Recent Progress and Perspectives of Solid State Na-CO₂ Batteries

Abstract

Solid state Na-CO₂ batteries are a kind of promising energy storage system, which can use excess CO₂ for electrochemical energy storage. They not only have high theoretical energy densities, but also feature a high safety level of solid-state batteries and low cost owing to abundant sodium metal resources. Although many efforts have been made, the practical application of Na-CO₂ battery technology is still hampered by some crucial challenges, including short cycle life, high charging potential, poor rate performance and lower specific full discharge capacity. This paper systematically reviews the recent research advances in Na-CO₂ batteries in terms of understanding the mechanism of CO₂ reduction, carbonate formation and decomposition reaction, design strategies of cathode electrocatalysts, solid electrolytes and their interface design. In addition, the application of advanced in situ characterization techniques and theoretical calculation of metal–CO₂ batteries are briefly introduced, and the combination of theory and experiment in the research of battery materials is discussed as well. Finally, the opportunities and key challenges of solid-state Na-CO₂ electrochemical systems in the carbon-neutral era are presented.

1. Introduction

Metal–CO₂ batteries (such as Li/Na/Zn/K-CO₂ batteries) are a high energy density energy storage and power supply technology that enables CO₂ fixation and conversion [1,2]. Among various kinds of metal–CO₂ batteries, Na-CO₂ batteries have attracted more attention because of the abundant resources of sodium metal and similar physical and chemical properties to lithium. Na-CO₂ batteries exhibit better comprehensive performance, including high energy density (1.13 kWh kg⁻¹), and relatively high working voltage (2.35 V) [3]. In addition, the abundance of metal sodium is 1352 times that of metal lithium, so the cost is lower (the price of metal sodium is 20-times cheaper than that of metal lithium) [4,5]. In addition, a Na-CO₂ battery with the reaction of 4 Na + 3 CO₂ ↔ 2 Na₂CO₃ + C (∆_rG⁰_m = −905.6 kJ mol⁻¹) has a low reaction Gibbs free energy, which means that Na-CO₂ batteries may have a low charging voltage and inhibit electrolyte decomposition, which is conducive to improving round-trip efficiency and prolonging service life [6,7]. Compared with Li⁺, Na⁺ as a charge carrier has other advantages, for example, sodium has a larger ionic radius and atomic mass than lithium, but the Stokes radius of sodium is smaller than lithium, so it has higher mobility and ionic conductivity, resulting in a smaller polarization [8]. Therefore, Na-CO₂ batteries are widely considered to be a promising next generation energy storage power supply technology. However, the research of Na-CO₂ batteries is still in its infancy, and there are problems, for example, the reaction mechanism is still unclear, the types of cathodic materials and catalytic activity are relatively limited, the ionic conductivity and interfacial stability of the solid electrolyte still need to be improved, the impedance between the electrode and the interface is large and dendrites grow at the metal anode interface during the reaction. These problems lead to unsatisfactory electrochemical performance, such as cycling stability, overpotential, rate capability and specific full discharge capacity of Na-CO₂ batteries [9,10,11]. To solve these problems, it is necessary to comprehensively understand the reaction mechanism of Na-CO₂ battery and clarify the possible origins of the issues.

2. Mechanism of Na-CO₂ Electrochemistry

In 2011, Asaoka’s group found that adding CO₂ to Li/Na-O₂ batteries could increase its discharge capacity and energy density, demonstrating the feasibility of metal–CO₂ batteries for the first time [12]. In 2013, Das et al. first designed a rechargeable O₂-assisted Na-CO₂ battery [13]. Under different partial pressures of CO₂, the discharge capacity of rechargeable Na-CO₂ (O₂) battery with tetraethlene glycol dimethyl ether (TEGDME) and ionic liquid (IL) electrolyte increased by 2.6 times and 2.1 times, respectively (Figure 1b,c), the optimal CO₂/O₂ ratio of the maximum discharge capacity ranges from 40% to 70% (Figure 1c) [13]. The ex situ Fourier transformed infrared (FTIR) spectroscopy test and X-ray powder diffraction (XRD) analysis showed that Na₂CO₃ and Na₂C₂O₄ exist simultaneously in TEGDME-based electrolyte; Na₂C₂O₄ is the main discharge product in IL electrolyte. These results proved that the electrolyte solvent may determine the final discharge product by affecting the intermediate stability.

In 2016, Hu et al. first proposed and demonstrated the electrochemical reaction mechanism of rechargeable Na-(pure) CO₂ batteries:

3CO₂ + 4Na ↔ 2Na₂CO₃ + C

(1)

which is inconsistent with the conjecture proposed by Das that Na₂C₂O₄ or Na₂CO₃ and CO is obtained in pure CO₂ atmosphere [6]. The Na-CO₂ battery was designed based on NaClO₄/TEGDME electrolyte and treated multi-walled carbon nanotube (MWCNT) cathode, which was successfully cycled 200 times in a pure CO₂ atmosphere, and the reaction mechanism was verified by a series of tests including XRD and X-ray photoelectron spectroscopy (XPS). Combined with the research results of Li-CO₂ system similar to Na-CO₂ system [14], it is reasonable to speculate that the reaction mechanism of rechargeable Na-pure CO₂ system is as follows: CO₂ molecules directly capture e⁻ to form C₂O₄²⁻, the unstable C₂O₄²⁻ undergoes a two-step disproportionation reaction to form CO₃²⁻ and C, and finally the discharge products Na₂CO₃ and C are produced as follows:

2CO₂ + 2e⁻ → C₂O₄²⁻

(2)

C₂O₄²⁻ → CO₂²⁻ + CO₂

(3)

C₂O₄²⁻ + CO₂²⁻ → 2CO₃²⁻ + C

(4)

CO₃²⁻ + 2Na⁺ → Na₂CO₃

(5)

The schematic diagram of Na-CO₂ battery structure is shown in Figure 2. Its electrochemical reaction route can be proposed as the following chemical equations:

Cathode reactions: 4Na⁺ + 3CO₂ + 4e⁻ → 2Na₂CO₃ + C

(6)

Anode reactions: Na → Na⁺ + e⁻

(7)

Overall reaction equation: 4Na + 3CO₂ → 2Na₂CO₃ + C

(8)

In addition, a rechargeable Na-O₂ (CO₂) battery with an IL-propylene carbonate-based electrolyte supplemented with 10% SiO₂ nanoparticles was reported by Archer et al. in 2014 [15]. The discharge product was detected as NaHCO₃, while CO₂ and O₂ were released during charging. Although it was speculated that the H in NaHCO3 might come from the introduction of trace amounts of H₂O during electrolyte preparation, there was no sufficient experimental evidence to prove it. Obviously, the exploration of the Na-CO₂ reaction mechanism is rather limited and further studies are needed.

3. Cathode Material/Catalysts of Na-CO₂ Batteries

The charging and discharging process of Na-CO₂ batteries is accompanied by the adsorption and desorption of CO₂, as well as the deposition and decomposition of the insulated discharge product Na₂CO₃ on the cathode surface. Inadequate reaction kinetics is the main obstacle, resulting in large overpotential, poor reversibility, poor rate performance and poor cycling stability, etc. [16]. Therefore, exploring an efficient, stable and low-cost electrocatalyst to facilitate CO₂ reduction and carbonate decomposition is one of the key issues in the development of this technology. The following issues should be considered during the research of cathode materials/catalysts [17,18,19]: to realize proper CO₂ adsorption properties, reasonably designed porous and macroporous structures are essential elements for designing cathode materials, so as to facilitate Na⁺ and CO₂ diffusion, reduce the activation energy of the rate-controlling step and accommodate the insulation discharge product Na₂CO₃. In addition, the selection of efficient catalytic materials and the design of rich catalytic sites are the key factors to reduce the overpotential in the discharge process and improve the electrochemical performance. Finally, the factors such as abundant raw material resources, environmentally friendly preparation process and easy preparation are the prerequisites for the practical application of Na-CO₂ batteries. To realize the commercialization of Na-CO₂ batteries, the exploration of efficient catalytic materials is of broad significance. In the past decades, Na-CO₂ batteries have achieved rewarding results, especially in the design of efficient electrocatalysts. In this section, we review and discuss the research progress of cathode catalysts based on their chemical composition and microstructure from the three major categories: carbon and heteroatom-doped carbon materials, metal-loaded composite catalytic materials and single-atom catalysts, and then analyze their structure-activity relationship and the overall performance of Na-CO₂ batteries

3.1. Carbon Materials and Heteroatom-Doped Carbon Materials

Carbon materials have been widely used in various electrochemical energy storage devices, especially metal–O₂ batteries and metal–CO₂ batteries, due to their high electronic conductivity, large specific surface area, stable chemical and electrochemical properties, controllable pore structure and adjustable surface chemistry (defect engineering and heteroatom doping) [20,21,22,23,24]. Super P, Ketjen black (KB), activated carbon, carbon nanotubes, graphene, metal–organic framework materials (MOFs) and nitrogen-doped porous carbon materials have been widely used as cathodes for Li-CO₂ batteries, and they are also very suitable for Na-CO₂ batteries [13,15,25].

3.1.1. Commercial Carbon Materials

Commercial carbon materials can be used as cathode materials for metal–CO₂ batteries due to their good electronic conductivity, large surface area, relatively chemical stability, low cost, as well as their mature and scalable preparation process. Common commercial carbon materials, such as Ketjen black (KB) [12] and Super P [26,27], have been developed as porous cathode materials for Li-CO₂ batteries. However, the electrochemical performance of commercial carbon materials is not ideal owing to the inherent defects of low electrical conductivity, small pore volume, relatively small specific surface area and limited active sites [28].

Super P is first applied to the metal–CO₂ battery as the cathode material. Notably, the electrolyte plays a key role in determining the battery performance when the electrode materials are the same. As reported by Archer et al. Super P has almost no discharge capacity when it was used as a cathode material in ionic liquid electrolytes [26]; while Yang et al. applied a Super P cathode in combination with an ether-based electrolyte to a Li-CO₂ battery, showing a greatly enhanced discharge capacity at 100 mA g⁻¹, reaching 6062 mAh g⁻¹ (further increase in discharge capacity with the addition of Ru metal, Ru@Super P, 8229 mAh g⁻¹) [27]. The reason for the significant difference in discharge capacity was ascribed to the electrolyte solvent could affect the stability of the intermediate discharge products and the formation mechanism of the final discharge products, thus the final discharge products and discharge capacities differed with different electrolytes as reported in the earliest studies on Na-CO₂(O₂) batteries with Super P as the cathode material and ionic liquid electrolytes or ether-based electrolytes [13]. Therefore, the selection of a suitable electrolyte and cathode material plays a key role in determining the battery performance. Advanced characterization methods can be used to track these intermediate species and theoretical calculations can also further provide sufficient evidence.

3.1.2. Nanocarbon Materials

In addition to the above commercial activated carbon materials, nanostructured carbon materials (such as carbon nanotubes and graphene) are more widely explored in Na-CO₂ batteries [6,29,30]. They not only have novel structures, high electronic conductivity and high specific surface area, but also have excellent physical and chemical properties due to their unique quantum size effects and surface chemical states, so they have excellent electrochemical activity [31].

Chen’s group prepared activated multi-walled carbon nanotube (a-MWCNT) cathode with three-dimensional tri-continuous porous structure (Figure 3a), high electronic conductivity and good wettability to electrolyte by boiling MWCNT with TEGDME at 100 °C and coating it on Ni network, which effectively improved its reactivity and reduced electrochemical polarization [6]. The rechargeable Na-CO₂ battery assembled with this material as cathode electrode catalyst and ether-based liquid electrolyte has a maximum reversible capacity of 60,000 mAh g⁻¹ at 1 A g⁻¹, can be cycled 200 times and the initial discharge/charge voltage difference is only 0.6 V, and gradually increases to 1.3 V after 200 cycles (Figure 3b). Transmission electron microscopy (TEM) images (Figure 3b) shows grape-like discharge nanoparticles (50 nm) were randomly deposited on the t-MWCNT surface during the discharge process, and the grape-like discharge products disappeared after charging (Figure 4c), clearly indicating that the reaction is extremely reversible. A series of electrochemical test methods such as selected area electron diffraction (SAED), Raman spectroscopy (Figure 3d), XPS (Figure 3e), electron energy loss spectroscopy (EELS) (Figure 3f) similarly demonstrated the discharge products of Na-CO₂ batteries as Na₂CO₃ and C. Subsequently, this group reported solid-state Na-CO₂ batteries via acid-treated MWCNTs cathodes and NaF-modified anode [29]. The successful introduction of -COOH and -OH groups (Figure 3h) by acid treatment of MWCNTs before use can improve the adsorption capacity of CO₂ and thus the reaction kinetics of CO₂ electrochemical reduction was enhanced.

In general, carbon materials have remarkable characteristics such as large specific surface area, rich surface chemical properties, high intrinsic electronic conductivity, high chemical and electrochemical stability and low cost, and are the most commonly used cathode materials in Na-CO₂ batteries [17,29]. Although the incorporation of new nanostructured carbon materials such as graphene and carbon nanotubes has greatly improved the discharge capacity of Na-CO₂ batteries, their catalytic activity and resulting rate performance and cycling performance are still unsatisfactory, and require effective strategies, such as heteroatom doping to are needed to adjust the microstructure, porosity, defects and surface charge distribution.

3.1.3. Heteroatom-Doped Carbon Materials

It is widely accepted that the activation of CO₂ at the early stage of discharge depends mainly on the interaction of CO₂ with the chemically inert surface of the carbon material, i.e., the adsorption/desorption of CO₂. In the field of Li-CO₂ batteries, it has been demonstrated that heteroatom doping can modulate the charge distribution of nanostructured carbon material, strengthen nearby positively charged carbon atoms [32]. This has a significant impact on the adsorption mode and adsorption energy between the gas molecules (CO₂) and the carbon material surface, thus significantly enhancing the gas reduction kinetics and promoting the decomposition kinetics of the solid discharge products on the cathode surface of the metal–CO₂ batteries [20,33,34].

N element doping is an important strategy for catalyst design [35]. Due to the high electronegativity of nitrogen, N doping can break the charge neutrality of the carbon skeleton and cause charge redistribution, which gives the material excellent electronic and catalytic properties. For example, Hu et al. obtained nitrogen-doped carbon cathode nanomaterials with unique structures by calcining zeolite imidazolium salt framework (ZIF-8) at a certain temperature and then washing with dilute hydrochloric acid [25]. This material has higher CO₂ absorption and CO₂ adsorption properties, as well as better cycling stability. As shown in Figure 4b, the CO₂ absorption of the nitrogen-doped sample is much higher than that of carbon black, but its SSA is relatively small. According to the calculated results (Figure 4d), the surface of N-doped has stronger interaction with CO₂ bond, which can promote the reduction in CO₂ and the formation of discharge products. The electrochemical impedance spectrum shows that the resistance of the cell with the optimized nitrogen-doped nanocarbon (NC900) cathode increases only slightly after 80 cycles, which is much better than that of the cell with the carbon black cathode (about a six-fold increase after seven cycles), as shown in Figure 4c. Their assembled solid-state Na-CO₂ batteries Na || liquid-free PEO-based polymer electrolyte || optimized nitrogen-doped nanocarbon (NC900) cathode material exhibited better electrochemical performance: lower overpotential at 50 °C, higher discharge capacity of 10,500 mAh g⁻¹, the energy density of 180 Wh kg⁻¹ and stable cycling for 320 h (at a capacity of 1000 mAh g⁻¹), as shown in Figure 4a. In addition, X-ray photoelectron spectroscopy (XPS) revealed that the signal of CO₃²⁻ was detected after the NC900 cathode was discharged and disappeared after charging, confirming the reversible formation and decomposition of Na₂CO₃ (Figure 4e). These excellent properties are attributed to the efficient catalytic effect of porous and highly conductive N-doped nanocarbons. Furthermore, the optimized nitrogen-doped nanocarbons facilitate the formation of sheet-like discharge products, which are easily decomposed into CO₂ after charging.

In 2019, Sun’s group prepared nitrogen-doped single-walled carbon nanotubes (N-SWCNH) catalytic materials with a unique structure [36]. The excellent electrocatalytic performance of metal-free N-SWCNH for CO₂ reduction is mainly attributed to the unique structure of single-walled carbon nanotubes and nitrogen doping. The porous nature and unique conical structure of SWCNH provide sufficient storage space for discharge products; the highly dispersed nitrogen doping provides a large number of structural defect sites for CO₂ adsorption and electron transfer, which contributes to the electron affinity and CO₂ adsorption/desorption ability and improves catalyst activity and reversibility. They successfully prepared refillable hybrid Na-CO₂ batteries using N-doped single-walled carbon nanotubes (N-SWCNH) as the cathode catalyst and Na superionic conductor (NASICON) solid electrolyte as the separation medium for the hybrid electrolyte system. The use of aqueous electrolyte is also beneficial to the dissolution of discharge products, greatly improving the electrochemical reaction kinetics. As shown in Figure 4f,g, compared with the highly promising N-MWCNTs and Au NPs catalytic materials, N-SWCNH exhibited better performances and the prepared Na-CO₂ hybrid battery not only exhibited a low discharge/charge voltage difference of 0.49 V at a current density of0.1 mA cm⁻¹ (Figure 4f). It provides a high discharge capacity of 2293 mAh g⁻¹ at a cut-off voltage and a current density of 0.2 mA cm⁻², as shown in Figure 4g. It can be cycled for more than 100 times at a current density of 0.1 mA cm⁻¹ (Figure 4h).

Figure 4. (a) Schematic of the structure of an all-solid-state Na-CO₂ cell with N-doped carbon cathode. (b) Low-pressure CO₂ adsorption isotherm of N-doped nanocarbon and carbon black at 273 K. The adsorption amount at 1 atm is shown in parentheses. (c) Nyquist plots of the NC900 cathode and carbon black cathode of the Na-CO₂ batteries in different charge states. (d) Density function theory (DFT) calculations of the interaction of undoped, graphitic N-doped and pyridine N-doped nanocarbon with CO₂ molecules and the adsorption of one CO₂ molecule on undoped and doped nanocarbon binding energy calculations. (e) XPS characterization of the discharged products. Reproduced with permission from [25]. Copyright 2020, American Chemical Society. (f–h) Electrochemical performance of Na-CO₂ batteries. (f) Discharge–charge voltage curves of the cells with N-MWCNTs, AuNPs and N-SWCNH as catalysts at 0.1 mA cm⁻² current density. (g) Discharge capacity curves at a current density of 0.2 mA cm⁻². (h) Rate performances at different current densities. Reproduced with permission from [36]. Copyright 2020, Elsevier Ltd.

Figure 4. (a) Schematic of the structure of an all-solid-state Na-CO₂ cell with N-doped carbon cathode. (b) Low-pressure CO₂ adsorption isotherm of N-doped nanocarbon and carbon black at 273 K. The adsorption amount at 1 atm is shown in parentheses. (c) Nyquist plots of the NC900 cathode and carbon black cathode of the Na-CO₂ batteries in different charge states. (d) Density function theory (DFT) calculations of the interaction of undoped, graphitic N-doped and pyridine N-doped nanocarbon with CO₂ molecules and the adsorption of one CO₂ molecule on undoped and doped nanocarbon binding energy calculations. (e) XPS characterization of the discharged products. Reproduced with permission from [25]. Copyright 2020, American Chemical Society. (f–h) Electrochemical performance of Na-CO₂ batteries. (f) Discharge–charge voltage curves of the cells with N-MWCNTs, AuNPs and N-SWCNH as catalysts at 0.1 mA cm⁻² current density. (g) Discharge capacity curves at a current density of 0.2 mA cm⁻². (h) Rate performances at different current densities. Reproduced with permission from [36]. Copyright 2020, Elsevier Ltd.

In addition to N element doping, density function theory (DFT) calculations show that F doping and B doping also improve the performance [37]. In addition, multi-element co-doping, N/S co-doping [38] and N/B co-doping [39] are also used. Cheng et al. adopted the S/N co-doping strategy to improve the physical and chemical properties of the cathode material [40]. They designed electrophilic S vacancies and nucleophilic N-doped active centers on the surface of ReS₂, and used the synergistic coupling effect between heteroatoms to adjust the interaction of the catalyst with Li atoms and C/O atoms to show suitable adsorption during charging and discharging processes of Li-CO₂ batteries, respectively, thus reducing the activation energy of the rate-determining step and thus increasing the reaction rate.

In summary, the design of three-dimensional (3D) structures and heteroatom doping is an effective strategy to facilitate the diffusion of reactants and improve the catalytic activity of carbon nanomaterials for CO₂ reduction and Na₂CO₃ decomposition. However, the optimal design for carbon materials cannot yield the desired high catalytic activity. Therefore, an in-depth analysis of the active center of heteroatom doping and the related CO₂ reduction mechanism with the help of advanced characterization is beneficial for exploring the methods to achieve uniform doping and controlled preparation.

3.2. Metal-Loaded Composites

Despite the demonstrated applicability of nitrogen-doped carbon materials, the catalytic activity for the reversible reaction of the cathode in Na-CO₂ batteries is quite limited [41]. Therefore, various carbon loaded metal composites were designed as cathodes for activity modulation and thus to improve the performance of Na-CO₂ batteries. The results show that metal-loaded composite catalysts appear to be a good choice for reducing the battery charging potential and improving the electrochemical performance of the battery compared to pure carbon materials [38,42,43,44,45,46,47].

3.2.1. Precious Metals and Their Composites

Due to the inherent electronic configuration of half-filled anti-bonds and high electrical conductivity of noble metals, noble metals and their composites usually have the advantages of good adsorption, low resistance and low overpotential as electrocatalysts [2,48].

Among all precious metal elements, ruthenium (Ru) is one of the most widely studied precious metal catalysts [27,49]. For example, Guo et al. obtained Ru@KB composites by in situ reductions of RuCl₃ on porous Ketjen Black [16]. As shown in Figure 5a–c, compared with pure KB as the cathode of Na-CO₂ battery, Ru@KB can significantly improve the discharge capacity up to 11,537 mAh g⁻¹, cycle stability for more than 130 cycles and coulomb efficiency of 94.1% of Na-CO₂ battery. Their electrochemical performance and ex situ characterization confirmed that ruthenium nanoparticles can significantly reduce the charging overpotential, promote the reversible reaction between Na₂CO₃ and carbon, and further improve the cycling stability of Na-CO₂ batteries. Besides the excellent catalytic activity of metallic Ru, its corresponding oxide RuO₂ has also been shown to have significant catalytic ability for reversible metal–CO₂ batteries. In addition, Pt [50], Ag, Ir [51,52], Au [53], Pd [54] and other noble metals and their oxides have been shown to be very effective in bifunctional catalysis for metal–CO₂ batteries, and need further be explored.

Compared with pure carbon-based materials, the composite by loaded with noble metal-based catalysts help to reduce the active energy required for CO₂ reduction and decomposition of discharging products, exhibit excellent catalytic activity in facilitating the electrochemical reaction of Na-CO₂ batteries, allowing for smaller overpotentials and higher energy efficiency, and significantly increase the discharge/charge capacity of the batteries. Unfortunately, the expensive cost and limited resources seriously hinder the commercial application of noble metal-based catalysts. Therefore, reasonable strategies such as developing noble metal-based single-atom catalysts or other inexpensive and abundant transition metal-based catalysts should be explored to meet their industrial development.

3.2.2. Transition Metals and Their Composites

Transition metals Ni [21], Co [43,55], Mn [56], Cu [57] and Fe [58] are supported on carbon-based materials with high specific surface area and high electronic conductivity due to their unique adjustable structure and multivalent characteristics, providing rich active sites for electrochemical reactions. In recent years, single metal composites, alloy type composites and transition metal oxide composites of transition metals have been widely reported. Moreover, due to the advantages of rich reserves and low cost of transition metals, transition metal-based composites are a feasible solution for future controlled scale up production [17].

In 2020, Xu et al. obtained an efficient active material (Co/Co₉S₈ @SNHC) for hybrid system Na-CO₂ batteries by anchoring Co/Co₉S₈ active nanoparticles on biomass-derived S and N-doped graded porous carbon via a microporous/mesoporous domain-limited synthesis strategy, as shown in Figure 5d [45]. The Na-CO₂ battery with Co/Co₉S₈ @SNHC not only exhibits a low overcharge potential of ~0.32 V and a charge/discharge voltage difference of only 0.65 V (Figure 5e). As shown in Figure 5f, it also shows better rate performance, cycle stability (cycled for more than 200 cycles at a current density of 0.1 mA cm⁻¹) and higher specific area discharge capacity (~18.9 mA cm⁻²). These excellent electrochemical properties are attributed to the meso- and mesoporous-limited domains of the biomass carbon skeleton, which not only effectively inhibit the agglomeration of Co/Co₉S₈ nanoparticles, but also provide diffusion channels for CO₂ and Na⁺ as well as sufficient space for storing the discharge products. In addition, the effective synergistic interactions between the effective catalytically active sites (Co/Co₉S₈, C-N, C-S bonds) and the defect-rich carbon interfaces (S, N doping) similarly prevent the agglomeration and separation of Co/Co₉S₈ nanoparticles, enhance the catalytic activity and improve the stability.

In addition to monometallic composites, bimetallic composites are also potential candidates for CO₂ reduction reaction (CO₂RR) and CO₂ electroreduction reaction (CO₂ER). In 2021, Xu et al. obtained bimetallic nitrogen-doped carbon materials of Fe-Cu-N-C with dense bimetallic active sites as catalysts for Na-CO₂ batteries with mixed air by introducing Fe³⁺ and Cu²⁺ regulated in situ pyrolytic growth of carbon nanotubes via solid-phase reactions [58]. They suggested that the excellent electrocatalytic activity of Fe-Cu-N-C is attributed to the synergy between the N-doped carbon framework with more defects and a large number of active sites in the Fe-N_x, Cu-N_x and Fe/Fe₃C nanocrystals. Moreover, Fe³⁺ is the key to catalyze the conversion of g-C₃N₄ to CNT conformation, while Cu²⁺ gives the carbon nanotubes a good structure and uniform diameter. As shown in Figure 5g,h, the Fe-Cu-N-C materials synthesized at a pyrolysis temperature of 700 °C exhibit an ultra-low voltage gap of 0.44 V and a cycle efficiency of 83.2%, as well as a large discharge capacity of 8411 mAh g⁻¹ and a long-term cycling performance of 1550 cycles (over 600 h).

Figure 5. Na-CO₂ battery cycling behavior of KB and Ru@KB cathode at 200 mA g⁻¹: (a) KB cathode. (b) Ru@KB composite cathode. (c) Discharge–charge profiles of Na-CO₂ batteries with KB and Ru@KB composite cathodes at a current density of 100 mA g⁻¹ in the first cycle. Reproduced with permission from [16]. Copyright 2019, Royal Society of Chemistry. (d) Schematic illustration of catalytic cathode of hybrid Na-CO₂ battery. (e) Discharge–charge voltage curves at a current density of 0.2 mA cm⁻². (f) Discharge capacity curves of hybrid Na-CO₂ batteries with Co/Co₉S₈@SNHC or SNHC at a current density of 0.5 mA cm⁻². Reproduced with permission from [45]. Copyright 2021, Elsevier Ltd. (g–i) Electrochemical performances of the as-obtained Fe-Cu-N-C, catalysts: (g) Discharge–charge voltage curves at 0.05 mA cm⁻²; (h) Discharge–charge cycling curves; (i) Discharge–charge curves of Fe-Cu-N-C. Reproduced with permission from [58]. Copyright 2021, Royal Society of Chemistry.

Figure 5. Na-CO₂ battery cycling behavior of KB and Ru@KB cathode at 200 mA g⁻¹: (a) KB cathode. (b) Ru@KB composite cathode. (c) Discharge–charge profiles of Na-CO₂ batteries with KB and Ru@KB composite cathodes at a current density of 100 mA g⁻¹ in the first cycle. Reproduced with permission from [16]. Copyright 2019, Royal Society of Chemistry. (d) Schematic illustration of catalytic cathode of hybrid Na-CO₂ battery. (e) Discharge–charge voltage curves at a current density of 0.2 mA cm⁻². (f) Discharge capacity curves of hybrid Na-CO₂ batteries with Co/Co₉S₈@SNHC or SNHC at a current density of 0.5 mA cm⁻². Reproduced with permission from [45]. Copyright 2021, Elsevier Ltd. (g–i) Electrochemical performances of the as-obtained Fe-Cu-N-C, catalysts: (g) Discharge–charge voltage curves at 0.05 mA cm⁻²; (h) Discharge–charge cycling curves; (i) Discharge–charge curves of Fe-Cu-N-C. Reproduced with permission from [58]. Copyright 2021, Royal Society of Chemistry.

Due to the excellent catalytic activity, low cost and simple preparation methods, transition metal oxides such as NiO [59,60], MnO [37], MnO₂ [61], have been extensively studied in the field of Na-CO₂ batteries. Fang et al. obtained in situ grown (CMO@CF) composites of Co₂MnO₄ on carbon fibers by a simple hydrothermal method and high temperature annealing [55] which achieved a reversible charge/discharge process and remained stable after 75 cycles at 200 mA g⁻¹. In addition, XRD, XPS, Raman spectroscopy and SEM characterization demonstrated that the high catalytic activity of CMO@CF electrodes is mainly due to the homogeneous morphological, chemical and structural stability and the hybridized Co²⁺/Co³⁺ and Mn²⁺/Mn³⁺ redox pairs. ZnCo₂O@CNT materials were synthesized by the same hydrothermal method as that of preparing ZnCo₂O₄ porous carbon nanorod cathodes [43]. The discharge capacity of 500 mAh g⁻¹ at 100 mA g⁻¹ can be charged and discharged stably for at least 150 cycles. Theoretical calculations show that there is a strong adsorption energy for CO₂, Na and Na₂CO₃ on three surfaces of ZnCo₂O₄, namely the [001] surface of ZnCo₂O₄, the [111] surface with only Co atoms exposed and the [111] surface with Co and Zn atoms exposed. In addition, the exposed Co atoms on these three surfaces of the density of states (DOS) calculated surface are the real catalytically active sites for the CO₂ electrochemical reaction.

Apparently, the transition metal-based composites combine the high SSA and electronic conductivity of porous carbon materials and the unique multivalent characteristics of transition metals, which enable an effective catalytic effect in the application of Na-CO₂ batteries, which are considered as the most attractive catalytic cathodes. Therefore, in future research, the catalytic mechanism of transition metal materials should be understood in depth, and more attention should be paid to and design of composite carbon-based materials with synergistic effects of transition metals, transition metal oxides and multi-transition metals, and their controlled and large-scale production should be realized.

3.2.3. Other Types of Composite Materials

In addition to metal nanoparticles loaded on porous carbon materials, several composites have been applied to metal–CO₂ batteries, such as molybdenum-based electrode materials [62], metal–organic complexes, polymers and self-supporting freestanding cathodes.

Molybdenum carbide (Mo_xC) has a d-band electronic structure similar to that of precious metals, especially metallic palladium, and is considered to be a “similar catalyst to precious metals”. In early studies, Mo₂C was first applied to lithium–oxygen batteries and was shown to improve the Coulombic efficiency and cycle life of the batteries [63]. After that, Chen’s group prepared Mo₂C/CNT composites by a simple carbon thermal reduction method and applied them to Li-CO₂ batteries [64]. In this work, the Li₂C₂O₄ intermediate was stabilized by forming a Li₂C₂O₄-Mo₂C species on the surface of the catalyst. Finally, amorphous Li₂C₂O₄ final discharge products with thin film morphology were obtained instead of Li₂CO₃ species. In this case, Mo₂C makes the amorphous Li₂C₂O₄ components readily decomposable at very low potentials below 3.5 V. Moreover, it shows excellent cycling properties. However, the specific mechanism of Mo₂C catalyzing the formation of amorphous Li₂C₂O₄ remains unclear yet, and some advanced characterization tools as well as theoretical simulation calculations are needed in this regard to strongly elucidate the effects of Mo₂C on the performances of CO₂ batteries. In addition, MoS₂ is also an effective electrocatalyst for the decomposition of Li₂CO₃ [65]. In fact, we believe that it is a worthwhile effort to study the crystal structure of Mo_xC and to develop new molybdenum-based catalysts/materials for Na-CO₂ batteries.

Metal–organic backbone (MOF) is a crystalline porous organic-inorganic hybrid material with periodic network structure self-assembled by inorganic metal centers (usually metal ions or metal clusters) and bridging organic ligands [24]. Furthermore, MOFs not only combine the stiffness of inorganic materials with the flexibility of organic materials, but also have a porous structure and a metal–nonmetal structure, offering attractive prospects in the field of metal–air batteries research [66,67,68]. Successful applications of MOFs have been reported for Li-CO₂ batteries [24] and Na-CO₂ batteries [25]. The results show that MOFs have significant potential to increase the discharge capacity and reduce the polarization potential of Li-CO₂ batteries. However, the electronic conductivity of MOF is still insufficient, and the catalytic efficiency of the cathode needs to be further improved. We believe that the use of conductive MOFs with high electronic conductivity, along with metal ion doping to synergistically promote catalytic activity, can be attempted. In addition, MOF-derived single-atom catalysts are expected to be a promising new class of MOF-based materials in the laboratory due to their excellent activity in ORR [46,69].

Due to the insulator nature of discharge product, Na₂CO₃, of the Na-CO₂ batteries higher external voltages are typically required to decompose Na₂CO₃ during the charging process. However, high charging voltages typically lead to degradation of the battery components (e.g., electrolytes, carbon additives and binders), resulting in a decrease in overall battery performance [14]. Therefore, the development of carbon-free and binder-free cathodes is also an important task.

3.3. Single-Atom Catalysts

Large-scale applications of precious metals are hindered by their high price and scarce resources. Non-precious metal materials are promising candidates for noble metals, but with lower catalytic performance. In order to reduce the use of noble metals and further enhance the effective activity of metal atoms, single-atom catalysts (SACs) have emerged and attracted extensive attention [70]. Compared with nanomaterials, SACs can maximize the use of metal atoms; in addition, due to quantum confinement effects, these sub-nanometer catalysts have unexpected catalytic properties and can be used for various energy conversion reactions, and their applications in metal–CO₂ batteries are also highly anticipated [71,72,73]. Despite these advantages, SACs still suffer from some drawbacks, such as synthesis and characterization difficulties.

When metal particles are dispersed on the support at the atomic/cluster level, the properties of the catalyst, such as surface free energy, unsaturated coordination environment, quantum size effect and metal–support interaction change drastically, and SACs are easily agglomerated during preparation and application, leading to catalyst deactivation (Figure 6) [69]. Theoretically, increasing the loading of single-atom catalysts and avoiding atomic agglomeration are mainly achieved by increasing the surface area of the carrier and enhancing the interaction of the metal carriers [73]. Researchers have obtained “bottom-up” and “top-down” synthesis strategies for SACs [74]. Typical “bottom-up” strategies include atomic layer deposition (ALD), wet chemical methods, electrochemical deposition and chemical vapor deposition [75,76,77,78]. Atomic layer deposition (ALD) allows for precise control of the film thickness by evaporating the active precursor into a vapor and depositing it layer by layer on a carrier in a self-limiting manner using sophisticated equipment [79,80]. Wet chemical methods consisting of wet impregnation and co-precipitation are considered a promising technique for the synthesis of SACs because of their simplicity and the possibility of mass production of SACs [81]. However, it should be noted that the SACs obtained by this strategy usually exhibit the disadvantages of low metal loading and easy aggregation. In addition, electrochemical deposition has proven to be a universal method for the preparation of single-atom catalysts. Single-atom catalysts on various substrates such as transition group metals and oxides, sulfides, selenides and carbon materials were obtained using electrochemical deposition by Zeng et al. [82]. In contrast to the bottom-up strategy, top-down strategies are generally ball milling, high-temperature atom capture and pyrolysis [83,84,85]. Ball milling is a simple and versatile method that enables large-scale preparation by breaking/reconstructing the bonding, thus efficiently producing SA [83]. The pyrolysis strategy often uses various metal–organic complex precursors with molecular confinement and organic ligand coordination characteristics to obtain uniformly dispersed M-N-C site catalysts by pyrolysis, which are considered to be the most promising non-precious metal-based ORR catalysts [86]. Despite the consumption of organic reagents, it shows clear advantages: simpler preparation process, controllable tuning of the performance of SACs by controlling the pyrolysis environment and inducing defects [87,88].

With the development of nanotechnology and characterization science, advanced characterization techniques provide rich information and reliable evidence for studying the composition, structure and performance relationships of catalysts, which plays a crucial role in the rapid development of single-atom catalysts [89,90]. In addition, theoretical calculations can provide a reasonable reaction model for uniformly dispersed active sites and calculate the activation energy of the reaction. Combining advanced characterization techniques and data analysis of theoretical calculations, detailed information on the geometry and electronic structure of the active sites of single-atom catalysts can be obtained, thus elucidating the structure and performance relationships of single-atom materials in the catalytic process and guiding the design of single-atom catalysts [91,92]. The contents of characterization technology and theoretical calculation will be elaborated in the following section.

SACs are a promising redox electrocatalyst. However, the application of SACs in Na-CO₂ batteries is relatively new and much research is needed to explore the potential of SACs for practical applications in the next generation of rechargeable batteries. Theoretically, SACs ensure 100% atom utilization efficiency, which is very valuable for saving metal resources, especially precious metals. Recently, Zhu et al. successfully prepared single-atom Pt deposited on nitrogen-doped carbon nanotubes (Pt@NCNT) as a cathode material for Na-CO₂ batteries [44]. The single-atom dispersed Pt catalytic site, with a unique electronic structure and low coordination environment (Figure 7a,b), can achieve high activity and high selectivity due to the well-dispersed and fully exposed active site. Compared with the air batteries with pure NCNT as the positive electrode, the Pt single-atom catalyst can effectively improve the discharge reaction rate. The schematic diagram of Figure 7c illustrates the reaction mechanism based on Pt-SA in the discharge/charge electrochemical process of Na-CO₂ battery.

Given the two important factors of catalyst performance and cost, non-precious metal SACs with similar activity to noble metal SACs are more attractive [93,94]. Future work should target the combination of nitride or carbide supports with non-precious metal single atoms, which may provide unique electronic interactions with the metal and thus improve anode electrochemical catalytic performance [95]. In addition, SACs applied to Na-CO₂ batteries require harsh operating conditions, and the development of industrial fabrication methods for low-cost, stable, highly metal-loaded SACs is critical.

4. Progress of Solid Electrolyte and Interface Research

Many Na-CO₂ batteries reported still use the conventional organic liquid electrolyte [6,43], because Na-CO₂ batteries operate in a semi-open system, there are great safety hazards in actual application, such as the inherent volatility and flammability of organic solvents, and potential problems such as reduced battery life due to interfacial side reactions under high pressure [7]. Replacing the liquid electrolyte with a solid electrolyte can not only effectively solve the above safety hazards, but also inhibit the growth of dendrites and prevent CO₂ from corroding the metal Na anode during the long-term cycling process [11]. In addition, through scientific and reasonable structural design, solid electrolyte can also provide wider electrochemical window, higher energy density and longer cycle life [96,97,98].

For the solid electrolyte of the Na-CO₂ batteries, in addition to meeting the necessary energy requirements for a solid sodium ion conductor, e.g., high ionic conductivity, non-conductivity, interface compatibility, simplicity of preparation, low cost and environmental friendliness, it needs to be highly resistant to superoxide and peroxide to facilitate the smooth Na-CO₂ reaction [11]. It is difficult for the reported sodium ion solid electrolyte for Na-CO₂ batteries to meet all the above requirements at the same time, further research is needed [99,100]. In this section, we provide a brief overview of several important applications of solid electrolyte in the Na-CO₂ batteries, and analyze the research on solid electrolyte interface. Sodium-based electrolytes can be classified into three categories according to the chemical composition of the electrolyte and the transport mechanism of sodium ion, namely inorganic solid electrolytes (ISEs), polymer electrolytes (PEs) and composite polymer electrolytes (CPEs).

4.1. Inorganic Solid Electrolytes

Inorganic solid electrolytes (ISEs) are usually characterized by high ionic conductivity (10⁻⁵~10⁻² S cm⁻¹), high ionic mobility number, mechanical rigidity, non-flammability and non-flow characteristics due to their structural properties [101]. In order to understand the reasons for the high ionic conductivity of ISEs and to further improve its ionic conductivity, it is necessary to understand its ionic transport mechanisms and properties. The three principal migration mechanisms shown in Figure 8a are vacancy leap, gap-site leap and linkage leap [102]. According to the mechanism of vacancy leap and gapsite leap ion transport, the ionic transport is mainly related to the activation energy of the material and the number of defects in the vacancy, so ion doping can effectively improve the ionic conductivity of the material [103,104,105]. The ionic transport in linkage leap is not directly jumping between vacancies, but knocking the ions of adjacent sites to make them migrate to vacancies, so the ionic transport potential of linkage leap mechanism is lower than that of vacancy or gap site leap, so increasing the concentration of sodium ions in ISEs can improve the ionic conductivity [106,107,108,109]. ISEs can be classified according to their components and crystal structure types: NASICON, Na-beta-Al₂O₃, sulfide electrolytes and complex hydride electrolytes.

4.1.1. NASICON Structured Electrolytes

Na_1+xZr₂Si_xP_3−xO₁₂ (0 ≤ x ≤ 3) was proposed by Hong and Goodenough in 1976 and consists of NaZr₂(PO₄)₃ and Na₄Zr₂(SiO₄)₃ solid solution; the NASICON-structured fast ion conductor material with a 3D transport channel for Na⁺ was obtained by partially replacing the pentavalent P in NaZr₂(PO₄)₃ by the tetravalent Si and introducing Na⁺ to balance the charge [110,111]. Among them, Na_1+xZr₂Si_xP_3−xO₁₂ (0 ≤ x ≤ 3) fast ionic conductors are CO₂-stable ISEs. Due to its high ionic conductivity, good thermal and chemical stability, Na Superionic conductor (NASICON) structure electrolytes is the most widely studied inorganic material in the solid electrolyte of Na-CO₂ batteries [36,42,112,113]. Kim et al. [114] demonstrated that after NASICON materials contact with water, H₃O⁺ occupies part of the Na⁺ sites, causing a shift in the NASICON peak position, as revealed by XRD analysis. Although NASICON reacts with water, it does not cause the fracture and decomposition of the NASICON ceramic sheet, and there is no drastic decay in ionic conductivity.

According to the crystal structure analysis of this series of materials by Hong (Figure 8b) [111], the crystal structure of the material is monoclini with space group C2/c when 1.8 ≤ x ≤ 2.2 at room temperature; when x is outside this range, the crystal structure becomes tripartite (rhombic structure) with space group R-3c; with the change in temperature, the two crystal structures can be converted, and the phase transition temperature depends on the specific components of the material, usually at 150~200 °C. In the triangular structure (R-3c), SiO₄ or PO₄ tetrahedra are connected to ZrO₆ octahedra at the top corners, forming the three-dimensional skeleton for Na⁺ transport. Therefore, the most commonly used method to increase the ionic conductivity is doping. Heterogeneous element doping is used to enhance ionic conductivity by changing the local structure of the crystal, making the Na⁺ transport channel larger and lowering the ion diffusion potential barrier [107]. Among Na_1−xZr₂Si_xP_3−xO₁₂ (0 ≤ x ≤ 3), Na₃Zr₂Si₂PO₁₂ (x = 2) with the monoclinic phase structure is more stable. It exhibits the highest ionic conductivity of ~10⁻⁴ and ~10⁻¹ S cm⁻¹ at room temperature and 300 °C, respectively [110]. As shown in Figure 8c, Song et al. investigated the crystal structure and ionic conductivity of NASICON doped with different contents of alkaline earth metal ions (Mg, Ca, Sr, Ba) [105]. It was found that Mg ion doping improved its conductivity most significantly, and the ionic conductivity of Na_3.1Zr_1.95Mg_0.05Si₂PO₁₂ could reach 3.5 × 10⁻³ S cm⁻¹ at room temperature. However, Ba²⁺ ion substituted compounds by the elements with larger ionic radii exhibit narrower bottlenecks than the original compounds, suggesting that the ionic radius of the substituent plays an important role [105]. Lu et al. prepared Na_3.2Zr_1.9Mg_0.1Si₂PO₁₂ matrix materials by doping Mg²⁺ with Zr sites, showing an ionic conductivity of 1.16 mS cm⁻¹ at room temperature in order to obtain high-performance Na-CO₂ battery [113] (Figure 8d). In addition, heterogeneous elements with similar ionic radii to Zr (including calcium ions [115] and aluminum ions [116]) doped NASICON-based composite electrolytes can significantly improve the ionic conductivity of NZSP by changing the Na⁺ concentration and crystal localization structure, thus improving the Na⁺ transport channels [117].

Another way to improve the ionic conductivity of NASICON materials is to reduce the grain boundary resistance by adjusting the grain size and grain boundary chemistry [100]. Ihlefeld et al. investigated the size effect on the grain boundary resistance of Na_1+xZr₂Si_xP_3−xO₁₂ (0.25 < x < 1.0) [118]. The results showed that changing the Si/P ratio, increasing the process temperature and decreasing the annealing temperature all reduced its grain size and thus the ionic conductivity is increased. Hu et al. synthesized La³⁺ doped NASICON by adding La(CH₃COO)₃ to the precursor through a self-growth strategy [103]. XRD results showed that several new phases of Na₃La(PO₄)₂, La₂O₃ and LaPO₄ appeared at the grain boundaries of the final electrolyte material (Figure 8e). The newly appeared phases adjusted the chemical composition of the grain boundaries and increased the ionic conductivity at the grain boundaries.

4.1.2. Na-Beta-Al₂O₃

Na-beta-Al₂O₃ has become one of the most studied sodium ion solid electrolytes due to its high ionic conductivity and suitable mechanical properties, and was first applied in solid-state Na-S batteries [119,120,121]. Na-beta-Al₂O₃ has two crystal structure types as shown in Figure 8f, both are layered structures made of alternating stacks of spinel and sodium-conducting layers, with Na⁺ conducting in two dimensions between two adjacent spinel stacks, called sodium-conducting layers. one crystal structure of Na-beta-Al₂O₃ is a hexagonal crystal system structure made of two spinel structures stacked with space group P6m₃/mmc, labeled beta-Al₂O₃, composed of Na₂O-(8~11) Al₂O₃; the other is a tripartite crystal structure, consisting of three spinel structures stacked in the space group R3m, labeled β″-Al₂O₃, composed of Na₂O-(5~7) Al₂O₃.

A higher proportion of β″-Al₂O₃ phase is usually desired in the material because the ionic conductivity of the β″-Al₂O₃ phase is higher than that of the β-Al₂O₃ phase. This is not difficult to analyze: the two adjacent layers of spinel structure are connected by O²⁻ in the Na⁺ conducting sodium layer, forming an Al-O-Al bond and Na⁺ two-dimensional ion transport along the ab plane. For the β-Al₂O₃ phase, O²⁻ in the Na⁺ conduction layer has a higher electrostatic gravitational force on the surrounding Na⁺ and can accommodate a smaller amount of Na⁺, while for the β″-Al₂O₃ phase, O²⁻ in the Na⁺ conduction layer has a lower electrostatic gravitational force on the surrounding Na⁺ and can accommodate more Na⁺, thus the ionic conductivity of the β″-Al₂O₃ phase is higher than that of the β-Al₂O₃ phase. Single-crystal β″-Al₂O₃ can exhibit ionic conductivity up to 1 S cm⁻¹ at 300 °C, while polycrystalline structures exhibit 0.002 S cm⁻¹ at room temperature and 0.2–0.4 S cm⁻¹ at 300 °C. The lower ionic conductivity exhibited by polycrystalline structures is due to the high grain boundary resistance in polycrystalline β″-Al₂O₃. However, pure β″-Al₂O₃ is a thermodynamically sub-stable phase, which decomposes into Al₂O₃ and β-Al₂O₃ at 1500 °C and has relatively poor mechanical properties (200 MPa). Na-beta-Al₂O₃ is generally prepared by solid-phase method; however, the solid-phase synthesized Na-β″-Al₂O₃ powder has residual NaAlO₂ at the grain boundaries of the β″-Al₂O₃ phase and β-Al₂O₃ phase, which is unstable in air and easily reacts with CO₂ and H₂O [122,123]. Virkar et al. used steam-assisted method to obtain dense ceramic flakes by sintering Y-ZrO₂ and α-Al₂O₃ at 1450 °C to enhance their chemical stability [124].

The research of this material focuses on achieving stabilization in CO₂ atmosphere and reducing the impedance of grains and grain boundaries by suitable preparation methods, enhancing the ratio of β″-Al₂O₃ and reducing unwanted by-products. Doping is usually used to stabilize the β″-Al₂O₃ phases, such as Li⁺, Mg²⁺, Ni²⁺ and Ti⁴⁺; in addition, the overall mechanical strength can be enhanced by synthesizing mixed crystals of β″-Al₂O₃ and β-Al₂O₃ or by adding zirconium oxide [125].

4.1.3. Sulfide Electrolytes

Compared with oxide solid electrolytes, sulfides as electrolyte materials have higher ionic conductivities and lower grain boundary resistances, which is due to the intrinsic characteristics of sulfur [126]. S is less electronegative than O, which is less binding to Na⁺ and facilitates the free movement of Na⁺, thus it has lower grain boundary resistance; S has a larger ionic radius than O, and S replaces O to expand the lattice structure and form channels that facilitate the diffusion of Na⁺ and thus has a high ionic conductivity [106,127]. In addition, sulfide electrolytes also have the advantages of mild synthesis conditions, good mechanical strength, good ductility, etc. However, sulfide is unstable in humid air, easy to absorb water and easy to decompose with water in air, releasing toxic H₂S gas [128,129,130]. Therefore, it is necessary to improve the ionic conductivity and air stability by designing new sulfide-based electrolytes for the application in metal–CO₂ batteries.

Na₃PS₄ has two crystal structures, the tetragonal phase (P-42₁c, a = b = 6.9520 Ǻ, c = 7.0757 Ǻ) and the cubic phase (I-43m, a = b = c = 7.0699 Ǻ) [131,132]. The Figure 8g shows that in the cubic phase, Na⁺ is distributed in two distorted tetrahedral interstitial sites with space group I-43m, while in the tetragonal phase, Na⁺ is distributed in a tetrahedral site and an octahedral site with space group P-42₁c. Typically, Na₃PS₄ exists as a tetragonal phase, which can be transformed into a cubic phase at about 530 K. The cubic phase monomer has the lowest activation energy and has been studied the most. In 2012, the Hayashi and colleagues reported glass-ceramic sulfide electrolytes with an ionic conductivity of 2 × 10⁻⁴ S cm⁻¹ at room temperature, and the high ionic conductivity can be attributed to the stability of cubic Na₃PS₄ in microcrystalline glass electrolytes at room temperature [126]. It was found that the ionic conductivity of several crystalline phases of Na₃PS₄ sulfide electrolytes, has the following pattern: ionic conductivity of the cubic phase > ionic conductivity of the tetragonal phase > ionic conductivity of the glass-ceramic phase > ionic conductivity of the glass phase. Therefore, how to obtain a stable cubic phase Na₃PS₄ crystal structure at room temperature is one of the methods to improve the ionic conductivity of sulfide solid electrolytes.

Tuning the size of unit cell/channel by introducing Na vacancies, gaps or modulating the interaction between Na⁺ and anion skeleton in the lattice by means of elemental doping is an important method to improve the ionic conductivity of sulfide electrolytes. The doping of P sites with homologous As⁵⁺ of larger ionic radius expands the lattice and introduces Na vacancies while increasing the distance between Na-S [128,133]. In the tetragonal phase of Na₃PS₄, replacing S²⁻ with negative monovalent F⁻, Cl⁻, Br⁻ and I⁻ can also introduce Na vacancies according to the charge balance theory, thus increasing the migration probability of Na⁺ from one site to neighboring sites and thus increasing the ionic conductivity [134,135]. De Klerk and Wagemaker investigated the effect of Na⁺ vacancies on the ionic conductivity of Na₃PS₄ by molecular dynamics (MD) simulations: a significant increase at 300 K compared to pure cubic Na₃PS₄ (0.17 S cm⁻¹) [136]. Experimental results show that the ionic conductivity of chlorine-doped tetragonal sulfide (Na_2.9375PS_3.9375Cl_0.0625) at 303 K is 1.14 × 10⁻¹ S cm⁻¹, which is much higher than that of the pristine tetragonal Na₃PS₄ (5 × 10⁻⁵ S cm⁻¹) [135]. The first-principles study shows that the Na⁺ gap defect structure can also increase the Na₃PS₄ carrier density, i.e., in the cubic phase of Na₃PS₄, doping with tetravalent ions M⁴⁺ (M = Si, Sn, Ti, Ge) replaces the P⁵⁺ sites, while more Na⁺ will be introduced to maintain the electroneutrality and broaden the Na⁺ channel size, thus reducing the gap migration barrier and increasing its ionic conductivity [127]. The silica sulfide doped microcrystalline glass electrolytes (94Na₃PS₄·6Na₄SiS₄) also show higher ionic conductivity (7.4 × 10⁻⁴ S cm⁻¹) compared to the undoped state [120]. In addition, the interaction between Na⁺ and the anionic backbone has a significant effect on the ionic conductivity of the electrolyte. Theoretical calculations show that Se doping can increase the polarizability of the anionic framework, smooth the lattice and lowering the activation potential barrier [137].

However, to apply sulfide electrolytes in metal–air batteries, it is important to make them stable in air. In 2016, Liang et al. proposed an air-stable sulfide electrolyte [138]. When the P-site of Na₃PS₄ is completely replaced by Sb to obtain Na₃SbS₄, the unit cell and channels are expanded to further improve the ionic conductivity of Na₃PS₄, exhibiting a high ionic conductivity of 1 mS cm⁻¹ at 25 °C and a good compatibility with sodium metal anodes. More importantly, Na₃SbS₄ is stable when exposed to O₂ and H₂O, which can be well explained by the theory of soft and hard acids and bases [138,139]. In Na₃PS₄, the interaction between O²⁻ and P⁵⁺ (hard acid) is stronger, so S²⁻ (soft base) is easily replaced by O²⁻ (hard acid), while in Na₃SbS₄, Sb⁵⁺ (soft acid) and S²⁻. The interaction between them is strong and not easy to be destroyed by O₂. When exposed to air, pure Na₃SbS₄ tends to form Na₃SbS₄-xH₂O and H₂O may reversibly lose after heating at 150 °C for 1 h (Figure 8h). Despite the attractive stability and ionic conductivity of Na₃SbS₄ (~10⁻³ S cm⁻¹), the toxicity of Sb should be considered. At present, the research on sodium sulfide ionic solid electrolytes is still at the early stage and further studies are needed to improve their ionic conductivity, chemical stability and electrochemical stability.

4.1.4. Complex Hydride Electrolytes

In addition to the inorganic solid electrolytes mentioned above, complex hydride electrolytes are also good conductors of Na ions. In 2012, Orimo and colleagues first reported complex hydrides as Na⁺ ion solid electrolytes [140]. The reported complex hydride electrolytes have high ionic conductivities and ion mobility numbers, but complex hydrides are usually prone to water absorption, difficult to maintain stably in air. The electrochemical window is also not meet the requirement. If its chemical stability can be improved by modification, it may provide a new solution for the application of Na-CO₂ batteries.

Complex hydrogen compounds consist of a metal cation Na⁺ and a complex anion consisting of a central atom and a ligand hydrogen atom, which have been reported as Na₂(BH₄)(NH₂) [141], Na₂B₁₂H₁₂ [142,143], NaCB₁₁H₁₂ [144], NaCB₉H₁₀ [145], Na₂(B₁₂H₁₂)_0.5(B₁₀H₁₀)_0.5 [146,147], Na₂(CB₉H₁₀)(CB₁₁H₁₂) [144], Na₂B₁₀H₁₀ [148] and Na₃OBH₄ [149], etc. For example, Na₂B₁₂H₁₂ exhibits high ionic conductivity (>0.1 S cm⁻¹ at 573 K) due to its high-temperature disordered bulk-centered cubic phase (cation vacancy rich structure). However, due to the high phase transition temperature, the complex hydride electrolytes of large anions cannot meet the requirements of practical applications in different fields. It is necessary to reduce or eliminate the phase transition temperature of complex hydrides. It has been shown that the phase transition temperature can be significantly lowered after the introduction of C by chemical modification of the anion [145,150], for example, NaCB₁₁H₁₂ vs. Na₂B₁₂H₁₂: 380 and 529 K; NaCB₉H₁₀ vs. Na₂B₁₀H₁₀: 290 and 380 K, with ionic conductivity of 7 × 10⁻² S cm⁻¹ at room temperature (Figure 8i). In addition, mixing different anions helps to reduce or eliminate the transition temperature due to the introduction of geometric hindrance [140]. Tang and colleagues found that the transition temperature can also be effectively reduced by ball milling by reducing the grain size and disordering, and that a more homogeneous mixing of atoms is less prone to phase transitions [151]. Further studies should focus on investigating the stability and ionic conductivity of composite hydrides and applying them to solid-state Na-CO₂ batteries.

Figure 8. (a) Arrows indicate the three typical migration mechanisms: vacancy site leap, gap site leap and linkage leap. The circles indicate cations in stable (green) and substable (orange) positions in the model lattice. The dashed lines indicate the transition states of cation jumps imposed by the anion skeleton (not explicitly shown). Reproduced with permission from [102]. Copyright 2019, Springer Nature. (b) Crystal structures of representative NASICON (Na₃Zr₂Si₂PO₁₂) with rhombic and monoclinic crystalline phase. Reproduced with permission from [111]. Copyright 1976, Elsevier Ltd. (c) Lattice parameter, volume of unit cell and area of the bottleneck (marked as T1 in the inset) of Na₃Zr₂Si₂PO₁₂ and Na_3.1Zr_1.95M_0.05Si₂PO₁₂ (M = Mg, Ca, Sr, Ba). Reproduced with permission from [105]. Copyright 2016, Springer Nature. (d) Simplified view of the structure of Na_3.2Zr_1.9Mg_0.1Si₂PO₁₂ approximately along the [101] direction. Reproduced with permission from [113]. Copyright 2022, John Wiley and Sons. (e) X-ray diffraction patterns. The X-ray diffraction patterns of the nominal composition Na_3+xLa_xZr_2−xSi₂PO₁₂ (x = 0, 0.05, 0.10, 0.15, 0.20, 0.25, 0.30, 0.35, 0.40) solid electrolytes. The peaks in the circled by the dotted box represent the Na₃La(PO₄)₂ minor phase. The peaks marked by the black solid spheres and black asterisks denote the La₂O₃ and LaPO₄, respectively. Reproduced with permission from [103]. Copyright 2017, John Wiley and Sons. (f) Crystal structures of β-Al₂O₃ and β″-Al₂O₃. Reproduced with permission from [100]. Copyright 2018, Elsevier Ltd. (g) Crystal structures of Na₃PS₄ with cubic and tetragonal phase. Reproduced with permission from [131]. Copyright 2017, Reproduced with permission from. (h) XRD patterns of pristine Na₃SbS₄·9H₂O, as-prepared Na₃SbS₄, air-exposed Na₃SbS₄ and reheated air-exposed Na₃SbS₄ (150 °C for 1 h under vacuum). Reproduced with permission from [138]. Copyright 2016, John Wiley and Sons. (i) Relative geometries of the B₁₂H₁₂²⁻, B₁₀H₁₀²⁻, CB₁₁H₁₂⁻ and CB₉H₁₀⁻ anions. Reproduced with permission from [148]. Copyright 2014, John Wiley and Sons.

Figure 8. (a) Arrows indicate the three typical migration mechanisms: vacancy site leap, gap site leap and linkage leap. The circles indicate cations in stable (green) and substable (orange) positions in the model lattice. The dashed lines indicate the transition states of cation jumps imposed by the anion skeleton (not explicitly shown). Reproduced with permission from [102]. Copyright 2019, Springer Nature. (b) Crystal structures of representative NASICON (Na₃Zr₂Si₂PO₁₂) with rhombic and monoclinic crystalline phase. Reproduced with permission from [111]. Copyright 1976, Elsevier Ltd. (c) Lattice parameter, volume of unit cell and area of the bottleneck (marked as T1 in the inset) of Na₃Zr₂Si₂PO₁₂ and Na_3.1Zr_1.95M_0.05Si₂PO₁₂ (M = Mg, Ca, Sr, Ba). Reproduced with permission from [105]. Copyright 2016, Springer Nature. (d) Simplified view of the structure of Na_3.2Zr_1.9Mg_0.1Si₂PO₁₂ approximately along the [101] direction. Reproduced with permission from [113]. Copyright 2022, John Wiley and Sons. (e) X-ray diffraction patterns. The X-ray diffraction patterns of the nominal composition Na_3+xLa_xZr_2−xSi₂PO₁₂ (x = 0, 0.05, 0.10, 0.15, 0.20, 0.25, 0.30, 0.35, 0.40) solid electrolytes. The peaks in the circled by the dotted box represent the Na₃La(PO₄)₂ minor phase. The peaks marked by the black solid spheres and black asterisks denote the La₂O₃ and LaPO₄, respectively. Reproduced with permission from [103]. Copyright 2017, John Wiley and Sons. (f) Crystal structures of β-Al₂O₃ and β″-Al₂O₃. Reproduced with permission from [100]. Copyright 2018, Elsevier Ltd. (g) Crystal structures of Na₃PS₄ with cubic and tetragonal phase. Reproduced with permission from [131]. Copyright 2017, Reproduced with permission from. (h) XRD patterns of pristine Na₃SbS₄·9H₂O, as-prepared Na₃SbS₄, air-exposed Na₃SbS₄ and reheated air-exposed Na₃SbS₄ (150 °C for 1 h under vacuum). Reproduced with permission from [138]. Copyright 2016, John Wiley and Sons. (i) Relative geometries of the B₁₂H₁₂²⁻, B₁₀H₁₀²⁻, CB₁₁H₁₂⁻ and CB₉H₁₀⁻ anions. Reproduced with permission from [148]. Copyright 2014, John Wiley and Sons.

4.2. Polymer Electrolytes

The concept of polymer electrolytes (PEs) was first proposed in 1973 when Fenton et al. discovered that alkali metal salts dissolved in polyethylene oxide were found to form conductive complexes [152]. Subsequently, Feuillade et al. introduced organic plasticizers into the polymer–salt binary system to obtain a quasi-solid electrolyte [153]. Compared with inorganic solid electrolytes, electrolytes with polymer matrix usually have good flexibility, easy processing and high tolerance to vibration, shock and mechanical deformation and better interfacial contact and compatibility between electrodes and electrolytes.

For the ion mobility number and ionic conductivity of PEs, the methods of increasing amorphous regions and fixing anions are usually used. Soluble metal salts with anion stabilizing effect are generally preferred [11], such as NaCF₃SO₃, NaPF₆ and NaClO₄; in addition, a certain percentage of inorganic fillers are also added to enhance the amorphous region and ionic transport of organic electrolytes, thus improving ionic conductivity, mechanical strength and electrochemical properties. Polymer electrolytes are classified into solvent-free polymer electrolytes (SPEs) and gel polymer electrolytes (GPEs) based on their composition and physical form, and then explains the ion transport mechanisms, basic properties and their applications in Na-CO₂ batteries [99].

4.2.1. Solvent-Free Polymer Electrolytes

Solvent-free polymer electrolytes (SPEs) are usually consisted of only polymer matrix and lithium salt as solute without adding liquid solvent as plasticizer, and can be easily fabricated by solvent casting, thermoforming or extrusion techniques. SPEs that have been reported for Na-CO₂ batteries are mainly of the polyethylene oxide (PEO) type [25,30]. The carbonate electrolytes are susceptible to nucleophilic attack by superoxide radicals and are not suitable for metal–air batteries, while other types of SPEs are still to be developed.

PEO is the earliest and most studied, and its chemical structure is H-(O-CH₂-CH₂)_n -OH, a polyether compound with the advantages of good chemical stability, good compatibility with alkali metal negative electrode, good flexibility, good water solubility, low density, good viscoelasticity and easy film formation, and tolerance to superoxide radicals, which is a more suitable organic material for Na-CO₂ battery electrolyte. However, the current PEO-based solid electrolytes during application are often troubled by problems such as: the relatively high degree of room temperature crystallization of PEO (PEO chains are mainly crystallized at 65 °C), resulting in low room temperature ionic conductivity (~10⁻⁸ S/cm), thus requiring operation at higher temperatures; low upper limit of electrochemical stability potential (≤4.2 V), thus preventing the use of high voltage cathode materials; poor dimensional thermal stability (softening point of 55–64 °C); low mechanical strength (≤10 MPa).

According to the ionic transport mechanism of PEO-based PEs (Figure 9a), the metal salts dissociate and delocalize metal cation (Li⁺, Na⁺), metal ions with polar groups on the polymer chain, such as O and N. Under the action of electric field and above the glass transition temperature, the polymer molecular chain segments in the amorphous region (amorphous region) are able to vibrate, and with the movement of the polymer chain, the metal cations are continuously complexed and dissociated with the groups on the polymer chain segments. Metal cation jumps from one coordination site (usually composed of more than three electron-giving groups) to the next, thus enabling the transfer of metal ions [154]. However, since the movement of chain segments transfers both metal cation and anions, the metal cation migration number of PEs is usually less than 0.5 and their ionic conductivity is generally inferior to that of ISEs [155,156,157].

The following aspects can be used to improve the comprehensive performance of PEO-based polymer electrolytes: (1) to facilitate the migration of metal ions, the polymer-cation interaction must be a compromise between sufficient strength (ensuring salt solubility through cation solventization) and sufficient instability (facilitating ion hopping from one coordination site to another); (2) selecting a polymer matrix with a dielectric constant a polymer matrix with a large dielectric constant, a high dielectric constant facilitates effective charge separation of the metal salt and thus a high Na⁺ concentration; (3) the higher backbone flexibility and motility of the PEO chain facilitates the segmental movement of the polymer chain and (4) the high molecular weight of the polymer matrix is also desirable to obtain a polymer electrolyte with better mechanical strength.

In 2018, Chen’s team obtained a high-performance PEO/NaClO₄/SiO₂ all-solid polymer electrolyte by adding inorganic filler nano-SiO₂ and assembled an all-solid flexible Na-CO₂ batteries with sodium as the anode, PEO/NaClO₄/SiO₂ as the electrolyte and multi-walled carbon nanotubes as the cathode [30]. In the all-solid polymer electrolyte, PEO acts as a Na ion conductor to transport Na between the Na anode and the gas cathode, and the addition of nano-SiO₂ can reduce the crystallinity and promote the dissociation of NaClO₄, thus improving the ionic conductivity, in addition to improving its mechanical strength and thermal stability. The highest sodium ionic conductivity of 6.4 × 10⁻⁴ S cm⁻¹ and the highest ion transfer number of 0.56 were obtained at an operating temperature of 70 °C when the SiO₂ content in the polymer electrolyte was 3 wt% (200 μm thickness). Assembled flexible batteries with good bendability (21,000 times), foldability and shape customizability, as well as in the bending state (0~360°), stable operation time (80 h), high cycling stability (240 cycles with −0.4 V increase in overpotential), high capacity (450 mAh g⁻¹) and high energy density (173 Wh kg⁻¹) were obtained in electrochemical tests. Furthermore, Sun et al. developed a non-aqueous solution of expanded perfluorosulfonic ethylene carbonate and propylene carbonate (EC-PC) with a hexagonal structure of Na_0.67Ni_0.23Mg_0.1Mn_0.67O₂ as the cathode, and metal sodium as the anode. It proved to be a safe and durable all-solid-state Na-ion battery [158].

The research on SPEs for Na-CO₂ batteries is obviously insufficient, and subsequent studies need to further explore and investigate the Na⁺ transport mechanism of SPEs to elucidate the intrinsic connection between Na⁺ and polymer matrix, and to further develop polymer electrolyte matrix that can be applied to Na-CO₂ battery systems by simple blending, copolymerization, hyperbranching or cross-linking methods to find electrolyte materials with better overall performance.

4.2.2. Gel Polymer Electrolytes

Gel polymer electrolytes (GPEs) is a semi-solid electrolyte in a gel state, consisting of a polymer matrix, electrolyte salt and plasticizer. For GPEs, ion transport is the result of synergistic transport between the solid polymer matrix and the liquid electrolyte, and their room temperature ionic conductivity is higher than that of SPEs due to the presence of a certain amount of solvent; however, the safety in batteries is poorer than that of SPEs due to the presence of organic solvents. In addition, GPEs have both polymeric properties and are more flexible and easier to process than glass-ceramic type solid electrolytes. Therefore, GPEs exhibit good mechanical properties and good compatibility with electrodes. In quasi-solid polymer electrolytes, sodium ion transport occurs mainly in the liquid plasticizer containing dissolved lithium salts, while the polymer matrix provides mechanical strength to the GPEs and keeps it quasi-solid, thus minimizing the safety risk caused by leakage of liquid components. During battery charging and discharging, the plasticizer in the GPEs reacts on the electrode surface to form a solid electrolyte interface (SEI) film, similar to a liquid electrolyte. In contrast, electrochemically inert polymeric matrices are typically not involved in SEI formation.

PVDF-based gel electrolytes are the most widely studied matrix materials for GPEs due to their good film formation, large dielectric constant, high glass transition temperature and strong electron-absorbing groups. In 2017, Hu et al. [159] reported a quasi-solid Na-CO₂ battery with an applied PVDF-HFP-4% SiO₂/NaClO₄ -TEGDME CPE with advantages such as high ionic conductivity (1.0 mS cm⁻¹), strong toughness, non-flammability (automatically extinguished within 1 s even if ignited for 5 s) and operated at a capacity of 1000 mAh g⁻¹, 400 cycles at 200 °C, as shown in Figure 9b,c. The product has shown higher safety and good performance.

In addition to the above PVDF-based gel polymers, PEO-based and cyano polymer-based GPEs are widely used in lithium batteries, and related sodium ion gel electrolytes and their applications in sodium batteries are yet to be developed.

PEO-based GPEs have a high degree of crystallization at room temperature and low ionic conductivity at room temperature, making them difficult to apply in practical batteries. Suitable plasticizers can be added, such as polyethylene glycol [153] or crown ether [160], to reduce the crystallinity and enhance the room temperature ionic conductivity. Co-blending [161], cross-linking [162] or the addition of inorganic fillers to improve the mechanical strength and thermal stability.

Cyano (e.g., polyacrylonitrile) [153] cyano (such as polyacrylonitrile and cyanoethyl polyvinyl ether) is a polar group with a high dipole moment [163], and is a strong electron-absorbing group with a dielectric constant of about 30. Cyano has high oxidation resistance and its introduction into the polymer matrix can increase the oxidative decomposition voltage of the electrolyte. Feuillade et al. [153] prepared polyacrylonitrile electrolytes with ionic conductivity close to 1 × 10⁻³ S cm⁻¹ at room temperature. In polyacrylonitrile gel electrolytes, the ionic conductivity increases with increasing plasticizer and salt content. However, due to the strong polarity of the cyano group, there is a large passivation effect on the metal negative electrode leading to an increase in the interfacial resistance, thus requiring an effective negative interface protection.

4.3. Composite Polymer Electrolytes

Composite polymer electrolytes (CPEs) combining the soft mechanical properties of organic solid electrolytes and the high ionic conductivity of inorganic solid electrolytes offer a promising direction for highly stable Na-CO₂ battery electrolytes [97,115,164]. PEs have a generally low ionic conductivity. Inorganic powders incorporated into solid polymer electrolytes to obtain composite solid electrolytes can effectively reduce the orderly arrangement of polymer chains and prevent polymer crystallization; and more ion transport channels are formed in the surface region of nanoparticles, which helps to improve ionic conductivity. In 1988, Skaarup et al. added fast ionic conductor particles Li₃N as fillers to polyethylene oxide (PEO)-LiCF₃SO₃ matrix to obtain a composite polymer electrolyte, which not only improved its mechanical strength, but also had a higher ionic conductivity than the pure polymer electrolyte [165]. More interestingly, Wieczorek et al. found that the addition of nonionic fillers such as Al₂O₃ also improved the ionic conductivity of PEO-based polymer electrolytes, mainly due to the increase in the amorphous phase [166]. In addition, the addition of inorganic filler can improve the mechanical properties of electrolyte and maintain good flexibility, which improves the interfacial contact between electrolyte and electrode.

Inorganic fillers are mainly divided into two major categories: one is inert fillers, which do not have ion transport capability by themselves, such as Al₂O₃ [167], SiO₂ and TiO₂ [168], etc.; the other is active fillers, which have ion transport capability by themselves, such as Na₂SiO₃ [169], NASICON [170], β″-Al₂O₃ [171] and other inorganic solid electrolytes. Sun’s group prepared a high performance PVDF-HFP-Na_3.2Zr_1.9Mg_0.1Si₂PO₁₂ organic-inorganic composite solid electrolyte [113]. The Na_3.2Zr_1.9Mg_0.1Si₂PO₁₂ inorganic material with an ionic conductivity of 1.16 mS cm⁻¹ at room temperature was obtained by replacing the Zr ion in Na₃Zr₂Si₃PO₁₂ with Mg²⁺, which was then combined with structurally stable PVDF-HFP with excellent mechanical properties as a solid electrolyte for Na-CO₂ batteries. After 120 cycles at cut-off capacities of 200 mA g⁻¹ and 500 mAh g⁻¹, the intermediate gap voltage was below 2 V. The Coulomb efficiency, multiplicative performance and cycling performance of the batteries were significantly improved.

In addition to the above method of compounding inorganic particles directly with polymers, composite solid electrolytes can also be prepared by compounding functional polymers, functional inorganic particles, or polymer/inorganic particles together with a polymer matrix. For example, blending cellulose in sodium ionomer solid polymers [172] or glass fiber [173] can effectively enhance the mechanical properties of the electrolyte film; or adding functional inorganic particles to immobilize the anions on the polymer matrix chains and inhibit the formation of polarization centers, thus promoting the movement of cations [174].

Figure 9. (a) Schematic illustration of Li ion transport mechanism in PEO-based SPEs. Reproduced with permission from [154]. Copyright 2016, Royal Society of Chemistry. (b) Composition of the CPE. Inset: Transmission electron microscopy (TEM) image of fumed SiO₂. (c) Ionic conductivity of CPE with various contents of SiO₂. Reproduced with permission from [159]. Copyright 2017, the American Association for the Advancement of Science. (d) Charge/discharge curves tested at different current densities. (e) Cut-off capacity of NZM1SP-PVDF-HFP electrolytes at 200 mA g⁻¹ for 500 mAh g⁻¹. (f) Corresponding changes in intermediate charging and discharging voltages of cells with NZSP-PVDF-HFP and NZM1SP-PVDFHFP electrolytes. Reproduced with permission from [113]. Copyright 2022, John Wiley and Sons.

Figure 9. (a) Schematic illustration of Li ion transport mechanism in PEO-based SPEs. Reproduced with permission from [154]. Copyright 2016, Royal Society of Chemistry. (b) Composition of the CPE. Inset: Transmission electron microscopy (TEM) image of fumed SiO₂. (c) Ionic conductivity of CPE with various contents of SiO₂. Reproduced with permission from [159]. Copyright 2017, the American Association for the Advancement of Science. (d) Charge/discharge curves tested at different current densities. (e) Cut-off capacity of NZM1SP-PVDF-HFP electrolytes at 200 mA g⁻¹ for 500 mAh g⁻¹. (f) Corresponding changes in intermediate charging and discharging voltages of cells with NZSP-PVDF-HFP and NZM1SP-PVDFHFP electrolytes. Reproduced with permission from [113]. Copyright 2022, John Wiley and Sons.

4.4. Interfaces of Solid Electrolytes

During the electrochemical cycling of a solid-state Na-CO₂ batteries, the interface problems between the solid electrolyte and the electrode include two main ones: interfacial close contact and interfacial compatibility (no chemical reaction).

First, the close contact between the interfaces will facilitate the rapid transport of electrons or ions. In the conventional liquid electrolyte battery, the electrolyte has extremely high fluidity and good wettability to the electrode material, which ensures a low contact impedance inside the battery. In contrast, in solid-state batteries, the effective contact area is greatly reduced due to the point to point solid–solid contact between the electrolyte and the electrode, thus the interfacial impedance between the solid electrolyte and the electrode is very high (Figure 10a). Therefore, the current research work on solid-state batteries is mainly focused on improving the contact between solid electrolyte and electrode material and reducing the interfacial impedance. Sun et al. developed a monolithic symmetric cell for solid-state sodium batteries (SSSBs). A three-dimensional (3D) electronic and ionic conductive network is formed by integrating sodium anodes into NZSP-type monolithic structure. The interfacial resistance of the monolithic symmetric cell was significantly reduced, and exhibited a stable sodium plating/strip cycle with a low polarization of over 600 h [175].

Tong et al. artificially improved the interfacial stability and interfacial contact on the positive side of NASICON by obtaining a butanedionitrile-based plastic crystal electrolyte in situ on the cathode interface side (Figure 10b). This interfacial layer can be reversibly deformed with the change in particle volume, which reduces the interfacial charge transfer resistance and NASICON electrolyte is protected from successive interfacial side reactions, and thus the prepared solid-state Na-CO₂ battery can stably cycle for more than 50 cycles at a limited capacity of 500 mAh g⁻¹ [29]. In addition, solution casting method and in situ polymerization are also effective measures to enhance their interfacial contact.

In addition, chemical and electrochemical compatibility between different components will contribute to the stable operation of solid-state batteries (Figure 10c). Usually, the electrolyte and electrode will react at the interface, and dendrites may also form at the interface during the electroplating and stripping process, leading to short-circuiting of the battery [176,177]. Effective measures to solve the above problems include: (1) Optimizing the composition of the solid electrolyte or electrode, where the electrolyte reacts with the electrode to form an intermediate phase, but the interfacial products are stable and can reduce the interfacial impedance. For example, Tong et al. found that on the anode interface side, Na₃Zr₂Si₂PO₁₂ reacts with Na metal to form a thin passivation layer (indicated by NaO⁻), which not only prevents further reaction between electrolyte and Na metal, but also improves the interfacial contact between electrolyte and anode. (2) Optimizing the composition of electrode, by reducing the local current density and homogenizing the electric field distribution, so as to guide the sodium ion fluxes and improve the problem of battery short circuit caused by interface dendrites. Based on this idea, Sun et al. modified the surface oxidation functional groups of three-dimensional carbon cloth to make the carbon cloth sodium-friendly, and then injected molten sodium to obtain a stable Na @ CC anode which inhibited dendrite growth. The composite anode greatly improves the reversibility and safety of the electrode [178]. (3) To introduce artificial interfacial layers to reduce the interfacial impedance caused by the chemical or electrochemical reaction of the intermediate phase (Figure 10d). Goodenough et al. [179] heated sodium metal and NASICON to 380 °C, chemically reacted to get stable interface layer and then reduced to room temperature, and the negative metal Na showed good wettability to NASICON, and the symmetric batteries of sodium achieved good cycling at 65 °C and at current densities of 0.15 mA cm⁻² and 0.25 mA cm⁻². Sun et al. added 1,1,1,3,3,3-hexafluoroisopropylmethyl ether (HFPM) to 0.1 M NaPF₆ in 1,2-dimethoxyethane (DME) and fluoroethylene carbonate (FEC) electrolyte to form a new fluorine-containing organic layer, effectively stabilized the surface of Na anode, and showed excellent cycling performance [180].

Recently, a new approach can solve the above interfacial close contact and interfacial compatibility simultaneously. For example, Sui et al. and others constructed a thin self-reinforcing GPE in situ by polymerizing 1,3-dioxolane (DOL) in the nanofiber skeleton [181]. The framework is composed of polydopamine modified PVDF-HFP (PDA/PVDF-HFP) nanofiber membrane (Figure 10e). It has good affinity with PDOL and DOL, and can improve ionic conductivity. The DOL precursor solution can be firmly absorbed in the porous membrane, making the PVDF-HFP chain well swelled and forming good contact with the electrode interface, and be more stable to the negative metal Na [182]. After polymerization, there is no residual organic liquid GPE to ensure the full safety of the battery. Similarly, Sun et al. injected methylmethacrylate (MMA) into porous Na₃Zr₂Si₂PO₁₂-polymer vinylidene fluoride-hexafluoropropylene (NZSP-PVDF-HFP) composite membrane and obtained in situ polymerization polymethylmethacrylate (PMMA)-filled composite electrolyte membrane with excellent electrochemical performance [183].

Figure 10. (a) Schematic diagram of the “point-to-point” contact between the porous cathode and the NZSP. Reproduced with permission from [42]. Copyright 2021, Elsevier Ltd. (b) SN wetting on the porous cathode surface. Reproduced with permission from [29]. Copyright 2019, Royal Society of Chemistry. (c) Na dendrite growth process. (d) Contact model of ceramic solid electrolyte with metallic sodium in the process of sodium plating. Reproduced with permission from [184]. Copyright 2020, Elsevier Ltd. (e) Preparation of 3D PDOL@PDA/PVDF-HFP gel polymer electrolyte via in situ polymerization. Reproduced with permission from [182]. Copyright 2022, John Wiley and Sons.

Figure 10. (a) Schematic diagram of the “point-to-point” contact between the porous cathode and the NZSP. Reproduced with permission from [42]. Copyright 2021, Elsevier Ltd. (b) SN wetting on the porous cathode surface. Reproduced with permission from [29]. Copyright 2019, Royal Society of Chemistry. (c) Na dendrite growth process. (d) Contact model of ceramic solid electrolyte with metallic sodium in the process of sodium plating. Reproduced with permission from [184]. Copyright 2020, Elsevier Ltd. (e) Preparation of 3D PDOL@PDA/PVDF-HFP gel polymer electrolyte via in situ polymerization. Reproduced with permission from [182]. Copyright 2022, John Wiley and Sons.

5. In Situ Characterization Techniques and Theoretical Calculations

The electrochemical performance of metal–air batteries is closely related to the composition and structure of each component and the process of battery fabrication. A detailed understanding of the mechanism of action of each component is required to assemble a battery with excellent performance [185]. The combination of theoretical calculations and characterization techniques can help to better understand the electrochemical reaction mechanisms and the constitutive relationships of batteries materials in the cell system, which in turn can guide the design of safer batteries with higher energy density and longer life [186,187].

5.1. In Situ Characterization Techniques

During charge and discharge, the transport of Na⁺ ions between the cathode and anode usually induces various changes in the electrode materials and the electrode/electrolyte interface, such as material volume expansion and contraction, structural phase changes, morphological evolution and surface reconstruction phenomena. With the help of in situ characterization technology, the structural development process of Na-CO₂ batteries can be monitored in real time on multiple scales, providing a strong theoretical basis for studying the structural evolution of Na-CO₂ batteries [188,189,190,191]. This section summarizes typical in situ characterization techniques, including in situ diffraction (XRD, NPD and PDF), in situ microscopy (SEM, TEM and AFM), and in situ spectroscopy (FTIR, Raman and XAS), and gives some representative application examples. The key importance of each technique is highlighted in the study of metal–air batteries [185,189,191,192,193].

5.1.1. In Situ Diffraction Technique

The diffraction technique is a common method for characterizing the composition of crystalline materials and phases. The basic idea is as follows: any phase has its own characteristic diffraction spectrum; the diffraction patterns of any two phases cannot be identical; the diffraction peaks of multi-phase samples are the simple superposition of the phases [194,195]. Therefore, the position and intensity of the peaks are closely related to the composition, arrangement and type of atoms in the crystal. Typical in situ diffraction techniques include in situ XRD, in situ NPD and in situ PDF.

XRD originates from the interference of X-rays scattering crystals to produce diffraction patterns. In situ XRD can be used to monitor information on phase changes, lattice parameter changes and intermediate phase formation of the batteries’ material during charge and discharge [194,196]. The development of in situ XRD technology offers the possibility to explore the reaction mechanism of metal–CO₂ batteries [197,198,199]. The battery structure for in situ XRD is shown in Figure 11a [196]. Chen and colleagues used in situ XRD to observe the growth of Na₂O, the discharge product of Na-O₂ batteries in different electrolytes [198]. It proved that the electrolyte has an important influence on the reaction mechanism of Na-O₂ batteries, i.e., NaO₂ is the main discharge product in high and medium donor number electrolytes, but is rapidly converted to Na₂O₂ in low donor number electrolytes.

In addition to the conventional laboratory X-ray light source, the synchrotron X-ray light source can provide higher intensity and larger photon energy, which can collect higher signal-to-noise diffraction data in a short period of time and thus better study the lattice changes and structural evolution of battery materials during charging and discharging, reflecting more real-time change information [200].

Similar to in situ XRD techniques, in situ neutron diffraction allows for unique information for detecting the structural evolution of materials because neutrons are more sensitive to lighter elements (e.g., H [201,202], Li [97,203], O [204,205,206], Na [207], Al [208]) and can easily distinguish elements with similar atomic numbers (e.g., Fe and Mn) [209]. Specific applications include component determination, crystallinity, lattice constants, ion diffusion, expansion tensor, bulk modulus, phase transition and structure refinement and determination. However, the disadvantage of neutron diffraction is that it requires a strong neutron source, often a large sample size and a long data collection time [210]. Using NPD technology, Sun et al. found that there is an opportunity to create more sodium ion diffusion paths and improve the sodium ion diffusion capacity by applying perturbations, such as doping elements [113]. However, in situ XRD or NPD methods to study the long-range crystal structure of materials are unable to determine phase transitions in amorphous electrode materials or local structural changes during battery cycling.

Unlike the above diffraction techniques, the pair distribution function (PDF) can provide information about the local structure in the material, and can study both crystalline and amorphous materials [211]. The pair distribution function is also based on X-ray or neutron light source tests, but it provides information on the structure of the material at the atomic scale, especially on the atom–atom interactions in the structure. Fourier transform allows to obtain a real-space pair distribution function map, based on peak positions that can directly correspond to adjacent atomic spacings in real space; the peak intensity corresponds to the relative abundance of adjacent atomic spacings. For amorphous materials, local structural information, such as bond lengths and coordination numbers, can be extracted from the pair distribution familiarity map; for crystalline materials, the corresponding structural information can be fitted by taking the corresponding crystal structure model [212,213]. For crystalline materials, the corresponding crystal structure model can be adopted to fit the corresponding structural information. Unfortunately, there are few reports on the use of in situ PDF for characterization, and we believe that these studies will play a more important role in the design and development of solid state Na-CO₂ batteries in the future.

In situ diffraction techniques are important to obtain information on the structural changes in battery materials. The application of in situ diffraction techniques such as in situ XRD, in situ NPD and in situ PDF in the study of metal–air batteries are emphasized. However, some microscopy and spectroscopy techniques are still needed to obtain more comprehensive information about the structure and composition of cell materials. For example, while synchrotron radiation-based X-ray imaging techniques can provide morphological and chemical information over a wide length scale from tens to hundreds of nanometers. Although synchrotron X-ray imaging techniques can provide morphological and chemical information on scales from tens of nanometers to a few millimeters, imaging at the atomic length scale (submicron) still requires transmission electron microscopy (TEM) techniques. In addition, spectroscopic techniques can also provide important complementary information for structural analysis of cellular materials.

5.1.2. In Situ Microscopy Techniques

In situ microscopy techniques, such as Scanning Electron Microscopy (SEM), Transmission Electron Microscope (TEM) and Atomic Force Microscope (AFM), allow visual observation of particle size and morphological changes on the material surface and is often used to observe the fine structure of the electrochemical reaction process in metal–air batteries and thus to determine the (catalytic) reaction mechanism. In situ microscopy combined with other spectroscopic techniques, such as energy dispersive spectroscopy (EDS) and electron energy loss spectroscopy (EELS), gives information on the chemical composition, electronic structure and chemical bonding of the sample. TEM-EDS is essential for understanding the chemical composition and microstructure of materials. Scanning transmission electron microscopy-electron energy loss spectroscopy (STEM-EELS) also allows for selective imaging of various parts of the electron energy loss spectrum to obtain information on the structure, interfacial features, diffusion pathways and electronic structure at the atomic scale in the sample.

SEM uses a focused electron beam to scan and image the surface of a specimen point by point, and obtains information about the surface morphology of the test specimen by accepting, amplifying and displaying imaging of these signals. Usually, SEM collects secondary electrons from the sample surface, and its lining degree reflects the surface morphology and roughness of the sample. The spatial resolution of SEM can reach 10 nm, and the actual resolution is limited by the conductivity of the sample and the environment of the electron microscope cavity, with some time-resolved capability for in situ characterization of the battery. Neelam et al. studied the degradation mechanism of two sulfide-based solid electrolytes, b-Li₃PS₄ and Li₆PS₅Cl, during the operation of batteries by in situ SEM [214]. As shown in Figure 11b, compared with Li₆PS₅Cl, b-Li₃PS₄ shows huge plating, faster dendrite formation, cracks and ultimately cell failure.

TEM has extremely high spatial resolution, can realize the observation of diffraction patterns in tiny regions (several nanometers), and is suitable for the study of crystal structure of microcrystals, surfaces and thin films. Because electrons are easily scattered or absorbed by objects with low penetration, TEM must prepare ultra-thin samples, usually in thickness of 50–100 nm; on the other hand, the intense irradiation of the electron beam tends to damage the sample and bring artifacts. The in situ TEM technique has been applied to the study of the reaction mechanism and catalyst of metal–CO₂ batteries [44,215,216,217]. Zhu et al. directly studied the morphological changes in the cathode surface of single Pt atom-loaded nitrogen-doped carbon nanotubes (Pt@NCNT) during the charging and discharging of Na-CO₂ batteries using in situ ambient transmission electron microscopy (TEM) under CO₂ atmosphere (Figure 11c) [44]. Han et al. investigated the mechanism of the reversible redox reaction in Na-O₂ batteries with CuS cathodes using in situ TEM [215,216]. The results show that Na₂O₂ is the main final ORR product, which uniformly covers the whole linear cathode (Figure 11d). In addition, Huang et al. studied the reaction process of Na-O₂/CO₂ battery in situ using spherical aberration-corrected environmental transmission electron microscopy, and characterized the structure and composition of the discharging and partially charging products by using annular dark field images (ADF) and EELS [187]. Figure 12a,b showed the ADF diagrams characterizing the cathode surface: the discharge and charge products had similar spherical profiles in the low-loss and core-loss spectra. As shown in Figure 12c,d, EELS showed that: only Na and Na₂O₂ are present in the discharge and charge spheres, and the shell layer shows the formulation of Na₂CO₃.

Figure 11. (a) Schematic diagram of a typical battery device for in-situ XRD testing;. Reproduced with permission from [196]. Copyright 2019, John Wiley and Sons. (b) Plating, cracking and dendrite formation in b-Li₃PS₄ and Li₆PS₅C. Reproduced with permission from [214]. Copyright 2021, Cambridge University Press. (c) Morphology evolution of individual platinum-doped NCNTs (Pt@NCNTs) during electrochemical discharge and charge of Na-CO₂ nanobatteries. Reproduced with permission from [44]. Copyright 2020, Elsevier Ltd. (d) In situ TEM micrographs of the OER process in Na-O₂ batteries. Reproduced with permission from [215]. Copyright 2020, American Chemical Society.

Figure 11. (a) Schematic diagram of a typical battery device for in-situ XRD testing;. Reproduced with permission from [196]. Copyright 2019, John Wiley and Sons. (b) Plating, cracking and dendrite formation in b-Li₃PS₄ and Li₆PS₅C. Reproduced with permission from [214]. Copyright 2021, Cambridge University Press. (c) Morphology evolution of individual platinum-doped NCNTs (Pt@NCNTs) during electrochemical discharge and charge of Na-CO₂ nanobatteries. Reproduced with permission from [44]. Copyright 2020, Elsevier Ltd. (d) In situ TEM micrographs of the OER process in Na-O₂ batteries. Reproduced with permission from [215]. Copyright 2020, American Chemical Society.

Compared with SEM and TEM, AFM can more easily simulate cell environment conditioning, and be used for in situ operational characterization to monitor the microscopic morphology of the electrode surface in real time, provide physicochemical information of the electrode surface at the nanoscale, and provide an experimental basis for the optimization of electrode materials and electrolyte modification. In addition, unlike scanning electron microscopy which can only provide two-dimensional images, AFM can provide true three-dimensional images. At the same time, scanning electron microscopy needs to operate under high vacuum conditions, while AFM can work under atmospheric pressure or even in liquid environments. Cort’es et al. used improved AFM without putting AFM in glove box, visualized the formation and decomposition process on thin Li₂O₂ carbon cathode as shown in Figure 12e. This equipment provides a technical support for future studies of new cathode materials [218].

Figure 12. Annular dark field images (ADF) and EELS characterization of the structure and composition of the discharge and partial charging products: (a,b) ADF images of the time-dependent structural evolution of the air cathode of the Na-O₂/CO₂ batteries. During the discharge reaction, two balls emerged at 561 s at the CNT-Na substrate-O₂/CO₂ triple point, which then grew under a constant −3 V bias, showing the core-shell structure of the ball. (c,d) Low-loss and core-loss EELS spectra. Reproduced with permission from [187]. The low-loss and core-loss spectra of both the discharged (red profile) and charged (blue profile) balls show similar profile: in the Low-Loss region, there are three multiple plasma peaks at 5.7, 11.5 and 17.3 eV respectively, and another three plasma peaks at 22.9, 31.1 and 37.6 eV, and combined with side absorption peaks of O-K and Na-K in the Core-Loss region, indicating the presence of only Na and Na₂O₂ exist in the discharge sphere and charge sphere. The Core-Loss of the shell layer shows the presence of Na, C and O, indicating the formation of Na₂CO₃.Copyright 2020, American Chemical Society. (e) Schematic diagram of the flow electrochemical atomic force microscope (FE-AFM). Reproduced with permission from [218]. Copyright 2021, Elsevier Ltd. (f,g) In situ SERS characterization using LiCF₃SO₃-TEGDME with and without Ru catalyst electrodes during discharge and recharge in CO₂. Reproduced with permission from [27]. Copyright 2017, Royal Society of Chemistry.

Figure 12. Annular dark field images (ADF) and EELS characterization of the structure and composition of the discharge and partial charging products: (a,b) ADF images of the time-dependent structural evolution of the air cathode of the Na-O₂/CO₂ batteries. During the discharge reaction, two balls emerged at 561 s at the CNT-Na substrate-O₂/CO₂ triple point, which then grew under a constant −3 V bias, showing the core-shell structure of the ball. (c,d) Low-loss and core-loss EELS spectra. Reproduced with permission from [187]. The low-loss and core-loss spectra of both the discharged (red profile) and charged (blue profile) balls show similar profile: in the Low-Loss region, there are three multiple plasma peaks at 5.7, 11.5 and 17.3 eV respectively, and another three plasma peaks at 22.9, 31.1 and 37.6 eV, and combined with side absorption peaks of O-K and Na-K in the Core-Loss region, indicating the presence of only Na and Na₂O₂ exist in the discharge sphere and charge sphere. The Core-Loss of the shell layer shows the presence of Na, C and O, indicating the formation of Na₂CO₃.Copyright 2020, American Chemical Society. (e) Schematic diagram of the flow electrochemical atomic force microscope (FE-AFM). Reproduced with permission from [218]. Copyright 2021, Elsevier Ltd. (f,g) In situ SERS characterization using LiCF₃SO₃-TEGDME with and without Ru catalyst electrodes during discharge and recharge in CO₂. Reproduced with permission from [27]. Copyright 2017, Royal Society of Chemistry.

5.1.3. In Situ Spectroscopy Technique

Spectroscopic techniques such as XPS, FTIR, Raman and XAS are sensitive to changes in the local chemical environment, have elemental resolution and are suitable for compositional analysis of crystalline and amorphous phases.

X-ray photoelectron spectroscopy (XPS) is based on the photoelectric effect, using the interaction of X-rays with the surface of the sample to generate photoelectrons, and using an energy analyzer to analyze the kinetic energy of the photoelectrons, the X-ray photoelectron spectrum is thus obtained. Based on the measured kinetic energy of photoelectrons, it is possible to determine which elements are present on the surface and to know the content of an element on the surface based on the intensity of photoelectrons with a certain energy, with an error of about 20%. It is possible to determine both the relative concentration of elements and the relative concentration of different oxidation states of the same element. XPS not only determines the composition of all elements on the surface except H and He, but also gives information on the chemical state of each element, with high energy resolution and a certain spatial resolution (currently on the micron scale) and temporal resolution (on the minute scale). XPS also gives information on the surface, tiny areas and depth distribution information. In situ atmospheric pressure XPS (APXPS) is an effective tool to study the chemical state changes on the surface of solid electrolytes and catalysts. Wang et al. used synchrotron in situ APXPS to study the redox reaction of CO₂ on the surface of porous carbon electrode in ionic liquid electrolytes [219]. The results showed that the reduction in pure CO₂ at porous carbon electrodes has no electrochemical activity at room temperature. In contrast, when O₂ (CO₂:O₂ = 2:1) is added to the ionic liquid electrolyte of Li-CO₂ batteries, CO₂ is reduced to low-valent amorphous carbon and Li₂O₂/Li₂O. Compared with the reaction of Li₂CO₃, the charging reaction of amorphous carbon, Li₂O₂ and Li₂O is faster. The main reason why the ultimately discharge product is not Li₂CO₃ is that the strong solvation between ionic liquid electrolyte and Li⁺ stabilizes the intermediate anion of metastable CRR reaction.

In situ Fourier infrared reflectance absorption spectroscopy (FTIR) is particularly advantageous in identifying different functional groups and components during electrochemical reactions, and thus in situ FTIR techniques can be used to analyze reaction intermediates and determine reaction mechanisms. Although the application of in situ FTIR in metal–air batteries has not been widely reported, it can be expected to play an important role in future research in this field.

In situ Raman spectroscopy is based on the detection of laser-induced vibrational, rotational and other low-frequency leap patterns in electrode materials, and is a fingerprint spectrum of the structure of matter [192]. Raman spectroscopy is suitable for vibration mode measurements of symmetric vibrations, non-polar groups and homoatomic bonds, such as S=S, S-S, N=N, C=C and O=O. In addition to identifying substance types and chemical components, Raman spectroscopy mainly measures molecular vibration frequencies to quantitatively understand intermolecular forces and intra-molecular forces, and to infer information such as symmetry, geometry and arrangement of atoms in molecules. However, conventional Raman spectroscopy signals are relatively weak, and surface-enhanced Raman spectroscopy (SERS) is able to detect lower concentrations of species in the surface layers of bulk materials.

Zhou et al. used in situ SERS to study the catalytic activity of Ru in promoting CO₂ reduction to Li₂CO₃ and C on ruthenium-free sputtered gold electrodes and gold-ruthenium electrodes [27]. By observing figures f and g, it can be seen that the peak Raman intensity corresponding to Li₂CO₃ at 1080 cm⁻¹ decreases during the charging process and completely disappears at the end of the charging process. Comparing Figure 12f,g, the difference is mainly the Raman peak corresponding to the G-band of carbon at 1580 cm⁻¹. As shown in Figure 12f, the peak value of the carbon G band corresponding to 1580 cm⁻¹ in the gold electrode hardly changes during the whole charging process. In the Au-Ru electrode (Figure 12g), the peak intensity of the G-band in carbon (1580 cm⁻¹) decreases during the charging process and disappears completely at the end of the charging process, which is similar to the Raman peak of Li₂CO₃. This result shows that the ruthenium catalyst can significantly promote the reaction between Li₂CO₃ and carbon during charging (charging reaction), while the absence of ruthenium leads to the self-decomposition of Li₂CO₃ and leads to an irreversible discharge–recharging cycle of the battery.

However, SERS can only characterize information from a few atomic layers on the surface. In contrast, X-ray absorption spectroscopy (XAS) can effectively capture deeper signal variations in bulk materials. By analyzing and fitting the data to the X-ray absorption spectra, the most accurate characterization of the sample as a whole can be obtained, including interatomic distances, number and type of neighboring atoms and information on the average oxidation state of the elements in the coordination environment [189,220]. In addition, unlike in situ diffraction techniques, in situ XAS can provide real-time element-specific information on crystalline and amorphous evolution and phase transitions. XAS can be divided into X-ray absorption near edge structure (XANES) spectroscopy and extended X-ray absorption fine structure (EXAFS) [221]. EXAFS is commonly used in the range of 150–2000 eV to obtain quantitative local structural information such as bond length, coordination number and disorder of the central and coordination atoms. On the other hand, XANES can quickly identify the chemical state of elements in the low kinetic energy range of 5–150 eV. In addition, XANES can be used for time-resolved experiments and high-temperature in situ chemistry experiments because of its violent vibrations, short spectrum acquisition time and temperature dependence. The self-reinforced catalysis of LiCoO₂ as an electrocatalyst for Li-O₂ batteries was reported by Zhou et al. [222]. The evolution of Co electrons and local structure was investigated using XAFS, illustrating that the intercalation/extraction of Li⁺ in LiCoO₂ not only induces changes in Co valence and regulates the electron/crystal structure, but also the surface disorder, lattice strain and local symmetry, thus promoting bidirectional catalysis. The in situ technique XAFS has not been reported.

Photoelectronic information (XPS), which is sensitive to the material surface, can detect the changes in the element composition and chemical state information on the battery surface in real time. Raman spectroscopy is derived from the inelastic scattering of incident laser by the molecules on the surface of the object. In situ Raman spectroscopy can detect the changes in material composition and structure on the electrode or solid electrolyte surface of the all-solid battery. In situ XAS plays an irreplaceable role in the rapid and high precision analysis of the elements and their valence states and their respective distributions in solid state batteries. With the construction of more and more advanced synchronous light sources, in situ spectroscopy will play an increasingly important role in the research of solid state metal–air batteries and other energy materials.

Since each characterization method has its own advantages and disadvantages, combining the advantages of different characterization techniques to study the physicochemical changes and failure mechanisms of solid Na-CO₂ electrode materials and interfaces during operation can provide an important information for further understanding and optimization of material performance, and provide strong support for subsequent improvement of battery performance including energy density, cycle life and safety. With the development of modern technology, these in situ characterization techniques and data acquisition and analysis systems will be further improved and intelligent, and more in situ characterization techniques will be available for more systematic real-time inspection of all-solid-state metal–air batteries to guide the design of solid-state metal–air batteries.

5.2. Theoretical Calculations and Simulations

Theoretical calculations and simulations also play important roles in the development of battery technology. With the rapid development and wide application of computer simulation techniques in quantum and atomic scale-based materials science research, it provides the possibility to understand the reaction mechanism of microscopic metal–air batteries and to develop and design new battery materials efficiently. For catalytic materials, theoretical calculations and simulations can calculate adsorption energy, migration energy, charge distribution, electronic structure and even defect states to understand the reaction mechanism; for electrode materials, theoretical calculations and simulations can calculate the energy band gap, sodium storage voltage, bulk deformation and charge compensation mechanism of positive and negative electrodes and electrolyte materials. In addition, theoretical calculations and simulations can study the electrochemical window, ionic conductivity, ion mobility and ion diffusion barrier and transport mechanism. With the dramatic increase in computer computing power, methods such as high-throughput computing and machine learning have also started to develop rapidly, making the research and development of battery materials tend to be more efficient and intelligent. The combination of theoretical calculations and experiments helps to use the deep understanding of electrochemical theory to design and control electrochemical reactions in turn, thus advancing the battery research. This section outlines some common computational and simulation methods based on quantum mechanical theory in battery research, and briefly introduces the application of materials genetic engineering and machine learning in battery materials development.

5.2.1. Theoretical Calculation and Simulation Methods Based on Quantum Mechanics

Density Function Theory (DFT) was developed by Kohn, the 1998 Nobel Prize Chemistry awardee, which enables efficient solution of the many-body Schrödinger equation. In recent years, theoretical calculations based on density DFT and molecular dynamics (MD) simulations have greatly contributed to the understanding of the reaction mechanism of metal–air batteries at the atomic scale and to the efficient design of battery materials. The calculation of various properties of materials by inputting structural information of crystals has provided a theoretical basis for understanding the reaction mechanism, catalytic reaction kinetics and ionic transport in electrolytes in metal–air batteries.

Due to the multi-electron composite process of oxygen reduction reaction (ORR) and oxygen evolution reaction (OER), an in-depth understanding of the catalytic reaction process of cathode catalysts is essential to improve the reaction kinetics by changing the growth path and morphological evolution of the cathode surface [223]. In recent years, first-principles calculations have become a routine approach to understand complex catalytic phenomena and experiments at the electronic level, and have been further extended to find and design novel catalysts [25,61,224]. Jiang et al. used DFT calculations to investigate the contribution of point defects on carbon surfaces to the reaction of non-aqueous Li-O₂ batteries [225]. In this work, five representative defect structures were considered, including SV (point defects), DV5-8-5 (two pentagons and one octagon), DV555-777 (three pentagons and three heptagons), DV5555-6-7777 (four pentagons, one hexagon and four heptagons) and SW (Stone–Wales) defects. The calculated free energy results (Figure 13a–f) showed that the DV555-777, DV5555-67777 and SW type defects can increase the discharge voltage and decrease the charge voltage, while promoting the ORR and OER processes. In addition, the discharge voltage of DV555-777 is the highest, and the charging voltage of DV5555-6-7777 and SW is the lowest. Therefore, non-aqueous Li-O₂ batteries require carbon materials containing DV5555-6-7777 and SW type defects. Deng et al. obtained porous self-supporting cathodes for Li-CO₂ batteries by directly anchoring two-dimensional (2D) cobalt-doped CeO₂ nanosheets on graphene aerogels. Combined with the experimental results and DFT calculations, it was proved that co-doped CeO₂ nanosheets could effectively improve the conductivity and adsorption capacity of CO₂. In addition, they obtained the possible reaction path of non-aqueous Li-O₂ batteries from the perspective of thermodynamics, and the analysis results are shown in Figure 13g [224]. Recently, Sun et al. constructed a new autocatalytic system for lithium–air batteries using an in situ electrochemical doping strategy combined with theoretical computational simulations and systematically investigated its reaction mechanism and battery performance. The work was based on density function theory (DFT) for the theoretical study, and it was found that the batteries’ performance could be significantly enhanced by doping the discharge product Li₂O₂ with metal ions, and then the theoretical calculations were used to guide the experimental screening of the most effective doping structure [226].

For solid electrolytes, polymeric solid electrolytes are difficult to establish standard models for calculations due to their complex structures. In contrast, inorganic solid electrolytes, such as Na₃Zr₂Si₂PO₁₂, are easier to model due to their simple structure, and their corresponding theoretical simulations and calculations are more studied [107,227]. Park et al. investigated the effect of adding excess Na to standard NASICON and the mechanism of Na⁺ ion transport [107]. Both experimental and theoretical calculations demonstrated that the main mechanisms of Na⁺ ion transport in the NASICON structure are, grain boundary diffusion at low temperatures and grain diffusion at high temperatures, respectively (Figure 14a). In addition, it was found that after adding 10% excess Na, the expansion of the bottleneck of polycrystalline particle sodium diffusion channel was conducive to improving the bulk conductivity of NASICON. DFT calculation results show that by adding 10% excess Na, the minimum bottleneck area was expanded from 4.9922 Ǻ to 5.4086 Ǻ; the activation energy was significantly reduced. The minimum bottleneck area is the key factor to determine the conductivity of Na ion in NASICON electrolytes, which is basically consistent with the experimental results. Figure 14b shows that the ionic conductivity of NAICON has been significantly improved after adding excessive sodium, resulting from the enlargement of the bottleneck areas in the Na diffusion channels of polycrystalline grains.

The combination of DFT calculations and MD simulations has a strong potential in facilitating the development of metal–air batteries: the mass transfer dynamics of gas molecules and metal ions throughout battery system can be studied at the atomic scale. Notably, most of the current theoretical calculations are based on predictions under 0 K and vacuum conditions. However, powerful predictive tools should model the actual battery environment and conduct extensive experiments to verify theoretical simulations and improve accuracy.

5.2.2. Introduction to Material Genetic Engineering and Machine Learning

Materials genetic engineering is a new concept and method for materials research that has emerged in recent years and is at the forefront of materials science and engineering in the world today. Materials genetic engineering is a technological integration and synergy of high-throughput materials computing, high-throughput materials synthesis and characterization experiments and databases, which can at least double the speed of materials discovery, manufacturing and application, and reduce the cost by at least half. Materials genetic engineering is a leap forward in materials science and technology, replacing the traditional trial-and-error method of multiple sequential iterations with high-throughput parallel iterations, gradually changing from “experience-guided experiments” to a new model of “theoretical prediction and experimental verification” in materials research, and finally realizing materials design on demand.

There are some important steps in the process of high-throughput computational screening of materials [228]. First, a database of experimental crystal structures containing a large number of materials is needed as the seven points for high-throughput screening; second, structural information is input in bulk and data are screened; then the screened crystal structure information is generated into a data format that can be calculated using density functional theory and put into the server for calculation; then the calculation files are output and kept together to form a database after the calculation is completed; finally, the data from the database are used to analyze the various properties of the material, which can be used to guide the material design or to filter the data for the next cycle.

In recent years, the combination of machine learning and materials genetic engineering has led to the advancement of materials informatics and promoted the development of materials science. Currently, the use of data-driven machine learning algorithms to build material performance prediction models and then apply them to material screening and new material development has attracted much attention from research community. Machine learning can be seen as an umbrella term for a class of algorithms that can mine potential laws from large amounts of experimental data, acquire new knowledge or skills, build corresponding data analysis models and allow machines to repeatedly analyze the corresponding content and reorganize existing knowledge structures through the input of new data to continuously improve their performance. In the field of materials development, machine learning can show great potential and advantages, especially in the design of new catalysts, organic and inorganic battery materials, etc. [229]. A machine learning toolkit for genetic engineering improves the performance of solid-state Na-CO₂ batteries by searching for high-performance solid electrolytes.

Materials genetic engineering and machine learning aim to establish a new model of data-driven, computational simulation and theoretical prediction first, experimental validation second, thus replacing the existing empirical and experimental-based materials R&D paradigm, which is a powerful tool for the future development of metal–air batteries.

6. Outlook

Metal–CO₂ batteries are considered as a promising clean technology because of their potential for high energy density energy storage and power supply, as well as their ability to convert and immobilize CO₂ as well as mitigate global warming trends. After nearly a decade of development, although the research of Na-CO₂ battery has made great progress, compared with other energy storage systems, the research of Na-CO₂ battery is still in the infancy stage, facing the challenges of unclear electrochemical reaction mechanism, slow catalytic reaction of CO₂ gas at the cathode and the discharge product carbonate has difficulties in reversible transformation due to problems such as good thermodynamic stability, limited research on electrolyte materials, lack of suitable electrolytes and interface incompatibility. To solve these problems and realize the practical application of solid-state Na-CO₂ batteries operating at room temperature, we propose the following aspects for further research. First, the analysis of the electrochemical mechanism of Na-CO₂ batteries, especially the formation and decomposition process of the discharge products in the rate-determined step, is of great significance for the study of Na-CO₂ batteries, and should be considered in conjunction with in situ experimental studies, such as the electrochemical scanning microscope system combined with Fourier transform infrared spectroscopy, to further rationalize the design of the Li/Na-CO₂ cell system by analyzing the carbonate formation and decomposition pathways. Secondly, the sluggish reaction kinetics is a non-negligible challenge for Na-CO₂ batteries, which will lead to large overpotential and poor reversibility. Therefore, efficient catalysts with low cost, reasonable structure and high catalytic ability need to be developed to fill this gap. Considering the cost and performance, heteroatom doping, noble metal single-atom catalysts and transition metal complex catalysts are very promising.

Third, the applications of solid electrolytes in Na-CO₂ batteries are rather limited, and there is an urgent need for researchers to design and develop more solid electrolyte materials with high performance (e.g., high ionic conductivity and good chemical/electrochemical stability) based on fundamental principles and theoretical calculations, and to create stable interfaces with good contacts. In fact, ion transport in electrolytes and at interfaces is a multi-scale process, including atomic scale, micro and mesoscopic scale and device scale [102]. Monitoring the ion conduction and interfacial reactions at different scales, together with atomic-scale characterization techniques and theoretical modeling, are important for the design and fabrication of high-performance batteries. To improve the interfacial contact and form a three-dimensional (3D) electronic and ionic conducting network monolithic cell architecture may be a good choice. Finally, how to scale up the production of high-performance solid electrolytes by a simple, scalable and cost-effective method is also an important consideration for practical applications.

Finally, with the rapid development of in situ characterization technology and high performance computing, more comprehensive real-time inspection, data acquisition and analysis systems will be available and with the powerful predictive capabilities of computer simulation technology (especially theoretical calculations and simulations, genetic engineering of materials and machine learning), the development of battery materials will tend to be more efficient and intelligent, which is a powerful tool for the future development of metal–air batteries.

In summary, Na-CO₂ batteries demonstrate great application potential in CO₂ fixation as well as energy storage. They involve a cross multidiscipline. We believe that the development of practical metal–CO₂ batteries will bring new hope for achieving the strategic goal of carbon neutralization and carbon peaking.