, and designing core-shell structures, hollow structures or heterostructures has become a research hotspot in the field of PBAs [648,649]. Kang et al. [650] reduced the particle size of FeHCF to ~10 nm and constructed a mixed electronic/ionic conductor with reduced graphene oxide. In addition to the conventional storage mode, this magical structure demonstrates a new mechanism for space charge storage, which allows electrons and ions to be stored separately in the electronic conductor and ionic conductor of the space charge zone, respectively. The space charge storage not only provides additional capacity but also shows ultra-fast charge transfer kinetics, resulting in a high capacity of 79 mAh/g at 50 C for the mixed conductor.

  (5) Electrolyte design. To address the issue of interface stability between PBAs and electrolytes, strategies such as using high-concentration electrolytes, introducing electrolyte additives, or developing solid-state electrolytes can be employed. For instance, in response to the issue of unstable electrode/electrolyte interface in the FeMnHCF||HC full cell with ether-based electrolyte, Chou et al. [651] added sodium difluoro(oxalato)borate (NaDFOB) into the electrolyte, which can form NaF and B-species rich CEI with good mechanical and chemical stability. The durable CEI effectively inhibits the re-decomposition of CEI components and the dissolution of transition metals, realizing outstanding electrochemical performance of the full cells.

  4.1.6   Organic cathodes

  The research history of organic compound cathode materials is quite extensive, and their raw materials can be obtained from biomass. The absence of transition metals makes the material recycling process more environmentally friendly, and the structural design is quite flexible, allowing for good regulation of electrochemical performance. Despite numerous advantages, organic compound cathode materials did not receive enough attention after the commercialization of inorganic cathode materials for LIBs. In recent years, with the increasing demand for energy storage materials, these cathode materials have regained attention [652]. Organic compound cathode materials mainly include polymers, carbonyl compounds, Schiff base compounds, and azo compounds, among others. The main challenges faced by organic compound cathode materials are low electrical conductivity and solubility in organic electrolytes [653]. Wu et al. [654] designed a composite separator material to prevent the diffusion of soluble cathode material dissolution products to the vicinity of the anode, significantly improving the cycling performance of sodium-ion batteries using organic cathode materials. Kim et al. [655] encapsulated a polymer cathode material in carbon nanotubes, resolving the self-discharge issue and enhancing cycling and rate performance, with the modified cathode material still providing a reversible specific capacity of 190 mAh/g at a 5 C rate. Future research on organic compound cathode materials will focus on the development of new material systems and solutions to the aforementioned issues, thereby further promoting the application of these cathode materials in sodium-ion batteries.

  4.2 Anode materials

  4.2.1   Hard carbon anodes

  Sodium ion batteries (SIBs) have risen to prominence in recent years as a leading contender for next-generation energy storage technologies, offering benefits such as abundant resource availability, cost-effective materials, and environmental compatibility. Among SIB components, the anode material is instrumental in defining critical parameters, including energy density, power density, and cycling performance. Although graphite demonstrates excellent performance in LIBs, with a specific capacity of 372 mAh/g and an initial Coulombic efficiency (ICE) above 95%, it is less effective in SIBs due to the unfavorable thermodynamics, which inhibit smooth (de)intercalation within its layers. Hard carbon, an amorphous carbon material, stands out as a superior alternative, featuring wider interlayer spacing and a highly disordered structure ideal for accommodating sodium ions [656,657].

  Hard carbon is a material characterized by a coexistence of amorphous and graphitic regions, exhibiting a multi-scale disordered structure containing graphitic domains, amorphous carbon regions, nano-sized closed pores, functional groups and other defect sites. According to the ’house of cards’ model, hard carbon materials consist of carbon sheets, but they are not fully developed and are significantly smaller than those found in graphite. As a result, hard carbon sheets have more surface and edge sites, alongside various defects such as heteroatoms, functional groups, and missing carbon atoms. These surface and edge sites, together with defects, provide abundant active sites for sodium ion adsorption, while functional groups and heteroatoms enable chemical interactions with sodium ions. The carbon sheets align to form graphitic domains, but the arrangement is more disordered compared to graphite. Furthermore, the interlayer spacing in hard carbon is much larger than in graphite, and varies within a certain range from 0.37 nm to 0.40 nm, compared to the 0.34 nm spacing in graphite. This larger interlayer spacing offers storage space for sodium ions through intercalation. These graphite-like microcrystals are interconnected by amorphous carbon regions, forming closed nanopores that act as storage sites for sodium ions [658,659].

  As discussed above, there are several types of active sites for sodium storage (Fig. 20). (1) Sodium ions can physically adsorb onto the surface and within the open pores, a process primarily governed by the surface area and pore characteristics. (2) Sodium ions can chemically bond with defects or functional groups. (3) Sodium ions can intercalate into the disordered carbon layers, with this process depending on the interlayer spacing and structural disorder of the hard carbon. (4) Sodium ions can fill into the closed pores, potentially forming metallic sodium-like clusters. So far, the sodium storage mechanism in hard carbons remains a subject of debate, impeding progress in the creation of advanced carbon materials.

  Figure 20

  

  Figure 20.  Sodium storage sites in disordered carbon.

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  In 2000, Dahn’s research group firstly reported the "hard carbon" anode, which exhibits a characteristic sodium storage curve with a sloping region (>0.1 V vs. Na⁺/Na) and a plateau region (<0.1 V vs. Na⁺/Na) [660]. Noting striking similarities between that the Li-storage and Na-storage curves for this hard carbon, they proposed that the sodium storage follows a similar insertion/extraction mechanism. For the first time, they proposed the "house of cards" structural model for hard carbon materials, and based on this model, they suggested the "intercalation and pore filling" mechanism. According to this, the slope region corresponds to Na⁺ intercalation between the graphitic microcrystalline layers, while the plateau region is related to Na⁺ filling in the micropores. Komaba et al. [661] verified the above conclusions by means of ex-situ XRD, Raman and SAXS.

  Subsequent studies have refined and debated these mechanisms. Cao et al. [662] identified similarities in charge-discharge profiles of polyaniline-derived tubular hard carbon and graphite in Li-storage, proposing that sodium storage in the plateau region resembles Li intercalation in graphite. Combining in situ XRD, ex situ NMR and EPR data, they introduced the "adsorption-intercalation" mechanism. That is, the slope region corresponds to the adsorption of Na⁺ on defect sites and heteroatoms, while the plateau region is the intercalation of Na⁺ between graphitic microcrystalline layers. Moreover, theoretical calculation confirmed that a layer spacing of 0.37–0.47 nm is optimal for intercalation; too small a spacing prevents intercalation, while overly large spacing is equivalent to adsorption. In another study, Hu’s group analyzed the sodium storage behavior in cotton-derived microtubule hard carbon, finding negligible changes in d₀₀₂ spacing before/after discharge, while a higher photoelectron binding energy in the plateau region [663]. Therefore, they proposed the "adsorption-pore filling" mechanism, where Na⁺ adsorption occurs in the slope region at surface, edge and defect sites, while the plateau region reflects Na⁺ filling in nanopores without intercalation. The above mechanism was confirmed by Xu, Tarascon and other researchers [150]. In summary, the highly variable precursors, diverse preparation techniques, and poor crystallinity of hard carbon materials make it particularly challenging to ascertain the exact origins of the slope and plateau regions. Most of the debates concerning sodium storage center on the active sites for sodium storage, the origins of slope and plateau regions, and the states of sodium in closed pores.

  Hard carbon, with its excellent capacity characteristics and good cycling stability, provides critical support for the commercialization of SIBs. However, overcoming key bottlenecks such as production cost, ICE, and long-cycle stability is essential for the widespread application of hard carbon anode materials. In the future, through the close integration of fundamental scientific research and engineering technology, hard carbon is expected to play a significant role in next-generation energy storage technologies, contributing to the realization of global sustainable energy development goals.

  4.2.2   Alloy anodes

  In recent years, alloy anodes have become one of the research focuses because of their high theoretical specific capacity and suitable potential in sodium-ion battery. Metallic material (Sn, Sb, Bi, etc.) and their alloys store sodium through alloying reactions which can provide higher capacity compared to traditional carbon-based electrode. The metal Sn and Sb have the theoretical capacity of 847 mAh/g and 660 mAh/g, respectively, which is far better than the 372 mAh/g of graphite. However, the volume expansion of the alloy anodes is more serious in the cycle process, which leads to the electrode structure is easy to damage and the active material is seriously pulverized, thus affecting its cycle performance and rate capability.

  In order to remit the volume expansion and improve electrode performance of sodium storage process, researchers have reasonably designed the structure and composition of alloy anodes through various strategies. The various morphologies and structures are formed with different carbon materials alleviate the volume effect and produce coordination effect in alloy composites [664,665]. Kollu et al. [666] have systematically studied the synergistic effect between the alloy particles and the conductive matrix for the reduced graphene oxide anchored tin-antimony alloy composite (rGO-Sn_xSb_y (x ~ y = 1)), which is displayed the capacity of ~320 mAh/g at current density of 500 mA/g after 300 cycles in a sodium-ion battery and the potential for fast charge and discharge applications. Secondly, the sodium storage site is improved and the electrode structure is stabilized by doping and carbon-coated alloy [667,668]. Wang et al. [669] have designed the SbSn alloy nanoparticles implanted in honeycomb N-doped porous carbon (SbSn/NPC), which can effectively offset part of the volume effect and buffer the overall stress, and NPC can produce dozens of Na⁺ adsorption edges/defects and alleviate particle agglomeration to resist severe cracking and crushing during the Na⁺ storage process, thus making the outstanding performance as the anode. In addition, the mechanism of sodium storage for the alloy-based anodes formed by combining carbon materials with multi-element is a research hot spot and direction for the high performance or fast charge in sodium-ion batteries [670,671].

  4.2.3   Titanium-based anodes

  Titanium based materials, including titanium-based oxides [672,673], sulfides [674,675], fluorides [676], phosphates [677–680], and MXene materials [681–685], have advantages of good safety performance and environmental friendliness, and are considered a promising type of anode materials for sodium ion batteries. In 2011, Xiong et al. [686] first studied the performance of TiO₂ as an anode material for sodium ion batteries. The study showed that amorphous TiO₂ could deliver the reversible specific capacity of 150 mAh/g at the 15^th cycle [686]. Since then, the researches on TiO₂ as an anode material for sodium ion batteries has gradually attracted researchers’ attention. There are various crystalline phases of TiO₂, among which anatase TiO₂ has been studied the most as an anode material for sodium ion batteries [687,688]. Beyond that, amorphous TiO₂ [689,690], TiO₂ (B) [691,692], rutile TiO₂ [693,694], hollandite type TiO₂ [695], and dual phase TiO₂ [696,697] have also been reported one after another. Titanates can usually be regarded as mixed metal oxides, such as A₂Ti_nO_2n+1 (A = Na or K, n = 3–8), which can also be written as A₂O·nTiO₂. A₂Ti_nO_2n+1 materials have a monoclinic layered structure, Na⁺ could insert and deinsert between TiO₆ octahedral interlayers. In 2011, Senguttuvan et al. [698] first reported the sodium storage performance of Na₂Ti₃O₇. Na₂Ti₃O₇ can reversibly intercalate and deintercalate 2 mol Na⁺, with a theoretical specific capacity of 178 mAh/g and a low sodium insertion potential around 0.3 V. Because that the sodium storage mechanism is mainly attributed to the intercalation/deintercalation reactions of Na⁺, the main disadvantage of titanium-based oxides is their low specific capacity. For example, because the theoretical specific capacity of TiO₂ is 335 mAh/g, its actual specific capacity is usually less than 300 mAh/g. However, it is worth noting that some low-valence titanium oxides (TiO and Ti₂O) could achieve sodium storage through conversion reactions, so these low-valence titanium oxides could exhibit high capacity, such as Ti₂O electrode delivers a high reversible capacity of 515 mAh/g [699]. Titanium based phosphates, including NASICON (Na⁺ superionic conductor)-type A_xTi₂(PO₄)₃ [700,701], TiP₂O₇ [702], have excellent structural stability and are an attractive type of anode material for sodium ion batteries. A_xTi₂(PO₄)₃ has a three-dimensional network structure, and Na⁺ can migrate in its crystal structure and cause extremely small volumetric strain. However, the low electronic conductivity of 3-D phosphates, limited capacity based on insertion/deinsertion mechanism, and relatively high potential platform significantly hinder its practical application. Two-dimensional titanium-based MXenes show remarkable performance as anode materials for sodium-ion batteries due to their advantage of ultra-large interlayer spacing, excellent electrical conductivity, significant safety performance, large specific surface area, and high sodium-storage activities, so they are receiving increasing attention from researchers [703]. Titanium-based MXenes, include carbides, nitrides, and carbonitrides MXene. Among these, Ti₃C₂T_x MXene is the most eye-catching and has been studied the most extensively, and Er et al. calculated the capacity of Na on Ti₃C₂ to be 351.8 mAh/g [704]. However, titanium-based MXenes face challenges in commercial applications due to the difficulties in the synthesis, the low specific capacity and fast degradation rate upon cycling caused by the inevitably coupling with surface species during synthesis and the aggregation/self-stacking of MXene nanosheets. Unlike other titanium-based materials, the sodium storage mechanism of titanium-based sulfides and fluorides is based on conversion reactions, so they have high theoretical specific capacities. The most commonly reported titanium-based sulfides as anode materials for sodium ion batteries is TiS₂. TiS₂ has a layered structure analogous to graphite, with the Ti ions sandwiched between two packed sulfur layers, and these are held closely together by van der Waals forces. Because its sodium-storage mechanism is based on a four-electron transfer through electrochemical conversion of TiS₂ to metallic Ti, it has a theoretical capacity of 957 mAh/g [675]. However, it suffers from severe volume variation and shuttle effect of the intermediate polysulfides. Theoretically, TiF₃ can achieve three electron transfers when reacting with sodium, so its theoretical capacity is as high as 767 mAh/g. However, fluoride-based materials suffer from sluggish kinetics and poor capacity retention essentially due to low electric conductivity and severe volume variation.

  For improving the sodium storage performance of titanium based anode materials, recombination with carbon-based materials [705–707], heteroatom doping such as Co [708,709], Si [710], and Nb [693,711], engineering of defects [712,713], and designing novel nanostructures for instance nano-wires [714–716], nanotubes [717–719], micro/nano spheres [720–724], and fluffy nanofibers [725] have been widely focused and investigated by researchers. Carbon networks and unique nanostructures are favorable for the transfer/diffusion of electrons/Na⁺ and the structural stability. The engineering of defects, such as oxygen vacancies, provides a promising approach to address the poor electronic and ionic conductivities of titanium-based anode materials. Ion doping could essentially modify the electronic structure of crystalline samples. In addition to these, compositing with high sodium storage capacity materials, such as Sn [726], SnS₂ [727], Sb₂S₃ [682,728], VO₂ [729], MoS₂ [730,731], phosphorus [732,733], has been proved as an effective strategy to promote the specific capacity of titanium based oxides and MXenes. In fact, to promote the performance of materials, some comprehensive strategies are often adopted. A facile strategy was developed by Yao et al. to synergistically engineer the lattice defects (i.e., heteroatom doping and oxygen vacancy generation) and the fine microstructure (i.e., carbon hybridization and porous structure) of TiO₂-based anode, which efficiently enhances the sodium storage performance [710]. The Si-TiO_2-x@C exhibits a high sodium storage capacity (285 mAh/g at 0.2 A/g), excellent long-term cycling, and high-rate performances (190 mAh/g at 2 A/g after 2500 cycles with 95.1% capacity retention) [710]. Introducing oxygen vacancies into core-shell C@NaTi₂(PO₄)_3-x composites could tailor the electronic structures and enhance the Na storage performance, the C@NaTi₂(PO₄)_3-x with oxygen vacancies retaining 108.9 mAh/g after 10,000 cycles at 20 C [700]. TiF₃ subset of C core/sheath electrode exhibits a high capacity of 161 mAh/g at a high current density of 1000 mA/g over 2000 cycles [676]. Wang et al. [734] proposed Ti₃C₂-Sb₂S₃ composites, in which monodispersed Sb₂S₃ uniformly pinned on the surface of Ti₃C₂T_x MXene through covalent bonding of Ti-O-Sb and S-Ti, and Ti₃C₂T_x MXene serves as both charge storage contributor and flexible conductive buffer to sustain the structural integrity of the electrode. Ti₃C₂-Sb₂S₃ composite could deliver a high reversible capacity of 475 mAh/g at 0.2 A/g after 300 cycles, even retaining 410 mAh/g at 1.0 A/g after 500 cycles [734].

  4.2.4   Transition metal sulfides

  Transition metal sulfides (MS_x, M = V, Mo, Fe, Ni, etc.) have been widely studied as anodes for sodium-ion batteries due to their high theoretical capacity. However, due to the conversion reactions, these materials undergo huge volume change during cycling, leading to the pulverization of the active materials and rapid capacity decay. In addition, the low electrical conductivity of these materials usually leads to slow electrochemical kinetics. To address the above problems, researchers have proposed a variety of strategies to optimize the sodium storage performance of MS_x materials, including nanostructure design, composite with conductive materials (carbon, conductive polymer, etc.), element doping and heterostructure construction. Nanostructure design can shorten the migration path of ions/electrons and reduce the migration energy barrier. The nanomaterials have larger specific surface area and can provide more active sites, which is also conducive to the penetration of electrolytes [735]. Composite with conductive materials can improve the electrical conductivity of the transition metal sulfides. At the same time, the conductive material substrates show excellent mechanical stability and can effectively adapt to the drastic volume change during the charge and discharge process, thereby improving the cycling stability of the materials [736–738]. Element doping can regulate local lattices and introduce vacancies or electronic defects, which may promote the migration of Na⁺ in the host material and improve the electrical conductivity of MS_x materials to a certain extent [739–741]. In addition, constructing heterostructures has become an attractive strategy to improve their electrochemical performance. The generation of the built-in electric fields at the heterointerfaces improves the electrochemical reaction kinetics. Moreover, the different redox potentials of the building blocks can promote the structural stability of the heterostructures during the redox processes [742–745].

  Zhao et al. [741] designed a Fe single atom doped MoS₂ hollow multi-shell structure (Fe-M-HoMS) material. As shown in Fig. 21a, the Fe single atoms in MoS₂ promote the electron transfer and the unique HoMS structure shortens the charge diffusion path, thereby achieving excellent rate performance. The self-catalysis of Fe single atoms enhances the reversible transformation between 2H and 1T phases and conversion reaction from Na_x MoS₂ to Mo and Na₂S during the redox process. In addition, the buffering effect of HoMS structure during cycling improves the cycling stability of the material. Therefore, Fe-M-HoMS electrode exhibits a high specific capacity of 213.3 mAh/g at ultra-high current density of 30 A/g (Fig. 21b). Even at 5 A/g, 259.4 mAh/g capacity is retained after 500 cycles and the capacity retention rate reaches 83.68%. Zhang et al. [746] inserted PO₄³⁻ groups into MoS₂ (1T-P-MoS₂), expanding the interlayer spacing of MoS₂ and promoting the migration of Na⁺. In addition, after PO₄³⁻ embedded, it promotes the phase change of MoS₂ from the 2H phase to the 1T phase, thereby improving the electronic conductivity of MoS₂, the phase transition mechanism of MoS₂ is shown in Fig. 21c. The resulting material has ultrafast charging performance (up to 277.1 mAh/g at 40 A/g, Fig. 21d). Furthermore, Sun group [747] constructed the MoS₂/NiS₂ heterostructure material. The redistribution of charge occurred on the MoS₂/NiS₂ heterointerfaces, forming a built-in electric field pointing from NiS₂ to MoS₂, which enhanced the electrochemical reaction kinetics. In addition, the different redox potentials of MoS₂ and NiS₂ in the MoS₂/NiS₂ heterostructure promote the structural stability of the material (Fig. 21e). The MoS₂/NiS₂ heterostructure anode has excellent sodium storage performance, with a capacity of 339.4 mAh/g at 10 A/g and a capacity retention of 480.5 mAh/g after 350 cycles at 1 A/g.

  Figure 21

  

  Figure 21.  (a) Schematic diagram of the advantages of Fe-M-HoMS as anode for SIBs. (b) The rate performance of Fe-M-HoMS. (c) Schematic diagram of the phase transition mechanism of MoS₂. (d) The rate performance of 1T-P-MoS₂. (e) Schematic diagram of the excellent electrochemical performance of the MoS₂/NiS₂ heterostructure. (a, b) Reproduced with permission [741]. Copyright 2024, Wiley-VCH. (c, d) Reproduced with permission [746]. Copyright 2023, Wiley-VCH. (e) Reproduced with permission [745]. Copyright 2024, Elsevier.

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  4.2.5   Organic anodes

  Organic anode materials for sodium ion batteries are predominantly composed of elements such as carbon (C), nitrogen (N), oxygen (O), and hydrogen (H). These materials exhibit a broad array of structural frameworks, low synthesis costs, and substantial molecular structural versatility, which collectively enhance the insertion and extraction of sodium ions [748,749]. These characteristics have garnered significant interest among researchers. However, the performance of organic anode materials is hindered by intrinsic limitations inherent to the materials themselves, including low electronic conductivity, solubility in organic electrolytes, which compromises cycling stability, and other associated challenges. This section provides a comprehensive review of organic small molecules. It examines their energy storage mechanisms, summarizes recent advancements in addressing the inherent issues of organic anode materials, and concludes with a forward-looking perspective on the future development of these materials.

  Organic small molecules can be classified into three principal categories based on the functional groups that participate in reversible interactions with sodium ions during the charge and discharge cycles: Those containing C=O groups (e.g., quinones, ketones, carboxylates, acid anhydrides, and imide derivatives) [750,751], C=N groups (e.g., Schiff bases and pteridine derivatives) [752,753], and N=N groups (e.g., azo derivatives) [754]. Throughout the charge and discharge processes, electrons are reversibly stored in the frontier molecular orbitals of the organic molecules, involving electron removal from the highest occupied molecular orbital (HOMO) and electron insertion into the lowest unoccupied molecular orbital (LUMO). The charge redistribution resulting from the electron exchange is primarily accommodated within the π-conjugated framework of the organic molecules, thereby maintaining the overall structural integrity. During charging, sodium ions can rapidly intercalate into the negatively charged functional groups of the organic molecules, a process facilitated by the relatively large ion transport channels inherent to organic anode materials [755].

  Carbonyl compounds are extensively utilized in the anode materials of sodium-ion batteries, owing to their structural diversity, adjustable conjugation effects, and inductive effects [750,754]. Sodium terephthalate (Na₂C₈H₄O₄) is the most representative organic anode material for sodium ions [756]. During the charging process, each carboxylate group can accept one electron, while the benzene ring in Na₂C₈H₄O₄ transforms into a cyclohex-1,4-diene skeleton. The reduced [C₈H₄O₄]⁴⁻ maintains stability through π-conjugation effects. Notably, the advantage of Na₂C₈H₄O₄ lies in its theoretical reduction potential of only 0.3 V (vs. Na⁺/Na) and an ultra-high two-electron theoretical specific capacity of 255 mAh/g. Additionally, introducing electron-donating groups (such as NH₂, NO₂) can lower the reduction potential of Na₂C₈H₄O₄ and increase the discharge voltage. Here are the discharge voltage rankings of several common organic anode materials [757]: F-Na₂C₈H₄O₄ > NO₂-Na₂C₈H₄O₄ > (COONa)-Na₂C₈H₄O₄ > Br-Na₂C₈H₄O₄ > NH₂-Na₂C₈H₄O₄ > Na₂C₈H₄O₄. Among these, F-Na₂C₈H₄O₄ and NO₂-Na₂C₈H₄O₄ demonstrate higher capacities. However, the formation of irreversible NaF and NaNO₂ results in rapid capacity degradation, thereby limiting their practical application. Extending the π-conjugated system has been shown to enhance the chemical stability, conductivity, and reduce the solubility of organic anode materials. For example, sodium 4,4′-biphenyldicarboxylate delivers a discharge capacity of 100 mAh/g at a current density of 3740 mA/g [758]. Additionally, the anhydride derivative of 3,4,9,10-naphthalenetetracarboxylic acid, sodium salt (Na-PTCDA), has also been investigated as an anode material for sodium-ion batteries [759]. Due to its extended conjugated structure, Na-PTCDA exhibits a stable cycling performance over more than 300 cycles, demonstrating its potential for long-term use.

  Schiff bases and pteridine derivatives, characterized by C=N bonds, exhibit tunable electrochemical activity, with their planar structures and conjugated systems playing a crucial role in stabilizing their electrochemical properties [752,760]. However, compounds containing C=N bonds often encounter challenges such as low theoretical capacity and poor cycling stability. To enhance the capacity of Schiff base electrode materials, the incorporation of Schiff bases with carboxylic acid terminal groups has been proposed [753]. Studies suggest that the addition of carboxylic acid groups can effectively modulate the operating voltages of systems involving carboxylic acid, Schiff base, and carboxylic acid-Schiff base hybrids. Azo compounds containing N=N bonds have recently emerged as potential electrode materials for sodium-ion organic anodes, with the N=N bond acting as an electrochemically active group capable of reversibly binding with sodium ions. However, a major challenge associated with these azo compounds is their high solubility in electrolytes, which leads to poor stability during cycling [754]. To address this issue, the introduction of carboxyl groups has been proposed as an effective strategy to enhance the stability of these materials. While this approach shows promise, it remains in the early stages of research, requiring further exploration to fully understand its impact on the performance and longevity of azo-based organic anodes.

  Despite the significant progress made in the development of organic sodium-ion anode materials, it remains challenging to simultaneously achieve key performance characteristics, including low operating voltage, high specific capacity, structural stability, high initial coulombic efficiency, high tap density, high conductivity, air stability, low cost, and safety without toxicity. Several critical issues persist in the development of organic anode materials for sodium-ion batteries, notably poor chemical stability, high solubility in organic electrolytes, and low conductivity. The chemical stability of organic compounds is primarily governed by the strength of the chemical bonds within the molecular structure. Most organic compounds are covalent in nature, and during the charge-discharge cycles, the covalent bonds are prone to breaking, leading to the formation of free radicals that can interact with active groups along the main chain, thereby compromising the stability and performance of the material. Thus, side reactions contribute to the deactivation of organic electrode materials in sodium-ion batteries. Additionally, the significant volumetric expansion of organic materials during charge and discharge cycles causes severe particle fragmentation and leads to detachment from the current collector, resulting in poor cycling stability. These issues highlight the need for optimizing the molecular structure and morphology of organic materials to enhance their stability as electrode materials. Organic materials, particularly small organic molecules, often exhibit high intrinsic solubility in organic electrolytes, which contributes to rapid capacity fade and poor cycling performance. This solubility issue represents a major barrier to the widespread application of organic materials in sodium-ion batteries, although it also facilitates electrochemical reactions in the solution phase. To address this challenge, solutions may involve appropriate molecular design or composite strategies that incorporate inorganic materials, or the use of aqueous electrolytes, which may offer advantages over organic solvents.

  Moreover, many organic compounds are covalent in nature and lack free electrons or ions, which results in slow charge transfer rates and poor electrical conductivity. Consequently, a significant amount of carbon must be incorporated as a conductive additive during electrode preparation, which diminishes the overall energy density of sodium-ion batteries. To overcome this limitation, enhancing the electrical conductivity of organic electrode materials through strategies such as molecular design, doping, or compositing with conductive polymers or inorganic materials holds promise for improving the performance of organic sodium-ion batteries.

  4.3 Electrolyte chemistry

  4.3.1   Carbonate-based electrolytes

  The common carbonate solvents can be categorized into cyclic and linear types. Cyclic carbonates, characterized by their high dielectric constants, have the potential to weaken the interactions between cations and anions, thereby facilitating the dissociation of sodium salts. This leads to an increased number of free ions and ultimately enhances the ionic conductivity of the electrolyte [761]. The high dielectric constant also strengthens intermolecular dipole-dipole interactions, thereby elevating the freezing point and viscosity of the electrolyte [762]. EC, a cyclic carbonate, is distinguished by its high dielectric constant (89.7), exceptional salt dissolution capability, and stable film formation on the anode electrode. However, it exhibits solid-state behavior at room temperature with a melting point of 36.4 ℃. Linear carbonates demonstrate weaker dipole-dipole interactions and possess lower freezing points and de-solvation energies with Na⁺ in comparison to cyclic carbonates [763], rendering them more suitable for low-temperature electrolytes. Diethyl carbonate (DEC), ethyl methyl carbonate (EMC), and dimethyl carbonate (DMC) exhibit low dielectric constants, low viscosity, and wide liquid ranges. Employing binary or multi-component solvent mixtures to optimize the physicochemical properties of electrolytes is a widely employed strategy [764].

  Currently, widely used ester-based electrolyte combinations include EC/PC, EC/DMC, EC/EMC, and EC/DEC, among which EC/PC is particularly popular. Ponrouch et al. [765] focused on the selection of solvents, testing various parameters such as the voltage window, thermal stability, ionic conductivity, and cycling feasibility (Fig. 22a). Compared to other conventional solvents, the binary EC: PC mixture was identified as the optimal formulation, demonstrating a reversible capacity of 200 mAh/g with excellent rate performance and capacity retention when paired with hard carbon anodes. Shakourian-Fard et al. [766] calculated the binding energy (ΔE_b), solvation enthalpy (ΔH_Sol), and solvation Gibbs free energy (ΔG_Sol) values for Na⁺ interactions with carbonate-based solvents (EC/PC, EC/DMC, EC/EMC, and EC/DEC). The ΔH_(sol) values of the Na⁺-carbonate complexes indicated that the formation of all complexes is thermodynamically favorable. Among pure carbonate solvents and binary carbonate mixtures, the Na⁺ (EC: PC) complex exhibited the highest ΔG_(sol), suggesting that EC: PC is the optimal binary solvent for sodium-ion battery applications (Fig. 22b). The EC: PC binary system has a relatively high viscosity. Ponrouch et al. [767] optimized the electrolyte formulation by adding DMC, DEC, and DME as co-solvents. The ternary solvent system with an EC: PC: DMC ratio of 0.45:0.45:0.1 not only exhibited the highest ionic conductivity and lower viscosity but also derived a SEI film with low polarization, demonstrating excellent compatibility with the hard carbon anode. Based on this ternary formulation, an HC||NVPF full cell was assembled, achieving an average operating voltage of 3.65 V and a discharge capacity of 97 mAh/g over more than 120 cycles. Beyond EC/PC, EC/DEC, EC/DMC, and EC/EMC have also been applied in sodium-ion battery systems. The electrolyte containing the EC/DEC binary solvent and 1 mol/L NaTFSI exhibits a coulombic efficiency (CE) of up to 99.97% in Na||NaCrO₂ cells, demonstrating promising application potential [768].

  Figure 22

  

  Figure 22.  (a) Electrochemical potential window stability and temperature range of 1 mol/L NaClO₄ dissolved in different ester solvents. Reproduced with permission [765]. Copyright 2012, Royal Society of Chemistry. (b) Comparison of binding energy (ΔE_b), solvation enthalpy (ΔH_sol) and solvation-free energy (ΔG_sol) for different ester solvents. Reproduced with permission [766]. Copyright 2015, American Chemical Society.

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  In addition to exploring the impact of multi-component mixed systems on the physical properties of electrolyte systems, an increasing number of researchers are investigating the influence of solvation structures on electrolyte performance. Hu et al. [769] substituted low-melting-point PC (−48.8 ℃) for high-melting-point EC (36.4 ℃), pairing it with the low-melting-point, low-polarity solvent DEC (−74.3 ℃). This modification not only broadened the liquid range of the electrolyte but also reduced the desolvation energy at the interface, thereby enhancing the low-temperature performance of carbonate-based sodium-metal batteries. Ming et al. [770] studied the effect of solvent-solvent interactions on electrolyte performance. Their research revealed that the weak solvent-solvent interactions between EC and DEC could effectively strengthen the Li⁺-PF₆^– interaction, thereby increasing the energy difference between the orbitals of Li⁺(EC)_x(DEC)_y complexes. This, in turn, enhances the reduction resistance of the electrolyte. This strategy has also been validated in sodium-ion batteries, offering a novel approach to improving the reduction stability of electrolytes.

  Based on the above analysis, for carbonate-based solvents, EC, as a beneficial film-forming agent, is often selected for use in combination with other solvents. Factors such as the solvent’s electrochemical window, thermal stability, compatibility with sodium salts, its impact on the SEI, and methods for detecting the composition of the SEI film must be carefully considered. Additionally, the severe degradation of ester-based electrolytes at the anode surface, which compromises their electrochemical performance, remains a major challenge actively addressed by researchers.

  4.3.2   Ether-based electrolytes

  Ether-based electrolytes have attracted extensive attention in the field of secondary battery owing to the fascinating properties of low viscosity, good reductive stability, fast ion diffusion kinetics, and good compatibility with anodes [771]. Lei et al. [772] reported a high reversible K-storage capacity of ~200 mAh/g for dipotassium terephthalate with a capacity retention of 94.6% at 1000 mA/g after 500 long cycles in a 1,2-dimethoxyethane (DME)-based electrolyte (Fig. 23a), while the electrode presented limited capacities, low Coulombic efficiency, large polarization, and rapid capacity decay in the carbonate-based electrolytes. The difference is attributed to the stable and highly K⁺-conductive SEI film formed in the DME-based electrolyte. Subsequently, Li et al. [773,774] found the bulk Bi electrode is gradually developed into a three-dimensional porous network in the ether-based electrolytes upon battery operation (Fig. 23b). The formation of porous structure originates from the unprecedented movement of the surface Bi atoms owing to the chemical adsorption of ether-based molecules (Fig. 23c). It can buffer huge volume fluctuation. Thus, the bulk electrode exhibits high-capacity retentions of 94.4% after 2000 cycles in sodium-ion batteries and 86.9% after 300 cycles in potassium-ion batteries. However, the ether-based electrolytes still face the severe challenge of poor electrochemical stability at high voltage (>4.0 V), and then cannot be matched with the high-voltage cathodes [775]. A common strategy is to regulate the solvation structure of cations to construct anion-derived cathode-electrode interface (CEI) film [776]. Li et al. [777] developed a unique localized high-concentration ether-based electrolyte with diluent participated solvates (Fig. 23d). The participation of the fluorinated diluent in the solvation structure can not only accelerate the decomposition of anions and diluents, improving the oxidative stability of the electrolyte, but also reduce the desolvation energy of Na⁺. As a result, the P’2-Na_0.67MnO₂ cathode possesses a high-capacity retention of 87.3% under a high cut-off voltage of 4.2 V after 350 cycles. Lei et al. [778] reported an electrolyte with anion-dominated solvation at a moderate concentration of 1.5 mol/L by using a weakly coordinated cosolvent ethylene glycol dibutyl ether. The anion-reinforced solvation structure is beneficial to construct the inorganic-rich CEI film (Fig. 23e) and improve the oxidative stability of the ether-based electrolyte to 4.94 V. When applied in the K_0.67MnO₂ cathode, the designed electrolyte exhibits excellent cycle stability with a capacity retention of 83.0% after 150 cycles in the voltage range of 1.5–4.5 V, which is much higher than that of 18.4% in the traditional carbonate-based electrolyte (Fig. 23f). Nonetheless, the oxidative stability of the ether-based electrolytes still needs to be further improved. The improvement of wide-temperature and flame-retardant performance is also the focus of future research on the ether-based electrolytes. In addition, the volatility of the ether-based electrolytes can cause corrosion of the production equipment and increase the internal pressure of battery, thus limiting its commercial application.

  Figure 23

  

  Figure 23.  (a) Schematic illustration of synthesis and reversible K⁺ insertion/extraction of K₂TP. (b) Schematic illustration of morphological evolution of the Bi electrode in the ether-based electrolytes during cycles. (c) Three adsorption models of DME molecule on (012) crystal plane of Bi based on DFT calculation: Bridge, Top and Hollow. (d) Molecule structure of HFME in Na⁺ solvation configuration. (e) XPS spectra of C 1s, F 1s, and S 2p on KMO surface. (f) Cycle stabilities of K_0.67MnO₂ at 4.5 V in the traditional carbonate-based electrolyte and the designed electrolyte. (a) Reproduced with permission [772]. Copyright 2017, Royal Society of Chemistry. (b) Reproduced with permission [773]. Copyright 2017, Wiley-VCH. (c) Reproduced with permission [774]. Copyright 2018, Wiley-VCH. (d) Reproduced with permission [777]. Copyright 2024, Wiley-VCH. (e, f) Reproduced with permission [778]. Copyright 2024, Wiley-VCH.

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  4.3.3   Phosphate-based electrolytes

  Sodium-ion batteries (SIBs) hold great promise for large-scale energy storage applications. However, achieving high safety standards is a prerequisite for their widespread adoption. Conventional carbonate-based electrolytes are highly flammable, posing significant fire hazards. In this context, phosphate-based electrolytes have gained attention for their non-flammable properties (Fig. 24a) [779,780]. The flame-retardant mechanism of phosphates lies in their ability to release phosphorus-containing [P]^• radicals upon thermal decomposition. These radicals can scavenge essential H^• and O^• radicals involved in combustion reactions, effectively suppressing the combustion process [781]. Despite their excellent flame-retardant characteristics, phosphate-based electrolytes suffer from severe decomposition on the anode surface, which drastically undermines electrochemical performance.

  Figure 24

  

  Figure 24.  (a) Electrolyte flammability test for conventional carbonate electrolyte and phosphate-based Electrolyte. (b) Cycling performance of HC||NVP using phosphate-based electrolytes with and without VC. (c) Projected density of states of the 3.3 mol/L NaFSI/TMP electrolyte. (d) Cycling performance and Coulombic efficiency of the Na||HC half-cells using concentrated 3.3 mol/L NaFSI/TMP electrolyte and conventional 1 mol/L NaPF₆/EC: DEC (1:1, v/v) electrolyte [957]. (e) Energy level diagrams of TFEP and TEP. (f) The SEM and TEM images of the cycled HC electrodes the TMP-based (left) and TMP/TFEP-based (right) electrolyte. (g) Photo of the fully charged HC||NFPP pouch cells after package clipping and under a flame. (a) Reproduced with permission [784]. Copyright 2021, Elsevier. (b) Reproduced with permission [783]. Copyright 2022, Royal Society of Chemistry. (c, d) Reproduced with permission [790]. Copyright 2018, Springer Nature. (e) Reproduced with permission [795]. Copyright 2018, Elsevier. (f, g) Reproduced with permission [236]. Copyright 2024, American Chemical Society.

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  Encouragingly, the use of film-forming additives has proven effective in mitigating the above issue. For example, Zeng et al. [782] introduced 10 vol% FEC into a trimethyl phosphate (TMP)-based electrolyte, enabling stable cycling of Sb-based anodes and NaNi_0.35Mn_0.35Fe_0.3O₂ cathodes. Similarly, Du et al. [783] formulated a triethyl phosphate (TEP)-based electrolyte containing 3% vinylene carbonate (VC), which increased the first-cycle coulombic efficiency by 34.4% compared to a baseline electrolyte without VC, and the cycling stability was significantly improved (Fig. 24b). Zhu et al. [784] further demonstrated that the simultaneous incorporation of FEC and 1,3,2-dioxathiolane 2,2-dioxide (DTD) into a phosphate-based electrolyte facilitated the in-situ formation of a robust SEI layer on Na. These findings highlight that suitable additives can effectively suppress the reductive decomposition of phosphates on anodes, thereby enhancing electrochemical performance.

  Another viable approach involves replacing traditional conductive salts with those exhibiting superior film-forming capabilities, enabling the use of single phosphate solvents [785]. Younesi et al. [786,787] extensively explored this strategy by replacing conventional salts with sodium bis(oxalato)borate (NaBOB), enabling the stable operation of hard carbon (HC) electrodes in a 0.38 mol/kg NaBOB in TEP electrolyte [788]. To address the low conductivity of such electrolytes, amides or N-methylpyrrolidone (NMP) solvents were introduced, effectively doubling the conductivity while preserving the inherent non-flammability of phosphate-based electrolytes [789].

  Tuning solvation shell structures offers another effective solution to improve the compatibility of phosphate-based electrolytes with anodes. Wang et al. [790] increased the concentration of sodium bis(fluorosulfonyl)imide (NaFSI) to 3.3 mol/L, promoting the preferential decomposition of anions over TMP on HC (Fig. 24c). This electrolyte supported stable cycling of HC electrodes for at least 1200 cycles (Fig. 24d). Similarly, Jiang et al. [791] achieved stable cycling in sodium dual-carbon batteries by increasing the sodium bis(trifluoromethylsulfonyl)imide (NaTFSI)/TMP ratio to 1:2. By introducing non-coordinating hydrofluoroether diluents into phosphate-based electrolytes, the local salt concentration is increased, resulting in a solvated structure akin to that of high-concentration electrolytes, which effectively prevents the decomposition of TMP [792,793]. Additionally, Ma et al. [794] incorporated trifluorotoluene (PhCF₃) as a solvent reorganization additive, which reduced the number of free TEP molecules through intermolecular interactions, thus mitigating phosphate decomposition.

  A promising alternative to stabilizing anodes involves molecular-level redesign of phosphate solvents to impart intrinsic film-forming abilities, obviating the need for additives or high salt concentrations. Jiang et al. [795] proposed using fluorinated tris(2,2,2-trifluoroethyl) phosphate (TFEP) as a solvent, which facilitated the formation of NaF-rich SEI layers and enabled HC electrodes to operate effectively with just 0.9 mol/L NaFSI (Fig. 24e). Chen et al. [236] incorporated this fluorinated phosphate with non-fluorinated TMP, significantly enhancing the SEI on HC (Fig. 24f), thereby facilitating stable cycling of HC||NFPP pouch cells for 2000 cycles in a 1.22 mol/L NaClO₄ TMP/TFEP (3:2 by mol) electrolyte. Remarkably, this electrolyte-maintained functionality even under ignition conditions (Fig. 24g).

  In summary, phosphates, as non-flammable solvents, play a pivotal role in enhancing the safety of sodium-ion batteries. Their poor compatibility with anodes can be effectively addressed through additive strategies, adjustments to conductive salts, solvation structure tuning, and solvent fluorination. These promising strategies lay a solid foundation for the development of high-safety sodium-ion batteries.

  4.3.4   Polymer-based solid-state electrolyte

  Polymer-based solid-state electrolytes are carbon-based polymers that incorporate heteroatoms such as oxygen or nitrogen, which can solvate embedded Na⁺ ions [796,797]. While these materials exhibit a broad range of physical properties, polymers are generally more flexible and cost-effective compared to ceramics and glasses [798]. They can be further classified into fully solid polymer electrolytes (SPEs) and gel-polymer electrolytes (GPEs). Unlike conventional batteries that utilize liquid electrolytes and separators, GPEs do not contain free-flowing solvents; instead, the polymer matrix integrates with liquid solvents to form a cohesive material. This integration minimizes electrolyte leakage and mitigates associated flammability concerns [799,800].

  Considering the arrangement and structural configuration of various monomers within a copolymer, it can be inferred that there exist countless polymers capable of conducting Na⁺ ions [796,797]. Nevertheless, a comparative number of polymers have been explored as GPEs or SPEs for sodium metal batteries (SMBs). The most commonly used polymers are poly(ethylene oxide) (PEO), polymerized ionic liquid (PIL), polyvinylidene fluoride (PVDF), poly(methacrylate) (PMA), poly(ethylene glycol) (PEG), and so on [801–803]. PEO and PEO-based SPEs have been extensively investigated since Wright et al. demonstrated that PEO can solvate alkali metals [804]. Among previously reported work, Song et al. [805] designed a hybrid PEO-silica-ionic liquid polymer electrolyte with good conductivity at room temperature for SMBs. The cell capacity could reach 90 mAh/g at room temperature and 210 mAh/g at 60 ℃. Janakiraman et al. [806] prepared a PVDF-based GPE for SMBs, which exhibited a high ionic conductivity of 1.08 × 10⁻³ S/cm and a wide voltage window (≥5 V vs. Na⁺/Na). Tao et al. [807] designed a composite polymer electrolyte comprising of poly(methacrylate) (PMA) and poly(ethylene glycol) (PEG) and achieved an ionic conductivity of 1.46 × 10⁻⁴ S/cm at 70 ℃ and a wide electrochemical stability window ≥4.5 V.

  Currently, the performance of these polymer electrolytes still suffers low ionic conductivity, poor mechanical and thermal stability, and high interfacial resistance, influencing their commercial acceptance for energy storage devices. Common strategies to improve amorphicity include embedding inorganic nanoparticles, copolymerization, crosslinking, and blending with other polymers [798,801]. To tackle these burning issues, strategies to couple inorganic (e.g., SiO₂, TiO₂, or Al₂O₃, among many others) and polymer electrolytes as composite polymer electrolytes (CPEs) [808–810]. Fan et al. [811] developed an aluminum oxide (Al₂O₃) filler-modified composite solid electrolyte via in situ thermal polymerization of hexanediol diacrylate (HDDA) trimethylolpropane trimethacrylate (TMPTMA) and achieved a high ionic conductivity of 5.59 × 10⁻³ S/cm at 25 ℃. The addition of plasticizers in the electrolyte could improve the polymer chain mobility, and create liquid phases within the polymer thus enhancing the ionic conductivity. For example, Xin et al. [812] prepare a succinonitrile-plasticized polymer electrolyte with FEC additive, which effectively prevents the formation of sodium dendrites, thus enabling Na||Na₃V₂(PO₄)₃ full cell with a capacity retention of 93.0% at 1 C rate after 1200 cycles under room temperature. The high interfacial resistance between Na metal and electrolyte results in unstable SEI and irreversible sodium dendrite growth thus reducing Coulombic efficiency and thereby rapid capacity decay [813]. Recently, the in-situ polymerization method has garnered significant attention due to its ability to facilitate the in-situ polymerization of precursor solutions containing monomers, sodium salts, and initiators directly within battery systems [814,815]. The precursor solutions effectively penetrate electrodes while forming SPEs on their surfaces or interiors under external energy influences, such as heating, initiators, or ionizing radiation, thereby establishing optimal interface contact and promoting a continuous pathway for Na⁺ conductivity. Tao et al. [816] designed a poly(butyl acrylate)-based GPE via an in-situ polymerization method, which achieved ultrastable plating/stripping over 900 h at 0.2 mA/cm², and ultralow overpotential of 233 mV at 1 mA/cm².

  Up to now, the performance of polymer-based solid-state electrolytes has been greatly improved, but the large-scale practical application for SMBs still has a long way to go. In the future, researchers could make efforts in the following aspects: Improving ionic conductivity, developing new methods, avoiding the use of unnecessary organic solvents, simplifying manufacturing procedures, and reducing production costs.

  4.3.5   Inorganic solid-state electrolytes

  Analogous to lithium-ion inorganic solid-state electrolytes, sodium-ion inorganic solid-state electrolytes are primarily classified into four categories: Oxides, sulfides, halides, and hydrides. These categories typically exhibit considerably greater mechanical strength compared to polymeric solid-state electrolytes. Among the sodium-ion oxide solid-state electrolytes, Na-β″-alumina [817,818], tellurium-based layered oxides [819,820], and NASICON-type phosphates [821] are the most extensively investigated. Ceder et al. [822] utilized a high-entropy design to create Na_3.5Mg_0.1Sc_0.15In_0.15Ti_0.3Hf_0.3ZrSi₂PO₁₂, finding that the local disorder in high-entropy materials can effectively enhance site percolation and improve ionic conductivity, which shows a high bulk conductivity of 3.3 mS/cm at room temperature. Apart from phosphates, silicates could also be used as solid-state electrolyte, including Na₅YSi₄O₁₂ [823] and Na₅SmSi₄O₁₂ [824,825], which exhibit room-temperature conductivities exceeding 1 mS/cm (with Na₅Sm_0.8Si₄O₁₂ demonstrating a conductivity of 5.61 mS/cm [826]), which suggests a promising avenue for the development of dense and stable solid-state electrolytes. Despite the oxides demonstrating high conductivity and favorable air stability, their considerable hardness and brittleness present substantial challenges in the fabrication of all-solid-state batteries.

  Sulfides, halides, and hydrides, characterized by relatively softer mechanical properties and enhanced conductivity, are experiencing rapid advancements. Yao et al. [827] present a study on a series of oxysulfide glass solid electrolytes (Na₃PS_4-xO_x, 0 < x ≤ 0.60) characterized by a fully homogeneous glass structure, showing stronger and denser glass networks than pure sulfide. With the incorporated oxygen preferentially contributing to the formation of structure-bridging oxygen. Na₃PS_3.4O_0.6 can form a self-passivating solid electrolyte interphase at the sodium-electrolyte interface. Oxygen bridging strategy is also effective for halide. Hu et al. [828] replaced the chlorine of tetrachloroaluminates with oxygen (MAlCl_4–2xO_x, M = Li, Na, 0.5 < x < 1) and obtained the organic-polymers-like viscoelasticity and achieved a high conductivity (>1 mS/cm), with the bridging oxygen forming Al-O-Al networks. Various alternative synthetic pathways were investigated by modifying precursor materials, and kilogram-level electrolyte materials could be obtained [829–831]. Although halides generally exhibit stability with high-voltage cathodes, their inadequate reductive stability remains a pressing issue that requires urgent resolution. Compared to halides and sulfides, hydrides exhibit notable advantages in terms of reductive stability, even when using sodium metal as the anode material [832,833]. Meng et al. [834] successfully developed an anode-free all-solid-state sodium battery, which exhibited stable cycling performance over several hundred cycles, by employing Na_0.625Y_0.25Zr_0.75Cl_4.375 as the catholyte and Na₄B₁₀H₁₀B₁₂H₁₂ as the anolyte. Furthermore, they identified and elucidated four critical factors essential for the design of anode-free solid-state cells.

  5. Potassium-ion batteries

  5.1 Cathode materials

  5.1.1   Layered oxide cathodes

  In Fig. 25, among many cathode materials, layered transition-metal oxide (A_xMO₂, A = K; M = Mn, Fe, Ni, Cr, Co, and their combinations) stand out due to their low cost, layered framework, high specific capacity, and simple synthesis [835,836]. However, there are still some challenges such as irreversible phase transition, Jahn-Teller distortion caused by transition metal ions dissolution, electrode/electrolyte interface instability and poor air stability hinder their development [837–839]. Achieving superior cycling stability, high capacity, and excellent rate performance simultaneously remains difficult [840–842]. Traditional transition metal layered oxides relies on a single metal cation redox couple to provide capacity, its cycle stability and capacity need to be improved [843,844]. Effective structural regulation can inhibit irreversible phase transitions, hinder Jahn-Teller effect and enhance the electrochemical performance [845]. The stability and specific capacity of A_xMO₂ can be enhanced by adjusting the K⁺ ions content. Introducing water molecules into the interlayer of transition metal layered oxides to synthesize birnessite can effectively support the structure and increase the interlayer spacing, thereby enhancing ions migration kinetics and structural stability [846]. Based on different complementary characteristics of metal ion redox pairs, the partial substitution of electrochemical active/inert cations for transition metal sites can effectually regulate the stability of the crystal framework [847]. Specifically, the replacement of electrochemical active elements such as Co, Ni, Mn and Fe improves the cycle stability without sacrificing capacity [848]. The substitution of electrochemical inert elements such as Mg, Cu and Al can enhance the cycling stability of A_xMO₂ materials and obtain superior electrochemical performance [849]. Depending on the combination of elements, these materials can be categorized into binary, ternary, or high-entropy layered oxides [850]. The poor air stability also restricts the development of layered oxide materials. Many coating materials, such as Al₂O₃, FePO₄, K₂C₄O₄, and AlF₃, have been used to enhance the cycle ability of A_xMO₂ [851]. By modifying the electrode/electrolyte interface of A_xMO₂, the electrochemical performance can be stabilized, which increases its application potential [852]. The design of multiphase composites can show the synergistic effect of different phases and obtain superior electrochemical properties [853]. Besides, the cycle stability of layered oxide can also be optimized by composite phase, gradient structure and micro-nano structure [854]. The improvement of synthesis and processing technology, the deeply study of internal reaction principles, and the development of high-pressure electrolytes are also of great significance for the use of A_xMO₂ cathode materials [855–857].

  Figure 25

  

  Figure 25.  Summary about the mechanism, suppression strategies, and effective utilizations and prospects of the Jahn–Teller effect in sodium - ion batteries [848,849,855–857].

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  5.1.2   Polyanionic compounds

  In Fig. 26, polyanionic compounds, composed by polyanion groups and transition metal elements, have been widely investigated as the most promising cathode materials for PIBs [858–860]. The adjustable inductive effect of the redox electric pair allows for potential regulation, enabling high potential in these materials [861–863]. In addition, polyanionic compounds exhibit the relatively stable structure, superior thermal stability and high safety, making them suitable for long-term stable cycling performance and large-scale stationary energy storage systems [864,865]. Various polyanionic structures, including phosphate, fluorophosphate, pyrophosphate, sulfates, have been extensively excavated in recent years. The characteristic atomic structure of polyanionic compounds typically consists of tetrahedral XO₄ groups (X = P, S, Si, etc.) and corner-shared MO₆ octahedron (M = transition metal element) due to their diverse framework and the relatively low alkali metal ion migration barrier [866,867]. The strong inductive effect from adjacent anion groups contribute to higher operating potential and energy densities, resulting in satisfactory cycling performance. Representative polyanionic cathode materials such as K₃V₂(PO₄)₃, KFeSO₄, K₂MnP₂O₇, and KVPO₄F have garnered significant attention [868–870]. However, their inherently poor electrical conductivity and slow kinetics limit their performance [871]. Strategies to address these issues include surface modifications (e.g., carbon or oxide coatings) and electrolyte optimization to enhance ionic and electronic diffusion at the electrode/electrolyte interface [872]. The doping or substitution of electrochemically active and inert elements can effectively optimize the hierarchical structure of polyanionic compounds, thereby enhancing their ion diffusion kinetics and rate performance [873]. In addition, morphological design can also improve the charging and discharging environment of polyanionic material, thus improving the cycle performance [874]. Recent advances in polyanionic cathode materials show great promise in PIBs compared to other types of cathode materials [875,876].

  Figure 26

  

  Figure 26.  Summary about the mechanism, suppression strategies, and effective utilizations and prospects of the Jahn–Teller effect in sodium - ion batteries [863,870,875,876].

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  5.1.3   Prussian blue analogues

  Prussian blue analogues (PBAs) have emerged as the most competitive cathode materials for potassium-ion batteries, primarily because of their remarkable features, namely low cost, simple synthesis process, controllable structure, and excellent cycling stability [877,878]. The chemical formula of PBAs, denoted as K_xM[Fe(CN)₆]_1-y.wH₂O, presents an open three-dimensional framework structure. This distinctive structure generates open ion channels and spacious interstitial sites, which are highly favorable for the rapid insertion and extraction of large-sized K⁺ ions [879]. Fe-PBAs were among the initial cathode materials studied for potassium-ion batteries [880]. However, materials synthesized via conventional methods usually have problems of low crystallinity and high-water content (Figs. 27a-c). These deficiencies severely impair the electrochemical performance of Fe-PBAs. Fortunately, the introduction of chelating agents has proven to be an effective remedy. As shown in Fig. 27d, this method effectively slows down the precipitation rate of PBAs during synthesis, thereby reducing the number of defects and the amount of crystallization water within the crystals, slowing down the lattice expansion during the cycling process, and consequently enhancing the electrochemical performance of PBAs [881].

  Figure 27

  

  Figure 27.  (a) Defect-free crystal structure with 1/4 of the vacant sites occupied by water (green spheres) and 1/3 of the empty sites. (b) TEM of KFeHCF-E nanoparticles. (c) The corresponding charge/discharge and color-mapped curves for in situ XRD patterns for different angle ranges. (d) Schematic crystal structure and (e) galvanostatic charge-discharge voltage profiles of KMF-EDTA sample. (f) In situ substitution of Mn with Fe in KMnF. (g) The 40th charge-discharge curves of KMnF in the pure and modified electrolytes. (h) The ultralong cycling stability of the KMnF electrode in the modified electrolyte. (a) Reproduced with permission [879]. Copyright 2020, Springer Nature. (b, c) Reproduced with permission [881]. Copyright 2023, Wiley-VCH. (d, e) Reproduced with permission [882]. Copyright 2021, Springer Nature. (f-h) Reproduced with permission [883]. Copyright 2021, Springer Nature.

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  Compared with Fe-PBAs, Mn-PBAs have drawn more extensive research attention. This is due to several factors, including their lower cost, abundant resource reserves, and the existence of Mn²⁺/Mn³⁺ redox-active pairs with relatively higher potentials, which contribute to their higher energy density. Among various Mn-PBAs, K_1.94Mn [Fe(CN)₆]_0.994·0.08H₂O is a prominent candidate. As shown in Fig. 27e, it can provide an outstanding reversible capacity of approximately 154.7 mAh/g and a high-voltage discharge of 3.9 V (vs. K⁺/K) [882]. Nevertheless, during the battery cycling process, Mn in Mn-PBAs tends to dissolve into the electrolyte, resulting in an irreversible capacity loss. To tackle this issue, introducing Fe³⁺ into the electrolyte and enabling the formation of Fe-N₆ bonds within the framework during the discharge process has been proposed (Fig. 27f). This strategy successfully preserves the structural integrity of the material. As a result, K_1.82Mn[Fe(CN)₆]_0.96·0.47H₂O can withstand over 130,000 cycles with only a negligible capacity loss (Figs. 27g and h) [883]. Furthermore, doping PBAs with different transition metal ions has been shown to be an effective way to modulate their lattice parameters and redox properties. This, in turn, significantly improves the stability and rate capability of the materials [884]. Additionally, it is worth emphasizing that combining PBAs with other conductive carbon materials and accurately regulating the formulation of electrolytes are also crucial strategies for enhancing the cycling stability and charge-discharge capacity of these materials [885,886].

  5.2 Anode materials

  5.2.1   Carbon-based material

  Carbon materials are abundant, environmentally friendly, and have excellent stability, and are considered by researchers to be the most promising anode materials for potassium ion batteries. In the family of carbon materials, carbon quantum dots, graphite, graphene, carbon nanosheets, carbon nanospheres, and biomass-derived carbon are widely used in potassium ion batteries [887,888].

  Carbon quantum dots with high specific surface area and abundant active sites have natural advantage in potassium ion battery storage. The smaller size of quantum dots can alleviate volume expansion during charging and discharging, effectively enhancing the cycle life of potassium ion batteries [889]. Graphite as a stable carbon material is now widely used in commercial LIBs. Due to the large ionic radius of potassium ions, graphite faces serious volume expansion problems when used as an anode in potassium ion batteries. However, the potassium ion storage performance of expanded graphite prepared by modifying graphite is excellent. Wu’s team used highly efficient mechanical ball milling to prepare defect-rich expanded graphite carbon nanosheets, which had a capacity of 300.1 mAh/g after 1800 cycles at 1 C when used as the anode in a potassium-ion battery [890]. Graphene, a unique two-dimensional material, has also been applied to the anode of potassium ion batteries. Cary L. Pint group synthesized a nitrogen-doped multi-layer graphene that enables potassium ion batteries to reach a capacity of 350 mAh/g [891]. Some researchers have also modulated the layer spacing by controlling the degree of graphene graphitisation to improve rate properties. Besides, Cao groups adopted a one-step arc discharge method synthesized a N, P co-doped graphene to improve potassium ion storage performance [892]. To enhance the potassium ion storage properties of carbon materials, researchers often carefully design the structure of carbon materials. Carbon nanosheets and carbon nanospheres are two anode materials with structural advantages for potassium ion batteries. For examples, Tang groups synthesized a carbon nanosheets with ultrahigh content of nitrogen for potassium ion battery [893]. Carbon nanosheets exhibit high specific capacity and excellent cycle life when used as anode in potassium ion batteries due to their stable structure and abundant active sites. Carbon spheres as a unique structure used in potassium ion batteries can promote the penetration of electrolyte, more conducive to the formation of a stable solid electrolyte film. Liu’s group adopted polypyrrole coated polystyrene methods synthesized multiple active sites decorated porous hollow carbon nanospheres [894]. Hollow carbon spheres can effectively alleviate volume expansion during charging and discharging and enhance cycle stability. Another class of carbon materials with promising applications are biomass-derived porous carbon [895]. Biomass-derived porous carbon is inexpensive and structurally rich, and a variety of biomass-derived carbon is currently used as anode for potassium ion batteries, including potato, sisal, and lignin [896,897]. Widespread use of porous carbon from biomass will significantly accelerate commercialization of potassium ion batteries.

  5.2.2   Alloy materials

  Alloy-type materials are primarily composed by the metal components in IVA and VA groups and intermetallic compounds formed between these metal elements, especially refers to Si, Ge, Pb, P, Sn, Sb, Bi, and the intermetallic compounds formed among these elements such as Sn₄P₃, Se₃P₄. Since the discernible differences of involved redox reactions of alloy materials with those of embedded anode materials, in which large volume changes during alloying and de-alloying results in the pulverization of anode materials, where the electrode material can seriously be dislodged from the collector. Additionally, the sluggish reaction kinetics significantly affects the application of alloy materials in potassium ion batteries (PIBs). Up to present, the incorporation of high conductive matrix, the construction of interconnected framework, the dedicatedly fabricated heterojunction structure of alloys, the nanosize/nanocluster engineering and the optimization of geometric structure are affective protocols to resolve the prevalent challenges posed in PIBs [898,899].

  Among the tremendous alloy-type candidates, P-based materials has been heralded to be one of the promising PIB anodes with developmental potential due to its highest theoretical capacity and abundant reserves in resources [900]. Qin et al. [901] constructed a simple but efficient magnetic field assisted protocol to delicately fabricate hollow red P nanosheets encased well in 3D N-doped carbon nanosheets/nanotubes framework with trace Fe (Fe@NCNS/NCNT). The compositing of phosphorus and other elements into multivalloy complex are also witnessed to be a promising strategy to mitigate the volume expansion/extraction force and promote the kinetics behaviors for K⁺ storage, such as Sn and P hybrid anode Sn₄P₃/C, CuP₂/C [902], Se₃P₄@C [903]. Except for the large volume changes during repetitive cycling, low electrical conductivity, and detrimental electrolyte decomposition due to highly reactive phosphide species seriously hinder their potential application of phosphorus-based materials for alkali metal-ion batteries. The metric species with a lower reaction surface are introduced to address above problems. Qin et al. [904] designs a novel, scalable viscoelastic and K⁺ conducting SEI formed via UV-induced polymerization of phenyl carboxylic acid (PCA) with elastic external layers terminated by hydroxide and chloride groups. This interlayer composes a resilient ultrathin dihaliderich KF, KF_xCl_1−x quasibinary protective layer over nanosized black phosphorus (BP) complexes, conferring an ingenious long cycle stability, excellent safety and even broad temperature adaptability (Fig. 28).

  Figure 28

  

  Figure 28.  Schematic illustration of the SEI evolution for Co-NBC@BP@PCA. Reproduced with permission [904]. Copyright 2024, Wiley-VCH.

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  Silicon alloy anode materials are also among the most studied samples due to their high theoretical specific capacity, and the referable safety in cell devices. Identically, the massive volume changes and the fragile interface stability potentially cause poor stability of the electrode material and adversely affecting the cycling performance of the battery [905,906]. Exhausted endeavors were performed to expedite the K⁺ transport rate via the combination carbon matrix and the construction of pore structures [907]. Huang et al. synthesized a 3D silicon-diamond (Si-DY) composite, comprised by butadiynyl units and sp³-incorporated silicon atoms, demonstrating a rigid diamond-like backbone structure. Si-DY excels diffusion kinetics of alkali metal ions (Li, Na, K), beneficial for its affluent dienes and multiple internal channels, and shows a high reversible capacity. Therefore, the featured 3D amorphous C skeleton can effectively buffer volume expansion, increases ionic/electronic conductivity and improves the structural integrity.

  Germanium (Ge) has successfully been considered one of the promising anode materials derived from its abundant reserves and its relatively high theoretical specific capacity [899,908,909]. However, its long-term cycling stability, the energy density and K⁺ transport kinetics remains further promotion. He et al. [910] demonstrated firstly the electrochemical behavior of Ge anodes in PIBs. Nanoporous Ge anodes were prepared via chemical dealloying protocols based on the employment aluminum (Al) as a sacrificial template. Noteworthily, Ge electrodes with interconnected porosity and small ligaments delivers long-dural cycling stability, which is paramount important for surmounting the kinetic challenges posed by large K⁺. Furthermore, Moreover, Chen et al. [911] prepared Ge@C-CMK composites by encapsulating germanium nanospheres within a porous carbon matrix (Ge-CMK) and subsequently further modified another amorphous carbon layer (Ge@C-CMK). The unique channel-like pore structure, coupling with the amorphous carbon framework, not only ameliorates the volume variation of germanium particles but also considerably enhance the electrical conductivity of Ge and facilitates K⁺ diffusion, underscoring the importance of porous framework assisting in achieving remarkable stability and high capacity.

  Bismuth (Bi) metal, featured with its high theoretical capacity and a relatively low electrode potential (−2.93 V) has served as an ideal candidate for anode based on the alloying reaction mechanism. Bismuth demonstrates high multiplicity and stable K⁺ storage mechanism due to their layered structure advantages, low cost [912]. However, the distinct large volume change (~400%) experienced in practical applications poses significant challenges to bismuth metal. It has been widely considered that the incorporation antimony (Sb) with bismuth alloys can harvest an intentional synergistic effect. This synergistic avenue not only fuses the respective advantages of both elements but also significantly improves the stability of the electrode material. Liu et al. [913] proposed a strain-relaxation structure with well-designed Bi/Bi₂O₃ nanodots encased within amorphous carbon sheets (Bi/Bi₂O₃ NDs@CSs) by elaborately rendered Bi₂O₃ to tailor Bi/Bi₂O₃ nanodots to significantly boost K⁺ storage performance (Fig. 29). Such strain relaxation model serves as an efficient mitigation against stress and strain induced by volume expansion, and an outstanding rate capability and exceptional long-term cycling behaviors can be yielded.

  Figure 29

  

  Figure 29.  Synthesis steps for Bi/Bi₂O₃ NDs@CSs. Reproduced with permission [913]. Copyright 2023, Wiley-VCH.

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  Antimony (Sb) is extensively investigated as an anode material for potassium-ion batteries due to its abundant availability, inherent safety, and cost advantages. However, the Sb-based material undergoes drastic structural alterations and severe pulverization resulting excessive capacity decay and poor cycling stability, as documented in previous work. To address the notorious volume change issue, Lin et al. [914] encapsulated Sb nanoparticles into hollow porous nitrogen-doped carbon nanotubes (Sb@N-C nanotubes). Such 3D interconnecting network allows for rapid ionic/electronic transport, preventing the aggregation of Sb nanoparticles, and such Sb@N-C nanotubes delivers excellent reversible capacity and long cycling stability as a consequent. Liu et al. [915] propose HD-Sb@Ti₃C₂T_x-G composite, in which a highly conductive elastic network and a compact double-coated structure effectively mitigates the volume expansion of Sb, provides S good electrolyte penetration and fast ion/electron transport kinetics.

  Sn alloys has drawn considerable attention in PIBs because of its high theoretical capacities. However, as one of the alloy metallic anode materials, Sn alloys undergo massive volume expansion (typically up to 300%) when interacting with potassium ions [916], causing a rapid capacity fading and a poor structure integrity in PIBs. Identically with that Bi and Sb, Feng et al. [917] proposed a feasibility of the incorporation strategy between Bi and Sn to address the above-mentioned challenges, which is substantiated in K⁺ storage. The fabricated Bi-Sn anode with a 3D porous structure forms an abundant array of active interface, which are pivotal for K⁺ interactions and facilitates for mitigating volume change and swift electron transport. Wu et al. [918] synthesized porous F-Sn/SnO_x@C layered complex as a promising candidate or PIBs. The complex exhibits a remarkable specific surface area and uniformly weaved tin nanoparticles within the layered carbon layers, contributing to its excellent potassium intercalation/decalcination performance. Li et al. [919] prepared Sn@C@Na₂CO₃ composites for high-performance K⁺ store because of their unique structure and ultrastable electrochemical properties.

  5.2.3   Transformation materials

  Transformation anode materials, similar to alloy-based anodes, demonstrated high theoretical specific capacities and underwent substantial volume variations during charging and discharging. This class included diverse types like transition metal chalcogenides, oxides, and phosphides [920,921].

  For layered transition metal chalcogenides, their large interlayer spacing facilitates the intercalation and deintercalation of K⁺ as well as conversion reactions [922,923]. However, many of these compounds suffer from poor electrical conductivity and accompany large volume changes, prompting their frequent integration with carbon materials to enhance conductivity and mitigate volume effects [924]. For instance, Zhang et al. [922] synthesized Sb₂S₃@MXene nanocomposites by acid-induced self-assembly, encapsulating Sb₂S₃ with MXene. This material retained a high capacity of 422.1 mAh/g after 100 cycles at 0.1 A/g and delivered 119.0 mAh/g even at 2 A/g. The advantages lie in the MXene as a conductive framework that enhances electron conduction in Sb₂S₃@MXene; the formation of a heterostructure with interfacial Sb-O-Ti bonds during acid induction facilitates charge transfer at the interface and effectively alleviates volume expansion of Sb₂S₃; and the uniform anchoring of nanoscale Sb₂S₃ particles on MXene nanosheets prevents aggregation and favors electrolyte infiltration. Cheng et al. [925] prepared a heterostructured binary metal sulfide Ni₃S₄/Co₉S₈ encapsulated in nitrogen-doped carbon nanocubes (NCS@NC) using a coprecipitation followed by sulfurization method. NCS@NC maintained a capacity of 417.7 mAh/g after 1000 cycles at 2 A/g and exhibited a discharge capacity of 312.2 mAh/g at 5 A/g. This demonstrates that the heterointerfaces of metal sulfides accelerate charge transfer and enhance the redox reaction kinetics of the material. Additionally, the nitrogen-doped carbon coating significantly improves electrical conductivity and greatly alleviates volume changes during K⁺ insertion/extraction, thereby imparting exceptional potassium storage performance.

  Metal oxide materials have garnered extensive attention as pivotal anode candidates for potassium-ion batteries [926,927]. Gu et al. [928] successfully synthesized a binder-free 3D self-supporting CuO@copper foam (CuO@CF) composite, where CuO nanoparticles are stably embedded within the reticulated framework of copper foam (Fig. 30a). Experimental results revealed that CuO@CF maintained high discharge capacities of 207 mAh/g after 100 cycles and 307 mAh/g after 300 cycles at current densities of 0.05 and 1 A/g, respectively, significantly outperforming pure CuO (Figs. 30b and c). This 3D conductive network design effectively mitigated the volume expansion of CuO, prevented material aggregation, and facilitated electron transport, thereby enhancing potassium storage performance. In terms of oxide modification, carbon compositing is a commonly employed method. Adekoya et al. [929] fabricated Co₃O₄@N-doped carbon (Co₃O₄@N-C) composites (Fig. 30d), which delivered a discharge specific capacity of 448.7 mAh/g after 40 cycles at a current density of 50 mA/g and retained a capacity of 213 mAh/g after 740 cycles at 500 mA/g (Figs. 30e and f). The coating of Co₃O₄ with an N-doped carbon layer not only enlarged the lattice spacing of the carbon layer to accommodate more K⁺, but also improved the conductivity, facilitated charge transfer and K⁺ diffusion, and effectively prevented material aggregation. Furthermore, Huang et al. [930] synthesized SnO₂@C composites using a metal ion and dopamine heat treatment method. This material exhibited a discharge specific capacity of 217 mAh/g after 100 cycles at a current density of 0.05 A/g and maintained a discharge capacity of 156 mAh/g after 700 cycles at 0.5 A/g, demonstrating excellent potassium storage performance (Figs. 30g and h).

  Figure 30

  

  Figure 30.  Structure and electrochemical performance characterizations of CuO@CF, Co₃O₄@N-C, and SnO₂@C. CuO@CF: (a) SEM image. (b) Cycling performance at 50 mA/g. (c) Cycling performance at 1 A/g. Co₃O₄@N-C: (d) SEM image. (e, f) Cycling performance at 0.05 and 0.5 A/g, respectively. SnO₂@C: (g, h) Cycling performance at 0.05 and 0.5 A/g, respectively. (a-c) Reproduced with permission [928]. Copyright 2022, Elsevier. (d-f) Reproduced with permission [929]. Copyright 2020, American Chemical Society (ACS). (g, h) Reproduced with permission [930]. Copyright 2022, Elsevier.

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  Metal phosphides also belong to the category of transformation anode materials [931,932]. Yang et al. [924] employed a phosphidation-nanoboxing strategy to synthesize FeP@carbon nanoboxes (FeP@CNBs) composites with a core-shell structure. FeP@CNBs maintained a discharge specific capacity of 205 mA/g after 300 cycles at 0.1 A/g. The structural design facilitated electron/ion transport through the carbon nano-boxes, enhancing conductivity and preventing aggregation of the active material. Additionally, the hollow cavity structure effectively buffered volume expansion of the material. Tan et al. [932] designed a method combining pyrolysis film blowing and phosphating treatment to prepare composites embedding FeP nanoparticles (FeP@FGCS). This three-dimensional foam-like graphitic carbon material not only improved the conductivity of the material but also inhibited aggregation of FeP nanoparticles and mitigated volume changes. Furthermore, FeP nanoparticles were firmly embedded in the FGCS matrix through strong P-C chemical bonds, enhancing the stability of the material. FeP@FGCS delivered a discharge specific capacity of 382 mAh/g at 0.1 A/g and maintained a capacity of 213 mAh/g after 800 cycles at 2 A/g, exhibiting excellent cycling performance.

  5.2.4   Other materials

  Carbon materials, alloy materials, and transformation materials are the most widely used anode materials for PIBs, with carbon materials being the most prevalent. Research on these materials accounts for over 90% of studies on PIB anodes. Beyond traditional carbon-based, alloy, and conversion materials, research into alternative anode materials is relatively limited. Current investigations primarily focus on potassium metal, anode-free designs, silicon-based materials and composites.

  Potassium metal anodes, known for their exceptional theoretical capacity, have attracted significant attention from researchers. However, the growth of potassium dendrites remains a major challenge. To address this, a self-supporting electrode composed of bismuth and nitrogen-doped reduced graphene oxide (Bi80/NrGO) has been developed [933]. This design effectively suppresses dendrite formation, thereby enhancing the battery’s stability. A symmetric battery assembled with the K@Bi80/NrGO electrode has demonstrated stable cycling for over 3000 h at a current density of 0.2 mA/cm². Furthermore, a full cell with a Prussian blue cathode and K@Bi80/NrGO anode showed remarkable stability, with no degradation after 1960 cycles at 1000 mA/g and a Coulombic efficiency of 99%. The anode-free design, which directly deposits potassium metal onto the current collector, simplifies the battery structure and reduces costs [934]. For example, anode-free potassium metal batteries with high mass-loading cathodes of Prussian blue (9.8 mg/cm²) achieved a capacity of 82 mAh/g after 200 cycles [935]. However, challenges related to potassium dendrite formation and the formation of "dead potassium" must be addressed to ensure long battery life and high capacity. Silicon-carbon anodes, successfully applied in LIBs, have also been explored for use in PIBs due to the similarities in the working principles of these two types of batteries. For example, by precisely controlling the etching temperature, researchers successfully synthesized silicon carbide-derived carbon materials (SiC-CDC) with controllable pore structures [936]. Notably, SiC-CDC demonstrated a maximum capacity of up to 284.8 mAh/g after 200 cycles at a current density of 0.1 A/g, and it maintained a high reversible capacity of 197.3 mAh/g after 1000 cycles at a current density of 1 A/g. In addition, research on anode materials for potassium-ion batteries also covers other composite materials, which integrate the advantages of different materials to further enhance the potassium storage performance. For instance, iron-based metal-organic frameworks (Fe-MOFs) serve as precursors to ingeniously disperse FeS₂ nanostructures layer-by-layer on the surface of graphene, forming a composite material (FeS₂@C-rGO) [937]. FeS₂@C-rGO exhibits a high capacity of 550 mAh/g at a current density of 0.1 A/g. Even at a high rate of 2 A/g, after 500 cycles, the capacity is still maintained at 171 mAh/g. Furthermore, bismuth (Bi)-based electrodes, renowned for their high electrochemical potential, are considered optimal for PIBs. However, the larger size of potassium ions can cause instability in the solid electrolyte interphase (SEI) film during the ongoing processes of potassiation and depotassiation, which in turn diminishes the battery’s cycle life. To address this, researchers produced the ultrathin carbon film@carbon nanorods@bismuth nanoparticle (UCF@CNs@BiN) for anodes with extended cycle life in PIBs [938]. This UCF@CNs@BiN anode stands out for its exceptional electrochemical performance, including a high specific capacity of approximately 425 mAh/g at a current density of 100 mA/g, making it a promising candidate for enhancing the performance of PIBs.

  5.3 Electrolytes

  The development of electrolytes for PIBs has been a journey of continuous exploration and innovation. In the early stages, limited by technology and understanding, research on electrolytes for PIBs was scarce, often borrowing from electrolytes for LIBs systems. The larger K⁺ radius and other traits posed challenges like slow ion migration and electrode incompatibility, affecting battery performance [939]. As research delves deeper, scientists are developing electrolytes specifically for PIBs [940,941]. On one hand, the modification of traditional organic solvents is carried out, such as selecting different carbonate ester solvents for mixing. By optimizing their ratios, the dielectric constant and viscosity of the electrolyte can be balanced, enhancing ionic conductivity [942,943]. On the other hand, the development of new solvents, such as ether [944], sulfone [945], nitrile and phosphate [946,947] organic solvents, provides a more suitable migration environment for K⁺, enhancing the charging and discharging efficiency of the batteries. Compared to ester-based electrolytes, ether solvents have higher ionic conductivity, excellent compatibility with potassium metal anodes, and can form a stable SEI on the surface (Fig. 31a) [948,949]. The research team led by Professor Bingan Lu from Hunan University has achieved significant results in the field of electrolytes for PIBs. They have not only developed a new electrolyte solvent of dipropylene glycol dimethyl ether (DMM) [944], tris(2,2,2-trifluoroethyl)phosphate (FTEP) [950], 1,2-diethoxyethane (DEE) [951], but also created a cyclic anion electrolyte based on hexafluoropropane-1,3-disulfonimide (HFDF⁻) [952]. These innovative achievements have effectively reduced the "dead potassium" phenomenon and significantly enhanced the stability of potassium metal batteries under high voltage. In the modification of electrolytes, the development of additives is an important research direction. Film-forming additives, such as FEC [953], can form a stable SEI on the electrode surface, effectively reducing side reactions of K⁺ at the electrode interface and significantly enhancing the cycle life. Recent studies have found that additives can change the structure of K⁺ solvation, which in turn determines the interfacial behavior of potassium solvents at the electrode interface [954]. The concentration of the electrolyte has a significant impact on PIBs [955,956]. High-concentration electrolytes help form a stable and dense SEI, but they can impede ion migration, resulting in reduced conductivity (Fig. 31b) [957]. Using weakly coordinating solvents to adjust the solvation structure of K⁺ can lower the solvation energy while inhibiting dendrite growth (Fig. 31c) [958–960]. For example, Guo et al. [961] used 2-methyltetrahydrofuran (MTHF) as a solvent, with its reduced potassium chelation effect, can diminish the interactions between MTHF molecules and potassium ions (Fig. 31d). Thereby enhancing the solvation capability of potassium and imparting excellent low-temperature performance to the electrolyte. Localized high-concentration electrolytes (LHCEs) adjust the concentration of the electrolyte by introducing a diluent [962,963], which ensures ion transport while reducing viscosity and improving the battery’s performance at both high and low temperatures (Fig. 31e). Qin et al. [964] first reported that LHCE can eliminate co-intercalation of solvents and achieve highly reversible K⁺ intercalation/deintercalation into graphite interlayers. There is also research exploring the use of ionic liquids as additives or cosolvents [965,966]. They have good thermal stability and electrochemical windows, which can enhance the safety and stability of the electrolyte. However, issues of cost and viscosity still need to be addressed. Additionally, research on gel polymer electrolytes (GPE) has made certain progress [967,968]. Due to their solid-state characteristics, GPE has a lower risk of combustion compared to liquid electrolytes, offering possibilities for high energy density and high safety applications in PIBs. Currently, although the research on electrolytes for PIBs has achieved significant results, it still faces challenges. Such as enhancing high-voltage stability for high-energy cathodes and improving compatibility with potassium anodes to prevent dendrites growth. Future advancements in understanding electrolyte-electrode interactions and the use of new solvent may lead to breakthroughs, advancing the commercial use of PIBs in energy storage and electric vehicles.

  Figure 31

  

  Figure 31.  (a) TEM images of graphite anodes from K/graphite cell with EC/DEC (left) and DEGDME after 20 cycles (right). Reproduced with permission [949]. Copyright 2020, Elsevier. (b) Molar concentration of different electrolytes versus ionic conductivity at 28±2 ℃. Reproduced with permission [957]. Copyright 2018, Royal Society of Chemistry. (c) Schematic diagrams of graphite electrodes with different intercalation behaviors in weakly or strongly ion-solvent interacting electrolytes and the corresponding solvated structures. Reproduced with permission [960]. Copyright 2022, Wiley-VCH. (d) Cycling performance of Gr//KPTCDA full cells at −20 ℃ and 0.1 C with different electrolytes. Reproduced with permission [961]. Copyright 2023, Wiley-VCH. (e) Linear sweep voltammetry curves of three electrolytes. Reproduced with permission [962]. Copyright 2024, Oxford University Press.

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  6. Multivalent metal-ion batteries

  6.1 Zinc-based batteries

  6.1.1   Zinc metal anode

  Currently, extensive and advanced research on high-performance zinc anodes has been carried out. Nevertheless, a series of challenges pertaining to zinc metal anodes still impede the application of aqueous zinc ion batteries (AZIBs). Briefly, as shown in Fig. 32, the zinc metal anodes in AZIBs primarily encounter issues such as the growth of zinc dendrites due to non-uniform zinc deposition, the water-induced hydrogen evolution reaction (HER), and severe metal corrosion and passivation. These problems lead to low coulombic efficiency, poor cycling stability, internal short circuits, and even cell failure.

  Figure 32

  

  Figure 32.  Challenges of Zn metal anodes encounter during electrochemical cycling.

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  The configuration and composition of the battery, as well as the charging and discharging conditions, including concentration polarization, current density, the flatness of the metal surface/edge, the zincophilicity of the substrate, operating temperature, the pH of the electrolyte, and external pressure, all induce the formation of zinc dendrites to varying extents. Zinc ions tend to deposit at regions with stronger electric field strength. Subsequent zinc ions continue to accumulate at these deposition sites, resulting in the formation of Zn dendrites. The ensuing "tip effect" further augments the number of dendrite protrusions, leading to capacity degradation or even short circuits. The loose and rough structure formed by zinc dendrites enlarges the surface area of the Zn metal, which further promotes side reactions such as corrosion and hydrogen evolution. The corrosion and passivation of the Zn metal can cause the unevenness of the zinc anode surface, exacerbating the non-uniform distribution of the electric and ion concentration fields. Additionally, the attachment of hydrogen to the Zn metal surface can inhibit the nucleation of zinc ions, leading to increased overpotential and non-uniform zinc deposition. Simultaneously, the hydrogen evolution reaction will, in turn, alter the pH near the anode and cause the accumulation of OH⁻ anions, further accelerating the generation of corrosion and passivation products. During the cycling process, the zinc metal is continuously consumed, and the formation of new corrosion products will cover the surface of the metal electrode, reducing its activity and hindering the further reduction of zinc ions, resulting in a decrease in the Coulombic efficiency. There exists an interrelated and interactive relationship among the various problems faced by the zinc metal anode.

  Inhibiting dendrite growth, the hydrogen evolution reaction, and mitigating corrosion are the keys to enhancing the performance of zinc metal anodes. Researchers have proposed a series of strategies to address the stability of zinc metal anodes in AZIBs, including the design of the metal ontology structure, zinc alloying, crystal surface modulation, the construction of lattice aptitude-inducing deposition layers, and interfacial modification methods as illustrated in Fig. 33.

  Figure 33

  

  Figure 33.  Strategies used to address the challenges of Zn metal anodes.

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  In terms of structural design, the ontological structure of the zinc metal can be engineered to homogenize the current density, increase the nucleation sites to influence the deposition behavior, and suppress the formation of dendrites [969,970]. Currently, the main structural designs encompass the utilization of zinc powder anodes, the construction of three-dimensional zinc metal, and the structural modulation of the zinc metal surface interface. Zinc powder is electrochemically active, industrially mature, and structurally adjustable, and it can be adapted to novel molding techniques to fabricate electrodes with unconventional structures. However, the high reactivity of zinc powder due to its high specific surface area can readily lead to more severe side reactions such as corrosion and hydrogen evolution, which necessitate mitigation through the modification of the zinc powder itself. Currently, the principal optimization strategies for zinc powder include: modification of the zinc powder structure, combination with other materials, the design of new suitable binders, collector modification, and electrolyte optimization.

  The construction of a Zn alloy coating layer on the Zn surface is another efficacious strategy to enhance the stability of the negative electrode structure and the durability of the battery [971,972]. The development of a homogeneous Zn-based alloy interfacial layer improves the corrosion resistance of the Zn metal, facilitates the uniform redistribution of Zn ions, regulates smooth nucleation, and enables dendrite-free Zn growth. The main optimization mechanisms of different types of Zn alloys as Zn anodes can be categorized into four aspects, namely, the improvement of corrosion resistance, the retardation of the hydrogen precipitation reaction, the reduction of by-products, and the inhibition of dendrite growth. Each type of zinc alloy possesses one or more mechanisms to enhance the performance of zinc metal anodes. For instance, ZnTi, ZnSnPb, and ZnAg alloys are optimized to improve corrosion resistance, retard the hydrogen evolution reaction, and inhibit dendrite growth.

  The unique arrangement of Zn atoms and the stacking structure in the three-dimensional direction determine the specific lattice and spatial symmetry of Zn metal. Therefore, the crystallographic characteristics of the metal itself are inherent factors influencing the orientation of the Zn deposition morphology. The selective realization of single crystal surface exposure, which induces the targeted deposition of Zn, is also one of the main directions of Zn metal anode modification. The electrochemical performance of the Zn anode is anticipated to be fundamentally enhanced by adjusting the crystal surface orientation of the Zn anode surface. Currently, the main methods for preparing Zn substrates with a single exposed crystal surface include chemical etching, mechanical treatment, heat treatment, and the design of electroplating solution components.

  A well-matched lattice fit can augment the charge density and zincophilicity, inhibit the formation of crystal defects, and simultaneously induce uniform epitaxial zinc deposition, which significantly suppresses dendrite growth. Based on the principle of lattice fitness, significant works have been reported [973,974]. Various lattice-adapted design schemes, such as heterogeneous superlattice pro-zinc ZnLi, low lattice mismatch (0.38%) vermiculite, (002) single-crystal Zn, Cu (111) crystalline-exposed Cu nanowires, and Cu₃(C₆O₆)₂ thin films as deposition-inducing layers of zinc ions, have been proposed. These schemes effectively guide the homogeneous zinc nucleation and dendrite-free zinc deposition, enabling the achievement of ultra-high rate and long-life zinc-metal batteries.

  Surface and interfacial modification of zinc metal is also one of the primary means of metal anode modification [975–977]. Well-designed solid electrolyte phases have been demonstrated to be an effective approach to modulating the deposition morphology. Constructing artificial coatings or electrolyte-induced in-situ interfacial layers on the surface of the zinc anode can prevent water or oxygen in the electrolyte from directly contacting the zinc metal and reduce corrosion or side reactions. Secondly, a multifunctional protective layer is applied to reduce the water activity of the electrolyte, inhibit the active water molecules from reaching the zinc metal, and improve the stability and cycle life of the zinc anode.

  In conclusion, the research progress on zinc-metal anodes indicates that the stability and battery performance of zinc anodes can be remarkably enhanced through strategies such as structural design, zinc alloying, the construction of lattice aptitude layers, single crystal surface exposure, and surface and interfacial modification. These studies offer valuable guidance for the development of reversible zinc-metal anodes and, simultaneously, provide an important scientific foundation and technical approach for the practical application of aqueous zinc-ion batteries.

  6.1.2   Zinc ion battery cathode

  Mn-based cathode: Mn-based oxides are one of the most promising cathode materials for aqueous Zn ion batteries (AZIBs) due to their low cost, environmental friendliness and abundant crystal structure [978]. Among numerous manganese oxides (MnO₂, MnO, Mn₂O₃, Mn₃O₄), MnO₂ is the most widely studied. MnO₂ is composed of [MnO₆] octahedrons which are connected by sharing edge or/and angle forming different crystalline phases (α-MnO₂, β-MnO₂, γ-MnO₂ and so on) [979]. The energy storage mechanisms of MnO₂ are complex and controversial. Based on reported literatures, there are four typical mechanisms: (1) Zn²⁺ insertion/extraction; (2) H⁺/Zn²⁺ co-insertion/extraction; (3) conversion reaction mechanism; (4) dissolution/deposition mechanism [980]. Ions insertion/extraction and dissolution/deposition mechanisms are mainly energy storage mechanism for MnO₂, while relatively few studies have been reported on conversion reaction mechanism.

  Ions insertion/extraction mechanism undergoes single electron transfer reaction (Mn⁴⁺/Mn³⁺) and MnO₂ with this mechanism has relatively low theoretical capacity (308 mAh/g) and voltage (~1.3 V vs. Zn²⁺/Zn) (Fig. 34a). During ion insertion/extraction process, complex multiphase transformation is commonly observed, such as the transformation of tunnel phase (β-MnO₂) into layer Zn-buserite, resulting in structure collapse of material and capacity degradation [981]. The ubiquitous J-T distortion caused by Mn³⁺-induced disproportionation reaction (2Mn³⁺ → Mn²⁺ + Mn⁴⁺) would lead to the dissolution of manganese, which is also a non-negligible issue on deteriorating performance. Moreover, Zn²⁺ have strong electrostatic interaction with host structure of MnO₂, which increases Zn²⁺ diffusion barrier. Furthermore, inferior electrical conductivity of MnO₂ limits electrons transfer rate. Above all, the application of MnO₂ cathode is still hampered by the discontented rate performance and capacity decay resulting from the sluggish kinetics and unstable structure. Recently, many strategies have emerged to improve electrochemical performance of MnO₂, such as defect engineering, doping engineering, constructing composite/hybrid structures and electrolyte optimization. Although some achievements have been made, the deeper information about phase transition and H⁺ insertion/extraction mechanism in MnO₂ with different crystalline phases need to be further explored. Dissolution/deposition mechanism is based on the reversible dissolution/deposition behavior of MnO₂ (MnO₂/Mn²⁺), involving two electrons transfer reaction. Therefore, MnO₂ with two electrons transfer reaction could offer high theoretical capacity and voltage (616 mAh/g and 1.99 V vs. Zn²⁺/Zn) [982]. This mechanism is related to pH value of electrolyte and requires acidic electrolyte to achieve adequate dissolution/deposition reaction. However, Zn corrosion and hydrogen evolution reaction (HER) would be aggravated in acidic electrolyte, deteriorating electrochemical performance. Therefore, many reported optimization strategies are mainly based on isolating Zn anode from electrolyte or promoting dissolution/deposition of MnO₂. For example, developing decoupled system and flow system battery or doping metal elements to catalyze dissolution/deposition of MnO₂ [983,984]. Currently, the complex reaction path of dissolution/deposition mechanism and none 100% reversible dissolution-deposition of MnO₂ are still focused. In addition, several researches have reported that both ions insertion/extraction and dissolution/deposition of MnO₂ mechanisms simultaneously exist in Zn//MnO₂ battery, which makes the energy storage mechanism of MnO₂ greater complex.

  Figure 34

  

  Figure 34.  (a) An overview of electrochemical properties for MnO₂ with single electron transfer reaction and two electrons transfer reaction. (b) An overview of electrochemical properties for vanadium oxides and vanadium phosphates.

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  Vanadium-based cathode: Vanadium oxides possess several variants with various oxidation states of vanadium, mainly including V₂O₅, VO₂, V₆O₁₃ and V₂O_3. Among these, V₂O₅ could deliver high theoretical capacity owing to its open diffusion channels between layers and a two-electron oxidation–reduction reaction between V⁵⁺ and V³⁺. Therefore, V₂O₅ and its derivatives are dominant in the research of cathode material for AZIBs. The common energy storage mechanisms for V₂O₅-based cathodes include Zn²⁺ insertion/extraction and H⁺/Zn²⁺ coinsertion/extraction [985]. During Zn²⁺ migration process, Zn²⁺ performs sluggish transport kinetics due to its high desolvation penalty at electrode-electrolyte interface and strong electrostatic in the host structure of materials. Meanwhile, lattice spacing undergoes periodic expansion/shrinkage during repeated Zn²⁺ insertion/extraction which inevitably causes volume change and structural pulverization of cathode. Many researches devote great effort to construct stable crystal structure and electrode-electrolyte interface to facilitate Zn²⁺ transport by guest pre-intercalation, surface modification, electrolyte optimization and so on. H⁺ with small ionic radius and special Grotthuss shuttle mechanism has good transport kinetics. Therefore, most researches tend to promote H⁺ insertion. However, proton insertion consumes electrolyte to form layered double-hydroxide salts (LDHs) attaching to electrode surface by weak electrostatic interaction, which would hinder H⁺/Zn²⁺ across electrode-electrolyte interface. Moreover, LDHs may fall off upon prolonged resting which would cause irreversible behavior and cell failure [986]. At present, relevant researches on electrochemical behavior of H⁺ is still lacking, such as the relationship of H⁺ insertion and limited lifetime of cathode, the subtle relation between H⁺ and electrolyte.

  Vanadium phosphates materials as important derivatives of V₂O₅ are composed of corner-sharing [VO₆] octahedra linking to [PO₄] tetrahedra. Monoclinic-type structure (Li₃V₂(PO₄)₃), NASICON (Na₃V₂(PO₄)₃) and layered-type structure (VOPO₄) are typical structure of vanadium phosphates cathodes applying to AZIBs [987]. The introduction of polyanionic group leads to a higher voltage making up for the shortage of vanadium oxides (Fig. 34b). However, [VO₆] octahedrons are separated by phosphate causing that vanadium phosphates exhibit poor electronic conductivity. Due to limitation of electronic conductivity, the practical capacity and rate capability cannot be unsatisfied. There are three typical energy storage mechanisms for vanadium phosphates cathodes: (1) Zn²⁺ insertion/extraction; (2) H⁺/Zn²⁺ coinsertion/extraction; (3) bimetal ion co-insertion/extraction (Zn²⁺ and Na⁺, Li⁺, Mg²⁺) [987]. The Na₃V₂(PO₄)₃ and Li₃V₂(PO₄)₃ are mainly bimetal ion co-insertion/extraction. Their stable 3D ion transport channel can offer fast pathway for Zn²⁺ diffusion. However, the limited capacity and poor cycle performance restrict their application. Layer VOPO₄ is attractive cathode, showing Zn²⁺ or H⁺/Zn²⁺ co-insertion/extraction. During ions insertion/extraction process, redox platform of oxygen could be activated, thus VOPO₄ possesses high voltage and capacity [988]. However, after repeated charge/discharge, the PO₄^3- would extract from bulk structure, leading to disappearance of high-voltage platform and transferring into redox platform of V₂O₅. Although the lifetime and voltage of vanadium phosphates could be improved effectively by making tremendous efforts on electrolyte optimization, structure design, interfacial engineering and transition-metal substitution, the design and applications of vanadium phosphates cathodes are still urgently improved.

  6.2 Magnesium-based batteries

  Rechargeable magnesium batteries (RMBs) are a promising candidate in post-LIBs era, due to the abundant resource (~28,104 ppm vs. lithium’s 16 ppm) [989], low potential (−2.37 V vs. SHE) [990], high volumetric capacities (3832 mAh/cm³ vs. lithium’s 2061 mAh/cm³) [991], and dendrite-free deposition of magnesium metal [992].

  6.2.1   Cathode materials

  Cathode materials fall into three categories: intercalation and conversion-type inorganic materials, and organic electrode materials with unique redox mechanisms [993]. Among them, intercalation inorganic materials are the most widely used in RMBs due to their variety, stable structure and reversible intercalation/de-intercalation of Mg²⁺. These include Chevrel phase, layered and spinel-structured materials, and polyanionic compounds. Mo₆S₈ of the Chevrel phase is one of the most successful cathode materials [994]. Improvements in synthesis methods [995], smaller particle size [996], replacement of anions (e.g., S to Se) [997], and higher test temperature [998] can have been extensively investigated. Notably, the Chevrel phase Mg₂Mo₆S₈ was used for the first time in the anode-free, boosting the volumetric energy density [999]. Layered materials offer fast Mg²⁺ diffusion and reversible embedding/de-embedding due to their two-dimensional structure, with increased layer spacing reducing the ion diffusion barrier [1000]. Spinel-type materials provide more stable crystal structures, favoring cycling stability [1001]. Polyanionic compounds, in spite of offering high structural stability and operating voltage, suffer from lower conductivity and theoretical capacity density [1002,1003]. Currently reported intercalation materials are mainly based on cationic redox chemistry for Mg²⁺ storage. To further increase the capacity, cathode materials with anionic redox chemistry were recently focused [1004,1005]. Conversion-type materials, which involve chemical bond breaking and structural rearrangement during Mg²⁺ storage, tend to have higher theoretical energy densities. Mg-based oxides have open channel structures that provide larger ion transport channels but suffer from structural collapse due to the conversion of Mg²⁺ [1006]. Sulfide and selenide materials exhibit more stable cycling with lower capacities [1007]. Organic cathode materials have attracted widespread attention due to their structural designability and skeleton flexibility. Their active functional groups, such as carbonyl [1008], C=N double bonds [1009], and radicals [1010], allow for compatibility with various RMB electrolytes and efficient Mg²⁺ and MgCl⁺ ion intercalation, offering promising alternatives to traditional inorganic materials.

  The advances of RMBs are restrained by the strong interaction between Mg²⁺ and the host material, resulting in slow Mg²⁺ diffusion and poor electrochemical properties [1011]. Strategies to enhance Mg²⁺ diffusion include reducing particle size, modifying host anions, and employing molecular or ionic pre-intercalation techniques [1012]. Conversion-type materials face structural instability and low ion mobility, which can be improved by incorporating highly conductive carbon materials and optimizing particle size [1013]. Organic materials, especially small molecules, suffer from high solubility and poor conductivity, causing rapid capacity decay. These issues can be addressed by polymerizing organic molecules, compositing with conductive carbon, and matching functional groups with electrolyte solvents.

  6.2.2   Electrolytes

  The low redox potential of Mg metal leads to high reaction activity with common organic solvents, salts, and impurities in electrolyte, forming an ion- and electron-insulating passivation layer that impedes Mg²⁺ deposition/striping [994,1014]. Addressing this issue in non-aqueous electrolytes has been the research spotlight, which is the main focus of this review [1014,1015]. Mg deposition was first demonstrated with Grignard reagent in early 1920s [1016] and boron-based anionic Mg salt electrolyte by 1990 [1017]. In 2000, to enhance the oxidation resistance of Grignard’s reagent, Aurbach et al. proposed a Mg organohaloaluminate salts electrolyte as the "first-generation" RMBs electrolyte, which enhanced oxidative stability to 2.64 V vs. Mg²⁺/Mg (vs. 1.95 V for Grignard’s reagent) and achieved nearly 100% Mg deposition/stripping efficiency [1018]. Replacing alkyl groups with phenyl groups further increased the oxidative stability to 3 V vs. Mg²⁺/Mg and improved ionic conductivity to ~4–5 S/cm, due to the strong binding energy between phenyl structures and aluminum nuclei [1019,1020]. Chlorine in these electrolytes facilitates the removal of passivation layers but poses challenges such as corrosion of cell components and the formation of Mg-Cl⁺ or Mg₂Cl₃⁺ [1021,1022]. Chlorine-free electrolytes, such as Mg(BH₄)₂, can support reversible Mg²⁺ deposition/stripping. In contrast, most inorganic Mg salts, including BF₄⁻, ClO₄⁻, SO₃CF₃⁻, and PF₆⁻, are ineffective [1023,1024].

  It is worth noting that the solvents of these electrolytes are mainly ether molecules, due to their low viscosity conducive to ionic conductivity, the coordination of ether oxygen to stabilize Mg ions, and compatibility with Mg metal anodes. Among these, THF is the most widely used solvent [1025,1026]. However, it is low boiling point (~66 ℃) brings evaporation issues, leading to the more research interest in long-chain ethers, such as DME (~85 ℃), G2 (~162 ℃), G3 (~216 ℃) and G4 (~275 ℃) [1015]. The high polarity of the ether-oxygen group in the chain ether has strong interaction with Mg²⁺, enhancing coordination, solubility for Mg salts, and antioxidant property by lowering the highest occupied molecular orbital energy level [1027,1028]. Amine-modified ethers further improve the affinity to Mg²⁺ by 6-to-41-fold times and decrease the de-solvation energy barrier [990]. Despite these advantages, ether electrolytes face challenges in pairing with high-voltage cathodes [1029,1030]. This underscores the need to explore alternative solvents, including aqueous or non-aqueous systems, more complex co-solvent/co-salt combinations [1031], additives [1032], and even solid electrolytes [1033,1034]. Such innovations could deepen understanding of RMBs and accelerate their commercialization efforts.

  6.2.3   Anode materials

  In addition to developing electrolytes compatible with Mg metal anodes, alternative "beyond-Mg metal" anodes have also explored, including Bi- [1035,1036], Sn- [1037], and Ti-based anodes [1038]. These alternatives can insert Mg²⁺ at higher potentials than Mg²⁺/Mg, reduce reactivity with simple electrolytes, and avoid passivation layer formation [1039]. Specifically, Bi-based anode, one of the most studied candidates after Mg metal anodes, offers a high theoretical capacity (~385 mAh/g) as each Bi atom binds up to three Mg atoms to form an alloy with alloy potential of 0.23 V vs. Mg²⁺/Mg [1040]. Bi and its alloys show good cycling stability and compatibility with simple electrolytes [1040]. Furthermore, the morphology and size of the anode also affects the diffusion and deposition of Mg. For example, micrometer-sized Bi anode can release values close to the theoretical capacity (~300 mAh/g) at 2 C [1041], while Bi nanotubes have smaller deposition overpotentials than micrometer particles [1042]. Sn-based anode is another promising alternative because of a higher theoretical specific capacity (903 mAh/g), and a lower alloy potential (0.15 V vs. Mg²⁺/Mg) than Bi. They also show lower hysteresis of the alloy and de-alloy process (50 and 90 mV in Sn and Bi anode, respectively) [1037]. However, the instability of pure Sn anodes leads to poor cycling performance [1043]. The SnSb alloy outperforms both pure Sn and Sb in terms of rate performance and cycling stability [1044]. Other metals like In and Pb have been tested as anodes but suffer from poor Coulombic efficiency due to volume changes during cycling [1045,1046]. Ti-Based anodes are gaining attentions for their compatibility with simple electrolytes and nearly unchanged volume during alloying [1047]. To summarize, the development of anodes for RMBs should focus on ensuring high reversible Mg²⁺ deposition/stripping and compatibility with simple electrolytes.

  6.3 Calcium-based batteries

  Electrochemical energy storage and conversion systems are the key technologies for constructing a renewable clean energy structure. The wide application of LIBs in the fields of portable electronic devices and electric vehicles has also proved the feasibility of this solution. However, the relatively low energy density and high cost of traditional LIBs have hindered their development on a larger scale [1048]. To solve this problem, the development of new battery systems with high energy density and abundant resource reserves has become a research hotspot [1049]. Among them, calcium metal batteries use calcium metal, which ranks third in the reserves of metal elements in the earth’s crust, as the negative electrode. Moreover, calcium metal has a relatively low reduction potential (−2.87 V vs. SHE) and a high theoretical specific capacity (~1337 mAh/g, ~2072 mAh/cm³) [1050]. Compared with other multivalent ion systems, the relatively small charge density of calcium ions makes them more suitable for high-energy and high-rate application scenarios [1050,1051]. Although calcium metal batteries exhibit obvious advantages, due to the relatively high reactivity of calcium metal and the relatively high charge density of Ca²⁺, calcium metal is easily passivated in conventional electrolytes, forming a passivation layer or solid electrolyte interphase (SEI) on the surface of calcium metal. As a result, it is very difficult or even impossible for Ca²⁺ with a relatively large charge density to be effectively transported, and thus the effective electrochemical deposition/stripping process of calcium metal cannot be achieved [1052]. Similarly, the relatively large charge density of Ca²⁺ is also a major challenge for the development of calcium storage electrode materials. Therefore, the development of efficient electrolytes, the construction of stable interfaces for the rapid transport of calcium ions, and the development of stable and reversible calcium storage electrode materials are the key focuses of the research on calcium metal batteries.

  6.4 Aluminum-based batteries

  Aluminum (Al) is the most abundant metal in the earth’s crust (8.2 wt%) and the lowest cost raw material ($1.8 kg⁻¹) [1053]. Al metal can offer an ultrahigh theoretical volumetric capacity of 8046 mAh/cm³, nearly four times of lithium (2060 Ah/cm³). The theoretical gravimetric capacity of Al is 2980 mAh/g, which is much higher than that of Mg (2200 mAh/g), Zn (820 mAh/g) or Ca (1341 mAh/g) [1054]. Capitalizing on its intrinsic safety, light weight, ease of handling, and eco-friendliness, the direct application of aluminum metal as the anode strongly drives the development of aluminum-based batteries. Aluminum as an anode was first used in the Al/HNO₃/C Buff cell in 1857 [1055]. Early works on aluminum anode were mainly limited to primary batteries, such as Al-O₂ [1056,1057], Al-MnO₂ [1058], and Al-H₂O₂ [1059], due to the irreversible plating process of Al³⁺ in aqueous electrolytes. Until 1985, Auborn et al. [1060] proposed chloroaluminate ionic liquid electrolyte to break the passivation layer and achieve reversible plating/stripping of Al metal at room temperature. The turning point of room-temperature RABs came in 2015, when both Dai and Jiao’s group proposed a new type of aluminum battery with a graphite cathode and an Al anode, which exhibited a high-voltage (close to 2 V) and stable cycling performances (over 7500 cycles) [1061,1062]. Subsequently, this battery system has inspired numerous studies investigating rechargeable Al-ion batteries (RABs) [1063–1065], aluminum-sulfur batteries (ASBs) [1066–1068], and Al-graphite dual-ion batteries (DIBs) [1069,1070]. Extensive research has been conducted on Al metal anode, cathode materials, and electrolytes.

  6.4.1   Anode materials

  To date, most RABs have been assembled with Al plate as an anode. However, achieving stable and efficient plating/stripping of Al³⁺ at the Al anode is challenging due to the spontaneous formation of a passivating oxide layer on the Al surfaces, growth of dendrites, continuous chemical corrosion side reactions, and volume evolution [1054,1071]. The electrochemical performance of Al anode has been summarized to be significantly influenced by the purity, grain size, surface composition, micro morphology, and the internal crystalline phase of the aluminum plate [1072,1073]. To enhance the Coulombic efficiency, research has focused on optimizing the solid-electrolyte interphase (SEI) layer by constructing the thiophdiyne and Bi-based artificial interphase layer on Al plate [1074,1075]. A series of Al-based alloy anodes have also been employed to replace pure Al plate to decrease the self-corrosion rate, such as Al-Sn, Al-Ga, Al-Cu [1076,1077]. Advanced electrode designs, such as 3D structured aluminum foams or nanostructured surfaces, have been proposed to mitigate dendrite growth by increasing the surface area and improving the diffusion of ions [1078].

  6.4.2   Cathode materials

  The cathode material is a critical component that directly influences the energy density, power output, and cycle life of RABs. However, most cathode materials that can reversibly intercalate Li⁺ hardly show appropriate performance in RABs. The main reason is that the trivalent Al³⁺ suffers a high charge density due to the small radius (0.535 Å), leading to strong electrostatic interactions with both the lattice anions in cathodes and the solvents or anions in the electrolyte [1079]. The charge storage mechanisms of aluminum ions in various cathodes based on chloroaluminate IL electrolytes can be basically classified into three categories. (1) Al³⁺ intercalation and conversion. Numerous cathode materials, including transition metal oxides (V₂O₅, SnO₂, etc.) [1080,1081], transition metal chalcogenides (TiS₂, a-TiS₄, Mo₆S₈, CoSe₂ etc.) [1082–1084], MXenes (V2CTx) [1085], and Prussian blue materials (K_0.2Fe[Fe(CN)₆]_0.79·2.1H₂O) [1086] have been reported to enable reversibly intercalate Al³⁺. Besides, the typical conversion S and Se cathodes that exhibit high energy density of in Al-based batteries [1087,1088]. (2) AlCl₄⁻ intercalation. The widely reported graphite-based cathode materials with different morphology have been demonstrated to storage AlCl₄⁻ [1061,1089]. (3) AlCl_y^(3-y)+ (y = 1, 2) and X⁺ (solvent molecules) intercalation. The organic materials with macrocyclic structures, carbonyl or amine functional groups are the representative cathode under AlCl_y^(3-y)+ and X⁺ intercalation (PQ, TPBQ) [1090].

  6.4.3   Electrolyte

  The pursuit of the optimal electrolyte for Al-based batteries is an ongoing investigation. Current Al-ion electrolytes can be categorized into three main groups: aqueous electrolytes, molten salts electrolyte, and room-temperature organic electrolyte. Aqueous electrolytes exhibit high ionic conductivity, high safety, low cost, and environmentally friendly merits [1091]. However, the standard electrode potential of Al³⁺/Al (−1.68 V) is beyond the stable electrochemical window of water (1.23 V), whereas hydrogen evolution reaction become the main challenge [1092]. Recently, the "water-in-salt" Al-ion electrolytes have been proposed as a method to shift the hydrogen evolution potential negatively and suppress hydrogen evolution by improving the concentration of Al(NO₃)₃, Al(ClO₄)₃, Al(OTF)₃ and AlCl₃ solutions [1086]. The molten salt electrolyte containing AlCl₃ and alkali metal chloride salt has also been investigated in RABs under a temperature range of 70–300 ℃, showing favorable characteristics of inherent low cost and low flammability [1093,1094]. Up to now, the most widely investigated electrolyte system for RABs is the chloroaluminate and imidazole halides ionic liquid (IL) electrolyte, which exhibit an ionic conductivity of (10–25 mS/cm) at room temperature [1095]. Additionally, other ionic liquid electrolytes such as AlCl₃/Et₃NHCl, AlCl₃/urea, and AlCl₃/acetamide were developed to reduce costs of imidazolium halide [1096–1098]. Meanwhile, organic and chlorine-free electrolytes of Al(TFSI)₃/AN, Al(OTF)₃/G2 are under development to mitigate Cl⁻ corrosion [1099,1100].

  7. Dual-ion batteries

  Dual-ion batteries (DIBs), as a new type of energy storage technology in which anions and cations participate in the electrochemical reaction together and the positive electrode mainly relies on anionic intercalated graphite materials, have attracted much attention in recent years because of their cost-effectiveness, high operating voltage, low cost, and environmental friendliness [1101–1103]. The development of DIBs originated from research on graphite intercalation compounds, and the development of DIBs was based on the study of graphite intercalation compounds. The development of DIBs originated from the study of graphite intercalation compounds. In 1938, Rüdorff and Hofmann reported for the first time the reversible de-embedding of HSO₄⁻ anionic ligands into graphite electrode materials, a process similar to that of lithium-ion "rocking chair" batteries. In 1989, McCullough et al. proposed a "dual-carbon battery" based on the dual-embedding mechanism and applied for the first patent on DIBs, which became a milestone in the development of DIBs. In 2000, Seel and Dahn [1104] investigated the reaction mechanism of the PF₆⁻ de-embedded graphite electrode material using in situ X-ray diffraction. This work laid the foundation for the study of the energy storage mechanism of the anionic intercalation reaction. Since then, a lot of research has been conducted on the improvement of the performance of DIBs by different anionic ligands and different solvents [1105]. In 2016, Tang et al. [1106] reported a new type of aluminum-graphite DIBs. The battery used graphite as the anode material, and the aluminum foil served as both the anode material and the current collector. During the charging process, an anionic intercalation reaction occurs at the graphite cathode and an aluminum-lithium alloying reaction occurs at the aluminum cathode, which significantly improves the energy density of the battery and dramatically reduces the manufacturing cost of the battery. In addition, various new types of DIBs, such as aqueous and all-solid-state, have been reported successively [1107].

  8. Aqueous rechargeable batteries

  8.1 Aqueous alkali metal-ion batteries

  8.1.1   Aqueous lithium-ion batteries

  For making intrinsically safe LIBs, a reasonable proposal is to replace the commercial organic electrolytes with aqueous electrolytes and match them with cathode and anode materials similar to the commercial LIBs, which is the aqueous lithium-ion batteries (ALIBs) [1108,1109]. In 1994, the ALIB was first reported [1110]. The ALIBs used LiNO₃ as the main salt with a concentration of 5 mol/L, and the pH of the electrolyte was adjusted with 0.001 mol/L LiOH to the pH of 11, which made hydrogen evolution reaction (HER) hard to occur. The cathode was LiMn₂O₄, the anode was VO₂, and the output voltage of the battery was 1.5 V.

  In recent years, great efforts were made to improve the performance of ALIBs [1111]. Wang et al. [1112] reported that the output voltage of the full cell was 1–1.1 V with LiMn₂O₄ as the cathode, Li₂Mn₄O₉ or Li₄Mn₅O₁₂ as the anode, and 6 mol/L LiNO₃ + 0.0015 mol/L OH⁻ as the electrolyte. Kohler et al. [1113] reported that the output voltage of the full cell was 1–1.2 V with LiV₃O₈ as anode, LiNi_0.81Co_0.19O₂ as the cathode, and 1 mol/L Li₂SO₄ or LiCl as the electrolyte. The full battery capacity was 45 mAh/g, and the capacity retention rate was only 70% after 30 cycles. In 2007, reported several works on ALIBs from Institute of Physics, Chinese Academy of Sciences were reported, including the TiP₂O₇ anode and LiTi₂(PO₄)₃ anode, the two anodes were respectively matched with LiMn₂O₄ cathode, and 5 mol/L LiNO₃ electrolyte. The output voltage and capacity of the two full cells were 1.4 V, 42 mAh/g and 1.5 V, 45 mAh/g, respectively. The second work used polypyrrole (PPy) coated LixV₂O₅ as the anode, LiMn₂O₄ as the cathode, and 5 mol/L LiNO₃ as the electrolyte. The output voltage was about 1 V. The third work used polyaniline (Pan) coated LixV₂O₅ as the anode, 5 mol/L LiNO₃ as the electrolyte, and LiNi_1/3Mn_1/3Co_1/3O₂ as the cathode. The charge and discharge curves are shown in. Xia et al. [1114] found that removing oxygen from the electrolyte can significantly improve the performance of ALIBs (1000-cycle capacity retention rate greater than 90%). In summary, early ALIBs had the disadvantage of low output voltage (≤1.5 V), which greatly limited the output energy density of ALIBs and hindered the practical application of ALIBs.

  One of the challenges of ALIBs is the anode interface. Compared with traditional commercial organic electrolytes, the electrolyte and lithium salt can be reduced within their operating voltage range, forming a dense solid electrolyte interphase (SEI) at the anode interface (Fig. 35a), thereby providing protection for the anode and preventing the continuous side reactions [1115]. In traditional aqueous electrolytes, due to the low thermodynamic decomposition potential of water (the stable electrochemical window of water is only 1.23 V) [1110], the decomposition potential of water can be higher than the reduction potential of the lithium salt, so it is impossible to form a dense SEI through the reduction of the salt. Meanwhile, water can continue to decompose, and the product is H₂, which cannot provide support for the formation of SEI, but can greatly affect the performance of the battery (Fig. 35b).

  Figure 35

  

  Figure 35.  (a) The anode interface of the conventional commercial Li-ion battery. (b) Anode interface of ALIBs. (c) Ion pairs in WIS electrolyte. (d) Schematic of WIS aqueous lithium-ion battery. (e) Important time point in the development of ALIBs. (c, d) Reproduced with permission [1116]. Copyright 2015, American Association for the Advancement of Science.

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  In 2015, Wang proposed the concept of aqueous electrolyte Water-in-Salt (WIS) [1116], which successfully improved the electrochemical window of aqueous electrolyte to >3 V and has been widely used in fields other than batteries [1117]. The electrolyte is the 21 mol/kg lithium bis(trifluoromethanesulfonyl)imide (LiTFSI) aqueous solution (LiTFSI has high water solubility) [1118]. With such a high concentration of electrolytes, there can be unique solvent-ion and ion-ion interactions in the electrolyte (Fig. 35c): On the one hand, through the solvation of lithium ions and water, the activity of water is thermodynamically reduced, slowing down the hydrogen evolution of water. On the other hand, due to the strong interaction between anions and cations, anions can be preferentially reduced to form SEI before water decomposition. The interface of the anode is protected because of the SEI (Fig. 35d). It can be concluded that WIS can attract widespread attention, the essential reason is that the WIS can solve the challenge of the anode interface in ALIBs.

  After the WIS electrolyte was proposed, various derivative concepts were generated [1119–1122], and the mechanism of widening the window of WIS electrolytes was also deeply explored [1123–1125]. These results have driven the development of ALIBs towards higher output voltages [1126,1127]. Along with the introduction of the WIS concept, Wang and other authors proposed an aqueous battery with LiMn₂O₄ as cathode, WIS as the electrolyte, and Mo₆S₈ as anode, with an output voltage of about 1.8 V. In 2016, Wang et al. [1128] proposed the concept of water-in-bisalt (WIBS, 21 mol/L LiTFSI + 7 mol/L LiOTF), further increasing the concentration of lithium salt and broadening the electrochemical window of the aqueous electrolyte. The anatase TiO₂ anode with a lower lithiation potential can be utilized in ALIBs. With the LiMn₂O₄ cathode, the output voltage of the full cell is 2.1 V. In the same year, Atsuo Yamada et al. [1129] proposed the concept of hydrated-melt electrolytes. The molecular formula of the electrolyte is Li(TFSI)_0.7(BETI)_0.3·2H₂O. In the full cell, LiNi_0.5Mn_1.5O₄ is cathode and Li₄Ti₅O₁₂ is used as the anode. The output voltage of this battery system is close to 3 V. In 2017, Wang et al. [1130] proposed to gel the WIS electrolyte (gel polymer form, GPE) to form WIS GPE, and used LiVPO₄F as the cathode and anode of the full cell, making the output voltage of the full battery 2.4 V. In 2019, a cathode material (LiBr)_0.5(LiCl)_0.5-graphite (LBC-G) was reported [1131] that can undergo conversion-intercalation reactions. The average potential of the cathode is 4.2 V and the capacity is 243 mAh/g. This high capacity comes from the first-order graphite intercalation compound C_3.5[Br_0.5Cl_0.5] that can be reversibly formed in the WIBS electrolyte. Meanwhile, the authors assembled an ALIB using WIBS as a gel electrolyte, LBC-G as cathode, and graphite protected by a fluorinated polymer gel as anode. The average voltage of the full battery at 0.2 C is 4.1 V, and a stable discharge capacity of 127 mAh/g can be obtained. After 150 cycles, the capacity retention rate is 74%, and the average coulombic efficiency is 99.8%. The important time points in the development of ALIBs are shown in Fig. 35e.

  In the future, the development direction of ALIBs will be towards high energy density, high output voltage, and long cycle life (Fig. 35e). Among them, high output voltage is conducive to improving the energy density of ALIBs, which requires that cathodes with higher output voltage can be utilized in ALIBs, such as Li_0.5Mn_1.5O₄. It is required that lower lithiation potential should be used in ALIBs, such as Li₄Ti₅O₁₂, graphite, silicon. Long cycle life is a more challenging target in ALIBs, which requires a very stable anode interface environment to reduce the decomposition of water. The key technology may lie in the construction of artificial SEI. Meanwhile, high-capacity cathode and oxidative sacrificial compensators added to the electrolyte can prevent the loss of electrons due to the HER, which is also beneficial to the long cycle life of ALIBs.

  8.1.2   Aqueous sodium-ion batteries

  Aqueous sodium-ion batteries are based on aqueous electrolytes, which are in metric of low cost, high safety, and high sustainability, however, restricted by the inherent properties of narrow electrochemical window and high freezing point. Thus, various strategies have been proposed to solve the above issues, including highly concentrated electrolytes [1132,1133] and organic additives [1134,1135], to generate robust SEI layer and reduce the activity of water for widening the electrochemical window, and to module the hydrogen bond linked water structure [1136,1137] for lowering the freezing point of aqueous electrolytes.

  Since 2017, Suo et al. [1138] first reported the water-in-salt electrolyte (WISE) of 9.26 mol/kg NaOTF, which can form NaF-based SEI and stabilize the anode, extending the electrochemical window to 2.5 V. The full battery coupled with Na_0.66[Mn_0.66Ti_0.34]O₂ as cathode and NaTi₂(PO₄)₃ as anode shows high Coulombic efficiency (>99.2%) at a 0.2 C. The electrochemical window of NaClO₄ electrolyte was expanded from 1.2 V to 2.8 V with high concentration of 17 mol/kg, which enables the 2.6 V operation of Na₂Mn[Fe(CN)₆]||NaTi₂(PO₄)₃ full battery [1139]. Bi-salt electrolyte can further widen the windows and improve stability. Jin et al. [1132] designed 19 mol/kg electrolyte composed of NaClO₄ and NaOTF, stabilizing the symmetric Na₃V₂(PO₄)₃ battery for 100 cycles. The introduction of 22 mol/kg tetraethylammonium triflate widens the electrochemical window of 9 mol/kg NaOTF electrolyte from 2.7 V to 3.3 V, supporting Na_1.88Mn[Fe(CN)₆]_0.97·1.35H₂O||NaTiOPO₄ full battery with high energy density of 71 Wh/kg [1140]. Liu et al. [1134] reported that the ternary eutectic phase of SN−NaClO₄−H₂O features a water-locked bipolar environment, which breaks the hydrogen-bond network and intensify the O−H bond, thus delivering 3.4 V electrochemical window. Polymers can also stabilize water molecules by confining them in the ionogel network through hydrogen bond, show high operational cell voltage of 3.0 V [1135]. The high concentration and organic additives are effective way for widening the electrochemical windows of aqueous electrolyte above 3 V.

  To improve low-temperature performance degradation, lowering the freezing point of electrolytes is the principal. The freezing point is highly related to the hydrogen bond network of water [1141]. The more solute, the more disruption occurs to hydrogen bond. On the one hand, the concentration of solute affects the hydrogen bond structure. The freezing point generally decreases with concentration at the initial stage. The highly concentrated electrolyte can serve as low-temperature electrolyte as well [1133]. On the other hand, the intrinsic properties of the solute also play a role in the freezing process. When introducing solutes with strong interaction ability related to water molecules, the freezing point of electrolytes drops sharply. For example, the addition of Al(OTF)₃, Ca(OTF)₂, and ethylene glycol into NaOTF electrolyte can reduce the low freezing point to −86.1, −72.6, and −114.5 ℃, due to their strong interaction with water molecules, which destroy the original hydrogen bond network [1136]. Mg(ClO₄)₂ [1142] and dimethyl sulfoxide [1137] can also serve as antifreeze additives, which enables the batteries operation under the ultralow temperature of −60 ℃.

  8.1.3   Aqueous potassium-ion batteries

  Compared to traditional organic battery systems, aqueous alkali-metal-ion (Li⁺, Na⁺, and K⁺) batteries (AAIBs) offer several advantages: non-flammable, cost-effective, and fast ion-diffusion kinetics [1143]. In the AAIB systems, potassium (K) is more abundant in the earth’s crust compared to lithium (Li), and it has a lower redox potential and smaller hydrated ion radius than sodium (Na). Therefore, K is particularly advantageous for large-scale energy storage systems [1144]. However, up until now, the development of aqueous K-ion batteries (AKIBs) still faces numerous challenges: (1) Traditional aqueous electrolytes are prone to oxygen and hydrogen evolution reactions during charging/discharging processes; (2) There are limited electrode materials available for K systems, and few K-based electrode materials are suitable for aqueous systems [1145]; (3) The high dielectric constant and polarity of water cause many electrode materials to dissolve; (4) Aqueous electrolytes can lead to battery failure at low temperatures, especially below the freezing point of water [1146,1147].

  In recent years, significant progress has been made in addressing these challenges in the AKIB systems, particularly in electrode materials and electrolytes. For electrode materials, there has been substantial research, including a series of cathode materials of Prussian blue analogues (PBAs), polyanionic compounds, and metal oxides [1147–1149], as well anode materials including metal oxides and organic materials [1150–1152]. The development of electrolytes has seen improvements with high-concentration electrolytes and hydrogel electrolytes, significantly enhancing battery performance [1153–1155], as shown in Fig. 36, which outlines some of the recent research on cathode and anode materials and electrolytes in AKIB systems.

  Figure 36

  

  Figure 36.  Schematic illustration of cathode [883,1147–1149,1155–1157], anode [1150,1151,1159–1162] and electrolyte [1158] of AKIBs in recent years.

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  Although AKIBs have been developing for several years and have made progress in electrode materials and electrolytes, there are still certain drawbacks: (1) For the cathode, even in high-concentration electrolytes, PBAs still experience a degree of dissolution, which significantly affects the stability of batteries [1156]. Further research is needed to improve electrode stability and develop new electrode materials [1157]; (2) Currently, stable and single-plateau K-based anode materials with appropriate potentials have not been identified. Organic materials with high-potential plateaus can severely affect the effective energy density of full cells [1158–1160]. In contrast, polyanionic compounds like KTi₂(PO₄)₃ (KTP) have large polarization and varying electrochemical behaviors reported from different studies. The voltage issue in anode requires urgent resolution [1161,1162]; (3) Although high-concentration aqueous electrolytes have largely alleviated the narrow potential window issue of AKIBs, the high costs associated with these electrolytes cannot be ignored.

  In summary, the development of new high-potential, high-capacity, and non-dissolvable cathode materials; the development of novel anode materials with low polarization and low charging voltage; and the development of new hydrogel electrolytes while controlling electrolyte costs to widen the stable potential window of electrolytes will likely become important research directions for the further development of AKIBs.

  8.2 Aqueous proton batteries

  Although aqueous metal-ion batteries exhibit high theoretical specific capacities, their ion-diffusion kinetics is limited, leading to a high demand for battery systems with fast charging capabilities. In recent years, a new category of aqueous batteries, known as aqueous proton batteries (APBs), has emerged to address this issue. These batteries utilize proton charge carriers and operate by either releasing or absorbing protons (H⁺) and/or hydronium ions (H₃O⁺) during the charging and discharging processes. These ions are supplied by aqueous acidic electrolytes, such as solutions of H₂SO₄ and H₃PO₄. Protons, being the smallest and lightest cations, enable rapid transport kinetics while causing minimal structural stress during (de)intercalation processes. Moreover, protons can be transported via a fast Grotthuss mechanism, similar to Newton’s cradle, by breaking and forming hydrogen bonds with neighboring oxygen atoms in a network of water molecules (Figs. 37a and b). This unique feature improves the rate capabilities and cycling performance of APBs. For instance, Ji et al. [1163] showed that defective Cu-Fe bimetallic Prussian blue analogues (CuFe-PBA, Fig. 37c) serve as effective cathode materials with exceptional rate performance at 4000 ℃ (380 A/g), maintaining 50% of their capacity at 1 C (Fig. 37d) and demonstrating remarkable cycling stability over 0.73 million cycles. Similarly, Lu et al. [1164] demonstrated rapid proton insertion in H_1.75MoO₃ anode material (Fig. 37e), achieving a high capacity of 111 mAh/g at a current density of 2500 C (500 A/g). The pursuit of Faradic proton storage electrode materials with both high capacity and extended lifespan is a formidable challenge. This difficulty arises largely from the rapid dissolution and failure of most high-theoretical-capacity electrode candidates in acidic electrolytes. Significant efforts have been undertaken to develop novel electrode materials and tune the electrode-electrolyte interface, all with the aim of enabling rapid proton storage chemistry while maintaining structural robustness in acidic conditions. Recent advancements have spotlighted several potential electrode candidates, including Prussian blue analogues [1165,1166], transition metal oxides [1167–1170], organic compounds [1171–1180], and two-dimensional MXenes [1181,1182].

  Figure 37

  

  Figure 37.  (a) A schematic of the Grotthuss mechanism, in which proton conduction is operated by rearranging bonds along a water chain. (b) A schematic of Newton’s cradle. (c) Schematic structures of a defective Prussian blue analogue (PBA). (d) GCD curves at various current rates of defective Prussian blue analogue electrode. Inset: relationship between the hysteresis of the potential profiles and the applied current rate. The polarization between the charge and discharge profiles increases linearly with the current rate, suggesting that the rate performance of CuFe-PBA is limited more by the testing cells’ electrical resistance than the proton transport and the reaction kinetics (e) GCD curves at various current rates of H_1.75MoO₃ electrode. (a-d) Reproduced with permission [1163]. Copyright 2019, Springer Nature. (e) Reproduced with permission [1164]. Copyright 2022, American Chemical Society.

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  Research on aqueous batteries is still in its early stages. The key to advance aqueous proton batteries for practical applications lies in the development of electrode materials that combine acid resistance, stable hydrogen bonding networks, and high densities of proton storage sites. It is imperative to note that, for the present, APBs are not anticipated to supplant current commercial batteries like lithium-ion and nickel-metal hydride batteries. Rather, APBs intrinsically possess merits in rapid diffusion kinetics and elevated rate capacities, making them potentially revolutionary in bridging the energy gap between capacitors and power limitations of conventional batteries. APBs are particularly suitable for ultrafast charging applications, such as electronic devices that charge in seconds and special batteries with flexible power supply for intermittent grid energy storage. Additionally, by leveraging the low freezing points of acidic aqueous electrolytes, APBs have the potential to revolutionize energy storage at sub-zero temperatures [1183,1184].

  8.3 Aqueous zinc-based batteries

  8.3.1   Aqueous zinc-iodine batteries

  Zinc-iodine batteries (ZIBs) hold the merits of high safety, low cost and high capacity [1185–1187]. Currently, I⁻-containing electrolyte and iodine-loading cathode are two main sources of active iodine for ZIBs [1188]. For the I⁻-containing electrolyte, it is hard to evaluate the active iodine mass, and high concentration of I⁻ ions always induce the generation of polyiodides, leading to low coulombic efficiency and inferior cyclic stability. In general, ZIBs driven by the iodine-loading cathode, deliver similar work mechanism to lithium-sulfur batteries and great potential as large-scale electrochemical energy storage devices. Due to the high reactivity of iodine component in iodine-loading cathode, the I⁻ ions generated during the initial discharge process (I₂ + 2e⁻ → 2I⁻) tend to react with I₂ component to form polyiodides (I⁻ + I₂ → I₃⁻; I₃⁻ + I₂ → I₅⁻), thus leading to the active iodine dissolution [1189,1190]. With the increase of I₃^-/I₅^- concentration, these polyiodides will transfer from iodine-loading cathode to zinc anode, resulting in the severe self-discharge reaction, fast capacity attenuation, and zinc surface corrosion (Zn + I₃⁻ → Zn²⁺ + 3I⁻) [1191,1192]. Therefore, the practical application of ZIBs is greatly hindered by the issues of active iodine dissolution and polyiodide shuttle behavior.

  To address these issues, many effective strategies were implemented around the design of chief components of ZIBs, such as iodine host materials [1193,1194], functional separators [1195,1196], and non-liquid electrolytes [1197–1199]. For the iodine host materials, the common porous carriers immobilize the active iodine species by physical adsorption, while the catalytic carriers anchor the iodine species through chemisorption effect and facilitate the iodine conversion reaction kinetics [1200,1201]. For the functional separators, the surface coating materials always contain active groups to allow or block the entrance of specific iodine species [1202]. For the non-liquid electrolytes, the content of free H₂O molecules is greatly reduced to change the dissolution environment of active iodine component, simultaneously suppressing the polyiodide shuttle and zinc corrosion [1190]. Additionally, designing functional binders with active groups and strong iodine chemisorption is also considered as valid method to inhibit the iodine dissolution and polyiodide shuttle of ZIBs, including the chitosan-based binders [1203], polyacrylonitrile copolymer (LA133) [1204], liner-chain polysaccharides [1205], and sulfonated lignin [1206].

  Facing the future application of ZIBs, high energy density and durable cyclic stability are two essential requirements. In current stage, the long-term cycling stability of ZIBs is effectively achieved, and it is a critical need but still great challenge to break the limitation of two electron transfer reaction mechanism, realize multiple electron transfer iodine conversion reaction [1207,1208], and finally improve the energy density of ZIBs.

  8.3.2   Aqueous zinc-manganese bioxide batteries

  Aqueous zinc-ion batteries (AZIBs) have emerged as a competitive option due to their inherent safety, low cost, eco-friendliness, and the cathode materials are critical to the electrochemical performance of AZIBs [1209]. Manganese-based oxides possess notable advantages, including high theoretical capacity, high operating voltage, diverse crystal structures, and environmental friendliness, which are regarded as a promising cathode material for AZIBs [1210,1211]. As the representative one, MnO₂ has been confirmed a favorable ability for Zn²⁺ storage, and the crystal structure of MnO₂ depends on the connection mode of [MnO₆] octahedron units, mainly including α, β, γ, δ, λ types [1212]. Among them, α-MnO₂ with tunnel structure and δ-MnO₂ with layered structure are more favorable for the diffusion of Zn²⁺ [1213]. The main energy storage mechanisms for MnO₂ are reversible H⁺ insertion/extraction, Zn²⁺ insertion/extraction, H⁺ and Zn²⁺ co-insertion/extraction, chemical conversion reaction, and Mn dissolution/deposition reaction [1214].

  Notwithstanding these advances, the development of MnO₂ cathode still faces great challenges, such as poor durability, low electrical conductivity, and sluggish kinetics. During the charging/discharging processes, the unstable crystal structure of MnO₂ cathode leads to volume changes, structural damage, and the dissolution of active materials [1215]. For example, the disproportionation reaction of Mn³⁺ causes the dissolution of manganese into the electrolyte. Additionally, irreversible phase transitions give rise to a rapid capacity decline [1216]. Low electrical conductivity and slow ionic transport influence the reaction kinetics, which will result in poor rate performance. To address these challenges, numerous modifying strategies have been implemented, including guest species incorporation, composites construction, defect engineering, and nanostructure design, to inhibit manganese dissolution and strengthen crystal structure [1217]. The intercalation of guest species into the interlayer can be used as interlayer pillars to expand the interlayer spacing and weaken the strong electrostatic interactions between the host and Zn²⁺, thus promoting Zn²⁺ diffusion and stabilizing the crystal structure of the cathode material [1218]. Up to now, such as metallic ions (Li⁺, Na⁺, K⁺, Fe²⁺, Co²⁺, Ni²⁺, Mn²⁺, etc.), nonmetallic ions (NH₄⁺), and organic molecules (polypyrrole and polyaniline) have been introduced into the manganese-based compounds [1219,1220]. Composites construction is another effective method to enhance the conductivity and stability of electrode materials by combining MnO₂ with carbon-based materials or organic compounds [1221]. Defect engineering can enhance the electrical conductivity of materials by modifying the local electronic structure of transition metal elements and accelerating the charge transfer process [1214]. Additionally, the absence of atoms results in the formation of lattice cavities and distortions, which will optimize the diffusion paths and structures, thereby improving the reaction kinetics [1222]. Nanostructure design represents an effective method to provide more ionic/electronic transport paths [1221]. A reasonable design of nanostructured materials can increase the specific surface area, expose additional active sites, and reduce the ions diffusion distance, thereby improving the reactions kinetics and enhancing the rate capability of the cathode [1223].

  8.3.3   Aqueous zinc-vanadic anhydride batteries

  Vanadium-based oxides are considered as another appropriate cathode candidate for AZIBs due to their multiple valence states (V²⁺, V³⁺, V⁴⁺, and V⁵⁺), high theoretical capacity, and superior rate capability, which are composed of V-O coordination polyhedral with various crystal structures [1224–1226]. Among them, the layer-structured orthorhombic V₂O₅ composed of [VO₅] square pyramids with shared corners is beneficial for the storage of Zn²⁺. However, with the repeated Zn²⁺ insertion/extraction between the interlayer, the lamellar structure of vanadium oxide is destroyed along with the dissolution of vanadium, which will lead to the deterioration of the cyclic performance [1227].

  To achieve satisfactory electrochemical performance, it is imperative to address the issues of low ionic/electrical conductivity and vanadium dissolution of vanadium oxides. Guest pre-intercalation is an effective strategy to enlarge the interlayer spacing, regulate the electronic structure and lower the electrostatic interaction between vanadium oxides and Zn²⁺, thus further optimizing the electrochemical properties [1228]. Besides the water molecule introduction, the pre-intercalation of cations, such as K⁺, Ba²⁺, Ca²⁺, Al³⁺, La³⁺, NH₄⁺, can redistribute the local electron density, boost the structural stability and inhibit the vanadium dissolution, enabling the better cycling performance [1229,1230]. Nazar’s group reported a V₂O₅ nanobelts cathode pillared by Zn²⁺ and H₂O, which shows satisfactory rate and cycling performance for AZIBs [1231]. Moreover, compared to the single component vanadium-based oxides, the combination of two or more materials to construct composites can not only generate synergistic effects through the coupling interaction between different components, but also serve as a barrier between the vanadium oxide and the electrolyte to protect the cathode [1232]. Defect engineering is a common method to adjust the electron distribution of vanadium oxides, which can further improve the ionic/electronic conductivity [1221]. Peng et al. [1233] designed a new type of vanadium oxide cathode with defects that can significantly shorten the Zn²⁺ transport path and lead to better reaction kinetics.

  8.3.4   Aqueous zinc-organic batteries

  Organic materials are held together by weak van der Waals forces, and their soft molecular structure allows for rapid diffusion of divalent cations [1234]. Moreover, by altering their molecular structure, they can accommodate the intercalation and deintercalation of multivalent cations with minimal volume change. Zhao et al. [1235] found that quinone compounds exhibit a high capacity of 335 mAh/g in a Zn(OTf)₂-based aqueous electrolyte [1236]. Successively, tetrachloro-1,4-benzoquinone and (9,10-bis(1,3-dithiol-2-ylidene)-9,10-dihydroanthracene) have been reported for use as zinc-ion cathodes, demonstrating excellent capacity and rate performance [1237,1238]. However, the conductivity of organic cathodes is relatively poor, often requiring a significant number of conductive additives to enhance the electrochemical performance of the battery.

  8.3.5   Aqueous electrolytes

  The aqueous electrolyte is a crucial component of zinc batteries, complementing the cathode, anode, and separator by acting as a conductive bridge between the electrodes. Compared to organic electrolytes, aqueous electrolytes offer several advantages: (1) High ionic conductivity due to water’s high dielectric constant; (2) Eco-friendliness, as they are non-toxic and biodegradable; (3) Intrinsic safety and thermal stability, attributed to their non-flammable nature.

  However, aqueous electrolytes face several challenges (Fig. 38a), including: (1) A limited electrochemical stability window, making them incompatible with high-voltage cathodes; (2) Parasitic hydrogen evolution reactions, which consume the electrolyte and reduce coulombic efficiency; (3) The formation of passivation layers, hindering ion transport; (4) Reactivity with zinc metal, leading to the loss of active materials; (5) Low ionic conductivity at subzero temperatures. Despite these challenges, aqueous electrolytes remain a popular choice for zinc batteries.

  Figure 38

  

  Figure 38.  Solvation chemistry and Zn anode/electrolyte interfaces of (a) traditional aqueous electrolyte, (b) anion-receptor aqueous electrolyte and (c) lean-water aqueous electrolyte and physicochemical properties of different electrolytes. (d) Schematic illustration of symmetric anion zinc salt and asymmetric anion zinc salt. (e) The dominated solid-liquid transition temperature of KOH solutions with different C_KOH and the phase composition of KOH solution at different temperatures and concentrations. (d) Reproduced with permission [1246]. Copyright 2024, Wiley-VCH. (e) Reproduced with permission [1249]. Copyright 2023, Elsevier.

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  To address these challenges, several strategies have been developed [1239]. One approach is the use of an anion receptor of imidazolidinyl urea (IU), which selectively binds with anions of zinc salt to form a stable, compact, and ion-conductive SEI layer [1240]. This promotes the preferential growth of Zn (002) plane, which offers exceptional reversibility of Zn over 2000 cycles with a coulombic efficiency > 99% (Fig. 38b). Another strategy involves the establishment of the water-poor interface to suppress side reactions. The utilization of 1,3-propanediol (PDO) to replace a majority of water molecules in the solvation shell of Zn²⁺ ions significantly promote the formation of solvent-separated ion pairs (SSIP) and contact ion pairs (CIP), which stabilizes Zn²⁺ and prevents the generation of free water (Fig. 38c) [1241]. Thus, the ZnǀǀZn batteries cycle over 4000 h at 1 mA/cm², remarkably outperforming that using traditional electrolyte (698 cycles).

  Although organic additives with high-reduction potential are commonly used to stabilize the Zn anode, the introduction of organic species into aqueous electrolytes compromises the safety and environmental compatibility, reduces the ion conductivity, and increases the desolvation barrier of Zn deposition [1242,1243]. Aqueous electrolytes based on conventional Zn salt, such as ZnSO₄, ZnCl₂, Zn(OTf)₂ and Zn(TFSI)₂ often fail to provide good cycling stability because their salt anions are difficult to decompose, hindering the formation of a robust SEI [1244,1245]. An effective approach to improving interfacial stability involves modulating the zinc salt. A notable example is the design of Zn(DFTFSI)₂ salt that can not only achieve good ion conductivity, high transference number, and but also promote the formation of robust SEI for Zn electrodes by decomposition of DFTFSI anion (Fig. 38d) [1246]. As a result, the batteries achieved a ZnǀǀZn cell with long-term cycling stability of 2500 h with 1250 mAh/cm² cumulative capacity, and outperforms many other conventional Zn salt electrolytes.

  To improve the performance of aqueous electrolytes in subzero conditions, a Zn(BF₄)₂ salt-based aqueous electrolyte with a symbiotic water-acetamide solvation sheath has been developed. This electrolyte reduces viscosity and freezing point, enhances conductivity and ion transfer, and forms a ~15 nm SEI to prevent side reactions such as hydrogen evolution and Zn anode corrosion [1247]. The amide group in the acetamide-water electrolyte breaks the hydrogen bond network between water molecules, lowering the freezing from 12 ℃ to −45 ℃.

  In more extreme conditions, particularly in alkaline electrolytes, achieving antifreezing ability and stability is more challenging. The addition of dimethyl sulfoxide (DMSO) as an antifreezing additive in alkaline aqueous electrolytes represents a significant advancement for improving the low-temperature performance of Zn||Ni batteries. DMSO reduces the freezing point of the electrolyte to −90 ℃, but its effectiveness is limited in the highly corrosive alkaline environment [1248]. Nevertheless, when combined with an 8 mol/L KOH electrolyte, the freezing point can be dramatically lowered to −120 ℃ (Fig. 38e), ensuring electrolyte functionality in extreme cold conditions [1249].

  8.4 Other aqueous multivalent metal-ion batteries

  Aqueous multivalent metal ion batteries use multivalent metal ions such as zinc, magnesium, calcium, and aluminum as carriers, and the reaction mechanism of the positive and negative electrodes includes embedded reaction, alloying reaction, and conversion-type reaction, etc. [1250]. Generally speaking, aqueous multivalent metal ion batteries have the advantages of low cost, abundant raw materials, good safety, and high-power density. The electrolytes used in aqueous multivalent metal ion batteries are mostly near-neutral or alkaline solutions, which provide better environmental compatibility, and the theoretical energy density of aqueous multivalent metal ion batteries is high. In addition, the natural abundance of metal elements such as zinc, magnesium, calcium, and aluminum are high and the production cost is low, so the high-performance aqueous multivalent metal-ion batteries have the potential to meet the future demand for energy storage on a large scale. However, aqueous multivalent metal-ion batteries also suffer from a number of problems such as complex reaction processes, narrow electrochemical windows, low actual energy density, and serious side reactions, which limit their practical applications. In recent years, researchers at home and abroad have made many efforts to address these problems and have achieved a series of progress [1251].

  The advantages of aqueous multivalent metal ion batteries are (1) Good safety: compared with organic electrolytes, aqueous electrolytes have the advantages of being nontoxic, nonflammable, and nonvolatile, which can greatly improve the safety of the batteries [1252]. (2) Low cost: multivalent metal elements (Mg, Ca, and Al) are abundant in resources and at a low cost. In addition, the use of aqueous electrolytes eliminates the complex battery assembly and manufacturing requirements, effectively reducing the production cost of the battery. (3) High power density: the ionic conductivity of aqueous electrolytes can reach ~1 S/cm, which is two to three orders of magnitude higher than the ionic conductivity of commonly used organic electrolytes (10⁻³–10⁻² S/cm). As a result, aqueous batteries typically have fast charge/discharge rates [1253]. On the other hand, multivalent metal cations have higher charge densities and stronger interactions with water molecules and material lattices than monovalent lithium, sodium, etc. Multivalent metal cations are present in aqueous electrolytes as hydrated ions. During many electrode reactions, multivalent metal ions are embedded as hydrated (or partially hydrated) ions, and the shielding effect of water molecules can overcome electrostatic interactions with the electrode material and promote interfacial diffusion, allowing for faster charge and discharge rates and higher power densities [1254]. (4) Theoretically high volumetric energy densities: more than one charge equivalent can be transferred per mole of a multivalent metal cation and the density of multivalent metals is generally high, the theoretical volumetric specific capacity advantage is obvious [1255]. However, it should be noted that the use of metal Mg and Ca anodes in aqueous batteries is still very difficult. In addition, factors such as cathode materials and electrolytes should also be considered when comparing the actual energy densities, and the calculation and evaluation of the actual energy densities of multivalent metal-ion batteries are still inaccurate and need to be further explicitly verified [1255].

  The challenges faced by aqueous multivalent metal ion batteries are: (1) The working mechanism is still unclear: unlike the mature aqueous monovalent metal ion (lithium/sodium ion, etc.) battery energy storage system, the working mechanism of aqueous multivalent metal ion batteries is more complex, and the specific reaction processes of certain anodes have not been thoroughly investigated. Strong interactions between multivalent ions and water molecules result in the formation of a tightly solvated structure, which is highly susceptible to charge transfer under the polarization potential to form protons and hydroxides, and the electrochemical embedding of protons likewise interferes with the reaction process. Polarization potential is highly susceptible to charge transfer and the formation of protons and hydroxides, while the electrochemical embedding of protons likewise interferes with the reaction process. Therefore, it is necessary to use advanced characterization techniques to probe deeply into the reaction mechanism of aqueous multivalent metal-ion batteries and to establish the dynamic course of the charging/discharging process to determine the relationship between the physical/chemical changes in the surface and bulk phases of the cathode material and the electrochemical performance [1256]. (2) Narrow electrochemical window of the electrolyte: The hydrolysis reaction consists of the precipitated oxygen reaction on the cathode side (neutral/alkaline solution for 4OH⁻ → 2H₂O + O₂ + 4e⁻, while 2H₂O → 4H⁺ + O₂ + 4e⁻ in acidic solutions) and hydrogen precipitation reactions on the negative side (2H₂O + 2e⁻ → 2OH⁻ + H₂ in neutral/alkaline solutions, and 2H⁺ + 2e⁻→H₂ in acidic solutions). The water decomposition process limits the electrochemical window of a typical aqueous electrolyte to 1.23 V, which also limits the operating voltage of aqueous multivalent metal-ion batteries, resulting in insufficient energy density [1257]. In addition, O₂ and H₂ produced on the surface of the electrodes not only disrupt the electrode structure and segregate the electrolyte, resulting in large overpotentials (polarization) and cell swelling, but they also continuously consume the electrolyte, leading to poor cell Coulombic efficiency and cycling stability decreases [1258]. The side effects of water decomposition are often overlooked under test conditions where the electrolyte is used in excess, whereas the effects of water decomposition on the stability of the battery will be particularly prominent in practical battery systems where the amount of electrolyte is limited, and further research is needed. (3) Dendrite growth of metal negative electrode, aqueous multivalent metal-ion batteries can use multivalent metals as negative electrodes, mainly including Zn, Mg, Al, etc. However, except for Mg metal, the vast majority of metal-negative electrodes generally suffer from the problem of dendrite growth during the dissolution/deposition process. The continuous growth of dendrites can puncture the diaphragm and lead to a short circuit of the battery, which not only reduces the cycle life of the battery but even poses a safety issue. Despite the similarity of dendrite growth caused by kinetic factors, there are differences in the thermodynamic origins of dendrite formation during the deposition of different metal anodes, and the mechanical properties of the dendrites are characterized differently, and these differences make it difficult to solve the problems caused by metal dendrites by a unified approach [1259].

  Currently, the multivalent metal ions reported in the study for aqueous battery research include Mg²⁺, Ca²⁺, and Al³⁺, etc., but the choice of cathode materials for multivalent ions is still relatively limited [1260]. Due to the strong interaction of multivalent metal ions with the crystal lattice, the intercalation-type cathode materials suitable for monovalent cations are often unable to realize the reversible de-embedding of multivalent metal ions, which has become an important factor restricting the development of aqueous multivalent metal ion batteries. As for the anode, structural design and surface modification of the metal anode to improve its cycling performance and inhibit the growth of dendrites is an important research direction for the multivalent metal anode. On the other hand, adjusting the composition (solvent and salt) of the aqueous electrolyte can change the solvation structure, which in turn regulates the improvement of the bulk phase migration kinetics of multivalent metal ions, and changes the interfacial composition and properties of the positive and negative electrodes. The optimally designed electrolyte can effectively broaden the electrochemical window, stabilize the cycling of anode materials, and inhibit the generation of dendrites at the metal anode. Therefore, the electrolytes for multivalent metal ion batteries have also received extensive attention.

  9. Redox flow batteries

  9.1 Vanadium redox flow batteries

  A vanadium flow battery (VFB) is a type of rechargeable battery that stores electrical energy by electrochemical reactions of vanadium ions in liquid electrolytes. VO₂⁺/VO²⁺ and V³⁺/V²⁺ redox couples are used for the positive and negative electrode respectively. Sulfuric acid is used as the supporting electrolyte in general. The electrochemical reactions of VFBs during charge and discharge process are shown as follows (Eqs. 5–7).

  Positive:VO2++H2O−e−↼⇁VO2++2H+Eθ=1.00Vvs.SHE

  (5)

  Negative:V3++e−↼⇁V2+Eθ=−0.26Vvs.SHE

  (6)

  Overall:VO2++H2O+V3+↼⇁VO2++V2++2H+Eθ=1.26V

  (7)

  VFBs were first proposed in 1984 by Maria Skyllas-Kazacos’ team at the University of New South Wales, Australia. And then Maria team constructed a 1 kW VFB stack system in 1991. VFBs were gradually concerned about by many companies and research institutes with the increasing demand for large-scale energy storage. In 1993, Sumitomo electric company in Japan got authorization to develop VFBs. Sumitomo electric company made a great breakthrough and installed the largest system of a 15 MW/60 MWh VFB at that time in 2015. Recently, Dalian Rongke power company and Dalian Institute of Chemical Physics (Chinese Academy of Sciences) had built 100 MW/400 MW VFB system in 2022, that is the largest flow battery station all over the world. This system accelerates the industrialization of flow battery and the utilizations of renewable energy, which has established China’s world leading position in the basic research and engineering application of vanadium flow batteries.

  VFBs have attracted much interest owing to the advantages of long cycle life, high safety, and independent regulation of power and energy. In addition, the recycling and reusing of the electrolyte are simple due to the same vanadium ion solutions in anolyte and catholyte, which is another important advantage of VFBs. VFB cell stack is mainly layer-by-layer assembled of key components, including membranes, electrodes and bipolar plates. It has passed the scale-up stage of cell stack from kilowatts to tens of kilowatts, and entered the development stage of high power density cell stack (e.g., 42 kW cell stack, up to 180–220 mA/cm²). The structure optimization and materials innovations are the main ways to improve the cell stack power density in the future. Among the key materials, membranes are used to separate the vanadium ions while selectively transport carriers to complete the internal circuit. The most widely used membrane is Nafion membrane from Chemours company. That exhibits high conductivity and stability, but the low selectivity needs to be improved. The modification of Nafion membranes, other ion exchange membranes, and new-type porous membranes are attracting massive researches. The electrodes provide the places where electrochemical reactions of positive and negative redox couples occur. At present, carbon felt and graphite felt are mainly used as electrode. These materials display the good electrical conductivity, but the electrochemical activity need to be further improved. Researchers have tried the methods, such as oxidation treatment, doping with heteroatoms, and introducing catalysts, to improve the activity of electrodes. Bipolar plate is another important component to conduct the electrons. The graphite bipolar plate with high electrical conductivity is mainly used in laboratory studies, but it is fragile and expensive. For large-scale VFB stacks, carbon plastic bipolar plate is a top priority, whose properties can be optimized by the ratio of carbon materials and organic materials. It might be the key direction of bipolar plate to design the flow channel structures or integrate electrode and bipolar plate in the future. Vanadium electrolyte provides the reactant and determines the energy density of the vanadium battery. Until now, the sulfate acid vanadium electrolyte, hydrochloric acid vanadium electrolyte and mixed acid electrolyte are proposed. But the amount of vanadium used per kilowatt-hour has not decreased. The development of hybrid electrolytes reducing the amount of vanadium, is a future research direction.

  VFB technology is going through the stages from scientific research to practical application after the development of nearly 40 years. Although the initial investment cost is relatively high due to the expensive price of vanadium ore, whole life cycle cost of VFB system can be accepted by the market. VFBs with high safety show the great advantages especially for energy storage systems with large capacity. However, VFB technology still faces the challenges of the low power and energy density. The technical breakthroughs of the key components (such as membranes, electrode, bipolar plates, electrolytes) and the optimization of high-power density cell stack are expected to promote the industrialization of VFB technologies.

  9.2 Iron-based redox flow batteries

  All-iron redox flow batteries (AIRFBs) have been received substantial attention owing to low price and abundant supply of raw materials. The first AIRFB was pioneered in 1981 in an acidic environment and employed Fe²⁺/Fe and Fe³⁺/Fe²⁺ as redox pairs [1261]. In the charging/discharging cycle of acidic AIRFB, the reduction/oxidation reactions and their corresponding reaction potentials are as follows:

  AnodeFe2++2e−⇌Fe(0)E0=−0.44V

  (8)

  CathodeFe3++e−⇌Fe2+E0=0.77V

  (9)

  The anode redox reaction of the acidic AIRFBs involves the process of Fe plating/stripping, unlike the catholyte which consists of a solvated aqueous Fe³⁺/Fe²⁺ redox couple. In this sense, the acidic all-iron RFB can be defined as a hybrid-type RFB, where the power and capacity are not completely decoupled, thus presenting significant challenges including poor Fe anode reversibility, hydrogen evolution reaction (HER), and metal dendrite formation. To improve the reversibility of the Fe anode, Tang et al. [1262,1263] introduced sodium citrate, and dimethyl sulfoxide (DMSO), which can intimately coordinate the Fe²⁺ and can change the preferred crystal plane, respectively. In addition, both of them modified the solvation structure of Fe[(H₂O)₆]²⁺, greatly enhancing the reversibility of the plating/stripping reaction at the Fe anode. Consequently, it achieved nearly 100% Coulombic efficiency and extended the cycle life to 300 cycles. They also introduced a surface-engineered carbon felt with abundant defects, which improves Fe²⁺/Fe kinetics and promotes uniform Fe deposition and dissolution [1264]. The modified electrode also enhanced the reversibility of the Fe anode, as confirmed by both theoretical calculations and practical tests, achieving high power density (80 mW/cm²) and stable operation over 250 cycles with 99% Coulombic efficiency. Furthermore, Lu et al. adopted high load solid conversion electrode Fe₃O₄/Fe(OH)₂ instead of metallic deposit electrode to realize high cycle stability under high surface capacity or current density [1265]. The solid conversion electrode eliminates dendrite issues and the limitations of metal areal capacity, and demonstrates 215 mAh/cm² at 60 mA/cm² for 100 cycles (700 h) without capacity decay. Narayanan et al. [1266] demonstrated a high-performance rechargeable iron electrode composed of carbonyl iron and bismuth sulfide, which achieved a ten-fold reduction in hydrogen evolution rate, a discharge capacity of 0.3 Ah/g, and a twenty-fold increase in capacity at the 2-h discharge rate. The high purity of carbonyl iron, combined with in situ-produced bismuth electro-deposits, suppressed the wasteful hydrogen evolution, while the in-situ formation of mixed-valent conductive iron sulfides promoted high discharge rates.

  Alternatively, all-soluble AIRFBs, constructed via the chelation effect between ligands and iron ions and employing soluble iron(Ⅲ/II)-ligand complexes in both anolyte and catholyte, offers a more promising candidate for large-scale energy storage. Chelation can significantly alter a metal ions’ pH stability, redox potential, and solubility. Therefore, all-soluble AIRFBs can be constructed by selecting suitable ligands as long as two different iron-ligand complex pairs provide appropriate potential difference. Based on the different pH levels of the electrolyte, all-soluble AIRFBs can be classified into near-neutral and alkaline types. Noemí et al. combined the Fe-N, N’-ethylene-bis-(o-hydroxyphenylglycine) complex (Fe(racEDDHA)) as anolyte and Fe(CN)₆ as catholyte, the near-neutral Fe(racEDDHA)/Fe(CN)₆ AIRFB displayed a voltage of 0.848 V [1267]. Under the operating pH 8–9, all the electrochemical-active redox substances exist in the form of stable water-soluble iron complexes and the cell achieved a capacity retention of nearly 100% over 75 cycles. Gabriel et al. [1268] developed a novel AIRFB that operates in near-neutral conditions by pairing Fe-(chloride and nitrilotri(methylphosphonic acid)) complex (Fe(NTMPA)₂) and Fe(CN)₆ complex. The designed near-neutral AIRFB demonstrated high cycling stability over 1000 cycles, and noteworthy performances, including 96% capacity utilization, a minimal capacity fade rate of 0.0013% per cycle, high Coulombic efficiency and energy efficiency near 100% and 87%, respectively. To further enhance the battery voltage, Fe-N, N-bis(phosphonomethyl)glycine (Fe(BPMG)₂) was proposed and used in near-neutral AIRFB. Compared to Fe(NTMPA)₂, Fe(BPMG)₂ demonstrated a negatively shifted redox potential, resulting in ≈ 0.07 V (≈ 10%) increase in battery output voltage [1269].

  Alkaline AIRFBs typically employ Fe(CN)₆ complexes and iron-polyhydroxy ligand complexes as the catholyte and anolyte, enabling operation in highly alkaline environments with pH ≥ 14. The first alkaline AIRFB was proposed by Yan et al. [1270], which utilized Fe(CN)₆ complex (0.48 V vs. SHE) and Fe-triethanolamine complex (Fe(TEA), −0.86 V vs. SHE) as the catholyte and anolyte, respectively. The Fe(CN)₆/Fe(TEA) AIRFB exhibited a cell voltage of 1.34 V, a Coulombic efficiency of 93%, and an energy efficiency of 73% at 40 mA/cm². Unfortunately, Fe(TEA) was prone to decomposition during charge-discharge cycling due to weak complexation between Fe²⁺ and TEA. Kwon et al. [1271,1272] employed some TEA derivatives, such as 3-[bis(2-hydroxyethyl)amino]-2-hydroxypropanesulfonic acid (DIPSO) and bis-(2-hydroxyethyl)amino-tris(hydroxymethyl)methane (BIS-TRIS), as the iron ions ligand in catholyte, which offered a similar redox potential of Fe(TEA). The constructed alkaline AIRFB exhibited enhanced stability with a Coulombic efficiency of 99.5%. However, although the molecular structure modification was used to improve the stability, the alkaline AIRFBs usually demonstrated a short lifespan (<250 cycles). To further enhance battery life, Li et al. [1273] proposed a low-cost alkaline AIRFB by coupling iron/ferrous cyanide redox pairs with iron/ferrous gluconate complex redox pairs. The six coordinated iron core makes the redox pair of ferric iron/ferrous gluconate complex highly stable. The designed AIRFB demonstrated an energy efficiency of 83% and a Coulombic efficiency of more than 99% at 80 mA/cm². It can operate continuously for more than 950 cycles for continuous charging/discharging (530 h) without significant capacity degradation. Cheng et al. [1274] employed a double-ligand strategy to prepare a six-coordinated Fe(TEA)MM complex by introducing two ligands (triethanolamine (TEA) and 2-methylimidazole (MM)) on a single iron center, which significantly increased the binding energy of the complex and achieved 1400 stable cycles. Duan et al. [1275] systematically analyzed the capacity attenuation mechanism in alkaline AIRFBs using two unbalanced batteries and spectroscopy techniques. Their research suggests that, in addition to the issue of complex decomposition, ligands crossover issues is another root cause of capacity decay. Building on this degradation mechanism, they addressed both ligand crossover and complex decomposition issues by designing several new poly-sulfonated ligands, such as (3,3′-((2-hydroxyethyl)azanediyl)bis(2-hydroxypropane-1-sulfonic acid) (TEA-2S) and sodium 3,3′,3″,3″′-(ethane-1,2-diylbis(azanetriyl))tetrakis(2-hydroxypropane-1-sulfonate) (EDTS) [1276,1277]. On the one hand, by skillfully designing the ligands, the binding energy of the iron complexes increased significantly. On the other hand, the poly sulfonated-ligands with high charge density and large volume effectively mitigates crossover through the size exclusion and the Donnan effects. Paired with the Fe(CN)₆ electrolyte, the Fe(CN)₆/Fe(EDTS) AIRFB and Fe(CN)₆/Fe(TEA-2S) AIRFB demonstrated high stability with capacity retention of 96.08% after 3000 cycles and 97.2% after 2000 cycles, respectively.

  In summary, AIRFBs offer great potential for large-scale energy storage due to their low cost and abundant materials. However, several challenges remain that must be addressed. Future improvements can be carried out in these directions: (1) For the acidic AIRFBs, future research should explore strategies for addressing the HER issue while maintaining the electrochemical performance of AIRFBs, such as adding hydrogen evolution inhibitors. (2) For the near-neutral AIRFBs, the battery voltage is relatively low (<1 V), and future research should focus on exploring iron complexes for the catholyte and anolyte with a larger voltage difference. (3) For the alkaline AIRFBs, the overall energy density is limited by the low solubility of Fe(CN)₆. Therefore, future research should focus on designing high solubility positive iron active species to improve the overall energy density of alkaline AIRFBs.

  9.3 Zinc-based redox flow batteries

  Flow batteries, as one of the optimal technologies for electrochemical energy storage, are playing an increasingly important role in storage of intermittent renewables such as wind and solar power [1278]. The flow batteries can be divided into all-liquid flow batteries such as vanadium flow battery, iron chromium flow battery and organic-based flow batteries, and hybrid flow batteries such as zinc-based flow batteries, all iron flow batteries and lead-based flow batteries [1279].

  As the representative hybrid flow batteries, the zinc-based flow batteries are suitable for stationary energy storage applications in virtue of their high energy density and low-cost advantages. Based on the structure of zinc species in anolyte, the zinc-based flow batteries can work efficiently across a wide pH range [1279]. Compared with state-of-the-art vanadium flow battery, the redox reaction for anode of zinc-based flow batteries is a plating-stripping process of zinc species. This process will bring the zinc-based flow batteries many critical issues, i.e., dendritic zinc in anode, limited areal capacity and poor voltage consistency for series-wound single cells or cell stacks. In addition to anode, the redox couples in catholyte also suffer from their intrinsic issues. For example, the Ni(OH)₂/NiOOH or complexed bromine species shown inferior kinetics, the ferro/ferricyanide redox couple suffered from low solubility in alkaline media [1280]. All these issues hinder the progress of zinc-based flow batteries applications.

  To accelerate the application of zinc-based flow batteries, current researches focus on the scientific understandings of the fundamental design of the key materials including electrodes, electrolytes, membranes, and their chemistry in relation to the battery performance. For example, to address the issues of zinc dendrite in anode, membrane structure design or solvation structure manipulating of zinc species have been carried to adjust the membrane-electrode interface’s properties, such as the distribution of temperature, mass transport process, which are critical factors affecting the plating process of zinc redox couple. In recent years, we have developed many composite membranes whose functional layer (e.g., boron nitride nanosheets (BNNSs) [1281], montmorillonite (MMT) [1282] and layered double hydroxides (LDHs) [1283]) can be designed flexibly to manipulate the plating of metallic zinc for zinc-based flow batteries, which provides an effective way for zinc morphology control. Based on the plating process of zinc redox couple, we recently engineer an artificial bridge between anode surface and anolyte enabled by ethylenediaminetetraacetic acid (EDTA) organic ligand to construct directional three-dimensional transport channel for zinc species, which realizes the fast mass transfer of zinc species from bulk solution and facilitates a uniform concentration distribution of zinc species at interface of anode [1284]. Benefiting from this artificial bridge between anode surface and anolyte, the growth angle of zinc can be controlled, favoring resolving the issue of zinc dendrite. These strategies provide route to address the issue of zinc dendrite and offer path for developing high-reliability and sustainable zinc-based flow batteries.

  Different from the zinc redox couple, most of redox couples in catholyte will not experience complicated phase transition process, however, their reaction on cathode is still confronted with challenges, such as the low kinetics of complexed bromine species, the low solubility of Fe(CN)₆⁴⁻/Fe(CN)₆³⁻ redox couple. Taking zinc–ferricyanide flow battery as an example, the ferri/ferro-cyanide demonstrates low solubility (~0.4 mol/L at 25 ℃) in alkaline media and can easily precipitate below room temperature. Although the mild alkalescent electrolyte can enable a zinc-ferricyanide flow battery using 0.8 mol/L ferrocyanide working stably at 0 ℃ [1285], the distribution of renewables in cold regions, such as north-western China, necessitates a more anti-freezing (below 0 ℃) electrolyte for this kind of battery if additional heating is to be avoided. The ferri/ferro-cyanide in this alkalescent media can still form precipitates below 0 ℃ and cause the battery failure. To realize stable operation below 0 ℃ for ferri/ferro-cyanide redox couple, a lithium-based supporting electrolyte was used to break the compact adsorption configuration between cations and Fe(CN)₆⁴⁻, and incorporate more water molecules among Fe(CN)₆⁴⁻. This enables 0.8 mol/L ferri/ferro-cyanide redox couple to access a broad range of physicochemical and electrochemical properties at −10 ℃ [1286].

  The materials chemistries also accelerate the forward progress of zinc-based flow battery cell stack and systems, which have been summarized in our previous work [1287]. Currently, significant advances have been achieved for the zinc-based flow battery technologies. However, it is still insufficient to realize the large-scale and commercialized application of the zinc-based flow battery technologies and ongoing challenges still remain. The research and development of advanced materials are of very great importance to promote the applications of zinc-based flow battery technologies, where they typically supplement each other within a system. The structure design and assembling of cell stack, and system integration for demonstration will advance the progress of zinc-based flow battery technologies.

  9.4 Polysulfide-based redox flow batteries

  Aqueous polysulfide-based redox flow batteries have garnered significant attention in recent years owing to the extremely low cost and high earth-abundance of sulfur [1288]. The fundamental chemistry of polysulfide-based negolytes involves the conversion of long-chain polysulfide (S₄^2–) and short-chain polysulfide (S₂^2–) species (Eq. 10) [1289].

  S42−+2e−↔2S22−

  (10)

  The first implementation of aqueous polysulfide-based RFBs was the polysulfide-bromine (S-Br) RFB. However, its commercialization has been hindered by significant crossover issues involving both polysulfide and bromide species. To circumnutate the toxicity and corrosion issue associated with bromine, Lu et al. [1290] proposed a high-energy polysulfide-iodide (S-I) RFB, achieving a high energy density of 43.1 Wh/L_{posolyte+negolyte} and a low chemical cost ($85.4 kW/h). Wang et al. [1291] reported a polysulfide ferrocyanide RFB (S-Fe RFB) with a cobalt-decorated graphite electrode, showing an energy efficiency of 74% at 20 mA/cm². This S-Fe RFB stably operated for 100 cycles with a high coulombic efficiency (CE) of 99%. The crossover issue is the main obstacle that prevents polysulfide-based RFB from long-term operation. Surprisingly, the charge-reinforced ion-selective (CRIS) membrane, made by a very simple strategy of coating a carbon layer on the Nafion cation exchange membrane, could effectively mitigate this issue and maintain a long-term operation [1292]. Another challenge of polysulfide-based RFB is the poor kinetics. Various solid catalysts like NiS_x, CoS_x, CuS_x, etc. have been reported to improve the kinetics [1293,1294]. Additionally, introducing soluble catalysts into the negolyte solution via homogenous catalysis presents another effective approach. Lei et al. [1295] reported a molecular catalyst riboflavin sodium phosphate (FMN-Na), an isoalloxazine derivative, to improve the kinetics of polysulfide species by transferring the sluggish electrochemical reaction into a fast electrochemical-chemical route ((11), (12)).

  FMN3−+2e−→FMN5−

  (11)

  FMN5−+S42−→FMN3−+2S22−

  (12)

  This homogenous catalysis strategy provides a new approach to address the bottleneck of polysulfide species and it can be universally applied to other redox-active organic molecules [1296,1297].

  One promising approach to enhance the energy density of polysulfide redox flow batteries (RFBs) is to adopt a hybrid RFB system that incorporates solid or gas on the positive side [1298,1299]. Manganese is an attractive candidate for the active material due to its low cost and high potential [1300]. Lei proposed a sulfur manganese RFB [1301] based on the deposition-dissolution of Mn²⁺/MnO₂ reaction, which demonstrated an ultra-high areal capacity of 100 mAh/cm² and a low chemical cost of $11 kW/h. MnO₄^2–/MnO₄^– is another promising posolyte with high potential, while its intrinsic stability is a challenge now [1302,1303]. Chiang et al. [1304] reported a sulfur oxygen RFB with a ceramic solid electrolyte with lower chemical cost, while its operating current density and energy efficiency are limited because of the huge resistance from solid electrolyte. Xia et al. [1305] successfully improved the power capacity by implementing a three-chamber system utilizing both a cation-exchange membrane and an anion-exchange membrane to replace the solid electrolyte. Nonetheless, challenges remain, including low areal capacity and high overpotential from gas evolution reactions, which must be addressed before commercialization can proceed. Exploring high-energy and high-potential positive electrolytes is believed to further enhance the performance of polysulfide-based RFBs and meet commercialization requirements.

  9.5 Organic redox flow batteries

  Aqueous organic redox flow batteries (AORFBs) utilize low-cost and water-soluble organics compounds composed of earth-abundant elements such as C, H, O, and N as energy storage medium, demonstrating significant advantages of environmentally friendly, high safety, and design flexibility. Currently, the development of AORFBs with high power/energy density, stability, and long lifetime remains a critical challenge. Tailoring redox-active organic molecules (ROMs) provides effective solutions by achieving high solubility, highly positive/negative redox potential, multi-electron transfer capability, and robust chemical/electrochemical stability. In this review, we review the latest advancements in redox-active molecular engineering, focusing on advanced strategies designed to achieve above objectives.

  (1) High solubility

  The energy density of AORFBs is proportional to the solubility of ROMs. Theoretically, molecules with highly hydrophilic structures exhibit better solubility in water, whereas hydrophobic molecular are more soluble in organic solvents [1306]. Therefore, the incorporation of hydrophilic functional groups, such as hydroxyl, sulfonic, carboxylic, and amino groups were frequently employed due to their high polarity and hydrophilicity [1307]. Additionally, addition of functional additives can also improve solubility of ROMs through the hydrotropic effect. For instance, urea can be utilized as a hydrotropic agent that induced the location-specific intensified solvation [1308].

  (2) Multi-electron transfer

  Developing multiple electrons system fundamentally enhance the capacity and energy density of AORFBs. According to the nature of reaction mechanism, withdrawing groups such as sulfonate and carboxylate were introduced to decrease the p K_a of reduction product. This approach prevented the irreversible protonation of radical anion formed upon reduction, thereby enabling a reversible two-electron redox reaction under battery operating conditions [1309]. Another strategy involves expanding the active centers. For example, by designing molecules with six N atoms as redox-active sites, Hu et al. achieved a reversible three-step six-electron reaction, resulting an impressive full battery capacity of 37.2 Ah/L [1310].

  (3) Highly positive/negative redox potential

  The redox potential of ROMs is a crucial factor influencing the battery voltage that correlates with the power and energy density of AORFBs. Achieving a higher battery voltage necessitates a more positive potential for the catholyte and a more negative potential for the anolyte. This can be accomplished by tuning the energy level of frontier HOMO (oxidation) and LUMO (reduction) orbitals of organic molecules. Researches have focused on functional substituent strategies that grafting electron-donating groups (e.g., –OH, –NH₂, –CH₃) leading to a negative potential shift, or electron-withdrawing groups (e.g. –SO₃Na) leading to a positive potential shift [14]. Additionally, Fan et al. proposed a novel perspective that lower-radical charge population sum (CPS) contributed to higher redox potential [1311]. Building on this concept, researchers further developed a five-membered ring nitroxide radical with a decreased CPS of −0.441 a.u., which corelated to a higher redox potential of 0.76 V (vs. Ag/AgCl) [1312].

  (4) Strong stability

  The stability of ROMs is pivotal in determining the cycling lifetime of AORFBs. Instability of these molecules mainly arises from various interactions, including side reactions, aggregation, precipitation, and crossover. Molecular design engineering has shown significant advantages in addressing these issues through: (ⅰ) Grafting of functionalized groups [1313], (ⅱ) extension of conjugate π-bond systems [1314], and (ⅲ) spatial structure regulation [1315]. For example, Li et al. [1316] demonstrated a spatial transformation from a sigmoid to a rod shape by confining the movement of the alkyl chain and the sulfonic anion. This modification resulted in weakened charge attraction and larger molecular dimension, leading to an ultralow membrane permeability rate of 1.25 × 10⁻¹⁰ cm²/s, which is merely 14.7% that of the typical sigmoid structure.

  In summary, we reviewed advanced strategies of ROMs based on their intrinsic features, aiming to achieving AORFB system with high power/energy density, high stability and long cycling lifetime. Nowadays, machine learning (ML) has revolutionized various fields by enabling data-driven predictions and optimizations. The integration of ML with molecules engineering presents a promising avenue for accelerating the development of next-generation AORFBs. By leveraging large datasets and advanced algorithms, ML can predict the electrochemical properties of potential active molecules. This predictive capability allows for the rapid screening and identification of optimal molecules, reducing the time and cost associated with experimental trial and error. Therefore, future research should focus on refining ML models and expanding their applications within AORFB technology to fully realize these benefits.

  10. Anion batteries

  10.1 Chlorine-ion batteries

  10.1.1   Concept and advantages

  Rechargeable chloride ion batteries (RCIBs) have attracted attention in recent years due to their unique merits, including high theoretical energy density (2500 Wh/L, 1100 Wh/kg), abundant sources (natural minerals and seawater), and environmental friendliness [1317,1318]. The working principle of RCIBs similarly to cation shuttle-based batteries, but use chloride ions as the charge carrier. Chloride anions usually form hybrid covalent-ionic bonds within the host materials, which allows them perform conversion or conversion-insertion reactions to provide additional redox charge transfer. The concept of RCIBs were first developed by Zhao et al. [1319] in 2014 using metal oxychlorides (CoCl₂, VCl₃, and BiCl₃) as cathode materials and [OMIM][Cl]/[BMIM][BF₄] as electrolyte. Since then, various nonaqueous/aqueous RCIBs have been extensively investigated.

  10.1.2   Chloride storage materials

  The exploration of chloride storage materials is one of the major challenges facing the development of RCIBs. Although RCIBs is successfully realized for a decade, only a few materials have been shown to be able to achieve reversible chloride ion storage, including metal chlorides/oxychlorides, layered double hydroxides (LDH), organic polymers, and metal carbides and nitrides (MXenes).

  Based on the conversion reaction mechanism, metal chlorides (CoCl₂, BiCl₃, VCl₃, and CuCl, etc.) [1320,1321], and oxychlorides (BiOCl, VOCl, FeOCl, WOCl₄, and Sb₄O₅Cl₂, etc.) [1322,1323] have high theoretical capacities. However, they suffer from severe capacity decay during the cycling processes due to irreversible material dissolution and large volumetric expansion. To alleviate this problem, Zhao et al. [1324] proposed a nanoconfinement strategy to improve the structural stability of FeOCl (Fig. 39a). By utilizing porous carbon (CMK-3) as a matrix, the electronic conductivity and structural integrity of FeOCl/CMK-3 composites are greatly enhanced. As a result, the as-prepared nanocomposite cathode shows a high reversible capacity of 202 mAh/g and superior cycling stability (Fig. 39b). LDHs are a typical class of layered materials with large interlayer spacing, and their open 2D channels are conducive to the rapid transport of chloride ions. Han et al. [1325,1326] synthesized a series of LDHs with chloride ions in the interlayer as the cathodes of RCIBs (Fig. 39c). Among them, the trimetallic Ni₂V_0.9Al_0.1-Cl LDH exhibited the highest discharge capacity (312.2 mAh/g@200 mA/g) and the longest cycle life (1000 cycles) [1327]. Detailed characterization demonstrates that the high chloride storage capacity originates from redox reaction of vanadium metal. In addition, no changes in XRD diffraction peaks were observed during cycling suggesting that the layered structure of LDHs remained stable after repeated insertion/extraction of chloride ions.

  Figure 39

  

  Figure 39.  (a) Schematic illustration of the preparation process of FeOCl/CMK-3. (b) Galvanostatic curves of the FeOCl/CMK-3. (c) Crystal structures of CoFe-Cl and Ni₂V_0.9Al_0.1-Cl LDHs. (d) Cycling performance of PPyCl. (e) Schematic of the chloride ion storage mechanism in PPyCl. (f) HAADF-STEM images and atomic structural models of the Ti₃C₂Cl_x. (g) Schematic of the charge storage mechanism of the seawater-based ACIBs. (a, b) Reproduced with permission [1324]. Copyright 2017, American Chemical Society. (d, e) Reproduced with permission [1330]. Copyright 2017, American Chemical Society. (f, g) Reproduced with permission [1331]. Copyright 2024, American Chemical Society.

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  Organic polymers have been widely studied as electrodes for metal ion batteries. Zhao et al. [1328,1329] first applied them to RCIBs, including polypyrrole (PPy), polyaniline (PANI), and polymerized polyxylylviologen chloride (PXVCl₂). They found that the N⁺ species in polypyrrole chloride (PPyCl) was able to bind to Cl⁻ ions to achieve charge transfer, thus providing a discharge capacity of 118 mAh/g (Figs. 39d and e) [1330]. Recently, 2D metal carbides and nitrides (MXenes) were demonstrated as a new chloride storage material for RCIBs. Yang et al. [1331] synthesized a Ti₃C₂Cl_x MXene with Cl surface terminations by molten salt method as a Cl⁻-ions storage electrode for aqueous chloride ion batteries (Fig. 39f). By coupling with metal chlorides and oxychlorides (AgCl, CuCl, and FeOCl) cathodes, this novel RCIBs enables highly reversible chloride ion storage in NaCl solution and natural seawater with high specific capacity and long cycle life (Fig. 39g). Overall, although numerous advances have been made in RCIBs, many challenges still remain to tackle problems such as capacity fading or low energy density.

  10.2 Fluoride-ion batteries

  The novel secondary battery, which utilizes fluoride ions as carriers to shuttle between the anode and cathode of the battery, is referred to as a fluoride-ion battery (FIB). Currently, the majority of fluoride-ion batteries operate based on the conversion reaction between metals and metal fluorides. The schematic diagram of its discharge reaction and the corresponding equation are as follows:

  Cathode:M1Fn+ne−→M1+nF−

  (13)

  Anode:M2+nF−→M2Fn+ne−

  (14)

  Compared to LIBs, fluoride-ion batteries have remarkable advantages. Firstly, fluorine is abundant in reserves, with a content in the Earth’s crust that is 50 times that of lithium [1332]. Secondly, fluoride-ion batteries exhibit excellent safety, free from issues such as metal dendrite segregation. Additionally, as the element with the strongest electronegativity, fluoride ions (F⁻) possess stable chemical properties and a wide electrochemical stability window, theoretically enabling higher operating voltages. By integrating the attributes of high operating voltage and multi-electron transfer reactions of transition metal compounds, fluoride-ion batteries can attain a high theoretical energy density, thus emerging as a promising new type of battery. Based on the physical state of fluoride-ion electrolytes, they can be categorized into two types: Solid-state fluoride-ion batteries and liquid fluoride-ion batteries. This chapter provides a comprehensive review of research on fluoride ion electrolytes and electrode materials.

  11. Characterization technology

  11.1 Nuclear magnetic resonance

  Compared to the diffraction methods focusing on long-range structure, solid-state nuclear magnetic resonance (ssNMR) is a powerful tool to probe the local chemical structure (through chemical shift, dipole-dipole interaction, and electric quadrupole interaction) and the local electronic/magnetic structure (through hyperfine interaction) of atoms/ions both in crystalline or non-crystalline solids [1333,1334]. Unlike surface-sensitive techniques, NMR can provide quantitative information from the whole sample in a nondestructive manner. Furthermore, the abundant NMR-active nuclei cover the most of constituent elements of battery materials, bringing about highly compatibility between ssNMR and secondary batteries [1335,1336]. Particularly, ^6/7Li and ²³Na as the charge carriers, are the most frequently probed nuclei in the field of LIBs and SIBs. Moreover, the skeletal atoms (including ¹³C, ¹⁷O, ¹⁹F, ²⁹Si, ³¹P, ³³S, ⁵¹V, ⁵⁹Co, ⁹⁵Mo, ¹¹⁹Sn) or doping atoms (including ¹¹B, ²⁵Mg, ²⁷Al, ⁶⁷Zn, ⁷⁷Se), can also provide important information regarding the local structural and electronic properties.

  According to different research objectives, combined with different methods and/or pulse sequences, ssNMR has many important applications in the study of secondary batteries. For paramagnetic materials with strong electron-nuclei dipolar coupling, pjMATPASS (MAT = magic-angle turning, PASS = phase-adjusted sideband separation, pj = projection) technique was frequently utilized to accomplish effective sideband separations and explicit identification of various atom/ion sites [1337,1338]. For nuclei with large quadrupole coupling constant (¹⁴N, ^47/49Ti, ⁹⁵Mo, etc.), WURST-CPMG (WURST = wideband uniform rate smooth truncation, CPMG = Carr−Purcell−Meiboom−Gill) was frequently used for signal acquisition with an ultra-wide spectrum [1339,1340]. ssNMR technique can provide fruitful information on the ion transport/diffusion kinetics by measuring the spin-lattice relaxation time (T₁) and the spin-spin relaxation time (T₂), and through variable-temperature (VT), 2D homonuclear correlation and exchange, and pulsed-field gradient (PFG) experiments [1341,1342]. Of particular, chemical exchange saturation transfer (CEST) has recently been demonstrated as a powerful tool to directly detect the Li exchange across the solid electrolyte interphase (SEI) on metal anodes [1343,1344].

  Moreover, in situ or operando ssNMR has been developed to capture the intermediates and short-lived metastable states or reaction products during electrochemical cycling, unmasking the real-time reaction mechanism inside the secondary batteries [1345,1346]. Magnetic resonance imaging (MRI) is able to map the distribution of chemical species during battery operation with ample spatial and temporal resolution [1347,1348]. Established approaches and new developments in the field of in situ ssNMR and MRI characterization of batteries have been summarized in several review articles [1335,1349].

  In addition to ssNMR, solution NMR was often used to identify different electrolyte decomposition byproducts (ethylene glycol, methanol, etc.) [1350,1351]. Solution NMR has an innate advantage in resolution since anisotropic NMR interactions are averaged out by rapid random flipping in the solution state. Complementary to Raman and FTIR (sensitive to the strong cation-anion interaction), solution NMR is more sensitive to structural changes in the electrolyte induced by weak interactions, and has been intensively used to study the solvent-solvent and solvent-anion interactions and assists in designing better electrolytes on the fundamental level [1352,1353].

  11.2 Transmission electron microscopy

  Understanding the structure-performance relationship of materials is a fundamental topic in the field of functional materials, especially in secondary battery research. The performance of materials is closely related to their microstructure, and conventional characterization methods provide only average structural information [1354,1355], which cannot reveal the diversity of microscopic structures. Transmission electron microscopy (TEM), with its atomic-scale spatial resolution, plays a crucial role in capturing subtle changes in microstructure [1356,1357].

  TEM techniques generally include three major modes: imaging, diffraction, and spectroscopy. In recent years, a variety of imaging techniques, including high-angle annular dark-field (HAADF) [1358,1359], annular bright-field (ABF) [1360,1361], high-resolution TEM (HRTEM) [1362], differential phase contrast (DPC) [1363,1364], and electron holography (EH) [1365,1366], along with diffraction techniques such as electron diffraction (ED) [1367,1368] and spectroscopy techniques like energy-dispersive X-ray spectroscopy (EDS) [1369,1370] and electron energy loss spectroscopy (EELS) [1371,1372], have expanded the application of TEM in secondary battery research. In particular, with the support of in situ TEM [1373,1374], real-time monitoring of the microstructural evolution, reaction kinetics, phase transitions, chemical changes, mechanical stress, and atomic-level structural and compositional evolution at the surface/interface of electrode materials under working conditions is now possible [1375]. This provides precise data for understanding the dynamic behavior and reaction mechanisms of electrode materials, offering new insights for the design and performance optimization of high-performance electrodes.

  This section summarizes the important research progress in applying various TEM techniques to analyze the microstructural dynamics and failure mechanisms of key electrode materials in secondary batteries during charge-discharge processes. This includes in situ and ex situ TEM studies of various cathode materials, high-capacity anode materials, and even all-solid-state batteries, particularly focusing on their microstructural, compositional, and phase changes during electrochemical cycling. Furthermore, the challenges currently faced in TEM characterization techniques are discussed, and the future directions of using TEM to advance secondary battery research are explored.

  (1) Imaging techniqu

  TEM and scanning transmission electron microscopy (STEM) offer a variety of imaging modes to reveal the microstructural features of materials. In TEM, information is typically obtained through amplitude contrast and phase contrast, with sub-angstrom spatial resolution, it can reveal fine structural features such as lattice spacing, orientation, defects, and phase composition. In contrast, STEM mode scans a focused electron beam over the sample and, combined with different imaging modes like HAADF and ABF, can effectively provide detailed information on elemental distribution and local structural details. HAADF is particularly suitable for imaging heavy elements, while ABF is more sensitive to light elements.

  In recent years, emerging techniques such as DPC and EH have further enhanced phase imaging capabilities. DPC allows for precise measurements of electric and magnetic field distributions, providing insights into interface structures, while EH uses interference effects to reveal crystal defects, local electric fields, and stress distributions. These advanced imaging techniques not only enrich our understanding of internal material structures but also play a crucial role in situ observations and dynamic process analyses, offering more comprehensive microstructural information. Chen et al. [1376] investigated metallic lithium deposition and stripping in solid-state batteries using in situ HRTEM, focusing on the mechanical stress generated during electrochemical reactions and how to maintain the system’s stability. Zhou et al. [1377] used HAADF and ABF imaging to study the behavior and atomic structural changes of lithium in different crystal planes during the charging and discharging of MnO@C anode materials. Wang et al. [1363] directly observe lithium-ion accumulation at the electrode/electrolyte interface due to the space charge layer (SCL) effect in sulfide-based all-solid-state LIBs using in-situ DPC. Liu et al. [1364] used in-situ DPC-STEM to investigate the initial activation mechanism of Li_1.2Ni_0.13Co_0.13Mn_0.54O₂ (LLO) in high energy density all-solid-state LIBs. Yamamoto et al. [1378] used EH to directly observe the potential distribution resulting from lithium-ion diffusion in all-solid-state LIBs, particularly the potential variation at the electrode/electrolyte interface. Yang et al. [1379] investigated Li⁺ transport at the interfacial region using EH to reveal the origin of the high Li⁺ transfer impedance in a LiCoO₂(LCO)/LiPON/Pt ASSLIB. These studies have advanced our understanding of lithium-ion dynamics and interfacial behaviors in rechargeable batteries using advanced imaging and characterization techniques.

  (2) Diffraction technique

  Unlike imaging techniques that provide only 2D projected amplitude information of materials, diffraction techniques are based on the de Broglie hypothesis, which suggests that electrons exhibit wave-like properties during propagation. When an electron beam encounters obstacles in the sample, wave interference occurs, leading to diffraction. This interference effect, manifested as discontinuities in spatial distribution, results in electron scattering. By analyzing the scattering patterns, diffraction techniques can provide detailed information on the sample’s phases and structure, enabling precise phase analysis and structural identification.

  Moreover, combining diffraction techniques with in situ electrochemical methods significantly enhances the ability to study dynamic processes. This combination allows for real-time monitoring of structural changes in electrode materials during the charge and discharge cycles, revealing phase transitions, strain evolution, and other critical information. This approach provides a deeper understanding of the material’s behavior during electrochemical reactions. Lewis et al. [1380] studied the (electro)chemical reaction processes at the interface between Li_1.4Al_0.4Ge_1.6(PO₄)₃ (LAGP) electrolyte and lithium using in situ ED and ex situ techniques. Zhou et al. [1381] demonstrate a lithium-enrichment strategy to revive lithium nickel oxide (LNO), a high-energy cathode material (>900 Wh/kg) long plagued by poor cycle performance and thermal instability. Cheng et al. [1382] propose using a chemically inert and mechanically robust boron nitride (BN) film as an interfacial protection layer to prevent the reduction of Li_1.3Al_0.3Ti_1.7(PO₄)₃ (LATP) solid electrolyte by lithium in solid-state Li-metal batteries. Yang et al. [1383] produced air-stable lithium spheres (ASLSs) by electrochemical plating under a CO₂ atmosphere inside an advanced aberration-corrected environmental transmission electron microscopy. These studies collectively demonstrate the power of diffraction techniques in uncovering critical insights into electrochemical processes and advancing battery materials design.

  (3) Spectroscopy technique

  EELS and EDS are two primary techniques widely used for analyzing elemental distribution and chemical composition in materials. The principle of EELS involves the interaction between the incident electron beam and the sample, where a portion of the electrons undergo elastic scattering without energy loss, while another portion experiences inelastic collisions with the atoms in the sample, resulting in energy loss. The greater the energy loss, the larger the scattering angle. These inelastically scattered electrons are collected by an energy analyzer, and a spectrum is generated with energy loss on the x-axis and electron intensity on the y-axis. This spectrum provides valuable information on the material’s electronic structure, chemical composition, oxidation state, and local coordination environment. EELS offers the advantage of high sensitivity to light elements, making it particularly well-suited for analyzing elements such as lithium. In contrast, while EDS also provides elemental composition information, it is primarily used for the analysis of heavier elements.

  Lin et al. [1369] explore the potential of solid-state lithium-metal batteries for electric vehicle applications, particularly the use of solid polymer electrolytes (SPEs) in batteries. Using cryo-electron microscopy imaging and EDS, the authors investigated the structure and chemistry of the interface between lithium-metal and a polyacrylate-based SPE (Fig. 40a). Contradicting conventional knowledge, they found that no protective interphase forms due to sustained reactions between lithium dendrites and polyacrylate backbones and succinonitrile plasticizer. Zhou et al. [1381] prepared a trigonal-structured, slightly Li-enriched LNO (Li_1.04Ni_0.96O₂) material. The study shows that in the slightly Li-enriched LNO prepared by a specially designed molten-salt synthesis method, Ni migration occurs within the layers during delithiation, leading to the formation of vacancy clusters. These vacancies trap the electrochemically oxidized oxygen, effectively suppressing the detrimental effects of lattice oxygen release (LOR). EELS line scan detects the presence of O₂, oxygen vacancies (V_O), and NiO in the Li-enriched LNO, while O₂ was not detected on the surface of conventional LNO (Fig. 40b).

  Figure 40

  

  Figure 40.  (a) HAADF-STEM images and EDS maps of plated Li, indicates the formation of uniform SEI. Reproduced with permission [1369]. Copyright 2022, Springer Nature. (b) HAADF image and EELS line scan, showing the existence of O₂ in the near-surface lattice of LRLNO. Reproduced with permission [1381]. Copyright 2022, Elsevier. (c) Mn, Co, and Ni EELS mapping for a charged NMC particle. (d) Li-concentration EELS mapping of NCA polycrystalline particles during different charge and discharge state. Reproduced with permission [1384]. Copyright 2016, Royal Society of Chemistry. Reproduced with permission [1385]. Copyright 2020, American Chemical Society.

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  Liu et al. [1384] systematically investigate the chemical evolution and structural transformation of the LiNi_xMn_yCo_1-x-yO₂ (NMC) cathode material to understand the deterioration in battery performance, particularly the impact of cathode degradation during electrochemical cycling (Fig. 40c). Using EELS, the study clarifies the role of transition metals in the charge compensation mechanism, especially the behavior of Ni²⁺ (active) and Co³⁺ (stable) ions at different states-of-charge (SOC) under a 4.6 V operating voltage. Nomura et al. [1385] investigated the Li-ion dynamics in bulk-type solid-state Li-ion batteries under low- and high-rate charge–discharge conditions using operando EELS (Fig. 40d). The study found that the diffusion rate of Li ions in bulk-type cathodes is limited, primarily due to the barriers at the nanocrystal grain boundaries, which is the main factor restricting Li conduction.

  Transmission electron microscopy has shown unique advancements and superiority in characterizing the structural evolution of electrodes, as well as the interface structure, composition, and potential distribution in secondary batteries. This not only enriches researchers’ understanding of energy storage mechanisms, interface evolution, and ion transport dynamics in secondary batteries but also provides new methods and ideas for addressing the interface challenges of batteries. However, during high-resolution imaging, the interaction between the electron beam and the electrolyte, alkali metals, and interface phases often causes damage to the sample, making it difficult to characterize all types of secondary batteries at the atomic level, such as solid-state lithium batteries based on sulfide or polymer electrolytes. A traditional method to mitigate electron damage is to reduce the electron dose, but elements with low atomic numbers, such as lithium and oxygen, scatter electrons, leading to a deterioration in the signal-to-noise ratio. Therefore, balancing electron beam damage and signal-to-noise ratio becomes a key issue limiting the broader application of TEM in the study of solid-state lithium battery interfaces. This also presents new challenges and research opportunities, making it an innovative area of work.

  11.3 In-situ spectroscopy

  As demand for more efficient, durable, and sustainable energy storage solutions rises, understanding the fundamental electrochemical processes within these batteries is of great importance. A key approach to gaining deeper insight into these processes is the use of in-situ spectroscopy, a powerful analytical technique that enables real-time monitoring of structural and chemical changes at the atomic and molecular levels during battery operation.

  In the context of rechargeable batteries/capacitors, the extensively used spectroscopic techniques include Fourier-transform infrared (FTIR) spectroscopy, Raman spectroscopy, and ultraviolet-visible (UV–vis) absorption spectroscopy. These spectroscopic techniques can provide valuable spectroscopic data to unveil the electrochemical mechanisms, such as the intercalation/deintercalation of ions, structural transformations, and the evolution of surface/interface interactions [1386–1388]. For instance, Lian et al. [1389,1390] employed in-situ FTIR technique to thoroughly elucidate the energy storage mechanism of the hydrogen-rich carbon nanoribbon (HCNR), that is, the protonated aromatic sp²-hybridized carbon (C(sp²)–H) undergoes a highly reversible rehybridization to sp³-C for efficient Li⁺/Na⁺ ions uptake (Fig. 41).

  Figure 41

  

  Figure 41.  Real-time vibrational evolutions of HCNR during electrochemical processes monitored by in situ FTIR technique. (a) 3D colormap surface plot with contour projection and the corresponding discharge–charge profile at 0.1 A/g in the initial two cycles. (b) 2D contour maps of main peak evolutions with the corresponding second-cycle discharge–charge profile. (a, b) Reproduced with permission [1389]. Copyright 2024, American Chemical Society.

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  In short, in-situ spectroscopy provides in-depth insights into the fundamental behavior of rechargeable batteries, offering a comprehensive understanding of their electrochemical processes and guiding the development of more efficient and durable energy storage technologies. As these in-situ spectroscopic techniques continue to evolve, their application is expected to play a crucial role in advancing battery technologies and meeting the growing energy demands of modern society.

  11.4 In-situ X-ray diffraction

  X-ray diffraction (XRD) is a widely-used technique that enables detailed characterization of various materials, revealing crucial characteristics such as crystal structure, lattice parameters, and the composition of different phases. Its application is particularly significant in the research of battery electrode materials [1391]. Particularly, in-situ XRD involves real-time collection of XRD data during the charging and discharging processes of the batteries, allowing for the monitoring of structural changes in electrode materials under operational conditions [1392,1393]. This investigation is crucial for understanding battery performances such as cycling stability and capacity retention, and it plays a vital role in guiding the optimization of material design.

  The first research of in situ XRD for the study of operando lithium-ion cells dates back to 1978 by Chianelli and co-workers [1394]. The researchers conducted a study on a Li/TiS₂ cell configured in a parallel plate arrangement. During the discharge process, the TiS₂ electrode exhibited topochemical structural changes in its crystal lattice owing to the ordering of lithium ions. Subsequent historical developments of electrochemical cells used for in situ XRD were comprehensively summarized by Tarascon [1395–1400]. As for the information at electrode-solution interface during operando electrochemical processes, in 1985, Fleischmann et al. [1401] investigated the adsorption of iodine on graphite and the formation of hydrides on a nickel electrode using in situ XRD. The pioneering studies established an important foundation for the widespread application of in situ X-ray techniques.

  There are two main types of X-rays, including conventional X-rays and synchrotron X-rays. Compared to conventional X-rays, synchrotron X-rays offers high brightness, exceptional collimation, and a wide tunable energy range, making them particularly valuable for time-resolved studies. Thus, synchrotron XRD can penetrate battery components and collect data with high spatial and temporal resolution, thereby providing more detailed information [1402].

  Qiao et al. [1403] have demonstrated that the nanoscale size of in situ-formed Bi crystals enhances Mg²⁺ diffusion kinetics and significantly improves the efficiency of Mg–Bi alloying and de-alloying processes in a magnesium battery, which is drawn through a combination of operando synchrotron X-ray diffraction and other characterization techniques. Graae and Norby have employed in situ XRD to map phase transformations and redistribution within the anodes of Zn-air batteries. It indicated that the primary challenge for enhancing the performance of secondary cells is the morphological changes in the anode, which restrict access to ZnO nucleation sites and contribute to increased overpotential during discharge [1404]. Villevieille and co-workers have observed lithium polysulfides directly in lithium-sulfur batteries using in-situ XRD, identifying the structural signatures of Li₂S_n species adsorbed onto glass fiber surfaces. This research revealed that only long-chain polysulfides adhere to the glass fiber and persist during repeated lithiation and delithiation cycles [1405]. Dai conducted an operando study using in-situ XRD on the intercalation behavior of chloroaluminate ions with graphite cathodes in aluminum batteries. The results revealed the changes in graphite structure during the charge and discharge processes of aluminum batteries at low temperatures. By comparing theoretical simulations with experimental results, important experimental data and theoretical support was provided for understanding the working principles and performance optimization of aluminum batteries [1406].

  Despite the significant advantages of in situ XRD, it also has certain limitations. XRD is primarily suitable for crystalline materials, however, it is inadequate for amorphous and organic compounds, which are also important components of the solid electrolyte interphase in batteries. Besides, the time resolution of in situ XRD may be insufficient to capture very rapid structural changes, especially during high-rate charge and discharge processes. Additionally, sample preparation and environmental control present challenges, requiring specialized electrochemical cell designs to ensure data quality. To overcome these limitations, integrating multiple characterization techniques, such as in situ TEM and in situ NMR, can provide more comprehensive information.

  11.5 X-ray photoelectron spectroscopy

  X-ray photoelectron spectroscopy (XPS) is one of the important surface analysis methods that not only allows for the investigation of elemental composition, chemical states and concentrations on battery material and electrolyte interfaces, but also enables the study of spatial distributions of elements and chemical states on a microscale, as well as the dynamic evolution of batteries under in-situ conditions [1407,1408].

  Based on the principle of the photoelectric effect, XPS utilizes X-rays to excite samples and generate photoelectrons. By analyzing the emitted photoelectrons, XPS spectra are collected that shows the changes in photoelectron intensity as a function of binding energy. The binding energy of XPS peaks reflects the atomic ionization levels and chemical shifts induced by changes in the chemical environment, providing information on elemental composition, chemical states, and relative concentrations of components and chemical states [1409]. It is worth noting that XPS is one of the few techniques capable of performing chemical state analysis on lithium element [1410,1411]. Spectrum types can be classified into survey spectra and narrow spectra. Survey spectra usually cover a broad energy range (typically around 1100 eV), encompassing the primary characteristic photoelectron peaks of elements. In contrast, narrower spectra have a narrower energy range (typically around 20 eV) with higher energy resolution for chemical state analysis. The typical XPS testing procedure involves first conducting a survey scan of the sample, followed by high-resolution narrow scans targeting specific characteristic peaks in the survey spectrum.

  Conventional XPS often relies on large-area and high-power X-ray sources to enhance signal intensity. However, this method typically yields average information from the surface using a large-area beam spot, as a result it is difficult to accurately reveal the uniform distribution of surface features. As the demand for micro-area XPS analysis has become increasingly significant, the advanced scanning XPS microprobe was developed, which enables the precise focusing of the X-ray beam spot to the micrometer scale, achieving a leap from point analysis to two-dimensional surface analysis. This technology not only captures spatial distribution images of elements but also further reveals the spatial distribution characteristics of chemical states, providing a powerful tool for analyzing micro-area characteristics in battery systems [1412].

  XPS based on Al Kα source typically has a kinetic energy range for emitted photoelectrons between 100 eV and 1000 eV. Due to its small inelastic mean free path, XPS demonstrates extremely high surface sensitivity, capable of probing information within 10 nm of the sample surface. However, in the study of battery materials, material characteristics are not limited to the surface but also involve depth distribution. XPS depth profile is conducted by layer-by-layer removal of surface atoms through sputtering by ion sources such as Ar⁺, Xe⁺, or Ar⁺ GCIB, which is a commonly used method to study the distribution of SEI layer components and chemical states in the depth direction [1413]. It is noteworthy that preferential sputtering maybe happens due to the differences in sputtering yields among different elements, resulting in changes in the chemical state of the sample. For instance, common oxide species in battery materials like CoOx, NiOx, MnOx, and FeOx are prone to preferential sputtering during depth analysis, causing the chemical states of elements such as Co, Ni, Mn, and Fe to be reduced [1414]. To address this issue, hard X-ray photoelectron spectroscopy (HAXPES) technology has emerged. With its greater probing depth, HAXPES enables non-destructive depth analysis, playing a crucial role in the depth analysis of SEI and CEI [1415].

  11.6 Time-of-flight secondary ion mass spectrometry

  Time-of-flight secondary ion mass spectrometry (TOF-SIMS) plays a crucial role in secondary batteries research by virtue of its excellent surface analysis capability. The principle of TOF-SIMS is to bombard the surface of the sample with a primary ion beam, which excites the secondary ions on the surface of the sample, and then inversely deduces the mass-to-charge ratio of the secondary ions according to their time-of-flight to reach the mass analyzer, so as to obtain the elemental composition and chemical composition information of the most superficial layer and chemical composition information. This technique provides spatial resolution at the submicron level and high sensitivity elemental analysis at the ppb to ppm level, which is particularly suitable for the detection of light elements such as lithium, hydrogen, carbon and oxygen.

  TOF-SIMS demonstrates significant advantages in the qualitative and quantitative analysis of the surface composition of battery materials, including the analysis of elements, isotopes, chemical states, and molecular structures, which are critical to understanding the electrochemical behavior of battery materials. The high sensitivity of TOF-SIMS enables the detection of trace impurities and trace elements on the surface, which is of great significance for improving the formulation of battery materials and enhancing the performance of the batteries. This is of great significance for improving the formulation of battery materials and enhancing the performance of batteries. In addition, TOF-SIMS is capable of depth profiling and 3D imaging to reveal the internal structure and interfacial properties of battery materials, such as the formation mechanism of SEI and CEI. In electrode material modification studies, TOF-SIMS explores how surface modification affects battery performance, including the enhancement of cycling stability and multiplicity performance of electrode materials through surface coating or doping. In interfacial characterization, TOF-SIMS analyzes the interfacial reactions between the electrode and the electrolyte, such as the composition and distribution of the SEI film, which is critical for enhancing the safety and extending the lifetime of the battery.

  TOF-SIMS has been widely used in the study of secondary batteries, covering everything from the analysis of the surface composition and elemental distribution of the electrode materials, to the surface modification of the lithium metal negative electrode, to the composition of SEI and CEI membrane composition and interfacial stability studies of solid-state electrolytes. Through these applications, TOF-SIMS technology provides researchers with an important tool for in-depth understanding of the surface and interfacial properties of battery materials, and provides valuable scientific data for optimizing battery design and improving battery performance and safety, thus promoting the advancement of high-performance secondary batteries technology.

  11.7 Synchrotron radiation technology

  To develop high energy battery systems, it is essential to achieve deep understanding for degradation mechanism, such as irreversible phase transformation, micro-crack formation, unstable interphase, and gas evolution. High-brilliance synchrotron X-ray sources with advantages of higher focus, high penetration depth and beam size/energy tunability become powerful technique to reveal the charge compensation mechanism, structural/morphological evolution as well as correlation between the physiochemical properties of battery’s key components with the overall electrochemical performances [1416]. As shown in Fig. 42, multi-model X-ray diagnostic techniques show unique capability for investigating the structural, morphological and elemental evolution of battery components. Among them, X-ray diffraction (XRD) is a versatile technique that can provide information on crystal structure, lattice parameter, phase transformation mechanism, grain size and crystalline orientation. In addition, in situ time-resolved X-ray diffraction (TR-XRD) with ultrafast data collection capability has been widely used to study the dynamic process during charge/discharge.

  Figure 42

  

  Figure 42.  Synchrotron X-ray diffraction, total scattering and absorption techniques applied for battery research and the key information provided by each method.

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  As a typical example, Zhou and coworkers [1417] investigated current rate dependent structural evolution processes of LiNi_1/3Mn_1/3Co_1/3O₂ (NMC333) cathode using TR-XRD and proposed the nonequilibrium phase transition behavior of NMC333 cathode cycling at high current rate. As shown Fig. 43a, NMC333 cathode demonstrated a typical "solid solution-two phase-solid solution" phase transition behavior at low current rate of 0.1 C, while, when the current rate increases to 10 C, a newly formed intermediate phase was clearly observed between two solid solution reaction regions and became more pronounced at higher current rates of 30 and 60 C. By combining multi-scale diagnostic techniques, they proposed that newly formed intermediate phase acts as a buffer between Li-poor and Li-rich phases to reduce the local strain during charging. Mn-based P2-type layered cathode material with reversible oxygen redox have attracted more attention due to high capacity contributed by both TM and Oxygen redox, high voltage and fast kinetics. However, structure degradation induced by irreversible phase transformation, Jahn-Teller distortion and lattice oxygen loss during charging/discharging processes hindered their practical application [1418,1419]. To reveal correlation between composition and phase evolution mechanism of P2-type layered cathodes, Liu et al. synthesized a series of P2-type Na_2/3Mn_xNi_x-1/3Co_4/3–2xO₂ (1/3 ≤ x ≤ 2/3) and systematically investigated the dynamic phase structure transitions using in situ high energy X-ray diffraction (HEXRD). Their result showed that P2-Na_2/3Mn_1/2Ni_1/6Co_1/3O₂ cathode demonstrates stable structure during initial several cycles even at the high voltage without phase transition from P2 to O2, which benefits the long-term cyclic stability. They also utilized ex situ XRD to investigate the phase transformation mechanisms of Na_2/3Mn_xNi_x-1/3Co_4/3–2xO₂ cathodes with x = 5/12, 7/12 and 1/3. It is interesting that, x = 1/3 sample showed a single P2 phase during the initial charging, while, the structure cannot be maintained well at the charging state of 20^th cycle. As for the x = 7/12 sample, a newly formed hydrate phase was observed when it is charged to 4.5 V and the P2 phase almost disappeared after 20 cycles, indicating structural degradation in high voltage region. These results suggested that uniformly distributed sodium ions in the structure effectively suppresses the segregation of vacancy and distortion of structure, thus enhances the long-term cyclic stability even at high voltage operation.

  Figure 43

  

  Figure 43.  (a) In situ XRD of LiNi_1/3Mn_1/3Co_1/3O₂ during the first charge. Contour plot of the (003) diffraction peak of Li_1−xNi_1/3Co_1/3Mn_1/3O₂ with as x is increased between x = 0 and 0.7 during the first charging process at different C rates (0.1, 1, 10, 30, and 60 C). Data were collected at X14A at NSLS with a wavelength of 0.7747 Å. Reproduced with permission [1417]. Copyright 2016, Wiley-VCH. (b) SEI XRD of low concentration electrolytes (LCEs) and low concentration electrolytes (HCEs) using LiFSI as the salt and PC, DMC and DME as the solvents. The light grey pattern belongs to LiF_(SEI). The wavelength used is 0.18323 Å. Reproduced with permission [1420]. Copyright 2021, Springer Nature. (c) The electrochemical lithiation (discharge) and delithiation (charge) curve of an FeOF electrode (ⅰ). (ⅱ) A series of PDF data can be obtained at fine reaction intervals (0.1 Li steps here). Structural models can be fitted to the data to show how (ⅲ) the phases and (ⅳ) particle size evolve during the reaction, or (ⅴ) the changes in the Fe–O and Fe–F features can be evaluated to show how the chemistry changes. Reproduced with permission [1425]. Copyright 2013, American Chemical Society.

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  In addition to investigating the bulk structure of electrode materials, XRD is one of the promising diagnostic tools for analyzing the solid electrolyte interphase (SEI) components quantitatively and qualitatively. Shadike and Hu et al. [1420] used synchrotron based XRD to quantitatively analyzed the crystalline phase of lithium metal anode (LMA) SEI formed in different electrolyte systems. By comprehensively investigating the crystal structures (Fig. 43b), lattice parameters and grain size, they differentiated and quantified two elusive components LiH and LiF in SEI, and proposed possible formation of LiH_xF_1−x solid solution, which has higher ionic conductivity than LiF. More interestingly, they discovered a SEI-LiF, different from the conventional bulk LiF, with grain size around 3 nm, which benefits Li⁺ transport in SEI. These results explained why an ionic insulator compound LiF has been considered as important component of highly stable LMA SEI. More recently, Hu et al. [1421] further utilized XRD to investigated the formation mechanism of LiH, evaluation of SEI during the cell cycling and possible formation mechanism of LiF through anion decomposition in low-concentration electrolytes.

  XRD method provides valuable information of bulk/interphasial properties of electrode materials, however, it can only elucidate the structure properties of crystalline materials and hard to obtain the properties of amorphous phases. Pair distribution function (PDF) analysis, which is a total scattering technique includes both diffuse and Bragg scattering, provides local structure information of battery materials with short-range order, such as nanosized and amorphous organic/inorganic materials as well as the solvation structure of electrolytes [1422]. Taking the unique advantages of PDF analysis, Shadike et al. [1420] further quantitatively investigated the amorphous phases in SEI to reveal the correlation between the electrolyte solvation structure and SEI composition. SEI components formed in low and high concentration electrolytes with different electrolyte solvents have been investigated, while molecular dynamic simulation has been used for demonstrating the electrolyte solvation structures. According to the solvation structure and fitting result of PDF, they proposed that the stability of SEI is highly depending on the electrolyte concentration and solvent properties. The PDF results also directly confirmed formation of alkyl carbonates-dominant and Li₂(FSI(_−F))₂-dominant amorphous phases in low and high concentration electrolyte, respectively. Another good example of PDF application in battery research is investigating the local structure of solid-state electrolyte. Elemental substitution in garnet type solid state electrolyte is an effective approach to enhance the overall conductivity, stability and interface compatibility with anode/cathode. However, it is critical to understand the effects of dopant of substituted elements on the local structure, Li⁺ transport kinetics and electrochemical performance of solid-state electrolyte. Slater et al. [1423] comprehensively investigated the long-range and short-range structures of cerium doped Li₅La₃Nb₂O₁₂ electrolyte coupling XRD and PDF methods. Based on the XRD patterns of series Li_5+xLa₃Nb_2−xCe_xO₁₂ electrolyte, the solid solution limitation is estimated around x = 0.85 and impurity phases occurs when Ce content is further increased. The PDF result further confirmed the incorporation of Ce on the Nb site by analyzing the variation of Nb/Ce-O bond length. Conversion-type electrode materials deliver high energy density due to the multi electron transfer during charging/discharging. However, huge volume change after lithiation/sodiation induces formation of amorphous or nano-sized and highly disordered components, which brings great challenge for understanding the charge storage mechanism and structural properties of conversion-type electrodes. Grey and coworkers [1424] applied operando and ex situ PDF analysis of total scattering data to investigate the changes in local atomic environments during initial sodiation/desodiation processes of card carbon anode. Wiaderek et al. [1425] conducted in situ PDF to measurement using Argonne’s multipurpose in situ X-ray (AMPIX) cell to investigate the reaction mechanism of FeOF (Fig. 43c). The phase evolution, particle size variation and chemical evolution achieved from Rietveld refinement PDF confirmed conversion of FeOF to metallic Fe through rock-salt-type intermediate phase. In addition, the charge compensation mechanism was revealed by analyzing bond length, coordination number and peck intensity of Fe-Fe and Fe-O/F species. To correlated the atomic local structure of sulfurized inorganic-organic hybrid polymer (S-HYB) cathode with its electrochemical properties, Liao et al. [1426] specified the sulfur conversion pathway using ex situ PDF. Their results clearly showed consumption and re-formation of S-S dimer during discharge and charge, respectively. Moreover, irreversible translocation of partial S-S from C-S to P-S sites after initial cycle is another important information provided by PDF analysis, which is also coincident with formation of sulfur trimers via dynamically re-bonding of irreversibly released sulfur species.

  Charge compensation mechanism of electrode materials and surface/interphase evolution during electrochemical cycling are also essential for designing high performance battery systems. X-ray absorption spectroscopy (XAS) is an elemental specific diagnostic technique and able to probe the local atomic arrangement as well as electronic structure of electrode materials and interphases. The XAS can be divided into soft (<2 k eV), tender (2–5 k eV) and hard (>5 k eV) XAS according to energy range (wavelength) of X-ray source. The hard XAS are used to investigate the valance state and coordination environment of various transition metals (TMs) by probing the K-edge feature. In situ hard XAS is commonly used to investigate the charge compensation mechanism and local structural changes of electrode materials in both crystalline and amorphous materials. The valance state and partial local structures of specific elements can be achieved from X-ray absorption near edge structure (XANES), which covers from a few eV below absorption edge to about 50 eV above the edge, while extended X-ray absorption fine structure (EXAFS) provides unique information about the average local structure in terms of inter-atomic distance and the neighboring atomic species [1427].

  As one of the most promising cathodes for LIBs, Ni-rich cathodes exhibit the distinctive advantage of the high energy density. However, critical issues such as oxygen loss and Li/Ni(Ⅱ) intermixing hinder its practical application. To solve these issues, it is of great necessity to investigate the complicated charge compensation mechanisms during the charge/discharge process and further seek for the corresponding strategies. XAS techniques can be exactly used to detect the charge compensation mechanisms through distinguishing the energy variations of the elemental K-edge or L-edge, whose reversible shift can be regarded as the reversible charge compensation process. Yang et al. [1428] proposed a Te-doping strategy to stabilize the lattice oxygen in an ultrahigh-nickel layered LiNi_0.94Co_0.05Te_0.01O₂ cathode and utilized XAS techniques to investigate the influence of Te-doping on the Ni and O charge compensation process. The XANES of Ni K-edge in LiNi_0.94Co_0.05Te_0.01O₂ (NC95T) monotonically shifts to the higher energies during the whole charge process (Fig. 44a), indicating the oxidative process of Ni. However, the energy shift of Ni K-edge is not monotonical in undoped LiNi_0.95Co_0.05O₂ (NC95), which exhibits the lower K-edge energy at 4.6 V charged state than that of doped one, attributed to the oxygen release during the charge process from 4.4 V to 4.6 V. Soft X-ray absorption spectroscopy (sXAS) can distinguish surface and bulk features by probing the chemical species with few nanometers to hundreds of nanometers, inherently sensitive to the element, orbital and different atomic sites in the materials. More importantly, sXAS is mostly used to measure K-edges for low-Z elements (e.g., C, N, O, F) and a more direct probe of the transition metal oxidation states and chemical bonds through the diploe allowed 2p-3d transitions through collecting the L-edge spectra of various metal elements. To verify the oxygen charge variations, Yang et al. further performed sXAS to investigate the electronic structure variation of oxygen by detecting the change of O pre-edge. The intensity ratio of t_2g and e_g exhibits a sharp decrease from 4.4 V to 4.6 V in NC95, indicating the participation of oxygen in charge compensation. Conversely, the relative stable t_2g in NC95T was observed, demonstrating the Te-doping strategy effectively stabilizes the lattice oxygen and suppresses oxygen loss.

  Figure 44

  

  Figure 44.  (a) Ni K-edge XANES of charging state of NC95 and NC95T. Reproduced with permission [1428]. Copyright 2024, Springer Nature. (b) Characterization of the cathode surface by the total electron yield mode of soft XAS: Mn L3 edge, Co L3 edge and Ni L3 edge. The probing depth of soft XAS in total electron yield mode is around 10 nm. Reproduced with permission [1430]. Copyright 2022, Springer Nature. (c) The difference spectra from WT-EXAFS spectra highlight the changes of TM−O and TM−TM bonds when charged to 4 V vs. OCV (top row), and when charged to 4.3 V vs. charged to 4 V (bottom row). Reproduced with permission [1432]. Copyright 2024, American Chemical Society (ACS). (d) The Fourier transformed EXAFS spectra extracted from hard XAS spectra of Fe K-edge at different SOCs in Na_0.6Li_0.1Fe_0.3Mn_0.6O₂ (denoted as Fe_0.3), and Na_0.6Li_0.1Co_0.15Fe_0.15Mn_0.6O₂ (denoted as Fe_0.15Co_0.15). Reproduced with permission [1433]. Copyright 2024, Wiley-VCH.

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  In addition to the bulk stability, interphasial properties Ni-rich NMCs is also important for enhancing cyclic stability. Fu’s group [1429] studied the stability of cathode electrolyte interphase (CEI) and examined the effectiveness of LiPON artificial CEI on interphasial stability of NMC811 cathode cycled in large scale pouch cell. P K-edge XAS spectra collected on bare-NMC811 clearly demonstrated oxidation decomposition of CEI layer upon electrochemical cycling, while LiPON protected interphase was relatively stable without further degradation. For Ni-rich NMCs, formation of rock-salt or spaniel phase are usually accompanied with deteriorated interphase, attributed to the electrochemical performance degradation. sXAS is widely used to examine the surface reconstruction and CEI formation combining with electron microscopic techniques. To reveal the degradation mechanism of NMC76 cathode and understand the role of LiDFP film forming additive, Hu and collaborators [1430] designed multi-model synchrotron X-ray measurement to investigate the bulk and surface properties before and after electrochemical cycles. For the baseline electrolyte without additive, as shown in Fig. 44b, the Mn and Co L3-edge sXAS with total electron yield mode demonstrated reduction of Mn in the beginning of electrochemical cycling, which is gradually intensified during long cycles. Although, Co was more stable than Mn during initial cycles, reduction after 200 cycles was still observed. The Ni L3-edge sXAS spectra showed negligible change even after 200 cycles, probably due to the stability of divalent nickel dominant surface. By strong contrast, Mn, Co and Ni L3-edge sXAS of NMC76 cycled in electrolyte with LiDFP additive show little changes, indicating surface reconstruction and undesired side rection have been effectively suppressed by LiDFP-derived CEI.

  In addition to the information about elemental valence, the EXAFS derived from XAS can also convey the details of local structure variations. In the cathode materials for lithium or sodium-ion batteries, the 3d-orbital of the TMs splits in the octahedral coordination, forming TMO6 octahedra with oxygen. Mn-based oxides are prone to undergo Jahn–Teller local structural distortion due to the asymmetric t_2g³e_g¹ orbital of Mn³⁺, leading to the degradation of the structure and performance during cycling. EXAFS can reflect the changes in Mn−O and Mn−TM, offering an excellent approach to investigate the Jahn-Teller effect. Wang et al. [1431] employed XANES and EXAFS of Mn to explore the local structural distortion of MnO6 at different charge states. EXAFS fitting results indicate two types of Mn−O lengths, 1.86 and 1.99 Å, which are attributed to the Jahn-Teller effect of Mn³⁺. Qiao et al. [1432] compared the variations of Fe−O, Mn−O, and Ni−O in the O3-type NaNi_1/3Fe_1/3Mn_1/3O₂ cathode using wavelet-transformed EXAFS to investigate the different degrees of the Jahn-Teller effect (Fig. 44c). Significant changes in NiO6 at 4.0 V can be observed, while the local structures of FeO6 and MnO6 show minimal variations, indicating the high activity of the NiO6 Jahn-Teller effect. EXAFS provides average variation local structure information of TMO6, which can not only demonstrate the effectiveness of the strategy but also explore the performance degradation mechanisms. Cai et al. [1433] also utilized Fe K-edge EXAFS to investigate the local structure change of P2-type Na_0.6Li_0.1Co_xFe_0.3-xMn_0.6O₂ before and after initial two cycles. As shown in Fig. 44d that π-interaction in CoO6 octahedron effectively mitigates the local structure distortions and phase evolution during electrochemical process, benefits Na⁺ migration.

  12. Recycling spent power batteries

  The burgeoning energy storage industry has precipitated a concomitant surge in the global demand for resources [1434]. Beyond the extraction of metals from primary sources, the recycling of valuable elements from secondary resources have garnered heightened attention, for the sustainable development of energy storage industry [1435,1436]. Despite the existence of a multitude of spent battery types, such as alkaline zinc-manganese batteries, lead-acid batteries, and waste LIBs, recycling technologies can be broadly classified into three categories based on their technical characteristics: pyrometallurgy, hydrometallurgy, and direct repair and regeneration [1437,1438]. The following discussion will provide an overview of the development of these three recycling technologies.

  12.1 Pyrometallurgy

  Pyrometallurgical recovery technology is a pivotal component of spent battery recycling methodologies. Traditional pyrometallurgical techniques typically involve the smelting process, which transforms valuable metals into alloys, while impurity metals are incorporated into the slag phase [1439,1440]. In recent years, solid-phase roasting and pyrolysis pretreatment methods have also been categorized under pyrometallurgical technologies [1441].

  Smelting technologies employed for the recycling of spent batteries typically draw upon the established methodologies used in mineral metallurgy. The lead paste derived from spent lead-acid batteries can be directly processed in a smelting furnace at temperatures exceeding 1000 ℃ through a direct smelting method [1439]. In the case of spent LIBs, the pyrometallurgical recycling process aims to enrich nickel, cobalt, and copper elements within the alloy through high-temperature smelting in a furnace. During this process, manganese and aluminum are incorporated into the slag phase, while lithium is partially captured in the smelting slag and partially in the smelting dust [1440]. The advantages of smelting are simplicity and mature process, for this, the smelting technologies have been adopted by several industrial companies [1442,1443]. However, there are also some disadvantages such as emission of toxic gases and high energy consumption [1443]. Furthermore, it should be noted that since the purity of the product obtained by pyrometallurgy is generally low, it is still necessary to further combine the hydrometallurgical methods. For instance, the alloy and smelting slag produced during the pyrometallurgical processing of spent LIBs require further leaching, separation, and purification steps to yield high-purity products [1444].

  In addition to smelting technologies, several roasting techniques involving salt, carbon, and hydrogen gas as additives have been proposed for the recovery of valuable metals from spent batteries [1445,1446]. The roasting methods are generally designed to enhance the extraction efficiency of subsequent hydrometallurgical processes, and can thus be considered as pretreatment steps for hydrometallurgy. Different extraction effects can be achieved by selecting different additives. By selecting different additives, various extraction effects can be achieved. For instance, when (NH₄)₂SO₄ is used as a roasting additive, it can be mixed with spent LiCoO₂ cathode materials and roasted at 400 ℃, allowing the Co and Li contained in the roasted products to be leached by water [1447]. In contrast, the traditional hydrometallurgical process requires acids and reductants to leach valuable metals from spent LiCoO₂ cathode [1442]. When NaHSO₄, MnSO₄, or CoSO₄ are employed as roasting additives, the subsequent roasting process enables the selective extraction of lithium (Li) from the roasted products using water [1448,1449]. During this roasting process, the lithium phase in spent NCM transitions from Li(Ni_xCo_yMn_z)O₂ to Li₂SO₄, which exhibits high solubility. Concurrently, these transition metals remain in the form of oxides with low solubility. Consequently, lithium can be selectively extracted based on the differential solubility of these compounds.

  Pyrometallurgical technologies are generally characterized by their simplicity and maturity of operation. Due to the requirement for higher reaction temperatures, pyrometallurgical methods typically exhibit higher energy consumption compared to other routes [1450,1451]. Meanwhile, since pyrometallurgical processes often yield products with low purity or intermediate products, the majority of pyrometallurgical processes require integration with hydrometallurgical processes to ultimately achieve the recovery and utilization of valuable metals. Furthermore, the combination of pyrometallurgical and hydrometallurgical processes is advantageous for enhancing the adaptability of the process to various raw materials and improving the recovery efficiency of valuable metals. This integration represents a frontier direction in technological development and holds promise for extensive industrial applications.

  12.2 Hydrometallurgy

  The recycling of spent lithium batteries mainly targets ternary cathode materials and lithium iron phosphate cathode materials. Currently, the main processes are full leaching recovery process and selective lithium extraction recovery process, which include three steps: leaching, purification and material reproduction [1452,1453]. The full leaching hydrometallurgical recovery process can recover lithium in the last step, and can also achieve deep separation and recovery of nickel, cobalt, manganese or ferrophosphorus slag. However, the recovery process of this process is cumbersome, consumes a lot of chemical reagents, and causes lithium loss. The selective lithium extraction recovery process can achieve preferential separation of lithium in the front leaching stage, thereby avoiding the loss of lithium in the back recovery stage after all components of the cathode material are leached, and improving the lithium recovery rate [1454]. The process flow diagram was descripted in Fig. 45.

  Figure 45

  

  Figure 45.  Hydrometallurgical recovery process of spent battery cathode materials.

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  For ternary lithium batteries, commonly used leaching acids include sulfuric acid, hydrochloric acid and phosphoric acid. Co³⁺ and Mn⁴⁺ are difficult to be dissolved by acid solution in the cathode material, so the original agent H₂O₂ is also introduced into the acid solution to reduce its valence state to enhance the leaching effect of Co and Mn [1455], and finally the leaching rate of lithium and nickel can exceed 90%, and the leaching rate of cobalt and manganese is about 70% and 30%, respectively [1456]. Subsequently, impurities are removed by coprecipitation or co-extraction, and the corresponding metal salts are added to adjust the ratio of each metal ion in the solution, and then a precipitant is added to coprecipitate to synthesize the ternary precursor, and finally the precursor is mixed with lithium salt and regenerated into a new cathode material by high temperature solid phase method [1457,1458]. For example, Sun et al. [1459] used low-concentration acid (1 mol/L H₂SO₄) at 50 ℃ and a solid-liquid ratio of 60 g/L for 30 min, and the leaching rates of Li, Ni, Co, and Mn were 94.47%, 99.36%, 99.28%, and 99.56%, respectively. Then, the molar ratio of Ni²⁺-Co²⁺-Mn²⁺ in the solution was adjusted to 6:2:2 by adding trace amounts of NiSO₄⋅6H₂O, CoSO₄⋅7H₂O, and MnSO₄⋅H₂O. Finally, 2 mol/L NaOH (precipitant) and 0.5 mol/L NH₃⋅H₂O (chelating agent) were used to induce the generation of the precursor Ni_0.6Co_0.2Mn_0.2(OH)₂. The precursor was calcined with Li₂CO₃ to synthesize the new cathode material LiNi_0.6Co_0.2Mn_0.2O₂, which has a capacity of 163 mAh/g at a current density of 1 C and maintains a capacity of 85.08% after 300 cycles.

  However, in the hydrometallurgical recovery of lithium iron phosphate batteries, sulfuric acid, hydrochloric acid and formic acid are used as leaching agents, hydrogen peroxide, air and sodium hypochlorite are added as oxidants, and the solid-liquid ratio is controlled at 50–200 g/L to achieve 95%-99% lithium selective preferential leaching [1460–1462]. Fu et al. [1463] used the H₂SO₄-H₂O₂ system to leach spent lithium iron phosphate cathode powder, setting the conditions as liquid-solid ratio of 40 g/L, leaching temperature of 60 ℃, and leaching time of 60 min. The leaching rates of lithium and iron elements reached 98.79% and 94.97%, respectively. The pH value was accurately adjusted by chemical precipitation method, and the products FePO₄ and Li₂CO₃ were recovered as precursors. The raw material with a lithium-iron molar ratio of 1.03:1 was used to regenerate the LiFePO₄/C material. The first discharge capacity at 0.1 C was 160.1 mAh/g, and the Coulombic efficiency of the first cycle was 96.70%. After 100 cycles at 1 C, it can still maintain 99.7% of its good electrochemical performance.

  At present, the research hotspot of leaching process is the selection and use of green and efficient leaching agents. The main leaching agents are inorganic acids and organic acids. Organic acids have a milder reaction system than inorganic acids, but the higher price of organic acids will lead to higher processing costs [1464]. At the same time, since the leaching process of the two recovery processes consumes a large amount of acid reagents, reducing agents or oxidizing agents can be added during acid leaching to enhance the leaching effect, thereby reducing the consumption of acid reagents [1465,1466].

  In the separation stage, the main difficulties are the removal of impurities and the deep separation of metals. For example, elements such as magnesium and manganese will reduce the initial capacity of the subsequent electrode materials [1467], while copper impurities will cause battery self-discharge, and aluminum impurities are inert substances that are difficult to achieve Li⁺ deintercalation and will reduce the specific capacity of the material. The improvement method is mainly to improve the separation accuracy, such as screening high-efficiency selective precipitants and optimizing liquid phase conditions. Zhang et al. [1468] used N₂H₄·H₂O as a reducing agent to cyclically leach copper, and the precipitation rate of Cu was 96%. Finally, Cu was reduced to Cu₂O and entered the precipitate residue. At the same time, Al was selectively precipitated from the leaching solution using C₇H₅NaO₂, with a precipitation rate of 96%, and Al entered the precipitate in the form of Al(C₇H₅O₂)₃. During the extraction process, the metal cations M²⁺ (M−Ni and Co) are solvated to form stable hydrated ions M(H₂O)₆²⁺. However, Ni(H₂O)₆²⁺ and Co(H₂O)₆²⁺ show different activities when reacting with the acidic extractant (HA), forming different complexes NiA₂·2H₂O and CoA₂. Yuan et al. [1469] used this extraction mechanism to develop a method to reduce solvation (addition of lactic acid), which increased the maximum separation factor of Ni and Co by 192 times to 345, opening up a new path for the efficient and selective extraction of Co.

  In the process of material reproduction, the electrochemical performance of the material can be improved by synergistic doping and interface modification. Zhang et al. [1470] used the impurity element Al and introduced the slow diffusion element Ti to upgrade the spent LiCoO₂ cathode to high-voltage LiCoO₂, enhancing its cycle stability and rate performance at a high cut-off voltage (4.6 V). Therefore, material reproduction can adjust the morphology of the material by controlling the synthesis conditions and using impurity elements to make it have better electrochemical performance [1471].

  With the development of new batteries, the recycling technology of new batteries such as solid-state batteries and sodium-ion batteries has gradually attracted attention, and there have been reports on the recycling of solid electrolytes [1472,1473]. Kirstin Schneiderdeng et al. [1474] used inorganic acids (H₂SO₄, HCl) and organic acids (HCOOH, CH₃COOH, C₂H₂O₄, C₆H₈O₇) to completely or selectively dissolve LLZO, and proposed a two-stage leaching process using water and sulfuric acid to achieve a total leaching efficiency of 100% for Li and Al, 95% for La, 94% for Zr, and 91% for Ta. For hydrometallurgical process recycling of solid-state batteries, the solid electrolyte contains many types of elements and the materials are very complex, and the element behaviors vary greatly, resulting in low material recovery rate and very high separation costs during the recycling process [1475]. The recycling of sodium-ion batteries has also attracted the attention of some researchers. Chen et al. [1476] recovered cathode materials and all-solid-state electrolytes from spent all-solid-state sodium-ion batteries using deep eutectic solvents (DES). The results show that the highest leaching efficiency of Na in all-solid-state sodium-ion batteries cathode materials and all-solid-state electrolytes using DES is 88.3% and 56.9%, respectively. At present, hydrometallurgical recycling is still the best choice for battery recycling.

  12.3 Direct recycling

  Direct recycling technology is a closed-circuit strategy that can regenerate spent electrodes into fresh ones within a short process with less consumption of chemical reagents and energy [1477], which is an environment-friendly and economical technology for the recycling of spent anode and cathode, including hydrothermal method [1478], eutectic-salts methods [1479], and electrochemical method [1480], and solid-state methods [1481]. By using a simple flash joule heating technology, Chen et al. [1482] recovered waste graphite anode within seconds. The solid-electrolyte-interphase-derived carbon shell is coated on the initial graphite structure so that the flash-recycled anode reveals high initial specific capacity, superior rate performance, and cycling stability (Fig. 46a). The recovered graphite shows a high capacity of 351 mAh/g at 0.2 C, along with excellent cycling stability and rate performance when compared with the anode heated by traditional calcination. Zhang et al. [1483] proposed a continuously rolled-over heating method to separate and regenerate spent electrodes (Fig. 46b). Compared with the hydrometallurgical methods, transient recycling has significant advantages in energy consumption, time efficiency, and material purity. This approach generally applies to various anodes from the market, making it practical to directly recycle waste batteries with high economic and environmental benefits. The direct recycling technology also plays a vital role in the cathode recycling. Zhou et al. [1484] proposed a direct recycling method under mild conditions, whose innovation is the use of uniformly precoating lithium sources at room temperature and atmospheric pressure (Fig. 46c). The waste LiNi_0.83Co_0.12Mn_0.05O₂ with serious structure damage is repaired to a high-purity cathode with a high initial capacity of 181.6 mAh/g. Formulating some novel strategy to directly recover cathodes with different degradation levels has attracted much attention with the rapid development of direct recycling technology. Lv et al. [1485] developed a photocatalytic-assisted recovery method based on the different decommissioning degrees of cathode materials, showing rapid reaction rates, high recovery efficiency, and low-cost recovery (Fig. 46d). They repaired degraded LiCoO₂ in a lithium solution assisted with UV leaching. The repaired cathode delivered a high capacity of 179 mAh/g at 0.1 C. Overall, direct recycling technology as one of the recycling technologies shows great application potential in terms of its low energy consumption, high efficiency, and mild environmental effects.

  Figure 46

  

  Figure 46.  (a) Schematic diagram of flash recycling of waste anode. Reproduced with permission [1482]. Copyright 2023, Wiley-VCH. (b) Schematic of transient recycling of spent graphite from copper foils via rolled-overheating. Reproduced with permission [1483]. Copyright 2023, Royal Society of Chemistry. (c) Mechanism of the direct electrochemical leaching method for spent LIBs. Reproduced with permission [1484]. Copyright 2024, Wiley-VCH. (d) Recycling process flow of the spent lithium-ion battery. Reproduced with permission [1485]. Copyright 2023, American Chemical Society.

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  Declaration of competing interest

  The authors declare no conflict of interest.

  CRediT authorship contribution statement

  Gaojing Yang: Writing – review & editing, Writing – original draft, Funding acquisition, Data curation, Conceptualization. Zhimeng Hao: Writing – review & editing, Writing – original draft, Visualization, Conceptualization. Chun Fang: Writing – original draft, Data curation. Wen Zhang: Writing – original draft, Data curation. Xia-hui Zhang: Writing – original draft, Data curation. Yuyu Li: Writing – original draft, Data curation. Zhenhua Yan: Writing – original draft, Data curation. Zhiyuan Wang: Writing – original draft, Data curation. Tao Sun: Writing – original draft, Data curation. Xiaofei Yang: Writing – original draft, Data curation. Fei Wang: Writing – original draft, Data curation. Chengzhi Zhang: Writing – original draft, Data curation. Hongchang Jin: Writing – original draft, Data curation. Shuaifeng Lou: Writing – original draft, Data curation. Nan Chen: Writing – original draft, Data curation. Yiju Li: Writing – original draft, Data curation. Jia-Yan Liang: Writing – original draft, Data curation. Le Yang: Writing – original draft, Data curation. Shouyi Yuan: Writing – original draft, Data curation. Jin Niu: Writing – original draft, Data curation. Shuai Li: Writing – original draft, Data curation. Xu Xu: Writing – original draft, Data curation. Dong Wang: Writing – original draft, Data curation. Song Jin: Writing – original draft, Data curation. Bo-Quan Li: Writing – original draft, Data curation. Meng Zhao: Writing – original draft, Data curation. Changtai Zhao: Writing – original draft, Data curation. Baoyu Sun: Writing – original draft, Data curation. Xiaohong Wu: Writing – original draft, Data curation. Yuruo Qi: Writing – original draft, Data curation. Lili Wang: Writing – original draft, Data curation. Nan Li: Writing – original draft, Data curation. Bin Qin: Writing – original draft, Data curation. Dong Yan: Writing – original draft, Data curation. Xin Cao: Writing – original draft, Data curation. Ting Jin: Writing – original draft, Data curation. Peng Wei: Writing – original draft, Data curation. Jing Zhang: Writing – original draft, Data curation. Jiaojiao Liang: Writing – original draft, Data curation. Li Liu: Writing – original draft, Data curation. Ruimin Sun: Writing – original draft, Data curation. Zengxi Wei: Writing – original draft, Data curation. Xinxin Cao: Writing – original draft, Data curation. Kaixiang Lei: Writing – original draft, Data curation. Xiaoli Dong: Writing – original draft, Data curation. Xijun Xu: Writing – original draft, Data curation. Xiaohui Rong: Writing – original draft, Data curation. Zhaomeng Liu: Writing – original draft, Data curation. Hongbo Ding: Writing – original draft, Data curation. Xuanpeng Wang: Writing – original draft, Data curation. Zhanheng Yan: Writing – original draft, Data curation. Guohui Qin: Writing – original draft, Data curation. Guanghai Chen: Writing – original draft, Data curation. Yaxin Chen: Writing – original draft, Data curation. Ping Nie: Writing – original draft, Data curation. Zhi Chang: Writing – original draft, Data curation. Fang Wan: Writing – original draft, Data curation. Minglei Mao: Writing – original draft, Data curation. Zejing Lin: Writing – original draft, Data curation. Anxing Zhou: Writing – original draft, Data curation. Qiubo Guo: Writing – original draft, Data curation. Wen Luo: Writing – original draft, Data curation. Xiaodong Shi: Writing – original draft, Data curation. Yan Guo: Writing – original draft, Data curation. Longtao Ma: Writing – original draft, Data curation. Xiangkun Ma: Writing – original draft, Data curation. Jiangjiang Duan: Writing – original draft, Data curation. Zhizhang Yuan: Writing – original draft, Data curation. Jiafeng Lei: Writing – original draft, Data curation. Hao Fan: Writing – original draft, Data curation. Jinlin Yang: Writing – original draft, Data curation. Chao Li: Writing – original draft, Data curation. Tong Zhou: Writing – original draft, Data curation. Jiabiao Lian: Writing – original draft, Data curation. Jin Zhao: Writing – original draft, Data curation. Huanxin Ju: Writing – original draft, Data curation. Tinglu Song: Writing – original draft, Data curation. Zulipiya Shadike: Writing – original draft, Data curation. Weiguang Lv: Writing – original draft, Data curation. Jiawei Wen: Writing – original draft, Data curation. Lingxing Zeng: Writing – original draft, Data curation. Jianmin Ma: Writing – review & editing, Funding acquisition, Conceptualization.

  Acknowledgment

  This work was supported by the National Natural Science Foundation of China (Nos. U21A20311 and 22409147).

  
  
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• 

  Figure 1  Research advances in the field of rechargeable batteries.

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  Figure 2  (a) Crystal structure of lithium-rich manganese-based materials. (b) Typical first-cycle charge-discharge curves of lithium-rich manganese-based cathode materials. (c) Compositional phase diagram showing the electrochemical reaction pathways for a xLi₂MnO₃·(1-x)LiMO₂ electrode. (c) Reproduced with permission [87]. Copyright 2021, Wiley-VCH.

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  Figure 3  Design of DRX and DRXPS cathodes. (a) The partially disordered spinel phase as an intermediate between the DRX and fully ordered spinel structures. (b) Schematic of the structural evolution of Li_1.2Mn_0.65Ti_0.15O_1.9F_0.1 with each synthesis step. (c) The structure of M₂O₄, M_2−uO₄ and M_2−u[XO₄]_xO_4(1−x). (a-c) Reproduced with permission [115–117]. Copyright 2024, Springer Nature.

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  Figure 4  Schematic diagram of Rudörff model and Daumas-Herold model of lithium storage process in graphite.

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  Figure 5  (a) The structures and the calculated HOMO/LUMO energies of the solvents. Reproduced with permission [217]. Copyright 2023, Elsevier. (b) Schematic illustration of the reinforced mechanism of the LiTFA−LiNO₃ dual-salt additive on conventional carbonate electrolyte. Reproduced with permission [218]. Copyright 2024, Wiley-VCH. (c) Schematic illustration of the mechanism on the stabilized 4.6 V LCO batteries. Reproduced with permission [219]. Copyright 2024, Royal Society of Chemistry.

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  Figure 6  (a) Schematic of the contrast of Li⁺ insertion process between LHCE (1.5 mol/L LiFSI DME-BTFE) and HCE (4.5 mol/L LiFSI DME) in graphite layers. (b) The electrostatic potential maps distributions of DX and DOL. (c) Density functional theory calculation of the binding energy of Li⁺-DX, Li⁺-DOL, Li⁺-FSI⁻ and Li⁺-TFSI⁻. (d) Ionic conductivity and t_Li⁺ of LD (1.0 mol/kg LiFSI in DX), LDD (1.0 mol/kg LiFSI in the mixture of DX and DOL), and LTDD (0.75 mol/kg LiFSI and 0.25 mol/kg LiTFSI in the mixture of DX and DOL) electrolytes. (a) Reproduced with permission [225]. Copyright 2020, Wiley-VCH GmbH. (b-d) Reproduced with permission [227]. Copyright 2024, Elsevier.

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  Figure 7  Phosphate-based electrolyte. Flame-retardant test in the (a) blank electrolyte without PFPN (flammable), (b) PFPN-based electrolyte (non-flammable). (c) Cycling performance of Na||Na₃V₂(PO₄)₂O₂F cells in PFPN-based electrolyte. (d) Schematic illustration of interfacial chemistry and corresponding solvation sheath of the Tri-anion regulated phosphate-based electrolyte. (e) Long-term cycling of NCM||Li cell in CFTN electrolyte at 0.1 C. (f) Schematic of TDNE electrolyte solvation structure. (g) Cycling performance of Li||NCM 622 batteries using TDNE electrolytes at 1 C. (h) Long-term cycling of the HC||Na₄Fe_2.91(PO₄)₂(P₂O₇) pouch cell at 1 C. (a-c) Reproduced with permission [232]. Copyright 2024, Wiley-VCH. (d) Reproduced with permission [233]. Copyright 2024, Elsevier. (e) Reproduced with permission [234]. Copyright 2022, Wiley-VCH. (g) Reproduced with permission [235]. Copyright 2023, American Chemical Society. (h) Reproduced with permission [236]. Copyright 2024, American Chemical Society.

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  Figure 8  Ionic liquid-based electrolyte for Li based batteries: (a) The frontier molecular orbital energy level of different kinds of ionic liquids. Reproduced with permission [275]. Copyright 2021, Elsevier. (b) Solvation structure of ionic liquid-based electrolyte. Reproduced with permission [276]. Copyright 2020, Wiley-VCH. (c) Schematic illustration of the interphase of Li||LiCoO₂ full cell with novel ionic liquid electrolyte. Reproduced with permission [277]. Copyright 2020, Wiley-VCH. (d) Schematic illustration of the interphase of NCM811 cathode with ionic liquid electrolyte. Reproduced with permission [278]. Copyright 2024, Wiley-VCH. (e) Schematic illustration of Li||NCM811 full cell in the ionic liquid modified electrolyte. Reproduced with permission [279]. Copyright 2023, Wiley-VCH. (f) Synthesis procedures of difluorinated ionic liquid for Li metal batteries. Reproduced with permission [280]. Copyright 2024, Wiley-VCH.

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  Figure 9  (a) Schematic illustrations of the synthetic process for the Li-HA-F nanofibers and Li-HA-F CSE with lithium conductive mechanism and Radar plots which compare the comprehensive performance of nongel SPEs with different fillers. (b) Surface SEM image of the Li-HA-F membrane. (c) Comparison of typical lithium superionic conductors with the Li-HA-F nanofiber. (d) Mechanical strength of the Li-HA-F CSE and PEO-LiTFSI (inset: the optical image of the Li-HA-F CSE at bending state). (e) LSV curves for the Li-HA-F CSE and HA CSE. (f) Binding energy calculation of the (002) surface of HA with and without F doping. (g) Schematic diagram of the design principle for the EDTA-pH PSE. (h) Cross-section SEM image of the EDTA-pH electrolyte. (i) Stress-strain curves of the EDTA-pH and pH electrolyte membranes. (j) Stress and single-edge crack simulations of the EDTA-pH electrolyte. (k) The calculated binding energy of the Li⁺ or LiTFSI^- with the nearest solvent molecules (for the pH electrolyte) or polymer chain (for the EDTA-pH electrolyte). (l) Schematic illustration of the vapor-phase fluorinated approach. (m) The cross-section scanning electron microscope (SEM) image of the F-NBR-g-VEC matrix and corresponding elemental mapping images of C, F, and O (inset: 3D rendering of C-, F-, and O- secondary ion distributions of F-NBR-g-VEC electrolyte). (n) The 3D rendering of LiF^-, Li⁺, and O^- secondary ion distributions on the cycled lithium anode with the F-NBR-g-VEC electrolyte. (o) Schematic illustration of ultraconformal interface between lithium anode and the F-NBR-g-VEC electrolyte. (a-f) Reproduced with permission [286]. Copyright 2024, American Chemical Society. (g-k) Reproduced with permission [287]. Copyright 2024, Wiley-VCH. (l-o) Reproduced with permission [290]. Copyright 2023, Wiley-VCH.

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  Figure 10  Schematic of the intertwined phenomena inducing the impaired reaction activity and the fast degradation of the sulfur cathode and Li metal anode during cycling. Reproduced with permission [317]. Copyright 2023, Wiley-VCH.

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  Figure 11  Classification diagram of carbon material as cathode of lithium-sulfur batteries. Reproduced with permission [332]. Copyright 2023, Elsevier.

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  Figure 12  (a) The synthesis flowchart, SEM diagram and cycle performance of S/TiO₂@rGO. Reproduced with permission [338]. Copyright 2023, Elsevier. (b) The synthesis flowchart, SEM diagram and cycle performance of TiO₂-VO_x cathode. Reproduced with permission [337]. Copyright 2016, American chemical Society.

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  Figure 13  Schematic diagram of the solvation structures of lithium ions and lithium polysulfides in different electrolyte systems for Li–S batteries.

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  Figure 14  (a) SEM images of graphene air electrode with hierarchical and interconnected structure. Reproduced with permission [441]. Copyright 2015, Wiley-VCH. (b) Schematic illustration for the fabrication of free-stand CNTs air electrode. Reproduced with permission [442]. Copyright 2017, Elsevier. (c) Schematic illustration of the cable-type Li-O₂ batteries. Reproduced with permission [446]. Copyright 2018, Wiley-VCH. (d) In situ observation of the morphological evolution of the discharge product with CNTs-based electrodes. Reproduced with permission [449]. Copyright 2021, Wiley-VCH.

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  Figure 15  (a) Schematic illustration of the essential difference in electrochemical reactions on carbon-based and carbon free cathode surface. (b) Schematic illustration of the change in local electronic structure at Li_xO₂ molecular orbitals upon reactions on the surface of metal alloy cathode. (c) First charge−discharge curves of six types of metal-based cathodes at a current density of 1.0 A/g and a specific capacity limit of 3000 mAh/g. (d) Predicted catalytic activities of typical catalyst including metal oxide based on the established correlation of O₂ desorption and charging voltage with surface acidity. (a, b) Reproduced with permission [455]. Copyright 2023, Wiley-VCH. (c) Reproduced with permission [458]. Copyright 2017, American Chemical Society. (d) Reproduced with permission [460]. Copyright 2015, American Chemical Society.

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  Figure 16  Proposed mechanism for the decomposition of ether-based solvents during ORR process in non-aqueous Li-O₂ batteries. Reproduced with permission [505]. Copyright 2011, Wiley-VCH.

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  Figure 17  Work principles of RT Na-S batteries.

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  Figure 18  The charge/discharge profiles of Na_4/7[□_1/7Mn_6/7]O₂ cathode upon the 2nd cycle with the voltage range of 1.5–4.7 V. Reproduced with permission [613]. Copyright 2018, Wiley-VCH.

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  Figure 19  (a) The structure of Na₂Fe₂(SO₄)₃ projected along the c axis; and (b) local environment of two independent Fe sites. Green octahedra, yellow tetrahedra and blue spheres show FeO₆, SO₄ and Na, respectively. (c) Galvanostatic charging and discharging profiles of Na_2−xFe₂(SO₄)₃ cathode cycled between 2.0 V and 4.5 V at a rate of C/20 (2Na in 20 h). (d) Schematic illustration of the heterostructure and (e) Na⁺ transfer process in the heterostructure. (f) The structure of Na₂Fe(SO₄)₂ along the c-axis. (g) Cyclic voltammograms for the 1^st to 3^rd cycle for Na₂Fe(SO₄)₂/C at a scan rate of 0.1 mV/s and (h) galvanostatic charge–discharge profiles of Na₂Fe(SO₄)₂ and Na₂Fe(SO₄)₂/C at 0.1 C cycled between 1.5 V and 4.2 V. (i) Na₂Fe(SO₄)₂-HC pouch cell cycling performance. (a-c) Reproduced with permission [619]. Copyright 2014, Springer Nature. (d, e) Reproduced with permission [621]. Copyright 2023, Springer Nature. (g, h) Reproduced with permission [622]. Copyright 2019, Royal Society of Chemistry. (i) Reproduced with permission [623]. Copyright 2025 Elsevier.

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  Figure 20  Sodium storage sites in disordered carbon.

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  Figure 21  (a) Schematic diagram of the advantages of Fe-M-HoMS as anode for SIBs. (b) The rate performance of Fe-M-HoMS. (c) Schematic diagram of the phase transition mechanism of MoS₂. (d) The rate performance of 1T-P-MoS₂. (e) Schematic diagram of the excellent electrochemical performance of the MoS₂/NiS₂ heterostructure. (a, b) Reproduced with permission [741]. Copyright 2024, Wiley-VCH. (c, d) Reproduced with permission [746]. Copyright 2023, Wiley-VCH. (e) Reproduced with permission [745]. Copyright 2024, Elsevier.

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  Figure 22  (a) Electrochemical potential window stability and temperature range of 1 mol/L NaClO₄ dissolved in different ester solvents. Reproduced with permission [765]. Copyright 2012, Royal Society of Chemistry. (b) Comparison of binding energy (ΔE_b), solvation enthalpy (ΔH_sol) and solvation-free energy (ΔG_sol) for different ester solvents. Reproduced with permission [766]. Copyright 2015, American Chemical Society.

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  Figure 23  (a) Schematic illustration of synthesis and reversible K⁺ insertion/extraction of K₂TP. (b) Schematic illustration of morphological evolution of the Bi electrode in the ether-based electrolytes during cycles. (c) Three adsorption models of DME molecule on (012) crystal plane of Bi based on DFT calculation: Bridge, Top and Hollow. (d) Molecule structure of HFME in Na⁺ solvation configuration. (e) XPS spectra of C 1s, F 1s, and S 2p on KMO surface. (f) Cycle stabilities of K_0.67MnO₂ at 4.5 V in the traditional carbonate-based electrolyte and the designed electrolyte. (a) Reproduced with permission [772]. Copyright 2017, Royal Society of Chemistry. (b) Reproduced with permission [773]. Copyright 2017, Wiley-VCH. (c) Reproduced with permission [774]. Copyright 2018, Wiley-VCH. (d) Reproduced with permission [777]. Copyright 2024, Wiley-VCH. (e, f) Reproduced with permission [778]. Copyright 2024, Wiley-VCH.

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  Figure 24  (a) Electrolyte flammability test for conventional carbonate electrolyte and phosphate-based Electrolyte. (b) Cycling performance of HC||NVP using phosphate-based electrolytes with and without VC. (c) Projected density of states of the 3.3 mol/L NaFSI/TMP electrolyte. (d) Cycling performance and Coulombic efficiency of the Na||HC half-cells using concentrated 3.3 mol/L NaFSI/TMP electrolyte and conventional 1 mol/L NaPF₆/EC: DEC (1:1, v/v) electrolyte [957]. (e) Energy level diagrams of TFEP and TEP. (f) The SEM and TEM images of the cycled HC electrodes the TMP-based (left) and TMP/TFEP-based (right) electrolyte. (g) Photo of the fully charged HC||NFPP pouch cells after package clipping and under a flame. (a) Reproduced with permission [784]. Copyright 2021, Elsevier. (b) Reproduced with permission [783]. Copyright 2022, Royal Society of Chemistry. (c, d) Reproduced with permission [790]. Copyright 2018, Springer Nature. (e) Reproduced with permission [795]. Copyright 2018, Elsevier. (f, g) Reproduced with permission [236]. Copyright 2024, American Chemical Society.

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  Figure 25  Summary about the mechanism, suppression strategies, and effective utilizations and prospects of the Jahn–Teller effect in sodium - ion batteries [848,849,855–857].

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  Figure 26  Summary about the mechanism, suppression strategies, and effective utilizations and prospects of the Jahn–Teller effect in sodium - ion batteries [863,870,875,876].

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  Figure 27  (a) Defect-free crystal structure with 1/4 of the vacant sites occupied by water (green spheres) and 1/3 of the empty sites. (b) TEM of KFeHCF-E nanoparticles. (c) The corresponding charge/discharge and color-mapped curves for in situ XRD patterns for different angle ranges. (d) Schematic crystal structure and (e) galvanostatic charge-discharge voltage profiles of KMF-EDTA sample. (f) In situ substitution of Mn with Fe in KMnF. (g) The 40th charge-discharge curves of KMnF in the pure and modified electrolytes. (h) The ultralong cycling stability of the KMnF electrode in the modified electrolyte. (a) Reproduced with permission [879]. Copyright 2020, Springer Nature. (b, c) Reproduced with permission [881]. Copyright 2023, Wiley-VCH. (d, e) Reproduced with permission [882]. Copyright 2021, Springer Nature. (f-h) Reproduced with permission [883]. Copyright 2021, Springer Nature.

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  Figure 28  Schematic illustration of the SEI evolution for Co-NBC@BP@PCA. Reproduced with permission [904]. Copyright 2024, Wiley-VCH.

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  Figure 29  Synthesis steps for Bi/Bi₂O₃ NDs@CSs. Reproduced with permission [913]. Copyright 2023, Wiley-VCH.

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  Figure 30  Structure and electrochemical performance characterizations of CuO@CF, Co₃O₄@N-C, and SnO₂@C. CuO@CF: (a) SEM image. (b) Cycling performance at 50 mA/g. (c) Cycling performance at 1 A/g. Co₃O₄@N-C: (d) SEM image. (e, f) Cycling performance at 0.05 and 0.5 A/g, respectively. SnO₂@C: (g, h) Cycling performance at 0.05 and 0.5 A/g, respectively. (a-c) Reproduced with permission [928]. Copyright 2022, Elsevier. (d-f) Reproduced with permission [929]. Copyright 2020, American Chemical Society (ACS). (g, h) Reproduced with permission [930]. Copyright 2022, Elsevier.

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  Figure 31  (a) TEM images of graphite anodes from K/graphite cell with EC/DEC (left) and DEGDME after 20 cycles (right). Reproduced with permission [949]. Copyright 2020, Elsevier. (b) Molar concentration of different electrolytes versus ionic conductivity at 28±2 ℃. Reproduced with permission [957]. Copyright 2018, Royal Society of Chemistry. (c) Schematic diagrams of graphite electrodes with different intercalation behaviors in weakly or strongly ion-solvent interacting electrolytes and the corresponding solvated structures. Reproduced with permission [960]. Copyright 2022, Wiley-VCH. (d) Cycling performance of Gr//KPTCDA full cells at −20 ℃ and 0.1 C with different electrolytes. Reproduced with permission [961]. Copyright 2023, Wiley-VCH. (e) Linear sweep voltammetry curves of three electrolytes. Reproduced with permission [962]. Copyright 2024, Oxford University Press.

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  Figure 32  Challenges of Zn metal anodes encounter during electrochemical cycling.

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  Figure 33  Strategies used to address the challenges of Zn metal anodes.

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  Figure 34  (a) An overview of electrochemical properties for MnO₂ with single electron transfer reaction and two electrons transfer reaction. (b) An overview of electrochemical properties for vanadium oxides and vanadium phosphates.

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  Figure 35  (a) The anode interface of the conventional commercial Li-ion battery. (b) Anode interface of ALIBs. (c) Ion pairs in WIS electrolyte. (d) Schematic of WIS aqueous lithium-ion battery. (e) Important time point in the development of ALIBs. (c, d) Reproduced with permission [1116]. Copyright 2015, American Association for the Advancement of Science.

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  Figure 36  Schematic illustration of cathode [883,1147–1149,1155–1157], anode [1150,1151,1159–1162] and electrolyte [1158] of AKIBs in recent years.

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  Figure 37  (a) A schematic of the Grotthuss mechanism, in which proton conduction is operated by rearranging bonds along a water chain. (b) A schematic of Newton’s cradle. (c) Schematic structures of a defective Prussian blue analogue (PBA). (d) GCD curves at various current rates of defective Prussian blue analogue electrode. Inset: relationship between the hysteresis of the potential profiles and the applied current rate. The polarization between the charge and discharge profiles increases linearly with the current rate, suggesting that the rate performance of CuFe-PBA is limited more by the testing cells’ electrical resistance than the proton transport and the reaction kinetics (e) GCD curves at various current rates of H_1.75MoO₃ electrode. (a-d) Reproduced with permission [1163]. Copyright 2019, Springer Nature. (e) Reproduced with permission [1164]. Copyright 2022, American Chemical Society.

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  Figure 38  Solvation chemistry and Zn anode/electrolyte interfaces of (a) traditional aqueous electrolyte, (b) anion-receptor aqueous electrolyte and (c) lean-water aqueous electrolyte and physicochemical properties of different electrolytes. (d) Schematic illustration of symmetric anion zinc salt and asymmetric anion zinc salt. (e) The dominated solid-liquid transition temperature of KOH solutions with different C_KOH and the phase composition of KOH solution at different temperatures and concentrations. (d) Reproduced with permission [1246]. Copyright 2024, Wiley-VCH. (e) Reproduced with permission [1249]. Copyright 2023, Elsevier.

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  Figure 39  (a) Schematic illustration of the preparation process of FeOCl/CMK-3. (b) Galvanostatic curves of the FeOCl/CMK-3. (c) Crystal structures of CoFe-Cl and Ni₂V_0.9Al_0.1-Cl LDHs. (d) Cycling performance of PPyCl. (e) Schematic of the chloride ion storage mechanism in PPyCl. (f) HAADF-STEM images and atomic structural models of the Ti₃C₂Cl_x. (g) Schematic of the charge storage mechanism of the seawater-based ACIBs. (a, b) Reproduced with permission [1324]. Copyright 2017, American Chemical Society. (d, e) Reproduced with permission [1330]. Copyright 2017, American Chemical Society. (f, g) Reproduced with permission [1331]. Copyright 2024, American Chemical Society.

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  Figure 40  (a) HAADF-STEM images and EDS maps of plated Li, indicates the formation of uniform SEI. Reproduced with permission [1369]. Copyright 2022, Springer Nature. (b) HAADF image and EELS line scan, showing the existence of O₂ in the near-surface lattice of LRLNO. Reproduced with permission [1381]. Copyright 2022, Elsevier. (c) Mn, Co, and Ni EELS mapping for a charged NMC particle. (d) Li-concentration EELS mapping of NCA polycrystalline particles during different charge and discharge state. Reproduced with permission [1384]. Copyright 2016, Royal Society of Chemistry. Reproduced with permission [1385]. Copyright 2020, American Chemical Society.

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  Figure 41  Real-time vibrational evolutions of HCNR during electrochemical processes monitored by in situ FTIR technique. (a) 3D colormap surface plot with contour projection and the corresponding discharge–charge profile at 0.1 A/g in the initial two cycles. (b) 2D contour maps of main peak evolutions with the corresponding second-cycle discharge–charge profile. (a, b) Reproduced with permission [1389]. Copyright 2024, American Chemical Society.

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  Figure 42  Synchrotron X-ray diffraction, total scattering and absorption techniques applied for battery research and the key information provided by each method.

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  Figure 43  (a) In situ XRD of LiNi_1/3Mn_1/3Co_1/3O₂ during the first charge. Contour plot of the (003) diffraction peak of Li_1−xNi_1/3Co_1/3Mn_1/3O₂ with as x is increased between x = 0 and 0.7 during the first charging process at different C rates (0.1, 1, 10, 30, and 60 C). Data were collected at X14A at NSLS with a wavelength of 0.7747 Å. Reproduced with permission [1417]. Copyright 2016, Wiley-VCH. (b) SEI XRD of low concentration electrolytes (LCEs) and low concentration electrolytes (HCEs) using LiFSI as the salt and PC, DMC and DME as the solvents. The light grey pattern belongs to LiF_(SEI). The wavelength used is 0.18323 Å. Reproduced with permission [1420]. Copyright 2021, Springer Nature. (c) The electrochemical lithiation (discharge) and delithiation (charge) curve of an FeOF electrode (ⅰ). (ⅱ) A series of PDF data can be obtained at fine reaction intervals (0.1 Li steps here). Structural models can be fitted to the data to show how (ⅲ) the phases and (ⅳ) particle size evolve during the reaction, or (ⅴ) the changes in the Fe–O and Fe–F features can be evaluated to show how the chemistry changes. Reproduced with permission [1425]. Copyright 2013, American Chemical Society.

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  Figure 44  (a) Ni K-edge XANES of charging state of NC95 and NC95T. Reproduced with permission [1428]. Copyright 2024, Springer Nature. (b) Characterization of the cathode surface by the total electron yield mode of soft XAS: Mn L3 edge, Co L3 edge and Ni L3 edge. The probing depth of soft XAS in total electron yield mode is around 10 nm. Reproduced with permission [1430]. Copyright 2022, Springer Nature. (c) The difference spectra from WT-EXAFS spectra highlight the changes of TM−O and TM−TM bonds when charged to 4 V vs. OCV (top row), and when charged to 4.3 V vs. charged to 4 V (bottom row). Reproduced with permission [1432]. Copyright 2024, American Chemical Society (ACS). (d) The Fourier transformed EXAFS spectra extracted from hard XAS spectra of Fe K-edge at different SOCs in Na_0.6Li_0.1Fe_0.3Mn_0.6O₂ (denoted as Fe_0.3), and Na_0.6Li_0.1Co_0.15Fe_0.15Mn_0.6O₂ (denoted as Fe_0.15Co_0.15). Reproduced with permission [1433]. Copyright 2024, Wiley-VCH.

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  Figure 45  Hydrometallurgical recovery process of spent battery cathode materials.

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  Figure 46  (a) Schematic diagram of flash recycling of waste anode. Reproduced with permission [1482]. Copyright 2023, Wiley-VCH. (b) Schematic of transient recycling of spent graphite from copper foils via rolled-overheating. Reproduced with permission [1483]. Copyright 2023, Royal Society of Chemistry. (c) Mechanism of the direct electrochemical leaching method for spent LIBs. Reproduced with permission [1484]. Copyright 2024, Wiley-VCH. (d) Recycling process flow of the spent lithium-ion battery. Reproduced with permission [1485]. Copyright 2023, American Chemical Society.

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