Electronic modulation strategies for enhanced cathode catalysis in lithium–oxygen batteries: challenges, advancements, and future perspectives

Abstract

Lithium–oxygen (Li–O₂) batteries have emerged as a promising rechargeable energy storage technology, offering exceptionally high theoretical energy density, which makes them attractive for applications in electric vehicles, renewable energy storage, and portable devices. Despite their potential, the practical application of Li–O₂ batteries is significantly hindered by challenges such as high overpotential, low discharge capacity, poor rate capability, and unstable cycle life. These issues primarily arise from the insulating nature of the discharge product (Li₂O₂) and the sluggish electrochemical kinetics of the oxygen reduction reaction (ORR) and oxygen evolution reaction (OER). Addressing these challenges requires the development of efficient and stable cathode catalysts. Recent advancements in cathode catalyst design have provided valuable insights into the structure–performance relationship, especially regarding how electronic modulation can enhance catalytic efficiency. This review systematically examines the challenges, recent advancements, and future perspectives of cathode catalyst design for Li–O₂ batteries. Particular emphasis is placed on understanding how the electronic properties of catalysts influence battery performance and exploring innovative strategies, such as external-field-assisted catalysis, to optimize catalytic activity. Furthermore, the review discusses future research directions, including precise molecular and atomic-level modulation of catalyst electronic structures, advanced in situ characterization techniques, and the synergistic optimization of electrode/electrolyte interfaces, aiming to provide theoretical guidance for developing high-performance Li–O₂ batteries.

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Mengyao Huang

Meng-Yao Huang received her master's degree from the Department of Chemistry at Zhengzhou University in 2025. She is currently pursuing her Ph.D. at the Institute of Applied Physics and Materials Engineering, University of Macau, under the supervision of Prof. Kwun Nam Hui. Her research focuses on the design and mechanistic study of cathode catalysts and electrolyte additives in metal–air batteries.

Kwan San Hui

Kwan San Hui is a Full Professor in the Department of Mechanical Engineering at the College of Engineering, Prince Mohammad Bin Fahd University. He earned his Ph.D. in Mechanical Engineering from the Hong Kong University of Science and Technology in 2008. His research interests include advanced materials for energy storage, energy conversion, and electrocatalysis. He is a Fellow of both the Royal Society of Chemistry (UK) and the Higher Education Academy (UK).

Qingchao Liu

Qingchao Liu received his M.E. and Ph.D. degrees in Materials Science from Jilin University, Changchun, China, in 2012 and 2015, respectively. He is currently working at Zhengzhou University, China. His research focuses on the design and synthesis of functional materials for energy storage and catalysis.

Fuming Chen

Fuming Chen is a Professor at South China Normal University. He received his Ph.D. from Nanyang Technological University, Singapore, in 2011. His current research focuses on advanced functional materials for electrochemical energy storage and desalination.

Sambasivam Sangaraju

Sambasivam Sangaraju is a Senior Researcher at the National Water and Energy Center, United Arab Emirates University (UAE). He received his Ph.D. in Physics from Sri Venkateswara University, India, in 2008. He has held postdoctoral and associate professor positions in Taiwan, South Korea, Japan, the USA, China, India, and the UAE. His research mainly focuses on the development of high-performance nanostructured and nanocomposite materials based on super-thin, thin, and thick films, using novel and advanced material synthesis methods for applications in solar cells, energy storage, optoelectronics, magnetic materials, and hydrogen production.

Kwun Nam Hui

Kwun Nam Hui is an Associate Professor at the Institute of Applied Physics and Materials Engineering, University of Macau. He obtained his Ph.D. from The University of Hong Kong in 2009. His research focuses on developing new materials and understanding their properties for next-generation energy storage and conversion systems. He is a Fellow of the Royal Society of Chemistry (UK).

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Broader context

Lithium–oxygen (Li–O₂) batteries, with their remarkably high theoretical energy density, are at the forefront of next-generation energy storage technologies. As the demand for sustainable energy solutions grows, particularly for electric vehicles and grid-scale storage, Li–O₂ batteries offer a promising alternative to conventional lithium-ion systems. However, their practical application remains hampered by fundamental challenges, including high overpotential, poor cycle stability, and low efficiency due to the insulating nature of lithium peroxide (Li₂O₂) and sluggish oxygen redox reactions. Recent progress in catalyst development, particularly through electronic modulation strategies, has opened new avenues for enhancing catalytic activity, stability, and overall battery performance. This review provides a comprehensive analysis of these advancements, offering insights that could pave the way for more efficient, durable, and scalable Li–O₂ battery systems. By addressing the key challenges and identifying future research directions, this work contributes to the broader effort of making high-energy-density batteries viable for real-world applications.

1. Introduction

Lithium–oxygen (Li–O₂) batteries have garnered significant attention as a next-generation energy storage technology due to their exceptionally high theoretical energy density (up to 3500 W h kg⁻¹). This remarkable energy density positions them as transformative solutions for electric vehicles, grid-scale energy storage systems, and portable electronics.¹ Compared to conventional lithium-ion batteries (LIBs),^2–4 Li–O₂ batteries offer a promising pathway towards sustainable and efficient energy storage. However, the practical deployment of Li–O₂ batteries faces several critical challenges that hinder their large-scale application. These challenges primarily stem from the complex electrochemical processes at the cathode, where the formation of insulating lithium peroxide (Li₂O₂) during discharge and its sluggish decomposition during charging result in high overpotential and limited cycle life.^5–10

Despite the impressive theoretical energy density of Li–O₂ batteries, their practical application is significantly constrained by the intricate interplay of multi-scale electrochemical processes (Fig. 1). At the cathode interface, the wide band gap and insulating nature of the discharge product Li₂O₂ pose substantial obstacles to electron and ion transport during both charge and discharge cycles, leading to interface passivation.^11,12 Consequently, the kinetics of the ORR and OER in Li–O₂ batteries are relatively slow, leading to increased charge/discharge polarization and a decline in energy efficiency (<65%). Existing cathode materials struggle to maintain high activity over long cycles. Additionally, high costs and limited material availability, particularly the reliance on noble metal catalysts^13–15 and scarce carbon materials,^13,16–18 hinder large-scale application. During dynamic cycling, the disproportionation of Li₂O₂ and the generation of singlet oxygen (¹O₂) can cause electrolyte decomposition, forming byproducts such as Li₂CO₃ and carboxylic acids in ether-based electrolytes.^19,20 This process leads to cathode pore blockage and interfacial failure. At charging voltages exceeding 3.4 V, the electrochemical decomposition of LiO₂ or Li₂O₂ further generates ¹O₂, and the auto-oxidation of superoxide species and O₂ has been identified as a common mechanism of electrolyte degradation. Solid-state electrolytes have garnered considerable attention due to their enhanced chemical stability; however, they continue to face challenges related to poor ionic conductivity. A critical issue concerning the lithium anode is the undesirable morphological evolution during repeated lithium plating and stripping.^21,22 In particular, at higher current densities, lithium metal surfaces can develop mossy (dense, porous microstructures) and dendritic (sharp, outward-protruding) growth. Dendritic growth can lead to electrical short circuits in the battery, posing severe safety risks if the lithium dendrites penetrate the separator and contact the O₂ electrode. Additionally, the formation of mossy structures results in low coulombic efficiency due to the continuous formation of new solid electrolyte interphase (SEI) layers on the increased lithium surface area.²³ Moreover, the emergence of electrically isolated lithium structures renders portions of the lithium metal electrochemically inactive. These challenges fundamentally stem from the suboptimal regulation of electrode/electrolyte interfacial reaction pathways.

Fig. 1 Major challenges for the rechargeable Li–O₂ battery.

In recent years, with advancements in materials science and electrochemical theory, significant breakthroughs have been made in the design and optimization of Li–O₂ batteries.^5,42–45 Notably, substantial progress has also been achieved in the field of cathode catalysis (Fig. 2). The rational design and precise control of catalysts with excellent reaction kinetics, reversibility, and high ORR/OER activity are crucial for advancing the practical application of Li–O₂ batteries. Breakthroughs in density functional theory (DFT) calculations and in situ characterization techniques have enabled researchers to progressively elucidate the structure–activity relationship between catalyst electronic properties and ORR/OER kinetics.^46,47 This provides critical theoretical support for advancing Li–O₂ battery technology and its practical applications. This review begins by examining the fundamental charge–discharge mechanisms of Li–O₂ batteries and the critical role of cathode catalysts. It discusses the charge transfer mechanisms at the electrolyte-Li₂O₂-cathode interface and integrates recent advancements in cathode catalyst design, with a particular focus on the influence of catalyst electronic properties on catalytic activity (Fig. 3). A systematic analysis is provided on various tuning strategies, including d-band center modulation in metals, optimization of p-band non-metallic sites, occupancy of e_g orbitals, orbital hybridization, exposed crystal facet engineering, and external-field-assisted innovations, all of which contribute to enhancing catalytic performance. Finally, the current state of cathode catalyst design for Li–O₂ batteries is discussed, along with perspectives on future research. This review provides theoretical insights and guidance for the design and further optimization of advanced cathode catalysts.

Fig. 2 Timeline of the development of Li–O₂ battery cathode catalysts.^24–41

Fig. 3 A schematic illustration outlining the key topics discussed in this work, including a comprehensive analysis of the reaction mechanisms in Li–O₂ batteries,^48,49 the understanding of reaction interfaces and charge transfer mechanisms,^20,50,51 and the investigation of the electronic states of cathode catalysts.^39,52,53

2. Mechanism and charge transport path of Li–O₂ battery

2.1 Mechanism for Li–O₂ battery electrochemistry

The fundamental working mechanism of Li–O₂ batteries relies on the oxygen reduction reaction (ORR) and oxygen evolution reaction (OER),^54,55 which alternately occur during discharge and charge cycles to enable energy storage and release (eqn (1)).

2Li + O₂ → Li₂O₂ (E° = 2.96 V vs. Li/Li⁺) (1)

The discharge process of Li–O₂ batteries relies on the ORR at the cathode, wherein oxygen molecules are reduced to solid lithium peroxide (Li₂O₂) through a series of multi-step electron transfers (Fig. 4).⁵⁶ The specific reaction begins with the adsorption of O₂ molecules onto the cathode surface, followed by the acceptance of an electron, forming superoxide ions (O₂⁻) (eqn (2)). The kinetics of this step are significantly influenced by the oxygen concentration in the electrolyte and the surface activity of the catalyst. O₂⁻ ions then combine with Li⁺ from the anode to form an intermediate product, LiO₂. The stability of LiO₂ is contingent upon the donor number (DN) of the electrolyte solvent^57,58 and the adsorption strength (ΔE_ads) of the cathode catalyst. In low-DN solvents (e.g., TEGDME), the solvation of Li⁺ in the electrolyte is weak, leading to a tendency for LiO₂ to undergo secondary electrochemical reduction or disproportionation reactions at the electrode surface (eqn (3)–(5)). Conversely, in high-DN solvents (e.g., DMSO), facilitate the effective solvation of Li⁺, enabling LiO₂ to dissolve into the electrolyte and undergo disproportionation reactions, resulting in the formation of toroidal Li₂O₂ particles (eqn (6) and (7)).

O_2(g) + e⁻ → O₂⁻_(sol) (2)

Li⁺_(sol) + O₂⁻_(sol) → LiO_2(sol) (3)

2LiO_2(sol) → Li₂O_2(sol) + O_2(g) (4)

Li⁺_(sol) + e⁻ + O₂ → LiO_2(ads) (5)

Li⁺_(sol) + e⁻ + LiO_2(ads) → Li₂O_2(ads) (6)

2LiO_2(ads) → Li₂O_2(ads) + O_2(g) (7)

Fig. 4 Oxygen reduction mechanistic pathways in nonaqueous Li–O₂ batteries.

During the charging process, the decomposition of Li₂O₂ plays a critical role in energy storage (Fig. 5). This process involves the OER, whose sluggish kinetics lead to a high overpotential (typically >1 V), which is a fundamental issue limiting the efficiency of the battery.^55,59 The mechanism of OER is complex, and the reaction pathway, as well as the presence of intermediates, remains a subject of ongoing debate. Peng et al. used in situ surface-enhanced Raman spectroscopy (SERS) to monitor the charging process and found that the oxidation of Li₂O₂ did not exhibit characteristic peaks for the LiO₂ intermediate (such as the O–O vibrational peak at 1130 cm⁻¹).⁶⁰ They proposed that Li₂O₂ decomposes through a direct two-electron transfer pathway:

Li₂O₂ → 2Li⁺ + O₂ + 2e⁻. (8)

Fig. 5 Oxygen evolution mechanistic pathways in nonaqueous Li–O₂ batteries.

Further studies indicate that, similar to the discharge mechanism, oxidation pathway of Li₂O₂ during the charging process is influenced by the DN of the electrolyte solvent. During charging, Li₂O₂ first undergoes a delithiation reaction to form the Li_2−xO₂ intermediate (eqn (9)). In high-DN solvents, Li_2−xO₂ dissolves into LiO₂(sol) and subsequently undergoes disproportionation to generate O₂ (eqn (10) and (11)). In contrast, in low-DN solvents, the Li_2−xO₂ intermediate undergoes further oxidation via a solid-state mechanism to release O₂ (eqn (12)).

Li₂O_2(s) → Li_2−xO_2(s) + xLi⁺_(sol) + xe⁻ (9)

Li_2−xO_2(s) → LiO_2(sol) + (1 − x)Li⁺_(sol) + (1 − x)e⁻ (10)

2LiO_2(sol) → Li₂O_2(s) + O_2(g) (11)

Li_2−xO_2(s) → O_2(g) + (2 − x)Li⁺_(sol) + (2 − x)e⁻ (12)

2.2 Reaction interfaces and electron transport pathways in charging and discharging processes

In Li–O₂ batteries, the electrolyte–catalyst interface is the core region where charge transport and ORR/OER occur, involving the adsorption, transport, dissociation of reactants, as well as the nucleation and decomposition of Li₂O₂.^61,62 The wide band gap characteristic of the Li₂O₂ discharge product leads to an extremely low intrinsic conductivity (σ_ea = 2.2 × 10⁻¹³ S cm⁻¹), and as discharge proceeds, both electron and ion transport become increasingly restricted.^63,64 Significant efforts have been made to study the active reaction interfaces and transport limitations during the charging and discharging processes of Li–O₂ batteries.^63,65–67

Wang et al. conducted staged discharge experiments in tetraglyme electrolyte using isotope labeling, sequentially introducing ¹⁶O₂ and ¹⁸O₂.⁴⁸ Analysis using scanning electron microscopy (SEM) and time-of-flight secondary ion mass spectrometry (ToF-SIMS) revealed that the ring-shaped Li₂O₂ product exhibits a layered distribution, with Li₂¹⁸O₂ enriched on the outer surface (electrolyte interface) and Li₂¹⁶O₂ concentrated in the inner layer (electrode interface) (Fig. 6a). This finding supports the hypothesis that Li₂O₂ preferentially grows at the Li₂O₂-electrolyte interface. By reversing the isotope introduction sequence (introducing ¹⁸O₂ first, followed by ¹⁶O₂), the Li₂O₂ enrichment order was correspondingly reversed, further validating the correlation between oxygen evolution sequence and isotope distribution within Li₂O₂ particles. Wang et al. employed an Au electrode | Li₂O₂ thin film | Li⁺/DMSO electrolyte system to in situ SERS. By discharging at a high current density (100 μA cm⁻²), an ultrathin Li₂O₂ film was formed on the Au surface via a surface growth mechanism. Initially, the battery was pre-discharged in a ¹⁸O₂ atmosphere, forming a Li₂¹⁸O₂ film. Subsequently, switching to ¹⁶O₂ for discharge, as shown in Fig. 6b, SERS observations revealed that Li₂¹⁶O₂ formed from ¹⁶O₂ gradually replaced Li₂¹⁸O₂ at the Au electrode interface, indicating that the ORR occurred at the buried Au|Li₂O₂ interface.⁶⁵ In this scenario, O₂ and Li⁺ can penetrate the previously deposited Li₂O₂ film, likely through grain boundaries or defect channels, reaching the underlying gold electrode surface. The transport-limiting mechanism suggests that electron transport is the primary limiting factor responsible for the sudden failure of the battery during discharge.⁶⁵ The extremely low electronic conductivity of Li₂O₂ (σ_ea = 2.2 × 10⁻¹³ S cm⁻¹) prevents efficient electron transport to the reaction interface, leading to a sharp voltage drop. Although the ionic conductivity of the Li₂O₂ film is relatively low (σ_io = 3.1 × 10⁻¹² S cm⁻¹), experimental results indicate that Li⁺ and O₂ can slowly diffuse to the interface through defect channels in the film, suggesting that ion transport is not the primary limiting factor. Therefore, increasing the porosity and amorphous nature of Li₂O₂ can enhance defect density, facilitate Li⁺/O₂ transport, and mitigate electron transport limitations.

Fig. 6 (a) 3D distribution of ¹⁸O in discharge products reconstructed from ToF-SIMS depth profiling, with three selected depth layers.⁴⁸ (b) In situ SERS collected at open circuit potential (OCP, black) and at the end of discharge under a ¹⁸O₂ atmosphere (red), followed by spectra recorded during discharge in a ¹⁶O₂ atmosphere at 2.0 V vs. Li/Li⁺ over 0–5 hours.⁶⁵ (c) Charging curves of a Li–O₂ battery using a VC electrode, along with OEMS analysis of oxygen evolution.⁴⁸ (d) Schematic illustration of charge transfer during Li₂O₂ decomposition in the charging process. Decomposition occurs at the Li₂O₂-electrolyte interface, where charge transfer is mediated by (i) Li⁺ vacancy diffusion and (ii) hole polaron hopping; (iii) reactions take place at the Li₂O₂-cathode interface.⁶⁸

During the charging process, the charge transport pathway at the electrolyte–catalyst interface is significantly influenced by the charging rate and overpotential.^69,70 At low rates and low overpotentials, the solid-solution decomposition and solvent-mediated steps during charging initiate at the Li₂O₂-electrolyte interface. Wang et al. revealed the oxygen release behavior of Li₂O₂ particles through isotopic tracing experiments.⁴⁸ As shown in Fig. 6c, for micron-sized particles with the outer layer of Li₂¹⁸O₂ (in contact with the electrolyte) and the inner layer of Li₂¹⁶O₂ (in contact with the electrode), OEMS monitoring during charging reveals that ¹⁸O₂ is released first (accounting for the first quarter of the total charge capacity), followed by ¹⁶O₂ becoming dominant. This indicates that oxygen loss originates at the Li₂O₂/electrolyte interface rather than the Li₂O₂/electrode interface. A comparison between the pure carbon electrode (VC) and the ruthenium/carbon catalyst electrode (Ru/VC) reveals that both electrodes follow the same oxygen release sequence. However, the Ru/VC electrode exhibits a reduction in charging overpotential by approximately 500 mV. This difference is attributed to the Ru catalyst, which stabilizes lithium-deficient phases on the Li₂O₂ surface (e.g., Li_2−xO₂-Ru) through interfacial chemical adsorption, maintaining contact, facilitating charge transfer, and enhancing reaction kinetics. If O₂ primarily evolves at the Li₂O₂-electrolyte interface, electrons and ions must penetrate Li₂O₂ during charging, despite Li₂O₂ being a well-known wide-bandgap insulator. In fact, numerous studies have demonstrated that Li⁺ vacancies and hole polarons are the primary charge carriers in Li₂O₂.^67,70,71 Perfectly ordered Li₂O₂ is a wide-bandgap insulator, with an HSE-calculated bandgap of 4.2 eV. However, under non-stoichiometric conditions or in the presence of defect states (e.g., lithium vacancies), polaron conduction pathways can be activated. In Li₂O₂, hole polarons exhibit an in-plane migration barrier of only 68 meV, significantly lower than the interlayer hopping barrier.⁷² This suggests that controlling crystal orientation can optimize charge transport efficiency, providing valuable insights for electrode material design, such as tuning exposed crystal facets. Ball milling introduces grain boundaries and dislocation defects, increasing the overall conductivity of Li₂O₂ by approximately two orders of magnitude, from 3.4 × 10⁻¹³ S cm⁻¹ to 1.1 × 10⁻¹⁰ S cm⁻¹ at room temperature. Additionally, it generates a high-mobility Li⁺ “spin reservoir”, significantly enhancing ionic transport kinetics.⁷³ Amorphous Li₂O₂ exhibits unique transport properties due to its disordered structure. Its ionic conductivity (2 × 10⁻⁷ S cm⁻¹) is 12 orders of magnitude higher than that of its crystalline counterpart, effectively mitigating lithium-ion transport limitations during charge–discharge cycles in Li–O₂ batteries.⁶⁹ By optimizing electrode design (e.g., incorporating a porous structure), the synergy between the high ionic conductivity of the amorphous phase and rapid O₂ diffusion channels can alleviate electron transport bottlenecks and extend the activity of the reaction interface. Koffi et al. proposed that transition metals and their oxides can act as reaction promoters by chemically converting Li₂O₂ into lithium metal oxides, which subsequently facilitate oxygen release via an electrochemical delithiation process.⁷⁴

In summary, defect engineering, including grain boundary modulation, amorphization, and nanostructuring of Li₂O₂, can alleviate charge transport limitations during the charge–discharge cycles of Li–O₂ batteries.^75–77 At low charging rates, the decomposition of Li₂O₂ can be sustained through bulk and surface Li⁺ vacancy diffusion (Fig. 6d(i)) as well as hole polaron migration (Fig. 6d(ii)). However, when the transport rate fails to keep up with the charging current, at high charging rates or high potentials, the formation of a thick Li₂O₂ film impedes charge transport, shifting the reaction site to the Li₂O₂-cathode interface (Fig. 6d(iii)).⁶⁸ Nevertheless, the stepwise decomposition at the interface can cause the structural collapse of Li₂O₂ particles from the base, leading to void formation. Detached Li₂O₂ loses electrical contact, preventing its participation in subsequent reactions and accelerating capacity degradation.⁷⁸ Therefore, enhancing the adhesion between Li₂O₂ and the cathode or designing hierarchical porous electrodes, where micropores trap reaction products and nanopores limit particle detachment, can mitigate electrochemical detachment and improve cycling performance. Notably, although the formation of a dense Li₂O₂ film restricts bulk transport, it helps maintain local electrical contact, highlighting the need for morphology control to balance transport properties and interfacial stability. Highly active solid-state catalysts (e.g., Ru) enhance charge transport by forming metal oxide interfaces and stabilize Li₂O₂ intermediates through direct solid–solid interactions, thereby lowering the lithium extraction energy barrier and reducing overpotential.⁷⁹ On one hand, the coupling interaction between the electrode surface and Li₂O₂ can regulate the interfacial delithiation process. On the other hand, this necessitates that the substrate material exhibits strong resistance to Li₂O₂ oxidation. The oxidation tendency or multivalent capability of the catalyst determines its ability to stabilize Li₂O₂ intermediates, thereby influencing OER activity.⁷⁴ The position of the catalyst's d-band center directly influences the adsorption strength and decomposition pathway of Li₂O₂. Interfacial hybridization can create delocalized electronic channels, leading to charge redistribution. Qian et al. designed a nickel-based catalyst encapsulated in carbon nanotubes (CNTs), denoted as Ni/Ni₃C-CNT.⁶⁶ The electron delocalization of the inner Ni/Ni₃C material induces interfacial charge interactions, redistributing the electron density and surface electrostatic potential of the outer carbon shell. This modulation of surface adsorption states governs the kinetics of Li–O₂ chemistry. The Ni/Ni₃C-CNT catalyst exhibits moderate adsorption strength toward superoxide intermediates, effectively lowering the reaction energy barrier and achieving remarkable rate performance.

2.3 Rational design of cathode catalysts

In Li–O₂ batteries, the precise tuning of the cathode catalyst's electronic structure directly determines catalytic activity and reaction kinetics.^80,81 The position of the d-band center in transition metal catalysts significantly influences the reaction energy barrier by modulating the adsorption free energy of reaction intermediates (e.g., O₂⁻, LiO₂, Li₂O₂). Guo et al. employed a high-entropy strategy to design a series of alloy catalysts with widely distributed d-band centers.¹¹⁶ Their study revealed that catalytic activity follows a volcano-shaped curve as a function of adsorption strength, adhering to the Sabatier principle of heterogeneous catalysis. It was demonstrated that an excessively high d-band center leads to overly strong interactions between the catalyst and intermediate species, resulting in slow desorption kinetics and an increased reaction energy barrier. Conversely, an extremely low d-band center provides insufficient binding strength, causing product loss and surface passivation. As a demonstration of the trade-off between adsorption and desorption strength, HEA-PtIr exhibited the highest catalytic activity in Li–O₂ batteries, displaying an optimal adsorption capability. Additionally, a high surface electron density enhances the accessibility of active sites, facilitating oxygen molecule adsorption and dissociation. Intrinsic defects in catalysts, such as non-metal and metal vacancies, can introduce unsaturated coordination sites, enhance electronic conductivity, and lower reaction energy barriers. Heteroatom doping can modulate local electron density through orbital hybridization. For example, Co₃O₄ with a higher concentration of surface oxygen vacancies promotes the formation of Li₂O₂ morphologies that are more easily decomposable during charging, thereby enhancing battery reversibility and cycling stability.⁸² A cation vacancy-rich composite, VCo-CoSe₂@MXene, exhibited enhanced covalency upon Co vacancy introduction, facilitating electron transfer from Ti₃C₂ MXene to CoSe₂ and optimizing the electronic structure of interfacial Co sites.⁸³ Notably, the secondary Co site adjacent to the Co vacancy served as an active center for the oxygen reduction reaction, leading to a low overpotential (0.35 V) and excellent cycling stability (250 cycles at 500 mA g⁻¹) in Li–O₂ batteries. Heteroatoms, such as metal dopants, also serve as an effective strategy for modulating local electron density. Loading various metal elements onto Co₃O₄ hollow spheres to substitute Co³⁺ centers in the octahedral structure establishes electron transfer pathways that stabilize Co³⁺ centers during charge–discharge cycles.⁸⁴ This strategy maximizes the number of available active sites, facilitates oxygen adsorption and Li₂O₂ nucleation, and spatially separates active Co³⁺ centers, directing the growth of well-dispersed Li₂O₂ nanosheets.

The design of the air cathode directly influences charge transfer efficiency at the electrolyte/catalyst interface, thereby determining overall battery performance.^85,86 First, the air cathode must exhibit high electrical conductivity. The deposition of Li₂O₂ can lead to electrode surface passivation. At low overpotentials, electrons must transfer from the insoluble Li₂O₂ to the cathode surface, necessitating a cathode material with high electronic conductivity. Studies have shown that integrating catalysts with conductive materials (e.g., carbon nanotubes,⁸⁷ graphene⁸⁸) can significantly enhance cathode conductivity and improve charge–discharge performance. Catalytic activity is a key performance metric for the air cathode. The cathode catalyst must effectively catalyze the ORR and OER reactions to reduce overpotential and enhance battery energy efficiency.^89,90 The electronic structure of the catalyst, particularly the electron density of the metal d-band, surface electron density, and adsorption capacity for intermediates, directly influences catalytic performance. Proper tuning of the d-band center can optimize the catalyst's adsorption energy for reaction intermediates, thereby enhancing catalytic activity. Additionally, defects and doped sites on the catalyst surface can increase the number of active sites, further facilitating electrochemical reactions.⁹¹ Furthermore, the air cathode must exhibit high chemical stability. During long-term battery operation, the formation and decomposition of Li₂O₂ is a recurring process, requiring the catalyst to maintain structural and functional stability. Side reactions, such as the formation of LiOH and Li₂CO₃, may block cathode pores, hindering long-term battery performance. Therefore, air cathode materials should possess strong corrosion resistance and durability to suppress side reactions and ensure the long-term stability of catalytic active sites. Structurally, the air cathode should have a high porosity to provide sufficient space for oxygen adsorption and Li₂O₂ deposition.⁹² Studies have shown that mesoporous and microporous structures facilitate efficient Li₂O₂ storage and gas diffusion, thereby improving discharge capacity and rate performance.

3. Modulating cathode catalyst electronic properties

In Li–O₂ batteries, the electronic structure of the cathode catalyst plays a decisive role in the adsorption/desorption behavior of reaction intermediate, charge transfer efficiency, as well as the reaction pathways, energy barriers, and overall performance of the ORR and OER.^93–95 In recent years, with advancements in DFT calculations and in situ characterization techniques, researchers have gradually unveiled the structure–activity relationships between catalyst electronic properties and catalytic activity. This has led to the development of a multiscale design approach guided by electronic regulation. Therefore, this paper systematically reviews the current primary strategies for optimizing electronic structures, including the modulation of the metal d-band center, optimization of non-metallic p-orbital sites, e_g orbital occupancy, orbital hybridization, facet-directed exposure, and innovative strategies assisted by external fields. It also explores their role and potential in enhancing the electrochemical performance of Li–O₂ batteries (Fig. 7). Table 1 provides a summary of key parameters for several representative cathode catalysts in Li–O₂ batteries reported over the past three years, including categories, electrolytes used, overpotential, discharge capacity, and cycle life, for readers to gain a comprehensive understanding. This review aims to deepen the understanding of catalytic mechanisms and provide theoretical and practical guidance for the rational design of high performance cathode catalysts.^{52,66,83,84,96–115}

Fig. 7 This work discusses strategies for modulating the electronic properties of catalysts, including modulation the metal d-band center,¹¹⁶ non-metallic p-orbital sites,¹¹⁷ e_g orbital occupation,¹⁰⁶ orbital hybridization,¹¹⁸ optimization of exposed crystal facets,¹¹⁹ and innovative approaches assisted by external fields.¹²⁰

Table 1 Summary of recent progress in different cathodic catalysts and the corresponding electrochemical performance for Li–O₂ batteries

Cathode catalyst	Electrolyte	Overpotential (V)/[current (mA g^−1)]	Capacity (mA h g^−1)/[current (mA g^−1)]	Cycles [current (mA g^−1)]/limited capacity (mA h g^−1)	P.Y	Ref.
MnCr0.5Co1.5O4	1 M LiTFSI/TEGDME	0.67/100	16 388/100	329/(200/1000)	2024	94
WCoFe@Ni	1 M LiTFSI/DMSO	1.24/200	5800/200	>400/(500/1000)	2023	95
Mo3P@MO	—	1.05/100	4720/500	500/(500/500)	2022	96
Ni-NG SAC	1 M LiTFSI/TEGDME	—/200	24 248/200	>500/(300/1000)	2024	97
Ti0.87O2/MXene	—	0.69/200	13 596/100	407/(1000/600)	2024	81
SPC-2	1 M LiTFSI/TEGDME	1.04/200	20 287/200	226/(200/600)	2022	98
CeO2−δ	1 M LiTFSI/DMSO	1.09/250	—	>35/(250/10 000)	2023	99
GMS-sheets	1 M LiTFSI + LiNO3/TEGDME	0.82	>6200 mA h g^−1/0.4 mA cm^−2	260/(0.4 mA cm^−2)/4 mA h cm^−2	2023	100
bpyN-MOF	1 M LiTFSI/TEGDME	0.487/200	17 275/100	270/(2000/1000)	2024	51
NiCo2O4@CeO2	1 M LiTFSI/TEGDME	1.07/500	5586/500	400/(500/500)	2022	101
GMS-CP	0.5 M LiTFSI/TEGDME		6727/100	307/(0.1 mA cm^−2/0.25 mA h cm^−2)	2023	102
Ir1/Co3O4	—	0.305/200	16 861/200	180/(200/500)	2023	82
UIO-66-NH2	—	0.87/100	12 261/100	169/(200/1000)	2023	103
NiIII-NCF	1 M LiTFSI/TEGDME	0.92/500	16 800/500	>200/(500/1000)	2022	104
PtAu	1 M LiTFSI/TEGDME	0.36/100	5049/100	220/(500/1000)	2022	105
Ce1/CoO	1 M LiTFSI/TEGDME	0.46/250	37 462/250	136/(250/1000)	2024	106
FeCo-N4/HCS	1 M LiTFSI/TEGDME	0.64/200	1637/200	275/(200/1000)	2023	102
Ru0.8Pd0.2/Fe-CNCs	1 M LiTFSI/TEGDME	0.84/100	31 211/100	300/(200/1000)	2023	107
Ni/Ni3C-CNT	1 M LiTFSI/25% EMIM-BF4 + 75% DMSO	0.43/200	21 612/0.5 mA cm^−2	170/(0.5 mA cm^−2/0.5 mA h cm^−2)	2022	65
NiRu-HTP	—	0.76/200	18 280/500	>200/(500/1000)	2022	108
Na–Pb-MOF	—	0.52/100	6247/100	140/(100/1000)	2023	109
TA-ZIF@Ru-280	1 M LiTFSI/TEGDME		7000/200	300/(200/500)	2023	110
NiCO2O4/MnO2	1 M LiTFSI/TEGDME	0.73/0.5 mA cm^−2	7000/0.2 mA cm^−2	800/(0.2 mA cm^−2/0.5 mA h cm^−2)	2024	111
Pd3Pd	1 M LiTFSI/TEGDME	0.45/100	7746/—	175/(500/1000)	2023	112

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3.1 Modulation of the d-band center

Research on the regulation of catalyst electronic structures, particularly d-band center optimization, has been continuously emerging, providing a solid theoretical foundation and experimental evidence for improving the performance of Li–O₂ batteries across various material systems and design strategies.^121–125 Guo et al. systematically investigated the effect of d-band center position on the adsorption strength of oxygen intermediates by designing a series of alloy catalysts with a broad distribution of d-band centers (a wide range of adsorption strengths) through a high-entropy strategy.¹¹⁶ The binding energies of LiO₂/Li₂O₂ with five different alloys exhibit corresponding energy conversion efficiencies, as shown in Fig. 8a and b. As the d-band center decreases in the order of HEA > PtIr > HEAPtIr > HEAIr > HEAPt, it confirms that the relationship between oxygen binding energy on the catalyst and energy conversion efficiency is not linear but follows a volcano-type trend. The peak performance is observed for HEAPtIr (d-band center = −2.64 eV), which exhibits the lowest overpotential of 0.38 V and an energy conversion efficiency exceeding 80%. The projected and integrated crystal orbital Hamilton population (pCOHP and iCOHP) values further quantitatively evaluate the interaction between O1 and O2 bonds (Fig. 8c). From HEAPt to HEAPtIr, as the binding energy increases, a greater number of antibonding states become occupied, indicating that HEAPtIr further weakens the interaction between O1 and O2 compared to HEAIr and HEAPt. Notably, the -iCOHP values of HEA and PtIr are close to zero, indicating minimal interaction between oxygen atoms. This weak interaction makes the decomposition of LiO₂ (or Li₂O₂) more challenging, ultimately affecting the kinetics of the charging process. Fig. 8d and e summarize the relationship between the d-band center of cathode catalysts and catalytic activity in LiO₂ batteries. An excessively high d-band center results in overly strong interactions with oxygen species, leading to slow desorption kinetics, which hinders subsequent reactions and deactivates active sites. Conversely, an excessively low d-band center provides insufficient binding strength, leading to product loss and sluggish charge transfer, which is detrimental to rate performance and stability. An optimal d-band center ensures appropriate adsorption strength with oxygen species, balancing adsorption and desorption, thereby achieving the ideal point in the Sabatier relationship. In the study of high-entropy perovskite oxide (LaSr(5TM)O₃), five transition metal elements (TM = Cr, Mn, Fe, Co, Ni) occupy the B-site metal lattice.⁸¹ As shown in Fig. 8f, the compressive strain in LaSr(5TM)O₃ effectively modulates the occupancy state of the 3d orbitals at active cobalt sites (Co²⁺ → Co³⁺). Consequently, the d-band center of Co shifts from −2.08 eV to −1.37 eV (Fig. 8g and h), enhancing the adsorption capability of LiO₂ intermediates at cobalt sites. Additionally, the high electron-withdrawing capability of Cr sites ensures sufficient electron exchange, further strengthening the adsorption of LiO₂. The charge–discharge overpotential is reduced to 0.79 V, extending the cycle life to 226 cycles.

Fig. 8 (a) Binding energy of LiO₂ with cathode catalysts (HEA, PtIr, HEAPtIr, HEAIr, HEAPt) and the corresponding energy conversion efficiency. (b) Overpotentials of different alloy catalysts at a current density of 500 mA h g⁻¹. (c) -iCOHP values of the corresponding systems. (d) Orbital interactions between different d-band center catalysts and LiO₂/Li₂O₂. (e) Corresponding catalytic effects.¹¹⁶ (f) Charge density differences at Co sites in LaSr(5TM)O₃ and LaSr(4TM)O₃ (yellow regions represent charge accumulation, and blue regions represent charge depletion). PDOS and DOS of (g) LaSr(5TM)O₃ and (h) LaSr(4TM)O₃.⁸¹

By employing a heteroatomic doping strategy, such as Ni-doped RuO₂ catalysts (Ni-RuO₂), electron redistribution effectively lowers the d-band center of RuO₂, thereby weakening the adsorption strength of oxygen intermediates.¹²⁶ The Ni-RuO₂ catalyst exhibits a low overpotential of 0.79 V and an extended cycle life exceeding 791 cycles. Additionally, non-metal doping effectively induces electron delocalization, leading to changes in the d-band center at metal sites. As illustrated in Fig. 9a, moderate N doping in Ni₅P₄ (N-Ni₅P₄/5N-Ni₅P₄) induces a downward shift in the d-band center, whereas higher doping levels (10N/15N-Ni₅P₄) lead to an upward shift. This suggests that appropriate doping can optimize the adsorption/desorption equilibrium of oxygen intermediates during the redox cycle. As the d-band center shifts downward, the antibonding states induced by O 2p–Ni 3d coupling move further away from the Fermi level, weakening the interaction between intermediates and the catalyst surface. Meanwhile, the significant contribution of P 2p orbitals to the TDOS near the Fermi level (Fig. 9b) ensures that the system retains high conductivity after oxygen intermediate adsorption, facilitating the growth of discharge products. Theoretical calculations indicate that a downward shift in the d-band center enhances antibonding orbital electron occupancy (Fig. 9c), thereby reducing the dissociation energy barrier of oxygen species, providing a key descriptor for catalyst design.¹²⁷ A novel bimetallic conductive metal–organic framework (MOF) catalyst, NiRu-HTP, significantly enhances the oxygen reduction/evolution reaction (ORR/OER) kinetics in Li–O₂ batteries by synergistically tuning the d-band center and interfacial orbital hybridization.¹¹¹ As shown in Fig. 9d and e, the d-band centers of Ni-HTP and NiRu-HTP are −1.82 eV and −1.43 eV, respectively. The upward shift of the d-band center in NiRu-HTP enhances its interaction with O₂ molecules. DFT calculations indicate that at the Ni–N₄ site, the d_xz and d_yz orbitals of Ni overlap with the π orbitals of O₂, forming (d_xz/d_yz–π)* orbitals (Fig. 9f). At the Ru–N₄ site, the d_z² and d_xz orbitals of Ru couple with the σ and π orbitals of O₂, forming (d_z²–σ)* and (d_xz–π)* orbitals, respectively (Fig. 9g). In NiRu-HTP, the (d_z²–σ)* and (d_xz–π)* orbitals exhibit lower electron occupancy and higher energy levels, indicating enhanced adsorption of O₂. As shown in Fig. 9h and i, differential charge density maps further confirm that O₂ adsorption energy on NiRu-HTP (−1.47 eV) is significantly higher than that on Ni-HTP (0.28 eV). Additionally, NiRu-HTP transfers more electrons to O₂ (0.47 e⁻ compared to 0.22 e⁻), indicating more effective O₂ activation. By tuning the d-band center to regulate the oxygen reduction and evolution reaction pathways, the NiRu-HTP cathode significantly reduces charge–discharge polarization to 0.76 V and maintains stable performance over 200 cycles without significant degradation.

Fig. 9 (a) d orbits PDOS of Ni₅P₄ with different levels of N doping and after adsorption of oxygen-containing intermediates. (b) P 2p and O 2p orbital distributions before and after adsorption of oxygen-containing intermediates on Ni₅P₄ and N-Ni₅P₄. (c) Schematic diagram of bonding interactions between the valence band of surface adsorbates and the metal d-band.¹²⁷ TDOS and PDOS of Ni 3d and O 2p orbitals in (d) NiN₄-O₂ within Ni-HTP and (e) Ru N₄–O₂ within NiRu-HTP. PDOS distribution of (f) Ni 3d orbitals and (g) O 2p orbitals corresponding to (d) and (e). Differential charge density of (h) Ni-HTP and (i) NiRu-HTP after O₂ adsorption, where cyan and yellow regions indicate charge depletion and accumulation, respectively.¹¹¹ (j) Differential charge density map of Co₃O₄@NiFe₂O₄. (k) Planar-averaged charge density difference (CDD) distribution on the surfaces of Co₃O₄, NiFe₂O₄, and Co₃O₄@NiFe₂O₄ (right); numerical values on the left indicate specific surface charge densities of each material. (l) Schematic diagram of valence band interactions between adsorbed LiO₂ and the Ni 3d orbitals on NiFe₂O₄ and Co₃O₄@NiFe₂O₄ surfaces.¹²⁸

Through heterointerface engineering,^103,129,130 as demonstrated in the Co₃O₄@NiFe₂O₄ composite catalyst (Fig. 9j and k), the introduction of a heterogeneous interface can modulate the occupancy state of d orbitals at active metal sites via electron redistribution, leading to an upward shift of the d-band center.¹²⁸ This enhances oxygen intermediate adsorption and reaction kinetics (Fig. 9l). Additionally, high-entropy alloy (HEA) heterostructures (HEA@Pt) optimize the d-band center position of catalysts through the synergistic effects of multiple metal components and interfacial electron redistribution.¹³¹ Single-atom catalysts also exhibit significant advantages in d-band center modulation.^{53,84,132,133} A single-atom Ir catalyst (Ir@CeO₂) supported on the (111) facet of CeO₂ leverages the electronic buffering effect of CeO₂, which shifts the d-band center of Ir sites closer to the Fermi level.¹³³ This enhances Lewis acidity and effectively optimizes the adsorption of LiO₂ intermediates.

In recent years, significant progress has been made in optimizing the reaction kinetics of Li–O₂ batteries by tuning the d-band center of catalysts.^{111,116,134,135} Studies have shown that the position of the d-band center directly determines the adsorption strength of oxygen intermediates (LiO₂, Li₂O₂), thereby influencing the energy barriers and cycling stability of the ORR/OER. Strategies such as high-entropy alloy design, heteroatomic doping, interface engineering, and single-atom catalysis achieve precise modulation of the d-band center at active sites through electronic structure rearrangement (e.g., orbital hybridization, charge transfer, and strain effects). These approaches enable a balance between the adsorption and dissociation kinetics of oxygen intermediates, significantly reducing charge–discharge polarization and extending cycle life, thereby elucidating the complexity and feasibility of d-band center regulation.^66,136,137

3.2 Modulation of p-band non-metallic center

Research on optimizing p-band non-metallic sites primarily focuses on modifying non-metallic sites, such as nitrogen doping,^138,139 phosphorization,^100,117 sulfidation,¹¹³ and non-metal vacancy introduction,^113,140 to enhance the electronic structure and reaction kinetics of catalysts. For instance, non-metal doping reconstructs the electronic structure of carbon-based catalysts. Surface phosphorization of sawdust-derived carbon catalysts (Fig. 10a) effectively mitigates the corrosion between carbon materials and the electrolyte/products.¹⁰⁰ As shown in Fig. 10b, the SPC cathode features an open and orderly microchannel structure, providing efficient mass transfer pathways, Li₂O₂ storage space, and a highly active three-phase interface during the charge–discharge process in Li–O₂ batteries. As illustrated in Fig. 10c–e, the doped P atoms preferentially substitute one N atom in pyrrolic-N sites, forming graphitized P–N active sites that act as reaction kinetics promoters. Compared to untreated carbon electrodes, SPC electrodes accelerate reaction kinetics through high O₂/LiO₂ adsorption energies (O₂: −1.15 eV, LiO₂: −2.03 eV) and enhanced conductivity. Even under long-term cycling in the presence of highly oxidative discharge products (Li_2−xO₂), SPC effectively suppresses byproduct formation and corrosion of sawdust-derived carbon catalysts. Further studies reveal that introducing sulfur anions into nitrogen–carbon catalysts (SNC) forms N–S pairs, effectively tuning the p-band center of pyridinic nitrogen groups. This accelerates the formation and decomposition of Li_1–3O₂ intermediates, reducing the battery overpotential to 0.35 V.¹¹⁷

Fig. 10 (a) Schematic illustration of SPC catalyst synthesis. (b) Schematic diagram of a Li–O₂ battery using an SPC cathode. The SPC cathode features an open and orderly arranged microchannel structure, facilitating Li⁺ transport and oxygen diffusion. (c) Electronic structure of carbon and ELF. ELF of graphitized (d) P–C sites and (e) P–N sites.¹⁰⁰ (f) Ex situ SEM images of (a₁–a₃) S-Co₃O₄ and (b₁–b₃) B-Co₃O₄ cathodes at different stages: (a₁ and b₁) initial state; (a₂ and b₂) fully discharged state; (a₃ and b₃) fully charged state. (g) Schematic illustration of the Li₂O₂ formation mechanism on S-Co₃O₄ and B-Co₃O₄ surfaces.⁸² (h) ELF of CuS_1−x. (i) TDOS of CuS and CuS_1−x. (j) Schematic illustration of the mechanism by which CuS_1−x catalyst enables high-performance catalysis.¹⁴¹

Oxygen vacancies, as key regulators of catalyst electronic structures, exhibit significant interfacial reaction optimization effects in Li–O₂ batteries. Studies indicate that precisely controlling the concentration of oxygen vacancies can effectively direct the nucleation and decomposition pathways of Li₂O₂.^3,82,142,143 In cobalt oxide-based studies, Co₃O₄ nanoboxes with varying surface oxygen vacancy concentrations were synthesized as cathode materials for Li–O₂ batteries by controlling catalyst particle size.⁸²Fig. 10f a_1–3 and b_1–3 depict the electrode surfaces of the samples in their initial state, after full discharge, and after charging, respectively. After discharge, small-sized S-Co₃O₄, due to its high specific surface area and abundant oxygen vacancies, induces the uniform deposition of Li₂O₂ in a thin-film morphology with smaller and evenly distributed particles. In contrast, large-sized B-Co₃O₄, with insufficient oxygen vacancies, results in the aggregation of Li₂O₂ into micron-sized clusters (∼4 μm). The surface of S-Co₃O₄ contains more oxygen vacancies, providing additional nucleation sites for Li₂O₂. This facilitates the formation of highly dispersed and smaller-sized Li₂O₂, ensuring better contact with the catalyst surface. During subsequent charging, the efficient transport of Li⁺ and electrons at the interface facilitates the rapid decomposition of Li₂O₂, enhancing the battery's reversibility and cycling stability. However, Li₂O₂ catalyzed by B-Co₃O₄ continues to migrate and aggregate on the catalyst surface, forming larger particles that reduce decomposition efficiency during charging. As cycling progresses, Li₂O₂ continuously accumulates (Fig. 10g). This provides valuable insights into the structural design of cathode catalysts and the regulation of Li₂O₂ morphology in Li–O₂ batteries. The hollow catalyst TA-ZIF@Ru-280, synthesized via chemical etching, significantly enhances electron transport capability and reaction kinetics by optimizing active sites and oxygen vacancy distribution.¹¹³ Similarly, CuS_1−x nanoflowers, designed and synthesized based on a defect engineering strategy, serve as cathode catalysts for Li–O₂ batteries, effectively optimizing reaction kinetics.¹⁴¹ Sulfur vacancies induce changes in the valence state distribution of Cu and trigger charge redistribution (Fig. 10h). This shifts the Fermi level of CuS_1−x toward the conduction band, significantly enhancing its electronic conductivity (Fig. 10i). Additionally, sulfur vacancies create electron-deficient regions, providing unsaturated Cu active sites that facilitate stable LiO₂ adsorption. This promotes the formation of highly conductive LiO₂ intermediates via surface pathways, thereby lowering the decomposition energy barrier of Li₂O₂ (Fig. 10j). This modification not only enhances the electronic conductivity of the catalyst but also strengthens charge transfer at Cu active sites, improving the adsorption stability of intermediate species.

3.3 e_g orbital occupation

The occupancy of the e_g orbitals serves as a key descriptor for the electronic structure of transition metal ion d-orbitals, directly influencing the adsorption–desorption equilibrium between catalyst surfaces and reaction intermediates.^106,144,145 According to molecular orbital theory, tuning the e_g orbital occupancy modulates the filling degree of antibonding orbitals (σ and π), thereby adjusting the adsorption energy of intermediates such as *O₂⁻, *LiO₂, and *Li₂O₂. The design of an ideal catalyst requires precise regulation of electron distribution within an optimal e_g orbital occupancy range to balance adsorption and desorption, ultimately achieving high-performance ORR and OER activities.

The introduction of highly electronegative metal elements (e.g., Mo or Au) effectively modulates the e_g orbital occupancy of neighboring metal atoms. For example, in the PtAu alloy, the highly electronegative Au captures more eg electrons from Pt, resulting in Pt atoms in PtAu having fewer e_g electrons compared to those in PtRu.¹⁰⁷ As shown in Fig. 11a, the Pt 4f peak of PtAu exhibits a noticeable blue shift relative to that of PtRu. PDOS calculations reveal that the number of e_g electrons on the Pt 5d orbitals in PtAu is 2.84, lower than the 2.96 in PtRu (Fig. 11b). The reduced e_g occupancy of Pt in the PtAu alloy leads to an upward shift of the Pt d-band, significantly decreasing antibonding electrons and thereby enhancing the adsorption strength between LiO₂ and PtAu. Consequently, during the decomposition of Li₂O₂, PtAu exhibits a lower energy barrier (0.842 eV) compared to PtRu (1.01 eV) (Fig. 11c). Specifically, in the oxygen reduction reaction ORR process, the PtRu catalyst exhibits weak affinity toward LiO₂, leading to the formation of ring-shaped Li₂O₂via a solution-growth model. The insufficient Li₂O₂/cathode interface results in sluggish OER kinetics and undesired side reactions. In contrast, the strong interaction between LiO₂ and the PtAu cathode facilitates the deposition of Li₂O₂ nanosheets on the PtAu electrode surface via a surface-growth model. Benefiting from the superior contact interface between Li₂O₂ nanosheets and the PtAu electrode, PtAu demonstrates a lower OER energy barrier and outstanding catalytic activity. Consequently, Li–O₂ batteries based on the PtAu cathode exhibit a low charge overpotential of 0.36 V and excellent cycling stability over 220 cycles (with a capacity limit of 1000 mA h g⁻¹), ranking among the best-performing noble metal-based cathodes reported to date.

Fig. 11 (a) XPS spectra of the Pt 4f region for PtAu and PtRu cathodes. (b) PDOS of the Pt 5d orbitals in PtRu and PtAu. (c) Free energy profiles of the PtAu cathode under different potentials. The inset illustrates the optimized crystal structures of PtAu with specific adsorbates at the corresponding reaction steps.¹⁰⁷ (d) Schematic representation of the interaction between adsorbed O₂ and the d_z² and d_xy/yz orbitals of the Ni site in NiIII-NCF and NiII-NCF. Differential charge density distribution for O₂ adsorption on (e) NiII-NCF and (f) NiIII-NCF.¹⁰⁶ (g) Schematic diagram illustrating different spin states of Co³⁺ and the formation of the Co–O–Co spin channel with increasing Co³⁺ spin state.¹⁴⁴ (h) Work functions of MoS₂ and NiS₂, along with a schematic depiction of the formation of an internal electric field.¹⁴⁶

Spin-state modulation alters the electron distribution of transition metal ions, thereby influencing the occupancy of the e_g orbitals.^147–149 In two-dimensional MOFs, oxidative treatment of NiII-NCF nanowire arrays adjusts the spin state of some Ni²⁺ metal centers (t_2g⁶e_g²) to a higher valence state, Ni³⁺ (t_2g⁶e_g¹), leading to a significant increase in e_g orbital occupancy.¹⁰⁶ As illustrated in Fig. 11d, the orbital hybridization between Ni and O₂ is analyzed. Compared to the antibonding orbitals (d_z²–σ)* and (d_xz/d_yz–2π)* in NiII-NCF, the (d_z²–σ)* orbital in NiIII-NCF exhibits a higher energy level with lower electron occupancy, which favors oxygen adsorption and enhances the Ni³⁺–O₂ interaction. Bader charge analysis reveals that NiIII-NCF (Fig. 11f) transfers more electrons (0.26 e⁻) to O₂ compared to NiII-NCF (0.21 e⁻) (Fig. 11e), thereby enabling stronger O₂ activation and an increased electron transfer rate. The high affinity of Ni³⁺ sites toward the LiO₂ intermediate facilitates the formation of nanosheet-like Li₂O₂ within the voids between NiIII-NCF nanowires. As a result, Li–O₂ batteries based on the NiIII-NCF cathode exhibit a remarkably high specific capacity of 16 800 mA h g⁻¹ and stable cycling performance exceeding 200 cycles. Similarly, in ZnCo₂O₄ spinel structures, high-temperature calcination induces a transition of cobalt ions from a low-spin state (t_2g⁶e_g⁰) to a high-spin state (t_2g⁴e_g²), significantly increasing the delocalization of e_g orbital electrons (Fig. 11g). This enhancement facilitates spin-selective charge transfer during the ORR/OER process and accelerates reaction kinetics through the formation of spin channels.¹⁴⁴

Heterostructures regulate the electronic distribution of catalysts through interfacial effects, serving as another critical strategy for optimizing e_g orbital occupancy. In Ni/Mn-MOF, Mn doping further enhances the e_g orbital occupancy of Ni, significantly improving electronic coupling between Ni and redox intermediates.¹⁴⁵ As a result, this battery achieves an exceptionally high specific capacity of 28 464 mA h g⁻¹ at 100 mA g⁻¹ while maintaining long-term cycling stability over 524 cycles. A hierarchical NiS₂–MoS₂ heterostructure nanorod, synthesized via hydrothermal and chemical vapor deposition methods, features abundant defects and a mesoporous structure.¹⁴⁶ The work function difference between NiS₂ (4.192 eV) and MoS₂ (4.801 eV) induces charge transfer from NiS₂ to MoS₂, leading to the formation of an internal electric field (Fig. 11h). This internal field optimizes the e_g orbital occupancy of Mo atoms, as electrons are partially captured by highly electronegative Ni atoms, thereby reducing antibonding orbital occupancy. Consequently, the system achieves a balanced adsorption strength for oxygen intermediates (LiO₂, Li₂O₂). The Li–O₂ battery based on the NiS₂–MoS₂ cathode exhibits a coulombic efficiency of 99.65% and maintains stable performance even after 450 cycles at 1000 mA g⁻¹.

3.4 Orbital hybridization

In recent years, catalyst design based on orbital hybridization strategies has achieved breakthrough progress in the field of Li–O₂ batteries.^118,150,151 In the spinel-structured MnCr_0.5Co_1.5O₄, the introduction of Cr³⁺ at octahedral sites extends the Co³⁺–O bond length and reduces lattice symmetry, leading to enhanced Co–O bond covalency.⁹⁶ This increased metal–oxygen covalency accelerates charge transfer in oxygen-related reactions, thereby promoting the reversible deposition and decomposition of Li–O₂ battery discharge product. Orbital hybridization can also be effectively tuned by modifying the coordination environment of metal centers.^112,134,152 In the single crystalline naphthalene lead MOF (Na–Pb-MOF), which features seven-coordinated PbO₇ nodes, the Pb–O bond length (2.88 Å) is longer than that in the six-coordinated PbO₆ nodes of the tetramethoxy lead MOF (4OMe-Pb-MOF) (2.37 Å).¹¹² This elongation weakens the coupling between Pb 5d and O 2p orbitals, optimizing orbital interactions and improving the adsorption properties of intermediates such as LiO₂ (Fig. 12a and b). Consequently, this enhancement lowers the activation energy for both LiO₂ reduction to Li₂O₂ and LiO₂ oxidation to O₂ (Fig. 12c), significantly reducing the overall battery overpotential (0.52 V) and enabling stable cycling performance over 140 cycles.

Fig. 12 (a) Bader charge analysis of LiO₂ adsorption on Na–Pb-MOF and (b) 4OMe-Pb-MOF. (c) Gibbs free energy curves of Na–Pb-MOF and 4OMe-Pb-MOF. The inset illustrates the optimized crystal structures of adsorbates at different reaction steps.¹¹² (d) Surface-enhanced in situ SERS of Ce₁/CoO and (e) CoO cathodes. (f) Crystal orbital Hamilton population (COHP) analysis of the Co–O bond in Ce₁/CoO and CoO.¹⁰⁸ (g) DFT calculations and schematic representation of CeO₂/Ce₂O₃/Li₂O₂ formation during discharge. (h) Schematic representation of charge transfer from CeO₂ to the O₂²⁻-π* energy level in Li₂O₂.¹⁰¹

Rare-earth elements play a unique role in band hybridization due to their inner-shell orbitals and the strong localization of 4f electrons.^{101,108,153,154} In the Ce single-atom-anchored CoO system (Ce₁/CoO), Ce 4f and Co 3d orbitals undergo hybridization mediated by oxygen bridges, forming a dynamic electron reservoir that flexibly regulates the adsorption–desorption equilibrium of reaction intermediates.¹⁰⁸ As illustrated in Fig. 12d, in situ SERS reveals reversible changes in the Raman signals of the Ce₁/CoO cathode during charge–discharge cycles. The characteristic O–Co–O bending (E_g) and Co–O stretching (A_1g) vibrations gradually weaken, suggesting increased surface disorder and distortion of the CoO₆ octahedral structure, thereby maintaining a highly active state. This indicates that Ce₁/CoO facilitates self-regulating catalysis throughout the Li–O₂ battery cycling process. In contrast, after a single charge–discharge cycle, the E_g/A_1g Raman peak intensities of CoO fail to return to their initial state, suggesting potential surface reconstruction or loss of active sites (Fig. 12e). As shown in Fig. 12f, compared to CoO (ICOHP = −0.66 eV), the Ce₁/CoO system exhibits a richer bonding state between Co and coordinated oxygen atoms (ICOHP = −0.70 eV). This indicates an increased number of Co–O bonding orbitals and their enhanced occupancy below the E_F, ensuring the stability of Co–O bonds under redox conditions. Benefiting from the enhanced orbital coupling, the Ce₁/CoO cathode delivers an ultra-high discharge capacity of 37 462 mA h g⁻¹ at a current density of 250 mA g⁻¹, representing a 125% and 750% increase compared to CoO (16 629 mA h g⁻¹) and CeO₂ (4408 mA h g⁻¹) cathodes, respectively (Fig. 12g). An in situ electrochemical strategy was employed to synthesize a metastable CeO_2−δ catalyst, where the cerium–oxygen reaction (Ce³⁺ + O₂ + e⁻ → CeO_2−δ) occurs between Ce³⁺ in the electrolyte and O₂, forming oxygen vacancy-rich CeO_2−δ nanocubes with high-energy (311) crystal facets (Fig. 12g).¹⁰¹ By optimizing the O–Ce hybridized electronic structure, an electronic transport bridge was established between lattice oxygen and neighboring cerium atoms, creating a direct charge transfer channel from the Ce-4f orbitals in CeO₂ to the O₂²⁻–π* orbitals in Li₂O₂ (Fig. 12h). This facilitated Li–O₂ reactions, leading to the formation of submicron Li₂O₂ shell layers, thereby unlocking the capacity limitations of Li–O₂ batteries. As a result, the discharge capacity increased significantly to 32 015 mA h g⁻¹ at 250 mA g⁻¹, far surpassing that of conventional CeO₂ based cathodes (5703 mA h g⁻¹).

The study of intermetallic compounds¹¹⁵ and single-atom alloy^40,155 catalysts has introduced innovative pathways in orbital hybridization strategies. Pd₃Pb nanocubes synthesized via a wet-chemistry method exhibit an ordered intermetallic structure, where Pd single atoms are uniformly dispersed within a Pb matrix, forming stable Pd–Pb covalent bonds.¹¹² Intra-molecular hybridization between the Pd 4d orbitals and Pb 6p orbitals lowers the d-band energy level of Pd (Fig. 13a), leading to an increase in the oxidation state of Pd. As shown in Fig. 13b and c, the influence of intra-molecular d–p hybridization results in a reduced overlap between the bonding and anti-bonding states of the Pd 4d_z²–5σ and 4d_xy/d_yz–2π* orbitals upon LiO₂ adsorption. This weakens the orbital interaction with LiO₂. By leveraging both intra-molecular p–d orbital hybridization (direct coupling between Pd 4d and Pb 6p orbitals within Pd₃Pb) and inter-molecular orbital hybridization (interaction between Pd 4d and LiO₂ 5σ/2π* orbitals), a cascade electron regulation system is established, effectively optimizing the electronic structure of the catalyst (Fig. 13d). This cascade-oriented orbital hybridization in Pd₃Pb enables an ultra-low OER overpotential of 0.45 V in Li–O₂ batteries, achieving outstanding cycling stability over 175 cycles at a fixed capacity of 1000 mA h g⁻¹. A Pt₁Pd single-atom alloy catalyst was synthesized via a one-step wet-chemistry method. As shown in Fig. 13e, Pt atoms are uniformly dispersed within ultrathin Pd hexagonal nanosheets, occupying Pd lattice sites in a single-atom form and forming stable Pt–Pd covalent bonds.⁴⁰ The ELF (Fig. 13f) and Bader charge analysis (Fig. 13g) reveal that Pt single-atom sites exhibit electron enrichment (−0.23e), while Pd sites experience electron depletion (+0.03e), resulting in a strong electronic localization effect. This electronic redistribution facilitates the activation of the O₂/Li₂O₂ redox couple, effectively lowering the energy barriers for the ORR and OER. As illustrated in Fig. 13h, O₂ preferentially adsorbs at Pt sites, promoting the two-step reduction process of LiO₂ → Li₂O₂ while suppressing the formation of byproducts such as Li₂CO₃. These findings provide an important theoretical framework for the design of highly efficient Li–O₂ battery cathode catalysts based on orbital engineering. By precisely tuning metal oxygen bonding characteristics,^{101,108,151,156} local electronic structures at active sites,^96,150,157 and orbital coupling strength,^40,112 this strategy offers a pathway to overcoming limitations in reaction pathway control and stability that challenge conventional catalysts.

Fig. 13 (a) DOS of Pd d_z², d_xz/yz orbitals and Pb p_x, p_y, p_z orbitals after intra-molecular p–d orbital hybridization in Pd₃Pb. (b) DOS of Pd 4d d_yz/d_z²/d_xz orbitals in Pd. (c) DOS of Pd 4d d_yz/d_z²/d_xz orbitals in Pd₃Pb with adsorbed LiO₂. (d) Schematic representation of cascade orbital-oriented hybridization from intra-molecular Pb–Pd p–d hybridization in Pd₃Pb to inter-molecular orbital hybridization between Pd₃Pb (Pd 4d_xz/yz and 4d_z²) and LiO₂ (2π* and 5σ), using the d_z² orbital along the z-axis as an example.¹¹⁵ (e) AC-HAADF-STEM image of the Pt₁Pd alloy (single-atom Pt marked with green circles). (f) ELF values of Pt₁Pd. (g) Charge density difference distribution (regions of electron enrichment and depletion) in Pt₁Pd. (h) Schematic representation of the discharge pathway of a Li–O₂ battery on the Pt₁Pd surface.⁴⁰

3.5 Crystal face engineering design

Regulating the deposition pathway of discharge products through rational catalyst design has been widely employed to control the morphology and structure of Li₂O₂. Given that the ORR/OER reactions occur at the surface of oxygen catalysts, cathode catalyst facet engineering is considered an effective approach to enhance catalytic activity and modulate the deposition/decomposition pathway of discharge products by exposing specific crystal facets.^99,158–160 Studies have shown that two-dimensional Mn₃O₄ nanosheets supported on graphene preferentially expose the (101) facet. The adsorption energy of Li₂O₂ on the Mn₃O₄ (101) facet (−2.16 eV) is significantly lower than that on the (211) facet (−5.67 eV), indicating that the (101) facet facilitates the desorption of Li₂O₂ and reduces the decomposition energy barrier.¹¹⁹ By exposing the (101) facet and introducing oxygen vacancies, Mn₃O₄ NS/G effectively reduces the adsorption energy of Li₂O₂, lowers the charge overpotential to 0.86 V, and achieves an ultrahigh specific capacity of 35 583 mA h g⁻¹ at 200 mA g⁻¹. Similarly, two-dimensional layered structures with chemical bonding, such as stacked Ag–Te tetrahedral layers, predominantly expose the Ag-terminated (200) crystal facet.¹⁶¹ The (200) facet exhibits adsorption energies of 0.487 eV and 0.963 eV for LiO₂ and Li₂O₂, respectively, indicating physical adsorption. This facilitates the efficient formation of Li₂O₂via a two-electron pathway, as supported by the ΔG–U phase diagram, which reveals a voltage window of 2.27–2.31 V (Fig. 14a–c). In contrast, the (102) facet exhibits stronger adsorption (weaker chemical adsorption), which suppresses Li₂O₂ conversion. The narrow-bandgap telluride Ag₂Te successfully mitigates the facet-dependent catalytic anisotropy, with both the 2D stacked surface (200) plane and the side-plane exhibiting exceptional catalytic activity. As a result, Li–O₂ batteries based on the Ag₂Te cathode demonstrate an ultra-long cycling lifespan of over 300 cycles at a high current density of 500 mA g⁻¹, with a fixed capacity of 1000 mA h g⁻¹.

Fig. 14 (a) Reaction phase diagram of the Ag₂Te cathode on the (200) and (102) crystal facets. (b) Gibbs free energy diagram of the Ag₂Te(200) and (c) (102) crystal facet during the ORR/OER process.¹⁶¹ (d) Comparison of adsorption energies of the LiO₂ intermediate on Ni₂P(0001) and (1010) crystal facets. (e) DOS analysis of Ni₂P(0001) and (1010) crystal facets. (f) Schematic representation of the nucleation and growth mechanisms of discharge products on Ni₂P(0001) and (1010) crystal facets.¹⁶⁰ (g) Synthesis pathway of Co_0.85Se@CNT and schematic illustration of the lattice matching between the Co_0.85Se(101) crystal facet and the Li₂O₂(100) crystal facet.¹⁶⁵

Crystal facet engineering, as an important strategy to regulate the deposition pathways of discharge products and enhance catalytic performance in Li–O₂ batteries, demonstrates significant application potential.^162–164 By precisely controlling the exposed orientation of catalyst facets and their surface electronic structures, it is possible not only to modulate the adsorption behavior of intermediates such as LiO₂ and Li₂O₂ but also to induce the formation of discharge products with specific morphologies, thereby alleviating electrode passivation. Moreover, integrating oxygen vacancy introduction, lattice matching effects, and porous structure design can further optimize the distribution of active sites, reduce polarization, and improve cycling stability. This approach provides both theoretical insights and practical pathways for the development of highly efficient and stable cathode materials for Li–O₂ batteries.

Crystal-facet-controlled Ni₂P nanorods were synthesized on Ni foam via a hydrothermal-phosphorization strategy, effectively regulating the nucleation, growth, and morphology of LiO₂ and Li₂O₂.¹⁶⁰ DFT calculations reveal that the adsorption energies of LiO₂ on the Ni₂P(0001) and (1010) facets are −0.42 eV and −1.87 eV, respectively (Fig. 14d). Compared to the (1010) facet, the (0001) facet exhibits a higher DOS near the Fermi level, indicating a stronger electron transport capability for the Ni₂P-0001 cathode (Fig. 14e). As illustrated in Fig. 14f, Ni₂P-0001, which preferentially exposes the (0001) facet, induces the formation of large-sized Li₂O₂ sheet-like and porous nanostructures due to its weak LiO₂ adsorption energy and high electron transport rate. This mitigates the passivation effect caused by insulating discharge products. Therefore, the deposition behavior and morphology of discharge products can be effectively tuned by controlling the preferentially exposed facets of the Ni₂P catalyst.

The lattice matching effect is also an important strategy for regulating facet exposure properties. As illustrated in Fig. 14g, the grain-refined Co_0.85Se@CNT catalyst exhibits a (101) facet with a lattice spacing of 2.69 Å, which is highly matched with the (100) facet of Li₂O₂ (2.72 Å).¹⁶⁵ This strong lattice matching serves as a template for the epitaxial growth of Li₂O₂, promoting its oriented formation while suppressing the generation of side products. As a result, Li–O₂ batteries based on the Co_0.85Se@CNT cathode demonstrate exceptional cycling stability of over 2400 hours at a current density of 100 mA g⁻¹.

4. Innovative strategies in enhancing catalytic performance

External field-assisted strategies offer innovative approaches to overcoming the intrinsic activity limitations of cathode catalysts in Li–O₂ batteries. By introducing energy inputs such as light,^166–168 electricity,¹⁶⁹ magnetism,^41,120 or heat,⁷⁵ it is possible to dynamically regulate the local electronic structure at the catalyst interface, carrier distribution, and reaction kinetics. These modulations significantly enhance the activity and stability of both the ORR and OER. Unlike conventional chemical modifications, external fields enable real-time and directional control of the catalytic process through non-contact intervention. This approach has shown particular advantages in addressing key challenges such as high overpotentials and poor cycling stability. In recent years, the emergence of novel strategies, including photo-electrochemical synergistic catalysis, magnetothermal-induced lattice vibrations, and piezoelectric field-enhanced charge separation, has not only deepened the understanding of multiphysics coupling mechanisms but also accelerated the transition of Li–O₂ batteries from passive energy storage systems to intelligent, responsive platforms.

Photo-assisted strategies have emerged as an innovative approach to addressing Li–O₂ battery challenges, as they can enhance the kinetics of redox reactions through the synergistic effect of photo-generated electron–hole pairs.^170–174 A study on facet-controlled WO₃ photoanodes revealed that by increasing the proportion of exposed (002)/(020) crystal facets in the WO₃ photoelectrode, a controllable transition from the solution-growth mode to the surface-growth mode of Li₂O₂ formation can be achieved (Fig. 15a).¹⁷⁵ Furthermore, the Li₂O₂ layer deposited during the discharge process can form a Z-type heterojunction with WO₃. This photo-assisted discharge process occurs in three stages, as shown in Fig. 15b: (I) initial stage: the photo-generated electrons from WO₃ are consumed to reduce O₂, forming a thin Li₂O₂ film on the electrode surface; (II) thickening stage: the Li₂O₂ layer formed interacts with WO₃ to create a Z-type heterojunction, which promotes exciton dissociation and generates free electrons, further accumulating discharge products; (III) termination stage: as Li₂O₂ continues to accumulate, its enhanced exciton binding energy inhibits carrier separation, ultimately leading to the termination of the discharge process. Ultimately, the WO₃ photoanode, dominated by the (002) facet, achieves an extremely low overpotential of 0.07 V and maintains high cycling stability for over 200 hours.

Fig. 15 (a) Facet-controlled Li₂O₂ growth pathways and the photocatalytic activity of WO₃ in photo-assisted Li–O₂ batteries. (b) Proposed mechanism of the photo-assisted discharge process, consisting of three distinct stages.¹⁷⁵ (c) Reaction mechanism of the AB-OV/CC photo-cathode in photo-assisted Li–O₂ batteries. (d) Discharge/charge profiles of AB-OV/CC and AB/CC photo-cathodes at a current density of 50 mA g⁻¹ under illuminated (solid lines) and dark (dashed lines) conditions.¹⁷⁶ (e) Schematic illustration of directional charge transfer induced by the internal electric field at the Fe₂O₃/C₃N₄ interface.¹⁷⁷

Interface engineering and heterojunction construction have further enriched the electronic modulation strategies for photocatalysts. A composite cathode containing oxygen vacancies, Ag/Bi₂MoO₆ (denoted as AB-OV/CC), demonstrates enhanced full-spectrum light absorption from ultraviolet to visible regions through the synergistic effects of Ag's localized surface plasmon resonance (LSPR) and the oxygen vacancies in Bi₂MoO₆.¹⁷⁶ The LSPR-generated hot electrons from Ag are injected into the oxygen vacancies band of Bi₂MoO₆, rather than its conduction band, enabling efficient charge separation (Fig. 15c). The accelerated oxygen reaction kinetics lead to the formation of amorphous Li₂O₂, resulting in an elevated discharge plateau voltage of 3.05 V. During charging, photoexcited holes rapidly decompose the amorphous Li₂O₂, lowering the charge plateau voltage to 3.25 V. This enables a high initial round-trip efficiency of 93.8% (Fig. 15d). An S-type heterojunction photoelectrode based on Fe₂O₃/C₃N₄ was constructed, in which interfacial Fe–N bonds (bond length: 2.42 Å) and an internal electric field (IEF) synergistically promote directional separation of photo-generated charge carriers.¹⁷⁷ This structure establishes a unique electron transport pathway with spatially separated redox centers, where photoelectrons are concentrated on C₃N₄ and photogenerated holes on Fe₂O₃. During discharge, electrons in the conduction band of C₃N₄ reduce O₂ to form LiO₂ intermediates, which further convert into Li₂O₂, while holes in the valence band of Fe₂O₃ participate in the oxidation reaction. During charging, holes in Fe₂O₃ oxidize Li₂O₂ to O₂, while electrons in C₃N₄ reduce Li⁺ to metallic lithium (Fig. 15e). This design achieves a high discharge voltage of 3.13 V and a low charge voltage of 3.19 V, resulting in a round-trip efficiency approaching 98%.

Photocatalytic technologies also offer new perspectives for dynamically regulating discharge products in Li–O₂ batteries. By employing a cathode composed of light-responsive gold nanoparticles exhibiting SPR, combined with reduced graphene oxide (Au/rGO), recent studies have revealed the critical role of matching ORR kinetics with lithium-ion transport rates in determining battery performance.¹⁷⁸ Under light irradiation, this system significantly alleviates the phenomenon of “sudden battery death” caused by Li⁺ concentration polarization, introducing a novel concept of dynamic control. Furthermore, a porphyrin-based MOF incorporating (Fe₂Ni)O(COO)₆ clusters has been used as a photo-cathode to substantially enhance the performance of photo-assisted Li–O₂ batteries.¹⁷⁹ The coupling between Ni 3d and Fe 3d orbitals strengthens ligand-to-metal cluster charge transfer (LMCT), thereby promoting the separation of excitons into free charge carriers. This process effectively suppresses recombination of photo-generated electron–hole pairs and activates the generation of O₂⁻, rather than ¹O₂. Under light irradiation of 100 mW cm⁻², the FeNi-TCPP-based photo-assisted Li–O₂ battery achieves a remarkably low overall overpotential of 0.28 V and a round-trip efficiency of 92%.

In addition to photo-assisted catalysis,^{174,181–183} researchers have explored a broader range of external field-assisted strategies.^120,180 One such innovative approach utilizes a stress field, with BaTiO₃ (BTO) piezoelectric material as the core.¹⁶⁹ As shown in Fig. 16a, Raman scattering peaks at 249 cm⁻¹ and 306 cm⁻¹ correspond to the symmetric and asymmetric crystal structures of BTO. During discharge, Li₂O₂ products rapidly accumulate on the surface of BTO (Fig. 16b–d), and the Raman scattering peaks at 249 cm⁻¹ and 306 cm⁻¹ are significantly enhanced, indicating increased lattice distortion of BTO. This is attributed to the intrinsic stress generated by the continuous accumulation of Li₂O₂, which causes Ti⁴⁺ ions to shift along specific crystal axes toward the O²⁻ direction. Consequently, the O²⁻ ions undergo electronic polarization displacement, generating a spontaneous polarization electric field that drives carrier separation and transport. This allows for the continuous separation and accumulation of electrons and holes, while the conduction and valence bands are tilted to favor enhanced ORR and OER kinetics (Fig. 16e). This coupling strategy between intrinsic stress and electric field at the surface effectively accelerates the interfacial transport rate of Li⁺ and improves the redox reaction kinetics in Li–O₂ batteries. The study on Co₃O₄ grown on nickel foam as a cathode explores the impact of acoustic fields on the performance of Li–O₂ batteries.¹⁸⁰ Ultrasonic charging plays a positive role in enhancing mass transfer processes, suppressing concentration polarization, promoting the rapid decomposition of Li₂O₂, and increasing the exposure of active sites (Fig. 16f). Ultrasonic energy, which can be converted into mechanical and thermal energy, aids in the decomposition of Li₂O₂, exposing more active sites and enhancing the rate of the OER. Experimental results demonstrate that under ultrasonic charging conditions with a duty cycle of 5 : 5 and a power of 675 W, the overpotential is reduced by 25.9%, and the battery's cycle life is extended from the typical 67 cycles to 321 cycles.

Fig. 16 (a) In situ electrochemical Raman spectra collected on the BTO-CNT cathode at a current density of 300 mA g⁻¹. (b) SEM image of the BTO-CNT cathode at a discharge depth of 1500 mA h g⁻¹ and a current density of 300 mA g⁻¹. (c and d) Low- and high-magnification SEM images of the BTO-CNT cathode at a discharge depth of 3000 mA h g⁻¹ under a current density of 300 mA g⁻¹. (e) Schematic illustration of Li₂O₂ formation and decomposition on the BTO surface.¹⁶⁹ (f) Schematic diagram of the ultrasonic charging mechanism on the electrode surface.¹⁸⁰

Magnetic field-assisted strategies have shown promising potential among multi-field coupling technologies due to the unique Lorentz force effect, which can suppress the rapid recombination of photoexcited electrons and holes. A three-dimensional porous NiO nanosheet on nickel foam (NiO/FNi) was employed as a photoelectrode.¹²⁰ As a p-type semiconductor, NiO/FNi generates photoinduced electrons and holes under illumination. As illustrated in Fig. 17a, electrons in the conduction band (CB, 2.04 V vs. Li⁺/Li) facilitate the ORR, while holes in the valence band (VB, 5.16 V vs. Li⁺/Li) accelerate the oxidation of Li₂O₂ during the OER, thereby reducing both discharge and charge overpotentials. When an external magnetic field (5 mT) is applied, the Lorentz force acts in opposite directions on the photoexcited electrons and holes, effectively suppressing their recombination (Fig. 17b). Serving as an integrated bifunctional photo-cathode, NiO/FNi combines magnetic effects with photoelectric conversion capabilities, enabling effective regulation of internal electrochemical properties, charge and lithium-ion transport, and redox reaction kinetics within the battery (Fig. 17c). The device operates stably at an ultra-low charging voltage of 2.73 V and achieves a high energy efficiency of 96.7% along with excellent cycling stability.

Fig. 17 (a) UV-vis absorption spectrum of NiO/FNi. Inset: energy band diagram of NiO/FNi. (b) Schematic illustration of the induced electromotive force (EMF) in NiO under relative motion within a magnetic field, along with a COMSOL simulation of the surface EMF distribution on NiO under magnetic assistance. (c) Schematic diagram of the working mechanism of a Li–O₂ battery under combined light and magnetic field assistance.¹²⁰

Field-assisted strategies, by dynamically modulating reaction kinetics and interfacial behaviors in batteries, hold theoretical promise for overcoming the intrinsic performance limitations of conventional catalysts.^184,185 However, their practical implementation still faces multiple challenges, including device integration, energy efficiency, cost control, and engineering scalability. At present, the coupling of external fields with electrochemical systems is not yet suitable for large-scale fabrication or industrial deployment. For instance, magnetic field-assisted strategies often require embedding permanent magnets or electromagnetic coils within the battery, which not only increases the device's volume and weight but may also induce variations in reaction site selectivity due to magnetic field gradients within porous electrodes. This necessitates complex magnetic shielding or gradient field designs for effective control. In the case of light-field strategies, the introduction of transparent windows (e.g., quartz or flexible polymers) into the battery casing can compromise system sealing, increasing the risk of electrolyte leakage and ambient gas permeation. Moreover, the uneven distribution and attenuation of light within porous structures can result in non-uniform activation of catalytic sites, limiting the overall reaction efficiency. Traditional photocatalysts are also prone to activity loss under prolonged illumination due to lattice oxygen depletion induced by photogenerated holes. More critically, the application of external fields typically requires additional energy inputs, such as light sources or electromagnetic devices, leading to relatively high power consumption and poor overall energy conversion efficiency.^186,187 Therefore, key directions for future research include the design of highly responsive materials, the development of energy recovery mechanisms, and the use of green energy sources to power the external fields, all aimed at improving system efficiency and promoting practical viability.

5. Conclusion and future perspectives

In recent years, significant progress has been made in the design and performance optimization of cathode catalysts for Li–O₂ batteries, laying a critical foundation for enhancing energy efficiency, cycling stability, and practical application potential. This review begins with an overview of the basic reaction mechanisms in Li–O₂ batteries, systematically summarizing the key role of cathode catalyst electronic properties in modulating the ORR and OER kinetics, as well as the corresponding structure–activity relationships based on current core research findings and innovative strategies. Through breakthroughs in DFT calculations and in situ characterization techniques, research has revealed the intrinsic relationships between catalyst electronic structures, such as the metal d-band center, non-metallic p-orbital sites, e_g orbital occupancy, orbital hybridization, optimization of exposed facets, and external field-assisted strategies. These insights progressively highlight the close connection between catalyst electronic structure and reaction activity, providing important theoretical support for the advancement of Li–O₂ battery technology and the rational design of efficient cathode catalysts. Notably, innovative strategies, including external field-assisted technologies such as light, sound, electricity, and magnetism, offer new dimensions for improving catalytic efficiency, achieving performance breakthroughs beyond traditional thermodynamic limitations.

Despite the remarkable progress in the design of cathode catalysts for Li–O₂ batteries, several critical areas still require further investigation to advance the practical application of this technology.

Future research should focus on the precise modulation of catalyst electronic structures at the molecular and atomic levels, with the goal of establishing quantitative structure–activity models that correlate adsorption free energy, electron transfer numbers, and catalytic activity. Such models will provide a fundamental basis for enhancing the kinetics of the ORR and OER and guide the rational design of high-performance cathode catalysts. Electronic structure optimization should not be limited to the modification of individual materials; instead, future efforts can explore the integration of different catalytic components. Designing composite materials or heterointerfaces to leverage synergistic effects between components can further improve catalytic performance. For example, metal–carbon composite materials can significantly enhance electrical conductivity and reaction kinetics. Moreover, emerging strategies such as external field-assisted catalysis and solid-state battery systems have shown great potential for improving both catalytic efficiency and long-term stability.

Although correlations between the electronic states of cathode catalysts and their catalytic activity can be established, such structure–activity relationships do not fully capture the actual reaction mechanisms. Under operating conditions, the nanoscale structures and coordination environments of active sites often undergo dynamic changes, diminishing the relevance of performance relationships derived solely from the catalyst's initial state. Therefore, the integration of advanced in situ characterization techniques is essential. For example, in situ X-ray absorption fine structure (XAFS) spectroscopy can identify the true active sites on electrode surfaces; in situ SERS enables the investigation of interfacial reactions and the capture of reaction intermediates during charge–discharge cycles; in situ UV-vis absorption spectroscopy allows monitoring of ¹O₂ dynamics and associated parasitic reactions. Additionally, TOF-SIMS provides precise compositional and three-dimensional spatial analysis of discharge products. Moreover, the development of theoretical modeling methods under realistic working conditions, such as molecular dynamics simulations, enables real-time tracking of the evolution of catalyst surfaces, active sites, and intermediates during reactions. These approaches help unravel reaction mechanisms and the dynamic nature of active sites. By integrating theoretical simulations with experimental studies, a deeper understanding of the formation and decomposition kinetics of Li₂O₂ can be achieved, offering a solid theoretical foundation for the design of next generation high efficiency catalysts.

The optimization of cathode, electrolyte, and anode interfaces must be pursued in a coordinated manner, encompassing the development of stable solid–liquid interfaces, enhancement of ionic conductivity, and improvement of interfacial durability and resistance to contamination. For example, electrolytes with low donor numbers tend to form thin Li₂O₂ films, making them suitable for applications requiring high stability, whereas high donor number electrolytes promote the formation of Li₂O₂ toroids, which are beneficial for achieving high discharge capacities. Therefore, it is essential to develop electrolytes that simultaneously offer high solubility and high diffusion coefficients. Catalyst degradation induced by reactive oxygen species, such as O₂⁻ and ¹O₂, along with the accumulation of side products (Li₂CO₃ and LiOH), can significantly reduce charging efficiency and increase energy loss. Designing catalysts capable of converting ¹O₂ into its less reactive triplet state (³O₂) is key to suppressing these undesirable side reactions. The lithium metal anode, due to its high reactivity, readily reacts with electrolytes and oxygen, leading to dendrite formation and increased risk of short circuits. Strategies such as constructing artificial solid electrolyte interphase (ASEI) layers, introducing lithium host materials, and optimizing lithium stripping/plating mechanisms have proven effective in improving the cycling stability and safety of the anode.

Data availability

The data that support the findings of this study are available from the corresponding author upon reasonable request. All relevant data generated or analyzed during this study have been properly cited and referenced within the manuscript. Any additional data or materials can be made accessible to interested researchers, in accordance with the data sharing policies of EES Batteries.

Conflicts of interest

There are no conflicts to declare.

Acknowledgements

This work was supported by the Science and Technology Development Fund (FDCT) of Macao S.A.R (0033/2023/ITP1, 0022/2023/RIB1, 046/2019/AFJ and 0070/2023/AFJ), the Multi-Year Research Grants (file no. MYRG2022-00223-IAPME and MYRG-GRG2024-00166-IAPME) from the Research Services and Knowledge Transfer Office at the University of Macau.

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