Recent advances in metal–organic frameworks for Li–O₂ batteries: advantages, challenges, and innovative design

Abstract

Metal–organic frameworks (MOFs) have emerged as promising candidates for cathode materials in Li–O₂ batteries due to their exceptional properties such as high surface area, periodic porous structures, tunable pore sizes, and controllable chemical compositions. This review highlights the innovative designs of MOFs, focusing on the central metal contributing to the oxygen reduction reaction (ORR)/oxygen evolution reaction (OER) kinetics. In addition, bimetallic/heterometallic node engineering, the structural optimization of MOFs, impact of organic ligand design, defect introduction, and hybrid material strategies are discussed, each providing new avenues to improve catalytic activity, electron/mass transport, and long-term durability. With the advancement of photo-assisted Li–O₂ batteries and the semiconducting properties of MOF materials, this review also underscores the application of MOFs in photo-assisted Li–O₂ batteries. Furthermore, the latest progress in MOF-based solid-state electrolytes for Li–O₂ batteries is covered, demonstrating how their hierarchical porous structure and strong mechanical/chemical stability substantially boost lithium-ion conductivity and battery safety. We have also incorporated a dedicated discussion to highlight recent advances in the preparation methods for MOFs and MOFs-derived materials for cathode materials in Li–O₂ batteries, emphasizing the practical use in this field. Finally, future research directions are proposed, stressing the necessity for integrated approaches that combine MOFs with other materials to further enhance the efficiency and longevity of Li–O₂ batteries.

---

Wider impact

This review discusses recent breakthroughs in the application of metal–organic frameworks (MOFs) for Li–O₂ batteries, covering their roles as advanced cathode catalysts, photo-assisted materials, and solid-state electrolytes. Key developments highlighted include the rational design of MOF metal centers and organic ligands for improved catalytic activity, introduction of hierarchical porous architectures for enhanced mass transport, and creation of MOF-based hybrids and semiconductor composites that push the boundaries of electrochemical performance and stability. This area is of broad significance as Li–O₂ batteries promise energy densities far beyond current lithium-ion technology, potentially transforming large-scale renewable energy storage, electric vehicles, and portable electronics towards a more sustainable future. The field's evolution towards photo-assisted processes and robust solid-state electrolytes directly addresses challenges of intermittency and safety, advancing the prospects for grid integration of renewables. Insights from this review-particularly the molecular-level tailoring of MOF structures and interfaces-will guide next-generation materials design, accelerating the practical deployment of high-capacity, long-life, and safe energy storage devices, and thus play a pivotal role in shaping the future landscape of materials science and green energy technologies.

1. Introduction

Over the past few decades, the exploitation of fossil fuels and the large-scale popularization of electricity have profoundly transformed human life. Currently, over 70% of the global energy demand is met by fossil fuels, predominantly coal, oil, and natural gas. However, the combustion of these fossil fuels and the subsequent release of greenhouse gases and pollutants (e.g., CO₂, SO₂, and NO_x) have triggered a cascade of environmental crises, including global warming, severe air pollution, and ecological degradation.¹ These issues have underscored the urgent need to reduce our reliance on fossil fuels and accelerate the transition to renewable energy sources, such as wind and solar power. Implementing this energy transition hinges on robust technological progress and support. However, renewable energy sources like wind and solar power are inherently intermittent and unstable due to their dependence on seasonal variations, weather conditions, and geographical factors. This intermittency poses significant challenges to the stability of the power grid, necessitating the development of effective energy storage systems to balance supply and demand. In this context, electrochemical power sources and energy storage systems have emerged as vital solutions, bridging the gap between renewable energy generation and grid integration. Among these systems, lithium-ion batteries (LIBs) have achieved remarkable commercial success, revolutionizing portable electronics due to their lightweight design and long-lasting energy supply. Nevertheless, the global push for electric vehicles has set new benchmarks for battery energy density. Current LIBs, with a maximum specific energy of approximately 250 Wh kg⁻¹, still fall short of matching the energy density and range capabilities of traditional gasoline-powered internal combustion engines.² This shortfall highlights the urgent need for next-generation battery systems with significantly higher energy density to meet the evolving demands of modern transportation and energy storage.

The concept of lithium–air batteries was first proposed in 1976 by Littauer and Tsai, who explored their potential in alkaline aqueous solutions with organic solvent additives. A major breakthrough came in 1996, when Abraham et al.³ developed the first non-aqueous lithium–air battery by using a carbon–cobalt composite cathode and an organic polymer electrolyte. They also speculated that the discharge product of this battery is likely to be lithium peroxide. The concept of lithium–air batteries has revolutionized the traditional understanding of battery systems. Among them, lithium–oxygen batteries, with their ultra-high specific energy density (5200 Wh kg⁻¹), are pushing the energy density of batteries to a new height.⁴ The unique architecture of Li–O₂ batteries consists of a lithium metal anode, a metal salt electrolyte, and an “oxygen/air” cathode. During charge and discharge cycles, the cathode continuously absorbs oxygen from the surrounding air, which reacts with lithium ions from the anode, enabling continuous energy conversion. From 2002 to 2005, the research conducted by Read et al.^5–7 gradually illustrated that the discharge capacity, rate performance, and cyclability of Li–O₂ batteries are largely dependent on the electrolyte solution formulation and the material used for the air cathode. In 2006, Ogasawara et al.⁸ provided the first evidence of the reversible formation of Li₂O₂ in Li–O₂ batteries, opening up a new chapter of the research.

Despite their potential as next-generation energy storage systems, Li–O₂ batteries still face several significant challenges that hinder their practical application. First, the sluggish kinetics of the oxygen reduction reaction (ORR) and oxygen evolution reaction (OER) lead to slow formation and decomposition of Li₂O₂. This issue is further exacerbated by the lack of stable and efficient bifunctional catalysts for the cathode, resulting in poor rate performance and low round-trip efficiency.⁹ Second, the complex multiphase (solid–liquid–gas) reaction interface at the cathode complicates the reaction kinetics.¹⁰ The discharge product, Li₂O₂, is both insulating and insoluble, which can block active sites and gas channels. This blockage increases the charging overpotential and triggers side reactions that degrade the electrolyte and carbon-based materials. Third, the instability of cathode catalysts during charge–discharge cycles further diminishes cycling stability, as irreversible chemical reactions or phase transformations often occur.

In recent years, metal–organic frameworks (MOFs) have emerged as a promising class of materials in energy storage systems, including LIBs, Li–S batteries and Li–O₂ batteries.¹¹ MOFs are crystalline porous materials constructed from metal ions coordinated with organic linkers, forming highly ordered and tunable frameworks.¹² The porous structure of MOFs facilitates extensive contact with electrolytes and active species, providing abundant active sites and ensuring efficient electron and mass transfer during electrochemical reactions, make them ideal candidates for enhancing battery performance.¹³ In 2014, Wu et al.¹⁴ pioneered the use of MOFs as cathodes in Li–O₂ batteries. This breakthrough heralded a new era in the application of MOFs in energy storage systems. The inherent advantages of MOFs lie in their ability to systematically arrange metal cations, creating highly dispersed catalytic active sites that enhance both catalytic activity and stability.¹⁵ Their tunable nature allows for the design of bimetallic systems, which can serve as highly efficient bifunctional catalysts for ORR and OER, outperforming traditional metal oxides or noble metals.¹⁶ The precise engineering of pore size, morphology, and distribution further enables the creation of hierarchical porous architectures, which accommodate discharge products while maintaining optimal contact with the electrolyte and oxygen.^13,17,18 Furthermore, MOFs have also been explored as semiconductor materials in photo-assisted Li–O₂ batteries. Their intrinsic semiconducting properties, combined with strategies such as constructing multimetallic MOFs to broaden light absorption ranges and forming heterojunctions with other semiconductors, have opened new avenues for enhancing ORR/OER activity in photocatalytic processes.^19–21 Beyond their role as cathode materials, MOFs have been utilized as electrolyte materials due to their hierarchical porous architecture, wide electrochemical window, and robust mechanical properties. Their porous structure and tunable chemical composition significantly improve lithium-ion transport efficiency, while their ability to adsorb anions through metal active centers enhances lithium-ion transference numbers.^22–24

This review begins with a concise overview of the fundamental working principles of Li–O₂ batteries and their photo-assisted counterparts. It then provides a detailed analysis of the applications of MOFs in Li–O₂ batteries, categorizing their roles as cathode materials, photocathodes, and electrolyte materials: (1) Central metal effect: the catalytic activity of MOFs is primarily driven by their central metal ions, which effectively catalyze the ORR and OER processes, improving overall reaction kinetics. (2) Structural optimization of MOFs: the highly ordered and tunable porous structures of MOFs, combined with functionalized cavities, create an optimal environment for catalytic reactions, enhancing both activity and stability. (3) Directed design of organic ligands: the deliberate design of organic ligands in MOFs further promotes catalytic reactions, enabling optimization for specific applications. (4) Strategies for MOF hybrids: rational integration of MOFs with functional components enhances charge/mass transport and stability, boosting Li–O₂ battery performance. (5) Photoactive MOF components for photo-assisted Li–O₂ batteries: the role of MOFs as semiconductor materials in photo-assisted Li–O₂ batteries is explored, highlighting their potential to enhance ORR/OER activity through strategies such as constructing multimetallic MOFs and forming heterojunctions with other semiconductors. (6) MOFs in solid-state electrolytes: the role of MOFs as electrolyte materials in Li–O₂ batteries is comprehensively discussed, emphasizing their hierarchical porous architecture, wide electrochemical window, and robust mechanical properties, which significantly improve lithium-ion transport efficiency and interfacial stability. (7) Other supplementary content: the preparation methods for MOFs and MOFs-derived materials for cathode materials in Li–O₂ batteries. By leveraging their unique properties, MOFs hold great promise for advancing the performance and longevity of next-generation Li–O₂ battery systems.

2. Charge and discharge reactions of Li–O₂ batteries

2.1 Li–O₂ batteries with Li₂O₂ as the discharge product

Generally speaking, the fundamental reactions of a Li–O₂ battery can be described by eqn (1).

2Li⁺ + O₂ + 2e⁻ ⇌ Li₂O₂ (1)

Typically, in a pure oxygen environment, Li₂O₂, an insulating and insoluble solid, is generally recognized as the primary discharge product in Li–O₂ batteries. During the discharge process, lithium metal at the anode is oxidized, generating Li⁺ ions in the electrolyte and releasing electrons into the external circuit. Meanwhile, O₂ at the cathode undergoes reduction to form Li₂O₂. The charging process is the reverse, where Li₂O₂ is oxidized to release O₂. To elaborate, the discharge mechanism involves the initial adsorption of O₂ molecules onto the cathode surface, where they capture an electron to form O₂˙⁻. This O₂˙⁻ then reacts with Li⁺ ions in the electrolyte to produce LiO₂.²⁵ The mainstream view nowadays is that in high donor number (high DN) electrolytes, LiO₂ tends to dissolve and subsequently undergoes a disproportionation reaction in solution to yield Li₂O₂ as eqn (2).²⁶

LiO_2(sol) + LiO_2(sol) → Li₂O_2(sol) + O₂ (2)

In low donor number (low DN) electrolytes, LiO₂ is more likely to adsorb onto the electrode surface (* denotes the adsorbed state). It then undergoes electrochemical reduction (eqn (3)) or disproportionation (eqn (4)) to form Li₂O₂ via a surface growth pathway.²⁷

LiO₂* + Li_(sol)⁺ + e⁻ → Li₂O₂ (3)

2LiO₂* → Li₂O₂ + O₂ (4)

The issue of Li₂O₂ decomposition into Li⁺ and O₂ during the charging process remains a topic of ongoing debate. In early studies, it was proposed that the decomposition of Li₂O₂ occurred via a direct two-electron process without the formation of intermediates (eqn (5)).²⁸

Li₂O₂ → 2Li⁺ + 2e⁻ + O₂ (5)

However, the high activation energy associated with the 2e⁻ electrochemical reaction pathway, combined with the poor electronic conductivity of Li₂O₂, makes the decomposition of Li₂O₂ during the charging process appear implausible. Ganapathy et al. demonstrated through experimental characterization and theoretical calculations that, during the initial stage of charging, electrons and lithium ions are preferentially removed from Li₂O₂ to form a lithium-deficient solid solution. The formation of a relatively conductive Li_2−xO₂ intermediate follows, which is subsequently oxidized to release O₂ (eqn (6) and (7)).²⁹

Li₂O₂ → Li_2−xO₂ + xLi⁺ + xe⁻ (6)

Li_2−xO₂ → (2 − x)Li⁺ + (2 − x)e⁻ + O₂ (7)

This stepwise oxidation and decomposition mechanism of Li₂O₂ is more kinetically feasible compared to the direct pathway.

2.2 Li–O₂ batteries with LiOH as the discharge product

It is worth noting that the discharge product Li₂O₂ is highly sensitive to moisture and carbon dioxide in the air, making it more likely to be converted into the more stable LiOH and eventually into Li₂CO₃.³⁰ In earlier studies, the formation of LiOH as a secondary discharge product was often attributed to parasitic reactions between reactive oxygen intermediates (such as O₂, LiO₂, Li₂O₂, and HO₂) and organic electrolyte components.^29,31–34 Adding water to the non-aqueous electrolyte of Li–O₂ batteries can promote the conversion of Li₂O₂ to LiOH (2Li₂O₂ + 2H₂O ⇌ 4LiOH + O₂). However, without a suitable catalyst, this process is only partially complete.^35,36 Since the first proposal of Li–O₂ batteries based on LiOH discharge products in 2015, numerous studies have been reported on this topic.^37,38 Recent research has demonstrated the feasibility of reversible cycling in non-aqueous electrolytes through LiOH formation/decomposition (4e⁻ O₂ reduction [ORR]/evolution reaction [OER]).^39,40

Liu et al.³⁷ proposed a reaction mechanism for the formation of LiOH involving the participation of LiI redox mediators. During the discharge process, O₂ is reduced on the electrode surface via an electrochemical reaction to form LiO₂. LiO₂ subsequently undergoes a chemical reaction through a solution mechanism to transform into LiOH. This transformation occurs because LiOH is observed to grow on the cathode and the insulating glass fiber separator, neither of which involves electron injection.

4Li⁺ + 4O₂ + 4e⁻ → 4LiO₂ electrochemical (8)

As previously discussed, LiOH can be formed through the 4e⁻/O₂ ORR (E⁰ = 3.32 V), which endows it with higher chemical stability compared to LiO₂ and Li₂O₂. However, the decomposition of LiOH involves the reorganization of O–O bonds, rendering it more kinetically challenging compared to the decomposition of Li₂O₂ to release O₂. Considering the use of LiI redox mediators, Liu et al.³⁷ revealed that during the charging process, I⁻ is initially directly electrochemically oxidized to I₃⁻. Subsequently, LiOH is chemically oxidized by I₃⁻, resulting in the formation of O₂ and H₂O.

6I⁻ → 2I₃⁻ + 4e⁻ electrochemical (10)

4LiOH + 2I₃⁻ → 4Li⁺ + 6I⁻ + 2H₂O + O₂ chemical (11)

Additionally, inspired by enzyme-catalyzed oxygen reduction/oxi-dation reactions, Wang et al.⁴⁰ introduced the copper(I) complex 3N–Cu^I (3N = 1,4,7-trimethyl-1,4,7-triazacyclononane) into Li–O₂ batteries and proposed a reaction mechanism involving 3N–Cu^I-catalyzed processes based on O–O bond breaking and formation, with LiOH as the primary discharge product. These studies demonstrate that the decomposition of LiOH to generate O₂ is feasible in several catalytic systems thereby highlighting the potential for designing rechargeable secondary batteries with LiOH as the primary discharge product. Future research will necessitate more detailed mechanistic investigations to elucidate how factors such as water content and electrolyte solvent influence the efficiency of the 4e⁻/O₂ ORR/OER catalytic reactions.

2.3 Photo-assisted Li–O₂ batteries

Despite the ultra-high theoretical energy density of Li–O₂ batteries, the slow kinetics of the Li₂O₂ redox reactions and the resulting large overpotentials have significantly hindered their development.⁴¹ In 2014, Wu et al. first proposed that TiO₂, a semiconductor material, could be used to catalyze the photoelectrochemical oxidation of I₃⁻ to I⁻ when excited by solar energy. These ions subsequently participate in the oxidation of Li₂O₂, effectively reducing the charging overpotential of Li–O₂ batteries.⁴² As an integrated device, the photo-assisted Li–O₂ battery utilizes photocatalysts to harvest solar energy and convert it into chemical energy stored via redox reactions at the cathode. From a thermodynamic standpoint, the suitability of a semiconductor for photo-assisted Li–O₂ batteries is primarily determined by the positions of its conduction band (CB) and valence band (VB). The working principle is illustrated in Scheme 1.

Scheme 1 The theoretical working principle of photo-assisted Li–O₂ batteries.

For semiconductor materials, the corresponding factors are the lowest unoccupied molecular orbital (LUMO) and highest occupied molecular orbital (HOMO) energy levels. The LUMO and HOMO energy levels should be positioned relative to the theoretical equilibrium potentials of the O₂/Li₂O₂ redox couple such that the excited e⁻ and h⁺ possess enhanced reducing and oxidizing capabilities. This enables their participation in the ORR and OER reactions, thereby altering the reaction kinetics in Li–O₂ batteries.⁴³ Taking MOFs materials as an example, upon light excitation, MOFs generate e⁻ and h⁺. The electrons in the LUMO orbit reduce O₂ to O₂⁻, which subsequently combines with Li⁺ to form LiO₂, eventually converting to Li₂O₂. Meanwhile, the holes in the HOMO orbit continuously recombine with electrons from the external circuit. During the charging process, the more oxidizing holes in the HOMO orbit decompose Li₂O₂ into O₂ and Li⁺, while the electrons in the LUMO orbit travel through the external circuit to the anode, where they reduce Li⁺ back to Li. The relevant photoelectrochemical reactions are described as follows:

Discharge process:

O₂ + e_(LUMO)⁻ + Li⁺ → LiO₂ (13)

LiO₂ + Li⁺ + e_(LUMO)⁻ → Li₂O₂ (14)

Li → Li⁺ + e⁻ (15)

Charge process:

Li₂O₂ + 2h_(HOMO)⁺ → O₂ + 2Li⁺ (17)

Li⁺ + e⁻ → Li (18)

2.4 MOFs for Li–O₂ batteries

The operational principles of Li–O₂ batteries dictate that cathode materials should possess high specific surface area and excellent electrical conductivity to ensure abundant catalytic active sites, efficient mass transport, and effective storage of discharge products. Unlike traditional monofunctional catalysts combined with carbon materials, MOFs offer significant advantages due to their tunable metal active sites, uniformly dispersed catalytic centers, and structural stability, making them promising candidates for cathode catalysts in Li–O₂ batteries.¹⁴ As illustrated in Fig. 1, metal nodes usually serve as the active sites and organic ligands act as framework in MOF materials. In line with the Sabatier principle,⁴⁴ appropriate electronic coupling is more favorable for accelerating reaction kinetics. The tunability of MOFs allows for precise design to improve the ORR and OER kinetics, including the coordination structure and the incorporation of bimetallic or trimetallic centers.⁴⁵ Their hierarchical porous structure enable numerous active sites to make full contact with O₂ and the electrolyte in Li–O₂ batteries, ensuring the continuous progress of the catalytic reactions.⁴⁶

Fig. 1 MOFs for Li–O₂ batteries: (1) adsorbing oxygen; (2) central metal effect for ORR and OER; (3) MOFs-derived materials; (4) MOFs in solid-state electrolytes; (5) strategies for MOF hybrids; (6) photoactive MOF components for photo-assisted Li–O₂ batteries.

Most MOFs are semiconductor materials and typically require compounding with conductive carbon materials when used as cathodes. Only a few specific ligand structures possess inherent conductivity, while in situ growth of MOFs on conductive substrates is a relatively effective method for preparing Li–O₂ battery cathodes. Compared to pristine MOFs, pyrolyzed MOF-derived materials that retain some structural characteristics exhibit significantly enhanced electrical conductivity, resulting in superior electrocatalytic performance. Notably, the semiconducting properties of MOFs enable them to serve not only as traditional cathode catalysts in Li–O₂ batteries but also as key components in photo-assisted Li–O₂ batteries. By leveraging their light-responsive properties, photoexcited electrons and holes can be utilized to increase discharge voltage or reduce charging potential. Their light absorption range and band structure can further be tuned by modifying metal clusters and ligands. Additionally, their photoactivity can be enhanced by encapsulating photoactive materials within their pores or constructing heterojunctions to further improve light-responsive performance.

With the development of high-performance solid-state electrolytes (SSEs), solid-state electrolytes based on porous materials (such as MOFs) are increasingly attracting attention.⁴⁷ The ion transport mechanisms of MOF solid-state electrolytes can be summarized into two points: one is the MOF-based single-ion conductor, which lacks ion solvation, and cation hopping occurs in its orderly close packing and periodic channels.⁴⁸ The other is to activate MOFs with ionic liquids (IL) or immerse them in polymer electrolytes. Since the ionic conductivity in MOFs depends on the presence of certain liquid electrolytes to further facilitate ion conduction, the ion conduction mechanism in MOFs is sometimes referred to as the coexistence of the liquid carrier/Grotthuss mechanism and the single-ion hopping/cooperative diffusion mechanism.⁴⁹

3. MOFs as electrocatalysts for the cathode

3.1 Central metal effect

For Li–O₂ batteries, the overall performance is largely determined by the kinetics of ORR and OER on the cathode. Highly dispersed active sites are crucial for accelerating reaction rates and improving efficiency, while the hierarchical porous structure of the electrode facilitates mass transfer and provides sufficient storage space for the solid discharge product of Li₂O₂. The uniformly dispersed metal active sites and designable bimetallic active sites of MOF-based catalysts, which exhibit different activities in various catalytic reactions,⁵⁰ makes MOFs key materials for advancing metal–air battery technology.

In 2014, Wu et al.¹⁴ pioneered the use MOFs as cathodes in Li–O₂ batteries. The research team selected three typical MOFs as research objects: MOF-5 (Zn₄O(BDC)₃, BDC: 1,4-benzenedicarboxylic acid), HKUST-1 (Cu₃(BTC)₂, BTC: 1,3,5-benzenetricarboxylic acid), and M-MOF-74 (M₂(DOT), M = Mg, Mn, Co, DOT: 2,5-dihydroxyterephthalic acid), as shown in Fig. 2(a). Low-temperature O₂ adsorption experiments showed that MOFs with open metal sites (such as HKUST-1 and the M-MOF-74 series) had significantly higher O₂ uptake (15.0–18.4 mg g⁻¹) at 1 atm compared to MOF-5 without open sites (6.6 mg g⁻¹) and pure Super P (<0.7 mg g⁻¹). Notably, the O₂ concentration within the pores of Mn-MOF-74 reached 18 times that of pure oxygen, attributed to the strong adsorption of O₂ by Mn²⁺ sites. Under a current density of 50 mA g⁻¹, the Mn-MOF-74/Super P composite electrode achieved an initial discharge capacity of 9420 mAh g⁻¹, 4.3 times that of pure Super P (2170 mAh g⁻¹) and far exceeding other MOFs like HKUST-1 (4170 mAh g⁻¹), as shown in Fig. 2(b). DEMS analysis revealed O₂ release as the dominant reaction with minimal CO₂ generation, confirming Li₂O₂ decomposition as the major pathway. XRD showed the MOFs retained their crystalline structure after cycling. When H₂O occupied the Mn sites (forming Mn₂(DOBDC)(H₂O)₂), discharge capacity dropped to 2820 mAh g⁻¹, and non-porous Mn(II) acetate yielded only 1920 mAh g⁻¹. This work pioneers the application of MOFs as cathode materials in Li–O₂ batteries, offering a new paradigm for designing next-generation high-energy-density batteries.

Fig. 2 (a) Schematic illustration of MOFs based Li–O₂ batteries. (b) Discharge profiles of the Li–O₂ batteries with various MOFs cathodes.¹⁴ Reproduced with permission from ref. 14. Copyright 2014, Wiley-VCH.

Seeking to improve the catalytic activity of Mn-MOF-74, Kim et al.⁵¹ utilized a straightforward hydrothermal synthesis approach to fabricate the bimetallic MnCo-MOF-74. In initial full discharge and charge tests conducted at a current density of 200 mA g⁻¹, MnCo-MOF-74 based Li–O₂ batteries achieved an impressive discharge capacity of 11 150 mA g⁻¹, significantly outperforming monometallic Mn-MOF-74 (6040 mAh g⁻¹) and Co-MOF-74 (5630 mAh g⁻¹). Furthermore, the MnCo-MOF-74 bimetallic system demonstrated markedly improved cycle life relative to its monometallic counterparts. Both the XPS and XRD analysis revealed that the discharge products of the MnCo-MOF-74 cathode are a mixture of Li₂O₂ and LiOH. In contrast, the discharge products of Mn-MOF-74 are a mixture of Li₂CO₃ and LiOH, with LiOH still present after charging. These exceptional electrochemical performances can be attributed to the synergistic effects of the bimetallic active centers of Mn and Co, which effectively modulate the catalytic activity and enhance the overall performance of the material.

By optimizing the bimetallic ratio, the catalytic performance of the bimetallic MOF can be further improved. Employing microwave reaction techniques, Wang et al.⁵² successfully synthesized NiMn-MOF where both Mn and Ni cations form octahedral coordination with the O anion in terephthalic acid. Theoretical calculations revealed that Mn doping in NiMn_0.05-MOF induced electronic coupling with Ni sites, increasing Ni–O covalency and activating Ni centers. This enhanced LiO₂/Li₂O₂ adsorption and reduced reaction barriers by facilitating electron transfer with oxygen species. The NiMn_0.05-MOF cathode exhibited the lowest charge–discharge overpotential and an ultralong cycle life of up to 540 cycles. In contrast, the Ni-MOF and NiMn_0.1-MOF electrodes displayed lower cycle lives of 210 and 247 cycles, respectively. These results clearly indicate that the excellent catalytic activity of NiMn_0.05-MOF for ORR/OER reactions contributes to its optimized Mn amount. In addition, Miao et al.⁵³ synthesized the bimetallic MOF Ni_1.5Cu_1.5(HHTP)₂ on carbon paper via a simple hydrothermal method. At a current density of 50 mA g⁻¹, Ni_1.5Cu_1.5(HHTP)₂/CP achieves a discharge capacity of 8367 mAh g⁻¹, while Ni₃(HHTP)₂/CP and Cu₃(HHTP)₂/CP only exhibit around 4000 mAh g⁻¹. The electronic redistribution at the bimetallic Ni and Cu sites creates a favorable electron cloud density for ORR and OER. Meanwhile, Pan et al.⁵⁴ grew Ni_1.5Cu_1.5(HHTP)₂ on hydroxylated graphene (G-OH) via an in situ method. This composite exhibited a remarkable capacity of 12 542 mAh g⁻¹ at 50 mA g⁻¹ and maintained stable performance over 40 cycles, highlighting the application potential of bimetallic MOFs in Li–O₂ batteries.

Typically, the coordination number (CN) of metal centers in MOFs is 4, 6, or 8. The stable and uniform distribution of metal centers enhances catalytic activity. However, the stable coordination of metal centers with an unfavorable electronic structure environment limits the progress of OER and ORR reactions.⁵⁵ In principle, MOFs with abnormal coordination numbers can be constructed by appropriately changing the coordination environment of the central metal to alter the adsorption strength of the reaction intermediate LiO₂, thereby improving the electrocatalytic activity of Li–O₂ batteries. For example, Zhou et al.⁵⁶ prepared single-crystal MOFs with different coordination environments centered on Pb. In the Na–Pb-MOF with 2,6-naphthalenedicarboxylic acid ligands, the unique PbO₇ metal node exhibits a longer Pb–O bond length compared to the PbO₆ metal node in 4OMe–Pb-MOF with 2,3,5,6-tetramethoxyterephthalic acid ligands. In the single-crystal naphthalene lead MOF (referred to as Na–Pb-MOF), the unique PbO₇ metal node has a longer bond length between Pb and O than the PbO₆ metal node in tetramethyl lead-MOF (referred to as 4OMe–Pb-MOF). This results in a weakened coupling of the Pb-5d–O-2p orbitals in the PbO₇ node. When employed as cathode materials in Li–O₂ batteries, Na–Pb-MOF exhibits superior performance than 4OMe–Pb-MOF. Na–Pb-MOF achieves a higher discharge capacity (6247 mAh g⁻¹ vs. 4765 mAh g⁻¹), lower overpotential (0.52 V vs. 1.29 V), and longer cycling stability (140 vs. 45 cycles at 0.5 A g⁻¹ and 1000 mAh g⁻¹ cutoff capacity). Electronic structure analyses reveal that the longer Pb–O bond length in PbO₇ (2.88 Å) lowers the Pb oxidation state, weakens Pb-5d–O-2p orbital coupling, and reduces LiO₂ adsorption energy (1.21 eV vs. 1.47 eV). These features contribute to lower ORR (0.43 eV) and OER (0.64 eV) activation energies compared to 4OMe–Pb-MOF. This study underscores the potential of engineering MOFs with abnormal coordination numbers to enhance reaction kinetics via tailored orbital coupling.

3.2 Structural optimization of MOFs

The advantages of MOFs as catalyst materials lie in their highly porous structures. When applied in Li–O₂ batteries, the large specific surface area allows for full contact between the electrolyte, catalyst, and O₂, effectively addressing the issue of low reactivity at the solid–liquid–gas triphase interface.⁵⁷ Good breathability and full impregnation with the electrolyte, as well as the timely decomposition of insulating Li₂O₂, ensure a stable triphase interface, facilitating the rapid kinetics of ORR/OER in high-performance lithium–oxygen batteries.⁴⁶ Compared to other porous materials, the unique aspect of MOFs is their designable pore structures and sizes given specific metal cations and ligands. Moreover, unique hollow structures can be designed through self-template transformation methods, providing additional space for Li₂O₂ deposition and facilitating gas transport and electrolyte penetration.⁵⁸ The integration of 2D MOFs further amplifies these advantages, as their planar structures enhance the exposure of active sites, improve gas diffusion rates, and optimize mass transfer kinetics. These structural and functional advancements highlight MOFs, particularly 2D MOFs, as promising materials for high-performance and scalable energy storage technologies.

Reducing the crystal size and tuning the morphology of MOFs can increase active site density and enhance electron and mass transport, thereby boosting electrochemical performance in Li–O₂ batteries. Yan et al.⁵⁹ prepared Co-MOF-74-X (X = 1400, 800, or 20, where X represents the average width of the nanorods and nanofibers) with three different sizes and morphologies by modulating the solvent composition. At a current density of 100 mA g⁻¹, the batteries with Co-MOF-74-1400 and Co-MOF-74-800 delivered discharge capacities of 4710 mAh g⁻¹ and 4880 mAh g⁻¹, respectively. However, the battery using Co-MOF-74-20 as the cathode exhibited an ultrahigh capacity of 11 350 mAh g⁻¹. Obviously, the small and uniform nanofiber structure of Co-MOF-74-20 facilitate electrolyte penetration and Li⁺ diffusion, thereby achieving better electrochemical performance.

Compared with common transition metals (Ni, Co, Mn), the trivalent ions of lanthanide elements possess 4f orbitals that are progressively filled.⁶⁰ Moreover, when the physical size is reduced to the nanoscale, due to their nanosize effect, exhibit superior heterogeneous catalytic performance.⁶¹ Liu et al.⁶² successfully synthesized Dy-BTC nanospheres via a solvothermal method. In contrast to Dy-BTC with microchannels and nanopores, Dy-BTC nanospheres, which have a narrow size range (3–15 nm) of mesoporous bulk crystals, feature abundant micropores and a wide size range (10–35 nm) of mesopores. Additionally, compared to the specific surface area of bulk Dy-BTC crystals (403.6 m² g⁻¹), the specific surface area of Dy-BTC nanospheres significantly increased to 568.1 m² g⁻¹. The abundant micropores effectively enhance gas exchange with O₂, and the wide size range of mesopores not only promotes electrolyte impregnation but also accommodates discharge products like Li₂O₂, thereby effectively improving battery capacity. This hierarchical pore structure is highly advantageous for enhancing battery capacity.⁶³ Dy-BTC nanospheres exhibited excellent electrochemical performance with significantly higher discharge capacity, rate performance, and cycling stability compared to bulk Dy-BTC crystals and Super P.

Uniformly dispersed and exposed catalytic active sites are critical for enhancing catalytic activity, and subnanoparticles offer significant advantages in this regard.⁶⁴ Unlike metal-dispersed carbon composites via high-temperature annealing, which often suffer from metal aggregation issues, MOFs provide uniformly distributed metal active sites. Choi et al.⁶⁵ designed a multishelled MOF structure (H-ZIF-8[nS]) that effectively prevented aggregation and enabled precise subnanoparticle loading. This structure effectively reduced the overpotential for OER and achieved decent performance in Li–O₂ batteries. As shown in Fig. 3, He et al.⁵⁸ initially synthesized hollow Co-MOF74 nanocages (Co-MOF74-H) through a self-template transformation of dodecahedral Co-ZIF67 particles using 2-hydroxyterephthalic acid. The hollow structure, formed via the Kirkendall effect, maintained the original template shape and provided a stable “solid–liquid–gas” triphase interface, facilitating continuous catalytic reactions in Li–O₂ batteries. Compared with solid Co-MOF-74-S and KB, Co-MOF-74-H exhibited smaller charge–discharge overpotentials (0.73 V vs. 1.24 V and 1.53 V) and larger discharge capacities (19.64 mAh cm⁻² vs. 16.32 mAh cm⁻² and 10.75 mAh cm⁻²) at 0.05 mA cm⁻², demonstrating superior catalytic performance. Zhang et al.⁶⁶ investigated the use of H-MIL-53 and MW-MIL-53, synthesized via hydrothermal and microwave-assisted methods respectively, as cathodes for Li–O₂ batteries. Both materials exhibited similar capacities of around 1000 mAh g⁻¹ in full discharge tests.

Fig. 3 (a) Schematic diagram for the preparation of H-ZIF-8[nS]. (b) Discharge and charge curves of H-ZIF-8[1S] and ZIF-8 at a current density of 50 mA g⁻¹. (c) Discharge and charge curves of H-ZIF-8[1S] and H-ZIF-8[5S] at a current density of 0.1 mA cm⁻². (d) Full discharge curves at a current density of 0.1 mA cm⁻².⁶⁵ Reproduced with permission from ref. 65. Copyright 2020, Wiley-VCH.

In recent years, two-dimensional (2D) MOF materials have gained significant attention due to their unique advantages, such as enhanced surface area, abundant active sites, and tunable interlayer spacing, which facilitate rapid ion transport and improve catalytic performance.⁶⁷ For instance, Yuan et al.⁶⁸ synthesized 2D Co-MOF, Ni-MOF, and Mn-MOF using BDC ligands and metal ions via ultrasonication. These 2D MOFs demonstrated nanosheet morphologies and higher specific surface areas compared to their 3D counterparts. Among these, 2D Mn-MOF exhibited exceptional performance in Li–O₂ batteries, achieving a high discharge capacity (9500 mAh g⁻¹) and cycling stability (>200 cycles). Its ability to catalyze the formation and decomposition of Li₂O₂ and LiOH, as confirmed by OER activity tests, effectively reduced charge–discharge overpotentials and prolonged battery life, positioning 2D Mn-MOF as a highly promising cathode catalyst for Li–O₂ batteries.

3.3 Directed design of organic ligands

Generally, the metal cations in MOFs act as the active centers for catalytic chemical reactions, while the catalytic activity can be further enhanced by adjusting the ratio of the central metal to the organic ligand. Lang et al.⁶⁹ synthesized Mn-based MOF catalysts with varying Mn-to-terephthalic acid molar ratios (1.2 : 1, 1.3 : 1, 1.4 : 1) using a solvothermal method. Structural analysis showed that the molar ratio significantly influenced the crystallinity and orientation of the MOFs, with optimal orientation observed at 1.3 : 1. Electrochemical testing of Li–O₂ batteries revealed that the 1.3 : 1 ratio provided the best cycling stability (35 cycles), outperforming the 1.2 : 1 and 1.4 : 1 ratios (13 and 15 cycles, respectively). These results demonstrate that an appropriate metal-to-ligand ratio enhances the electrode's electrochemical stability and reduces polarization resistance.

The introduction of defect engineering through ligands substitution has also proven effective in enhancing the electrochemical performance of MOF-based materials. Wang et al.⁷⁰ introduced various proportions of benzoic acid (BA) (0–40%) to replace BDC in Co-MOF synthesis, creating Co-MOF-BA defect states that enhanced electron exchange between coordination-unsaturated Co sites and oxygen. Among these, Co-MOF-BA0.1 showed the best performance, with the lowest overpotential (0.86 V) and highest discharge capacity (14 011 mAh g⁻¹ at 100 mA g⁻¹) due to its larger surface area and optimal LiO₂ adsorption. Additionally, Co-MOF-BA0.1 displayed excellent cycling stability, supporting 215 cycles, surpassing Co-MOF (60 cycles) and Co-MOF-BA0.2 (138 cycles). A volcano plot derived from DFT calculations showed that the defect states in Co-MOF-BA0.1 enhanced active site coverage while preventing blockage, resulting in low overpotentials and outstanding lithium–oxygen battery performance.

In addition to adjusting the ratio of central metals to organic ligands, the targeted design of organic ligands can further enhance the catalytic activity by increasing the affinity for Li⁺ binding, thereby promoting the uniform nucleation and growth of Li₂O₂ within the MOF framework. This ultimately leads to high capacity and stable cycling in Li–O₂ batteries. Kang et al.⁷¹ developed a bpyN-MOF/graphene (bpyN-MOF/g) cathode for Li–O₂ batteries, demonstrating the critical role of lithophilic nitrogen sites in facilitating LiO₂ adsorption and stabilizing Li₂O₂ formation, as shown in Fig. 4. DFT calculations showed nitrogen sites in bpyN-MOF exhibit stronger LiO₂ adsorption (−2.555 eV) compared to biphenyl linkers (−2.173 eV), while 3D STEM analysis revealed uniform nanoscale Li₂O₂ growth within the MOF, preventing particle aggregation. Electrochemical tests highlighted the superior performance of bpyN-MOF/g, with a charge–discharge overpotential of 0.487 V at 200 mA g⁻¹, significantly lower than bph-MOF/g (0.857 V) and Pt/C (1.028 V). Even at higher current densities (2000 mA g⁻¹), bpyN-MOF/g maintained low overpotentials (0.304 V discharge, 0.938 V charge) and exhibited excellent cycling stability (up to 270 cycles with 99.9% coulombic efficiency). The 3D porous structure of bpyN-MOF/g provided efficient pathways for Li ion and O₂ transport while enabling a surface-mediated, film-like growth of Li₂O₂. This study underscores the potential of optimized MOF cathodes for advancing high-performance Li–O₂ batteries through the targeted design of organic ligands.

Fig. 4 The synthesis of bpyN-MOF materials and their application in Li–O₂ batteries.⁷¹ Reproduced with permission from ref. 71. Copyright 2024, Wiley-VCH.

Poor conductivity is a common limitation of MOFs, which hinders efficient electron transfer, reducing their catalytic performance and overall battery efficiency. Enhancing the conductivity of MOFs enables better charge transport, improved ORR and OER kinetics, and more stable cycling performance, making conductive MOFs highly attractive for Li–O₂ batteries. Majidi et al.⁷² developed a conductive metal–organic framework Cu-THQ, which, in synergy with the redox mediator (RM) 0.1 M InBr₃ and 1 M LiNO₃, enabled the battery to stably cycle for 150 cycles under test conditions of a current density of 1 A g⁻¹ and a cutoff capacity of 2000 mAh g⁻¹. Additionally, when the current density was increased to 2 A g⁻¹, the battery still maintained stable cycling for 100 cycles, with a charging potential of less than 3.7 V. The results demonstrated the reversible discharge performance of the battery at different rates, with a high specific capacity of up to 3500 mAh g⁻¹.

3.4 Strategies for MOF hybrids

To overcome the limitations of MOFs in Li–O₂ batteries, such as poor conductivity and insufficient catalytic activity, integrating conductive polymers, metal nanoparticles, and other functional nanostructures has emerged as a promising strategy. By leveraging the intrinsic advantages of MOFs and the complementary properties of these materials, such hybrids significantly enhance electron and mass transfer while improving the stability and catalytic efficiency of Li–O₂ batteries. Zhang et al.⁷³ developed a 2D Co-MOF/Ti₃C₂ hybrid membrane using an electrostatic assembly method, as shown in Fig. 5(a), which prevented the self-stacking of Ti₃C₂ layers and enhanced ion transport, conductivity, and O₂ diffusion. By employing vacuum filtration, they prepared binder-free, self-standing membranes with improved cathode catalyst dispersion and a porous structure. The hybrid membrane leveraged the open metal active sites of Co-MOF and the increased interlayer spacing of Ti₃C₂, significantly enhancing ORR/OER catalytic activity. The 10 wt% Co-MOF/Ti₃C₂ hybrid membrane exhibited exceptional performance, achieving a discharge capacity of 34 763 mAh g⁻¹ and 278 cycles, at a current of 1000 mA g⁻¹. Far surpassing the 5 wt% hybrid (16 278 mAh g⁻¹, 208 cycles) and pure Ti₃C₂ (8147 mAh g⁻¹, 84 cycles), along with superior rate performance (Fig. 5b and c). Zheng et al.⁷⁴ also confirmed the synergistic effect of the Co-MOF/Ti₃C₂ hybrid structure in Li–O₂ batteries, demonstrating significantly enhanced discharge capacity and cycling stability compared to single-component counterparts.

Fig. 5 (a) Schematic of the synthesis of the 2D Co-MOF/Ti₃C₂ hybrid. (b) Cycling performance of a limiting specific capacity of 1000 mAh g⁻¹ of Li–O₂ batteries. (c) Full discharge/charge profiles of Li–O₂ batteries at a current of 1000 mA g⁻¹.⁷³ Reproduced with permission from ref. 73. Copyright 2022, Elsevier.

In addition, Zhang et al.⁷⁵ developed Mn-MOF-74@CNTs hybrids via an additive-mediated solvothermal method, with Mn-MOF-74 nanocrystals grown on CNTs. These hybrids exhibit excellent catalytic activity, enabling the reversible formation and decomposition of LiOH in humid oxygen environments. This advancement addresses a major challenge in the practical application of Li–O₂ batteries, which is hindered by the stringent testing conditions of dry oxygen. The hybrid catalyst demonstrated superior catalytic activity for H₂O₂ decomposition, enabling the formation of LiOH nanosheets as the discharge product in oxygen containing 200 ppm water, verified by PXRD and SEM analyses. In a humid environment, the Mn-MOF-74@CNTs cathode achieved a significantly extended cycle life of 60 cycles at a current density of 125 mA g⁻¹, compared to just 14 cycles in dry oxygen and less than 20 cycles for traditional carbon-based cathodes. This enhanced performance is attributed to reduced side reactions and the catalytic role of redox-active Mn²⁺/Mn³⁺ centers, offering a promising approach for Li–O₂ batteries to operate under more realistic, humid conditions.

Notably, the pore space of MOFs can be utilized to encapsulate other catalytically active substances. The encapsulated active substances are in full contact with the MOF materials, ensuring rapid electron and mass transfer, which in turn ensures rapid ORR/OER kinetics. Moreover, the active substances confined in the pore space can avoid aggregation, thereby ensuring as many active catalytic sites as possible. He et al.⁵⁸ utilized the hollow porous structure of Co-MOF74-H nanocages to encapsulate pyrrole monomers, which, through thermal initiation and iodination, formed iPPM nanostructures capable of releasing I⁻/I₃⁻ redox mediators. This iPPM structure serves as a cathode catalyst, anode protective layer, and redox mediator carrier, stabilizing the anode/electrolyte interface and promoting high Li⁺/Li redox kinetics. In full battery tests, the iPPM-based system exhibited an initial overpotential of 0.2 V at 0.05 mA cm⁻² and achieved 311 cycles at a cutoff capacity of 0.1 mAh cm⁻², compared to 106 cycles with Co-MOF74-H using 5% LiI in the electrolyte. This “trinity” design significantly reduces polarization, enhances discharge capacity, and improves cycle stability in Li–O₂ batteries.

4. Photoactive MOF components for photo-assisted Li–O₂ batteries

The incorporation of light energy has brought a paradigm shift to the development of Li–O₂ batteries, and the advent of photo-assisted technologies opens new possibilities for improving their efficiency and performance.⁷⁶ Extensive research has focused on exploring various catalysts, including noble metals, oxides, sulfides, polymers, g-C₃N₄.^77,78 Since the discovery of photocatalytic activity in MOF-5, photoactive MOF-based composite materials have undergone rapid development in applications such as CO₂ reduction catalytic reactions, photocatalytic water splitting for hydrogen production, and organic pollutant degradation.⁷⁹ Rational design of MOF architectures through functional group grafting, metal node engineering, ion doping and ligand modification provides an effective strategy for tailoring visible-light absorption properties and boosting photon utilization efficiency.^80,81 Additionally, by combining other semiconductor materials with MOFs to form semiconductor heterojunctions with interfacial barriers, the recombination of photogenerated electron–hole pairs can be inhibited, thereby achieving efficient charge separation and high redox capabilities.⁴³ Moreover, the high porosity and specific surface area, tunable metal active sites, and dispersed catalytic active centers of MOFs endow them with excellent optical, electrocatalytic properties, and surface reaction rates. Therefore, MOF materials demonstrate tremendous potential for application in photo-assisted lithium–air batteries.

Pioneering the use of metal–organic materials in photo-assisted Li–air batteries, Lv et al.⁸² synthesized a novel Co-TABQ nanosheet on carbon paper, achieving a band structure with a CB value of 2.03 V (vs. Li/Li⁺) and a VB value of 4.23 V (vs. Li/Li⁺). These values align with the redox potential of Li₂O₂/O₂, meeting the thermodynamic requirements for photocatalytic ORR and OER. Under illumination, Li–O₂ batteries with Co-TABQ cathode achieved a discharge voltage of 3.12 V and a charge voltage of 3.32 V at a high current density of 0.1 mA cm⁻², with a round-trip efficiency as high as 94%. Moreover, their photo-assisted Li–O₂ batteries stably cycled for 50 cycles with an overpotential of less than 0.5 V, in sharp contrast to the dark condition. DFT calculations revealed that during discharge, Co atoms adsorb O₂, facilitating its activation to LiO₂ intermediates and subsequent reduction to Li₂O₂. During charging, VB holes oxidized Li₂O₂ into Li⁺ and O₂, while CB photoelectrons reduced Li⁺ at the anode driven by the external circuit. The discharge voltage drop upon SCN⁻ poisoning experiments proved Co atoms are the ORR-active centers in Co-TABQ, as predicted by DFT. The design of this novel metal–organic polymer has propelled the application of MOFs in photo-assisted Li–O₂ batteries for energy storage.

The visible-light absorption of MOFs can be systematically red-shifted through strategic functionalization of organic ligands with electron-donating/withdrawing groups (–NH₂, –CH₃, –NO₂, –SO₃H, –SH).⁸³ Specifically, amino-functionalized MOFs demonstrate enhanced visible-light response due to the introduction of low-energy n → π* electronic transitions (compared to conventional π → π* transitions), effectively reducing the bandgap by 0.3–0.5 eV.⁸⁴ Tao et al.⁸⁵ found that MIL(Fe)-101-NH₂ extended light absorption from 450 to 750 nm due to dual excitations of –NH₂ and Fe–O clusters. Photo-assisted Li–O₂ batteries with MIL(Fe)-101-NH₂ achieved a discharge voltage of 2.93 V and a charge voltage of 3.15 V at a current density of 0.05 mA cm⁻², which is significantly lower than the charge/discharge overpotential of 0.82 V without illumination. The MIL-101(Fe)-NH₂ photocathode demonstrated exceptional electrochemical performance under illumination, maintaining 85% capacity retention at current densities up to 0.5 mA cm⁻² and achieving remarkable cycling stability of 195 cycles. In addition, Yu et al.⁸⁰ assembled photo-assisted Li–O₂ batteries by in situ growth of UIO-66-NH₂ on carbon cloth, which increased the round-trip efficiency of the batteries from 63.2% to 81.3% under light, and was able to cycle stably at high current densities.

For MOFs containing photoactive ligands (e.g., porphyrins, viologens), the photoresponsive properties can be further enhanced through rational engineering of metal nodes. Wen et al.⁸⁶ synthesized FeNi-TCPP porphyrin MOF by a one-step solvothermal reaction, where Fe³⁺ and Ni²⁺ coordinate with the carboxyl groups on the TCPP linker. Electron paramagnetic resonance (EPR) spectroscopy revealed that FeNi-TCPP generates stronger O₂⁻ but weaker ¹O₂ compared to Fe-TCPP under illumination, indicating enhanced O₂ activation through a charge transfer pathway. Theoretical calculations showed that Ni increases the density of states near the Fermi level and facilitates Ni 3d and Fe 3d orbital overlap, improving charge transfer and conductivity. In photo-assisted Li–O₂ batteries, FeNi-TCPP achieved higher efficiency (92%) and lower overvoltage compared to Fe-TCPP, maintaining stable performance over 45 cycles at a high current density. Overall, the introduction of Ni improves electron transfer, O₂ activation, exciton dissociation, and charge carrier lifetimes in FeNi-TCPP.

MOF materials can effectively enhance visible-light absorption properties through the modification of metal nodes or the incorporation of dopants, providing a promising strategy for precisely tuning light absorption characteristics. We noticed that Ce-UiO-66 extends its absorption range of UV-visible light to the visible light region due to the low-energy 4f orbitals of the central metal Ce.⁸⁷ Using a straightforward solvothermal approach, we successfully prepared a Ce-UiO-66-integrated photocathode.⁸⁸ UV-Visible spectroscopy indicated an absorption edge at 458 nm, showing strong absorption of visible light. Under illumination at 0.018 mA cm⁻², the photo-assisted Li–O₂ battery with Ce-UiO-66 showed improved performance, with the discharge voltage rising from 2.7 V to 3.1 V, and charge voltage dropping from 4.5 V to 3.6 V. At 0.09 mA cm⁻², the Ce-UiO-66 photocathode achieved 160 cycles, compared to 40 cycles in the dark. The Ce^3+/4+ redox centers enhance the photoelectrocatalytic process in the photo-assisted Li–O₂ battery, as confirmed by DFT calculations, while the integrated photocathode structure could further facilitate the charge transfer efficiency.

To enhance the visible-light harvesting capability and catalytic activity of MOF materials, we further developed a bimetallic Fe-UiO-66 MOF cathode through a facile solvothermal method by incorporating Fe³⁺ into the UiO-66 framework (Fig. 6a).⁸¹ When employed as a photocathode in photo-assisted Li–O₂ batteries, Fe-UiO-66 achieved a discharge voltage of 3.42 V and a charge voltage of 3.56 V, resulting in a high round-trip efficiency of 96%. The Fe-UiO-66 constructed photo-assisted Li–O₂ battery also exhibited superior discharge capacity and stability, cycling for more than 1000 hours without degradation (Fig. 6b and c). The exceptional performance is attributed to Fe³⁺ nodes, which regulate photogenerated electrons, expand visible light absorption (200–680 nm as shown by UV-vis spectra), and enhance ORR/OER catalytic activity. EPR analysis revealed electron transfer under illumination, where Fe³⁺ donates electrons to Zr-oxo clusters, facilitating O₂ activation into O₂⁻. DFT calculations demonstrated that the presence of Fe as an active site in Fe-UiO-66 significantly enhanced the adsorption of O₂ molecules, which subsequently reacted with Li⁺ ions to form LiO₂ and eventually transformed into a film-like Li₂O₂. Illumination also promotes electron transfer and hole-driven oxidation of Li₂O₂ during charge, reducing the potential required from the external circuit. This work demonstrates that engineering MOF metal nodes to enhance visible light absorption and catalytic activity under illumination is a promising strategy for efficient photo-assisted Li–O₂ batteries.

Fig. 6 (a) Schematic of the preparation process for Fe-UiO-66. (b) Discharge and charge profiles of the Li–O₂ batteries (c) cycling performance of Li–O₂ batteries.⁸¹ Reproduced with permission from ref. 81. Copyright 2024, Wiley-VCH.

MOF materials can be integrated with semiconductor nanoparticles or quantum dots within their porous structure to form light-responsive composites.⁸⁹ This design enables efficient separation of photogenerated electron–hole pairs and significantly extends the carrier lifetime, enhancing their photoelectric properties. Qiao et al.²¹ developed CsPbBr₃@PCN-333(Fe) composites by growing CsPbBr₃ perovskite quantum dots (QDs) inside the mesoporous cages of PCN-333(Fe) using a simple sequential deposition method. The hierarchical porous structure of PCN-333(Fe), featuring mesopores (∼4.2 and 5.5 nm) for encapsulating perovskite QDs and micropores (∼1.1 nm) for mass transfer, ensures the photoelectrocatalysis performance. Photogenerated electrons efficiently transfer from CsPbBr3 to PCN-333(Fe), while holes migrate in the opposite direction under illumination, enabling effective electron–hole pair separation, as confirmed by XPS and computational results. As a photoelectrode in Li–O₂ batteries, CsPbBr3@PCN-333(Fe) achieved a discharge voltage of 3.19 V, a charging potential of 3.44 V, and a round-trip efficiency of 92.7% at 0.01 mA cm⁻², outperforming individual CsPbBr₃ (82.8%) and PCN-333(Fe) (85.2%). Its cycling stability reached 200 hours, far exceeding that of pure components. This enhanced performance is attributed to efficient electron–hole pair separation, extended carrier lifetime, and structural stability provided by PCN-333(Fe). This work highlights a promising strategy for designing photoresponsive MOF composites as cathode materials in photo-assisted Li–O₂ batteries.

Tao et al.²⁰ synthesized a Zr-MOF/BiOIBr/CC heterojunction composite photoelectrode on carbon paper by in situ growth using ZrCl₄ and NH₂-BDC. The heterojunction photoelectrode broadened the light absorption range, while the built-in electric field facilitated the separation and directional movement of photogenerated carriers, thereby accelerating the reaction kinetics of photo-assisted Li–O₂ batteries. At a current density of 0.05 mA cm⁻², the Zr-MOF/BiOIBr/CC composite photoelectrode exhibited a charge/discharge overpotential of 0.15 V under illumination, lower than the 0.41 V of the Zr-MOF/CC photoelectrode. When cycled at 0.05 mA cm⁻² with a cutoff capacity of 0.05 mAh cm⁻², the Zr-MOF/BiOIBr/CC based photo-assisted Li–O₂ batteries achieved 255 cycles, surpassing the 170 cycles achieved by the Zr-MOF/CC photoelectrode. This highlights the potential application value of MOFs/semiconductor heterostructures in photo-assisted Li–O₂ batteries.

The comparative performance of various MOFs in Li–O₂ batteries is summarized in Table 1, including key metrics such as detailed structure of the MOF materials, discharge capacity, charge/discharge plateaus, current density, cycling stability, etc. Notably, the performance metrics highlight that photoactive MOF components further amplify the potential of MOFs in Li–O₂ systems by enabling photo-assisted charge/discharge processes, which mitigate overpotentials and enhance energy efficiency.

Table 1 Examples of MOFs for various types of Li–O₂ batteries

Sample ID	Metal	Organic linker	Specific surface area/m^2 g^−1	Synthesis methods	Capacity mAh g^−1	Cycle performance	Round-trip efficiency (%) 1st	Ref.
a H2BDC = 1,4-benzenedicarboxylic acid.b H3BTC = 1,3,5-tricarboxylic acid.c DHTA = 2,5-dihydroxyterephthalic acid.d L3 = 2,6-naphthalenedicarboxylic.e L5 = 2,3,5,6-tetramethoxyterephthalic acid.f Solvent mixtures DMF : H2O = 4 : 1.^75g 2-mim = 2-methyleimidazole.h H4DOBDC = 2-hydroxyterephthalic acid.i bpyN = bipyridine.j THQ = tetrahydroxy-1,4-benzoquinone hydrate.k TABQ = tetramino-benzoquinon.l NH2-BDC = 2-aminoterephthalic acid.m TCPP = 5,10,15,20-tetrakis(4-carboxyphenyl) porphyrin.n H3TATB = 4,4′,4′′-s-triazine-2,4,6-triyl-tribenzoic acid.
MOF-5	Zn	H2BDC^a	3622	Hydrothermal	1780	N/A	N/A	90
					50 mA g^−1
HKUST-1	Cu	H3BTC^b	1944	Hydrothermal	4170	N/A	N/A	90
					50 mA g^−1
Co-MOF-74	Co	DHTA^c	1446	Hydrothermal	3630	N/A	N/A	90
					50 mA g^−1
Mg-MOF-74	Mg	DHTA	1582	Hydrothermal	4560	N/A	N/A	90
					50 mA g^−1
Mn-MOF-74	Mn	DHTA	1213	Hydrothermal	9420	30 cycles 240 h	N/A	90
					50 mA g^−1	250 mA g^−1
MnCo-MOF-74	Mn, Co	DHTA	N/A	Hydrothermal	11 150	44 cycles 440 h	∼68.9%	51
					200 mA g^−1	200 mA g^−1	200 mA g^−1
NiMn0.05-MOF	Mn, Ni	H2BDC	89.5	Microwave reaction	28 464	540 cycles	∼89.6%	52
					100 mA g^−1	1000 mAh g^−1	100 mA g^−1
Ni1.5Cu1.5(HHTP)2/CP	Ni, Cu	HHTP	N/A	Hydrothermal	8367	43 cycles 860 h	∼50.9%	53
					50 mA g^−1	50 mA g^−1	50 mA g^−1
Ni1.5Cu1.5(HHTP)2/G-OH	Ni, Cu	HHTP	248.3	Hydrothermal	12 542	40 cycles 800 h	∼70.3%	54
					50 mA g^−1	50 mA g^−1	50 mA g^−1
4OMe–Pb-MOF	Pb	L3^d	N/A	Hydrothermal	6247	45 cycles 180 h	∼74.5%	56
					—	500 mA g^−1	100 mA g^−1
Na–Pb-MOF	Pb	L5^e	N/A	Hydrothermal	4765	140 cycles 560 h	∼84.5%	56
					—	500 mA g^−1	100 mA g^−1
Co-MOF-74-20^f	Co	DHTA	874	Hydrothermal	11 350	8 cycles 64 h	∼75.6%	59
					100 mA g^−1	250 mA g^−1	100 mA g^−1
Dy-BTC	Dy	H3BTC	568.1	Self-assembly	7618	60 cycles	∼61.3%	62
					50 mA g^−1	1000 mAh g^−1	50 mA g^−1
H-ZIF-8[5S]	Co, Zn	2-mim^g	N/A	Self-assembly	∼2.95 mAh cm^−2	∼20 cycles 200 h	∼75.6%	65
					0.1 mA cm^−2	0.1 mA cm^−2	0.1 mA cm^−2
Co-MOF-74-H	Co	H4DOBDC^h	722.2	Self-assembly	19.64 mAh cm^−2	112 cycles 448 h	∼76%	58
					0.05 mA cm^−2	0.05 mA cm^−2	0.05 mA cm^−2
MW-MIL-53	Al	H2BDC	1391	Microwave reaction	∼1200	∼7 cycles	∼57%	66
					—	—	—
2D Mn-MOF	Mn	H2BDC	140	Microwave reaction	9500 mAh g^−1	200 cycles 4000 h	68.5%	68
					100 mA g^−1	100 mA g^−1	100 mA g^−1
2D Co-MOF	Co	H2BDC	121	Microwave reaction	6960 mAh g^−1	25 cycles 500 h	66.2%	68
					100 mA g^−1	100 mA g^−1	100 mA g^−1
2D Ni-MOF	Ni	H2BDC	113	Microwave reaction	5367 mAh g^−1	17 cycles 340 h	66.1%	68
					100 mA g^−1	100 mA g^−1	100 mA g^−1
Mn-MOF	Mn	H2BDC	—	Hydrothermal	—	35 cycles	∼62%	69
						500 mAh g^−1	0.1 mA cm^−2
Co-MOF-BA0.2	Co	H2BDC	97.6	Microwave reaction	14 021 mAh g^−1	215 cycles	∼75.1%	91
					100 mA g^−1	1000 mAh g^−1	100 mA g^−1
bpyN-MOF/g	Zr	bpyN^i	—	Hydrothermal	17 275 mAh g^−1	270 cycles	∼85.2%	71
					100 mA g^−1	1000 mAh g^−1	200 mA g^−1
Cu-THQ	Cu	THQ^j	—	Self-assembly	—	150 cycles 600 h	∼70.8%	72
						1000 mA g^−1	1000 mA g^−1
10 wt% Co-MOF/Ti3C2	Co	H2BDC	274	Hydrothermal	34 763 mAh g^−1	278 cycles 1112 h	∼74.1%	73
					1000 mA g^−1	1000 mA g^−1	1000 mA g^−1
Mn-MOF-74@CNTs	Mn	H2BDC	462	Hydrothermal	∼3200 mAh g^−1	60 cycles 480 h	∼64.9%	75
					50 mA g^−1	125 mA g^−1	125 mA g^−1
					(200 ppm)	(200 ppm)	(200 ppm)
iPPM	Co	H4DOBDC	3.3	Hydrothermal	28.41 mAh cm^−2	311 cycles 1244 h	93%	58
					0.05 mA cm^−2	0.05 mA cm^−2	0.05 mA cm^−2
Co-TABQ	Co	TABQ^k	137	Hydrothermal	—	50 cycles 500 h	94%	82
						0.1 mA cm^−2	0.1 mA cm^−2
MIL(Fe)-101-NH2	Fe	NH2-BDC^l	1136	Hydrothermal	9.22 mAh cm^−2	195 cycles 390 h	93%	85
					0.1 mA cm^−2	0.05 mA cm^−2	0.05 mA cm^−2
MIL(Ti)-101-NH2	Ti	NH2-BDC	774	Hydrothermal	8.15 mAh cm^−2	136 cycles 272 h	90%	85
					0.1 mA cm^−2	0.05 mA cm^−2	0.05 mA cm^−2
MIL(Zr)-101-NH2	Zr	NH2-BDC	929	Hydrothermal	7.14 mAh cm^−2	154 cycles 308 h	89%	85
					0.1 mA cm^−2	0.05 mA cm^−2	0.05 mA cm^−2
UIO-66-NH2	Zr	NH2-BDC	608.7	Hydrothermal	4.8 mAh cm^−2	5 cycles 40 h	81.3%	80
					0.5 mA cm^−2	0.1 mA cm^−2	0.025 mA cm^−2
FeNi-TCPP	Fe/Ni	TCPP^m	—	Hydrothermal	—	45 cycles	92%	86
						0.5 mA cm^−2	0.1 mA cm^−2
Ce-UiO-66	Ce	p-Phthalic acid	—	Hydrothermal	—	160 cycles ∼320 h	82.2%	88
						0.09 mA cm^−2	0.09 mA cm^−2
Fe-UiO-66	Fe	H2BDC	—	Hydrothermal	2.5 mAh cm^−2	500 cycles 1000 h	96%	81
					0.02 mA cm^−2	0.01 mA cm^−2	0.01 mA cm^−2
CsPbBr3@PCN-333(Fe)	Fe	H3TATB^n	—	Hydrothermal	—	100 cycles 200 h	92.7%	21
						0.01 mA cm^−2	0.01 mA cm^−2
Zr-MOF/BiOIBr/CC	Zr	NH2-BDC	—	Hydrothermal	—	250 cycles 500 h	95.3%	20
						0.05 mA cm^−2	0.05 mA cm^−2

---

5. MOFs in solid-state electrolytes

In Li–O₂ batteries, the electrolyte is vital for Li⁺ transport, O₂ dissolution and transport, and the formation and decomposition of Li₂O₂.⁹² MOFs, as novel porous materials, offer significant advantages over traditional liquid electrolytes. Their non-flammability, low volatility, and high mechanical strength enhance battery safety and structural stability, minimizing risks from high temperatures, overcharging, or short circuits. Unlike liquid electrolytes that require separators to maintain structure, MOF-based electrolytes simplify battery design and reduce costs. MOFs can also effectively prevent lithium dendrite growth,⁹³ and their chemical stability allow operation within wide electrochemical windows, improving cycling stability and safety. In addition, MOFs are stable against air components (H₂O, CO₂, O₂) and oxygen reduction intermediates, reducing side reactions and corrosion.⁹⁴ Particularly, their tunable pore structure provides efficient channels for Li⁺ transport, while unsaturated metal sites immobilize anions, enhancing lithium ion transference. Surface modification and doping further boost ionic conductivity, making MOF-based solid electrolytes superior for Li–O₂ batteries.

Liu et al.⁹⁶ ingeniously crafted a LiIL-MOF solid-state electrolyte (SSE) for Li–O₂ batteries by encapsulating the ionic liquid [Py₁₃][TFSI] blended with lithium salt LiTFSI within the pores of ZIF-67. This advanced electrolyte boasts a high ionic conductivity (6.74 × 10⁻⁴ S cm⁻¹ at 30 °C), excellent thermal stability (>330 °C), and oxidation stability (>5.0 V). When tested at a current density of 100 mA g⁻¹, the LiIL-MOF-based battery achieved a discharge voltage of 2.76 V and a charge voltage of 3.26 V, significantly lower than the 2.69 V discharge and 3.70 V charge voltages of batteries using conventional liquid electrolytes. It maintained a lifespan exceeding 100 cycles, outperforming liquid electrolyte batteries that failed after 55 cycles, and operated reliably even at elevated temperatures (80 °C). Cycling tests revealed stable discharge/charge voltages, along with a protected lithium anode due to a durable MOF particle layer. With superior electrochemical performance, long cyclability, and structural stability, the LiIL-MOF electrolyte represents a major advancement toward practical commercial Li–O₂ battery applications.

Wang et al.⁹⁷ developed a solid-state Li–O₂ battery using lithium-doped UiO-67 (UiO-67-Li) as a SSE and a 3D hierarchical porous reduced rGO aerogel as a solid-state cathode. UiO-67-Li exhibited superior ionic conductivity (0.64 mS cm⁻¹ at room temperature) compared to UiO-66-Li (0.069 mS cm⁻¹) and MCM-48-Li (0.011 mS cm⁻¹), owing to its porous structure and open metal sites. Furthermore, UiO-67-Li demonstrated excellent chemical and electrochemical stability, with an extended stability window (4.41 V vs. Li⁺/Li) and resilience against H₂O, O₂⁻, and air exposure. By integrating UiO-67-Li SSE nanoparticles onto the UiO-67-Li@rGO aerogel, the solid-state Li–O₂ batteries demonstrated superior discharge capacity and prolonged cycling stability. This study highlights the importance of continuous triple-phase boundaries and efficient interfaces in enhancing the performance and longevity of solid-state Li–O₂ batteries, advancing MOF-based energy storage systems.

In addition, MOF materials can be further optimized for ionic conductivity through doping, surface modification with functional groups, and other modification strategies. For instance, Miao et al.⁹⁵ introduced positively charged –N⁺(CH₃)₃ functional groups into the framework of UiO-66 to enhance its ionic conductivity. The modified UiO-66, termed CMOF, exhibited an ionic conductivity of 6.45 × 10⁻⁴ S cm⁻¹ at room temperature (25 °C), significantly higher than the original UiO-66's conductivity of 4.33 × 10⁻⁵ S cm⁻¹ and surpassing most reported MOF-based SSEs.^101–104 The lone pair electrons in the amino group interact with Li⁺, regulating the uniform distribution of Li⁺ and accelerating selective Li⁺ transport.⁹⁸ Further investigation showed that excessive functional groups of –N⁺(CH₃)₃ hindered ionic transport, while CMOF achieved a high lithium transference number (0.59) and enhanced conductivity by facilitating PF₆⁻ adsorption and Li⁺ dissociation. Additionally, CMOF displayed excellent stability against H₂O, Li₂O₂, and O₂⁻, and was used to construct a CMOF@rGO solid cathode for Li–O₂ battery performance evaluation. As depicted in the Fig. 7, the Li–O₂ battery with CMOF SSE demonstrated a significantly reduced initial full discharge/charge overpotential, superior rate capability, and excellent cycle stability. In conclusion, the directional modification of ligand functional enabled the design of CMOF solid-state electrolytes with high ionic conductivity and stability, providing a new strategy for MOF-based solid-state electrolytes in Li–O₂ batteries. Integrating MOFs with polymers provides another effective approach to optimize solid-state electrolytes for Li–O₂ batteries. Wang et al.⁹⁹ integrated defect-rich NH₂-UiO-66 with a polymer-based solid electrolyte PVDF. The ordered porous structure, Lewis acidic sites, and functional groups of NH₂-UiO-66 enabled strong anchoring of residual DMF in the PVDF-based electrolyte. This resulted in a MOF@PVDF electrolyte with high Li⁺ conductivity (3.6 × 10⁻⁴ S cm⁻¹), low activation energy (0.20 eV), a broad electrochemical window (4.6 V), and excellent electrochemical stability. The redox reactivity of DMF was significantly suppressed, promoting the formation of an inorganic-rich SEI layer for uniform Li deposition and effectively inhibiting Li dendrite growth and electrolyte volatility. Consequently, the MOF@PVDF-based Li–O₂ battery exhibited a long lifespan of over 1100 hours at 200 mA g⁻¹. This work offers a novel strategy for regulating solvation structures in polymer electrolytes and paves the way for designing advanced high-voltage polymer-based Li–O₂ batteries.

Fig. 7 The applications of CMOF-based SSE for Li–O₂ batteries.⁹⁵ Reproduced with permission from ref. 95. Copyright 2023, Wiley-VCH.

6. Other supplementary content

6.1 The preparation methods for MOFs

Among various MOF synthesis strategies including mechanochemistry, electrochemistry, and self-assembly methods, hydrothermal synthesis maintains its position as the most extensively employed technique to date.^63,100 This method typically involves reacting metal salts with organic linkers at a preset temperature for a certain period to achieve crystallization. This process is particularly suitable for the preparation of milligram quantities of MOF powders and has thus become the preferred synthetic strategy in most research articles.¹⁰⁰ However, several drawbacks of this method make it challenging to meet the demands of low-cost industrial production of MOFs. The crystallization process is usually slow and requires heating, which significantly increases the energy consumption and time cost of the production process. When scaling up production, the need to handle large quantities of chemicals in a single batch poses safety hazards. Moreover, this method is not conducive to continuous production.

Microwave heating, which utilizes the interaction between electromagnetic fields and molecular dipole moments to achieve more efficient energy transfer, offers a faster reaction rate compared to conventional heating and has thus become an attractive option for synthesizing MOFs.¹⁰¹ One of the early attempts to use microwaves for MOF synthesis was reported by Chang et al.¹⁰² It is worth noting that the rapid crystallization induced by microwave heating often results in smaller crystal sizes. However, microwave synthesis has not yet been widely adopted for large-scale MOF production, primarily due to the limited penetration depth of microwaves, which can lead to uneven heating when scaling up the reactor size.

Flow chemistry offers an industrial alternative to hydrothermal synthesis.¹⁰³ Compared to batch production, the main advantage of flow chemistry is its ability to enable continuous production, which is highly beneficial for large-scale industrial manufacturing. As shown in Fig. 8, driven by pumps, reactants are continuously fed into the device from the left, flow through the internal tubing for reaction, and the products are continuously discharged from the right. Additionally, due to more effective mixing of reactants and efficient energy transfer, flow chemistry can significantly reduce synthesis time and energy consumption. Compared to batch reactions, continuous flow reactions have distinct advantages, such as higher safety when handling large quantities of chemicals and the elimination of downtime between batches. All these factors make the process more cost-effective and easier to operate. In 2014, Rubio-Martinez et al.¹⁰⁴ conducted one of the early studies on using flow chemistry to synthesize MOFs such as HKUST-1, UiO-66, and NOTT-400. The space–time-yield (STY) was calculated to be 592 to 4533 kg m⁻³ d⁻¹, with a production rate of about 60 g h⁻¹. Furthermore, Taddei et al.¹⁰⁵ investigated the effect of using microwave radiation instead of conventional heating in the flow chemistry synthesis of MOFs (such as UiO-66, MIL-53(Al), and HKUST-1). Microwave radiation can more effectively penetrate the narrow tubing used in flow chemistry devices. The STY obtained in this study reached as high as 7204 kg m⁻³ d⁻¹, 3618 kg m⁻³ d⁻¹, and 66 650 kg m⁻³ d⁻¹ for UiO-66, MIL-53(Al), and HKUST-1, respectively. Flow chemistry, combined with more efficient heating methods and continuous product output, further reduces costs and provides a feasible pathway for large-scale MOFs production. Although pipeline blockage is a typical challenge when using flow devices, this method is still considered an important direction for future industrial production of MOFs.

Fig. 8 Schematic illustration describing the synthesis of MOFs using continuous flow production.¹⁰⁴ Reproduced with permission from ref. 104. Copyright 2014, Springer Nature.

6.2 MOFs-derived materials

Benefiting from the retained partial framework structure of MOFs and enhanced electrical conductivity, pyrolyzed MOF-derived materials generally demonstrate superior catalytic activity compared to pristine MOFs when utilized as cathodes in Li–O₂ batteries. While our group has achieved advancements in the application of MOF-derived materials in metal air battery systems, leveraging the unique advantages of MOF precursors,^106–108 we still contend that the fundamental challenges of composition control and morphology preservation during conversion processes remain largely unexplored. Despite these challenges, MOF-derived materials continue to offer exceptional potential as cathode catalysts for Li–O₂ batteries, warranting further systematic investigation.

The calcination of MOFs in inert gases at high temperatures to produce MOF-derived carbon not only effectively enhances electrical conductivity but also retains the original porous structure and dispersed metal active sites.⁴ For example, Lyu et al.¹⁰⁹ used 3D printing technology to fabricate self-supported hierarchical frameworks embedded with Co-MOF (ZIF-67), and then annealed them at 800 °C in a N₂ atmosphere to obtain the 3DP-NC-Co cathode. The self-supported carbon framework not only has good electrical conductivity but also ensures the necessary mechanical stability and porosity. The dispersed Co nanoparticles are embedded in the N-doped mesoporous carbon sheets, exhibiting good catalytic activity. It has a high discharge capacity of 1124 mAh g⁻¹ and can stably cycle for 160 hours. Compared with metal nanoparticle catalysts, single-atom catalysts (SAS) have the advantages of ultrahigh active atom utilization efficiency, unsaturated atomic coordination sites, and uniformly dispersed active centers, which significantly outperform traditional metal nanoparticle catalysts. Wang et al.¹¹⁰ successfully prepared the Co-SAS/N-C catalyst by pyrolyzing Zn-hexamine MOFs in an argon atmosphere and capturing the sublimated CoCl₂·6H₂O. The atomically dispersed Co atoms bonded with N atoms can serve as active catalytic centers, which greatly enhance the absorption capacity of LiO₂ and fundamentally regulate the nucleation and growth pathways and spatial distribution of Li₂O₂. This uniformly distributed Li₂O₂ can fully activate the catalytic activity of the Co–N₄ group in the oxygen evolution reaction (OER), endowing the Co-SAs/N-C electrode with outstanding electrochemical performance: an ultralow charge–discharge overvoltage (0.40 V), an ultrahigh rate discharge capacity (11 098 mAh g⁻¹ at 1 A g⁻¹), and excellent long-cycle stability (stable cycling for 260 cycles at 400 mA g⁻¹).

In addition to carbonization, MOFs can also be calcined in oxidative gas atmospheres to produce metal oxides, which also possess potential catalytic activity.¹¹¹ Zhao et al.¹¹² annealed a Mn-based MOF with the ligand 1,2,3-triazolate at 350 °C in an air atmosphere, transforming it into Mn₂O₂ with a porous nanocage structure. The material fully retained the three-dimensional framework characteristics of the precursor, and its hierarchical porous structure generated an extremely high specific surface area, which effectively promoted the transport of Li⁺ ions and the diffusion kinetics of O₂. It is particularly worth noting that the abundant oxygen defect sites on the material surface significantly enhanced the electrocatalytic activity. This synergistic effect of structural advantages and defect engineering enabled it to exhibit outstanding performance in Li–O₂ batteries, which achieved an ultra-high specific capacity of 21 070.6 mAh g⁻¹ at a current density of 500 mA g⁻¹, while maintaining excellent cycling stability.

7. Conclusions and perspective

MOFs have emerged as a revolutionary class of materials for advancing the performance of Li–O₂ batteries, offering unparalleled advantages in terms of structural tunability, catalytic activity, and multifunctionality. This review has systematically explored their applications in Li–O₂ batteries, emphasizing MOFs' ability to serve as advanced cathode materials that enhance ORR and OER kinetics. Key mechanisms facilitating these improvements include the catalytic activity of central metal ions, the crucial role of tailored porous structures, the design flexibility of organic ligands to optimize reaction interfaces, and hybrid material strategies. Additionally, by integrating MOFs as semiconductor materials in photo-assisted Li–O₂ batteries, further enhancements in photocatalytic efficiency have been realized. Remarkably, the hierarchical porous structure of MOFs enables efficient lithium-ion transport and the accommodation of discharge products, thereby improving overall cycling stability and energy density. Besides, we have also discussed the preparation methods of MOFs and MOF-derived materials for cathode materials in Li–O₂ batteries, with a focus on their practical applications in this field.

Despite these significant milestones, there remain challenges in scaling MOF technologies for practical use in Li–O₂ systems. Laboratory experiments have consistently demonstrated exceptional performance, yet the transition to industrial-scale applications requires additional breakthroughs, particularly in cost-effective and eco-friendly synthesis protocols. Furthermore, stability concerns, specifically under complex, real-world operating conditions such as air exposure and moisture, limit current utilization. To fully harness the potential of MOFs in energy storage systems, future research should prioritize the following strategies:

(1) Advanced structural design: researchers should further optimize MOF structures, focusing on hierarchical architectures, tunable pore chemistry, and targeted ligand design. Innovations like the incorporation multimetal nodes and functional moieties are pivotal.

(2) Integrative hybrid systems: developing MOF hybrid materials, such as coupling with conductive polymers, metal nanoparticles, or advanced functional carbon materials, will improve electrical conductivity, enhance mass transfer, and address catalytic limitations.

(3) Photo-electrochemical integration: exploration of advanced photo-electrocatalytic functionalities using MOFs should focus on constructing multispectral absorption systems to maximize light utilization, extending the efficiency of photo-assisted Li–O₂ batteries. Combining these with renewable energy integration could revolutionize green energy technologies.

(4) Industrial scalability: achieving cost-efficient, large-scale production of MOFs will require innovative methods like flow chemistry, green chemistry approaches, and the use of earth-abundant materials. Addressing these factors is critical for commercialization.

(5) Stability and real-life operating conditions: addressing stability under operational conditions is paramount. Exploring strategies like encapsulating active sites, defect passivation, and robust coating of MOFs could substantially enhance battery performance and lifespan in ambient air environments.

In conclusion, the integration of MOFs into Li–O₂ batteries not only addresses existing challenges but also unlocks new opportunities for innovation. As the field progresses, MOFs are poised to play a pivotal role in shaping the future of energy storage, contributing to a cleaner and more sustainable energy landscape.

Conflicts of interest

There are no conflicts to declare.

Data availability

No new primary data were generated or analysed in this review. All discussed findings are based on previously published studies cited throughout the manuscript.

Acknowledgements

This work was supported by the NSFC (21972070) in China.

Notes and references

1. D. M. D’Alessandro, B. Smit and J. R. Long, Angew. Chem., Int. Ed., 2010, 49, 6058–6082.
2. M. M. Thackeray, C. Wolverton and E. D. Isaacs, Energy Environ. Sci., 2012, 5, 7854–7863.
3. K. Abraham and Z. Jiang, J. Electrochem. Soc., 1996, 143, 1.
4. B. Zhu, Z. Liang, D. Xia and R. Zou, Energy Storage Mater., 2019, 23, 757–771.
5. J. Read, J. Electrochem. Soc., 2002, 149, A1190–A1195.
6. J. Read, K. Mutolo, M. Ervin, W. Behl, J. Wolfenstine, A. Driedger and D. Foster, J. Electrochem. Soc., 2003, 150, A1351.
7. J. Read, J. Electrochem. Soc., 2005, 153, A96.
8. T. Ogasawara, A. Débart, M. Holzapfel, P. Novák and P. G. Bruce, J. Am. Chem. Soc., 2006, 128, 1390–1393.
9. H. Wang, X. Wang, M. Li, L. Zheng, D. Guan, X. Huang, J. Xu and J. Yu, Adv. Mater., 2020, 32, e2002559.
10. F. Cheng and J. Chen, Chem. Soc. Rev., 2012, 41, 2172–2192.
11. X. Zhang, A. Chen, M. Zhong, Z. Zhang, X. Zhang, Z. Zhou and X.-H. Bu, Electrochem. Energy Rev., 2019, 2, 29–104.
12. Y. Qian, F. Zhang and H. Pang, Adv. Funct. Mater., 2021, 31, 2104231.
13. H. Y. Xu, P. B. Geng, W. C. Feng, M. Du, D. J. Kang and H. Pang, Nano Res., 2024, 17, 3472–3492.
14. D. Wu, Z. Guo, X. Yin, Q. Pang, B. Tu, L. Zhang, Y.-G. Wang and Q. Li, Adv. Mater., 2014, 26, 3258–3262.
15. T. Liu, F. Yang, G. Cheng and W. Luo, Small, 2018, 14, 1703748.
16. L. Liu, Q. He, S. Dong, M. Wang, Y. Song, H. Diao and D. Yuan, J. Colloid Interface Sci., 2024, 666, 35–46.
17. X.-Y. Yang, L.-H. Chen, Y. Li, J. C. Rooke, C. Sanchez and B.-L. Su, Chem. Soc. Rev., 2017, 46, 481–558.
18. Z. Lyu, G. J. Lim, R. Guo, Z. Kou, T. Wang, C. Guan, J. Ding, W. Chen and J. Wang, Adv. Funct. Mater., 2019, 29, 1806658.
19. D. Sun, W. Liu, M. Qiu, Y. Zhang and Z. Li, Chem. Commun., 2015, 51, 2056–2059.
20. Y. L. Tao, T. Wang, X. Y. Yu, K. Gong, H. Gong, H. X. Chen, X. L. Fan, A. D. Zhang, X. L. Huang, K. Chang and J. P. He, Chem. Sci., 2024, 15, 17073–17083.
21. G. Qiao, D. Guan, S. Yuan, H. Rao, X. Chen, J. Wang, J. Qin, J. Xu and J. Yu, J. Am. Chem. Soc., 2021, 143, 14253–14260.
22. P. Dong, X. Zhang, W. Hiscox, J. Liu, J. Zamora, X. Li, M. Su, Q. Zhang, X. Guo and J. McCloy, Adv. Mater., 2023, 35, 2211841.
23. Z. Wang, H. Zhou, C. Meng, W. Xiong, Y. Cai, P. Hu, H. Pang and A. Yuan, ACS Appl. Energy Mater., 2020, 3, 4265–4274.
24. H. Liu, H. Pan, M. Yan, X. Zhang and Y. Jiang, Adv. Mater., 2023, 35, 2300888.
25. Y.-C. Lu, Z. Xu, H. A. Gasteiger, S. Chen, K. Hamad-Schifferli and Y. Shao-Horn, J. Electrochem. Soc., 2010, 157, A1016.
26. S. Ding, X. Yu, Z.-F. Ma and X. Yuan, J. Mater. Chem. A, 2021, 9, 8160–8194.
27. Z. Lyu, Y. Zhou, W. Dai, X. Cui, M. Lai, L. Wang, F. Huo, W. Huang, Z. Hu and W. Chen, Chem. Soc. Rev., 2017, 46, 6046–6072.
28. C. O. Laoire, S. Mukerjee, K. Abraham, E. J. Plichta and M. A. Hendrickson, J. Phys. Chem. C, 2009, 113, 20127–20134.
29. S. Ganapathy, B. D. Adams, G. Stenou, M. S. Anastasaki, K. Goubitz, X.-F. Miao, L. F. Nazar and M. Wagemaker, J. Am. Chem. Soc., 2014, 136, 16335–16344.
30. W.-J. Kwak, Rosy, D. Sharon, C. Xia, H. Kim, L. R. Johnson, P. G. Bruce, L. F. Nazar, Y.-K. Sun and A. A. Frimer, Chem. Rev., 2020, 120, 6626–6683.
31. X. Zhang, L. Guo, L. Gan, Y. Zhang, J. Wang, L. R. Johnson, P. G. Bruce and Z. Peng, J. Phys. Chem. Lett., 2017, 8, 2334–2338.
32. H. Lim, E. Yilmaz and H. R. Byon, J. Phys. Chem. Lett., 2012, 3, 3210–3215.
33. K. C. Lau, J. Lu, X. Luo, L. A. Curtiss and K. Amine, ChemPlusChem, 2015, 80, 336–343.
34. R. Black, S. H. Oh, J.-H. Lee, T. Yim, B. Adams and L. F. Nazar, J. Am. Chem. Soc., 2012, 134, 2902–2905.
35. M. C. Policano, C. G. Anchieta, T. C. Nepel, R. M. Filho and G. Doubek, ACS Appl. Energy Mater., 2022, 5, 9228–9240.
36. T. Liu, J. T. Frith, G. Kim, R. N. Kerber, N. Dubouis, Y. Shao, Z. Liu, P. C. Magusin, M. T. Casford and N. Garcia-Araez, J. Am. Chem. Soc., 2018, 140, 1428–1437.
37. T. Liu, M. Leskes, W. Yu, A. J. Moore, L. Zhou, P. M. Bayley, G. Kim and C. P. Grey, Science, 2015, 350, 530–533.
38. Z. Gao, I. Temprano, J. Lei, L. Tang, J. Li, C. P. Grey and T. Liu, Adv. Mater., 2023, 35, 2201384.
39. X. Zhang, P. Dong, J.-I. Lee, J. T. Gray, Y.-H. Cha, S. Ha and M.-K. Song, Energy Storage Mater., 2019, 17, 167–177.
40. C. Wang, Z. Zhang, W. Liu, Q. Zhang, X. G. Wang, Z. Xie and Z. Zhou, Angew. Chem., Int. Ed., 2020, 59, 17856–17863.
41. G. R. Sun, D. M. Yang, Z. X. Zhang, Y. Wang, W. Lu and M. Feng, J. Adv. Ceram., 2023, 12, 747–759.
42. M. Z. Yu, X. D. Ren, L. Ma and Y. Y. Wu, Nat. Commun., 2014, 5, 5111.
43. Z. Zhu, K. Chan, F. Li and J. Chen, Chem. Soc. Rev., 2022, 51, 1846–1860.
44. A. J. Medford, A. Vojvodic, J. S. Hummelshøj, J. Voss, F. Abild-Pedersen, F. Studt, T. Bligaard, A. Nilsson and J. K. Nørskov, J. Catal., 2015, 328, 36–42.
45. S. Kitagawa, R. Kitaura and S. Noro, Angew. Chem., Int. Ed., 2004, 43, 2334–2375.
46. J. Zhou and B. Wang, Chem. Soc. Rev., 2017, 46, 6927–6945.
47. J. Liu, D. Xie, W. Shi and P. Cheng, Chem. Soc. Rev., 2020, 49, 1624–1642.
48. X. Hu, Q. Liu, K. Lin, C. Han and B. Li, J. Mater. Chem. A, 2021, 9, 20837–20856.
49. R. A. Kharod, J. L. Andrews and M. Dincă, Annu. Rev. Mater. Res., 2022, 52, 103–128.
50. L. Q. Wang, Y. Q. Zhu, C. L. Du, X. L. Ma and C. B. Cao, J. Mater. Chem. A, 2020, 8, 24895–24919.
51. S. H. Kim, Y. J. Lee, D. H. Kim and Y. J. Lee, ACS Appl. Mater. Interfaces, 2018, 10, 660–667.
52. X. X. Wang, D. Y. Du, H. Y. Xu, Y. Yan, X. J. Wen, L. F. Ren and C. Z. Shu, Chem. Eng. J., 2023, 452, 139524.
53. S. Miao, P. G. Pan, Y. Zhang, S. L. Huo, D. F. Wu and S. C. Yu, New J. Chem., 2025, 49, 3855–3860.
54. P. Pan, S. Miao, Y. Zhang, S. L. Huo, D. F. Wu and S. C. Yu, Dalton Trans., 2025, 54, 4437–4443.
55. W. Yan, Z. Guo, H. Xu, Y. Lou, J. Chen and Q. Li, Mater. Chem. Front., 2017, 1, 1324–1330.
56. Y. Zhou, Q. Gu, Y. Xin, X. Tang, H. Wu and S. Guo, Nano Lett., 2023, 23, 10600–10607.
57. J. Yuan, J.-S. Yu and B. Sundén, J. Power Sources, 2015, 278, 352–369.
58. Y. He, L. Y. Ding, J. Cheng, S. W. Mei, X. L. Xie, Z. Y. Zheng, W. Y. Pan, Y. Z. Qin, F. D. Huang, Y. Peng and Z. Deng, Adv. Mater., 2023, 35, 2308134.
59. W. J. Yan, Z. Y. Guo, H. S. Xu, Y. B. Lou, J. X. Chen and Q. W. Li, Mater. Chem. Front., 2017, 1, 1324–1330.
60. J. Rocha, L. D. Carlos, F. A. A. Paz and D. Ananias, Chem. Soc. Rev., 2011, 40, 926–940.
61. Z. Zhang and Z. Zheng, Lanthanide Metal-Organic Frameworks, 2015, pp. 297–367.
62. D. Liu, X. M. Zhang, Y. J. Wang, S. Y. Song, L. F. Cui, H. B. Fan, X. C. Qiao and B. Z. Fang, Nanoscale, 2020, 12, 9524–9532.
63. Y. Xu, H. Xue, X. Li, X. Fan, P. Li, T. Zhang, K. Chang, T. Wang and J. He, Nano Res. Energy, 2023, 2, e9120052.
64. C. Zhang, N. Dandu, S. Rastegar, S. N. Misal, Z. Hemmat, A. T. Ngo, L. A. Curtiss and A. Salehi-Khojin, Adv. Energy Mater., 2020, 10, 2000201.
65. W. H. Choi, B. C. Moon, D. G. Park, J. W. Choi, K. H. Kim, J. S. Shin, M. G. Kim, K. M. Choi and J. K. Kang, Adv. Sci., 2020, 7, 2000283.
66. Y. J. Zhang, G. Ben, H. Khodja, M. Gauthier, E. Foy, B. Goetz, C. Serre, S. C. Leconte, V. Pimenta and S. Surblé, Materials, 2021, 14, 4618.
67. D. Zhu, C. Guo, J. Liu, L. Wang, Y. Du and S.-Z. Qiao, Chem. Commun., 2017, 53, 10906–10909.
68. M. W. Yuan, R. Wang, W. B. Fu, L. Lin, Z. M. Sun, X. G. Long, S. T. Zhang, C. Y. Nan, G. B. Sun, H. F. Li and S. L. Ma, ACS Appl. Mater. Interfaces, 2019, 11, 11403–11413.
69. X. S. Lang, C. X. Dong, K. D. Cai, L. Li and Q. G. Zhang, Int. J. Energy Res., 2020, 44, 1256–1263.
70. Z. Sun, L. Lin, J. He, D. Ding, T. Wang, J. Li, M. Li, Y. Liu, Y. Li and M. Yuan, J. Am. Chem. Soc., 2022, 144, 8204–8213.
71. S. J. Kang, M. Hong, J. H. Won, B. Han, J. K. Kang and H. M. Jeong, Adv. Energy Mater., 2025, 15, 2402651.
72. L. Majidi, A. Ahmadiparidari, N. N. Shan, S. K. Singh, C. J. Zhang, Z. H. Huang, S. Rastegar, K. Kumar, Z. Hemmat, A. T. Ngo, P. Zapol, J. Cabana, A. Subramanian, L. A. Curtiss and A. Salehi-Khojin, Small, 2022, 18, 2102902.
73. W. X. Zhang, S. Tang, Z. H. Chen, X. C. Xiong, B. Chen, K. Wu, G. H. Xu, S. J. Cheng and Y. C. Cao, Chem. Eng. J., 2022, 430, 132489.
74. X. Z. Zheng, M. W. Yuan, P. Y. Su, J. S. Xu and G. B. Sun, Appl. Surf. Sci., 2025, 692, 162742.
75. X. H. Zhang, P. P. Dong, J. I. Lee, J. T. Gray, Y. H. Cha, S. Ha and M. K. Song, Energy Storage Mater., 2019, 17, 167–177.
76. Y.-C. Lu, Z. Xu, H. A. Gasteiger, S. Chen, K. Hamad-Schifferli and Y. Shao-Horn, J. Am. Chem. Soc., 2010, 132, 12170–12171.
77. M. I. Hossain, F. K. Tareq and S. Rudra, Energy Storage Mater., 2025, 104025.
78. G. Wang, Y. Yang, Q. Zhang, Z. Xie and Z. Zhou, Coord. Chem. Rev., 2024, 511, 215879.
79. Q.-G. Zhai, X. Bu, X. Zhao, D.-S. Li and P. Feng, Acc. Chem. Res., 2017, 50, 407–417.
80. H. Yu, D. Liu, Z. Fu, S. Wang, X. Zuo and Y. Zhang, ChemSusChem, 2024, e202401697,  DOI:10.1002/cssc.202401697.
81. Y. Yang, X. Hu, G. Wang, J. Han, Q. Zhang, W. Liu, Z. Xie and Z. Zhou, Adv. Funct. Mater., 2024, 34, 2315354.
82. Q. L. Lv, Z. Zhu, S. Zhao, L. B. Wang, Q. Zhao, F. J. Li, L. A. Archer and J. Chen, J. Am. Chem. Soc., 2021, 143, 9696.
83. A. Dhakshinamoorthy, A. M. Asiri and H. Garcia, Angew. Chem., Int. Ed., 2016, 55, 5414–5445.
84. W. Wang, X. Xu, W. Zhou and Z. Shao, Adv. Sci., 2017, 4, 1600371.
85. Y. L. Tao, X. L. Fan, X. Y. Yu, K. Gong, Y. J. Xia, H. Gong, H. X. Chen, X. L. Huang, A. D. Zhang, T. Wang and J. P. He, Small, 2024, 20, 2403683.
86. B. Wen, Y. Huang, Z. Jiang, Y. Wang, W. Hua, S. Indris and F. Li, Adv. Mater., 2024, 2405440.
87. X. Qiu, Y. Zhu, X. Zhang, Y. Zhang, L. T. Menisa, C. Xia, S. Liu and Z. Tang, Sol. RRL, 2020, 4, 1900449.
88. Z. Hu, Y. Yang, B. Ye, X. Hu, L. Chen, H. Yu, Y. Zhou, Z. Xie and Z. Zhou, Adv. Mater. Interfaces, 2023, 10, 2300074.
89. L. Y. Wu, Y. F. Mu, X. X. Guo, W. Zhang, Z. M. Zhang, M. Zhang and T. B. Lu, Angew. Chem., Int. Ed., 2019, 58, 9491–9495.
90. D. F. Wu, Z. Y. Guo, X. B. Yin, Q. Q. Pang, B. B. Tu, L. J. Zhang, Y. G. Wang and Q. W. Li, Adv. Mater., 2014, 26, 3258–3262.
91. X. X. Wang, D. Y. Du, Y. Yan, L. F. Ren, H. Y. Xu, X. J. Wen, T. Zeng, G. L. Tian, S. Liu, F. X. Fan and C. Z. Shu, Energy Storage Mater., 2023, 63, 103033.
92. W. Xu, X. Pei, C. S. Diercks, H. Lyu, Z. Ji and O. M. Yaghi, J. Am. Chem. Soc., 2019, 141, 17522–17526.
93. P. Albertus, S. Babinec, S. Litzelman and A. Newman, Nat. Energy, 2018, 3, 16–21.
94. C. Yuan, J. Li, P. Han, Y. Lai, Z. Zhang and J. Liu, J. Power Sources, 2013, 240, 653–658.
95. C. L. Miao, X. X. Wang, D. H. Guan, J. X. Li, H. F. Wang and J. J. Xu, Adv. Funct. Mater., 2024, 34, 2307150.
96. H. X. Liu, L. Y. Zhao, Y. Xing, J. N. Lai, L. Li, F. Wu, N. Chen and R. J. Chen, ACS Appl. Energy Mater., 2022, 5, 7185–7191.
97. X. X. Wang, X. W. Chi, M. L. Li, D. H. Guan, C. L. Miao and J. J. Xu, Adv. Funct. Mater., 2022, 32, 2113235.
98. L. Xu, X. Xiao, H. Tu, F. Zhu, J. Wang, H. Liu, W. Huang, W. Deng, H. Hou and T. Liu, Adv. Mater., 2023, 35, 2303193.
99. H. F. Wang, D. H. Guan, X. K. Chang, X. Y. Ma, X. Y. Yuan, X. X. Wang and J. J. Xu, Small, 2025, 21, 2411747.
100. W. L. Teo, W. A. Zhou, C. Qian and Y. L. Zhao, Mater. Today, 2021, 47, 170–186.
101. P. Priecel and J. A. Lopez-Sanchez, ACS Sustainable Chem. Eng., 2019, 7, 3–21.
102. S. H. Jhung, J. H. Lee, J. W. Yoon, C. Serre, G. Férey and J. S. Chang, Adv. Mater., 2007, 19, 121–124.
103. J. Wegner, S. Ceylan and A. Kirschning, Adv. Synth. Catal., 2012, 354, 17–57.
104. M. Rubio-Martinez, M. P. Batten, A. Polyzos, K. C. Carey, J. I. Mardel, K. S. Lim and M. R. Hill, Sci. Rep., 2014, 4, 5443.
105. M. Taddei, D. A. Steitz, J. A. van Bokhoven and M. Ranocchiari, Chem. – Eur. J., 2016, 22, 3245–3249.
106. M. Zhong, X. Zhang, D. H. Yang, B. Zhao, Z. J. Xie, Z. Zhou and X. H. Bu, Inorg. Chem. Front., 2017, 4, 1533–1538.
107. A. Li, X. Zhang, Z. J. Xie, Z. Chang, Z. Zhou and X. H. Bu, Inorg. Chem., 2018, 57, 14476–14479.
108. Y. Y. Xue, Y. B. Guo, Q. M. Zhang, Z. J. Xie, J. P. Wei and Z. Zhou, Nano-Micro Lett., 2022, 14, 162.
109. Z. Y. Lyu, G. J. H. Lim, R. Guo, Z. K. Kou, T. T. Wang, C. Guan, J. Ding, W. Chen and J. Wang, Adv. Funct. Mater., 2019, 29, 1806658.
110. P. Wang, Y. Y. Ren, R. T. Wang, P. Zhang, M. J. Ding, C. X. Li, D. Y. Zhao, Z. Qian, Z. W. Zhang, L. Y. Zhang and L. W. Yin, Nat. Commun., 2020, 11, 1576.
111. Y. T. Qian, F. F. Zhang, S. X. Zhao, C. Q. Bian, H. Mao, D. J. Kang and H. Pang, Nano Energy, 2023, 111, 108415.
112. L. W. Zhao, J. J. Feng, A. Abbas, C. L. Wang and H. C. Wang, Small, 2023, 19, 2302953.

---