Crystalline porous membrane devices: emerging architectures for carbon-neutral technologies

Abstract

Crystalline porous materials, including metal–organic frameworks (MOFs), covalent organic frameworks (COFs), and hydrogen-bonded organic frameworks (HOFs) are a class of functional materials with periodic extended frameworks, abundant pore structures, designable and adjustable chemical structures. The unique crystalline porous structures facilitate efficient ion transport, provide active sites, and enable molecular separation. Considering these properties, they are regarded as good candidates for fabricating membrane devices in energy and environmental fields, which are closely related to carbon emissions and carbon neutrality. In this review, we summarize the recent progress of crystalline porous membrane devices. Common membrane fabrication methods are systematically summarized, including hot/cold pressing, in situ solvothermal growth, seed-assisted growth, solution processing, interfacial polymerization (IP), and current-driven synthesis. Additionally, diverse applications of crystalline porous membrane devices are presented, including lithium–metal batteries (LMBs), catalytic electrodes, solar cells, and gas/liquid separation membranes. In particular, we discuss the relationship between micro-structures of membranes and the performance of membrane devices and point out the challenges of crystalline porous membrane devices.

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Zifeng Chen

Zifeng Chen received his PhD degree from Tianjin University in 2022 under the supervision of Prof. Yanhou Geng and Prof. Yunhua Xu. Currently, he works with Prof. Xian-He Bu at School of Materials Science and Engineering, Nankai University as a Postdoc. His research interests focus on the design and construction of porous crystalline membranes and applications in gas separation.

Chong Li

Chong Li received his Master's degree in Inorganic Chemistry from Fuzhou University in 2024. He is now a PhD candidate of School of Materials Science and Engineering, Nankai University under the supervision of Prof. Xian-He Bu. His scientific interests include the design and synthesis of crystalline framework membrane materials and applications in gas separation and organic solvent nanofiltration.

Zerong Ge

Zerong Ge received his Master's degree in 2022 from Hebei University of Technology under the supervision of Prof. Hongyan Li. He is now a PhD candidate at Hebei University of Technology, supervised by Prof. Zhiqiang Li, and is jointly trained at Nankai University under the supervision of Prof. Bin Liang. His current research focuses on the applications of porous materials in gas adsorption and separation.

Shicong Wu

Shicong Wu received his Bachelor's degree from Hebei University of Technology in 2023. Currently, he is now a Master's degree candidate of Hebei University of Technology under the joint supervision of Prof. Bin Liang and Prof. Zhiqiang Li. His research interest focuses on the construction of crystalline membrane devices.

Bin Liang

Bin Liang is now a professor at the School of Materials Science and Engineering of Nankai University. He received his PhD degree from the Beijing University of Chemical Technology in 2015 under the supervision of Prof. Bing Cao. He worked with Prof. Zhiyong Tang, Prof. Banglin Chen, and Prof. Shengqian Ma at National Center for Nanoscience and Technology, University of Texas at San Antonio, and University of North Texas, respectively, as a Postdoctoral Fellow from 2015–2024. His current research interest is focused on the development of novel multifunctional membrane materials and their applications in separation, conduction, and sensing.

Xian-He Bu

Xian-He Bu is a professor of Nankai University and a member of the Chinese Academy of Sciences. He obtained his BS and PhD from Nankai University in 1986 and 1992 and has been a full professor since 1995 at the same university. His research interests include functional coordination chemistry, crystal engineering, molecular magnetism, and materials chemistry.

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Wider impact

Over recent decades, crystalline porous materials have revolutionized materials science, offering unparalleled control over porosity, atomic-level design, and versatile functionality. Their incorporation into advanced membrane systems has unlocked groundbreaking opportunities, positioning these emerging architectures at the forefront of next-generation technologies for clean energy, ecological preservation, and precision separations. This review, informed by decades of pioneering research in porous frameworks, delves into cutting-edge developments in membrane-based applications, with a spotlight on breakthroughs in energy storage, renewable conversion systems, and eco-conscious remediation tools. By synthesizing design principles and forward-looking perspectives, this work aims to catalyze cross-disciplinary collaboration, equipping researchers and engineers with actionable strategies to accelerate the evolution of sustainable technologies, and thereby appealing to a broad interdisciplinary readership of Materials Horizons.

1. Introduction

Persistently high carbon emissions have contributed to extreme weather events, rising sea levels, and biodiversity loss. Consequently, to ensure the sustainable development of human society, mitigating carbon emissions has become a global imperative. The international community has established coordinated principles and targets for emissions reduction, among which carbon neutrality is emerging as an important goal in the middle of this century.¹ To achieve this goal, research efforts are primarily focused on four key areas. One is the renewable energy conversion technologies, such as solar cells, which can efficiently convert solar energy into electric energy,^2,3 reducing the carbon emissions accompanied by fossil fuel power generation. The second research focus is advanced electrochemical energy storage systems,^4–6 such as rechargeable lithium–metal batteries (LMBs). The high energy density and long-term cycling stability of LMBs will reduce the dependence on fossil fuels. The third is the catalytic processes for sustainable chemical production, particularly through the hydrogen evolution reaction (HER) and carbon dioxide reduction reaction (CO₂RR), which play important roles in producing green hydrogen energy and reducing CO₂ emissions, respectively.^7–10 The last aspect is low energy consumption separation technologies, which are considered as environmentally friendly technologies with low carbon footprint and energy consumption,^11–15 including gas and liquid separation.^16–20

Membrane-based technologies represent promising solutions that exhibit substantial advantages in addressing carbon emissions.^21–25 Researches have demonstrated that the membrane devices hold exceptional potential for sustainable development. Thus far, polymers have been widely adopted to fabricate functional membrane devices due to their easy-processability.^26–29 However, these polymer membranes deliver certain disadvantages in various devices. In LMBs, polymer membranes are often used as solid-state electrolytes to conduct ions. Most polymer-based electrolytes show relatively low ionic conductivity at room temperature, which severely hinder the practical application of solid-state batteries.^30–33 In the gas separation field, polymer membranes with micropores usually show uneven pore sizes, which are harmful to the precise sieving of gas molecules.^34–36 So it is necessary to develop alternative materials to construct membranes for the purpose of reducing carbon emissions.

Advanced membrane devices based on crystalline porous materials demonstrate superior performance over conventional polymer membranes, owing to their tailorable pore architectures, tunable chemical functionalities, and robust stability.^37–41 These materials have been extensively utilized in critical applications such as energy storage and conversion,^42,43 gas and liquid separation,^44,45 and electrocatalysis.⁴⁶ Crystalline porous materials are primarily classified into three structurally distinct categories: metal–organic frameworks (MOFs), covalent organic frameworks (COFs), and hydrogen-bonded organic frameworks (HOFs). MOFs are a type of inorganic–organic material with crystalline porous structure.^47–59 In secondary LMBs, the existence of open metal sites (OMSs) on the surface of MOFs can bond with anions of lithium salts to liberate Li⁺, which can enhance the transference number of Li⁺ and prolong the cycling lifetime of batteries.^60,61 When adding MOFs into polymer electrolyte matrix as filler, the crystallinity of polymer chains can be efficiently blocked, thus significantly enhancing the ionic conductivity of solid-state electrolyte.^62,63 Meanwhile, the abundant pore channels and large surface area are beneficial to the molecules/ions diffusion, and the functional groups of organic ligands (such as –NH₂ and –F) can interact with specific gas, thus facilitating the efficient separation of mixed gas.^64,65 COFs are another new class of crystalline materials with covalent bonded structure networks and ordered pore channels.^66–74 Decorating COF pore walls with oligomers can enhance the ionic conductivity by facilitating the dissociation of lithium salts.^75,76 The ordered pore channels also provide sufficient pathways for ion transport. In separation field, through adjusting packing mode of COF layers or introducing missing-linker defect can act as continuous facilitated gas transport carriers.^77,78 In comparison of MOFs and COFs, HOFs represent a distinct category with building units linked by hydrogen bonds.^79–83 The weak and reversible hydrogen bonding endows HOFs with numerous advantages.^84–86 For example, the synthesis conditions of HOFs are always mild and HOFs are easy solution processed. Moreover, the hydrogen bonds make HOFs inherit self-healing capability, which is beneficial to the stability of gas separation.

Currently, although several reviews have focused on specific applications of crystalline porous membrane devices,^87–93 there remains a scarcity of comprehensive and systematic reviews summarizing their implementations in energy and environmental fields towards reducing carbon emissions. In this article, we provide a concise overview of recent advancements in crystalline porous membrane devices in energy and environmental fields. In Section 2, we introduce the common methods of fabricating crystalline porous membranes, including hot/cold pressing, in situ solvothermal growth, seed-assisted secondary growth, solution processing, interfacial polymerization, and current-driven synthesis. Then we concentrate on the utilization and development of crystalline porous membranes (Section 3). In this section, we introduce several crystalline porous membrane devices, including LMBs, catalytic electrodes, solar cells, and separation membranes (Fig. 1). Finally, Section 4 summarizes the current research status of crystalline porous devices, pointing out the challenges and potential avenues for further development.

Fig. 1 Schematic of applications of MOF, COF, and HOF membrane devices.

2. Representative crystalline porous membranes fabrication strategies

2.1. Design principles

To date, various crystalline porous materials have been developed which possess functions in energy and separation fields. However, to fully take advantage of these functional crystalline porous materials in practical devices, it is necessary to make them into membranes through rational fabrication strategy. The selected process method should maintain the original structure and performance of materials to the greatest extent. For example, COF materials have advantages of ordered and closely arranged pore structures, so it is necessary to adopt appropriate membrane fabrication strategies to achieve the full potential of crystalline porous materials. Meanwhile, searching for an effective approach to eliminate defects in polycrystalline membranes is essential to enhance separation performance. In addition, improving the adhesion/bonding between the selective layer and substrate to ensure enough mechanical stability is of great importance. In this section, we will discuss common approaches in detail to fabricate high performance crystalline porous membranes (Fig. 2). Meanwhile, the summary of advantages, disadvantages, and applicable conditions of each method are listed in Table 1.

Fig. 2 Typical membrane fabrication methods. (a) Cold/hot pressing. Reproduced with permission.⁶³ Copyright 2019, Elsevier. (b) In situ solvothermal growth of ZIF-90 membranes. Reproduced with permission.⁹⁷ Copyright 2010, American Chemical Society. (c) Formation of ZIF-8@COF nanosheet seeds. Reproduced with permission.⁹⁹ Copyright 2023, Wiley-VCH. (d) Solution-processing of UPC-HOF-6 membranes. Reproduced with permission.¹⁰¹ Copyright 2020, Wiley-VCH. (e) Interfacial polymerization of COF membranes Tp-Bpy. Reproduced with permission.¹⁰² Copyright 2017, American Chemical Society. (f) Current-driven synthesis of ZIF-8 membranes. Reproduced with permission.¹⁰⁶ Copyright 2018, Science.

Table 1 Summary of advantages, disadvantages, and applicable conditions of various membrane fabrication methods

Methods	Advantages	Disadvantages	Applicable conditions
Hot/cold pressing	Simple operation; no solvents	Inhomogeneous morphology; limited scalability for industrial production	Solid-state electrolyte membranes
In situ solvothermal growth	High crystallinity and continuous membranes	Difficulty in controlling crystal alignment	Gas separation membranes
Seed-assisted secondary growth	Ultra-thin membranes; eliminating defects	Regulating secondary growth conditions	High-permeance gas separation membranes
Solution processing	Ideal for flexible membranes; controlling thickness	Residual solvent defects	Gas/liquid separation membranes
Interfacial polymerization (IP)	Large-area fabrication	Moderate crystallinity	Catalytic electrode membranes; gas/liquid separation membranes
Current-driven synthesis	Rapid room-temperature synthesis; self-healing defects	Additional conductive substrates	High-selectivity separation membranes

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2.2. Hot/cold pressing

Hot pressing is a thermally assisted process mainly employed in the fabrication of binder-containing ceramic solid-state electrolytes (SSEs) and polymer-based SSEs. For ceramic SSEs, sulfide SSEs were explored due to the high theoretical ionic conductivity (∼1 × 10⁻² S cm⁻¹). In 2024, Kozen et al. prepared argyrodite SSE Li₆PS₅Cl (LPSCl) using hot pressing method.⁶ Certain amounts of LPSCl powder was placed in a temperature-controlled die set and various pressures were applied through a hydraulic press. The mold assembly was heated to a predetermined temperature, after which the heating program was terminated to allow for natural cooling. The applied pressure was sustained until the mold temperature decreased below 60 °C. The obtained LPSCl SSE membranes showed high ionic conductivity, which can be attributed to the synergistic effects of diminished grain boundary resistance and improved densification induced by the hot pressing treatment.

Polymer SSE membranes have been studied for many years due to the easy-processability of polymers, such as poly(ethylene oxide) (PEO). The applied hot pressing method enables the melting of polymers while eliminating the use of solvents. However, the inherent low ionic conductivity of polymers severely hinder the practical applications in LMBs. To enhance the ionic conductivity, additive fillers were adopted to fabricate polymer-based composite electrolyte membranes. The conventional fabrication protocol typically comprises an initial ball-milling and homogenization step for the uniform dispersion of polymeric matrices, lithium salts, and functional additives within the composites. Subsequently, the composite undergoes thermal extrusion processing under elevated temperature and pressure conditions. The fabricated composite electrolyte membranes by hot pressing can ensure the excellent interfacial contact between composite membranes and electrodes, which is beneficial to the reduction of impedance and uniform stripping and plating of Li⁺. Common additive fillers include lithium lanthanum zirconium oxide (Li_6.4La₃Zr_1.4Ta_0.6O₁₂, LLZTO) and MOFs. In 2018, Goodenough et al. introduced LLZTO into PEO matrix to prepare PEO–LLZTO electrolyte membranes by hot pressing method.⁹⁴ Precisely measured quantities of LLZTO powder, PEO, and lithium bis(trifluoromethanesulfonyl)imide (LiTFSI) were dry-grounded in a mortar to form a preliminary membrane. Subsequent hot-pressing at 70 °C under 10 MPa for 3 min yielded a homogeneous composite electrolyte membrane with an approximate thickness of 70 μm. The obtained membranes delivered high ionic conductivity of 10⁻⁴ S cm⁻¹ at 55 °C, and the assembled Li|LiFePO₄ (LFP) batteries showed excellent cycling stability. The incorporation of MOFs as functional fillers in solid electrolyte membranes demonstrated enhanced ionic conductivity and optimized interfacial characteristics. In 2019, Sun's group prepared PEO and cationic MOF (CMOF) composite electrolyte (denoted as P@CMOF) by hot pressing method.⁶⁰ First, the calculated volumes of PEO, LiTFSI, and CMOF were homogeneously mixed and further grounded in a mortar to obtain a rough ball. Subsequently, the rough ball was sandwiched between two pieces of polytetrafluoroethylene (PTFE) plates and subjected to compressive pressure of 20 MPa for 2 hours at 100 °C (Fig. 2a). A membrane with a uniform thickness of about 40 μm was successfully prepared. Assembled Li|LiFePO₄ and Li|LiFe_0.15Mn_0.85PO₄ batteries both exhibited superior rate and cycling performances at 60 °C. Stephan et al. once synthesized micro-particles with amine-functionalized zirconium-based MOF (UiO-66-NH₂) anchored on silica (SiO₂) (denoted as UiO-66-NH₂@SiO₂).⁶¹ Then UiO-66-NH₂@SiO₂ was used as additive filler to prepare PEO-based composite electrolyte membranes. Precise quantities of PEO and LiTFSI were first dissolved in anhydrous acetonitrile under continuous magnetic stirring for 8 hours. Subsequently, the nanofiller UiO-66-NH₂@SiO₂ was uniformly dispersed in the polymer solution, which was then solution-cast into a film and subjected to hot-pressing. Membranes with an average thickness of 50–60 μm can be obtained. The MOF-containing electrolyte membranes demonstrated good electrochemical performance due to the better contact between electrodes and electrolytes and the additional ion pathway in the electrolyte membranes.

In comparison with hot pressing, cold pressing at room temperature has advantages of mild preparation conditions and low energy consumption. Guo's group previously combined UiO-66 and Li-containing ionic liquid (Li-IL) to obtain a high ionic conductivity electrolyte.⁹⁵ The synthesized UiO-66 was mixed with various contents of Li-IL using mortar grinding. The resulting UiO/Li-IL composite powders (0.15–0.2 g) were placed into a PTFE mold (15.5 mm inner diameter) and compressed under 3 T pressure for 2 min using stainless steel dies to fabricate solid electrolyte pellets. For higher Li-IL loading UiO/Li-IL composite, the viscous composite was processed into thick films by pressing and rolling on a PTFE substrate, which were then punched into 15 mm diameter disks. The ionic conductivity of the prepared electrolyte membranes was 3.2 × 10⁻⁴ S cm⁻¹ at ambient temperature. Low resistance of electrodes/electrolyte interfaces can be obtained, and the fabricated batteries showed 94% capacity retention after 380 cycles at 1C.

Hot/cold pressing holds advantages of simple operation as well as no solvents. This method is particularly suitable for fabricating compact thin membranes, which are commonly employed as electrolyte membranes in solid-state batteries to enhance Li⁺ conductivity while reducing volumetric energy density. However, constrained by the fabrication process, this method has limited scalability for industrial production.

2.3. In situ solvothermal growth

In situ solvothermal growth is a widely applied method to fabricate polycrystalline porous membranes including MOF and COF membranes. In 2009, Lai's group first reported MOF-5 membrane by in situ solvothermal method on a porous alumina support.⁹⁶ The obtained MOF-5 polycrystalline membrane showed continuous morphology. Subsequently various MOF membranes were prepared by this approach. However, the challenges of direct growing crack-free membranes on substrates remained due to insufficient nucleation sites. To enhance the interactions between MOF and substrate, Caro's group developed a novel method of metal-targeted modification of the substrate⁹⁷ (Fig. 2b). By using 3-aminopropyltriethoxysilane (APTES) as an organic covalent linker as well as metal-targeted modifier, the ethoxy groups of APTES reacted with hydroxyl groups on the alumina surface. Then the amino end groups of APTES reacted with aldehyde groups of ligands of ZIF-90 via imine condensation to form a ligand layer, which promoted nucleation and crystallization of ZIF-90. Consequently, continuous ZIF-90 molecular sieve membranes with high H₂ separation performance were obtained. In addition, Caro and coworkers took advantage of the same method to prepare continuous and high-quality COF–LZU1 membranes on commercial ceramic tubes.⁹⁸ APTES was used to modify the tube surface and then reacted with 1,3,5-triformylbenzene (TFB) for 1 hour. Vertically placing the modified alumina tube into an autoclave with another monomer p-phenylenediamine (PDA) and reactant solvent, well-intergrown tubular COF membranes with good dye removal ability from water were obtained. Furthermore, adopting the same strategy, COF–LZU1–ACOF-1 bilayer membranes for H₂ separation and COF–MOF composite membranes were prepared and studied.

In situ solvothermal growth fabrication methods are usually adopted to fabricate high crystallinity and continuous membranes. Several gas separation membranes were prepared by this method. However, the membranes fabricated via this approach typically exhibit anisotropic properties, due to the difficulty in controlling crystal alignment during in situ growth.

2.4. Seed-assisted secondary growth

Although much progress has been made in membranes fabricated by in situ solvothermal method, the thickness of obtained membranes was larger than 1 μm which can be attributed to the lack of nucleation sites. Therefore, the thick membranes prepared by in situ solvothermal way often exhibit limited gas/liquid permeance. To tackle this issue, seed-assisted secondary growth was developed to fabricate high flux membranes. In 2023, Jiang's group designed a series of negatively charged COF nanosheets as porous substrates, which preferentially adsorbed Zn²⁺ (Fig. 2c).⁹⁹ Then the deprotonated 2-methylimidazole (2-MIM) ligands coordinated with the enriched Zn²⁺ on the COF surface and underwent heterogeneous nucleation. ZIF-8 crystals subsequently grew and ZIF-8@COF nanosheet seeds formed. The resulting ZIF-8 membranes exhibited an ultra-high propylene/propane (C₃H₆/C₃H₈) separation performance and superior stability with a thickness of only 100 nm. To obtain high H₂ separation membranes, Liu's group once prepared MIP-177-LT MOF membranes with uniform alignment of pore channels.¹⁰⁰ They fabricated b-oriented MIP-177-LT membranes through oriented epitaxial growth. First, they synthesized rod-shaped MIP-177-LT crystals with an aspect ratio of around 5.8 as seeds through adjusting the chemical composition of the precursor solution. Next, by spin-coating the aqueous suspension containing proper seed concentration and trace amount of surfactant polyvinylpyrrolidone (PVP), a closely packed b-oriented MIP-177-LT seed layer was constructed. Well-intergrown MIP-177-LT membranes with a thickness of 1.3 μm could be obtained via the optimization of the chemical composition of the precursor solution. Through the preferred orientation control of MOF membranes, an ultra-high H₂ selectivity was obtained, even surpassing the upper bound limit. Such superior performance can be attributed to the 0.3 nm-sized window aligned vertically to the substrate. In comparison, when choosing randomly oriented seed layer, c-oriented or randomly oriented membranes were obtained with low separation factors (SFs) of H₂/N₂ and H₂/CH₄.

Ultra-thin membranes usually adopt this fabrication process in order to prepare high-permeance membranes, especially in the field of gas separation. The secondary growth process can effectively eliminate the defects caused by the polycrystals. However, the formation of the seed layer (e.g., crystal seed coating, orientation control) requires precise regulation, increasing process complexity. Moreover, secondary growth conditions (e.g., temperature, duration, precursor concentration) significantly influence membrane quality, making optimization challenging. Furthermore, the compatibility between seed layer and substrate has to be deliberately considered.

2.5. Solution processing

Solution processing method is suitable for crystalline porous materials with good dispersion as well as stability in solvents. HOFs are a class of important crystalline materials which are self-assembled by hydrogen bonds. Some polar groups (–NH₂ or –COOH) without participating in bonding remained on the surface of the crystals, which can interact with polar solvents (H₂O, dimethyl sulfoxide (DMSO), or N,N-dimethylformamide (DMF)) and facilitate the homogeneous dispersion of HOFs in solvents. This property promotes the application of solution processing HOFs. In 2019, Sun's group reported the first example of HOF membranes for gas separation by an optimized solution processing¹⁰¹ (Fig. 2d). Triangular block 4′,4′′,4′′′-nitrilotris(([1,1′-biphenyl]-4-diaminotriazine)) (NBP-DAT) was used to construct flexible and stable HOF (UPC-HOF-6). The synthesized UPC-HOF-6 was first dissolved in DMSO with various concentrations, then the solution was dipped onto the anodic alumina oxide (AAO) substrate. At different temperatures, the solvent was evaporated and corresponding HOF membranes were obtained. Through adjusting the concentration of precursors as well as the evaporating temperature, a continuous UPC-HOF-6 membrane could be obtained with a thickness of 1 μm and showing excellent H₂ separation performance. Jiang and coworkers synthesized HOF nanosheets using 1,2,4,5-tetrakis (4-carboxyphenyl)-benzene (H₄TCPB) as the molecule motif.⁴⁰ In order to prepare HOF membranes, H₄TCPB nanosheet colloidal suspension was filtered under vacuum on a Nylon support. Subsequently, the membrane was dried and carefully delaminated from the Nylon support to obtain free-standing HOF membranes.

Solution processing is an ideal process for fabricating flexible self-standing membranes, and the thickness of membranes can be controlled by evaporation of solvents. However, residual solvent defects may be harmful to the performance of membrane devices.

2.6. Interfacial polymerization (IP)

IP is a facile method to construct large-scale thin crystalline porous membranes. Banerjee's group developed a bottom-up interfacial crystallization method to fabricate COF membranes¹⁰² (Fig. 2e). They dissolved an aldehyde organic linker in dichloromethane while the amine monomer in water. Polycondensation occurred at the liquid–liquid interface and COF crystalline membranes formed. This method can simultaneously control the crystallization and morphology of membranes. The obtained COF membranes Tp-Bpy delivered splendid acetonitrile permeance of 339 L m⁻² h⁻¹ bar⁻¹. Xu and coworkers presented a novel diffusion/modulator dual-mediated solid–liquid/vapor interfacial synthesis strategy.¹⁰³ They used the inner wall of the reaction vessel as a robust substrate to construct COF membranes. Through elaborate tuning temperature and solvent, a supersaturated liquid vapor layer can form on the vessel wall on which nuclei are cultivated. A sustained solid–liquid/vapor interfacial synthesis occurred which was driven by the diffusion of monomers in solution as well as vapor phases. Through an aniline-mediated three-interface exchange reaction, highly ordered crystalline imine-linked COF membranes were obtained. Although the IP method for constructing COF membranes has received much progress, the preparation of MOF membranes using this method still focuses on traditional materials such as ZIF-8.¹⁰⁴ Fabricating high-valent cluster-based MOF membranes remains a challenge due to the high energy barrier and sophisticated reaction mechanism. Recently, Sun’ s group first proposed preprocessed monomer IP approach (PMIP) to fabricate a series of high-valent cluster-based MOF membranes.¹⁰⁵ The pre-constructed metal clusters avoided the transformation from metal ions to clusters, which could reduce the energy barrier during membrane formation. Meanwhile, they adopted roll-to-roll techniques with low raw materials and time costs to prepare Zr-fum MOF membranes, which showed high H₂/CH₄ SF and stable separations.

Interfacial polymerization is well-suited for fabricating large-area flexible membranes and exhibits good compatibility with flexible substrates. However, it typically yields membranes with limited crystallinity, thereby restricting their molecular sieving precision. Catalytic electrode membranes, gas separation, and liquid separation membranes were often fabricated by this method.

2.7. Current-driven synthesis

Current-driven synthesis emerges as a novel approach which can quickly construct defect-free ultra-thin membranes. Upon the inner electric field between two electrodes, metal ions can move toward the cathode and accumulate on the support surface. Deprotonated ligands also form under the effect of an electric field, which can react immediately with metal ions due to the negative charges. In 2018, Caro and coworkers reported the current-driven synthesis of ZIF-8 membranes for the first time¹⁰⁶ (Fig. 2f). Through this method, they successfully fabricated defect-free and ultra-thin ZIF-8 membranes during the prolonged reaction time. The high quality of ZIF-8 membranes can be attributed to the defects self-elimination characteristics of current-driven methods: any defect areas will expose the conductive substrate, where further formation of ZIF-8 crystals will eliminate the defects. Eddaoudi's group introduced electrochemistry to control the ligand deprotonation and exchange process, obtaining continuous and defect-free fcu-MOF polycrystalline membranes.¹⁰⁷ The membrane growth process utilized pre-formed hexanuclear clusters with steadily supplied deprotonated linkers, which played a crucial role in constructing high-quality membranes. Later, they used the same method to construct ultra-thin membranes with mixed linkers for nitrogen removal from raw natural gas.¹⁰⁸

Current-driven synthesis holds advantages of rapid room-temperature synthesis and self-healing process. High-selectivity membranes can be prepared by this method. However, additional conductive substrates were required, which will add to the costs of membrane fabrication.

3. Typical crystalline porous membranes devices

3.1. Lithium metal batteries (LMBs)

LMBs utilizing metallic lithium (Li) anodes have been recognized as efficient energy storage solutions.^109,110 Among the reported anodes, Li has been extensively studied due to its exceptional theoretical specific capacity (3860 mA h g⁻¹) and extremely low redox potential (−3.04 V vs. the standard hydrogen electrode).¹¹¹ However, the application of lithium anodes in LMBs faces severe challenges due to the safety issue, which is caused by the flammability of organic liquid electrolytes.^112,113 To alleviate the potential safety hazard, solid-state batteries were developed rapidly in which solid-state electrolytes played a key role in determining the whole electrochemical performance of batteries.

3.1.1. Recent progress of MOFs-based solid-state electrolytes (SSEs) for LMBs.
3.1.1.1. Neat MOFs. Alongside the progress in solid-state electrolyte development, MOFs as new types of porous materials have attracted more and more attention due to the merits summarized below:^114,115 (i) high specific surface which promotes ion transmission; (ii) highly adjustable pore structures such as pore size and pore polarity that can provide unique transport channels for ions; (iii) OMSs can anchor anions of lithium salts through coordination and enhance the transference number of electrolytes.

In 2011, Long's group first reported a MOF-based SSE material, which combined lithium isopropoxide (LioⁱPr) with Mg₂(dobdc) (dobdc⁴⁻ = 1,4-dioxido-2,5-benzenedicarboxylate) followed by soaking in a typical electrolyte ethylene carbonate (EC)/diethylcarbonate (DEC) (1 : 1 vol%).¹¹⁶ The framework of Mg₂(dobdc) features one-dimensional (1D) hexagonal channels with a diameter of approximately 14 Å. The lithium ions can move along the channels freely while the alkoxide anions might bind the Mg²⁺ of the framework. The obtained SSE delivered a conductivity of 3.1 × 10⁻⁴ S cm⁻¹ at 27 °C. Since then, other MOFs have been studied as SSEs by incorporating various lithium salts, such as lithium halides (lithium chloride, lithium bromide, lithium iodide),^117,118 lithium tertbutoxide,⁶⁰ and lithium perchlorate (LiClO₄).^119–121 To further enhance the conductivity of SSEs, the pathway for transport of ions should be explored. Jiang et al. constructed a liquid–electrolyte–laden MOF (HKUST-1) for lithium batteries.¹²² The hierarchical pore architecture in this system encompasses both macroscopic and mesoscopic structural regimes. An extraordinarily high ionic conductivity of 1.02 mS cm⁻¹ was observed, which can be attributed to the abundant interstices and cracks in the MOF electrolyte (Fig. 3a). The interstices and cracks can accommodate liquid electrolyte and provide extra channels for ion transport together with MOF inner channels. The fabricated Li|LFP batteries showed excellent cycling stability with a capacity of 93% after 210 cycles at 1C.

Fig. 3 Representative MOF electrolyte membranes. (a) Illustration of Li⁺ transport mechanism of HKUST-1 electrolyte. Reproduced with permission.¹²² Copyright 2023, Wiley-VCH. (b) Design strategy of MOF electrolyte with bilayer zwitterionic nanochannels. Reproduced with permission.¹²³ Copyright 2023, Wiley-VCH. (c) Crystal structures and pore apertures of MOF-74, HKUST-1, and MOF-5. Reproduced with permission.¹²⁷ Copyright 2023, Wiley-VCH. (d) Orbitals and orbital interactions between Zr 3d of NH₂-UiO-66 and NO₂-UiO-66. Reproduced with permission.¹³² Copyright 2024, American Chemical Society. (e) Process of preparing 2D amino-functionalized MOF nanosheet and PEO composite electrolyte membrane. Reproduced with permission.¹⁴² Copyright 2023, Wiley-VCH. (f) Schematic of the ion transport in composite electrolyte. Reproduced with permission.¹⁴⁶ Copyright 2024, Royal Society of Chemistry.

Although MOFs incorporating with electrolyte solvents show high conductivity, the flammable properties of liquid bring safety issues of batteries.¹¹⁵ The low intrinsic conductivity of MOFs limits their application in all-solid-state electrolytes. Huang et al. developed a MOF with bilayer zwitterionic nanochannels (MOF-BZNs) that function as efficient solid-state electrolytes.¹²³ Soft multicationic oligomers (MCOs) were grafted onto the channels of a rigid anionic MOF framework (Fig. 3b). Lithium salt dissociation was facilitated by competing cationic MCOs and anionic sites in MOF-BZNs. The obtained MOF-BZN SSE exhibited high ionic conductivity of 8.76 × 10⁻⁴ S cm⁻¹ and Li⁺ transference number of 0.75. Solid-state full cells were fabricated to demonstrate the applicability of MOF-BZNs. Remarkably, the LFP|SSE|Li full cells maintained stable operation at 30 °C despite the challenging conditions of high LiFePO₄ loading (17.0 mg cm⁻²) and low N/P ratio (3.8), while delivering an impressive energy density of 259.2 W h kg⁻¹. Meanwhile, the assembled NCM-811|SSE|Li can work steadily with an energy density of 419.6 W h kg⁻¹, even under high NCM-811 loading (20.1 mg cm⁻²) and low N/P ratio of 2.4.

In order to rationally design MOFs structures to obtain high performance lithium batteries, the relationship between structures and performances must be systematically investigated. Although several studies have been reported before,^124–126 it remains ambiguous as to the relationship between MOF structures and ionic conductivity and electrochemical stability window. Also, the influence of OMSs on the interactions between guest electrolytes should be further studied. In 2023, Song's group synthesized a series of nanocrystalline MOFs (MOF-74, HKUST-1, and MOF-5) with various topological structures¹²⁷ (Fig. 3c). The researchers systematically investigated the influence of pore aperture and OMSs on both ionic transport characteristics and electrochemical stability. The results demonstrated that MOF-based electrolytes containing non-redox-active metal centers exhibited significantly broader electrochemical stability windows compared to their redox-active counterparts. Furthermore, larger pore apertures enable high ionic conductivity due to the incorporation of more lithium salts. Therein, Zn–MOF-74 electrolyte with pore apertures of 10 Å delivered a high conductivity of 1.73 × 10⁻⁴ S cm⁻¹ at ambient temperature. Additionally, MOFs with OMSs can achieve high Li⁺ mobility, which can be attributed to the effective dissociation of lithium salts and immobilization of anions. In conclusion, the obtained Zn–MOF-74 electrolyte demonstrated splendid electrochemical performance with LFP cathode, with good cycling stability exceeding 500 cycles at 0.5C and high voltage LiCoO₂ cathodes (up to 4.6 V vs. Li⁺/Li) at 30 °C. Designing MOFs at the molecular level plays an important role in enhancing the ionic conductivity, such as introducing halogen elements,^128,129 implementing grafting techniques,^130,131 and enabling reversible transitions between neutral and anionic phases in MOFs.¹¹⁸ The understanding of relationships between electronic structures and performances at the orbital level is crucial to rationally redesign MOF structures. Zhao's group used UiO-66 as a platform to study the structure–performance relationship.¹³² They introduced –NH₂ (electron-donating group) and –NO₂ (electron-withdrawing group) to UiO-66, in order to precisely tune the electronic structure of Zr central metal. The strong electron-withdrawing property of –NO₂ caused an increase in the energy splitting of Zr 3d_xz, 3d_yz, and 3d_z² orbitals, which can be attributed to the high valence state on the Zr 3d orbitals (Fig. 3d). As a consequence, the interactions between high energy Zr orbitals and anions were enhanced, resulting in the improvement of Li⁺ conductivity (2.72 × 10⁻⁴ S cm⁻¹, higher than that of original UiO-66-Li of 3.40 × 10⁻⁵ S cm⁻¹ and NH₂-UiO-66-Li of 2.14 × 10⁻⁵ S cm⁻¹). The assembled Li|LFP batteries adopting NO₂-UiO-66-Li electrolyte delivered excellent rate performance. At current densities of 0.1, 0.2, 0.5, and 2.0C, specific capacities of 163, 161, 146, and 99 mA h g⁻¹ can be obtained, respectively. When the current density comes back to 0.5C, the discharge specific capacity can be stable at about 155 mA h g⁻¹. In contrast, batteries based on NH₂-UiO-66-Li only exhibited a capacity of 112 mA h g⁻¹ at 0.5C, due to the low ionic conductivity. Moreover, NO₂-UiO-66-Li SSE also delivered a great cycling performance of 93.4% capacity retention after 150 cycles.

The OMSs and pore structures are two key factors for the electrochemical performance of LMBs. Through elaborately tuning pore size and the interactions between OMSs and guest electrolytes of MOF electrolyte membranes, high ionic conductivity and uniform Li stripping can be obtained, thus enhancing cycling stability and rate performance.

3.1.1.2. MOFs/polymers. Polymer electrolytes possess the advantages of light weight, superior processability, good flexibility, and chemical compatibility with electrodes, rendering them optimal choices for ionic conductive materials.^133–136 However, the inherent low conductivity (<10⁻⁴ S cm⁻¹) at room temperature and poor mechanical properties seriously hinder them from industrial applications. The crystalline regions of polymers always block chain segment motion, which further obstruct ion transport. In recent years, MOFs as fillers in polymer electrolytes have attracted the attention of researchers.^63,137,138 The incorporation of MOFs can reduce the crystallinity of polymers, and the OMSs of MOFs can effectively immobilize anions, thus facilitating the dissociation of lithium salts and enhancing the ion conductivity. Moreover, the unique tunable pore structures of MOFs can provide additional pathways for ions, thereby enabling the formation of homogeneous ion flux, thus promoting the uniform deposition of lithium ions.

Wang et al. employed nano-sized ZIF-8 as inorganic filler into PEO to prepare ZIF-8/PEO composite electrolyte.¹³⁹ The crystallinity of PEO decreased due to the introduction of ZIF-8, which inhibited the reorganization of PEO chain segments, thus an ionic conductivity of 2.2 × 10⁻⁵ S cm⁻¹ was obtained, one order of magnitude higher than that of pure PEO (3.6 × 10⁻⁶ S cm⁻¹) at 30 °C. Besides, the Lewis acid surface of ZIF-8 facilitated the dissociation of lithium salts through binding the anions, thus enhancing the Li⁺ transference number and mitigating the concentration polarization during charging/discharging. As a result, the fabricated Li|PEO/ZIF-8|LFP batteries delivered excellent cycling performance at 60 °C at 0.5C. Kim et al. simultaneously introduced ZIF-67 and ionic liquid electrolyte (ILE) as a filler and plasticizer into PEO electrolyte.¹⁴⁰ By optimizing the preparation technology, an ultra-thin composite electrolyte membrane with a thickness of ≈32 μm was obtained. The membrane was beneficial to enhancing energy density and facilitating the transportation of Li⁺. The introduction of ZIF-67 particles and ILE enhanced immobilization of anions and shortened the Li⁺ pathways. Consequently, the high ionic conductivity can be obtained as 1.19 × 10⁻⁴ S cm⁻¹ at room temperature and high Li⁺ transference number of 0.8. Assembling with LFP cathodes, the batteries delivered an initial specific discharge capacity of 166.4 mA h g⁻¹ and excellent cycling stability of 83.7% capacity retention after 1000 cycles at 3C under 60 °C.

Two-dimensional (2D) MOFs were used as solid-state electrolyte fillers due to the large specific area and abundant exposed OMSs. Zou's group explored 2D-MOF Cu-BTC as fillers to prepare novel Cu-BTC/PEO composite electrolyte CPE-5.¹⁴¹ Owing to the interactions between unsaturated sites in Cu-BTC and oxygen atoms in PEO chains, the mechanical properties of the composite electrolytes were improved. Meanwhile, the anions of lithium salts TFSI⁻ were anchored by electrostatic forces, thus suppressing the lithium dendrites in batteries. Interestingly, the introduced 2D-MOF provided multiple ion transport pathways. Li⁺ can jump along the benzene center of Cu-BTC based on the density functional theory (DFT) calculation results, and the interface between Cu-BTC and PEO chains can also provide a rapid Li⁺ transport channel. As a result, the prepared composite electrolyte showed a high ionic conductivity of 4.6 × 10⁻⁵ S cm⁻¹, and the fabricated Li|CPE-5|LFP batteries delivered good stability after 500 cycles at 0.5C under 60 °C. They also constructed a composite solid-state electrolyte using amino-functionalized 2D MOF nanosheets and PEO¹⁴² (Fig. 3e). 2D MOF nanosheets can provide fast lithium-ion transport interfaces. Meanwhile, the electron-donor amino functional groups (–NH₂) can act as ionic sieves to enhance Li⁺ transference number from 0.36 to 0.64. High ionic conductivity of 6.5 × 10⁻⁵ S cm⁻¹ was obtained under ambient conditions. As a result, the fabricated Li|LiFePO₄ batteries delivered highly reversible specific capacities of 148.8 mA h g⁻¹ after 200 cycles.

Although various MOFs as fillers and plasticizers were introduced into polymer electrolytes to enhance ionic conductivity, the decoupling of ionic conductivity from mechanical properties still remains a serious challenge.^143,144 Hu's group constructed a three-dimensional (3D) covalently crosslinked solid-state electrolyte.¹⁴⁵ In this structure, amino-terminated UiO-66-NH₂ MOFs were employed as chemical crosslinking nodes, poly(tetrahydrofuran) (PTMG) and hexamethylene diisocyanate (HDI) served as soft and hard segments, respectively. The reduced crystallinity of polymer and abundant C–O–C and C=O polar sites facilitated the dissociation of lithium salt. As a result, the electrolyte exhibited a high ionic conductivity of 1.48 × 10⁻⁴ S cm⁻¹. Meanwhile, due to the incorporation of covalent bonds as well as the reversible inter/intramolecular hydrogen bonds, the obtained composite electrolytes showed superb mechanical strength (5.12 GPa) and excellent resilience. These properties were conducive to suppressing the growth of lithium dendrites. The fabricated pouch cell with LFP positive electrode can maintain stable cycling for 30 cycles, successfully powering a red LED without internal short-circuiting or burning, even under harsh operating conditions of bending, cutting, and folding. He's group prepared heat-treated polyacrylonitrile fiber network coated with interconnected MOFs (h-PAN@MOF) with a 3D continuous Li⁺ transport pathway¹⁴⁶ (Fig. 3f). When composing h-PAN@MOF with poly(vinylidene fluoride) (PVDF), a high ionic conductivity of 1.03 × 10⁻³ S cm⁻¹ can be exhibited. This can be attributed to the strong interactions between C=O of DMF and MOF crystal surface, which weaken the coordination of Li⁺ with DMF, thus enhancing the Li⁺ conductivity. Also, the network contributed to the high tensile strength of 20.84 MPa. The fabricated Li|NCM811 batteries delivered superior cycling stability for 1000 cycles at 5C.

In MOFs/polymers composite electrolyte membranes, the key issue is the tuning of interactions between MOFs and polymers. The O atoms in PEO chains can coordinate with the unsaturated sites in MOF skeletons, then enhancing the dissociation of lithium salts and the ionic conductivity. It remains a challenge to explore new types of MOF fillers to further enhance ionic conductivity of composite electrolyte membranes.

3.1.2. Recent progress of COFs-based solid-state electrolytes for LMBs.
3.1.2.1. Neat COFs. In 2018, Jiang's group first proposed a strategy of chemically grafting the oligo(ethylene oxide) onto the COFs pore walls.¹⁴⁷ The flexible side chains created a polyelectrolyte interface upon complexation with Li⁺, providing additional pathways for Li⁺ transportation. The bare pore walls of TPB-TP-COF only delivered an ionic conductivity of 2.27 × 10⁻⁷ S cm⁻¹ at 40 °C, while anchored with oligo(ethylene oxide) side chains of COFs exhibited high ion conductivity of 1858-fold than that of TPB-TP-COF. Horike and coworkers designed three PEO-functionalized hydrazone-linked COFs with different lengths of PEO side chains (PEO-x, x = 3, 6, or 9)¹⁴⁸ (Fig. 4a). The segmental motion of PEO in its bulky, flexible, and glassy state within rigid 2D COF architectures can significantly enhance Li⁺ conductivity. Moreover, the conductivities of COF-PEO-x delivered PEO lengths dependency. At 200 °C, the conductivities of COF-PEO-3-Li, COF-PEO-6-Li, and COF-PEO-9-Li were measured as 9.72 × 10⁻⁵, 3.71 × 10⁻⁴, and 1.33 × 10⁻³ S cm⁻¹, respectively. The fastest COF-PEO-9-Li can maintain original crystal structure and electrochemical stability even at 200 °C.

Fig. 4 Representative COF electrolyte membranes. (a) Synthesis of COF-PEO-x (x = 3, 6, or 9). Reproduced with permission.¹⁴⁸ Copyright 2019, American Chemical Society. (b) Two-step synthesis of holistically oriented COF solid-state electrolyte membranes. Reproduced with permission.¹⁵⁰ Copyright 2021, Wiley-VCH. (c) Synthesis of Q-COF by the cyclization reaction of I-COF. Reproduced with permission.¹⁵⁰ Copyright 2021, Wiley-VCH. (d) Fabrication of lithiated COF nanosheets. Reproduced with permission.¹⁵⁴ Copyright 2020, American Chemical Society. (e) PXRD patterns of PEG-Li⁺@COF-M electrolyte membranes. Reproduced with permission.¹⁵⁸ Copyright 2022, Elsevier.

In order to achieve high ionic conductivity COF SSEs, elaborately tuning the electronic structure of COF skeletons is regarded as an effective method. Guo's group once designed a series of COF skeletons, which combined electron-acceptors incorporated different electronegativity atoms (C, N, and F) with electron-donor monomer tetra(p-amino-phenyl)porphyrin (TAPP).¹⁴⁹ Gauss calculation revealed the distribution of electronic orbits. Notably, both the highest occupied molecular orbital (HOMO) and the lowest unoccupied molecular orbital (LUMO) in C-COF and N-COF were predominantly localized on the TAPP moiety. In sharp contrast, the LUMO of F-COF was transferred to the linkers, suggesting the pronounced D–A property due to the strong electronegativity of F atoms. This is beneficial to the Li⁺ adsorption and conduction in electrolyte membranes. As a result, the interfacial stability between electrodes and electrolytes were very different. For C-MOF and N-MOF, the dendrite and dead Li can be obviously observed. In contrast, the F-COF-based solid-state battery exhibited a uniform and dendrite-free Li metal morphology, demonstrating exceptional interfacial stability. When considering the electrochemical performances, F-COF exhibited the highest ionic conductivity among the three membranes, and the fabricated Li|LFP cells using F-incorporated COF SSEs showed high capacity retention of 90.8% after 300 cycles at 5C.

Xu and coworkers designed a high-voltage stable COF by introducing electron-withdrawing groups as well as electrochemically stable quinolyl aromatic ring linkage¹⁵⁰ (Fig. 4b). They first synthesized highly crystalline imine-linked COFs, then the reversible imine linkages were converted into more stable quinoline aromatic ring linkages through Povarov reaction (Fig. 4c). The introduced triazine and polyfluorobenzene groups into the skeleton reduced HOMO level of COFs to an ultra-low value (−6.2 eV under vacuum), which broadened the band gap and contributed to the high oxidative stability (5.6 V vs. Li⁺/Li). Furthermore, after cold-pressing of COF powders, SSE thin films were obtained with holistically oriented arrangement along the (001) facet, which provided ion migration channels. A remarkable ionic conductivity of 1.5 × 10⁻⁴ S cm⁻¹ was obtained at 60 °C of COF SSE, with an excellent mechanical strength of 10.5 GPa Young's modulus. Consequently, the assembled Li|NMC811 full cells delivered stable cycling performance over 400 cycles and a high Coulombic efficiency over 99%.

Although various COFs have been reported as SSEs for lithium batteries, they are often fabricated by cold-pressing with powders, this approach severely hindering their large-scale practical applications. Zhang's group proposed a feasible electrospinning strategy to prepare large-sized (28 cm × 2 cm), ultra-thin (20 μm), and self-supported COF SSEs.¹⁵¹ They featured a COF skeleton with lithiophilic polyethylene glycol (PEG) chains. The incorporated flexible PEG segments facilitated the dissociation of lithium salts, thereby promoting Li⁺ transportation. The as-prepared electrolyte exhibited an excellent lithium dendrite inhibiting ability and a high ionic conductivity of 0.153 mS cm⁻¹ at 30 °C. ⁷Li NMR analysis indicated that the superior conductivity can be attributed to the loose ion pairs and the existence of free Li⁺. Adopting this SSE, the fabricated Li|LFP pouch cells exhibited good cycling stability even when folded at 180°. The strategy provides a universal method for large-scale fabrication of SSEs for practical applications.

Single ion COF SSEs have been reported to enhance Li⁺ transference number. Weakly coordinated organic anion groups have been grafted onto COF skeletons (imidazolate, sulfonate, carboxylate, and sulfonimide anions) to eliminate the mobility of Li salt anions. In 2019, Lee and coworkers demonstrated a single Li⁺ conducting COF which involved lithium sulfonated groups on the pore walls (TpPa-SO₃Li).¹⁵² A high Li⁺ transference number of 0.9 can be obtained at room temperature which can be attributed to the covalently grafted anion groups. Moreover, benefiting from directional ion channels, the SSE showed an ionic conductivity of 2.7 × 10⁻⁵ S cm⁻¹. In 2020, Zhang's group adopted a pre-synthetic strategy to construct single ion conductive COF SSE by using lithium(2,5-diaminobenzenesulfonate, PaSO₃Li) as monomer.¹⁵³ Then PaSO₃Li and triformylphloroglucinol (Tp) were condensed to obtain lithium salted COF Tp–PaSO₃Li–COF. A high Li⁺ conductivity of 1.6 × 10⁻³ S cm⁻¹ was obtained at ambient temperature. Furthermore, the SSEs of Tp–PaSO₃Na–COF and Tp–PaSO₃K–COF were also synthesized by using sodium and potassium salts as monomers, respectively. The prepared SSEs both delivered high ionic conductivity, verifying the universality of this pre-synthetic strategy. Loh and coworkers synthesized SD-COF-3 by post-synthetic modification through photochemical thiol–ene click reaction.¹⁵⁴ Then, lithiation was conducted with aqueous lithium carbonate and COF powders were exfoliated into nanosheets (Li-CON-3) due to the Coulombic repulsion between layers (Fig. 4d). The presence of well-ordered 1D channels made Li-CON-3 a candidate COF SSE at low temperature. The ionic conductivity can be up to 10⁻⁵ S cm⁻¹ at −40 °C and the transference number of Li⁺ was 0.92. When using 1,4-benzoquinone (BQ) as cathode material and Li-CON-3 as SSE, the all-solid-state cells delivered high cycling performance of 500 cycles with no obvious capacity decay at a current density of 500 mA g⁻¹.

Neat COF solid-state electrolyte membranes show potential to construct high-performance batteries, due to the easy modifiability and high stability. Polar functional groups can be easily grafted on the pore walls of COF skeletons to enhance the ionic conductivity as well as transference number. Meanwhile, elaborating the electronic structures of COF skeletons through choosing proper monomers can also promote the electrochemical performance of batteries. However, the conductivity of COF electrolyte should be further enhanced in order to reduce the content of conductivity agent. Meanwhile, the application of COF electrolyte membranes in lithium–sulfur (Li–S) batteries should be further explored due to the weaker catalytic conversion of polysulfides compared with that of MOF electrolytes.

3.1.2.2. COFs/polymers. Polymers such as polyacrylonitirle (PAN),¹⁵⁵ PEO,¹⁵⁶ and PVDF¹⁵⁷ are often used as electrolyte materials. However, the low ionic conductivity, especially at ambient temperature, severely hinders the application of SSEs. Preparing hybrid materials with polymers and COFs has attracted much attention, which can both utilize the processability of polymers as well as the advanced pore structures of COFs. In 2022, Zhang and coworkers developed a confinement approach to encapsulate PEG–Li⁺ chains within COFs, designated as PEG–Li⁺@COF-M.¹⁵⁸ In composite electrolyte, the crystalline of PEG was significantly reduced due to the existence of COF (Fig. 4e). In particular, when the content of bulk PEG was 69%, the materials showed a totally amorphous state. After confining into 1D channel of COF membrane, the relaxation time of PEG was shortened from 6.1 × 10⁻³ s to 6.4 × 10⁻⁶ s. Remarkably, the composite electrolyte achieved ionic conductivities of 2.2 × 10⁻⁵ S cm⁻¹ (20 °C) and 1.9 × 10⁻³ S cm⁻¹ (120 °C). Benefiting from the confinement effect, the lithium salt dissociation was reinforced. The assembled Li|LFP full cells delivered a high specific capacity of 135.7 mA h g⁻¹ after 80 cycles with a 97.6% Coulombic efficiency and stabilized the evolution of the electrode–electrolyte interface. Dong's group proposed a strategy to combine keto-enamine COF (TPBD) with PEO to prepare SSEs called TPBD–LiPF₆@PEO.¹⁵⁹ The strong interactions between C=O in COF skeletons and Li⁺ facilitated the dissociation of lithium salts, thus enhancing the ionic conductivity to 0.543 mS cm⁻¹ at room temperature and 3.23 mS cm⁻¹ at 70 °C. As a consequence, the assembled Li|LFP full batteries using TPBD–LiPF@PEO electrolyte offered a specific capacity of 140 mA g⁻¹ at 0.2C and can operate smoothly after 200 cycles at 0.2C.

3.2. Catalytic electrodes

3.2.1. Electrocatalytic HER electrodes. Electrocatalytic HER refers to the process of promoting the reduction of H⁺ through electron transfer on the electrode surface to produce H₂. It is one of the key technologies for realizing the hydrogen economy, because H₂, as a high-energy-density and environmentally friendly energy carrier, is an important part of the future energy structure. Moreover, water electrolysis presents great prospects as a strategic approach to green hydrogen generation using renewable energy. In this process, electrocatalysts play a vital role. They can accelerate the electron transfer process, thereby improving the reaction rate and efficiency. MOFs have broad application prospects in the field of electrocatalytic hydrogen evolution due to their advantages such as high porosity and surface area, structural controllability, and multifunctionality. They are expected to displace Pt-group electrocatalysts, reduce the cost of hydrogen production, and promote the widespread application of hydrogen energy.

Although the unique topological structure of MOFs offers numerous advantages as a catalyst, MOFs have long faced challenges in practical electrocatalytic applications for a long time. This is because MOFs have the inherent disadvantage of low conductivity and still show instability in the reaction and lack long-term stability.¹⁶⁰ Many researchers have contributed to improving the conductivity and stability of MOFs.^161–166 Yin's group introduced CoP species with better conductivity on the surface of Co–MOF through partial phosphating. CoP, as a conductive “coat”, covered the surface of Co–MOF, establishing a conductive path, which effectively improved the overall conductivity of the material. Furthermore, partial phosphating allowed CoP species to connect with Co–MOF through N–P/N–Co bonds to form a stable hybrid structure. This bonding can enhance the structural stability of Co–MOF and reduce its structural damage during the electrolysis process. CoP species covered the surface of Co–MOF to form a protective layer, which can effectively isolate the electrolyte from the erosion of Co–MOF and reduce the risk of dissolution and structural collapse of Co–MOF in acidic and alkaline environments.¹⁶¹ Peng's group¹⁶² adopted a simple spontaneous redox reaction strategy to anchor noble metal nanoparticles (NPs) on nickel MOF nanohybrid materials (M@Ni–MOF, M = Ru, Ir, Pd) (Fig. 5a). This method cleverly utilized nickel foam (NF) as both a reducing agent to participate in the reaction and a conductive substrate to support the catalyst, realizing the in situ synthesis and fabrication of the catalyst. The optimized Ru@Ni–MOF catalyst exhibited excellent HER activity under alkaline conditions, with a low overpotential of only 22 mV to achieve a current density of 10 mA cm⁻², exceeding commercial Pt/C and physically mixed Ru/C and Ni–MOF. At the same time, the catalyst also showed good catalytic activity and stability under acidic and neutral conditions, indicating its application potential in the full pH range. The 3D freestanding architecture facilitated rapid transport of electrons, mass, and charge while expanding the electrochemically active surface area and enhancing electrolyte accessibility, thereby significantly improving catalytic activity. The formation of Ni–O–Ru bonds strengthened the interaction between Ru nanoparticles and Ni–MOF, optimizing the electronic structures of both Ru and Ni atoms. This synergistic effect promoted the adsorption and dissociation of reaction intermediates while reducing the energy barrier for the HER, ultimately enhancing catalytic performance. Cheng's group¹⁶³ prepared a novel hierarchical composite catalyst Ni₃S₂@2D Co–MOF/CP by in situ growing 2D Co–MOF nanosheets on carbon paper and further electrodepositing Ni₃S₂ nanosheets to construct a unique core–shell structure. They cleverly combined the high specific surface area and abundant active sites of 2D MOF with the high conductivity and adsorption capacity of Ni₃S₂ for hydrogen intermediates, achieving a synergistic effect and effectively improving the catalytic performance. The defects and oxygen vacancies on the surface of 2D Co–MOF promoted the adsorption and dissociation of water, reduced the energy barrier of water decomposition, and thus improved the efficiency of hydrogen evolution. The Ni₃S₂@2D Co–MOF/CP catalyst exhibited excellent alkaline HER catalytic performance. At a current density of 10 mA cm⁻², the overpotential was only 140 mV and the Tafel slope was only 90.3 mV dec⁻¹, which was lower than those of Co–MOF/CP and Ni₃S₂/CP catalysts. At the same time, the catalyst has excellent stability. In 2000 consecutive cyclic voltammetry scans and 12 hours of constant current chronopotentiometry, the catalytic performance has almost no attenuation. Fujita's group¹⁶⁴ used HOFs as templates and achieved the controllable synthesis of various transition metal MOF nanotubes by modulating metal–ligand coordination environments. This strategy effectively avoids the collapse or overgrowth of MOF nanostructures during the conversion process by forming a protective shell. Using Ni, Ru, and Ir–MOF nanotubes as precursors, Ni, Ru, and Ir@NGTs (NGTs = nitrogen-doped graphene nanotubes) catalysts with high conductivity, abundant active centers, and stable structures were prepared by simple carbonization treatment (Fig. 5b). Ni, Ru, and Ir@NGTs exhibited excellent HER catalytic activity under alkaline conditions, and their overpotential, Tafel slope, and TOF values were superior to those of commercial Pt/C catalysts and other catalysts reported in the literature previously. Zhu's group¹⁶⁶ achieved uniform dispersion and size reduction of Ru nanoparticles by confining Ru nanoparticles in the 3D pores of Ni-BPM (Fig. 5c). There was a strong metal–support interaction between Ru nanoparticles and Ni-BPM, which promoted electron transfer and optimized the Gibbs free energy of water molecule dissociation and hydrogen atom adsorption, thereby improving HER kinetics and stability. The Ru@Ni-BPM/NF electrode exhibited excellent HER performance. In 1 M KOH solution, only 12 mV overpotential was required to reach a current density of 10 mA cm⁻², and the Tafel slope was only 40.3 mV dec⁻¹, which was better than that of Ru@Ni-phenyldicarboxylic (BPDC)/NF and commercial Pt/C catalysts. The Ru@Ni-BPM/NF electrode also exhibited excellent stability. After continuous testing at a current density of 100 mA cm⁻² for 100 hours, the performance had almost no attenuation, while the performance of the Ru@Ni-BPDC/NF electrode dropped significantly after 24 hours.

Fig. 5 (a) Synthetic pathway illustration of the formation of Ru@Ni–MOF nanosheets on conductive Ni foam scaffolds for HER. Reproduced with permission.¹⁶² Copyright 2021, Wiley-VCH. (b) Proposed schematic illustration of the synthetic route of NiRu–MOF nanotubes. Reproduced with permission.¹⁶⁴ Copyright 2022, American Chemical Society. (c) Synthetic pathway illustration of the formation of Ru@Ni–MOFs. Reproduced with permission.¹⁶⁶ Copyright 2024, Elsevier.

In summary, the porous architecture of MOFs provides abundant exposed metal active centers that significantly enhance HER reaction kinetics. Furthermore, strategic hybridization with noble metal NPs can concurrently improve electronic conductivity and catalytic activity, positioning MOF composites as promising electrocatalytic materials for hydrogen evolution. However, most MOFs exhibit intrinsic limitations as insulators or semiconductors with poor native conductivity, while also suffering from structural dissolution/collapse in acidic/alkaline electrolytes. These challenges can be effectively addressed through: (i) integration of conductive phases (like metal phosphides) to establish efficient electron transport pathways, and (ii) engineered protective coatings to enhance structural stability collectively advancing the practical implementation of MOF materials in electrocatalytic HER applications.

3.2.2. Electrocatalytic CO₂RR electrodes. Electrocatalytic CO₂RR is considered as an effective approach to reduce atmospheric CO₂ concentrations and convert CO₂ into valuable chemicals such as methanol and methane. Crystalline porous framework materials have demonstrated significant potential in enhancing the activity and selectivity of CO₂ electrochemical reduction reactions due to their unique structural and compositional properties. These materials have the characteristics of high specific surface area, high porosity, adjustable pore structure, and excellent crystallinity, which can provide sufficient active sites for reactants, hereby facilitating significant progress in the field of electrocatalytic CO₂RR.¹⁶⁷ Moreover, the unique structural and compositional properties of crystalline porous framework materials offer significant opportunities for low-concentration CO₂ reduction applications. These materials can effectively convert CO₂ into high-value chemicals under complex conditions, thereby playing a crucial role in the sustainable carbon economy.^168,169

Recent advancements in the field of electrocatalytic CO₂ reduction have been significantly propelled by the development and application of various COF and MOF-based catalysts, which have demonstrated remarkable improvements in selectivity, efficiency, and stability for converting CO₂ into valuable chemicals. In terms of improving electrocatalytic efficiency, directionally grown crystalline framework film materials can enhance catalytic efficiency by increasing charge transfer rate and electron utilization efficiency. Sun et al.¹⁶⁹ reported depositing a Re-based MOF thin film with high electrochemical efficiency onto a conductive fluorine-doped tin oxide (FTO) electrode using liquid-phase epitaxy (Fig. 6a). Thanks to the directional arrangement and growth of MOF membranes, the electron transfer speed and efficiency have been accelerated. This epitaxially grown MOF film exhibited an extremely high Faraday efficiency of 93.5% when used as an electrocatalyst to reduce CO₂ to CO. This is the highest efficiency reported so far, surpassing COF films and other MOF films.

Fig. 6 (a) Epitaxially grown MOF film and electrocatalytic reduction CO₂ to CO performance. Reproduced with permission.¹⁶⁹ Copyright 2016, Royal Society of Chemistry. (b) Reported COF films and using for electrocatalytic CO₂ reduction. Reproduced with permission.¹⁷⁰ Copyright 2018, American Chemical Society. (c) Schematic diagram of PCN-222(Fe)/C electrocatalytic reduction of CO₂. Reproduced with permission.¹⁷¹ Copyright 2018, American Chemical Society. (d) Solvent-free method for preparing COF films on carbon cloth electrodes and application for CO₂ reduction. Reproduced with permission.¹⁷³ Copyright 2019, American Chemical Society. (e) Schematic diagram of Al₂(OH)₂TCPP-Co thin film for CO₂RR. Reproduced with permission.¹⁷⁴ Copyright 2015, American Chemical Society. (f) Preparation of NU-1000 by solvent thermal deposition method. Reproduced with permission.¹⁷⁶ Copyright 2020, American Chemical Society. (g) MFM-300(In) on indium foil to improve the efficiency of CO₂RR. Reproduced with permission.¹⁷⁹ Copyright 2020, American Chemical Society.

Directional arrangement of crystalline framework materials can not only increase electron utilization efficiency but also improve selectivity. Yaghi et al.¹⁷⁰ reported the synthesis of a series of COF films on the surface of a highly ordered pyrolytic graphite (HOPG) electrodes (Fig. 6b). These films were prepared by modifying the benzene ring with different functional groups (BDA-(R1)₂(R2)₂) attached to a tetramine porphyrin (Co(TAP)) coordinated with cobalt. Then the resulting COF films were employed for electrocatalytic CO₂ reduction. The results demonstrated that this series of oriented COF films exhibited high selectivity (Faraday efficiency of 87%) and high current density (65 mA cm⁻²) in reducing CO₂ to CO at a low overpotential (550 mV), far exceeding other reported catalysts in terms of selectivity and efficiency. Molecular catalysts can exist stably for more than 12 hours without a decrease in reaction activity. Hupp et al.¹⁶⁸ reported that electrophoretic deposition of Fe–MOF-525 grains can effectively achieve catalyst heterogeneity and increase surface concentration, thereby enhancing the electrochemical reduction of CO₂. Further research found that this method can install the known CO₂ reduction catalyst Fe–TPP on the electrode surface at a high surface concentration equivalent to ∼900 monolayers, which is nearly ten times higher than the highest reported concentration of heterogeneous molecular catalysts currently available. The nanoscale pore structure of MOFs not only promoted the contact of solvents, reactants, and electrolytes with catalytic sites, but also facilitated the use of metalloporphyrin linkers as catalysts and channels for the transfer of reducing equivalents through electron hopping.

The introduction of metal active sites plays an important role in improving the catalytic activity of materials. Li et al.¹⁷¹ reported the preparation of a highly stable porous porphyrin-based MOF catalyst supported by PCN-222(Fe) using a simple dip coating method, and applied it to electrocatalytic conversion of CO₂ to CO (Fig. 6c). This composite catalyst exhibited excellent catalytic performance in a 0.5 M KHCO₃ saturated CO₂ solution, with an overpotential of 494 mV and 91% CO selectivity, and maintaining good crystallinity and stability after 10 hours of electrolysis process. Further research found that PCN-222(Fe)/C achieved significant electrocatalytic CO₂ reduction through the intrinsic activity of pore porphyrin molecules, excellent CO₂ adsorption capacity, and high conductivity of carbon black. Yaghi and his coworkers¹⁷² reported the construction of COFs cobalt porphyrin catalysts connected by imine bonds through a modular approach. The prepared COFs were used as catalytic materials for the electrochemical reduction of CO₂ to CO in aqueous phase. The experimental results showed that at pH = 7, the catalyst exhibited high Faraday efficiency (90%) and conversion number (up to 290 000, initial turnover frequency of 9400 h⁻¹), overpotential of −0.55 V, equivalent to 26 times the activity of the molecular cobalt complex and did not degrade within 24 hours. Kubiak et al.¹⁷³ reported a solvent-free method for preparing COF films on carbon cloth electrodes using 5,10,15,20-tetra-(4-aminophenyl)-porphyrin Fe(III) chloride (FeTAPPCl) and 2,5-dihydroxyterephthalaldehyd (Dha), and applied it for CO₂ reduction (Fig. 6d). The COF films exhibited good conversion frequency (>600 h⁻¹ mol⁻¹ electroactive Fe sites) and reasonable Faraday efficiency of CO (average 80%) at −2.2 V (vs. Ag/AgCl) for 3 hours.

Researchers have also conducted a series of studies to improve the catalytic selectivity of materials. Yaghi et al.¹⁷⁴ introduced nanoscale cobalt porphyrin MOF Al₂(OH)₂TCPP-Co (TCPP = tetra-(4-carboxyphenyl) porphyrin) thin film materials as efficient and selective catalysts for the reduction of CO₂ to CO in aqueous electrolytes (Fig. 6e). The selectivity of CO generation in this MOF film exceeded 76% and remained stable within 7 hours, with a per-site turnover number (TON) of 1400. Eddaoudi et al.¹⁷⁵ reported the synthesis of MOF thin films on gold nanostructured microelectrodes using both layer-by-layer (LbL) deposition and solvothermal deposition methods. They also studied the CO₂RR activity of these films. Meanwhile, electrocatalytic experiments revealed that this well-coated structure could completely suppress CO production while generating small amounts of CH₄ and C₂H₄. This work represented the first reported use of MOF films to regulate the CO₂RR activity of underlying catalysts.

Various forms of nanomaterials have also been introduced into the electrodes of crystalline framework membrane materials to enhance their catalytic performance. Hupp and his coworkers¹⁷⁶ reported the solvothermal deposition of Cu(II) in MOFs, followed by electrochemical reduction of Cu(II) to metallic Cu and embedding copper nanoparticles into a solvothermal grown thin film of zirconium MOF (NU-1000) (Fig. 6f). The confinement effect of MOF nanopores enabled precise electrochemical modulation of the synthesized Cu nanoparticles, which exhibited enhanced electrocatalytic performance for CO₂ reduction in aqueous electrolytes. Gu et al.¹⁷⁷ reported the preparation of a series of MOF Ni(Im)₂ nanosheets with different thicknesses using Ni²⁺ and imidazole (Im) ligands. These nanosheets exhibited varying electrocatalytic CO₂ reduction activities, with nanosheets 2–5 nm thick having more active sites and exhibiting higher catalytic selectivity. Buonsanti et al.¹⁷⁸ reported the preparation of Al-PMOF ([Al₂(OH)₂(TCPP)], tetra-(4-carboxyphenyl)porphyrin (TCPP)) on the surface of Ag nanocrystals (Ag NCs), forming Ag@Al-PMOF hybrid materials with excellent conductivity, which can effectively inhibit HER and promote CO₂RR. Han et al.¹⁷⁹ developed a MOF-based electrode prepared by electro-synthesis of MFM-300(In) on indium foil to improve the efficiency of CO₂ electroreduction (Fig. 6g). The MFM-300(In)-e/In electrode achieved a current density of 46.1 mA cm⁻² at a voltage of −2.15 V and a Faraday efficiency of 99.1%, surpassing all other MOF systems specifically used for synthesizing formic acid. The simple preparation and excellent electrochemical stability of MFM-300(In)-e/In electrode provided a new method for developing efficient CO₂ reduction electrocatalysts.

In summary, crystalline framework materials have demonstrated excellent electrocatalytic performance due to their well-defined/tunable chemical structure, high surface area, and easily accessible active sites, resulting in membranes and composite membranes. At the same time, the crystalline solid state of this membrane can provide a convenient, recyclable, and robust substrate that is resistant to chemical and physical treatments. Moreover, by designing the membrane reasonably, it can maximize exposure to active sites and mass transfer, and thus provide high catalytic activity. However, research on electrocatalysis of these materials is currently mainly focused on the fabrication of MOFs in thin film form, while COFs and HOFs have not been thoroughly explored so far. It is necessary to further investigate this area in the future.

3.3. Solar cells

Crystalline porous materials, due to their porosity and large surface area, have been widely used in solar cells. These materials can remarkably enhance light absorption rates and improve carrier separation efficiencies. Meanwhile, they can also reduce both the reflection and refraction of solar radiation, thereby enabling more optical energy to be converted into electrical energy.¹⁸⁰ Certainly, crystalline porous framework materials exhibit high chemical and thermal stability, which play crucial roles in enhancing the stability and durability of solar cells. Moreover, their structural tunability enables researchers to optimize solar cell performance to meet the specific requirements of various applications. This capability provides solid support for the green and sustainable development of solar cells. These materials can be incorporated into various types of solar cells, such as dye-sensitized solar cells (DSSCs), perovskite solar cells (PSCs), and organic solar cells (OSCs).

3.3.1. Dye-sensitized solar cells (DSSCs). Crystalline porous framework materials have demonstrated remarkable advantages in the application of DSSCs. Their highly porous structures facilitate the separation and transport of charge carriers, providing abundant reaction surfaces and active sites. This enhances the adsorption capacity and light absorption efficiency of dye molecules, thereby boosting energy conversion efficiency. Additionally, DSSCs can benefit from their stable structures and optimize light absorption and conversion efficiency, which further enhance their performance enhancements. Several reports have already been published on the application of crystalline framework membrane materials in the field of DSSCs.¹⁸¹

Crystalline porous framework materials can not only suppress interface charge recombination but also improve the performance of DSCCs by enhancing their light absorption ability and prolonging carrier lifetime. Wei et al.¹⁸² coated a 2-nm-thick ZIF-8 film on TiO₂ to enhance the open-circuit voltage (V_oc) of DSSCs. Electrochemical tests revealed that the increase in V_oc is due to the ZIF-8 shell material, which effectively suppressed interfacial charge recombination. This study offered a novel strategy for improving the photovoltaic performance of DSSCs. Yin et al.¹⁸³ proposed a method for synthesizing Cu₂ZnSnS₄ (CZTS) nanoparticles in porous TiO₂ derived from sensitized MOFs via thermal injection and hydrothermal synthesis. This material significantly broadened the light response range of TiO₂ in the visible light region by enhancing its light absorption capacity and extending the carrier lifetime. As a result, the photocurrent and photoelectric conversion efficiency (PCE) of TiO₂ photoelectrodes derived from CZTS MOFs increased by nearly 1.93 and 2.21 times, respectively (Fig. 7a).

Fig. 7 (a) Energy band structure and photo-generated charge transfer mechanism in CZTS nanoparticles/MOFs-derived TiO₂. Reproduced with permission.¹⁸³ Copyright 2016, American Chemical Society. (b) Schematic diagram of PPF-11 film structure and its photovoltaic response curve. Reproduced with permission.¹⁸⁵ Copyright 2019, American Chemical Society. (c) Tp-Azo-COF into the spiro-OMeTAD HTL to improve the optoelectronic performance and stability of PSCs and reduce Pb leakage. Reproduced with permission.¹⁸⁴ Copyright 2024, Wiley-VCH. (d) Device structure of the inverted PSCs incorporating COFs and photovoltaic response curve. Reproduced with permission.¹⁸⁶ Copyright 2024, Wiley-VCH. (e) Schematic illustration of mobile defects and Li⁺ migration being suppressed by the addition of COF@TCNQ or COF into the spiro-OMeTAD layer, and their photovoltaic response curve. Reproduced with permission.¹⁸⁷ Copyright 2024, American Chemical Society. (f) Schematic illustration of the process to prepare 2D MOF nanosheets via simultaneous exfoliation and functionalization with PEIE and the device structure of the inverted polymer solar cell with photovoltaic response curve. Reproduced with permission.¹⁸⁸ Copyright 2018, Elsevier.

The structural and interfacial engineering of crystalline porous frameworks through composite electrode design (e.g., ZIF-8/PEDOT:PSS) or precisely oriented framework architectures (e.g., PPF-11) has proven critical for advancing the photovoltaic performance of DSSCs. Recent advancements in the field of DSSCs have focused on enhancing their performance through innovative electrode materials and fabrication techniques. Zhao et al.¹⁸⁹ studied the application of ZIF-8/PEDOT:PSS composite counter electrode in DSSCs. ZIF-8 was combined with conductive polymer PEDOT:PSS and counter electrodes were prepared by spin coating at low temperatures. This composite electrode exhibited high PCEs in both iodine-based and cobalt-based electrolyte DSSCs of 7.02% (3% ZIF-8 content) and 6.84% (1% ZIF-8 content), respectively. Saha et al.¹⁸⁵ reported the preparation of [100] columnar porphyrin framework-11 (PPF-11) thin films with precise orientation on annealed ZnO-FTO surfaces using a spontaneous solvothermal growth method (Fig. 7b). Owing to the precise orientation, the PPF-11/ZnO-FTO photoanodes demonstrated an excellent photovoltaic response, with a short-circuit current density (J_SC) of 4.65 mA cm⁻², a V_oc of 470 mV, and a PCE of 0.86%. These values are superior to those of all control devices and previously reported visible-light-harvesting MOFs based on porphyrin and Ru(bpy)₃²⁺.

Innovative strategies for enhancing the performance of DSSCs have been explored through the modification of photoelectrodes with MOFs and the incorporation of advanced materials such as reduced graphene oxide (RGO). He et al.¹⁹⁰ employed UiO-66 and ZIF-8 to modify photoelectrodes and incorporated RGO to fabricate all-carbon DSSCs based on UiO-66-RGO-TiO₂ photoelectrodes. The substantial specific surface area of MOFs notably enhanced dye loading and photo-induced electron generation, resulting in a high PCE of 7.67%, significantly surpassing the performance of RGO-TiO₂ and pure TiO₂ photoelectrodes. Their group also prepared DSSCs photoanodes co-modified with ZIF-8 and a 3D graphene network (3DGN).¹⁹¹ The large specific surface area of ZIF-8 significantly increased the dye loading and higher J_SC of the device, while the continuous graphene layer structure in 3DGN provided a fast transmission channel for photogenerated electrons. The obtained fill factors (FF), J_SC, V_OC, and PCE were 61.1%, 20.9 mA cm⁻², 681 mV, and 8.77%, respectively.

Crystalline porous frameworks have proven to be a versatile platform for enhancing DSSC performance through multiple mechanisms, including improved charge transport, suppressed recombination, and enhanced light absorption. Future research may focus on optimizing MOF structures, exploring new hybrid materials, and developing scalable fabrication techniques to further push the boundaries of DSSC efficiency.

3.3.2. Perovskite solar cells (PSCs). PSCs have emerged as a prominent photovoltaic technology, characterized by their high PCEs and scalable solution-based fabrication methods. The recent introduction of MOFs and COFs has proven to be an effective strategy to significantly enhance both the efficiency and stability of PSCs, as widely validated.¹⁸⁰ Yang et al.¹⁹² synthesized a Zn-based MOF (Zn-TTB) via self-assembly from Zn²⁺ and 1-(triazol-1-yl)-4-tetrazole-5-ylmethylbenzene (TTB), allowing the perovskite precursor to grow in conjunction with the MOF backbone. This not only improved the stability of the perovskite film but also provided excellent thermal stability under an N₂ atmosphere at 85 °C. Meanwhile, Zn-TTB exhibited an asymmetric configuration with unequal electronegativity, thereby generating electric dipoles. Consequently, the transfer of perovskite charge carriers was enhanced, and the degradation activation energy was increased to 174.01 kJ mol⁻¹. The device achieved a J_SC of 25.16 mA cm⁻² and a PCE of up to 23.13%.

The strategic integration of 2D-COFs into PSCs has emerged as a powerful approach to enhance interfacial engineering, mitigate stability challenges, and boost photovoltaic efficiency. Park et al.¹⁸⁴ doped a chemically modified 2D conjugated covalent organic framework (Tp-Azo-COF) into the 2,2′,7,7′-tetrakis-(N,N-di-4-methoxyphenylamino)-9,9′-spirobifluorene (spiro-OMeTAD) HTL to improve the optoelectronic performance and stability of PSCs and reduce Pb leakage (Fig. 7c). The research results indicated that Tp-Azo-COF effectively reduced lead leakage and Li⁺ migration, significantly improving the environmental safety and operational stability of PSCs. The optimized PSCs achieved an efficiency of 24.25% in a small size (0.12 cm²) and maintained an efficiency of 21.96% on a large area module (18 cm²). Wang et al.¹⁸⁶ synthesized two 2D-COFs (TABT-COF and BABT-COF) based on benzotrithiophene (BTT) with high crystallinity, high specific surface area, and excellent thermal stability (Fig. 7d). These 2D-COFs were doped into the perovskite layer of PSCs and used as efficient additives. The addition of BTT-based COF can serve as an effective nucleation template to regulate the crystallinity of perovskite. The PCEs of TABT-COF and BABT-COF can reach 18.10% and 18.56%, respectively, which are better than that of the control equipment (16.58%). Zhao et al.¹⁸⁷ prepared BPTA-TAPD-COF by reacting N,N,N′,N′-tetrakis(4-aminophenyl)-1,4-phenylenediamine (TAPD) with 2,5-bis(2-propynyloxy)terephthalaldehyde (BPTA) then integrated 7,7,8,8-tetracyanoquinodimethane (TCNQ) and gained the BPTA-TAPD-COF@TCNQ to enhance the optoelectronic performance and stability of perovskite solar cells and incorporated it into the spiro-OMeTAD HTL (Fig. 7e). The integration of COF and TCNQ into the COF framework accelerated the extraction of charge carriers, suppressed lithium migration, and captured mobile defects, thereby significantly enhancing the PCE and long-term stability. Tests have demonstrated that benefiting from the increase in conductivity, adding BPTA-TAPD-COF@TCNQ can significantly enhance device performance. The highest PCE of this device is 24.68%, and it has long-term stability under harsh conditions. Collectively, these studies underscore the transformative potential of COFs in addressing critical challenges in PSCs, from defect passivation and ion migration suppression to scalable high-efficiency device fabrication.

Porous framework materials represented by MOFs and COFs demonstrate remarkable potential in perovskite solar cells due to their tailorable pore architecture, high specific surface areas, and robust chemical properties. By precisely regulating perovskite crystallization, optimizing charge carrier transport, and suppressing ion migration pathways, these materials not only achieve record efficiencies (exceeding 24%) but also address environmental concerns (e.g., lead leakage mitigation), establishing a novel paradigm for high-performance stable perovskite devices. Future advancements through structural optimization and multidimensional functional integration promise to surpass current efficiency thresholds while fulfilling commercial viability requirements for long-term operational durability.

3.3.3. Organic solar cells (OSCs). OSCs harness organic semiconductors to convert sunlight into electrical energy, characterized by their lightness, flexibility, and potential for low-cost production. Due to the relatively weaker semiconducting properties of crystalline porous framework materials, the application in OSCs is limited primarily to interlayer engineering. Research in this field is still in its early stages, and the development of crystalline porous framework materials for OSCs applications requires further exploration. Huang et al.¹⁸⁸ developed a novel 2D MOF based on tellurophene for inverted polymer OSCs and successfully exfoliated it into single and few-layer nanosheets (Fig. 7f). The hybrid MOF-PEIE (polyethyleneimine ethoxylate), used as an electron extraction layer (EEL), exhibited higher PCE than that of traditional PEIE. Overall, the integration of crystalline porous frameworks like MOFs into OSCs highlights a promising pathway to enhance charge extraction efficiency and device performance, emphasizing the necessity for continued innovation in material design and interfacial engineering.

3.4. Separation membranes

3.4.1. Gas separation membranes. Gas separation plays a pivotal role in industrial production and environmental protection. Membrane-based separation has gained attention as an energy-efficient technology, owing to their inherent advantages of operational simplicity and potential for large-scale implementation. Recent advances in the rational design of crystalline porous separation membranes will be discussed, in the fields of hydrocarbon separation, H₂ purification, and CO₂ separation.
3.4.1.1. Hydrocarbon separation membranes. Olefin/paraffin separation is one of the indispensable separations in modern chemical industry, because light olefins including ethylene (C₂H₄) and C₃H₆ are the basic chemicals with the largest production and broadest application.^193,194 Currently, such separation is realized by an energy-intensive cryogenic distillation method. Membrane separation technology is regarded as one of the most promising alternatives because of its high energy efficiency and low carbon footprint. It is predicted that at least 80% energy can be saved if membrane technology replaces cryogenic distillation technology.^195–197 Thus, exploiting advanced membrane materials for efficient olefin/paraffin separation has become an essential pursuit.

MOFs represent one of the most promising membrane materials for olefin/paraffin separations. However, the thickness of most MOF membranes exceeds 1 μm, which significantly compromises gas permeation performance. In 2020, Zhong et al.¹⁹⁸ reported an innovative interfacial polarization-induced approach for fabricating ultra-thin low-crystallinity MOF (LC-MOF) membranes with exceptional C₃H₆/C₃H₈ separation performance (Fig. 8a). This study elucidates the relationship between the structural characteristics and gas separation performance of LC-MOF membranes (Cu-BTC, ZIF-8, and DZIF-8), highlighting their advantages for gas separation. The low-crystallinity structure, induced by interface layer polarization generated abundant OMSs that enhance π-bond selective adsorption and transport of C₃H₆, while the ultra-thin membrane thickness (45–150 nm) significantly improved gas permeance. Notably, the diethanolamine (DEA)-modified DZIF-8 membrane, through the introduction of synergistic ligands, further reduced ligand coordination capability, achieving both ultra-high permeance (2000–3000 GPU) and high selectivity (C₃H₆/C₃H₈ selectivity: 90–120). In contrast, high-crystallinity MOF membranes exhibited limited permeance due to their dense structure, and Cu-BTC membranes displayed weaker separation performance owing to strong hydrogen bonding interactions between carboxyl oxygen and C₃H₈. This comparative analysis demonstrated that synergistic optimization of structural parameters-such as crystallinity, OMSs, and membrane thickness effectively overcome the traditional trade-off between permeance and selectivity in membranes, enabling advanced gas separation performance. In 2022, Jiang et al.¹⁹⁴ developed a novel approach for fabricating ultra-thin ZIF-8 membranes via an inhibited Ostwald ripening (IOR) strategy to achieve highly efficient C₃H₆/C₃H₈ separation (Fig. 8b). By incorporating polymer-based inhibitors with electron-donating groups into the ZIF-8 precursor successfully suppressed crystal growth, reducing the membrane thickness to a remarkable 180 nm. This structural optimization enables simultaneous enhancement of permeance and selectivity. The resulting membrane demonstrated exceptional separation performance with a C₃H₆ permeance of 386 GPU and a SF of 120. The IOR strategy enables precise control over membrane thickness by tuning the functional groups, molecular weights, and concentrations of the inhibitors. In 2023, Liu et al.¹⁹⁹ successfully fabricated a highly c-axis-oriented Cu@NH₂-MIL-125 membrane with sub-100-nm thickness through synergistic pore micro-environment engineering and meso-structure manipulation (Fig. 8c). The immobilized coordinatively unsaturated Cu sites within the NH₂-MIL-125 framework facilitated strong π-complexation interactions with C₂H₄, achieving an exceptional C₂H₄/C₂H₆ selectivity of 13.6, representing a 9.4-fold enhancement compared to the pristine NH₂-MIL-125 membrane. In 2024, Zhong et al.²⁰⁰ reported a ligand back diffusion-assisted bipolymer-directed strategy for fabricating large-area (over 2400 cm²) ultra-thin (50–130 nm) MOF membranes on flexible polymer supports (Fig. 8d). By utilizing a cross-linked bipolymer network to uniformly distribute metal ions and controlling ligand diffusion, the method enables the scalable production of defect-free ZIF-8, BUC-112, and MIL-68 membranes. The ZIF-8 membrane demonstrated exceptional C₃H₆/C₃H₈ separation performance, achieving a permeance of 3979 GPU with a selectivity of 43.88. Notably, spiral-wound modules (4800 cm²) incorporating these membranes maintained high performance (3930 GPU for C₃H₆), confirming the industrial viability of this methodology for energy-efficient gas separation.

Fig. 8 Several porous membranes related to hydrocarbon separation. (a) Fabrication of a dense ultra-thin LC-MOF membrane onto the interface layer. Reproduced with permission.¹⁹⁸ Copyright 2020, Wiley-VCH. (b) Illustration of ZIF-8 and IOR-ZIF-8 membrane fabrication via inhibited Ostwald ripening. Reproduced with permission.¹⁹⁴ Copyright 2022, Wiley-VCH. (c) Schematic depiction of the fabrication process for a c-axis-aligned ultra-thin π-coordination Cu@NH₂-MIL-125 membrane. Reproduced with permission.¹⁹⁹ Copyright 2023, Wiley-VCH. (d) Development process for a large-scale MOF selective layer formation. Reproduced with permission.²⁰⁰ Copyright 2024, Wiley-VCH.

In the field of hydrocarbon separation, MOF membranes have demonstrated remarkable advantages owing to their designable pore micro-environments and controllable crystal growth. Through multi-scale modulation strategies including interfacial polarization, IOR strategy, coordination engineering, and oriented growth, highly oriented and defect-minimized membrane architectures were precisely constructed. These membranes leveraged open metal sites or π-complexation-active centers to enhance preferential adsorption and diffusion kinetics of alkenes (e.g., C₃H₆, C₂H₄), achieving synergistic improvements in both high permeance (C₃H₆ permeance up to 3979 GPU) and selectivity (C₂H₄/C₂H₆ separation factor of 13.6). Furthermore, by integrating flexible substrate interfacial engineering and modular assembly technologies, such membranes exhibited scalable fabrication potential while maintaining mechanical robustness. This advancement provides a low-energy-consumption and cost-effective innovative pathway for industrial-scale olefin/paraffin separations, bridging the gap between laboratory research and practical applications, and propelling membrane-based separation technologies toward sustainable and efficient industrial implementation.

3.4.1.2. H₂ purification membranes. H₂ is a crucial carrier to realize a sustainable, low-carbon energy future. Its potential to reduce emissions, store energy, and provide a versatile and clean energy source makes it a key player in addressing environmental and energy challenges of our time.^201–204 However, the purification of H₂ from the common contaminants (CO₂, N₂, CH₄, and H₂O) is a requisite for the practical applications such as hydrogen fuel cells.

Ultra-thin 2D covalent organic nanosheets (CONs) were developed to fabricate H₂ purification membranes. However, reported CONs mainly have two disadvantages: (1) the relatively weak forces between inter-layers of CONs result in inefficient stacking arrangement, and (2) the comparatively large pore dimensions of CONs exceeding 0.6 nm, which significantly surpass the molecular dimensions of gaseous species (typically below 0.4 nm). In order to address these problems, Zhao et al.²⁰⁵ achieved a significant breakthrough in membrane technology by successfully fabricating ultra-thin ionic covalent organic nanosheets (iCONs) with tunable pore sizes and opposite charges via LbL assembly technique (Fig. 9a). The resulting membranes, achieved through the staggered packing of iCONs with strong electrostatic interactions, simultaneously delivered structural benefits of narrowed effective pore apertures, optimized molecular stacking patterns, and enhanced structural density, while maintaining exceptional thickness precision. These unique characteristics rendered the membranes particularly suitable for molecular-sieving-based gas separation. For the purpose of studying the relationship between structures and performances, they compared the gas separation performances of single-phase TpEBr nanosheet membrane and TpPa-SO₃Na nanosheet membrane. Due to the relatively large apertures of single-phase membranes, moderate H₂ separation factors were obtained. In sharp contrast, hybrid membranes, TpEBr@TpPa-SO₃Na with a thickness of 41 nm, showed a H₂ permeance of 2566 GPU and H₂/CO₂ SF of 22.6 at 423 K. One of the hybrid membranes, TpEBr@TpPa-SO₃Na with a thickness of 41 nm, showed a H₂ permeance of 2566 GPU and H₂/CO₂ SF of 22.6 at 423 K. This strategy provided not only a high-performance H₂ separation membrane candidate, but also an inspiration for pore engineering of COF membranes. In 2020, Sun et al.¹⁰¹ reported the first HOF membrane (UPC-HOF-6) fabricated through a solution-processing method, demonstrating unique pressure-responsive gas separation properties (Fig. 9b). At 298 K and 1.2 bar, the membrane exhibited a H₂ permeance of 461 GPU and H₂/N₂ selectivity of 11.9. Single-crystal X-ray diffraction revealed that each structural unit integrates with four adjacent building modules through twelve directional N–H⋯N hydrogen-bonding interactions originating from diaminotriazine (DAT) functional groups, thereby propagating a covalently-assembled planar architecture through coordinated non-covalent connectivity. The 2D lamellar structure manifested two distinct geometric configurations: a rhomboidal aperture with precisely quantified dimensions of 5.7 × 9.4 Å, and a non-uniform hexagonal pore architecture exhibiting reduced geometric parameters of 2.8 × 3.6 Å. The present pore sizes are suitable for separation of H₂/N₂, resulting in the high separation performance. Remarkably, when the transmembrane pressure increased from 1.2 to 2.0 bar, the H₂ permeance significantly improved to 1051 GPU while the selectivity rose to 19.5, attributed to pressure-induced framework flexibility (pore size expansion from 2.8 Å to 3.6 Å). In situ powder X-ray diffraction (PXRD) and molecular simulation confirmed the pressure-driven structural dynamic response mechanism. Additionally, the membrane displayed excellent mechanical properties (elastic modulus of 17.5 GPa) and self-healing capability (performance recovery after 15-minute DMSO vapor treatment). In 2022, Zhao et al.²⁰⁶ proposed a multi-interfacial engineering strategy, which involved the direct LbL interfacial reaction of two COFs, TpPa-SO₃H and TpTGCl with differing pore sizes (Fig. 9c). At 423 K, the fabricated 155 nm-thick ultra-thin COF membrane displayed a H₂ permeance as high as 2163 GPU and H₂/CO₂ selectivity of 26, transcending the 2008 Robeson upper bound. In 2021, Kim et al.²⁰⁷ presented a method for tuning the pore size of a MOF membrane by directly growing a ZIF-8 layer on a graphene oxide nanoribbon (GONR)-modified polymer substrate. The oxygen-containing functional groups and carbon edge structures on the GONR surface facilitated the oriented growth of ZIF-8 crystals and enhanced their rigidity, resulting in an effective pore aperture of 3.6 Å. The membrane exhibited exceptional performance in H₂ separation, achieving an ultra-high H₂ permeance of 7.6 × 10⁻⁷ mol m⁻² Pa⁻¹ s⁻¹, while maintaining high selectivity for large hydrocarbon molecules. Experimental and computational simulations revealed that the strong interactions between GONR and ZIF-8 are the key to pore tuning and performance enhancement. In 2023, Li et al.²⁰⁸ proposed a simple in situ modulation strategy by introducing different amounts of acetic acid as modulator acid into Co-gallate MOF precursor solutions (Fig. 9d). A series of Co-gallate MOF membranes (denoted as M/L-x membranes, where M/L-x represents the addition ratio of modulator acid/linker acid) were fabricated for H₂ separation (Fig. 9d). The M/L-8.3 membrane demonstrated the best balance between crystallinity and film formation with a H₂/CO₂ SF of up to 82.5 and exceptional stability. In 2023, Yang et al.²⁰⁹ presented a modular customization strategy for fabricating defect-free MOF separation membranes. By employing four structurally and functionally distinct MOFs (NH₂-Zn₂(Bim)₄, Ni(HBTC)(4,4′-bipy), MOF-74(Mg), and MIL-68(AI)), they successfully achieved rapid membrane preparation and precise performance tuning. The fabricated membranes demonstrated exceptional H₂/CO₂ separation capabilities, exhibiting a maximum SF of 1656 and H₂ permeance reaching 2745 GPU. Furthermore, post-synthetic modification techniques, including non-destructive ligand exchange, were employed to further optimize membrane performance. Notably, this versatile strategy was successfully extended to water purification applications, demonstrating its broad potential for industrial separation processes. In 2024, Jin et al.²¹⁰ reported a confined-coordination-induced intergrowth strategy for fabricating defect-free, highly connected Zr–MOF (MOF-801) membranes. By precisely regulating the coordination space (size and shape) and environment (solvents: water and DMF), the nucleation and growth rates of MOF were effectively reduced, suppressing preferential crystal formation in the bulk solution and enabling the progressive intergrowth of MOF crystals into an integrated membrane layer with minimized defects. The resulting MOF membrane features angstrom-sized lattice apertures, demonstrating exceptional separation performance for H₂ permeance of ∼1200 GPU and H₂/CO₂ selectivity of ∼67.

Fig. 9 Several porous membranes related to H₂ purification. (a) Schematic illustration of distinct assembly configurations in covalent organic nanosheet (CON) membranes. Reproduced with permission.²⁰⁵ Copyright 2020, American Chemical Society. (b) Schematic representation of UPC-HOF-6:2D layered structures and 3D framework. Reproduced with permission.¹⁰¹ Copyright 2019, Wiley-VCH. (c) Multi-interfacial design strategy enabling ultra-thin COF membranes with narrowed pores for gas separation applications. Reproduced with permission.²⁰⁶ Copyright 2022, Wiley-VCH. (d) Graphical representation of in situ strategies to optimize crystallinity and film-formation ability of MOF membranes. Reproduced with permission.²⁰⁸ Copyright 2023, Wiley-VCH.

H₂ separation membrane technology has achieved remarkable advancements in the design and fabrication strategies of porous materials, with a focus on ultra-thin structure construction, dynamic pore size regulation, and interfacial engineering optimization. Researchers have developed diverse high-performance membrane materials through methods such as LbL assembly, solution processing, and in situ growth. For instance, electrostatic stacking strategies based on iCONs enable pore aperture narrowing and high-precision membrane architectures. HOFs exhibit pressure-responsive dynamic pore size adjustments, demonstrating flexible gas permeation modulation. Interfacial engineering strategies, including COFs multilayer reactions and MOFs composite structures, enhance separation efficiency through precise pore size control and optimized crystal orientation. Furthermore, innovative approaches such as acid modulator incorporation, modular MOF customization, and confined coordination-induced growth have surpassed conventional performance limits while integrating mechanical robustness, self-healing capabilities, and industrial scalability. These breakthroughs establish a synergistic membrane material system combining high permeability and selectivity for H₂ purification, advancing the practical application of molecular sieving mechanisms in clean energy technologies.

3.4.1.3. CO₂ separation membranes. Increased CO₂ emissions contribute to major environmental issues, primarily global warming. They also pose industrial challenges, causing infrastructure corrosion and reducing the calorific values of biogas and natural gas.^211,212 Currently, CO₂ separation is primarily achieved through amine scrubbing, pressure swing adsorption, or cryogenic distillation.²¹³ Although these conventional processes definitely have their respective advantages, membrane-based gas separation technology would be a potentially energy-efficient alternative, and considerable progress has been made.²¹⁴

In 2019, Agrawal et al.²¹⁵ developed an efficient single-step post-synthetic modification protocol that could be completed within seconds at 360 °C, which dramatically enhanced the molecular sieving performance of ZIF-8 membranes. The key role of lattice stiffness on gas separation was revealed by comparing the structure and properties of ZIF-8 membranes before and after rapid heat treatment (RHT). The pristine ZIF-8 membrane has insufficient selectivity for CO₂/CH₄ and CO₂/N₂, due to the dynamic change of pore size caused by the lattice flexibility and the “gate-opening” phenomenon, whereas, after the RHT treatment, the introduction of lattice distortion and strain significantly increased the lattice stiffness and suppressed pore size fluctuations, which shifted the gas transport into a temperature activation mode, and the activation energy rose with increasing molecular size. This structural optimization resulted in the CO₂/CH₄ and CO₂/N₂ selectivities of the treated membranes exceeding 30 and a complete blocking of C₃H₆, while maintaining the crystalline integrity and the coordination environment unchanged. The results showed that the molecular sieving ability of MOF membranes can be effectively enhanced by modulating the lattice flexibility, which provided a new idea for the design of high-performance carbon capture membranes. This rapid thermal treatment process resulted in exceptional CO₂ separation capabilities, achieving selectivity exceeding 30 for both CO₂/CH₄ and CO₂/N₂ gas pairs, demonstrating outstanding potential for carbon capture applications. The order-of-magnitude enhancement in separation performance originated from lattice stiffening effects induced by precisely controlled lattice distortion and strain engineering through rapid thermal treatment in the ZIF-8 framework. In 2021, Kang et al.²¹⁶ reported a highly CO₂-selective CAU-10-H MOF membrane with ideal selectivity of 42 for CO₂/N₂ and 95 for CO₂/CH₄, coupled with an outstanding CO₂ permeability of 500 Barrer. Through integrated experimental and molecular simulation approaches, the research revealed that the membrane's superior performance originated from a unique Coulombic effect-governed transport mechanism. In spite of the larger pore-limiting diameter of Structure 2 than Structure 1, Structure 1 exhibited a higher free energy barrier for CO₂ diffusion (Fig. 10a). By comparing the structure and gas separation performance of COF membranes before and after defect engineering, this study revealed the pivotal role of defect modulation in CO₂ separation. The study revealed the Coulombic interaction was a critical factor determining the diffusion of CO₂. In 2022, Jiang et al.⁷⁷ reported a defect engineering approach to fabricate amino-functionalized COF membranes with precisely controlled missing-linker defects (Fig. 10b). By strategically incorporating monoaldehyde terminators during COF nanosheet synthesis, the researchers achieved in situ generation of abundant free amino groups (up to 38.4%) while preserving structural integrity. The resulting membranes exhibited exceptional CO₂ separation performance under humid conditions, combining high permeance (371 GPU) with remarkable selectivity (80 for CO₂/N₂ and 54 for CO₂/CH₄). The ultra-thin membrane variant (57 nm) set a new benchmark with record CO₂ permeance (1068 GPU) while maintaining practical selectivity (35 for CO₂/N₂). The amino groups facilitated CO₂ transport through dipole-quadrupole interactions and reversible carrier-mediated reactions, while the introduced defects enhanced porosity to reduce diffusion resistance. Meanwhile, the ordered channels ensured structural integrity and separation stability. In 2022, Liu et al.²¹⁷ pioneered a novel approach for fabricating defect-rich Ti–MOF (MIL-125) membranes using a titanium-oxo cluster (Ti₈Ph) as the titanium source, combined with single-mode microwave heating and tertiary growth techniques. Compared with conventional titanium sources, Ti₈Ph enabled synthesis at significantly lower temperatures (100 °C) while increasing missing-linker defects within the framework, thereby enhancing CO₂ adsorption capacity and selectivity. The resulting MIL-125 membranes demonstrated exceptional gas separation performance, achieving ideal selectivities of 38.7 for CO₂/N₂, 64.9 for H₂/N₂, and 40.7 for H₂/CH₄. Moreover, the membrane maintained stable separation performance under simulated flue gas conditions. This work provides an innovative strategy for developing high-performance MOF membranes with promising applications in CO₂ capture and H₂ purification.

Fig. 10 Several porous membranes related to CO₂ separation. (a) Schematic illustration of Structures 1 and 2 of CAU-10-H. Reproduced with permission.²¹⁶ Copyright 2021, Wiley-VCH. (b) Defect engineering strategy in COF membranes for enhanced CO₂-selective transport. Reproduced with permission.⁷⁷ Copyright 2022, Wiley-VCH. (c) Fabrication method of highly (111)-oriented ultra-thin UiO-66 membranes. Reproduced with permission.²¹⁸ Copyright 2023, Wiley-VCH. (d) Schematic illustration of the ADR method to synthesize COF membranes. Reproduced with permission.⁴⁵ Copyright 2024, Wiley-VCH.

In 2023, Liu et al.²¹⁸ highlighted the critical role of micro-structural optimization in gas separation by contrasting the structural and gas separation performance between conventional Zr–MOF membranes and the ultra-thin, highly (111)-oriented UiO-66 membranes fabricated through an innovative approach combining anisotropic etching-derived 40 nm-thick UiO-66 nanosheet seeds with confined counter-diffusion-assisted epitaxial growth (Fig. 10c). Conventional UiO-66 membranes, characterized by larger thickness (≥400 nm) and random crystallographic orientation, exhibited limited CO₂ permeance (typically <500 GPU) and low selectivity (CO₂/N₂ < 20). In contrast, the novel membrane-constructed using a 40 nm-thick triangular nanosheet seed layer via anisotropic etching, ZrS₂ as a metal source to suppress crystal twinning, and confined counter-diffusion epitaxial growth-achieved a continuous, 165 nm-thick (111)-oriented architecture. This structural refinement significantly reduced diffusion resistance, yielding a CO₂ permeance of 2070 GPU (over fourfold higher than conventional membranes) and a CO₂/N₂ selectivity of 35.4. Simultaneously, the oriented framework minimized grain boundary defects, enabling enhanced H₂/CO₂ selectivity of 41.3. Structural analyses confirmed that the ultra-thin configuration shortened mass transfer pathways, while the high orientation strengthened molecular sieving efficiency, and defect suppression ensured membrane integrity. This work demonstrates that synergistic control of membrane thickness, crystallographic orientation, and grain boundary structure can overcome the permeability-selectivity trade-off, establishing a new paradigm for designing high-performance gas separation membranes. In 2024, Zou et al.⁴⁵ presented an innovative assembly-dissociation-reconstruction (ADR) strategy for fabricating highly continuous COF membranes (Fig. 10d). The methodology involves the dissociation of parent COF into high-aspect-ratio nanosheets followed by their ordered reconstruction under water evaporation, yielding a membrane with exceptional crystallinity and open porosity. The resulting membranes demonstrated outstanding CO₂/N₂ separation performance, achieving a CO₂ permeance exceeding 1060 GPU coupled with a selectivity above 30.6. Mechanistic investigations revealed that nitrogen-rich functional groups enable selective CO₂ adsorption, while the ordered reconstruction process played a pivotal role in forming defect-free membrane architecture. The ADR approach establishes a novel paradigm for developing high-performance gas separation membranes, demonstrating significant potential for carbon capture applications.

Recent advancements in membrane-based CO₂ separation technologies have shown innovative strategies in material design and fabrication approaches. Researchers have employed diverse techniques such as rapid thermal post-synthetic modifications, defect engineering, Coulombic interaction optimization, and novel synthesis methods (e.g., microwave-assisted growth and anisotropic etching) to enhance membrane performance. Key developments include the creation of functionalized frameworks (MOFs and COFs) with tailored pore environments, defect-controlled architectures, and oriented crystalline structures, which synergistically improve CO₂ adsorption selectivity and diffusion kinetics. These membranes demonstrated exceptional gas separation efficiency through mechanisms like lattice stiffening, Coulombic transport regulation, and nitrogen-rich group-enhanced adsorption.

3.4.2. Liquid separation membranes.
3.4.2.1. Ion separation membranes. Ion separation refers to the process of separating different ions in a mixture. It mainly exploits the difference of ion charge or size to selectively penetrate small ions with low valence, while retaining large ions with high valence. Ion separation technology has been widely used in batteries, mining industry, medicine, and other fields, and has great significance for environmental protection, resource recovery, and energy utilization. Crystalline porous membranes have great application potential in ion separation due to their chemical diversity and structural designability and have been utilized for the extraction and recovery of lithium ions, as well as other noble metal ions.²¹⁹

At present, many significant breakthroughs have been realized in the separation of mono/bivalent ions using crystalline porous membranes.^220–222 It is reported that Wang's group achieved high-selectivity screening of sub-nanopores and rapid ion migration kinetics of nanoscale channels by assembling UiO-66-(COOH)₂ crystals onto PET membranes to form asymmetrically structured MOF-based subnanochannels (Asy-MOFSNCs) (Fig. 11a). By adjusting the pH value, the K⁺/Mg²⁺ selectivity of Asy-MOFSNCs can be adjusted between 10² and 10⁴, achieving pH-adjustable highly selective transport. DFT calculations were used to reveal the transport mechanism of metal ions in MOF channels, and molecular dynamics (MD) simulations were used to show that the dehydration–rehydration effect and ion-carboxyl interactions work together, resulting in a high energy barrier for divalent metal ions to pass through MOFSNCs, thereby achieving high selectivity and high transport rates.²²⁰ The incorporation of carboxylate ligands was found to effectively lower the energy barrier while simultaneously generating sub-nanometer pore architectures, resulting in membranes with significantly improved selectivity and ion transport kinetics. Subsequently, Li's group prepared an array of membranes based on MOF sub-nano channels, among which the UiO-66-(OMe)₂ membrane achieved selective discrimination between monovalent and divalent cations, with a K⁺/Mg²⁺ selectivity of up to 1567.8. In this work, by adjusting the surface functional groups (such as –NH₂, –SH, –OH, and –OMe) and pore size of MOF sub-nano channels, the ion binding affinity and dehydration kinetics were coordinately regulated. Meanwhile, the energy barrier of ion transport was adjusted, thereby affecting the ion transport rate and selectivity.²²¹ Sun's group designed and prepared COF-170/PAN membranes by using the unique properties of π-electron system, achieving selectivity of 214 and 451 for Na⁺/Al³⁺ and NO₃⁻/PO₄³⁻, respectively. At the same time, the membrane's permeation flux for NaCl reached 89.7 mmol m⁻² h⁻¹, which was higher than that of traditional reverse osmosis membranes. In this work, the molecular stacking mode was optimized to enhance the interactions between ions and π-electron system, and highly selective ion separation was achieved. At the same time, since this method does not require specific modification of the membrane, the design and preparation process of the membrane is simplified.²²² At the same time, the application of COF membranes in the separation of rare metal ions has also been reported. Wang's group prepared HBAB-TAPA-COF using an interfacial synthesis method. This material is different from the previously randomly stacked COF nanorods. The vertically oriented nanorods increased the effective contact area, promoting the adsorption of water molecules and monovalent ions and eliminating structural imperfections in the lower layer. The upper layer of the membrane was vertically arranged nanorods, and the lower layer was an ultra-thin dense layer, forming an asymmetric structure. The ultra-thin dense layer ensured high permeability and selectivity, and was mechanically protected by the substrate, which strengthened the mechanical resilience of the membrane. The selectivity of the membrane for Cs⁺/La³⁺ in a single ion system was as high as 75.9, and it was slightly reduced to 69.8 in a binary ion system, proving that crystalline porous membrane devices have great application potential in the separation of rare ions.²²³

Fig. 11 (a) Fabrication of Asy-MOFSNC using a facilitated interfacial growth strategy, MOF  = UiO-66-(COOH)₂. Reproduced with permission.²²⁰ Copyright 2020, Nature. (b) Fabrication and morphology characterization of TpPa COF nanosheets and membranes. Reproduced with permission.²²⁵ Copyright 2022, Nature. (c) IP methodology and skin-layer development process for outer-selective TFC HF NF membranes. Reproduced with permission.²³⁰ Copyright 2021, Elsevier. (d) Green method to synthesize TMC/HPEI/MIL-303/P84 membranes. Reproduced with permission.²³¹ Copyright 2024, Elsevier.

Although considerable progress has been made in the study of crystalline porous membranes for the separation of mono/bivalent ions, there are relatively few studies on the separation of mono/monovalent ions. It eagerly awaits a great breakthrough to use crystalline porous membranes for achieving efficient monovalent ion separations.^224–226 MOF, as an ordered porous material, has been used for mono/monovalent ion separation. The pore size and chemical properties of MOF will affect the selectivity of alkali metal ions. Wang's group found that ZIF-8 and UiO-66 MOF membranes with sub-nanometer pores can achieve ultra-fast and selective transport of alkali metal ions. At the same time, the ion selectivity decreased with the increase of the pore window diameter. The ZIF-8 membrane showed the best selectivity and conductivity. The ZIF-8/GO/AAO membrane prepared on the AAO support by the nanoporous GO-assisted interfacial growth method had a selectivity of 4.6, 2.2, and 1.4 for Li⁺/Rb⁺, Li⁺/K⁺, and Li⁺/Na⁺, respectively.²²⁴ At the same time, COF membranes also have great application potential in the mono/monovalent ion separation. Jiang's group used COF membranes to separate mono/monovalent ions for the first time. Their research found that the hydration layer formed by the combination of acid groups and water molecules was a key factor affecting ion transport, and increasing the thickness of the COF membrane can enhance the quantity of cascade stages, thereby further improving the separation performance (Fig. 11b). The actual selectivity of the TpPa–PO₃H₂ membranes assembled from nanosheets prepared by the oil–water–oil three-phase interfacial polymerization method for K⁺/Li⁺ binary mixtures was 4.2–4.7, and the ideal selectivity for K⁺/Li⁺ is about 13.7.²²⁵ The incorporation of phosphate groups enhanced hydration energy, leading to the formation of thicker hydration layers that reduced effective channel dimensions and increased ion transport resistance. This modification significantly improved monovalent cation selectivity while maintaining favorable hydrophilicity that facilitated water molecule transport through the COF membrane, thereby enhancing overall ion transport rates.

The application and performance of MOF and COF membranes in ion separation primarily depend on their sieving pore size and chemical properties. Ligand modification can effectively regulate the pore size and ion-binding affinity to achieve desired separation performance. Simultaneously, structural design must consider the trade-off between selectivity and permeability. Some membranes exhibit insufficient mechanical strength, requiring reinforcement through substrates or cross-linking to enhance toughness. Currently, significant progress has been made in MOF and COF membranes for ion separation. However, these membranes still demonstrate low efficiency in monovalent–monovalent ion separation.

3.4.2.2. Heavy metal ion retention membranes. The accelerated industrialization has led to the inevitable discharge of heavy metal-laden wastewater from various sectors, including metallurgical processes, chemical manufacturing, electroplating operations, and battery production. Heavy metal ions not only pollute the natural environment such as soil and aquatic systems, but also exert profound adverse impacts on the human body, leading to various diseases. Unlike other organic pollutants, heavy metal ions cannot be degraded. Therefore, heavy metal ions must be effectively collected and treated prior to their discharge into the natural environment. However, existing water treatment technology is costly and requires high energy consumption levels, which is difficult to meet human needs. Membrane separation technology offers distinct advantages including high efficiency, low energy consumption, and high separation performance, while the demand for membrane devices with heavy metal separation capability is increasing progressively.

Crystalline porous membranes exhibit significant potential for application in heavy metal wastewater treatment owing to their designable and tunable pore structures. Zhuang's group have applied COF materials to ultra-filtration membrane modification. By introducing COF into PVDF-based membrane, TbBd-modified ultra-filtration membrane was prepared by solution casting, which granted the membrane excellent hydrophilicity and anti-fouling properties. Its pure water flux can reach 123 L m⁻² h⁻¹, and the Pd²⁺ retention rate can reach 92.4%.²²⁷ Luo's group first combined sulfonated-polyethersulfone (SPES) powder with COFs to prepare COFs/SPES mixed matrix membranes. This combination utilized the strong adsorption properties of the sulfonic acid groups in SPES powder and the high chemical stability as well as large specific surface area of COFs to achieve high-capacity adsorption of uranium ions (99.4 mg g⁻¹) under strong acidic conditions (pH = 1).²²⁸ Ou's group introduced a macroporous structure by loading COF onto a biomimetic honeycomb chitosan membrane, overcoming the limitations of the stacking structure of COF itself, improving the ion diffusion rate, and thus achieving rapid adsorption. The adsorption capacities of the prepared CM@COF for Cu²⁺ and Cr⁴⁺ ions were 144 mg g⁻¹ (pH = 7) and 388 mg g⁻¹ (pH = 6), respectively.²²⁹ Su's group used an in situ construction method to directly form a 2D COFs intermediate layer on the surface of the polysulfone (PSf) hollow fiber substrate. This method can better control the growth and distribution of COFs, thereby more effectively regulating the interfacial polymerization process and constructing a thin-film nanocomposite (TFN) hollow fiber nanofiltration membrane with a thinner and less defective separation skin (Fig. 11c). In comparison to the benchmark thin-film composite (TFC) hollow fiber nanofiltration membrane, the incorporation of a COF interlayer significantly enhanced the water flux and rejection efficiency of the membrane. The rejection rates of TFC HF NF membrane for Cr₂(SO₄)₃, CuSO₄, ZnSO₄, and MnSO₄ were 95.4%, 94.3%, 91.7%, and 90.9%, respectively. The water permeation was about 60 L m⁻² h⁻¹ MPa⁻¹.²³⁰ Not only COFs, but also crystalline porous membranes with MOFs as the core have wide application potential in the separation and rejection of heavy metal ions. Recently, Kang's group combined the MOF MIL-303 with P84 copolymer polyimide to prepare MIL-303/P84 nanocomposite mixed matrix membranes (Fig. 11d). The introduction of MIL-303 provided specific channels for water molecules, thereby improving the flux of the membrane and enhancing its hydrophilicity. This study used hyperbranched polyethyleneimine (HPEI) and trimethylol chloride (TMC) to cross-link the MIL-303/P84 membrane to form an ultra-thin selective layer. This surface cross-linking technology made the surface of the membrane positively charged and enhanced its hydrophilicity, thereby improving the removal efficiency of heavy metal ions. The TMC/HPEI/MIL-303/P84 membrane had a retention rate of more than 92% for all heavy metal ions and a pure water permeability of 12.9 L m⁻² h⁻¹ bar⁻¹.²³¹ The incorporation of MIL-303 created specific channels for water molecules, reducing transport resistance and consequently enhancing water flux. Simultaneously, MIL-303 introduction modified the membrane's pore size distribution, generating narrower pores that effectively retain heavy metal ions.

3.4.2.3. Organic solvent nanofiltration (OSN) membranes. OSN represents an advanced membrane separation platform that uses nanofiltration membranes to separate molecules in organic solvents. This method synergizes nanofiltration technology with the solvent properties of organic phases, resulting in the effective and selective separation of organic molecules. Compared with traditional organic solvent separation technologies (such as distillation and evaporation), OSN technology achieves separation that can significantly reduce energy consumption and achieve efficient and precise separation of organic solvents. OSN technology serves as a vital element in the framework of green chemistry and can effectively remove impurities in organic solvents and reduce environmental pollution. It enables more environmentally benign industrial chemistry.

Membrane separation technology has broad application prospects in the field of OSN. At present, many COF membranes or COF-based composite matrix membranes with excellent OSN performance have been reported. Wang's group prepared TFPM-HZ/PAN membranes with sub-nanopores and anti-swelling properties by adjusting the interfacial crystallization method, overcoming the shortcomings of traditional OSN membranes such as low porosity, uneven pore size distribution, and easy swelling. Its rejection rates for curcumin, tetracycline, rifampicin, and vitamin B12 were 91%, 100%, 95%, and 96%, respectively (Fig. 12a). Simultaneously, using porous support materials as regulators, the controllable crystallization of 3D COF membranes under mild conditions was achieved, confirming the feasibility of preparing large-area membranes.²³² Liu's group prepared covalent triazine framework (CTF) membranes with recyclable separation performance by interfacial polymerization, overcoming the shortcomings of traditional COF membranes that are easy to contaminate. By changing the aldehyde monomer, CTF membranes with single-pore and double-pore structures can be prepared, thereby adjusting the separation performance of the membrane. CTF membranes have ultra-fast permeation and molecular-level selectivity for polar organic solvents, enabling efficient separation. The permeance of the CTF-1 membrane decreased slowly from the initial 43.3 L m⁻² h⁻¹ bar⁻¹ to 30.5 L m⁻² h⁻¹ bar⁻¹ while maintaining high rejection above 98%. And the permeance of CTF-m membranes for methylene blue (MB)/methanol decreased from the initial 107.1 L m⁻² h⁻¹ bar⁻¹ to 79.8 L m⁻² h⁻¹ bar⁻¹, while maintaining a rejection rate of more than 97%.²³³ Lai's group prepared COF_6V-72 h membranes using an electric field-assisted interfacial synthesis method (Fig. 12b). This study quantitatively analyzed the relationship between the pore structure of COF membranes and their overall performance for the first time, revealing the importance of high-density through-holes in improving membrane performance. The synthesized COF membrane had high-density, vertically arranged through-holes and exhibited ultra-high solvent permeability. COF_6V-72 h had a methanol permeance of 282.8 L m⁻² h⁻¹ bar⁻¹, which can be used for paclitaxel enrichment.²³⁴ Similarly, COF Turing membranes have also been used for organic solvent nanofiltration. Shao's group used tannic acid (TA) and amino monomers to form a composite precursor to increase reaction activity and reduce diffusion rate, thereby regulating reaction-diffusion kinetics and achieving spatiotemporal controlled growth of COF membranes (Fig. 12c). The problems of irregular stacking and grain boundary defects of COF crystal blocks in traditional methods were solved, and the precise design and preparation of COF membranes were achieved. By regulating the TA concentration, the Turing structure was formed on the surface of the COF membrane. Compared with traditional COF membranes, the Turing structure COF membrane has higher permeability and molecular sieving efficiency, achieving efficient and rapid organic solvent nanofiltration separation. The water permeance of TA/TpPa-COF membrane was as high as 49.4 L m⁻² h⁻¹ bar⁻¹, and the rejection of MB molecules (1.62 nm × 2.03 nm) was 99.2%.²³⁵ With the increasing TA concentration, the COF membrane surface morphology transitioned from smooth to labyrinthine stripes or spot patterns, forming characteristic Turing structures. These topological features significantly increased the effective separation area, thereby enhancing membrane permeation flux. Similarly, MOF membranes can also be used in organic solvent nanofiltration technology. Liu's group deposited TA and Fe²⁺ with MOFs in aqueous solution and in situ etched MOFs to form a hollow structure, thereby improving the permeance of the membrane (Fig. 12d). The hollow MOF was combined with the metal phenol network (MPN) selective layer to achieve efficient organic solvent nanofiltration. Without compromising the barrier properties enabled by the additional transport channels within the hollow MOF structure, the MPN matrix exhibited excellent interfacial adhesion with the MOF framework and the permeance of the nanocomposite membrane of different organic solvents was significantly improved. The methanol permeance of the prepared MPN/ZIF-8-2 membrane reached 24.7 L m⁻² h⁻¹ bar⁻¹, and the retention rate of Alcian blue was greater than 99%.²³⁶

Fig. 12 (a) Synthesis and characterization of TFPM-HZ. Reproduced with permission.²³² Copyright 2022, Wiley-VCH (b) demonstration of paclitaxel enrichment using COF_6v-72 h membranes. Reproduced with permission.²³⁴ Copyright 2024, American Chemical Society. (c) Schematic illustration of the synthesis and characteristic surface morphologies of the TA/TpPa-COF membranes. Reproduced with permission.²³⁵ Copyright 2024, Science (d) fabrication schematic illustration of the assembly of MPN/MOF hybrid membranes coupled with ZIF-8 dissolution process. Reproduced with permission.²³⁶ Copyright 2022, Elsevier.

In OSN, the molecular sieving performance of MOF and COF membranes is primarily governed by their pore size, with sub-1nm porous membranes demonstrating particularly promising applications. This highlights the critical importance of precise pore size control in crystalline porous materials, which can be achieved through post-synthetic modification of functional groups to enhance rejection performance. However, such improvements in separation selectivity are often accompanied by reduced permeance flux. This trade-off can be mitigated by fabricating ultra-thin membranes via IP or by modulating the IP process to create Turing-like structures, both strategies being effective for enhancing membrane permeability.

4. Conclusion and outlook

Crystalline porous materials represent a distinct category of materials that integrate long-range ordered frameworks, tunable pore structures, and programmable chemical functionality. These unique characteristics make them good candidates for constructing membrane devices. In this review, three main categories of crystalline porous membrane devices are introduced and discussed, containing MOFs, COFs, and HOFs. To apply the crystalline porous materials into membrane devices, several membrane fabrication methods have been summarized, including hot/cold pressing, in situ solvothermal growth, seed-assisted secondary growth, solution processing, interfacial polymerization, and current-driven synthesis. The constructed membrane devices show great potential in the application of LMBs, catalytic electrodes, solar cells, and separation membranes.

In LMBs, the recent progress of neat MOF, COF, MOF/polymer, and COF/polymer solid-state electrolyte membranes has been highlighted. Neat MOF electrolytes usually possess high Li⁺ transference number due to the existence of OMSs. Also, the adjustability of pore polarity of MOFs makes them a good platform to study the relationship between electrolyte structures and relationships. For neat COF electrolyte, the decoration of pore walls can enhance the Li⁺ transference number as well as ionic conductivity, which is benefited from the enhanced dissociation of lithium salts. However, the trace amount of flammable liquid electrolyte in MOFs or COFs have safety hazards during use. Meanwhile, the lithium dendrite penetration will occur in neat MOF or COF electrolytes, especially at high current densities. MOF/polymer and COF/polymer composite electrolyte membranes composed of MOF and COF as filler have good flexibility and processability. One of the main obstacles for the practical applications of these membranes is the relative low ionic conductivity of composite electrolytes. In particular, the application of composite electrolyte membranes at low temperatures is severely hindered due to the low Li⁺ conductivity. Moreover, the interactions between fillers and polymers are still ambiguous due to the complex composite systems. The ion transport mechanism in composite membranes lacks thorough and systematically studied, which is not favorable to rationally design the structures of crystalline porous materials. In order to give further insights into the Li⁺ conduction mechanism, it is necessary to combine theoretical and experimental techniques. Advanced characterization methods such as X-ray tomography microscopy (XTM), in situ/operando neutron diffraction patterns (NDP), and X-ray fine structure spectroscopy (XAFS) should be developed. The effect of the morphology of crystalline porous materials on the electrochemical performances of batteries should be intensively studied.

In separation fields, the applications of crystalline porous membranes including MOFs, COFs, and HOFs are discussed. Beneficial to the abundant pore structures and designable pore micro-environment, these membranes show great potential in gas separation (such as hydrocarbon separation, H₂ purification, and CO₂ separation) and liquid separation (such as ion separation, heavy metal ion retention, and organic solvent nanofiltration). However, several main challenges have remained which severely hinder the industrial application of them. Firstly, the fabricated polycrystalline membranes show inevitable grain boundary defects, which are harmful to the separation performance. To fabricate separation membranes with no grain boundary defects, glassy membranes can be constructed which have long-range disordered and short-range ordered micro-structures and maintain the pore channels. The unique structures of glassy membranes make them good candidates for separation fields which can conquer the trade-off effect between permeability and selectivity. Secondly, the mass transfer mechanism of crystalline porous membranes still remains ambiguous, lacking deep understanding of the interactions between the transfer molecules and pore walls. It is of vital importance to develop in situ characterization techniques as well as theoretical models to systematically reveal the transfer mechanisms. Thirdly, gas separation and nanofiltration membranes require crystalline porous materials with pore sizes smaller than 1 nm. Such COFs with pore size smaller than 1 nm can be prepared by adopting short organic linkers which are limited now. In order to obtain COFs with narrow pore sizes, post-modification can be used by anchoring side groups in the inner walls. Finally, the emergence of artificial intelligence (AI) should be deeply integrated with materials and membranes science, which can accelerate the discovery, design, and optimization of porous crystalline membrane devices, thereby addressing critical global climate challenges related to carbon emissions.

Moreover, to satisfy the practical application of crystalline porous membrane devices, the scalability, cost-effectiveness, and industrial feasibility of these membranes should be considered: (1) for MOF-based membranes, several MOF materials such as ZIF-8, UiO-66, and HKUST-1 have been industrially synthesized at relatively low cost, which are widely applied in the field of SSEs, catalytic electrode membranes, and separation membranes. It is noteworthy that seed-assisted secondary growth technique has emerged as one of the most promising candidates for achieving scalable fabrication of defect-free crystalline porous membranes with tailored functionalities. In MOF membrane fabrication via seed-assisted secondary growth, the decoupling of nucleation and crystal growth stages simplifies the synthetic pathway, thereby enhancing preparation reproducibility. (2) Currently, the industrialization of COFs remains in its infancy, facing numerous challenges in scaling up from laboratory to industrial production. Most COF monomers require complex synthetic procedures and high production costs, making large-scale manufacturing economically unfeasible. Zhang's group developed a solvent-free molten salt synthesis enabling kilogram-scale COF production, showing promise for C₂H₄ separation and providing a fundamental platform for economical manufacturing of high-efficiency COF membranes, which have potential to fabricate COF membranes for industrial applications.²³⁷ (3) The practical adoption of HOF-based membranes in industrial applications encounters formidable obstacles. Up to date, the use of HOF-based membranes under actual environment still lacks. Hydrogen-bonding is the main interaction between HOF motifs, which is weaker than coordinate bonds and covalent bonds in MOFs and COFs, so the long-term stability of HOF-based membrane devices will face more serious challenges than that of MOF or COF-based membrane devices, which will affect the industrial application of them. Additionally, the high production costs and complex, labor-intensive synthesis processes present substantial barriers to industrial-scale manufacturing.

Data availability

The authors confirm that the data are available within the article or from the corresponding author B. L. upon reasonable request.

Conflicts of interest

The authors declare no conflict of interest.

Acknowledgements

This work was financially supported by the National Natural Science Foundation of China (22405133, 22035003, 22494631, and 22494633), and Fundamental Research Funds for the Central Universities (63243119).

Notes and references

1. Z. Zhuo, E. Du, N. Zhang, C. P. Nielsen, X. Lu, J. Xiao, J. Wu and C. Kang, Nat. Commun., 2022, 13, 3172.
2. Z. Liu, R. Lin, M. Wei, M. Yin, P. Wu, M. Li, L. Li, Y. Wang, G. Chen, V. Carnevali, L. Agosta, V. Slama, N. Lempesis, Z. Wang, M. Wang, Y. Deng, H. Luo, H. Gao, U. Rothlisberger, S. M. Zakeeruddin, X. Luo, Y. Liu, M. Grätzel and H. Tan, Nat. Mater., 2025, 24, 252–259.
3. C. Li, J. Song, H. Lai, H. Zhang, R. Zhou, J. Xu, H. Huang, L. Liu, J. Gao, Y. Li, M. H. Jee, Z. Zheng, S. Liu, J. Yan, X.-K. Chen, Z. Tang, C. Zhang, H. Y. Woo, F. He, F. Gao, H. Yan and Y. Sun, Nat. Mater., 2025, 24, 433–443.
4. Q. Xu, T. Li, Z. Ju, G. Chen, D. Ye, G. I. N. Waterhouse, Y. Lu, X. Lai, G. Zhou, L. Guo, K. Yan, X. Tao, H. Li and Y. Qiu, Nature, 2025, 637, 339–346.
5. P. Ding, H. Yuan, L. Xu, L. Wu, H. Du, S. Zhao, D. Yu, Z. Qin, H. Liu, Y. Li, X. Zhang, H. Yu, M. Tang, Y. Ren, L. Li and C. W. Nan, Adv. Mater., 2024, 37, 2413654.
6. Y. Wang, R. Lim, K. Larson, A. Knab, D. Fontecha, S. Caverly, J. Song, C. Park, P. Albertus, G. W. Rubloff, S. B. Lee and A. C. Kozen, ChemSusChem, 2024, 17, e202400718.
7. J. Zhao, R. Urrego-Ortiz, N. Liao, F. Calle-Vallejo and J. Luo, Nat. Commun., 2024, 15, 6391.
8. X. Zheng, X. Zheng, M. Gao, Y. Liu, H. Pan and W. Sun, Angew. Chem., Int. Ed., 2025, e202422062,  DOI:10.1002/anie.202422062.
9. H. Li, L. Fang, T. Wang, R. Bai, J. Zhang, T. Li, Z. Duan, K. J. Chen and F. Pan, Adv. Mater., 2024, 37, 2416337.
10. T. Hou, J. Zhu, H. Gu, X. Li, Y. Sun, Z. Hua, R. Shao, C. Chen, B. Hu, L. Mai, S. Chen, D. Wang and J. Zhang, Angew. Chem., Int. Ed., 2025, 64(15), e202424749.
11. Z. Yang, Y. Belmabkhout, L. N. McHugh, D. Ao, Y. Sun, S. Li, Z. Qiao, T. D. Bennett, M. D. Guiver and C. Zhong, Nat. Mater., 2023, 22, 888–894.
12. Q. Liu, Y. Miao, L. F. Villalobos, S. Li, H.-Y. Chi, C. Chen, M. T. Vahdat, S. Song, D. J. Babu, J. Hao, Y. Han, M. Tsapatsis and K. V. Agrawal, Nat. Mater., 2023, 22, 1387–1393.
13. B. Wang, W. t Zhao, X. Xu, C. Zhang, S. Y. Ding, Y. Zhang and T. Wang, Adv. Mater., 2024, 36, 2309572.
14. X. Tian, L. Cao, K. Zhang, R. Zhang, X. Li, C. Yin and S. Wang, Angew. Chem., Int. Ed., 2024, 64, e202416864.
15. S. Yu, S. Yang, M. Yang, J. Yang, Z. Song, D. Hu, H. Ji, Z. Jia and M. Liu, Angew. Chem., Int. Ed., 2025, 64, e202420086.
16. X. Dong, X. Ai, W. He, Y. Zhai, R. Guo, Y.-W. Li, Z.-Q. Ma, Y. Yang and K.-G. Zhou, ACS Nano, 2025, 19, 4233–4241.
17. Y. Su, K.-I. Otake, J.-J. Zheng, S. Horike, S. Kitagawa and C. Gu, Nature, 2022, 611, 289–294.
18. P. Ajayan, W. Wang, Y. Chen, X. Bu and P. Feng, Adv. Mater., 2024, 36, 2408042.
19. Y. Liang, G. Xie, K. K. Liu, M. Jin, Y. Chen, X. Yang, Z. J. Guan, H. Xing and Y. Fang, Angew. Chem., Int. Ed., 2024, 64, e202416884.
20. S. Hong, M. Di Vincenzo, A. Tiraferri, E. Bertozzi, R. Górecki, B. Davaasuren, X. Li and S. P. Nunes, Nat. Commun., 2024, 15, 3160.
21. J. Han, H. Wu, H. Fan, L. Ding, G. Hai, J. Caro and H. Wang, J. Am. Chem. Soc., 2023, 145, 14793–14801.
22. X. Liu, Y. Ye, X. He, Q. Niu, B. Chen and Z. Li, Angew. Chem., Int. Ed., 2024, 63, e202400195.
23. P. Zhang, X. Zou, J. Song, Y. Tian, Y. Zhu, G. Yu, Y. Yuan and G. Zhu, Adv. Mater., 2020, 32, 1907449.
24. S. Bai, B. Kim, C. Kim, O. Tamwattana, H. Park, J. Kim, D. Lee and K. Kang, Nat. Nanotechnol., 2020, 16, 77–84.
25. B. T. Liu, X. H. Pan, D. Y. Nie, X. J. Hu, E. P. Liu and T. F. Liu, Adv. Mater., 2020, 32, 2005912.
26. B. Liang, H. Wang, X. Shi, B. Shen, X. He, Z. A. Ghazi, N. A. Khan, H. Sin, A. M. Khattak, L. Li and Z. Tang, Nat. Chem., 2018, 10, 961–967.
27. C. Li, A. Qi, Y. Ling, Y. Tao, Y.-B. Zhang and T. Li, Sci. Adv., 2023, 9, eadf5087.
28. M. Zhang, X. Jing, S. Zhao, P. Shao, Y. Zhang, S. Yuan, Y. Li, C. Gu, X. Wang, Y. Ye, X. Feng and B. Wang, Angew. Chem., Int. Ed., 2019, 58, 8768–8772.
29. S. Yan, H. Liu, Y. Lu, Q. Feng, H. Zhou, Y. Wu, W. Hou, Y. Xia, H. Zhou, P. Zhou, X. Song, Y. Ou and K. Liu, Sci. Adv., 2025, 11, eads4014.
30. J. Zhu, P. Bian, G. Sun, J. Zhang, G. Lou, X. Song, R. Zhao, J. Liu, N. Xu, A. Li, X. Wan, Y. Ma, C. Li, H. Zhang and Y. Chen, Angew. Chem., Int. Ed., 2025, e202424685,  DOI:10.1002/anie.202424685.
31. Y. L. Liao, X. L. Wang, H. Yuan, Y. J. Li, C. M. Xu, S. Li, J. K. Hu, S. J. Yang, F. Deng, J. Liu and J. Q. Huang, Adv. Mater., 2025, 37, 2419782.
32. R. Liao, C. Li, M. Zhou, R. Liu, S. Liu and D. Wu, Chem. Sci., 2024, 15, 18327–18334.
33. Y. Li, S. Ma, Y. Zhao, S. Chen, T. Xiao, H. Yin, H. Song, X. Pan, L. Cong and H. Xie, Energy Environ. Mater., 2023, 7, e12648.
34. X. Q. Cheng, Z. X. Wang, X. Jiang, T. Li, C. H. Lau, Z. Guo, J. Ma and L. Shao, Prog. Mater. Sci., 2018, 92, 258–283.
35. L. Hu, V. T. Bui, S. Fan, W. Guo, S. Pal, Y. Ding and H. Lin, J. Mater. Chem. A, 2022, 10, 10872–10879.
36. M. Shan, X. Liu, X. Wang, Z. Liu, H. Iziyi, S. Ganapathy, J. Gascon and F. Kapteijn, J. Mater. Chem. A, 2019, 7, 8929–8937.
37. S. He, B. Zhu, X. Jiang, G. Han, S. Li, C. H. Lau, Y. Wu, Y. Zhang and L. Shao, Proc. Natl. Acad. Sci. U. S. A., 2021, 119, e2114964119.
38. D. Peng, X. Feng, G. Yang, X. Niu, Z. Liu and Y. Zhang, J. Membr. Sci., 2023, 668, 121267.
39. P. Wang, Y. Peng, C. Zhu, R. Yao, H. Song, L. Kun and W. Yang, Angew. Chem., Int. Ed., 2021, 60, 19047–19052.
40. Z. Yang, Y. Zhang, W. Wu, Z. Zhou, H. Gao, J. Wang and Z. Jiang, J. Membr. Sci., 2022, 664, 121118.
41. S. Zhao, C. Jiang, J. Fan, S. Hong, P. Mei, R. Yao, Y. Liu, S. Zhang, H. Li, H. Zhang, C. Sun, Z. Guo, P. Shao, Y. Zhu, J. Zhang, L. Guo, Y. Ma, J. Zhang, X. Feng, F. Wang, H. Wu and B. Wang, Nat. Mater., 2021, 20, 1551–1558.
42. R. H. Choi, J. So, Y. Kim, D. Lee and H. R. Byon, ACS Energy Lett., 2024, 9, 5341–5348.
43. J.-H. Lee, H. Lee, J. Lee, T. W. Kang, J. H. Park, J.-H. Shin, H. Lee, D. Majhi, S. U. Lee and J.-H. Kim, ACS Nano, 2023, 17, 17372–17382.
44. K. J. Nam, A. M. O. Mohamed, J. Seong, H. An, D. Y. Kang, I. G. Economou and J. S. Lee, Small, 2025, 2409515,  DOI:10.1002/smll.202409515.
45. H. Zhang, T. Shao, Z. Cheng, J. Dong, Z. Wang, H. Jiang, X. Zhao, T. X. Liu, G. Zhu and X. Zou, Angew. Chem., Int. Ed., 2024, 63, e202411724.
46. L. Huang, Z. Liu, G. Gao, C. Chen, Y. Xue, J. Zhao, Q. Lei, M. Jin, C. Zhu, Y. Han, J. S. Francisco and X. Lu, J. Am. Chem. Soc., 2023, 145, 26444–26451.
47. Y. Zhang, X. Zhang, J. Lyu, K.-I. Otake, X. Wang, L. R. Redfern, C. D. Malliakas, Z. Li, T. Islamoglu, B. Wang and O. K. Farha, J. Am. Chem. Soc., 2018, 140, 11179–11183.
48. W. Gong, P. Gao, Y. Gao, Y. Xie, J. Zhang, X. Tang, K. Wang, X. Wang, X. Han, Z. Chen, J. Dong and Y. Cui, J. Am. Chem. Soc., 2024, 146, 31807–31815.
49. F. Lang, L. Zhang, Y. Li, X. J. Xi, J. Pang, W. Zheng, H. C. Zhou and X. H. Bu, Angew. Chem., Int. Ed., 2025, e202422517.
50. Y. Y. Liu, J. R. Huang, H. L. Zhu, P. Q. Liao and X. M. Chen, Angew. Chem., Int. Ed., 2023, 62, e202311265.
51. H. S. Cho, H. Deng, K. Miyasaka, Z. Dong, M. Cho, A. V. Neimark, J. K. Kang, O. M. Yaghi and O. Terasaki, Nature, 2015, 527, 503–507.
52. M. Barsukova, A. Sapianik, V. Guillerm, A. Shkurenko, A. C. Shaikh, P. Parvatkar, P. M. Bhatt, M. Bonneau, A. Alhaji, O. Shekhah, S. R. G. Balestra, R. Semino, G. Maurin and M. Eddaoudi, Nat. Synth., 2024, 3, 33–46.
53. S. M. Wang, M. Shivanna, S. T. Zheng, T. Pham, K. A. Forrest, Q. Y. Yang, Q. Guan, B. Space, S. Kitagawa and M. J. Zaworotko, J. Am. Chem. Soc., 2024, 146, 4153–4161.
54. W. Chen, S. Li, L. Yi, Z. Chen, Z. Li, Y. Wu, W. Yan, F. Deng and H. Deng, J. Am. Chem. Soc., 2024, 146, 12215–12224.
55. Y. Huang, Y. Feng, Y. Li, K. Tan, J. Tang, J. Bai and J. Duan, Angew. Chem., Int. Ed., 2024, 63, e202403421.
56. W. Zhang, H. Jiang, Y. Liu, Y. Hu, A. S. Palakkal, Y. Zhou, M. Sun, E. Du, W. Gong, Q. Zhang, J. Jiang, J. Dong, Y. Liu, D. Li, Y. Zhu, Y. Cui and X. Duan, Nature, 2025, 638, 418–424.
57. H. Fang, X. Y. Liu, H. J. Ding, M. Mulcair, B. Space, H. Huang, X. W. Li, S. M. Zhang, M. H. Yu, Z. Chang and X. H. Bu, J. Am. Chem. Soc., 2024, 146, 14357–14367.
58. Z. Lv, R. Lin, Y. Yang, K. Lan, C. T. Hung, P. Zhang, J. Wang, W. Zhou, Z. Zhao, Z. Wang, J. Zou, T. Wang, T. Zhao, Y. Xu, D. Chao, W. Tan, B. Yan, Q. Li, D. Zhao and X. Li, Nat. Chem., 2025, 17, 177–185.
59. Z. Han, K. Y. Wang, R. R. Liang, Y. Guo, Y. Yang, M. Wang, Y. Mao, J. Huo, W. Shi and H. C. Zhou, J. Am. Chem. Soc., 2025, 147, 3866–3873.
60. H. Huo, B. Wu, T. Zhang, X. Zheng, L. Ge, T. Xu, X. Guo and X. Sun, Energy Storage Mater., 2019, 18, 59–67.
61. N. Angulakshmi, Y. Zhou, S. Suriyakumar, R. B. Dhanalakshmi, M. Satishrajan, S. Alwarappan, M. H. Alkordi and A. M. Stephan, ACS Omega, 2020, 5, 7885–7894.
62. Z. Zhang, J. H. You, S. J. Zhang, C. W. Wang, Y. Zhou, J. T. Li, L. Huang and S. G. Sun, ChemElectroChem, 2020, 7, 1125–1134.
63. H. Wang, S. Duan, Y. Zheng, L. Qian, C. Liao, L. Dong, H. Guo, C. Ma, W. Yan and J. Zhang, eTransportation, 2024, 20, 100311.
64. A. Sabetghadam, B. Seoane, D. Keskin, N. Duim, T. Rodenas, S. Shahid, S. Sorribas, C. L. Guillouzer, G. Clet, C. Tellez, M. Daturi, J. Coronas, F. Kapteijn and J. Gascon, Adv. Funct. Mater., 2016, 26, 3154–3163.
65. S. Biswas and P. Van Der Voort, Eur. J. Inorg. Chem., 2013, 2154–2160.
66. Z. Zhou, T. Ma, H. Zhang, S. Chheda, H. Li, K. Wang, S. Ehrling, R. Giovine, C. Li, A. H. Alawadhi, M. M. Abduljawad, M. O. Alawad, L. Gagliardi, J. Sauer and O. M. Yaghi, Nature, 2024, 635, 96–101.
67. H. Li, Z. Zhou, T. Ma, K. Wang, H. Zhang, A. H. Alawadhi and O. M. Yaghi, J. Am. Chem. Soc., 2024, 146, 35486–35492.
68. L. B. Xing, K. Cheng, H. Li, K. Niu, T. X. Luan, S. Kong, W. W. Yu, P. Z. Li and Y. Zhao, Angew. Chem., Int. Ed., 2025, 64(16), e202425668.
69. W. Zhang, Y. Zhang, W. Ma, X. Han, W. Gong, Y. Liu and Y. Cui, Chem, 2025, 11, 102398.
70. S. Li, S. Xu, E. Lin, T. Wang, H. Yang, J. Han, Y. Zhao, Q. Xue, P. Samorì, Z. Zhang and T. Zhang, Nat. Chem., 2025, 17, 226–232.
71. J. Chang, Z. Zhang, H. Zheng, H. Li, J. Suo, C. Ji, F. Chen, S. Zhang, Z. Wang, V. Valtchev, S. Qiu, J. Sun and Q. Fang, Nat. Chem., 2025, 17, 571–581.
72. Y. Chen, R. Liu, Y. Guo, G. Wu, T. C. Sum, S. W. Yang and D. Jiang, Nat. Synth., 2024, 3, 998–1010.
73. J. Han, J. Feng, J. Kang, J. M. Chen, X. Y. Du, S. Y. Ding, L. Liang and W. Wang, Science, 2024, 383, 1014–1019.
74. Z. Wang, Y. Zhang, J. Liu, Y. Chen, P. Cheng and Z. Zhang, Nat. Protoc., 2024, 19, 3489–3519.
75. Z. Guo, Y. Zhang, Y. Dong, J. Li, S. Li, P. Shao, X. Feng and B. Wang, J. Am. Chem. Soc., 2019, 141, 1923–1927.
76. Q. Xu, S. Tao, Q. Jiang and D. Jiang, Angew. Chem., Int. Ed., 2020, 59, 4557–4563.
77. Z. Guo, H. Wu, Y. Chen, S. Zhu, H. Jiang, S. Song, Y. Ren, Y. Wang, X. Liang, G. He, Y. Li and Z. Jiang, Angew. Chem., Int. Ed., 2022, 61, e202210466.
78. E. T. C. Vogt and B. M. Weckhuysen, Nature, 2024, 629, 295–306.
79. W. K. Qin, D. H. Si, Q. Yin, X. Y. Gao, Q. Q. Huang, Y. N. Feng, L. Xie, S. Zhang, X. S. Huang, T. F. Liu and R. Cao, Angew. Chem., Int. Ed., 2022, 61, e202202089.
80. K. Ma, P. Li, J. H. Xin, Y. Chen, Z. Chen, S. Goswami, X. Liu, S. Kato, H. Chen, X. Zhang, J. Bai, M. C. Wasson, R. R. Maldonado, R. Q. Snurr and O. K. Farha, Cell Rep. Phys. Sci., 2020, 1, 100024.
81. C. Jiang, J. X. Wang, D. Liu, E. Wu, X. W. Gu, X. Zhang, B. Li, B. Chen and G. Qian, Angew. Chem., Int. Ed., 2024, 63, e202404734.
82. W. Wang, Y. Shi, W. Chai, K. Wing, K. Tang, I. Pyatnitskiy, Y. Xie, X. Liu, W. He, J. Jeong, J. C. Hsieh, A. R. Lozano, B. Artman, X. Shi, N. Hoefer, B. Shrestha, N. B. Stern, W. Zhou, D. W. McComb, T. Porter, G. Henkelman, B. Chen and H. Wang, Nature, 2025, 638, 401–410.
83. L. Li, X. Zhang, X. Lian, L. Zhang, Z. Zhang, X. Liu, T. He, B. Li, B. Chen and X. H. Bu, Nat. Chem., 2025, 17, 727–733.
84. Y. Yang, L. Li, R.-B. Lin, Y. Ye, Z. Yao, L. Yang, F. Xiang, S. Chen, Z. Zhang, S. Xiang and B. Chen, Nat. Chem., 2021, 13, 933–939.
85. B. Yu, S. Geng, H. Wang, W. Zhou, Z. Zhang, B. Chen and J. Jiang, Angew. Chem., Int. Ed., 2021, 60, 25942–25948.
86. Y. Zhou, C. Chen, R. Krishna, Z. Ji, D. Yuan and M. Wu, Angew. Chem., Int. Ed., 2023, 62, e202305041.
87. C. Chen, L. Shen, H. Lin, D. Zhao, B. Li and B. Chen, Chem. Soc. Rev., 2024, 53, 2738–2760.
88. X. Shi, W. Gu, B. Zhang, Y. Zhao, A. Zhang, W. Xiao, S. Wei and H. Pang, Adv. Funct. Mater., 2025, 2423760.
89. W. Wang, D. Chen, F. Li, X. Xiao and Q. Xu, Chem, 2024, 10, 86–133.
90. S. Yuan, X. Li, J. Zhu, G. Zhang, P. V. Puyveldeb and B. V. Bruggen, Chem. Soc. Rev., 2019, 48, 2665–2681.
91. S. Dong, Y. Xue, S. Li, Y. Zhang, S. Wang and Z. Lin, TrAC, Trends Anal. Chem., 2025, 185, 118167.
92. Y. Cheng, O. Shekhah, S. Datta, S. Zhou and M. Eddaoudi, Chem. Soc. Rev., 2022, 51, 8300–8350.
93. C. Zhang, F. Li, T. Gu, X. Song, J. Yuan, L. Ouyang, M. Zhu and J. Liu, Prog. Mater. Sci., 2025, 152, 101455.
94. L. Chen, Y. Li, S.-P. Li, L.-Z. Fan, C.-W. Nan and J. B. Goodenough, Nano Energy, 2018, 46, 176–184.
95. J. F. Wu and X. Guo, Small, 2019, 15, 1804413.
96. Y. Liu, Z. Ng, E. A. Khan, H.-K. Jeong, C.-B. Ching and Z. Lai, Microporous Mesoporous Mater., 2009, 118, 296–301.
97. A. Huang, W. Dou and J. Caro, J. Am. Chem. Soc., 2010, 132, 15562–15564.
98. H. Fan, J. Gu, H. Meng, A. Knebel and J. Caro, Angew. Chem., Int. Ed., 2018, 57, 4083–4087.
99. Y. Pu, M. Zhao, X. Liang, S. Wang, H. Wang, Z. Zhu, Y. Ren, Z. Zhang, G. He, D. Zhao and Z. Jiang, Angew. Chem., Int. Ed., 2023, 62, e202302355.
100. S. Chen, M. Wahiduzzaman, T. Ji, Y. Liu, Y. Li, C. Wang, Y. Sun, G. He, G. Maurin, S. Wang and Y. Liu, Angew. Chem., Int. Ed., 2024, 64, e202413701.
101. S. Feng, Y. Shang, Z. Wang, Z. Kang, R. Wang, J. Jiang, L. Fan, W. Fan, Z. Liu, G. Kong, Y. Feng, S. Hu, H. Guo and D. Sun, Angew. Chem., 2020, 132, 3868–3873.
102. K. Dey, M. Pal, K. C. Rout, S. Kunjattu H, A. Das, R. Mukherjee, U. K. Kharul and R. Banerjee, J. Am. Chem. Soc., 2017, 139, 13083–13091.
103. S. Hou, G. Zhang, Z. Qiao, Y. Bai, H. Di, Y. Hua, T. Hao and H. Xu, Angew. Chem., Int. Ed., 2025, 64, e202421555.
104. Y. Li, L. H. Wee, A. Volodin, J. A. Martens and I. F. J. Vankelecom, Chem. Commun., 2015, 51, 918–920.
105. Y. Feng, Z. Kang, Z. Wang, Z. Liu, Q. J. Niu, W. Fan, L. Qiao, J. Pang, H. Chang, X. Cui, L. Fan, H. Guo, R. Wang, D. Zhao and D. Sun, J. Am. Chem. Soc., 2024, 146, 33452–33460.
106. S. Zhou, Y. Wei, L. Li, Y. Duan, Q. Hou, L. Zhang, L.-X. Ding, J. Xue, H. Wang and J. Caro, Sci. Adv., 2018, 4, eaau1393.
107. S. Zhou, O. Shekhah, J. Jia, J. Czaban-Jóźwiak, P. M. Bhatt, A. Ramírez, J. Gascon and M. Eddaoudi, Nat. Energy, 2021, 6, 882–891.
108. S. Zhou, O. Shekhah, A. Ramírez, P. Lyu, E. Abou-Hamad, J. Jia, J. Li, P. M. Bhatt, Z. Huang, H. Jiang, T. Jin, G. Maurin, J. Gascon and M. Eddaoudi, Nature, 2022, 606, 706–712.
109. D. Xie, Z. Wang, X. Ma, Y. Feng, X. Tang, Q. Gu, Y. Deng and P. Gao, Energy Environ. Mater., 2024, 7, e12746.
110. F. Wang, H. Liu, Y. Guo, Q. Han, P. Lou, L. Li, J. Jiang, S. Cheng and Y. Cao, Energy Environ. Mater., 2023, 7, e12497.
111. M. Burton, S. Narayanan, B. Jagger, L. F. Olbrich, S. Dhir, M. Shibata, M. J. Lain, R. Astbury, N. Butcher, M. Copley, T. Kotaka, Y. Aihara and M. Pasta, Nat. Energy, 2024, 10, 135–147.
112. N. Piao, J. Wang, X. Gao, R. Li, H. Zhang, G. Hu, Z. Sun, X. Fan, H.-M. Cheng and F. Li, J. Am. Chem. Soc., 2024, 146, 18281–18291.
113. Y. Wang, Y. Ni, S. Xu, Y. Lu, L. Shang, Z. Yang, K. Zhang, Z. Yan, W. Xie and J. Chen, J. Am. Chem. Soc., 2025, 147(12), 10772–10783.
114. S. Jiang, T. Lv, Y. Peng and H. Pang, Adv. Sci., 2023, 10, 2206887.
115. R. Zhao, Y. Wu, Z. Liang, L. Gao, W. Xia, Y. Zhao and R. Zou, Energy Environ. Sci., 2020, 13, 2386–2403.
116. B. M. Wiers, M.-L. Foo, N. P. Balsara and J. R. Long, J. Am. Chem. Soc., 2011, 133, 14522–14525.
117. S. S. Park, Y. Tulchinsky and M. Dincă, J. Am. Chem. Soc., 2017, 139, 13260–13263.
118. E. M. Miner, S. S. Park and M. Dincă, J. Am. Chem. Soc., 2019, 141, 4422–4427.
119. L. Shen, H. B. Wu, F. Liu, J. L. Brosmer, G. Shen, X. Wang, J. I. Zink, Q. Xiao, M. Cai, G. Wang, Y. Lu and B. Dunn, Adv. Mater., 2018, 30, 1747476.
120. S. Yuan, J. L. Bao, J. Wei, Y. Xia, D. G. Truhlar and Y. Wang, Energy Environ. Sci., 2019, 12, 2741–2750.
121. S. Ma, L. Shen, Q. Liu, W. Shi, C. Zhang, F. Liu, J. A. Baucom, D. Zhang, H. Yue, H. B. Wu and Y. Lu, ACS Appl. Mater. Interfaces, 2020, 12, 43824–43832.
122. H. Liu, H. Pan, M. Yan, X. Zhang and Y. Jiang, Adv. Mater., 2023, 35, 2300888.
123. Y. Ouyang, W. Gong, Q. Zhang, J. Wang, S. Guo, Y. Xiao, D. Li, C. Wang, X. Sun, C. Wang and S. Huang, Adv. Mater., 2023, 35, 2304685.
124. A. E. Abdelmaoula, J. Shu, Y. Cheng, L. Xu, G. Zhang, Y. Xia, M. Tahir, P. Wu and L. Mai, Small Methods, 2021, 5, 2100508.
125. Z. Wang, R. Tan, H. Wang, L. Yang, J. Hu, H. Chen and F. Pan, Adv. Mater., 2017, 30, 1704436.
126. X. Qi, D. Cai, X. Wang, X. Xia, C. Gu and J. Tu, ACS Appl. Mater. Interfaces, 2022, 14, 6859–6868.
127. P. Dong, X. Zhang, W. Hiscox, J. Liu, J. Zamora, X. Li, M. Su, Q. Zhang, X. Guo, J. McCloy and M. K. Song, Adv. Mater., 2023, 35, 2211841.
128. Z. Li, S. Wang, J. Shi, Y. Liu, S. Zheng, H. Zou, Y. Chen, W. Kuang, K. Ding, L. Chen, Y.-Q. Lan, Y.-P. Cai and Q. Zheng, Energy Storage Mater., 2022, 47, 262–270.
129. T. Wang, X. Zhang, N. Yuan and C. Sun, Chem. Eng. J., 2023, 451, 138819.
130. P. Chiochan, X. Yu, M. Sawangphruk and A. Manthiram, Adv. Energy Mater., 2020, 10, 2001285.
131. Q. Zeng, J. Wang, X. Li, Y. Ouyang, W. He, D. Li, S. Guo, Y. Xiao, H. Deng, W. Gong, Q. Zhang and S. Huang, ACS Energy Lett., 2021, 6, 2434–2441.
132. Y. Zhou, J. Chen, J. Sun and T. Zhao, Nano Lett., 2024, 24, 2033–2040.
133. Y. Wang, C. J. Zanelotti, X. Wang, R. Kerr, L. Jin, W. H. Kan, T. J. Dingemans, M. Forsyth and L. A. Madsen, Nat. Mater., 2021, 20, 1255–1263.
134. J. Mindemark, M. J. Lacey, T. Bowden and D. Brandell, Prog. Polym. Sci., 2018, 81, 114–143.
135. F. Wu, K. Zhang, Y. Liu, H. Gao, Y. Bai, X. Wang and C. Wu, Energy Storage Mater., 2020, 33, 26–54.
136. S. Randau, D. A. Weber, O. Kötz, R. Koerver, P. Braun, A. Weber, E. Ivers-Tiffée, T. Adermann, J. Kulisch, W. G. Zeier, F. H. Richter and J. Janek, Nat. Energy, 2020, 5, 259–270.
137. D. M. Reinoso and M. A. Frechero, Energy Storage Mater., 2022, 52, 430–464.
138. S. Choudhury, Z. Tu, A. Nijamudheen, M. J. Zachman, S. Stalin, Y. Deng, Q. Zhao, D. Vu, L. F. Kourkoutis, J. L. Mendoza-Cortes and L. A. Archer, Nat. Commun., 2019, 10, 3091.
139. Z. Lei, J. Shen, W. Zhang, Q. Wang, J. Wang, Y. Deng and C. Wang, Nano Res., 2020, 13, 2259–2267.
140. M. C. Nguyen, H. L. Nguyen, T. P. M. Duong, S. H. Kim, J. Y. Kim, J. H. Bae, H. K. Kim, S. N. Lim and W. Ahn, Adv. Funct. Mater., 2024, 34, 2406987.
141. H. Lian, R. Momen, Y. Xiao, B. Song, X. Hu, F. Zhu, H. Liu, L. Xu, W. Deng, H. Hou, G. Zou and X. Ji, Adv. Funct. Mater., 2023, 33, 2306060.
142. L. Xu, X. Xiao, H. Tu, F. Zhu, J. Wang, H. Liu, W. Huang, W. Deng, H. Hou, T. Liu, X. Ji, K. Amine and G. Zou, Adv. Mater., 2023, 35, 2303193.
143. K. He, S. H. S. Cheng, J. Hu, Y. Zhang, H. Yang, Y. Liu, W. Liao, D. Chen, C. Liao, X. Cheng, Z. Lu, J. He, J. Tang, R. K. Y. Li and C. Liu, Angew. Chem., Int. Ed., 2021, 60, 12116–12123.
144. S. Liu, W. Liu, D. Ba, Y. Zhao, Y. Ye, Y. Li and J. Liu, Adv. Mater., 2022, 35, 2110423.
145. S. Li, F. Pei, Y. Ding, X. Guo, X. Zhang, H. Tao, Z. He, H. Hu and L. Zhang, Adv. Funct. Mater., 2024, 35, 2415495.
146. Y. Ma, Y. Qiu, K. Yang, S. Lv, Y. Li, X. An, G. Xiao, Z. Han, Y. Ma, L. Chen, D. Zhang, W. Lv, Y. Tian, T. Hou, M. Liu, Z. Zhou, F. Kang and Y.-B. He, Energy Environ. Sci., 2024, 17, 8274–8283.
147. Q. Xu, S. Tao, Q. Jiang and D. Jiang, J. Am. Chem. Soc., 2018, 140, 7429–7432.
148. G. Zhang, Y.-l Hong, Y. Nishiyama, S. Bai, S. Kitagawa and S. Horike, J. Am. Chem. Soc., 2018, 141, 1227–1234.
149. G. Zhao, H. Ma, C. Zhang, Y. Yang, S. Yu, H. Zhu, Y. Sun and H. Guo, Nano-Micro Lett., 2024, 17, 21.
150. C. Niu, W. Luo, C. Dai, C. Yu and Y. Xu, Angew. Chem., Int. Ed., 2021, 60, 24915–24923.
151. T. Liu, Y. Zhong, Z. Yan, B. He, T. Liu, Z. Ling, B. Li, X. Liu, J. Zhu, L. Jiang, X. Gao, R. Zhang, J. Zhang, B. Xu and G. Zhang, Angew. Chem., Int. Ed., 2024, 63, e202411535.
152. K. Jeong, S. Park, G. Y. Jung, S. H. Kim, Y.-H. Lee, S. K. Kwak and S.-Y. Lee, J. Am. Chem. Soc., 2019, 141, 5880–5885.
153. J. Li, F.-Q. Zhang, F. Li, Z. Wu, C. Ma, Q. Xu, P. Wang and X.-M. Zhang, Chem. Commun., 2020, 56, 2747–2750.
154. X. Li, Q. Hou, W. Huang, H.-S. Xu, X. Wang, W. Yu, R. Li, K. Zhang, L. Wang, Z. Chen, K. Xie and K. P. Loh, ACS Energy Lett., 2020, 5, 3498–3506.
155. X. Chen, L. Peng, L. Wang, J. Yang, Z. Hao, J. Xiang, K. Yuan, Y. Huang, B. Shan, L. Yuan and J. Xie, Nat. Commun., 2019, 10, 1021.
156. G. Homann, L. Stolz, J. Nair, I. C. Laskovic, M. Winter and J. Kasnatscheew, Sci. Rep., 2020, 10, 4390.
157. Y. Zhang, H. Liu, F. Liu, S. Zhang, M. Zhou, Y. Liao, Y. Wei, W. Dong, T. Li, C. Liu, Q. Liu, H. Xu, G. Sun, Z. Wang, Y. Ren and J. Yang, ACS Nano, 2025, 19, 3197–3209.
158. Z. Liu, K. Zhang, G. Huang, S. Bian, Y. Huang, X. Jiang, Y. Pan, Y. Wang, X. Xia, B. Xu and G. Zhang, Chem. Eng. J., 2022, 433, 133550.
159. D. Cheng, C. Sun, Z. Lang, J. Zhang, A. Hu, J. Duan, X. Chen, H.-Y. Zang, J. Chen, M. Zheng and Q. Dong, Cell Rep. Phys. Sci., 2022, 3, 100731.
160. A. Hanan, M. N. Lakhan, F. Bibi, A. Khan, I. A. Soomro, A. Hussain and U. Aftab, Chem. Eng. J., 2024, 482, 148776.
161. T. Liu, P. Li, N. Yao, G. Cheng, S. Chen, W. Luo and Y. Yin, Angew. Chem., Int. Ed., 2019, 58, 4679–4684.
162. L. Deng, F. Hu, M. Ma, S. C. Huang, Y. Xiong, H. Y. Chen, L. Li and S. Peng, Angew. Chem., Int. Ed., 2021, 60, 22276–22282.
163. J. Cheng, X. Yang, X. Yang, R. Xia, Y. Xu, W. Sun and J. Zhou, Fuel Process. Technol., 2022, 229, 107174.
164. Z.-X. Cai, Y. Xia, Y. Ito, M. Ohtani, H. Sakamoto, A. Ito, Y. Bai, Z.-L. Wang, Y. Yamauchi and T. Fujita, ACS Nano, 2022, 16, 20851–20864.
165. N. Q. Tran, Q. M. Le, T. T. N. Tran, T.-K. Truong, J. Yu, L. Peng, T. A. Le, T. L. H. Doan and T. B. Phan, ACS Appl. Mater. Interfaces, 2024, 16, 28625–28637.
166. S. Li, Z. Zhou, G. Liu, Q. Zhang, Y. Gao, H. Zhu and S. Zhu, Chem. Eng. J., 2024, 493, 152820.
167. W. Li, S. Mukherjee, B. Ren, R. Cao and R. A. Fischer, Adv. Energy Mater., 2021, 12, 2003499.
168. I. Hod, M. D. Sampson, P. Deria, C. P. Kubiak, O. K. Farha and J. T. Hupp, ACS Catal., 2015, 5, 6302–6309.
169. L. Ye, J. Liu, Y. Gao, C. Gong, M. Addicoat, T. Heine, C. Woll and L. Sun, J. Mater. Chem. A, 2016, 4, 15320–15326.
170. C. S. Diercks, S. Lin, N. Kornienko, E. A. Kapustin, E. M. Nichols, C. Zhu, Y. Zhao, C. J. Chang and O. M. Yaghi, J. Am. Chem. Soc., 2018, 140, 1116–1122.
171. B.-X. Dong, S.-L. Qian, F.-Y. Bu, Y.-C. Wu, L.-G. Feng, Y.-L. Teng, W.-L. Liu and Z.-W. Li, ACS Appl. Energy Mater., 2018, 1, 4662–4669.
172. S. Lin, C. S. Diercks, Y.-B. Zhang, N. Kornienko, E. M. Nichols, Y. Zhao, A. R. Paris, D. Kim, P. Yang, O. M. Yaghi and C. J. Chang, Science, 2015, 349, 1208–1213.
173. P. L. Cheung, S. K. Lee and C. P. Kubiak, Chem. Mater., 2019, 31, 1908–1919.
174. N. Kornienko, Y. Zhao, C. S. Kley, C. Zhu, D. Kim, S. Lin, C. J. Chang, O. M. Yaghi and P. Yang, J. Am. Chem. Soc., 2015, 137, 14129–14135.
175. P. De Luna, W. Liang, A. Mallick, O. Shekhah, F. P. García de Arquer, A. H. Proppe, P. Todorović, S. O. Kelley, E. H. Sargent and M. Eddaoudi, ACS Appl. Mater. Interfaces, 2018, 10, 31225–31232.
176. C.-W. Kung, C. O. Audu, A. W. Peters, H. Noh, O. K. Farha and J. T. Hupp, ACS Energy Lett., 2017, 2, 2394–2401.
177. J.-X. Wu, W.-W. Yuan, M. Xu and Z.-Y. Gu, Chem. Commun., 2019, 55, 11634–11637.
178. Y. T. Guntern, J. R. Pankhurst, J. Vávra, M. Mensi, V. Mantella, P. Schouwink and R. Buonsanti, Angew. Chem., Int. Ed., 2019, 58, 12632–12639.
179. X. Kang, B. Wang, K. Hu, K. Lyu, X. Han, B. F. Spencer, M. D. Frogley, F. Tuna, E. J. L. McInnes, R. A. W. Dryfe, B. Han, S. Yang and M. Schröder, J. Am. Chem. Soc., 2020, 142, 17384–17392.
180. J. Cao, G. Tang and F. Yan, Adv. Energy Mater., 2024, 14, 2304027.
181. D. Inácio, A. L. Pinto, A. B. Paninho, L. C. Branco, S. K. S. Freitas and H. Cruz, Nanomater, 2023, 13, 1204–1220.
182. Y. Li, A. Pang, C. Wang and M. Wei, J. Mater. Chem., 2011, 21, 17259.
183. R. Tang, Z. Xie, S. Zhou, Y. Zhang, Z. Yuan, L. Zhang and L. Yin, ACS Appl. Mater. Interfaces, 2016, 8, 22201–22212.
184. F. Li, X. Huang, Y. Li, X. Du, E. Yang, Y. Ahn, B. R. Lee, B. Wu and S. H. Park, Adv. Funct. Mater., 2024, 34, 2409811.
185. M. A. Gordillo, D. K. Panda and S. Saha, ACS Appl. Mater. Interfaces, 2018, 11, 3196–3206.
186. N. Wang, S.-Q. Ye, M. Li, H. Pan and G.-W. Wang, Solar RRL, 2023, 8, 2300790.
187. S. Wang, T. Wu, J. Guo, R. Zhao, Y. Hua and Y. Zhao, ACS Cent. Sci., 2024, 10, 1383–1395.
188. W. Xing, P. Ye, J. Lu, X. Wu, Y. Chen, T. Zhu, A. Peng and H. Huang, J. Power Sources, 2018, 401, 13–19.
189. A. S. A. Ahmed, W. Xiang, I. Saana Amiinu and X. Zhao, New J. Chem., 2018, 42, 17303–17310.
190. Y. He, Z. Zhang, W. Wang and L. Fu, J. Alloys Compd., 2020, 825, 154089.
191. Y. He and W. Wang, J. Solid State Chem., 2021, 296, 121992.
192. J. Wang, J. Zhang, S. Gai, W. Wang, Y. Dong, B. Hu, J. Li, K. Lin, D. Xia, R. Fan and Y. Yang, Adv. Funct. Mater., 2022, 32, 2203898.
193. Y. Liu, Z. Chen, G. Liu, Y. Belmabkhout, K. Adil, M. Eddaoudi and W. Koros, Adv. Mater., 2019, 31, 1807513.
194. J. Wang, Y. Wang, Y. Liu, H. Wu, M. Zhao, Y. Ren, Y. Pu, W. Li, S. Wang, S. Song, X. Liang, G. He, Y. Han and Z. Jiang, Adv. Funct. Mater., 2022, 32, 2208064.
195. S. Li, W. Han, Q. F. An, K. T. Yong and M. J. Yin, Adv. Funct. Mater., 2023, 33, 2303447.
196. S. Kim, H. Wang and Y. M. Lee, Angew. Chem., Int. Ed., 2019, 58, 17512–17527.
197. B. Hosseini Monjezi, K. Kutonova, M. Tsotsalas, S. Henke and A. Knebel, Angew. Chem., Int. Ed., 2021, 60, 15153–15164.
198. Z. Qiao, Y. Liang, Z. Zhang, D. Mei, Z. Wang, M. D. Guiver and C. Zhong, Adv. Mater., 2020, 32, 2002165.
199. Y. Sun, S. Hu, J. Yan, T. Ji, L. Liu, M. Wu, X. Guo and Y. Liu, Angew. Chem., Int. Ed., 2023, 62, e202311336.
200. Y. Liang, Z. Zhang, A. Chen, C. Yu, Y. Sun, J. Du, Z. Qiao, Z. Wang, M. D. Guiver and C. Zhong, Angew. Chem., Int. Ed., 2024, 63, e202404058.
201. X. Ma, X. Wu, J. Caro and A. Huang, Angew. Chem., Int. Ed., 2019, 58, 16156–16160.
202. B. Ghalei, K. Wakimoto, C. Y. Wu, A. P. Isfahani, T. Yamamoto, K. Sakurai, M. Higuchi, B. K. Chang, S. Kitagawa and E. Sivaniah, Angew. Chem., Int. Ed., 2019, 58, 19034–19040.
203. P. Nian, H. Liu and X. Zhang, J. Membr. Sci., 2019, 573, 200–209.
204. Y. Shen, F. Jiang, Q. Liu, Z. Chen, K. Ge, J. Bai and J. Duan, Adv. Membr., 2025, 5, 100136.
205. Y. Ying, M. Tong, S. Ning, S. K. Ravi, S. B. Peh, S. C. Tan, S. J. Pennycook and D. Zhao, J. Am. Chem. Soc., 2020, 142, 4472–4480.
206. Y. Ying, S. B. Peh, H. Yang, Z. Yang and D. Zhao, Adv. Mater., 2022, 34, 202104946.
207. E. Choi, S. J. Hong, Y. J. Kim, S. E. Choi, Y. Choi, J. H. Kim, J. Kang, O. Kwon, K. Eum, B. Han and D. W. Kim, Adv. Funct. Mater., 2021, 31, 2011146.
208. Y. Liu, H. Chen, T. Li, Y. Ren, H. Wang, Z. Song, J. Li, Q. Zhao, J. Li and L. Li, Angew. Chem., Int. Ed., 2023, 62, e202309095.
209. L. Shu, Y. Peng, H. Song, C. Zhu and W. Yang, Angew. Chem., Int. Ed., 2023, 62, e202315057.
210. G. Liu, B. Mo, Y. Guo, Z. Chu, X. M. Ren, K. Guan, R. Miao, Z. Wang, Y. Zhang, W. Ji, G. Liu, H. Matsuyama and W. Jin, Angew. Chem., Int. Ed., 2024, 63, e202405676.
211. R. Sahoo, S. Mondal, D. Mukherjee and M. C. Das, Adv. Funct. Mater., 2022, 32, 2207197.
212. C. Ma and J. J. Urban, Adv. Funct. Mater., 2019, 29, 1903243.
213. R. Thür, N. Van Velthoven, S. Slootmaekers, J. Didden, R. Verbeke, S. Smolders, M. Dickmann, W. Egger, D. De Vos and I. F. J. Vankelecom, J. Membr. Sci., 2019, 576, 78–87.
214. H. Guo, J. Wei, Y. Ma, J. Deng, S. Yi, B. Wang, L. Deng, X. Jiang and Z. Dai, Adv. Membr., 2022, 2, 100040.
215. D. J. Babu, G. He, J. Hao, M. T. Vahdat, P. A. Schouwink, M. Mensi and K. V. Agrawal, Adv. Mater., 2019, 31, 1900855.
216. D. S. Chiou, H. J. Yu, T. H. Hung, Q. Lyu, C. K. Chang, J. S. Lee, L. C. Lin and D. Y. Kang, Adv. Funct. Mater., 2021, 31, 2006924.
217. C. Wang, Y. Sun, L. Li, R. Krishna, T. Ji, S. Chen, J. Yan and Y. Liu, Angew. Chem., Int. Ed., 2022, 61, e202203663.
218. Y. Sun, J. Yan, Y. Gao, T. Ji, S. Chen, C. Wang, P. Lu, Y. Li and Y. Liu, Angew. Chem., Int. Ed., 2023, 62, e202216697.
219. H. Dou, M. Xu, B. Wang, Z. Zhang, G. Wen, Y. Zheng, D. Luo, L. Zhao, A. Yu, L. Zhang, Z. Jiang and Z. Chen, Chem. Soc. Rev., 2021, 50, 986–1029.
220. J. Lu, H. Zhang, J. Hou, X. Li, X. Hu, Y. Hu, C. D. Easton, Q. Li, C. Sun, A. W. Thornton, M. R. Hill, X. Zhang, G. Jiang, J. Z. Liu, A. J. Hill, B. D. Freeman, L. Jiang and H. Wang, Nat. Mater., 2020, 19, 767–774.
221. R.-J. Mo, S. Chen, L.-Q. Huang, X.-L. Ding, S. Rafique, X.-H. Xia and Z.-Q. Li, Nat. Commun., 2024, 15, 2145.
222. Q.-W. Meng, J. Li, Z. Lai, W. Xian, S. Wang, F. Chen, Z. Dai, L. Zhang, H. Yin, S. Ma and Q. Sun, Sci. Adv., 2024, 10, eado8658.
223. Q. Liu, M. Liu, Z. Zhang, C. Yin, J. Long, M. Wei and Y. Wang, Nat. Commun., 2024, 15, 9221.
224. H. Zhang, J. Hou, Y. Hu, P. Wang, R. Ou, L. Jiang, J. Z. Liu, B. D. Freeman, A. J. Hill and H. Wang, Sci. Adv., 2018, 4, eaaq0066.
225. H. Wang, Y. Zhai, Y. Li, Y. Cao, B. Shi, R. Li, Z. Zhu, H. Jiang, Z. Guo, M. Wang, L. Chen, Y. Liu, K.-G. Zhou, F. Pan and Z. Jiang, Nat. Commun., 2022, 13, 7123.
226. R. Li, B. Lu, Z. Xie and J. Zhai, Small, 2019, 15, 1904866.
227. W. Xu, X. Sun, M. Huang, X. Pan, X. Huang and H. Zhuang, J. Environ. Manage., 2020, 269, 110758.
228. G. Wu, Y. Liu, Q. Zheng, Z. Yu and F. Luo, J. Solid State Chem., 2020, 288, 121364.
229. L. Zhang, Y. Li, Y. Wang, S. Ma, J. Ou, Y. Shen, M. Ye and H. Uyama, J. Hazard. Mater., 2021, 407, 124390.
230. Y. Jiang, S. Li, J. Su, X. Lv, S. Liu and B. Su, J. Membr. Sci., 2021, 635, 119523.
231. N. K. Hundessa, C.-C. Hu, D.-Y. Kang, E. G. Ajebe, B. A. Habet, W.-S. Hung, K.-R. Lee and J.-Y. Lai, J. Hazard. Mater., 2024, 480, 136221.
232. X. Shi, Z. Zhang, C. Yin, X. Zhang, J. Long, Z. Zhang and Y. Wang, Angew. Chem., Int. Ed., 2022, 61, e202207559.
233. G. Li, Y. Liu, Z. He, K. Shi and F. Liu, Chem. Eng. J., 2024, 484, 149488.
234. L. Cao, C. Chen, S. An, T. Xu, X. Liu, Z. Li, I. C. Chen, J. Miao, G. Li, Y. Han and Z. Lai, J. Am. Chem. Soc., 2024, 146, 21989–21998.
235. F. Yang, J. Guo, C. Han, J. Huang, Z. Zhou, S.-P. Sun, Y. Zhang and L. Shao, Sci. Adv., 2024, 10, eadr9260.
236. Y. Chen, Y. Bai, L. Meng, W. Zhang, J. Xia, Z. Xu, R. Sun, Y. Lv and T. Liu, Chem. Eng. J., 2022, 437, 135289.
237. K. Wang, X. Qiao, H. Ren, Y. Chen and Z. Zhang, J. Am. Chem. Soc., 2025, 147, 8063–8082.

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