Toward highly durable aqueous zinc ion batteries: a review of MOFs/MOF-derived cathode materials

Abstract

The key to developing aqueous zinc ion batteries (AZIBs) lies in designing advanced cathodes that match well with the zinc anode. Metal–organic framework (MOF)-based materials have emerged as a focal point in research due to their unique benefits in energy storage. However, researchers lack clear guidance on the use of MOFs and their derived materials as AZIBs electrode materials, as well as a systematic exploration of their Zn²⁺ storage mechanisms. Herein, we summarize the recent progress in the research of pristine MOFs and their derivatives for high-performance AZIBs. Moreover, we provide a detailed understanding of the energy storage mechanisms. Finally, we discuss the challenges and future perspectives regarding MOFs and their derivatives for next-generation aqueous energy devices. This review provides insights into innovations in MOF-based cathode materials and inspiration for the development of future efficient energy storage and conversion technologies.

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1. Introduction

As environmental pollution worsens, reducing carbon emissions and achieving carbon neutrality have become increasingly critical.^1–3 In recent decades, many clean energy sources have been developed to reduce dependence on traditional energy sources.^4,5 However, their inherent intermittency and variability prevent them from meeting the demand for stable grid operation. At present, the electronics market is largely driven by lithium-ion batteries (LIBs), which are favored for their high energy density.^6,7 Nevertheless, the flammability of organic electrolytes poses significant safety issues for LIBs. Consequently, new rechargeable batteries are being investigated worldwide, with a dual priority on enhancing safety and reducing costs. Aqueous zinc ion batteries (AZIBs) are considered an attractive alternative due to the abundant availability of zinc and superior ionic conductivity merits.^8–10 Compared to monovalent Li⁺ or Na⁺, bivalent Zn²⁺ can shift two electrons per ion in electrochemical processes, thus showing a high theoretical capacity.^11,12 However, AZIBs exhibit slow (de)intercalation kinetics and rapid capacity decay under intense electrostatic interactions.¹³

To overcome these challenges, the strategic design of cathode materials is important for optimizing AZIB performance. Metal–organic frameworks (MOFs) possess customizable topologies, abundant porosity and adjustable sites.^14–16 Many MOF materials have been synthesized by altering the types of metal ions/clusters and organic ligands. These materials show significant advantages in electrochemical energy storage.¹⁷ For example, their porous nanostructure facilitates charge storage capacity. Moreover, the pore size can be adjusted by modifying the length of the organic ligand, thereby optimizing ion transport channels.^18,19 In addition, MOF-derived materials, which are carefully controlled, can be tailored to enhance their electrochemical properties, sometimes surpassing those of the original MOFs.^20–22 Motivated by these strengths, significant research efforts have focused on utilizing MOFs and MOF-derived micro/nanomaterials for AZIB cathodes.^23–26 From Fig. 1a, MOF-based materials are categorized into four primary types: pristine MOFs, MOF composites, MOF derivatives, and MOF composite derivatives.

Fig. 1 (a) Classification of MOF-based materials; (b) schematic diagram of the energy storage mechanisms of MOF-based cathode materials for AZIBs.

Most existing reviews for AZIBs have concentrated on cathode materials, dendrite-free Zn anodes and electrolytes.²⁷ However, there is little comprehensive discussion on the structural and compositional aspects of MOFs/MOF derivatives for AZIBs. Recently, various MOF derivatives have been used for AZIBs.^28–30 Herein, we summarize the current breakthroughs in MOFs and their derived materials for AZIBs. Then, we introduce the relationship between MOF cathode materials and electrochemical properties, and their energy storage mechanisms (Fig. 1b). Finally, we propose the current issues in this direction and the expectation of designing and synthesizing novel MOFs for AZIBs.

2. MOFs and MOF derivative cathode materials

In recent decades, extensive research has been dedicated to the creation of next-generation MOFs. The physical–chemical properties of MOFs are manipulated in terms of molecular structure, as well as the micro/nanostructures of MOFs are customized to obtain the specific properties.^31–33 MOF-derived materials are substances with new compositions and structures obtained from MOF precursors through pyrolysis, chemical etching and ion exchange processes. They possess significant differences from their precursors.³⁴

2.1 Pristine MOFs

Considering appropriate interlayer effects, Prussian blue (PB) and its analogues (PBAs) are the most suitable MOFs for AZIBs.^35,36 Their crystal structure of large gap nanocavity endows effective insertion/extraction of Zn ions.³⁷ Liu's group prepared two types of zinc hexacyanoferrates with diamond-shaped structures through the high-temperature co-precipitation method.³⁸ AZIBs are assembled according to the process shown in Fig. 2a. As shown in Fig. 2b, the Zn//ZnHCF batteries deliver discharged capacities of 65.4 mA h g⁻¹ at 1 C rate. Moreover, the curves present clear discharge platforms at 1.6 V. To achieve the goal of increasing the capacity of ZnHCF electrodes, Huang et al. coated a polyaniline (PANI) layer on ZnHCF.³⁹ In Fig. 2c, the Zn//ZnHCF-PANI batteries show two flat charge–discharge platforms and a specific capacity of 130 mA h g⁻¹ (0.2 A g⁻¹). The enhanced electrochemical behavior results from the PANI coating's role in preventing dissolution.

Fig. 2 (a) Schematic diagram of the Zn//ZnHCF battery. (b) Rate capability of the Zn//ZnHCF battery; reproduced with permission from ref. 38. Copyright 2014, WILEY-VCH. (c) Galvanostatic charge and discharge curves at different current densities of the Zn//ZnHCF-PANI battery; reproduced with permission from ref. 39. Copyright 2022, The Royal Society of Chemistry. (d) Schematic of the Mn(BTC) cathode//ZIF-8@Zn anode ZIBs. (e) Cycling performance at 1 A g⁻¹ in 2 M ZnSO₄ + 0.1 M MnSO₄ electrolyte; reproduced with permission from ref. 43. Copyright 2020, Springer. (f) Rate capability of Zn//Mn-H3BTC-MOF-4 battery; reproduced with permission from ref. 44. Copyright 2021, American Chemical Society.

As the central node of MOFs, metal ions are uniformly distributed in the whole frame network.^40–42 Kang's team chose four types of metal ions to connect two organic ligands to form five MOFs.⁴³ Among them, Mn(BTC) possesses favorable Zn²⁺ storage capacity as a cathode material for AZIBs. Moreover, a ZIF-8@Zn anode was prepared using ZIF-8 coating to regulate the dissolution–deposition behavior. Fig. 2d presents the schematic of the Mn(BTC) cathode//ZIF-8@Zn anode ZIBs. They maintain 92% of the initial capacity after 900 cycles at 1 A g⁻¹ (Fig. 2e). According to the literature, the coordination unsaturated MOFs were synthesized as electrode materials,⁴⁴ which ensures efficient Zn²⁺ transport and electron exchange. The rate performance of three samples is tested in Fig. 2f. The Zn//Mn-H3BTC-MOF-4 batteries show a reduction in capacity from 138 to 98 mA h g⁻¹ (from 0.1 to 3 A g⁻¹). It is superior to the other two electrodes due to its appropriate unsaturated coordination, which enhances the electronic/ionic conductivity of MOFs.

Compared with traditional MOFs, conductive metal–organic frameworks (cMOFs) with unique crystal structures and excellent conductivity possess broad application prospects in related electrochemical energy fields.^45–47 Stoddart and co-workers employed two-dimensional (2D) cMOFs with convenient diffusion channels as cathode materials for AZIBs.⁴⁸ From Fig. 3a, it can be observed that the synthesized Cu₃(HHTP)₂ is composed of hexagonal 2D flakes arranged in a sliding parallel configuration along the c-axis (pores of ∼2 nm). The reduction-induced electron density changes in Cu₃(HHTP)₂ are shown in Fig. 3b. When Cu atoms provide 6.9 additional electrons to Cu₃(HHTP)₂, both Cu atoms and linkers gain additional electrons. It indicates that the Cu atoms and quinoid configuration in the Cu₃(HTTP)₂ provide active sites for redox reactions involved in zinc ion storage. The interface resistance (R_I) of Cu₃(HHTP)₂ electrodes in aqueous and organic electrolytes is 150 and 16 000 Ω cm², respectively (Fig. 3c and d). The low R_I value results in fast diffusion kinetics and durable cycling life.

Fig. 3 (a) Schematic diagram of the Zn//Cu₃(HHTP)₂ battery, structure and coordination unit of Cu₃(HHTP)₂ during the redox process. (b) Changes in electron density upon the reduction of Cu₃(HHTP)₂. EIS spectra of Cu₃(HHTP)₂ electrodes in (c) an aqueous and (d) an organic electrolyte; reproduced with permission from ref. 48. Copyright 2019, Springer Nature Limited. (e) Discharging/charging curves of Cu-BTA-H/Ni-BTA-H at 0.2 A g⁻¹; reproduced with permission from ref. 49. Copyright 2023, American Chemical Society. (f) Schematic representation of the one-pot synthesis and digestion of Cu-TBPQ MOF; grey, red, white, and blue spheres represent C, O, H, and Cu atoms, respectively. (g) Calculated MESP distribution of Cu-TBPQ MOF; reproduced with permission from ref. 50. Copyright 2020, WILEY-VCH.

Moreover, Liang et al. reported a type of 1D cMOFs (Cu-BTA) as the cathode for AZIBs.⁴⁹ It is synthesized by coordination ultrasmall 1,2,4,5-benzenetetramine (BTA) ligands with polyvalent copper ions. BTA can effectively increase the proportion of active groups (–C=N) in the material system, while the synergistic redox reaction of transition metal ions (Cu²⁺/Cu⁺) provides additional energy storage capacity as the second active site. Therefore, the high crystallinity Cu-BTA electrodes possess a reversible capacity of 330 mA h g⁻¹ at 0.2 A g⁻¹ (Fig. 3e). Quinone compounds are easily soluble in electrolytes, which hinders their rate performance and cycling life. Considering the above-mentioned limitations, Chen's group tailored the anthraquinone-based polythematic catechol ligands and successfully prepared quinone-containing copper-catecholate MOFs (Cu-TBPQ), as illustrated in Fig. 3f.⁵⁰ From Fig. 3g, it is confirmed that adding quinone groups in cMOFs significantly enhances the density of redox-active sites. The novel Cu-TBPQ MOFs with quinone groups enhance cation absorption behavior and stabilize electrochemical performance.

2.2 Derived MOFs

Thus far, most of the reported pristine MOFs possess problems with poor rate performance and cycling stability as cathode materials for AZIBs.⁵¹ Satisfactorily, the appearance of MOF-derived materials has paved the way for customizing structural and morphological characteristics.^52,53 In recent years, ones have investigated many MOF-derived materials, which were as follows:

2.2.1 Mn-based MOF-derived materials. Mn-based oxides have attracted extensive research attention, featuring their adjustable storage performance.^54,55 However, the Jahn–Teller effect may be generated with the (de)intercalation of Zn²⁺ during cycling, resulting in structure distortion.^56,57 Therefore, the rational design of the structure of manganese oxide is a feasible strategy to improve its long-term stability. Constructing hollow structures is an effective route to buffer large volume changes caused by ion insertion/extraction.^58,59 Lou et al. successfully prepared heterostructured Mn₂O₃–ZnMn₂O₄ hollow octahedra (MO-ZMO HOs) using Mn-MIL octahedra as precursors through cation exchange reaction and annealing calcination.⁶⁰ Fig. 4a shows the FESEM and TEM images of Mn-MIL SO, Zn/Mn-MIL SO and MO-ZMO HO materials. The MO-ZMO HO sample still retains the octahedral shape. However, the particle size decreased due to significant shrinkage during the calcination process. By utilizing the synergistic strategy of the double phases and the rational design of hollow heterostructures, MO-ZMO HO electrodes possess dense active sites. The Zn//MO-ZMO HO batteries deliver a specific capacity of 247.4 mA h g⁻¹ at 0.1 A g⁻¹ (Fig. 4b).

Fig. 4 (a) Schematic illustration of the synthetic process of MO-ZMO HOs and FESEM and TEM images of Mn-MIL SO, Zn/Mn-MIL SO and MO-ZMO HO samples. (b) Rate performance based on discharging curves; reproduced with permission from ref. 60. Copyright 2021, WILEY-VCH. (c) Illustration of the Mn valence state distribution in ZMO QD@C. (d) FTIR spectra and (e) long-term cycle test of ZMO QD@C at 1 A g⁻¹; reproduced with permission from ref. 63. Copyright 2022, WILEY-VCH.

The carbon framework preserves structural integrity during ion diffusion, mitigates metal compounds’ poor conductivity, and stabilizes the host material.^61,62 Niu and co-workers combined Mn₃O₄ quantum dots derived from Mn-MIL-100 with carbon.⁶³ ZnMn₂O₄ quantum dots (ZMO QDs) are integrated into the porous carbon framework via in situ electrochemical induction. The obtained ZMO QD@C electrode material presents a mixed manganese valence state. From Fig. 4c, Mn⁴⁺ is concentrated at the ZMO QD cathode–carbon matrix interface, while Mn³⁺ predominantly resides inside the ZMO QDs. Furthermore, the Mn–O–C bonds appear in the FTIR spectra of ZMO QD@C samples (Fig. 4d). They can restrain the Jahn–Teller effect and minimize the manganese dissolution of the discharge product. As shown in Fig. 4e, the capacity still remains at 143.9 mA h g⁻¹ after 1500 times of cycling, achieving a capacity retention of 86.4% at 1 A g⁻¹.

The oxygen defect (O_d) strategy is widely used in material modification to enhance the local electronic conductivity of electrodes. The O_d can form unsaturated bonds on the surface or inside of materials, increasing the number of active sites. Moreover, this strategy can regulate the redox potential of the cathode and improve the reversibility of electrochemical reactions.^64–66 In the literature, oxygen-deficient nanorod arrays (NAs) were synthesized by a solvothermal method.⁶⁷ As illustrated in Fig. 5a, the O_d-Mn₃O₄@C nanorods grow tightly on each carbon fiber. The octahedral MnO₆ transforms into a pyramid following the lack of one oxygen atom. The carbon derived from Mn-MOFs facilitates O_d incorporation. The formed carbon skeleton also endows the electrode materials with a conductive network. Electron energy loss spectroscopy (EELS) indicates that O_d is uniformly distributed throughout the nanorods (Fig. 5b). In addition, Fig. 5c shows the EPR spectra of the O_d-Mn₃O₄@C NA/CC and Mn₃O₄/CC samples. The former possesses a clear symmetric signal g = 2.0, which is attributed to the presence of O_d. From Fig. 5d, the O_d-Mn₃O₄@C NA/CC electrode delivers a reversible capacity of 396.2 mA h g⁻¹ at 0.2 A g⁻¹. This is superior to the other three samples, further proving that O_d and carbon skeleton optimizes the performance of the composite material. As for the cycling stability shown in Fig. 5e, the Zn//O_d-Mn₃O₄@C NA/CC battery maintains a capacity of 125.2 mA h g⁻¹ after 3000 times of cycling (2 A g⁻¹).

Fig. 5 (a) Schematic illustration of the O_d-Mn₃O₄@C NA/CC nanostructure. (b) Marked position and EELS spectra of O-K and Mn-L edges from the O_d-Mn₃O₄@C NA/CC samples. (c) EPR spectra of the O_d-Mn₃O₄@C NA/CC and Mn₃O₄/CC electrodes. (d) Specific capacity of batteries composed of various cathode materials at different current densities. (e) Cycling stability and coulombic efficiency of the Zn//O_d-Mn₃O₄@C NA/CC battery at 2 A g⁻¹; reproduced with permission from ref. 67. Copyright 2020, WILEY-VCH.

2.2.2 V-based MOF-derived materials. V-MOF derivatives have also been applied as cathodes for AZIBs. MIL-100(V) presents two mesoporous cages of different sizes (25 and 29 Å) and 5.5 and 8.6 Å accessible microporous windows. Mesoporous properties increase active site availability within the open framework.^68,69 Cai et al. designed carbon-modified porous vanadium compound microspheres (p-VO_x@C) using MIL-100(V) as a precursor.⁷⁰ As shown in Fig. 6a, most of the carbon atoms of MIL-100(V) thermal decomposition migrate to the surface of MIL-100(V), resulting in a carbon shell. It then leaves behind a porous core with rich oxygen vacancy. From the TEM images (Fig. 6b), the p-VO_x@C-700 samples possess a core–shell two-phase microstructure. Subsequently, the edge area of the electrode material is observed by HRTEM (dotted red frame). The d-spacings of 2.43 and 2.74 Å match well the (101) plane of VO₂ phase and the (002) plane of V₂O₃ phase, respectively. In addition, V⁴⁺/V³⁺ in the electrode material is oxidized to V⁵⁺ during the activation process (Fig. 6c). The p-HVOx@C is formed with H₂O molecules embedded into the intermediate layer of vanadium oxide. In Fig. 6d, the Zn//p-HVO_x@C-700 batteries deliver specific capacities of 464.3 mA h g⁻¹ at 0.2 A g⁻¹ and 395.2 mA h g⁻¹ at 10 A g⁻¹.

Fig. 6 (a) Preparation and electrochemical activation process diagram of the p-VO_x@C samples. (b) TEM and HRTEM images of p-VO_x@C-700 samples. (c) V 2p of the p-VO_x@C-700 samples during the potentiostatic activation. (d) Discharge and charge profiles at different current densities of p-HVO_x@C-700 electrodes; reproduced with permission from ref. 70. Copyright 2024, Elsevier.

Pristine MIL-88B(V) is composed of trinuclear metal carboxylic acid nodes connected to terephthalic acid (TPA) ligands.^71,72 Simple synthetic blends in the preparation of MOFs may result in their inherent potential not being optimally utilized. It may lead to large agglomeration and damage the structural integrity of the frame.⁷³ He's group reported the production of MIL-88B(V)@rGO composite materials.⁷⁴ The electrode possesses an in situ irreversible transition from MIL-88B(V) to amorphous V₂O₅ during initial charge/discharge. After the subsequent cycling process, the amorphous V₂O₅ phase serves as an electric energy storage host. As shown in Fig. 7a, the shape of CV curves in the first cycle is distinct from those in later cycles. This demonstrates that the MIL-88B(V)@rGO electrode is transformed into new materials. The results in Fig. 7b indicate that the new peaks are well indexed with the orthorhombic V₂O₅ phase (JCPDS no. 41-1426). Moreover, the V₂O₅ phase exhibits broad peaks rather than sharp peaks, which is consistent with the amorphous characteristic from the HRTEM images in Fig. 7c. It is confirmed by DFT calculation that the addition of rGO increases the electronic conductivity of the composites and decreases the diffusion energy barrier of Zn²⁺ (Fig. 7d and e). The MIL-88B(V)@rGO cathode possesses a reversible capacity of 479.6 mA h g⁻¹ at 50 mA g⁻¹. Additionally, Luo et al. synthesized the nanorod-shaped material with a porous carbon skeleton-coated heterostructure (C@VO₂@V₂O₅) using MIL-88B(V) as the precursor (Fig. 7f).⁷⁵ The heterostructure can effectively alleviate the erosion of electrolytes and improve the electron transport capacity of electrodes. From the rate performance in Fig. 7g, Zn//C@VO₂@V₂O₅ batteries achieve a higher capacity than that of the other three samples.

Fig. 7 (a) CV curves of the MIL-88B(V)@rGO sample at 0.1 mV s⁻¹. (b) XRD patterns before and after the first cycle. (c) HRTEM images of the MIL-88B(V)@rGO sample; insertion/extraction energy and possible migration pathway of Zn²⁺ in V₂O₅ (d) and V₂O₅@rGO (e); reproduced with permission from ref. 74. Copyright 2023, WILEY-VCH. (f) Schematic illustration of the synthesis of the C@VO₂@V₂O₅ sample. (g) Rate capability tests; reproduced with permission from ref. 75. Copyright 2023, Elsevier.

As MOFs undergo crystallization, the added mini-sized units will be encapsulated within the porous structure of MOFs. The enclosure effect restricts the aggregation of particles, thus improving the effective availability of materials.⁷⁶ For example, the Ag@MIL-88B(V) material was fabricated by introducing Ag nanoparticles into the solvothermal synthesis of MIL-88B(V).⁷⁷ Then, a layered porous heterostructure (Ag–V₂O₅) was generated in the air through heating (Fig. 8a). To investigate the bonding behavior of Zn²⁺ during cycling, the differential charge density of pristine V₂O₅ and Ag–V₂O₅ was calculated, and is presented in Fig. 8b. The apparent carrier migration between Ag and V₂O₅ may potentially create an internally generated electric field, thereby promoting electron transfer. Fig. 8c shows the Nyquist plots of the Ag–V₂O₅ and V₂O₅ electrodes. The former possesses a lower charge transfer resistance (31 Ω) than that of the latter (220 Ω), further indicating the enhanced electronic conductivity of the Ag–V₂O₅ samples. The hierarchical porous structure endows the Zn//Ag–V₂O₅ batteries with a specific capacity of 426 mA h g⁻¹ at 0.1 A g⁻¹ and a capacity retention of 96.5% after 100 cycles (Fig. 8d).

Fig. 8 (a) Schematic illustration of porous Ag–V₂O₅. (b) Differential charge density with Zn²⁺ intercalation in V₂O₅ and Ag–V₂O₅. (c) EIS plots. (d) Cycling performances of V₂O₅ and Ag–V₂O₅; reproduced with permission from ref. 77. Copyright 2023, Elsevier. (e) Schematic illustration of Zn²⁺ diffusion and Zn²⁺ (de)intercalation energy in crystalline and amorphous V₂O₅ from first-principles calculations. (f) Zn²⁺ diffusion coefficient during discharging. (g) Rate capacity of a-V₂O₅@C; reproduced with permission from ref. 78. Copyright 2020, WILEY-VCH.

Amorphous V₂O₅ and carbon composites (a-V₂O₅@C) have been successfully synthesized by in situ electrochemical transformation of V₂O₃@C crystals obtained from MIL-88B(V).⁷⁸ From Fig. 8e, the amorphous and porous carbon of composite materials endow them with abundant Zn²⁺ isotropic diffusion pathways. Compared with crystalline V₂O₅, amorphous V₂O₅ presents more reaction locations for zinc ions, which are conducive to the diffusion of Zn²⁺. As shown in Fig. 8f, the D_Zn value of a-V₂O₅@C samples was calculated to be 10⁻¹⁰ to 10⁻⁹ cm² s⁻¹. The Zn//a-V₂O₅@C battery exhibits superior rate performance. Moreover, it can continuously provide a capacity of 73 mA h g⁻¹ at a sizeable current density of 200 A g⁻¹ (Fig. 8g).

2.2.3 Other MOF-derived materials. MoS₂ has been widely researched as a potential storage electrode for AZIBs. However, the narrow interlayer spacing (0.62 nm) limits the reduplicative (de)intercalation of [Zn(H₂O)₆]²⁺, resulting in low zinc storage capacity.^79,80 Tang and colleagues constructed N-doped 1T MoS₂ nanoflowers using a Mo-MOF as the precursor through one-step hydrothermal sulfurization.⁸¹ From the EPR spectra shown in Fig. 9a, the N-doped 1T MoS₂ electrode possesses strong symmetric signals at g = 2.002, and it represents the presence of sulfur vacancies. Furthermore, nanoflowers formed by self-supporting ultra-thin MoS₂ nanosheets promote the rapid transfer of ions. The N-doped 1T MoS₂ electrode materials reach a capacity retention rate of 89% after 1000 times cycling at 3 A g⁻¹ (Fig. 9b).

Fig. 9 (a) EPR spectra of three MoS₂ samples. (b) Cycling stability data obtained at 3 A g⁻¹; reproduced with permission from ref. 81. Copyright 2019, American Chemical Society. (c) Cycling stability performance at 2 A g⁻¹ of the two batteries; reproduced with permission from ref. 82. Copyright 2022, Elsevier. (d) Schematic of the synthesis route of the CoO/Ni₂P–Co₂P nanosheet arrays supported on the Ni foam substrate. (e) O 1s spectra and (f) EPR spectra of the CoO/Ni₂P–Co₂P sample and NiCo₂O₄/NiO sample. (g) GCD curves at different current densities of the Zn//CoO/Ni₂P–Co₂P-30 battery; reproduced with permission from ref. 83. Copyright 2021, The Royal Society of Chemistry.

Recently, layered double hydroxides (LDHs) have become a topic of interest in the field of electrochemistry owing to their excellent charge storage ability. Pan et al. prepared a NiAl-LDH/Ni@C cathode with cross-linking nanosheet structures by using porous carbon derived from Ni-MOFs as the precursor.⁸² They expose many electrochemical active sites, accelerating Faraday redox reactions. The Zn//NiAl-LDH/Ni@C batteries provide a specific capacity of 314.9 mA h g⁻¹ with a retention ratio of 95.3% after 2000 cycles at 2 A g⁻¹ (Fig. 9c). In addition, Xu and co-workers constructed a stacked CoO/Ni₂P–Co₂P nanosheet array using surface chemical reconstruction and partial phosphorylation processes,⁸³ as illustrated in Fig. 9d. In the O 1s spectra (Fig. 9e), the concentration of oxygen vacancies in the CoO/Ni₂P–Co₂P sample significantly increases, which can be observed in the EPR test (Fig. 9f). As shown in Fig. 9g, the phosphating sample (CoO/Ni₂P–Co₂P) achieves a desired capacity of 323 mA h g⁻¹ at 2.0 A g⁻¹.

3. Energy storage mechanisms of MOFs and their derived materials

Progress has been made in the research of MOFs and their derivatives as AZIB cathodes. However, their potential energy storage mechanisms remain intricate for MOFs with both organic and inorganic characteristics.⁸⁴ Herein, the energy storage mechanism is discussed in detail, including Zn²⁺ insertion/extraction, H⁺/Zn²⁺ co-insertion and metal ion dissolution/deposition mechanisms.

3.1 Zn²⁺ insertion/extraction mechanism

The (de)intercalation of Zn²⁺ within a host matrix during the charging and discharging process is a universally approved energy storage mechanism. It is similar to the rocking chair mechanism in LIBs, which absorbs and releases energy through redox reactions.^85–87 The pore structure of MOFs provides a satisfying environment for the insertion of Zn²⁺.⁸⁸ Metal nodes can serve as binding sites in the electrochemical reaction process. For example, the CV curves of the MIL-100(V) cathode show two types of redox peaks accompanied by the redox transitions of V⁴⁺ ↔ V³⁺ and V⁵⁺ ↔ V⁴⁺ (Fig. 10a).⁸⁹ In Fig. 10b, the HRTEM results of MIL-100(V) reveal that the lattice spacing increases to 1 nm owing to the insertion of Zn²⁺. In addition, the Zn element is uniformly distributed in the MIL-100(V) electrode by observing the EDX image in a fully discharged state (Fig. 10c). When charged to 1.6 V, the Zn signal decreased significantly with the embedding of zinc ions. The ex situ XPS spectra in Fig. 10d also confirmed the above analysis.

Fig. 10 (a) CV curve of the MIL-100(V)//Zn ZIB at different scan rates. (b) Ex situ HRTEM images of MIL-100(V) at 0.2 V discharged state. (c) Zn elemental mapping of MIL-100(V) at 0.2 V discharged and 1.6 V charged state. (d) Zn 2p XPS spectra of the MIL-100(V) electrode; reproduced with permission from ref. 89. Copyright 2023, Elsevier. (e) Charge/discharge curves under a current density of 100 mA g⁻¹, in which the marked states are selected for ex situ tests and ex situ XRD patterns. (f) Ex situ XPS spectra of Cu 2p; reproduced with permission from ref. 90. Copyright 2022, WILEY-VCH. (g) Charge–discharge profiles and corresponding stages of cathode collection for ex situ investigation. XPS spectrum of (h) N 1s and (i) Ni 2p; reproduced with permission from ref. 91. Copyright 2022, WILEY-VCH.

Moreover, metallic copper sites have been ascertained as possible promoting factors for ion storage. The ex situ XRD is performed to research the energy storage mechanism of the Cu-HHTP/MX heterostructure.⁹⁰ From Fig. 10e, the characteristic peaks of (100) crystal planes exhibit a minor angle change during the (de)insertion of Zn²⁺. Ex situ XPS confirms the reversible Zn²⁺ storage in Cu₃(HHTP)₂/MX, alongside the redox conversion from Cu²⁺ to Cu⁺ (Fig. 10f). Multiple redox reactions appear in MOFs during energy storage, which indicates that various binding sites are involved in the synergistic action of zinc ion storage. Zhi's team proposed a thermal modification strategy to regulate the exposed surface of MOF electrodes.⁹¹ In Fig. 10g–i, the XPS spectra of N 1s and Ni 2p confirm the existence of a two-element participation mechanism in the PFC-8350 sample. Ni and N serve as active centers in the shuttle process of Zn²⁺, with N contributing 60% and Ni 40% to the total capacity, as quantified.

Sun et al. employed ex situ X-ray absorption near-edge structure (XANES) spectroscopy to investigate the oxidation state of Mn in the MnO_x@N–C electrode during the discharge/charge process.⁹² As present in Fig. 11a and b, the average Mn oxidation state in MnO_x@N–C ranges from +2 to +4 in the original state. The oxidation state of Mn increases during discharge as a result of Mn²⁺ solubilization effect. After charging, the Mn valence state shows a mild reduction, but it is still marginally higher than the pristine state. This is caused by the re-integration of dissolved Mn²⁺ into the cathode, leading to an elevated oxidation state of Mn. The MnO_x@N–C samples possess a highly reversible behavior.

Fig. 11 (a) Discharge–charge profile of the MnO_x@N–C electrode. (b) Mn K-edge XANES curves at selected discharge and charge states of the MnO_x@N–C electrode; reproduced with permission from ref. 92. Copyright 2018, WILEY-VCH. (c) CV plots of the ZnMn₂O₄/Mn₂O₃ electrode. (d) Ex situ XRD spectra and (e) ex situ Raman spectra of the ZnMn₂O₄/Mn₂O₃ electrode; reproduced with permission from ref. 93. Copyright 2023, Elsevier. (f) Ex situ XPS spectra of a-V₂O₅@C samples. (g) Two-dimensional contour image of Zn 2p; reproduced with permission from ref. 78. Copyright 2020, WILEY-VCH.

The storage mechanism of Zn²⁺ in biphasic composites can be investigated through the combination of CV test and ex situ measurement methods. For example, Xia et al. prepared ZnMn₂O₄/Mn₂O₃ biphasic nanorods using ZIF-8 as a precursor.⁹³ In Fig. 11c, the CV curves present two pairs of redox peaks (1.58/1.23 and 1.61/1.38 V) at scan rates from 0.2 mV s⁻¹. They correspond to the (de)intercalation of Zn²⁺ and the redox reaction of Mn(II), respectively. From the amplified XRD patterns (Fig. 11d), a series of highly reversible peaks demonstrate the reversibility of lattice expansion and contraction. The peak migration trend in ex situ Raman spectra is consistent with XRD results. As illustrated in Fig. 11e, the Zn–O stretching peak weakens with Zn²⁺ extraction and strengthens with Zn²⁺ insertion.

The Zn//a-V₂O₅@C batteries also possess the energy release/storage mechanism with the Zn²⁺ reaction.⁷⁸ Fig. 11f presents the charge/discharge curves and the corresponding ex situ XPS spectra of O 1s and V 2p. In the O 1s spectrum, the fluctuations of the H₂O peak are synchronized with the (de)insertion of Zn²⁺ due to the solvation effect. In addition, the V 2p spectrum shows a transition towards lower binding energies during the discharge process. Subsequently, they revert to the original state while charging. In Fig. 11g, the Zn signal progressively increases as discharge proceeds, suggesting the steady diffusion into amorphous V₂O₅.

3.2 H⁺/Zn²⁺ co-insertion mechanism

In some MOF derivatives, H⁺ and Zn²⁺ can be co-inserted into the framework of MOFs. This mechanism usually involves complex redox reactions and ion transport processes. The free H⁺ from the aqueous electrolyte plays a crucial role in providing high energy density and cycling stability.^94–96 However, a decrease in the concentration of H⁺ in the electrolyte leads to an increase in pH, which may result in Zn₄(OH)₆SO₄·nH₂O by-products.⁹⁷ The H⁺/Zn²⁺ co-insertion has been recognized as a reliable energy storage behavior in the field of Mn-based MOF-derived materials. The in situ XRD device can monitor the structural changes of cathode in real-time during the cycling process. It includes the expansion and contraction behavior of volume.^98,99 In Fig. 12a, the in situ XRD patterns show several diffraction peaks at 8.1°, which corresponds to the (001) crystal plane of [Zn(OH)₂]₃(ZnSO₄)·(H₂O)₅ (ZHS, JCPDS no. 78-0246).¹⁰⁰ During the discharge process, the signal of ZHS gradually enhances, and the peak intensity disappears in the subsequent charging process. This is consistent with the in situ contour map, indicating that the Zn//δ-MnO₂–C batteries follow the storage behavior of Zn²⁺/H⁺ co-insertion. In addition, the invertible occurrence of laminar ZHS during electrochemical reaction is demonstrated by observing XRD and ex situ SEM (Fig. 12b and c).

Fig. 12 (a) Discharge–charge curves and in situ XRD patterns of δ-MnO₂–C. (b, c) XRD patterns and SEM images at full discharge and full charge status of δ-MnO₂–C; reproduced with permission from ref. 100. Copyright 2022, WILEY-VCH. (d) Capacity–voltage profiles of initial three cycles at 0.05 A g⁻¹ of Mn-BTC-500. (e) Capacity–voltage profiles at different current densities of Mn-BTC-500. SEM images of Mn-BTC-500 cathodes at different voltage points: (f) original state; (g) discharge to plateau; and (h) discharge to 1.0 V; reproduced with permission from ref. 101. Copyright 2021, Elsevier.

The exploration of the energy storage behavior of Mn-BTC-500 (α-Mn₂O₃) reveals an obvious correlation between the charge storage mechanism and the impressed current density.¹⁰¹ As illustrated in Fig. 12d, H⁺ and Zn²⁺ are sequentially inserted into the α-Mn₂O₃ host material at low current densities, and a significant inflection point appears on the discharge voltage curve. However, the intercalation of Zn²⁺ gradually decays at large current densities. This leads to the disappearance of the low-pressure platform after the turning point (Fig. 12e). The results indicate that the H⁺ insertion is the main storage mechanism in the Mn-BTC-500 electrode material at high current densities. Furthermore, the discharge products undergo significant morphological changes during the step-wise insertion of H⁺ and Zn²⁺. From Fig. 12f–h, they transition from flower-like to a uniform layered structure.

The co-insertion mechanism of H⁺ and Zn²⁺ is also available for V-based MOF-derived materials. As shown in Fig. 13a, the ex situ XRD patterns demonstrate that the ZnVO-800 electrodes transform into Zn_xV₂O₅·nH₂O phase due to the in situ electrochemical activation.¹⁰² During the subsequent discharge process, Zn_xV₂O₅·nH₂O is converted into Zn_x+yH_mV₂O₅·nH₂O. Meanwhile, several peaks appear at 8.1/16.2/24.4°, demonstrating the production of a Zn₄SO₄(OH)₆·5H₂O by-product. Fig. 13b presents the energy mechanism of the ZnVO-800 cathode at multiple states of charge, which reveals the route of H⁺/Zn²⁺ co-insertion. Ex situ XPS spectra further prove the above assertion. The Zn 2p peak moves toward a high binding energy (Fig. 13c), which is caused by the incomplete extraction of Zn²⁺. From Fig. 13d, the signal strength of oxygen defects and H–O–H heightens distinctly after the first charge. This phenomenon is due to the successful insertion of H₂O into the host material, thereby enhancing diffusion kinetics.

Fig. 13 (a) Ex situ XRD patterns of ZnVO-800. (b) Schematic illustration of the reaction mechanism of ZnVO-800. (c) Ex situ XPS spectra of Zn 2p and (d) O 1s of ZnVO-800. Reproduced with permission from ref. 102. Copyright 2023, Elsevier.

In addition, Zheng et al. first observed the proton (H⁺) hysteresis phenomenon in electrochemical processes.¹⁰³ As shown in the C 1s XPS spectra of V₂O_3−x-CC, C–O shifts towards high binding energy when discharged to 0.2 V (Fig. 14a). This condition is strongly influenced by electron arrangement and ion diffusion. H⁺ co-intercalates into the electrode and adsorbs on the carbon layer, creating an electric double layer. The hysteresis phenomenon of H⁺ (de)adsorption significantly boosts the pseudo-capacitance of the V₂O_3−x-CC electrode. Subsequently, in situ pH monitoring is used to monitor the changes in electrolyte pH near the positive electrode during the cycling process. In Fig. 14b, the pH curve shows the same trend, indicating that the V₂O_3−x-CC electrode possesses excellent reversibility. Notably, the maximum voltages (H and J) do not match the minimum pH values (C and E). As the system discharges from 1.6 V, the pH value keeps reducing and H⁺ migrates from the cathode to the Zn(CF₃SO₃)₂ aqueous solution. It is speculated that H⁺ moving into the electrolyte is not from V₂O_3−x-CC, but also partly from the layered carbon layer surface through desorption.

Fig. 14 (a) XPS spectra of O 1s at different charging/discharging states. (b) Illustration of pH measurements and evolution of voltage and pH during the initial three charge–discharge cycles at 0.1 A g⁻¹. (c) Evolution of pH versus voltage during the initial three charge–discharge cycles. (d) Adsorption energy of H⁺ on different surfaces. (e) Schematic diagram of ion intercalation/de-intercalation and adsorption/desorption during the discharging/charging process in V₂O₃-CC and V₂O_3−x-CC electrodes. Reproduced with permission from ref. 103. Copyright 2023, Elsevier.

Fig. 14c plots the real-time floating change of proton with voltage. The abnormal drop in pH at the initial discharge process is called the “H⁺ hysteresis effect”. Moreover, the H⁺ adsorption energy on graphene and diamond surfaces is employed to simulate the hybridization modes of amorphous carbon sp² and sp³ in the prepared cathode materials (Fig. 14d). The results reveal that H⁺ possesses a strong adsorption effect on the surface of the amorphous carbon layer during diffusion. It leads to the belated desorption of H⁺ from the cathode to the electrolyte. Finally, the complete electrochemical reaction process is presented in Fig. 14e. The proton hysteresis in the outer carbon layer improves the electrochemical performance.

3.3 Metal ion dissolution/deposition mechanism

In some MOFs, zinc ions may dissolve into the electrolyte during the charging process. These released Zn²⁺ deposits into the pores or surfaces of MOFs during the discharge process, forming metallic zinc or zinc compounds. This mechanism usually involves a reversible redox reaction of zinc ions.^104,105 It is especially evident in Mn-based MOFs. For instance, ZnMn Square acid (ZnMn-SQ) possesses various intricate electrochemical reactions during Zn²⁺ storage processes.¹⁰⁶ The ex situ XRD pattern reveals the evolution in the ZnMn-SQ electrode of the dual-metal MOFs in the first and second cycles (Fig. 15a). After the first charging cycle, Zn²⁺ replaces Mn²⁺ in ZnMn-SQ and is converted into Zn-SQ. The dissociative Mn²⁺ is deposited on the electrode and oxidized to MnO₂. At discharge states of 1.37 and 1.18 V, Zn²⁺ and H⁺ are co-inserted into MnO₂ to generate ZnMn₂O₄ and MnOOH. In the fully discharged state, OH⁻, ZnSO₄, and H₂O react to form Zn₄SO₄(OH)₆·5H₂O (ZHS). In addition, MnOOH and ZnMn₂O₄ are transformed back to MnO₂ in subsequent charging. As shown in Fig. 15b, the whole mechanism can be succinctly described as follows:

ZnMn-SQ + Zn²⁺ → Mn²⁺ + Zn-SQ (1)

Mn²⁺ − 2e⁻ + 2H₂O → MnO₂ + 4H⁺ (2)

MnO₂ + H⁺ + e⁻ → MnOOH (3)

2MnO₂ + Zn²⁺ + e⁻ → ZnMn₂O₄ (4)

Zn²⁺ + 6OH⁻ + ZnSO₄ + 5H₂O → Zn₄SO₄(OH)₆·5H₂O (5)

Fig. 15 (a) Ex situ XRD patterns of ZnMn-SQ samples. (b) Electrochemical mechanism of ZnMn-SQ; reproduced with permission from ref. 106. Copyright 2023, American Chemical Society. (c) Mn 2p XPS spectra of the Mn(BTC) cathode at fully charged state. (d) Mn 3s XPS spectra of the Mn(BTC) cathode at original and fully charged states. (e) FTIR spectra of 1,3,5-H3BTC, original Mn(BTC) cathode, and Mn(BTC) cathode at 1.9 and 1.0 V. Reproduced with permission from ref. 43. Copyright 2020, Springer.

Xu et al. demonstrated the transition of Mn (BTC) to MnO₂ during charging by ex situ XPS and FTIR spectroscopy.⁴³ From Fig. 15c, the Mn 2p spectra at 642.1 eV and 653.8 eV are assigned to the Mn 2p_3/2 and Mn 2p_1/2 orbitals, respectively. As for the Mn 3s spectra, their peak gaps decrease from 6.4 eV to 4.7 eV after they are fully charged (Fig. 15d), indicating an increase in the Mn oxidation state. In the ex situ FTIR spectra, the symmetric stretching vibration of –COO⁻ at 1372 cm⁻¹ moves to 1383 cm⁻¹, corresponding to the Zn(BTC) MOF (Fig. 15e). The results show that Mn (BTC) transforms to Zn (BTC) and Mn²⁺ dissolves into the ZnSO₄ aqueous solution during initial charging. As the charging process progresses, these Mn²⁺ are oxidized to MnO₂ on the cathode surface through manganese deposition reactions. In the subsequent electrochemical reaction, MnO₂ acts as the main energy storage entity of this MOF.

4. Conclusions and perspectives

Various MOF-based cathodes are emerging for AZIBs due to their versatility and high-density active site. This review highlights the recent advances in MOFs and their derived materials as cathode materials for AZIBs. We introduce the relationship between its structure and electrochemical performance. Subsequently, we discuss and summarize the prevalent energy storage mechanisms. Although remarkable progress has been made in their research, there are still plenty of challenges that need to be addressed in this area. For instance: (1) the expensive raw materials and intricate production technology hinder their commercialization. (2) The influence mechanisms are unclear, including the effect of crystal size on Zn²⁺ diffusion and the crystallinity on cycling stability. (3) The pores of derivatives may be blocked owing to the uncontrollable collapse of MOF precursors during high-temperature pyrolysis. Therefore, their internal active sites cannot fully participate in electrochemical reactions. In order to further develop the electrochemical properties of MOFs and their derivatives, we propose the following research directions:

(1) The desired MOFs should possess the potential for low-cost preparation. Titanium-based MOFs have received widespread attention in lithium-ion and lithium–sulfur batteries. However, it has not yet been investigated in the system of AZIBs. Therefore, it is crucial to search for new active MOF species for AZIBs. Porous covalent organic materials (COMs) possess framework structures similar to MOFs, including covalent organic frameworks (COFs), porous polymer networks (PPNs) and conjugated microporous polymers (CMPs). They present a series of advantages such as highly ordered channel structures, customizable structures and multifunctionality. The research on COMs and their related materials can make significant contributions to the development of AZIBs. Moreover, the production efficiency and crystallinity can be improved by optimizing the synthesis process of MOFs. Compared with the traditional hydrothermal method, rapid heating of reactants by microwave radiation can shorten the synthesis time. At the same time, the MOF particles can be obtained with small particle sizes and high purity. This method is suitable for large-scale production.

(2) Advanced in situ characterization techniques can detect the morphology, structure, composition and chemical properties of materials in real time, for instance, in situ electrochemical quartz crystal microbalance (EQCM), in situ X-ray absorption near-edge structure (XANES), and in situ FTIR. They are expected to play important roles in exploring the energy storage mechanism of MOF-based cathode materials. In addition, research on the connection between the synthesis conditions and electrochemical properties of MOF is still at the surface level. The combination of theoretical calculation and empirical investigation is conducive to establishing a quantitative relationship of MOF derivatives. Research efficiency can be promoted through rich data repositories. The high-throughput machine learning algorithms are beneficial to achieve rapid screening of MOF materials, performance prediction and in-depth exploration of the structure–activity relationship.

(3) Most MOFs undergo structural collapse in aqueous environments, resulting in poor stability during battery cycling. To effectively serve as cathode materials, MOFs not only need to show large surface areas and elaborate customization nanostructures but also possess steady structures and strong conductivity. It can lead to the formation of complex structures by introducing two different types of ligands during the synthesis of MOFs. For example, linear chains, network structures, or columnar structures. These structures are connected by strong coordination bonds from layer to layer, which are more stable than traditional hydrogen bonding connections. In addition, the conductivity is crucial for the electrochemical performance of MOF-based cathode. The electrical conductivity of MOFs can be enhanced by the design strategies of “through-bond” and “through-space”. The former relies on the selection of metal–ligand pairs and functions through the overlap of metal–ligand orbitals and covalent bond formation via electron transfer pathways. The latter generates additional charge transfer pathways by non-covalent interactions to promote the conductivity of MOFs. The π–π stacking is a typical way to achieve “through-space” electron transfer.

In conclusion, MOF-based electrode materials have achieved immense progress in the structural design and applications in AZIBs. With the continuous improvement of the synthesis methods of micro–nano MOFs, internal energy storage mechanisms are continuously discovered. We believe that through reasonable structural design and performance control strategies, as well as the integration of composite materials and technology, MOFs and their derivatives present broad prospects and enormous potential in energy storage and conversion applications.

Author contributions

Xiang Wu: writing-review & editing, supervision, resources, project administration, methodology, investigation, funding acquisition, formal analysis, data curation, conceptualization. Yoshio Bando: resources, methodology, investigation, data curation, conceptualization. Yi Liu: collected research papers, writing-original draft, software, resources, methodology, investigation, formal analysis, data curation, conceptualization. All authors participated in manuscript discussion.

Data availability

The data and materials supporting the findings of this study are available from the corresponding authors upon request.

Conflicts of interest

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

Acknowledgements

This project was supported by the Natural Science Foundation of China (no. 52472227).

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