Nanostructured Ion Conductor Design for Highly Efficient Ion Transport

ABSTRACT

As a principal component of ion conducting devices (ICDs), solid‐state ion conductors (SSICs) possess distinctive merits, including no leakage, low flammability, no volatility, high thermal stability and electrochemical stability, offering a path to realize devices with high safety and high specific performance. Nonetheless, the rapid expansion of SSICs‐based ICDs is largely restricted owing to the low ionic conductivity and poor interface contact behavior of SSICs. One practicable solution, albeit challenging, is utilizing SSICs with nano morphology because of the fast ion transport rate in nanochannels and the tight interface contact of nanomaterials. This paper provides an overview of the new advances in developing nano ion conductors with high ionic conductivity. We describe the typical ion transport mechanisms in nanochannels as well as the development and regulation methods of nano ion conductors, with special emphasis on the influence of nanostructure on the ion transport of integrated ion conductors. Moreover, we also discuss the representative challenges in designing nano ion conductors with high ionic conductivity from personal perspectives, and afford some possible design directions for constructing exceedingly good nano ion conductors for current iontronic devices.

This review summarizes and evaluates the nano ion conductors and their specific regulation, and presents the potential development trends for nano ion conductors, which may inspire the design of nano ion conductors and promote the expansion of the ion conductor family.

1. Introduction

Ion conducting devices (ICDs) [1, 2, 3], especially for solid‐state energy storage devices [4, 5, 6] and energy conversion devices [7, 8], will develop into indispensable necessities in the foreseeable future, ranging from small‐scale electronic and transportation equipment to large scale energy generation and storage applications [7, 9, 10]. As an important component of these devices, solid‐state ion conductors (SSICs) act as a bridge to transport ions in their internal zones. Normally, the traditional SSICs critically rely on their inherent physicochemical properties for ion transport [11, 12], such as inorganic ion conductors that utilize active sites on the crystal structure to transport ions, and organic ion conductors that use polymer segments to transport ions. There ion transport processes, which depends on the intrinsic structure of the materials, are often suffered from significant ionic migration barriers [13, 14], which seriously affects the energy release capacity and energy release rate of ICDs. Therefore, the contemporary researches on traditional SSICs are mainly focused on modifying their internal structure and/or composition to reduce the ionic migration energy barrier [15, 16, 17].

Although the performance of traditional SSICs has been dramatically improved through the unremitting efforts of researchers, the development of them has encountered bottlenecks due to the difficulty of modifying their internal properties (e.g., crystal defect sites and polymeric organic fragments) [18, 19, 20]. More recently, given the numerous advantages of nanomaterials, the applications of nanotechnology in various fields have made noteworthy breakthroughs [21, 22, 23, 24, 25]. Therefore, the nanosizing of SSICs is also an important strategy to break through the low ionic conductivity of bulk materials. Generally speaking, the nanosizing of SSICs refers to the design of SSICs with special nanostructures, which can obviously overcome the ionic migration resistance on the thermodynamics and kinetics of the traditional materials by forming the ion transport channels on the nanoscale. Depending on the constituent elements, the design of SSICs with nanostructures, namely the construction of nano ion conductors, can be divided into devising inorganic nano ion conductors, organic nano ion conductors, and organic‐inorganic hybrid nano ion conductors (Figure 1). Inorganic nano ion conductors are chiefly the two‐dimensional (2D) ionic nanosheets with some layered structures, which usually have large interlaminar channels for cation and OH⁻ transport [26, 27, 28]. Organic nano ion conductors mainly refer to the 2D covalent organic frameworks (2D COFs) with crystalline thin sheets and three‐dimensional (3D) COFs with stereoscopic morphology, their organic units can be covalently linked to form periodic interconnecting channels to accommodate and transport various ions [29, 30, 31]. Whereas the organic‐inorganic hybrid nano ion conductors mainly include the 2D/3D metal‐organic frameworks (MOFs), which serves as coordination compound that can form a rich pore structure and large surface to transport ions [32, 33, 34].

FIGURE 1.

The types of nano ion conductors, including inorganic nano ion conductors (2D inorganic nanosheets), organic nano ion conductors (2D COFs and 3D COFs) and organic‐inorganic hybrid nano‐ion conductors (2D MOFs and 3D MOFs).

Benefiting from the unique ion transport mode generated by the nanosizing effect, nano ion conductors exhibit different ion transport mechanisms (e. g., Grotthuss mechanism and guest solvent transport mechanism), as compared to traditional SSICs [28, 35]. In addition, nano ion conductors can be regarded as a coveted ion conduction material due to the following merits. First, the large and ordered nanochannels inside the nano ion conductors can serve as rapid ion transport paths, thereby avoiding the thermodynamic and kinetic barriers of ion migration in traditional SSICs [27, 36]. More importantly, the nanosizing of ion conductors is capable of weakening the adverse ion transport restrictions at low temperatures, which is important for the stable operation of devices in unconventional environments [37]. Second, the ion transport channels of nano ion conductors have a certain flexibility, such superiority is favor for introducing some functional molecules and/or ions to tailor their specific ion transport properties [38, 39]. Meanwhile, the nanometer characteristics of ion conductors endows them with excellent internal contact modes and interface contact behaviors, which is significant for the fast output of the performance of ICDs. Third, the majority of nano ion conductors have stronger mechanical stability and mechanical flexibility than the traditional counterparts, which could avoid the mechanical failure of ion conductors [40, 41, 42]. At last, most of the nano ion conductors feature good liquid dispersion and liquid insolubility as well as the suitable liquid electrolyte compatibility [43, 44]; these characteristics are convenient for the construction of nano ion conductor‐based hybrid materials. Therefore, developing and applying nano ion conductors will bring new opportunities for ICDs, especially solid‐state energy storage devices.

Although many classical works of nanomaterial‐based ion conductors have been reported in previous years, a specialized paper related to the summarization of ion conduction performance in nano ion conductors is very rare. In this review paper, we first present the ion transport mechanisms of nano ion conductors. Thereafter, we introduce the design strategies of nano ion conductors, mainly including designing 2D inorganic ion transport nanosheets, nano COFs (N‐COFs) and nano MOFs (N‐MOFs), and meanwhile, describe the relationship between ionic conductivity and internal structures. Subsequently, we explored the current challenges in the field of nano ion conductors. Finally, we provide some insights for further designing nano ion conductors.

2. Ion Transport Mechanism in Nano Ion Conductors

As the name implies, the ion transport mechanism is an important intrinsic factor that determines the ion transport performance. Therefore, designing nano ion conductors with high ion mobility must first clarify their specific ion transport mechanisms. In traditional SSICs, ion migration on inorganic crystal defects/interstitials and organic polymer segments are their main ion transport mechanisms (Figure 2a,b), which makes ion transport face huge energy barriers, so that the traditional SSICs usually exhibit the low ionic conductivity [6, 15, 19, 20]. However, in nano ion conductors, ions are rapidly transported in the directional nanochannels with low energy barriers, thus achieving high ionic conductivity. Specifically, the ion transport mechanisms in nano ion conductors mainly include the interstitial hopping mechanism, the defect transport mechanism, the Grotthuss mechanism, the vehicle mechanism, the Lewis acid‐base mechanism, and the guest solvent transport mechanism.

FIGURE 2.

The ion transport mechanisms of traditional SSICs and nano ion conductors. (a) The ion transport mechanism of organic polymer. (b) The ion transport mechanism of inorganic materials. (c) The vehicle mechanism of nano ion conductors. (d) The Grotthuss mechanism of nano ion conductors. (e) The Lewis acid‐base mechanism of nano ion conductors. (f) The guest solvent transport mechanism of nano ion conductors.

Interstitial hopping mechanism. Normally, ion transport through the interstices of traditional SSICs encounters huge energy barriers. However, when some traditional SSICs are nanosized, the internal strain energy of the material and the ion migration energy in the corresponding interstitial sites are significantly reduced. Therefore, ions can rapidly hop between neighbouring interstices [45, 46]. It should be noted that only a few traditional SSICs can be easily nanosized. Therefore, this interstitial hopping mechanism only occurs in a very few inorganic nano ion conductors.

Defect transport mechanism. Similar to traditional SSICs, nano ion conductors are not perfect crystals, they also contain some defects, which can act as Lewis acid sites or ion capture sites for ion transport [47, 48, 49]. The most typical case of this mechanism is that N‐MOFs uses defects for H⁺ transport, but given the few defect sites inside the N‐MOFs, H⁺ cannot be transported rapidly. Recent findings have demonstrated that high‐entropy laminates (Li_x(Fe_1/5Co_1/5Ni_1/5Mn_1/5Zn_1/5)PS₃) with a substantial amount of cation vacancies can facilitate the Li⁺ transport [42]. This suggests that when the vacancies within nanomaterials reach a certain threshold, they can spontaneously form continuous paths for cation transport, thereby enabling nano ion conductors to exhibit enhanced ionic conductivity.

Vehicle mechanism. The vehicle mechanism typically involves the diffusion of hydrated H⁺ in a certain direction in the hydrophilic/solvophilic channel of nanomaterials [50]. However, the proton conductivity under this mechanism is much lower than that under the Grotthuss mechanism in nano ion conductors [51, 52] (Figure 2c). But for Li⁺ transport in nano ion conductors, nanoconfinement can alleviate the slow ion transport under vehicle mechanisms through open metal sites, pore size confinements and anisotropic nanochannels [53, 54]. First, open metal sites snatch the ion/solvent in a solvation sheath, which reduces the ion diffusion barriers. Second, pore size confinement reorganizes the liquid electrolytes into a superionic phase, which accelerates the ion transport. Third, interlayer anisotropic nanochannels guide the directional ion diffusion. Therefore, vehicle transport under nanoconfinement is sometimes not an inherent ion transport limitation, but a tunable mechanism controlled by pore chemistry, channel geometry, and host guest interactions.

Grotthuss mechanism. Grotthuss mechanism is a H⁺ hopping behavior that relies on the hydrogen bonding networks formed by water molecules and protons [55]. By analogy, the Grotthuss mechanism resembles a Newton's cradle, in which the H⁺ bonded to the water molecules are transported from one water molecule to another, and during the transport process, H⁺ will dissociate from the previously bonded water molecule and combine with the water molecule in front along the direction of motion (Figure 2d). Furthermore, the energy required to continuously form hydrogen bonds and break hydrogen bonds is relatively low, so the H⁺ transport by the Grotthuss mechanism has ultra‐high proton conductivity [56]. This explains why the proton transport mechanism of most currently designed ion conductors belongs to Grotthuss behavior. In nano ion conductors, nanoconfinement has fundamentally altered the Grotthuss mechanism. First, the hydrogen bonding network is no longer an isotropic bulk medium, but is templated by functional groups on pore walls/interlayer walls. For example, the spacing between ‐SO₃H groups determines that whether protons remain in isolated H₃O⁺ domains or penetrate through interconnected hydrogen bonding networks in sulfonated COFs [41]. Second, the hopping barriers not only depend on the water‐water interaction, but also rely on the skeleton electrostatics (e.g., electron withdrawing substituents) [33, 34], as the latter weakens the proton binding and reduces the activation energy. These two factors under nanoconfinement jointly establish that the Grotthuss mechanism is not only a reduced version of bulk proton hopping, but a chemically tunable phenomenon controlled by the interaction of pore wall/interlayer functionalization, electrostatic environment, and spatial confinement. This acts as the principle for supporting the rational design of high‐performance nano ion conductors.

Lewis acid‐base mechanism. Lewis acid‐base mechanism describes the interaction of anions‐cations or polar functional groups‐ions in directional nanochannels, it is mainly used for the ion transport under the solvent‐free conditions [13, 30] (Figure 2e). Specifically, for the 2D clay mineral nanosheets [57, 58], anionic N‐COFs and anionic N‐MOFs [13, 32], the ion bonds between negatively charged framework and interlayer/pore wall cation can be broken under the Lewis acid‐base dissociation mechanism, allowing the transport of alkali metal ions. However, these nano ion conductors typically exhibit low ionic conductivity due to the high ion dissociation energy. For non‐ionic N‐COFs/MOFs with polar functional groups and cationic N‐COFs/MOF, the polar functional groups or cations on their frameworks produce Lewis acid‐base binding effects with guest anions, thus promoting the dissociation and transport of guest alkali metal ions or H⁺ [13, 32]. Therefore, when designing these types of nano ion conductors, significant attention should be paid to the number and type of polar functional groups/cations in the nanochannels.

Guest solvent transport mechanism. Guest solvent transport mechanism refers to the ion transport process of solvents as carriers (Figure 2f). The triggering of this mechanism requires the nanochannels of nano ion conductors to be filled with guest solvents. These solvents, as carriers of cations, can weaken the interactions associated with cations and form a continuous nanofluidic for cation transport. The most important application examples of the guest solvent transport mechanism are the aforementioned nano ion conductors that utilize the Lewis acid‐base mechanism for alkali metal ion transport. When organic solvents fill their nanochannels [57, 59, 60], the ion pairing behavior between the cations and anions, or the cation solvation process, is weakened, which generates a substantial number of relatively free cations, and induces a cation hopping behavior with low ion traction in nanochannels, thereby optimizing the ion transport behavior [61, 62, 63]. Recently studies have found that some inorganic ion conductors (e.g., 2D CdSP₃‐Li [40, 64] and 2D Li_0.8Sn_0.8S₂ [65]) that can be swollen by water also have the guest solvent transport mechanism, the underlying reason is that the nano‐confined water molecules form a continuous water network, which facilitates the transport of insufficiently solvated cations with a lower energy barrier. While the guest solvent transport mechanism effectively facilitates ion transport, it is imperative to consider both the size matching effect of the solvent and nanochannels, as well as the cation solvation within the nanochannels when selecting the guest solvents.

In fact, exploring the ion transport mechanisms in nano ion conductors is a complex and challenging process, on the one hand, there may be a variety of ion transport mechanisms in a specific ion transport process, and these ion transport mechanisms are complementary. On the other hand, given the difficulties associated with characterizing the ion transport in nanochannels, many ion transport behaviors may remain undiscovered. Therefore, the future study of the ion transport mechanisms in nano ion conductors should combine with more advanced characterization techniques and complex theoretical calculation models.

3. Design of Nano Ion Conductors

Currently, the low ionic conductivity of polymer and inorganic ion conductors affects the performance of ICDs, especially for the fast output of the device performance, which conflicts with the fact that the usage of electrified equipment is increasing sharply and the corresponding ICDs require to output high performance in a short period of time [12, 21]. Therefore, the design of SSICs with fast ion transport is vital for developing high‐performance ICDs and meeting the demands of daily life. Compared to the traditional SSICs mentioned above, nano ion conductors can form ion channels on the nanoscale to greatly improve the ion transport performance. Apart from that, nanostructure characteristics endow nano ion conductors with better liquid compatibility and mechanical properties (Figure 3). Given these advantages, tremendous efforts have been devoted to developing and designing nano ion conductors. Current strategies to design nano ion conductors can be categorized into three types: (1) Developing inorganic nano ion conductors with different compositions and structures, namely the design of inorganic nano ion conductors. (2) Developing organic nano ion conductors with different units and structures, that is the design of organic nano ion conductors. (3) Preparing various organic‐inorganic complex‐based nano ion conductors with different ligands and structures, i.e., the design of organic‐inorganic hybrid nano ion conductors. In addition, the physicochemical properties of the nanochannels in these nano ion conductors are somewhat adjustable, so the micro‐level modification of nano ion conductors provides a strong possibility to optimize their ion transport performance.

FIGURE 3.

The requirements and advantages of nano ion conductors.

3.1. Design of Inorganic Nano Ion Conductors

Inorganic nano ion conductors are typically 2D inorganic materials with ion conduction function, which exhibit an exceptionally high horizontal‐to‐vertical ratio. In simple terms, 2D inorganic materials possess extremely large lateral dimensions and extremely small longitudinal dimensions. Their lateral dimensions can typically reach hundreds of nanometers or even several micrometers, while their thickness is only a single or several atomic thicknesses (less than 5 nanometers) [66, 67, 68]. Owing to this unique aspect ratio, 2D inorganic nanosheets can be stacked between layers to form regular and flexible interlayer nanochannels [27]. Unlike traditional SSICs, which rely on defect sites or interval sites to transport ions, the interlayer nanochannels of 2D inorganic nanosheets feature large dimensions and weak interaction forces, which ensures the rapid and efficient ion transport [28]. In addition, compared with traditional SSICs, the high aspect ratio of 2D inorganic nanosheets also endows nano ion conductors with greater flexibility, which can broaden their application fields. Furthermore, 2D inorganic nanosheets inherit some advantages of traditional SSICs, including the non‐flammability, good chemical stability, and good thermal stability, which guarantees that they can transport ions under certain harsh environments [69]. Based on the current design strategies, the design of 2D inorganic nanosheets with superior ion transport performance can be divided into two categories. (1) The development of 2D inorganic nanosheets with ion transport function, which mainly involves exfoliating bulk materials with layered structures to achieve ion conduction. (2) The optimization of ion transport performance of the developed 2D inorganic nanosheets, such a modification method is originated from the highly controllable and functional surface exposure area of 2D inorganic nanosheets.

3.1.1. Development of 2D Inorganic Nanosheets

In the past few decades or even centuries, some materials with 2D nanometer properties in crystal structure have been predicted, but it was not until 2004 that the unprecedented peeling of graphene made researchers realize that such materials could be nanostructured through exfoliation [68]. As a result, the formed 2D nanosheets, after exfoliating, shine brightly in various fields [68, 70]. In particular, when the liquid‐phase exfoliated layered materials exhibit ion conduction and electronic insulation phenomenon, the resulting 2D inorganic nanosheets are gradually developed as fast ion conductors [66]. Currently, according to the types of non‐metallic elements that make up the organic materials, 2D inorganic nanosheets with ion transport behavior can be divided into oxide‐based ion transport nanosheets, phosphate‐based ion transport nanosheets, and metal sulfur/phosphorus sulfide‐based ion transport nanosheets. These nanosheets can utilize the unique physicochemical properties of the nanolayers to transport ions.

A critical cross‐class comparison of 2D inorganic nanosheets reveals striking performance disparities, which exhibit theoretical guidance significance. Oxide‐based nanosheets span from clay minerals and perovskites to layered double hydroxides (LDHs) [51, 57]. Natural clay minerals (e.g., montmorillonite, vermiculite, halloysite, kaolinite, attapulgite, sepiolite) feature a super‐stable layered structure (Figure 4a‐i) [71]. After exfoliation and acidification, H‐montmorillonite relies on the vehicle mechanism to deliver the conductivity of ∼1—10 × 10⁻³ S cm⁻¹ (20°C–30°C) (Figure 4a‐ii,a‐iii) [51, 72]. However, among the entire clay mineral family, proton conductivities remain modest due to the narrow interlayer galleries and strong interlayer forces, which constrain both the Grotthuss and vehicle pathways. As for transporting alkali metal ions [78, 79], the situation is even more constrained, the ‐O⁻ functional groups form strong ionic bonds with cations, and the limited alkali metal ion sources within the interlayers typically restrict conductivities to 10⁻³ S cm⁻¹ [80, 81] (Figure 4a‐iv). This inherent conductivity ceiling explains why clay nanosheets have established their most impactful applications not as standalone conductors but as functional additives in polymer electrolytes (Figure 4a‐v), where their negatively charged surfaces can weaken Li⁺‐polymer and Li⁺‐anion interactions [57, 73, 82, 83], boosting the conductivity of composite electrolyte to ∼1.0 × 10⁻⁴–1.06 × 10⁻³ S cm⁻¹ (Figure 4a‐vi), this conductivity is an order of magnitude improvement over the pure polymer electrolytes.

FIGURE 4.

Oxide‐based nano ion conductors. (a) Clay minerals based 2D nanosheets. (i) The layer structure of 2D clay minerals. Reproduced with permission [71]. Copyright 2022, the Royal Society of Chemistry. (ii) The transport of proton in 2D montmorillonite nanosheets [72]. (iii) The proton conductivity of montmorillonite nanosheet. Reproduced with permission [72]. Copyright 2025, American Chemical Society. (iv) The ion bonds between vermiculite nanosheets and Li⁺. Reproduced with permission [73]. Copyright 2019, Wiley‐VCH. (v) The structure of 2D clay minerals nanosheets as an additive in polymer electrolyte. (vi) The conductivity of composite polymer electrolytes with 2D clay minerals nanosheets. (b) 2D perovskite materials and single metal oxides. (i) The structure of 2D perovskite materials and single metal oxides. Reproduced with permission [66]. Copyright 2022, AIP Publishing. (ii) The preparation process of 2D perovskite materials or single metal oxides. Reproduced with permission [74]. Copyright 2024, Wiley‐VCH. (iii) The Li⁺ conductivity of TiO₂ nanosheet. Reproduced with permission [75]. Copyright 2022, American Chemical Society. (c) LDH [76]. (i) The structure of LDH. (ii) The OH⁻ transport network of LDH. (iii) The OH⁻ conductivity of LDH. Reproduced with permission [76]. Copyright 2025, American Chemical Society. (d) ϒ‐AlOOH‐PBI. (i) The preparation of ϒ‐AlOOH‐PBI. (ii) The OH⁻ transport diagram in ϒ‐AlOOH‐PBI. (iii) The OH⁻ conductivity of ϒ‐AlOOH‐PBI. Reproduced with permission [77]. Copyright 2023, The Royal Society of Chemistry.

A mechanistically distinct paradigm emerges in exfoliated perovskite and monometallic oxide nanosheets in comparison with above 2D clay minerals nanosheets, layered perovskite materials (e.g., KLaNbO₅ and KCa₂Nb₃O₁₀) and single metal oxides (e.g., TiO₂, NbO₅ and TaO₃) with edge‐shared/corner‐shared MO₆ octahedra (Figure 4b‐i) [66, 84, 85] can be exfoliated into monodispersed nanosheets after acidification and organic cation‐assisted exfoliation [74, 75, 84] (Figure 4b‐ii), which theoretically are able to utilize the well‐ordered 2D water networks formed by interlayer water for rapid ion transport [86]. However, early studies using acidified layered perovskite or acidified layered oxide without an exfoliated structure as proton conductors to deliver low H⁺ conductivity (<0.01 mS cm⁻¹), underscoring that exfoliation is not merely a morphological change but a prerequisite for achieving the fast H⁺ conduction potential. The recent Li⁺ conduction in exfoliated TiO₂ nanosheets achieve the Li⁺ conductivity of 0.21 S and 0.62 S cm⁻¹ at T = 25°C (RH = 100%) and T = 90°C (RH = 100%), respectively (Figure 4b‐iii) [75], which represents a significant breakthrough, highlighting the exfoliation and ion exchange combined with guest water‐mediated transport can elevate ionic conductivity by orders of magnitude compared to non‐exfoliated counterparts.

LDHs present the most successful case of anion conduction among oxide nanosheets, which has the general formula of [M_1–x ²⁺M_x ³⁺(OH)₂]. Their 2D cationic layered structure is alternately composed of divalent (M_1–x ²⁺) and trivalent (M_x ³⁺) metal cations [87, 88], whereas OH⁻ is located between the cationic layers to balance their positive charges (Figure 4c‐i). Therefore, LDH nanosheets can be used as ion conductors for anion transport. With this understanding, single‐layer LDH nanosheets, such as Mg‐Al and Co‐Al LDH, achieve rapid OH⁻ conduction through the interlayer hydrogen bonding networks formed by hydroxyl groups, interlayer anions, and water molecules [89] (Figure 4c‐ii), and their OH⁻ conductivity can be as high as 10⁻¹ S cm⁻¹ (Figure 4c‐iii) [76], placing them among the highest‐performing OH⁻ conductors across all 2D inorganic materials. Alternately, benefiting from the large contact area and the definite hydrogen bonding path in nanochannels of ϒ‐AlOOH (Figure 4d‐i,d‐ii), ϒ‐AlOOH‐PBI outputs the OH⁻ conductivity up to 60 mS cm⁻¹ at 40°C (Figure 4d‐iii) [77]. Beyond anion transport, the interlayer anionic characteristics of LDHs show potential for small cation conduction after ion exchange with traditional salts (such as LiCl and LiBF₄) if they are not significant restacking, providing a direction for further exploration.

Phosphate nanosheets reveal that how the exfoliation process transforms modest proton conductors into ultra‐high performance ion conductors. Layered phosphates have an anionic layered structure composed of metal polyhedra and PO₄ ²⁻ tetrahedron alternately [90] (Figure 5a‐i). Co‐precipitated phosphates (e.g., HUO₂PO₄·4H₂O, α‐Zr(HPO₄)₂·nH₂O and γ‐Zr(PO₄)(H₃PO₄)·2H₂O) can transport proton under the Grotthuss mechanism or vehicle mechanism [91, 92, 93]. For example, layered polyvalent (Zr or Ta) hydrogen phosphate exhibits the conductivity of 0.2 mS cm⁻¹ at T = 25°C and RH = 60% [92, 93]. However, a genuine breakthrough in transport ion lies in the spontaneously exfoliating solid acid nanosheets formed by high‐temperature calcination [94, 95, 96]. Taking H_nSb_nP₂O_3n+5, which exhibits the spontaneous exfoliating behavior, as an example (Figure 5a‐ii), the proton conductivity of H₃Sb₃P₂O₁₄ and HSbP₂O₈ members can be as high as 1.02 S and 1.18 S cm⁻¹ at RH = 100% and T = 90°C, respectively (Figure 5a‐iii,a‐iv) [97], which is hundreds of times higher than that of the Zr(HPO₄)₂·nH₂O based materials (10⁻⁵–10⁻³ mS cm⁻¹). This dramatic enhancement in proton conductivity strongly supports that the principle of spontaneous exfoliation maximizes the accessible surface area and hydrogen bonding connectivity of the nanosheets, becoming a key structural feature for determining ultra‐high proton conductivity.

FIGURE 5.

Layered phosphates, layered sulfides, phosphorus sulfide and halides based nano ion conductors as well as the regulation of 2D inorganic nanosheets. (a–c) Layered phosphates, layered sulfides and phosphorus sulfide based nano ion conductors (a) Layered phosphates. (i) The structure of layered H₃Sb₃P₂O₁₄. (ii) The structure of H₃Sb₃P₂O₁₄ and HSbP₂O₈ nanosheets. (iii and iv) The proton conductivity of H₃Sb₃P₂O₁₄ nanosheet. Reproduced with permission [97]. Copyright 2024, Springer Nature. (b) Li_xSnS₂ nanosheet. (i) The positions of Li⁺ in Li_xSnS₂ nanosheet. (ii) The positions of H₂O in Li_xSnS₂ nanosheet. (iii) The Li⁺ conductivity of Li_xSnS₂ nanosheet. Reproduced with permission [65]. Copyright 2021, American Chemical Society. (c) CdSP₃‐X nanosheet. (i) The structure of Cd_0.85PS₃Li_0.15H_0.15 nanosheet. Reproduced with permission [40]. Copyright 2020, The American Association for the Advancement of Science. (ii) The structure of CdSP₃‐X nanosheet [101]. (iii) The conductivity of Cd_0.85PS₃Li_0.15H_0.15 nanosheet [40]. (iv) The conductivity of CdSP₃‐X nanosheet. Reproduced with permission [101]. Copyright 2023, Springer Nature. (d) BiOI nanosheet [102]. (i) The structure of BiOI_1‐x(OH)_x. (ii) Ion conductivity of BiOI_1‐x(OH)_x. Reproduced with permission [102]. Copyright 2025, The American Association for the Advancement of Science. (e, f) Regulation of the 2D inorganic nanosheets. (e) Inorganic nanosheet. (i) The exfoliation of an inorganic nanosheet by an intercalant with a short chain. (ii) The exfoliation of an inorganic nanosheet by intercalant with a long chain. (f) Alkanediamine cross‐linked vermiculite membrane. Reproduced with permission [104]. Copyright 2022, American Chemical Society.

Metal sulfide and phosphorus sulfide nanosheets have produced the most striking conductivity values among all 2D inorganic ion conductors, despite the narrow range of materials and greater synthesis challenges. The limited family of ion‐conducting sulfide nanosheets includes Li_xSnS₂ and AgCrS₂ [98, 99], while phosphorus sulfide nanosheets with ion conduction behavior merely contain FePS₃‐X, NiPS₃‐X, CdPS₃‐X, and MnPS₃‐X [40, 100]. Li_xSnS₂ nanosheet [65], which is formed by spontaneous exfoliation of Li_0.8Sn_0.8S₂ in water, generates sulfur vacancy structures that are then occupied by OH⁻. Meanwhile, water molecules will be immersed into Sn−S layers to expand their layer spacing (Figure 5b‐i,b‐ii). The resulting synergistic Li⁺/water diffusion yields Li_xSnS₂ with a conductivity up to 47 mS cm⁻¹ (Figure 5b‐iii). 2D AgCrS₂, formed by electrochemical exfoliation in organic solvents, reduce the migration energy of Ag⁺ at tetrahedral positions, endowing a high ionic conductivity of 33.2 mS cm⁻¹ at anhydrous condition [46]. Such high cation conductivity without relying on water‐assisted transport shows significant implications for all‐solid‐state batteries. As For MPS₃‐X, the vacancy‐engineered 2D Cd_0.85PS₃Li_0.15H_0.15 delivers the proton conductivity of ∼0.95 S cm⁻¹ (90°C, RH = 98%), while 2D Cd_0.85PS₃Li_0.3 and 2D Mn_0.77PS₃Li_0.46 achieve Li⁺ conductivities of 0.80 and 0.75 S cm^{− 1} under identical conditions. (Figure 5c‐i,c‐iii) [40]. Most impressively, systematic tuning of intercalated cations in MPS_3‐x (M = Fe, Ni, Cd) enables universal alkali metal ion conduction [100, 101] (e.g., the conductivity of 0.01–0.8 S cm⁻¹ for CdPS₃‐X, X = K, Na, Li, Ca, Mg, and Al) (Figure 5c‐ii,c‐iv) [101]. The versatility of MPS₃‐X cannot be unmatched by any other 2D inorganic system. The recent BiOI_1‐x(OH)_x exfoliated by layered BiOI delivers the high OH⁻ conductivity up to 168 mS cm⁻¹ at 90°C (Figure 5d) [102], which further expands the design space. In view of the abundance and high ionic conductivity potential of bulk transition metal sulfides and transition metal halides (e.g., MX₃, AMo₃X₃, NbX₃, TiX₃, and TaX₃, X = I, S, Se, or Te) [67, 103], the development of suitable liquid phase exfoliation technologies represents a high‐reward frontier, which can substantially broaden the scope of 2D inorganic ion conductors.

In summary, although the above‐mentioned 2D inorganic nanosheets with high ionic conductivity have been developed, their proportion in the entire SSICs is like a drop in the ocean. The key bottleneck is not the intrinsic conduction capability of the nanosheets themselves, but rather the scarcity of scalable exfoliation techniques and rational interlayer engineering strategies [105]. Therefore, the design of new liquid phase exfoliation technology and layered manufacturing process is vital for the development of 2D inorganic nano ion conductors.

3.1.2. Regulation of the 2D Inorganic Nanosheets

The optimization of the ion conduction performance for 2D inorganic nanosheets originates from the residual acidic functional groups in their layered nanochannels. Current strategies can be divided into two categories: regulating the interlayer environment of 2D inorganic nanosheets and tailoring the element ratio of 2D inorganic nanosheets by organic molecules or ammonium salts [27, 66, 85]. Although both approaches aim to enhance ionic conductivity, they operate at different structural levels and carry different practical implications. The former modifies the physicochemical properties of the interlayer, including layer spacing, hydrophilicity, and host–guest interactions, while the latter reconstructs the atomic arrangement of the 2D inorganic nanosheets, thereby redesigning the intrinsic ion transport channel. Interlayer regulation by ammonium salts represents the most versatile and widely strategy, however, the choice of intercalant critically determines the resulting ion transport properties. The exfoliation of Ca₂Nb₃O₁₀ ⁻ and TiNbO₅ ⁻ by AUA with a long chain can achieve the larger layer spacing and exfoliation rate than that of TBAOH [106] (Figure 5e‐i,e‐ii), the obtained large layer spacing is very conducive to transport cations. This demonstrates that the steric and electrostatic characteristics of the intercalating cation directly governs the interlayer space, thereby determining the ion conduction efficiency. Similarly, the modification of H₃Sb₃P₂O₁₄ by ammonium salts with different lengths of alkyl chains can simultaneously improve its water retention capacity and layer spacing, which is helpful to optimize the water network structure between layers [107]. In clay minerals, organic salts selectively converts some hydrophilic channels between its layers into tunable interlayer channels, such regulation is favor for improving the mass flux in clay minerals [108], which can enhance ion transport via tunning the interlayer environment.

In contrast, optimizing the interlayer environment by organic molecules mainly weakens the interaction between the anion frameworks and the cations. Such an effect is evidenced by the ammonium molecules (e.g., TBA, DMAE, TMA, HDAVM, BDAVM and EDAVM) intercalated 2D inorganic nanosheets (e.g., Ca₂Nb₃O₁₀ ⁻, TiO₂ and vermiculite) (Figure 5f) [104, 109]. However, this approach typically yields smaller enhancement in conductivity than ammonium salt intercalation due to it is hard to construct new ion transport pathways.

Regulating the elemental composition and proportions of 2D inorganic nanosheets offers greater potential for intrinsic ion conduction breakthroughs, perovskite and monometallic oxide can serve as examples, H₂Ca_n‐1Mn_n‐3Nb₃O_3n+1 [110] and KNb_nO_m (e.g., K₄Nb₆O₁₇ and KNb₃O₈) [111, 112] possess different interlayer arrangement and interlayer thickness, which influence the layer spacing and surface acidic functional groups in the exfoliation process, eventually affecting the ion transport performance and the mechanical properties. So, when selecting 2D perovskite and monometallic oxide nanosheets as ion conductors, their comprehensive effect should be considered.

At present, in view of the limitations of inorganic nano ion conductors that have been developed, their modification methods merely include the above two categories. Therefore, the enrichment of modification methods can only be premised on developing more 2D nanosheets that can transport ions. In addition, since the current exfoliation of 2D inorganic nanosheet based ion conductors is mainly performed in aqueous solvent, their corresponding interlayer modification is also carried out in water. If the exfoliation process or modification technology in organic solvents can be vigorously developed, then this is bound to promote the application of inorganic nano ion conductors in non‐aqueous environments.

3.2. Design of Organic Nano Ion Conductors

Organic nano ion conductors, i.e., carbon‐based nano ion conductors, mainly include nano hydrogen‐bonded organic frameworks (HOFs) and nano COFs (N‐COFs). Compared with nano HOFs, N‐COFs, as organic nano ion conductors, are currently the most researched and hottest topic due to their richness and diversity. Typically, N‐COFs are a class of 2D or 3D crystalline materials covalently crosslinked by organic monomers under Bottom‐up synthesis and Top‐down method strategies, in which Bottom‐up synthesis focuses on designing COFs with 2D and 3D structure, while Top‐down method concentrates on exfoliating stacked COFs [13, 30, 32]. Compared with bulk COFs, N‐COFs as ion conductors present the following advantages: (1) The connected nanolayered structure of 2D COFs and the interconnected nano networks of 3D COFs shorten the ion transport pathways, as well as promote orderly ion conduction [30]. Besides, the nanosizing of COFs increases the contact area between particles inside the ion conductors and reduces the ion transport resistance at the interface [13]. (2) N‐COFs feature high specific surface area and abundant exposed active sites, which is conducive to introducing the guest substances to promote ion transport [113]. Furthermore, the monomers of N‐COFs possess high substitutability and controllability, which provides infinite possibilities for regulating their physicochemical properties [114]. According to the charged condition, N‐COFs can be divided into ionic N‐COFs and non‐ionic N‐COFs, ionic N‐COFs not only rely on their ionic functional groups to transport ions, but also combine trace guest liquid substances for efficient ion transport. By contrast, non‐ionic N‐COFs with ion‐free properties can only be combined with liquid electrolytes for ion transport. Currently, the design of N‐COF based ion conductors mainly comprises developing N‐COFs with ion conduction function and regulating the ion transport performance of N‐COFs. The former is devoted to designing new monomers and/or processes to develop N‐COFs, while the latter puts emphasis on regulating the physicochemical properties of already developed N‐COFs.

3.2.1. Development of N‐COFs

The development of N‐COFs based ion conductors can be classified into developing new ionic N‐COFs and non‐ionic N‐COFs. Generally, ionic N‐COFs have ionic functional groups that promote ion transport. Hence, the characteristics of the ion skeleton are the main induced factor for their ion transport [37]. By contrast, non‐ionic N‐COFs can only utilize the confinement effect of nanochannels or the polarity effect of functional groups for ion transport [32]. Consequently, ionic N‐COFs usually exhibit higher ionic conductivity than non‐ionic N‐COFs. According to the charge type on ionic skeletons, ionic N‐COFs are divided into cationic N‐COFs and anionic N‐COFs. The copolymerization frameworks of anionic N‐COFs possess a great number of negatively charged groups (e.g., sulfonic acid, benzimidazole, carboxylic acid), so they could achieve the cationic conduction effect. In contrast, the frameworks of cationic N‐COFs have some quaternary ammonium salt cations, which realizes the anion conduction. Despite this, considering the limitations of conducting anions in practical applications, cationic N‐COFs need to be combined with guest electrolytes to transport cations in many cases.

In anionic N‐COFs, the choice of ionic groups not only influences the ion conduction mechanism but also determine the attainable ionic conductivity. Proton conduction, primarily mediated by ‐SO₃ ⁻ groups on the copolymerization framework, achieves enhanced performance under the Grotthuss mechanism. The ‐SO₃ ⁻ moieties serve as proton carriers, forming a hydrogen bonding network with water molecules inside the nanochannels (Figure 6a). Critically, the wide proton conductivity range reported for TpPa‐SO₃H underscores that synthesis strategy, rather than framework chemistry alone, governs the ultimate conductivity. The proton conductivity of TpPa‐SO₃H nanosheets synthesized by a single solution method is 167.3 mS cm⁻¹ at T = 30°C and RH = 98%, which increases to 364.1 mS cm⁻¹ at T = 80°C [115] (Figure 6b‐i,b‐ii). The simultaneously controlled nucleation and in‐plane growth for TpPa‐SO₃H lead IPC‐COF nanosheets to have large 1D nanochannels and abundant proton carriers, it delivers the proton conductivity of 0.38 S cm⁻¹ at T = 80°C [116] (Figure 6c‐i,c‐ii). The surface‐initiated condensation polymerization technology enables TpPa‐SO₃H nanosheet to exhibit the super‐high proton conductivity of 0.54 S cm⁻¹ in pure water at T = 80°C [117] (Figure 6d‐i,d‐ii), which is one of the highest values reported for COF‐based proton conductors. This progressive improvement across the above synthesis methods directly illustrates that maximizing channel order and carrier accessibility are equally important to acidic group in anionic N‐COFs.

FIGURE 6.

Anionic N‐COFs based ion conductors. (a) H⁺ transport in anionic N‐COFs. (b) TpPa‐SO₃H nanosheets synthesized by single solution method. (i) The molecular structure diagram of TpPa‐SO₃H. (ii) The conductivity of TpPa‐SO₃H under the single solution method. Reproduced with permission [115]. Copyright 2023, Wiley‐VCH. (c) IPC‐COF nanosheets. (i) The preparation diagram of IPC‐COF nanosheets. (ii) The ionic conductivity of the IPC‐COF membrane. Reproduced with permission [116]. Copyright 2020, Wiley‐VCH. (d) TpPa‐SO₃H synthesized by surface‐initiated condensation polymerization technology. (i) The surface initiated condensation polymerization technology for TpPa‐SO₃H. (ii) The ionic conductivity of TpPa‐SO₃H under this method. Reproduced with permission [117]. Copyright 2021, Wiley‐VCH. (e) DT‐COF nanosheets. (i) The molecular structure diagram of DT‐COF nanosheets. (ii) The ionic conductivity of DT‐COF. Reproduced with permission [118]. Copyright 2022, The American Association for the Advancement of Science. (f) Li⁺ transport in anionic N‐COFs. (g) Phosphorylated COF based nanosheets [119]. (i) The molecular structure diagram of phosphorylated COF. (ii) The Li⁺ transport mechanism of phosphorylated COF. (iii) The ionic conductivity of phosphorylated COF. Reproduced with permission [119]. Copyright 2024, Wiley‐VCH. (h) TPDBD‐CNa‐NaTFSI nanosheets. (i) The molecular structure diagram of TPDBD‐CNa‐NaTFSI. (ii) The Na⁺ transport diagram in TPDBD‐CNa‐NaTFSI. (iii) The ionic conductivity of PC@TPDBD‐CNa‐NaTFSI. Reproduced with permission [120]. Copyright 2023, Springer Nature.

Except for TpPa‐SO₃H nanosheets with proton conduction under the above direct synthesis method, anionic SDT‐COF nanosheets with ‐SO₃H also can be synthesized through the continuously grafting functional groups onto non‐ionic 3D DhaTab‐COF, which exhibits excellent water dispensability and delivers the proton conductivity of up to 95.3 mS cm⁻¹ at RH = 98% in the directional water channel formed by ‐SO₃H and water molecules [118] (Figure 6e‐i,e‐ii). Despite this post‐synthetic approach provides greater structural versatility, its conductivity is lower than the directly synthesized analogs, suggesting that grafting density and spatial uniformity remain limiting factors of H⁺ conductivity.

Alkali metal ion conduction in anionic N‐COFs exhibits a fundamentally different and more challenging mechanistic characteristic, cations are transported based on the Lewis acid‐base dissociation mechanism between the anionic framework and cations (Figure 6f), this is inherently less efficient than Grotthuss hopping. For example, TpMbh‐PO₃Li₂ exhibits definite anionic nano channels and high Li⁺ distribution density, which accomplish fast Li⁺ conduction at room temperature (1.7 mS cm⁻¹) [119] (Figure 6g). However, this Li⁺ conductivity remains approximately two orders of magnitude below the above proton‐conducting N‐COFs, which reflects the intrinsic kinetic difference between vehicular/Lewis acid–base transport and Grotthuss hopping. More commonly, the strong interaction between most 2D anionic framework and Li⁺ leads TpPaSO₃‐Li [121, 122] to deliver low ionic conductivity, such as 10⁻⁵ S cm⁻¹ for TpPaSO₃‐Li, which occasionally cannot meet the requirements of fast ion conductors.

Therefore, the anionic N‐COFs sometimes should combine a little guest solvents or electrolytes to realize the fast transport of alkali metal ions. The contrast between pristine and solvent‐modified anionic N‐COFs is obvious. Li‐ImCOF nanosheets [61] and TPDBD‐CNa‐NaTFSI nanosheets (Figure 6h‐i) [120] feature low conductivity (e.g., ∼5 × 10⁻⁷ S cm⁻¹ for Li‐ImCOF) owing to the strong interaction between Li⁺/Na⁺ and benzimidazole anion/–COO^–. Once a trace amount of PC is added to them, PC could weaken this interaction due to the solvation effect, thereby achieving fast Li⁺/Na⁺ transport in 2D sub‐nanometer channels. Eventually, PC@Li‐ImCOF and PC@TPDBD‐CNa‐NaTFSI (Figure 6h‐ii,h‐iii) deliver conductivity up to 5.3 × 10⁻³ and 1.30 × 10⁻⁴ S cm⁻¹, respectively, while maintaining their original flexibility and strength [61, 120]. The approximately 10⁴‐fold enhancement of conductivity achieved through the addition of a trace amount of solvent embodies a practical design principle, namely, although the strong framework‐cation interactions are detrimental in dry conductors, they can be effectively modified through the use of a cosolvent, which competitively solvates cations without disrupting the integrity of the COF framework. Three‐dimensional anionic COFs follow the same mechanistic logic. If well‐defined nanochannels are formed in anionic 3D COFs, fast ion transport can be achieved under the dissociation mechanism (e.g., 2.67 × 10⁻⁵ S cm⁻¹ for LiO₃S‐COF1 [123] and 2.68 × 10⁻⁴ S cm⁻¹ for NaOOC‐COF [124]), then they can be directly used as the ion conductors. Conversely, if 3D anionic COFs have poor cations or strong interaction between cations and the anionic framework, they should utilize their high specific surface to combine with guest liquid electrolytes to complete the rapid transport of cations.

In summary, the ionic conductivity of anionic N‐COFs that depend on the Grotthuss mechanism to transport proton is much higher than those that rely on the acid‐base dissociation mechanism to transport Li⁺. This disparity is not accidental but stems from its underlying mechanism, the activation energy barrier encountered by proton hopping in the hydrogen bonding network is significantly lower than that of solvated alkali metal cation transport. Therefore, the primary design challenge for N‐COFs with alkali metal ion conduction lies in weakening the interaction between cations and the COF framework without sacrificing structural clarity. Future ion monomer designs may achieve this balance through precise tuning of charge delocalization and spatial shielding of anionic sites. In addition, in view of the pore size confinement effect and nanochannel solvation effect of anionic N‐COFs, special attention should be paid to the transport effects of cations by different solvents when encapsulating guest liquid electrolyte into them.

Cationic N‐COFs are structurally opposite to their anionic counterparts, and the charge polarity of the framework determines fundamental differences in their ion transport properties. In cationic 2D COFs, their large charge density on the ionic framework prevents the stacking of nanosheets [125], forming the interlayer nanochannels with fast anion transport (Figure 7a‐i). This point is evidenced by the high OH⁻ conductivity in hydrazone‐linked COF (COF‐3OH) [126] (e.g., over 0.25 S cm⁻¹ for COF‐3OH at 80°C and 100% RH) (Figure 7a‐ii,a‐iii), confirming that appropriately designed cationic N‐COFs can support anion transport as effectively as Grotthuss type transport. However, apart from OH⁻ conduction, conducting other anions in cationic 2D COFs has limited practical application. Therefore, liquid electrolytes/solvents are usually encapsulated into them for cation transport (Figure 7b). In such hybrid nano ion conductors with liquid electrolytes, the cation framework of cationic 2D COFs could adsorb guest anions, which weakens the interaction between guest anions and cations, eventually completing the fast transport of proton. As proof, the cationic framework of QACOF nanosheet (namely quaternized covalent organic framework) can be paired with H₂PO₄ ⁻ to reduce the dissociation energy of H‐H₂PO₄ ⁻, which leads H₃PO₄@QACOFMs to deliver proton conductivity up to 379.7 mS cm⁻¹ at 200°C [127]. Crucially, this performance is achieved under anhydrous conditions, which are unattainable for water‐dependent proton conductors, and far surpasses the performance of most hydrated N‐COFs. This comparison suggests that electrostatic decoupling of the proton carrier from its conjugate framework, rather than merely maximizing water absorption capacity, may be a more effective approach to achieving high‐temperature proton conduction. For cationic 2D COFs that transport proton under a water environment, cationic covalent organic framework film (TG‐DFP‐COF) achieves a proton conductivity of 2.8 mS cm^{− 1} at 99% RH under a continuous hydrogen‐bonding network that is constructed by ion‐dipole interactions among confined water, Cl⁻ counterions and the guanidinium framework (Figure 7c) [128], surpassing all reported cationic COFs for proton transport under water environment. This validates the concept of engineering hydrated ion channels as a viable ion conductor design strategy. Future development of cationic COFs could exploit counterion screening and backbone polarity modulation to realize nano ion conductors that combine high conductivity with mechanical adaptability.

FIGURE 7.

Cationic N‐COFs based ion conductors. (a) Anion transport in cationic N‐COFs [126]. (i) The diagram of anion transport in the nanochannel of COF‐3OH. (ii) The preparation and structure diagram of COF‐3OH. (iii) The ionic conductivity of COF‐3OH. Reproduced with permission [126]. Copyright 2025, Wiley‐VCH. (b) Cation transport in cationic N‐COFs. (c) TG‐DFP COF based nanosheet. (i) The preparation and structure diagram of TG‐DFP COF. (ii) The proton transport mechanism of TG‐DFP COF. (iii) The ionic conductivity of TG‐DFP COF. Reproduced with permission [128]. Copyright 2024, Elsevier Ltd. (d) Li‐TpTGTFSI‐COF naosheets. (i) The schematic diagram of ion association in cationic N‐COFs and non‐ionic N‐COFs. (ii) The molecular structure diagram of Li‐TpTGTFSI‐COF. (iii) The ionic conductivity of Li‐TpTGTFSI‐COF. Reproduced with permission [129]. Copyright 2018, American Chemical Society.

As for alkali metal ion transport, the TpTGCl‐COF nanosheets (Tp: 1,3,5‐triformylphloroglucinol, TGCl: triaminoguanidinium chloride) with self‐exfoliation behavior offers a structurally instructive case. The strong coulombic interaction between the host framework and guest ions results in TpTGCl‐COF having a loose ionic framework, which enables the cationic framework of Li‐TpTGTFSI‐COF nanosheets to interact with more TFSI⁻ (Figure 7d‐i), endowing more freedom Li⁺ in the nanochannels (Figure 7d‐ii). Hence, Li‐TpTGTFSI‐COF delivers the ionic conductivity of 2.09 × 10⁻⁴ S cm⁻¹ at 70°C (Figure 7d‐iii) [129]. Although the absolute ionic conductivity remains modest, TpTGCl‐COF conceptually illustrates how strong host–guest Coulombic interactions can induce a self‐exfoliated, loose ionic framework to liberate Li⁺, offering a structurally instructive insight for future N‐COFs design.

Based on the aforementioned examples, cationic N‐COFs usually need to combine with ionic substances to finish the high cation conductivity. Meanwhile, the cationic N‐COFs as excellent ion conductor should meet the following requirements: (1) cationic N‐COFs should have a large specific area, which guarantees them to expose more ion sites to capture guest anions; (2) the cationic frameworks of cationic N‐COFs are expected to have strong interactions with guest anions, which is favorable to the dissociation of targeted cations. Furthermore, the currently developed cationic framework of N‐COFs mainly belongs to quaternary ammonium salts. If phosphonium cation and sulfonium cation‐based frameworks in N‐COFs can be designed, or phosphonium cations and sulfonium cations can be grafted onto the surface of cationic N‐COFs, this will promote the rapid advancement of N‐COFs.

As regards the non‐ionic N‐COFs, due to the non‐ionizing nature in directional nanochannels (Figure 8a), they typically output relatively low ionic conductivity in aqueous environments and organic solvents. For example, the proton conductivity of 2D COF with strong interaction with water (COF‐H) (Figure 8b‐i) is only 8 × 10⁻⁸ S cm⁻¹ at T = 80°C and RH = 85% [53]. This proton conductivity is several orders of magnitude lower than the threshold for practical applications, emphasizing that the neutral nanochannel surface cannot maintain meaningful ion transport, and external functionalization or guest binding is an absolute prerequisite for achieving competitive ion conduction. Hence, non‐ionic N‐COFs either combine with an ionic substance to achieve fast proton conduction, or encapsulates organic electrolytes to fulfill Li⁺ conduction. It is crucial that these strategies operate through different mechanisms and pathways, resulting in vastly different performance characteristics that require careful comparative evaluation. Given the nanochannel confinement effect of non‐ionic N‐COFs (Figure 8b‐ii), PA@TAP‐COF nanosheets (phosphoric acid@ 2,4,6‐triaminopyrimidine‐COF) exhibit the proton conductivity of 2.65 × 10⁻³ S cm ⁻¹ at T = 140°C under anhydrous conditions [134], and PEG@COF+LiClO₄ show the Li⁺ conductivity of 1.45 × 10⁻⁵ S cm⁻¹ at T = 303 K (Figure 8b‐iii) [130]. Although these improved ionic conductivities represent significant improvements over the original framework, they are still moderate in absolute terms, which indicates that simple physical limitations are not sufficient to fully utilize the ion transport potential.

FIGURE 8.

Non‐ionic N‐COFs based ion conductors. (a) Cation transport in non‐ionic N‐COFs without group. (b) N‐COFs without obvious functional groups. (i) The molecular structure diagram of COF‐H. Reproduced with permission [53]. Copyright 2024, Wiley‐VCH. (ii) Schematic representation of COF‐5, PEG‐800, and LiClO₄. (iii) The ionic conductivity of PEG@COF‐5‐ LiClO₄ composite. Reproduced with permission [130]. Copyright 2025, American Chemical Society. (c) Cation transport in non‐ionic N‐COFs with group. (d) 2D Py‐R−COFs [131]. (i) The structure of 2D Py‐R−COFs. (ii) The mass transport of 2D Py‐R−COFs. (iii) The ionic conductivity of 2D Py‐R−COFs. Reproduced with permission [131]. Copyright 2026, American Chemical Society. (e) Some typical non‐ionic N‐COFs with polar functional groups. (i) The structure of 3D TPB‐DMeTP‐COF. Reproduced with permission [132]. Copyright 2020, Springer Nature. (ii) The synthesis and structure of PA@PyTTA‐BMTP‐COF. Reproduced with permission [133]. Copyright 2023, the Royal society of chemistry. (iii) The ionic conductivity of these non‐ionic N‐COFs.

It is essential to note that once functional groups are assigned to the skeleton of non‐ionic N‐COFs, these polar functional groups in directional nanochannels could provide hopping sites to transport alkali metal ions or form water networks to conduct proton [36, 131] (Figure 8c), such as 2D Py‐OMe‐COF (pyrene‐based tetragonal COFs) rely on the ordered hydrogen‐bond network by −OMe group and water molecules to realize the fast H⁺ hopping (Figure 8d‐i,d‐ii), which delivers the higher conductivity (1.24 × 10⁻⁶ at 303 K) than Py‐OH‐COF (Figure 8d‐iii) [131]. Notably, this performance is achieved without additional acid doping, indicating that a high density of such polar functional sites can autonomously generate abundant proton transport pathways under optimal hydration conditions, which is a design principle worthy of wider application. However, for most non‐ionic N‐COFs, simply binding functional groups is not sufficient to achieve performance comparable to charged frameworks, and synergistic combination with guest substances is still necessary. The most impressive anhydrous proton conductivity in non‐ionic systems has been achieved through this hybrid method. 3D H₃PO₄@TPB‐DMeTP‐COF (TPB: 1,3,5‐tri(4‐aminophenyl)benzene, DMeTP: 2,5‐dimethylterephthalaldehyde) [132] (Figure 8e‐i) nanoparticles, 2D H₃PO₄@TPB‐DABI‐COF (DABI: 2,5‐diamino‐3,4‐dimethylbenzimidazole) [52] nanosheets and PA@PyTTA‐BMTP‐COF (BMTP: 2,5‐bis(2‐(2‐methoxyethoxy)ethoxy‐p‐phenylenedicarboxylic acid, DHTA: 2,5‐dihydroxyterephthalaldehyde) [133] (Figure 8e‐ii,e‐iii) achieve the ultrahigh anhydrous proton conductivity approach ∼10⁻²–10⁻¹ S cm⁻¹ at high temperatures by virtue of the strong guest salt confinement effect and ordered proton transport pathway promoted by polar functional groups in non‐ionic N‐COFs, representing a true breakthrough as they originate from frameworks that are essentially non‐conductive, also verifying that non‐ionic N‐COFs are no longer the passive containers, but active participants in reconstructing the solvation environment.

Therefore, for proton conduction, the guest confinement method of a polar functionalized framework is significantly better than the simple physical confinement of non‐functionalized N‐COFs, indicating that pore wall polarity is an important factor in organizing proton carriers. For Li⁺ conduction, the electrolyte encapsulation route provides the highest absolute conductivity. Hence, the key design implication for improving the ionic conductivity of non‐ionic N‐COF is maximizing the density of polar functional groups within the nanochannel, as it simultaneously enhances the host‐guest interaction, improves ion pair dissociation, and provides auxiliary hopping sites.

Based on these, the ionic conductivity of N‐COFs largely depends on the ion transport mechanism. Proton transport via the Grotthuss mechanism (e.g., in anionic N‑COFs) outperforms alkali metal ion transport through acid–base dissociation, owing to the much lower activation energy of proton hopping. For alkali metal ion transport, weakening the interaction between cation and MOF framework by charge delocalization and spatial shielding is critical. For non‑ionic N‑COFs, pore wall/interlayer polarity plays a dominant role in organizing ion carriers. To accelerate the discovery of N‑COFs, computational screening of ionic/non‑ionic nano frameworks and expansion to other alkali metal salts (e.g., Na⁺, K⁺, Mg²⁺) are urgently needed.

3.2.2. Regulation of the N‐COFs

The regulation of N‐COFs refers to tailoring the physicochemical properties (e.g., chemical environment and nanochannel size) in their directional channel by some in situ and ex situ methods, and it is mainly divided into polar functional group regulation, ionic regulation, and guest molecule insertion regulation. Generally, polar functional group regulation aims to increase the polar sites on the frameworks of N‐COFs, ionic regulation focuses on optimize the type and number of ionic functional groups in ionic N‐COFs, and guest molecular insertion regulation involves using guest molecules with special properties to change the chemical properties and/or nanoconfinement effect of N‐COFs. Of note, while these three regulation methods also may change other properties of the N‐COFs to facilitate ion transport, each strategy plays a dominant role in its respective regulation process.

As previously mentioned, regulation of the polar functional group in N‐COFs is capable of optimizing the hopping sites and hopping resistance of alkali metal ions as well as the hydrogen bonding network of proton transport [13, 113], ultimately boosting the ion transport performance. As for non‐ionic N‐COFs, α‐cyclodextrin (CD) group diminishes the stacking of 2D layers to optimize the proton hopping sites in CD‐TpAzo@H₃PO₄ [135], 1,2,4‐triazole group adsorbs the guest anion to weaken the proton hopping resistance in DAAQ‐COF (DAAQ: 2,6‐diaminoanthraquinone) [134], and all nitrogen units generates strong interlayer interactions and H₃PO₄ adsorption to enhance the proton hopping sites in 3D TPT‐COF (triazine‐corebased COF) (Figure 9a‐i) [136]. Therefore, high proton conductivity is achieved, such as 0.78 S cm⁻¹ for CD‐TpAzo@H₃PO₄ at 150°C and 1.27 × 10⁻² S cm⁻¹ for H₃PO₄@TPT‐COF at 160°C (Figure 9a‐ii) under anhydrous condition. Despite these impressive ionic conductivities, the improved performance in each case is primarily a function of guest loading rather than an intrinsic transformation of the ion transport mechanism. For ionic N‐COFs, CF₃‐Li‐ImCOF and CH₃‐Li‐ImCOF serve as the example, the ‐CF₃ with strong electron‐absorbing ability will delocalize the negative charge on nitrogen atoms in CF₃‐Li‐ImCOF, which weakens the ion pairing between Li⁺ and the anion framework. On the contrary, the ‐CH₃ with electron‐donating property has a poor effect on assigning negative charge in CH₃‐Li‐ImCOF, which will strengthen the ion pairing between Li⁺ and the anion framework. This functional group difference leads CF₃‐Li‐ImCOF/PC exhibits high Li⁺ conductivity than CH₃‐Li‐ImCOF/PC (e.g., 7.2 × 10⁻³ S cm⁻¹ for CF₃‐Li‐ImCOF/PC and 8.0 × 10⁻⁵ S cm⁻¹ for CH₃‐Li‐ImCOF) (Figure 9b) [61]. This result establishes a clear design principle that for ionic N‐COFs, electron withdrawing substituents are not only beneficial, but also crucial for achieving actual cation conductivity, owing to they directly reduce the dissociation energy of ion pairs. In addition to regulating the types of polar functional groups, the number of polar functional groups on the framework should also be considered when regulating the ionic conductivity of N‐COFs. Because, in general, the more polar functional groups in COFs, the higher their corresponding ionic conductivity [137, 138]. This phenomenon has already been demonstrated by PEO‐functionalized hydrazone‐linked COFs, with an increase in the PEO chain that is capable of solvating Li⁺, the number of Li⁺ transport sites within the COFs is also increased. Thus, the Li⁺ conductivity of COF‐PEO‐3‐Li, COF‐PEO‐6‐Li, and COF‐PEO‐9‐Li is 9.72 × 10⁻⁵ S cm⁻¹ 3.71 × 10⁻⁴ S cm⁻¹ and 1.33 × 10⁻³ S cm⁻¹ at 200°C, respectively [139].

FIGURE 9.

Regulation of the nano COFs. (a) 2D TPT‐based non‐ionic N‐COFs. (i) Schematic illustration of the growth of 3D TPT‐COF. (ii) The ionic conductivity of H₃PO₄@TPT‐COF. Reproduced with permission [136]. Copyright 2022, Wiley‐VCH. (b) Anion based 2D CF₃‐Li‐ImCOF and 2D CF₃‐Li‐ImCOF. (i) The molecular structure diagram of CF₃‐Li‐ImCOF and CF₃‐Li‐ImCOF. (ii) The ionic conductivity of CF₃‐Li‐ImCOF and CF₃‐Li‐ImCOF. Reproduced with permission [61]. Copyright 2019, American Chemical Society. (c) Anion based 2D LiCON‐2 and 2D LiCON‐3. (i) The molecular structure diagram of 2D LiCON‐2 and 2D LiCON‐3. (ii) The ionic conductivity of 2D LiCON‐2 and 2D LiCON‐3. Reproduced with permission [37]. Copyright 2020, American Chemical Society. (d) 2D ziCOFNs@Li. (i) The structure of ziCOFNs. (ii) The ionic conductivity of ziCOFNs@Li. Reproduced with permission [140]. Copyright 2024, Wiley‐VCH. (e) 2D TpPa‐SO₃H with short hydrogen bond network. (i) Proton transport via short‐H‐bond and normal H‐bond in the confined water domain. (ii) Proton transport in different separated water domains with/without short‐H‐bond. (iii) The probability density distribution of H₃O⁺ in TpPa‐SO₃H. (iv) The ionic conductivity of TpPa‐SO₃H. Reproduced with permission [41]. Copyright 2022, Springer Nature. (f) PIL decorated m‐TpPa‐SO₃H. (i) The molecular structure diagram of PIL0.5@m‐TpPa‐SO₃H. (ii) The molecule format of PIL0.5@m‐TpPa‐SO₃H. (iii) The proton conduction mechanism of PIL0.5@m‐TpPa‐SO₃H. (iv) The conductivity of PIL0.5@mTpPa‐SO₃H. Reproduced with permission [141]. Copyright 2023, Wiley‐VCH.

Regulation of the ionic functional groups in ionic N‐COFs is divided into ion type regulation and ion number regulation, in which the regulation of ion types aims to weaken the Coulomb interaction between the ionic framework and guest anions or promote the dissociation between the ionic framework and cations. For example, after the lithiation of ‐COOH modified HCON‐2 and ‐SO₃H modified HCON‐3 that date back to SD‐COF‐1 (Figure 9c‐i,c‐ii), LiCON‐2 and LiCON‐3 exhibit different Li⁺ conductivity, such as 4.36 × 10⁻⁶ S cm⁻¹ for LiCON‐2 and 3.21 × 10⁻⁵ S cm⁻¹ for LiCON‐3 (Figure 9c‐iii) [37], in which the higher Li⁺ conductivity of LiCON‐3 originates from the strong acidity of ‐SO₃ ⁻, which features a high dissociation constant with Li⁺. Similarly, when ‐COOH is introduced in the cationic COF (DhaTG_ClCOF), the locally‐zwitterionic covalent organic framework nanosheets (Dha_COOLiTG_NO3COF) utilize the abundant ‐COO⁻ units to accelerate Li⁺ dissociation (Figure 9d‐i), endowing Dha_COOLiTG_NO3COF@LiTFSI to delivers the higher ionic conductivity compared with DhaTG_ClCOF (0.15 mS cm⁻¹ for the former and 0.029 mS cm⁻¹ for the latter) (Figure 9d‐ii) [140]. Although these enhancements for ionic functional groups are important, they are fundamentally limited by the acid‐base balance of fixed functional groups in N‐COFs. They represent ionic conductivity optimizations within known mechanisms, rather than paradigm shifts.

Differing from the ion type regulation, the regulation of the ion numbers in anionic N‐COFs could optimize the number of proton/Li⁺ donors, ionic interaction sites, and the structure of the hydrogen bonding network in nanochannels. In the derivatives of TpPa‐SO₃H nanosheets, adjusting the number of ‐SO₃H ligands in their nanochannels enables the formation of a localized short hydrogen bond (SHB) network in TpBd‐SO₃H. Specifically, water molecules form localized H₃O⁺ regions around ‐SO₃ ⁻ (Figure 9e‐i,e‐ii). This SHB is different from covalent bonds and ordinary hydrogen bonds, which enhances the proton delocalization between donor and acceptor molecules, thereby achieving ultra‐fast proton transport, such as 1389 mS cm⁻¹ at T = 90°C and RH = 100% (Figure 9e‐iii,e‐iv) [41]. This value represents a true leap beyond the performance range of traditional N‐COFs and emphasizes the crucial role of ion spacing in determining the hydrogen bonding network structure. Apart from this example, when increasing –SO₃ ⁻ in LiO₃S‐COF, more easily accessible ion sites and more lubricated nanochannels are established in LiO₃S‐COF2, leading LiO₃S‐COF2 to exhibit higher ionic conductivity than LiO₃S‐COF1 (4.46 × 10⁻⁵ S cm⁻¹ for LiO₃S‐COF2 and 2.67 × 10⁻⁵ S cm⁻¹ for LiO₃S‐COF1) [123]. As for cationic N‐COFs, the increase in ion number can alleviate the electrostatic repulsion and steric hindrance in nanochannels, leading COF‐SDQA (DQA: disubstituted hydrazide monomers) with increased quaternary ammonium groups to deliver OH⁻ conductivity of 329.4 mS cm⁻¹ at 80°C and 100% RH) [137]. Taken together, these examples demonstrate that ion number optimization can induce qualitative changes in ion transport behavior when the critical density threshold for hydrogen bond network or SHB formation is exceeded.

Guest molecular insertion constitutes a complementary and simpler post‐synthesis strategy, while expanding the effective channel size and customizing the local chemical environment. aramid nanofibers, PEG [130], graphene quantum dots [142], and ionic liquids [141], have been the primary intercalants explored. The modification of TpPa‐SO₃H with a polymeric ionic liquid (PIL) creates a hydration layer along the 1D channels, which cooperates with –SO₃H groups to form a dense hydrogen‐bonding network (Figure 9f‐i‐f‐iii), yielding a proton conductivity of 1.02 × 10⁻¹ S cm⁻¹ at 90°C and 100% RH (Figure 9f‐iv) [141]. While these guest‐mediated enhancements are substantial and synthetically straightforward, they inherently introduce a non‐covalent component in N‐COFs. The long‐term retention of guest molecular under operating conditions remains a concern. Therefore, this method is most appropriately viewed as a rapid optimization route for concept demonstrations rather than a definitive solution for durable nano ion conductors.

In addition to the above dominant regulation methods, the pore size adjustment can accommodate more ions within directional nanochannels of N‐COFs, thus promoting the ion transport process. However, pore size regulation does not always lead to a high ion transport rate, because a larger pore size will cause ions to move and accumulate on one side of the pore walls under the action of the electric field, thereby reducing the ionic conductivity [134]. Similarly, the side chain length of functional groups also has no specific influence law on the ion conduction properties of N‐COFs. The shorter side chain functional groups make N‐COFs have a smaller spacer length, and the longer side chain lead N‐COFs to feature small water content and large dissociation energy [143]. Therefore, when screening the pore size or functional group side chains for N‐COFs, the multi perspective ion transport effects (e.g., steric accessibility, solvation, and electrostatic landscape) should be taken into account.

In brief, among the above regulation methods, guest molecules are the fastest, simplest, and very effective way to regulate the ion transport performance of N‐COFs. Notwithstanding, the variety of guest molecules used so far in N‐COFs remains very limited. Theoretically, if the nanochannels of N‐COFs have a larger size than external molecules, these functional molecules will be inserted into the N‐COFs to improve their ion transport performance. Therefore, more efforts should be focused on the search for functional guest molecules to comprehensively improve the ion transport performance of N‐COFs.

3.3. Design of Organic‐Inorganic Hybrid Nano Ion Conductors

Organic‐inorganic hybrid nano ion conductors mainly refer to N‐MOFs formed by metal center and organic ligand under solvothermal conditions [144], which usually presents typical 2D nanosheet and 3D polyhedral morphology. Unlike the MOFs with ion conduction function in mainstream crystal structure design, whose dimension refers to the crystal structure field, MOFs with nano morphology in shape scale as ion conductors have only attracted attention in recent years. Besides, compared with the MOFs in the field of crystal structure category, N‐MOFs have the following advantages for ion transport: (1) The nanosizing of MOFs shortens the directional ion transport channels, thereby enhancing the ion transport rate [145]; (2) The nanoscale MOFs feature large specific surface area and can accommodate more guest species, which enables them to become the hybrid ion‐rich conductors [146, 147]. Meanwhile, nanosizing improves the contact between the electrode and ion conductors, and optimizes interfacial ion transport by reducing the particle size of MOFs, which is important for the design of practical devices [62]; (3) The specific morphology design could affect the physicochemical properties within nanochannels of N‐MOFs, thus influencing the ion conduction pathways [148]. Beyond these advantages, the highly tunable pore structure and functional groups in N‐MOFs provide the possibility to regulate their ion transport behavior [148, 149]. Similar to the aforementioned 2D inorganic nanosheets and N‐COFs, the development of N‐MOFs is also divided into developing N‐MOFs with fast ion conduction function and regulating the ion transport performance of N‐MOFs. Nevertheless, compared with developing N‐MOFs based on ion conductors, regulating the ion conduction performance of N‐MOFs is more challenging, which is mainly due to the limited ligands and modification methods as well as the difficulties in morphological design.

3.3.1. Development of N‐MOFs

The development of N‐MOFs includes designing novel ionic N‐MOFs and non‐ionic N‐MOFs, which purpose is to achieve rapid proton conduction and alkali metal ion conduction. Ionic N‐MOFs are composed of an ionic nano framework and counterions, which mainly relies upon the ion effects in the directional nanochannel to achieve or assist the rapid transport of targeted ions [13]. However, for non‐ionic N‐MOFs, it primarily depends on polar functional groups, unsaturated coordination metal sites, and nanoconfinement effects in the directional nanochannel to accomplish guest ion transport [32].

At present, the development of non‐ionic N‐MOFs in the ion conduction field mainly concentrates on selecting suitable metal centers and organic ligands to synthesize MOFs with nano morphology. Owing to the ion‐free characteristics of most non‐ionic N‐MOFs, their frameworks generally lack the charge carriers required for autonomous ion conduction. Therefore, the rapid proton transport needs the assistance of external humidity or guest acidic substances, while the fast alkali metal ion transport needs to the encapsulation of liquid electrolytes or alkali metal salt. Concretely speaking, under humidified atmospheres, coordinated water molecules in non‐ionic N‐MOFs will interact with the functional groups (e.g., ‐OH and ‐NH₂) to produce a certain amount of protons and hydrogen bonding networks [33], then these protons will be transported on the formed hydrogen bonding networks (Figure 10a‐i). Representative non‐ionic N‐MOFs, including 3D Cr‐MIL‐88B [145] and PCMOF‐5 [150], deliver the proton conductivity of 1.0–7.0 × 10⁻³ S cm⁻¹ at T = 25–100°C and RH = 85∼98% under the influence of generated proton and hydrogen bonding networks (Figure 10a‐ii). The similarity of these proton conductivities across different MOFs reveals that the performance plateau of non‑ionic MOFs is determined by the general characteristics of the confined water network (i.e., proton concentration and hydrogen bonding network). In other words, the modification of proton conductivity in non‑ionic MOFs is constrained by the moderate pore wall acidity and limited water absorption, rather than by framework characteristics. Consequently, no original non‐ionic N‐MOFs exceeds the conductivity threshold of 10⁻² S cm⁻¹, and the proton conductivity cannot break this upper limit with incremental changes in pore size or cluster geometry of non‐ionic MOFs. Meanwhile, under the synergistic effect between functional groups of non‐ionic N‐MOFs and acidic polymers, the as‐constructed SPEEK/MOF‐808 (SPEEK: non‐fluorinated sulfonated poly (ether ether ketone) [151] (Figure 10b) exhibits the high conductivity of 0.0766 S cm⁻¹ at high temperature and high humidity. Although composite non‐ionic N‐MOFs exhibit significant synergistic effects to improve proton conductivity, the specific values are still inferior to state‐of‐the‐art perfluorinated sulfonic acid membranes (0.1 S cm⁻¹). Therefore, composite non‐ionic N‐MOFs with acidic polymers only represent an important progressive advancement that overcomes the proton conductivity limitations of pure non‐ionic N‐MOFs, rather than establishing a new H⁺ transport paradigm.

FIGURE 10.

Development of the non‐ionic N‐MOFs. (a) The transport of H⁺ in non‐ionic N‐MOFs. (i) The transport of H⁺ in non‐ionic N‐MOFs with external humidity. (ii) The conductivity range of non‐ionic N‐MOFs with external humidity. (b) MOF‐808 based nano ion conductors. (i) Schematic diagram of proton transport in MOF‐808. (ii) The conductivity of SPEEK/MOF‐808. Reproduced with permission [151]. Copyright 2021, Wiley‐VCH. (c) The transport of proton under acidic environments or the transport of alkali metal cation in non‐ionic N‐MOFs with/without polar functional groups. (d) 3D‐UIO‐66 based nano ion conductors. (i) Schematic illustration of Li⁺ transport through 3D‐UIO‐66/PAN/PEO/LiTFSI. Reproduced with permission [154]. Copyright 2022, Elsevier. (ii) Schematic illustration of 3D‐UIO‐66‐D4. (iii) The ionic conductivity of LiTFSI@UiO‐66‐D4. Reproduced with permission [155]. Copyright 2025, American Chemical Society. (e) UiO‐66‐FV based nano ion conductors. (i) The structure diagram of copolymerized UiO‐66‐FV. (ii) The conductivity of LiTFSI+LiDFOB@UiO‐66‐FV. Reproduced with permission [147]. Copyright 2026, Wiley‐VCH. (f) Schematic illustration of Li⁺ transport in LiClO₄+PC+HKUST‐1. Reproduced with permission [156]. Copyright 2023, Wiley‐VCH. (g) UIO‐67 based nano ion conductors. (i) The structure diagram of ([EMIm]Cl):AlCl₃@UIO‐67. (ii) The ionic conductivity of ([EMIm]Cl):AlCl₃@UIO‐67. Reproduced with permission [157]. Copyright 2022, Wiley‐VCH. (h) MOF‐2OH based ion conductors. (i) The Li⁺ transport mechanism inside the nanopores of MOF‐2OH@DMF@Li. (ii) The ionic conductivity of LiClO₄‐DMF@Cu‐MOF‐2OH. Reproduced with permission [158]. Copyright 2026, Creative Commons Attribution‐Noncommercial. (i) PCN‐777 based nano ion conductors. (i) The structure diagram of (EMI)[N(CN)₂]@PCN‐777. (ii) The ionic conductivity of (EMI)[N(CN)₂]@PCN‐777. Reproduced with permission [159]. Copyright 2019, Wiley‐VCH. (j) CuBTC‐PSS based nano ion conductors. (i) The structure and Li⁺ transport diagram of Li‐PC@CuBTC‐PSS. (ii) The conductivity of Li‐PC@CuBTC‐PSS. Reproduced with permission [62]. Copyright 2022, Springer Nature.

Apart from the water environment, when the acidic substance is encapsulated into non‐ionic N‐MOFs, the hybrid ion conductors with fast proton transport can also be constructed depending on the coordination and nanoconfinement effects of the guest acid groups in directional nanochannels [152] (Figure 10c). A striking example is the encapsulation of ionic liquid into the directional nanochannels of 3D MIL‐101 polyhedron, protonated‐SO₃H/HTFSA and unprotonated ‐SO₃ ⁻/TFSA⁻ of EIMS‐HTFSA (HTFSA: N,N‐bis(trifluoromethanesulfonyl)amide. EIMS: 1‐(1‐ethyl‐3‐imidazolium)propane‐3‐sulfonate) form an orderly ion transport channel, enabling EIMS‐HTFSA@MIL‐101 with the anhydrous conductivity of 2 × 10⁻⁴ S cm⁻¹ at 140°C, and this conductivity is 6 orders of magnitude higher than MIL‐101 [153]. Despite the significant improvement of proton conductivity compared with pure non‐ionic N‐MOFs, the absolute conductivity remains modest and falls below that of many sulfonated polymers or ionic COFs under similar anhydrous conditions. This indicates that nanoconfinement cannot compensate for the low H⁺ carrier concentration in acid‑loaded non‑ionic MOFs. Therefore, the primary bottleneck for non‐ionic N‐MOFs remains the introduction of a high density of acidic functional groups on their framework to create more continuous proton pathways.

Unlike the proton conduction in non‐ionic N‐MOFs, the quasi‐solid/solid ion conductors, formed by encapsulating liquid electrolytes or lithium salts into non‐ionic N‐MOFs, mainly rely on the defect/functional group coordination effect, porous enrichment effect, and nanopore size confinement effect to rapidly transport cations. The coordination effects rely on the open metal sites or grafted functional groups to immobilize anions, (e.g., 3D UIO‐66 [154] (Figure 10d‐i) and, 3D UiO‐66‐D4 [155] (Figure 10d‐ii), thereby weakening the Li⁺ coordination environment and enabling the rapid Li⁺ transport. However, the conductivity achieved solely through coordination rarely exceeds the 10⁻³ S cm⁻¹ (e.g., 1.42 × 10⁻³ S cm⁻¹ for LiTFSI@UiO‐66‐D4 at around 30°C, Figure 10d‐iii), as the number of active anion binding sites is limited by the framework stoichiometry in non‐ionic N‐MOFs.

The porous enrichment effect capitalizes on the high internal surface area and directional nanochannels of non‐ionic N‐MOFs to increase the density and continuity of the encapsulated guest ion‐conducting phase. In 3D Uio‐66 [147] and HKUST‐1 [156], the abundant directional nanochannels provide a great number of conduction pathways for Li⁺ (Figure 10e,f), causing LiTFSI+LiDFOB@UiO‐66‐FV and LiClO₄+PC+HKUST‐1 possess the Li⁺ conductivity of 6.8 × 10⁻⁴ and 1.02 × 10⁻³ S cm⁻¹, respectively. These ionic conductivity reveals that cation conductivity under porous enrichment effect equals that of bulk liquid electrolytes, confirming the non‐ionic N‐MOFs as an inert host to provide mechanical rigidity and avoid leakage, the improvement of cation conductivity originates from the guest ion‐conducting phase. Meanwhile, for the transport of multivalent ions, the sufficient directional nanochannels in UIO‐67 also provide transport paths for Al_xCl_y ⁻ anions (Figure 10g‐i), resulting ([EMIm]Cl):AlCl₃@UIO‐67 offers the Al³⁺ conductivity of 7.5 × 10⁻⁴ S cm⁻¹ at 20°C (Figure 10g‐ii) [157]. Although this value is noteworthy for multivalent ion conductors, it is still one order of magnitude lower than the Li⁺ conductivity achieved in similar systems, highlighting the big challenge in transporting highly charged ions under porous enrichment effect.

The nanopore size confinement effect constitutes a third pathway for ion transport in non‐ionic N‐MOFs, which is operated by restricting the spatial distribution and dynamic behavior of encapsulated ions. Cu‐MOF‐2OH can confine the N,N‐dimethylformamide (DMF) in their nanopore channels and form a long‐range ordered DMF layer (Figure 10h), which increases the ionic conductivity of LiClO₄‐DMF@Cu‐MOF‐2OH to 2.87 × 10⁻⁴ S cm⁻¹ and Li⁺ transference number to 0.81 [158]. Although the Li⁺ transference number is increased, the conductivity remains an order of magnitude below that of the conventional bulk liquid electrolyte, indicating that pure size/solvent‐sieving cannot obviously accelerate Li⁺ transport. A more dramatic phenomenon of nanopore size confinement is observed in 2D PCN‐777, which can enable ionic liquids to form bulk‐like regions within their pores (Figure 10i‐i); achieving the superconductive behavior for (EMI)[N(CN)₂]@PCN‐777 (e.g., 4.4 × 10⁻³ S cm⁻¹ at 299.15 K and > 10⁻² S cm⁻¹ above 343 K) (Figure 10i‐ii) [159]. This transition of ionic liquids in nanopore size represents a genuine breakthrough of confinement engineering, as it demonstrates that the conductivity of encapsulated electrolytes can increase by tuning the pore architecture.

Note here that sometimes the porous enrichment effect is more important for ion transport than the pore size confinement effect, due to the large pore aperture can embrace more guest lithium metal salt. Such phenomena is validated by the highest Li⁺ conductivity of MOF‐74/LiTFSI/DME (1.73 × 10⁻⁴ S cm⁻¹) in comparison with HKUST‐LiTFSI/DME (0.44 × 10⁻⁴ S cm⁻¹) and MOF‐5/LiTFSI/DME (0.53 × 10⁻⁴ S cm⁻¹), in which the pore size for MOF‐74, HKUST‐1, and MOF‐5 are 10, 6.9, and 8 Å respectively [160]. The above result implies that the larger the pore size, the higher the conductivity, because the free space of non‐ionic N‐MOFs determines how many charge carriers can be loaded. But the highest conductivity is only 10⁻⁴ S cm⁻¹, indicating that simply enlarging the pore size of non‐ionic N‐MOFs is not enough, and the interaction between non‐ionic N‐MOFs and electrolyte needs to be allowed to help the ion dissociation.

The integration of nanopore size confinement effect and functional group coordination effect produces a synergistic effect to enhance ion transport, In CuBTC‐PSS, the sub‐nanoconfinement of nanopores on LiTFSI/PC and the coordination effects of unsaturated metal sites on TFSI⁻ afford Li‐PC@CuBTC‐PSS to generate aggregated regions with fast ion transport (Figure 10j‐i), so that Li‐PC@CuBTC‐PSS provides the high conductivity of 4.47 mS cm⁻¹ at 25°C (Figure 10j‐ii) [62]. Such synergistic conductivity has reached the top level of non‐ionic N‐MOFs for Li⁺ transport and is close to ionic N‐MOFs. This illustrates a universal design principle for fast non‐ionic N‐MOFs, the confinement effect enhances the coordination effect by maintaining high concentrated active sites, while the coordination effect provides the anion binding function that the confinement effect lacks. Both are indispensable for generating high ionic conductivity.

All in all, encapsulating external salts or electrolytes into non‐ionic N‐MOFs can indeed utilize their special properties to achieve fast ion transport, but special attention should be paid to the following points when developing non‐ionic N‐MOFs with ion conduction function in the future. (1) Functional groups that can generate proton and adsorb anions should be introduced into non‐ionic N‐MOFs as much as possible to produce ion transport behavior, aiming to boost the mobility of proton and alkali metal ions. (2) The suitable ligands should be chosen to improve the porous enrichment effect and nanopore size confinement effect, intending to optimize the ion transport nanochannels. (3) Emphasizing the encapsulation amount of guest substances in non‐ionic N‐MOFs, because the content and gap filling of MOFs may have a significant influence on the nano flow and ion bundling [161], which significantly influences the ion conductivity. Ultimately, incorporating guest molecules into non‐ionic N‐MOFs is indeed a simple and universal method for creating functional ion conductors. But to achieve a true breakthrough in ion transport for such materials, non‐ionic N‐MOFs cannot just be used as passive containers, but must actively participate in ion transport, such as developing organic linkers that can sense external stimuli (e.g., voltage and light) to switch or regulate ion transport.

Depending on how the ions are produced in MOFs, the design of ionic N‐MOFs is classified into in situ design strategy, ionic functional group grafting strategy, and guest ion pairing strategy. The in situ design strategy refers to the usage of ionic organic ligands to synthesize ionic N‐MOFs, and the ionic functional group grafting strategy is to chemically graft ions onto non‐ionic N‐MOFs. In contrast, the guest ion pairing strategy utilizes a large number of polar functional groups on the frameworks to interact with guest anions to construct pseudo ionic N‐MOFs, thereby forming anionic‐oriented nanochannels with weak interaction to transport ions.

The in situ design of ionic N‐MOFs is based on the direct incorporation of charged functional groups into the framework backbone. Anionic N‐MOFs are typically synthesized from organic ligands with special groups (e.g., ‐SO₃ ⁻, (POO)²⁻, COO^‐, and TFSI⁻), while cationic N‐MOFs are constructed from organic ligands with cyclic cation. This strategy generates abundant mobile counterions or proton sites within the nanochannels, thereby circumventing the carrier‐deficiency limitation that is encountered by non‐ionic N‐MOFs. For proton conduction in anionic N‐MOFs, ‐SO₃H on the anionic framework of N‐MOFs not only exerts its high‐water adsorption capacity to form the hydrogen bonding (H₂O−H₃O⁺) network for proton transport, but also capitalizes ‐SO₃H as a proton donor to provide an abundant proton source [145, 146] (Figure 11a–i). The superiority of such dual effects of ‐SO₃H is reflected in the high proton conductivity of anionic N‐MOFs. MIL‐88B‐PESA [145] and MIL‐101(Cr)‐NH₂‐SO₃H [146] exhibit the proton conductivities of ∼10⁻⁴–10⁻¹ S cm⁻¹ at T = 70°C–100°C and RH = 85‐98% (Figure 11a‐ii). Notably, these proton conductivities are one to three orders of magnitude higher than those of non‐ionic N‐MOFs under identical conditions, which demonstrates ‐SO₃H greatly improves the bottom line of proton transport that is hardly achieved by regulation of the non‐ionic N‐MOFs. The most notable example of this principle is H₂SO₄@MIL‐101‐SO₃H, which delivers the proton conductivity of 1.82 S cm⁻¹ at 70°C and 90% RH under the cooperative effect of ‐SO₃H and guest H₂SO₄ with the hydrogen bonding formation and proton replenishment effect [146]. This value exceeds Nafion under the same conditions and is also one of the highest proton conductivities in crystalline porous materials, thus setting a clear upper limit for sulfonic acid group design in anionic N‐MOFs. Zwitterionic based N‐MOFs represent a further structural refinement of ionic N‐MOFs, in their internal framework, ‐SO₃H forms a unique long‐range ordered hydrogen bonding network by the dense ‐SO₃ ⁻ in the confined nanochannel, which reduces the activation energy for proton hopping (Figure 11b‐i) [162]. The resulting zwitterionic based N‐MOF with [MIMS][MSA] (EIMS: 1‐(1‐ethyl‐3‐imidazolio)propane‐3‐sulfonate, MSA: methanesulfonic acid) delivers the proton conductivity of 2.87 × 10⁻³ and 1.03 × 10⁻⁴ S cm⁻¹ at 20°C and ‐40°C, respectively (Figure 11b‐ii). Although this low‐temperature performance is relatively small, the zwitterionic framework can still conduct protons compared to traditional MOFs that cannot conduct ion when water freezes, providing a new idea for low‐temperature proton conductors.

FIGURE 11.

Development of the ionic N‐MOFs under in situ design. (a) The H⁺ transport in ionic N‐MOFs. (i) The H⁺ transport in ionic N‐MOFs with external humidity. (ii) The proton conductivity of some reported ionic N‐MOFs. (b) IL₁MOF. (i) The long‐rang ordered framework of IL₁MOF. (ii) The proton conductivity of IL₁MOF. Reproduced with permission [162]. Copyright 2021, Wiley‐VCH. (c) The dynamic structural changes of H₁₂‐M₂‐(DOBDP)₃ in the c‐axis upon dehydration. Reproduced with permission [163]. Copyright 2025, Elsevier. (d) BUT‐8(Cr)A. (i) Schematic diagram of the structural flexibility of BUT‐8(Cr)A. (ii) The structure of BUT‐8‐8(M) (M = Cr, Al). (iii) The proton conductivity of BUT‐8(Cr)A. Reproduced with permission [164]. Copyright 2017, Springer Nature. (e) D‐UiO‐66‐NH₂. (i) The structure of D‐UiO‐66‐NH₂. (ii) The ionic conductivity of D‐UiO‐66‐NH₂ based ion conductor. Reproduced with permission [165]. Copyright 2026, Wiley‐VCH. (f) MIL‐101‐SO₃Na. (i) The structure of ILs+Na salts@MIL‐101‐SO₃Na. (ii) The ionic conductivity of ILs+Na salts@MIL‐101‐SO₃Na. Reproduced with permission [63]. Copyright 2019, Wiley‐VCH.

A key weakness of the rigid ionic N‐MOFs mentioned above is that they require an excessively high hydration state. In the rigid framework, the spatial distribution of acidic groups is fixed by the lattice. Once the humidity is reduced, the hydrogen bonding network collapses, leading to a sharp decrease in proton conductivity. This limitation has motivated the development of flexible ionic N‐MOFs, whose lattices can undergo dynamic and environment‐responsive deformations to maintain the continuous hydrogen bonding network even under harsh working conditions (Figure 11c) [163]. The structural adaptability of 3D BUT‐8(Cr) [164] with ‐SO₃H and H₁₂‐M₂‐(DOBDP)₃ [163] with ‐POOH originates from the carboxylate and uncoordinated phosphonate function groups, which allows their microstructure to undergo favorable deformations to maintain a continuous and abundant hydrogen bonding networks at the harsh working environment (Figure 11d‐i), enabling fast proton transport. Therefore, 3D BUT‐8(Cr) deliver with the proton conductivity of 1.27 × 10⁻¹ S cm⁻¹ at 100% RH and 80°C, and 6.32 × 10⁻³ S cm⁻¹ at 65% RH and 80°C (Figure 11d‐ii,d‐iii) [164]. While the latter proton conductivity is 20 times lower than the former, it is still hundreds of times higher than that of rigid sulfonic acid based MOFs. This indicates that micro framework flexibility is currently the best means to overcome high humidity dependence and achieve proton conduction.

As for the alkali metal ion conduction under the in situ design in ionic N‐MOFs, cations on the backbone of cationic N‐MOFs could adsorb the guest anions, thereby weaking the interaction between guest anions and cations and promoting the alkali metal ion transport. This mechanism is directly analogous to that established for cationic N‐COFs. In D‐UiO‐66‐NH₂, the pyridine cations in directional nanochannels can absorb the TFSI⁻ to weaken the interaction between TFSI⁻ and Li⁺ (Figure 11e‐i), enabling D‐UiO‐66‐NH₂ based ion conductor exhibits the ionic conductivity of ∼1.5 × 10⁻⁴ S cm⁻¹ (Figure 11e‐ii) [165]. This ionic conductivity, even under such polymer electrolyte‐facilitated conditions, remains comparable to those of typical cationic COFs, fundamentally highlighting the insufficient density and accessibility of active pyridine sites. In contrast, cations inside anionic N‐MOFs can dissociate from the framework anions via acid–base dissociation, thereby enabling alkali metal ion transport. But given the intrinsically ion‐poor nature and strong ionic interactions in most ionic frameworks, anionic N‐MOFs generally require combination with organic electrolytes to achieve practical ion conduction, following a mechanism similar to that of anionic 2D COFs. Representative systems, including N‐UiO‐2COOLi [166] and Zr‐BPDC‐2SO₃M [167] with LiTFSI/PC, output the Li⁺ conductivity of 1.05 × 10⁻⁵ and 1.65 × 10⁻⁴ S cm⁻¹ at 25°C, respectively. These anionic N‐MOFs with different structures exhibit small conductivity when loaded with the same electrolyte, indicating that the pores of the anionic N‐MOFs play a major role, rather than the interaction between the framework and ions. This also reflects that the ion transport performance is mainly determined by the liquid electrolyte within the framework.

A critical operational parameter, which is often overlooked in electrolyte‐encapsulated ionic N‐MOFs, is the amount of guest substances. The encapsulated amount influences not only the ionic conductivity but also the leakage and combustibility of the devices, both of which are decisive for practical safety. As an example, ILs+Na salts@MIL‐101‐SO₃Na demonstrates that adjusting ILs content in MOFs can improve the pore‐filling faction to modify the ion transport in the nanochannels (Figure 11f‐i), endowing ILs+Na salts@MIL‐101‐SO₃Na with the pore‐filling faction of 96.94% to deliver the high Na⁺ conductivity (1.32 × 10⁻² S cm⁻¹ at 50°C, Figure 11f‐ii) [63]. This result indicates that almost filling the voids of ionic N‐MOFs with liquid substances can not only maximize the proportion of conductive substances, but also maintain the integrity and safety of the N‐MOFs skeleton structure, which presents the optimal balance between improving conductivity and practical usability.

The ionic functional group grafting strategy refers to graft guest groups (e.g., ‐SO₃H, ‐SH, ‐OH, and TFSI⁻) into the frameworks of non‐ionic N‐MOFs, their ion‐conduction types and mechanisms are identical to those of the in situ synthesized ionic N‐MOFs. For proton transport, if ‐SH and ‐SO₃H are grafted onto the 3D UiO‐66 octahedron (Figure 12a‐i,a‐ii) [168], the formed UiO‐66(SH)₂ and UiO‐66‐(SO₃H)₂ will exhibit more proton donors, and adsorb more water to reorganize hydrophilic regions with fast proton transport pathway in the confined nanochannels, significantly improving the proton conductivity, e.g., 8.4 × 10⁻² S cm⁻¹ at 80°C and 90% RH for UiO‐66‐(SO₃H)₂ (Figure 12a‐iii). This value is close to the best in situ synthesized ionic N‐MOFs, suggesting that high‐density functional group grafting can be comparable to the directly assembled ionic N‐MOFs. Its advantage are the wide applicability, simple ligand, and rapid screening of pore chemistry.

FIGURE 12.

Development of the ionic N‐MOFs under ionic functional group grafting strategy and guest ion pairing strategy. (a–c) Ionic functional group grafting strategy. (d, e) Guest ion pairing strategy. (a) UiO‐66(SH)₂ and UiO‐66(SO₃H)₂. (i) The structure of UiO‐66(SH)₂. (ii) The structure of UiO‐66(SO₃H)₂. (iii) The proton conductivity of UiO‐66(SO₃H)_2. Reproduced with permission [168]. Copyright 2015, Wiley‐VCH. (b) UiO‐LiTFSI. (i) The structure and ion transport mechanism of UiO‐LiTFSI. (ii) The ionic conductivity of UiO‐LiTFSI. Reproduced with permission [169]. Copyright 2019, American Chemical Society. (c) UiO‐66‐LiSS. (i) Li⁺ transport mechanism in UiO‐66‐LiSS. (ii) The EIS of UiO‐66‐LiSS/PC+EC. Reproduced with permission [170]. Copyright 2020, American Chemical Society. (d, e) Pseudo ionic N‐MOF based ion conductor under guest ion pairing strategy. (d) UiO‐66 with OMSs [171]. (i) The structure of UiO‐66 (Hf). (ii) The adsorption structures of LiTFSI on perfect UiO‐66 and UiO‐66 with OMSs. (iii) The ionic conductivity of UiO‐66@LiTFSI/DME‐IL. Reproduced with permission [171]. Copyright 2025, American Chemical Society. (e) ZIF‐67 with OMSs [172]. (i) The design of quasi‐ZIF‐67 with OMSs. (ii) The coordination structure of ZIF‐67 and quasi‐ZIF‐67 with OMSs. (iii) The ionic conductivity of solid polymer electrolyte with quasi‐ZIF‐67. Reproduced with permission [172]. Copyright 2025, Wiley‐VCH.

For alkali metal ion transport, similar to the ionic N‐COFs and N‐MOFs mentioned earlier, ionic N‐MOFs with ‐OH⁻, ‐SO₃ ⁻ and TFSI⁻ after functional group grafting also exhibit the low ion transport rate due to the strong Lewis acid‐base interaction [169, 170] (Figure 12b‐i,c‐i), such as the ionic conductivity of 2.07 × 10⁻⁴ S cm⁻¹ for UIO‐66‐NH₂ polyhedron with TFSI⁻ [169] (Figure 12b‐ii). Therefore, they also should combine with the guest liquid solvents/electrolytes to weaken the Lewis acid‐base interaction, eventually achieving fast alkali metal ions transport. The efficacy of this approach is demonstrated by UiO‐66‐LiSS, where encapsulation of EC/PC promotes the dissociation of Li⁺ from ‐SO_3–Li⁺ ionic pair, leading EC+PC@UiO‐66‐LiSS with the higher ionic conductivity (7.8 × 10⁻⁴ S cm⁻¹) in comparison with UiO‐66‐LiSS (6.0 × 10⁻⁵ S cm⁻¹) at room temperature (Figure 12c‐ii) [170]. The ionic conductivity obtained by encapsulating electrolyte in the above ionic N‐MOFs is 2‐3 orders of magnitude lower than the optimal values of non‐ionic N‐MOFs and intrinsic ionic N‐COFs filled with electrolyte. This indicates that adding solvents cannot overcome the fundamental limitations caused by the strong acidity of sulfonic acid groups. In addition, although this grafting method is convenient, the generated anionic sites are not easily dissociated. Therefore, alternative strategies, which introduce more charge‐delocalized anionic species (e.g., sulfonimide‐based groups), are needed to achieve higher cation conductivity.

Pseudo ionic N‐MOFs are essentially a class of non‐ionic MOFs that have a weak interaction with the guest anions; their formation is highly correlated with the intentionally created open metal sites (OMSs) in non‐ionic N‐MOFs through heat treatment. Generally speaking, as long as a great number of OMSs are designed in the non‐ionic N‐MOFs, these OMSs will adsorb abundant guest anions to form a pseudo‐anionic framework with densely packed anions in the pores, thereby weakening the interaction between guest anions and cations, and eventually allowing pseudo ionic N‐MOFs to feature efficient ion transport pathways (Figure 12d,e) [171]. Experimentally, such a principle has been validated by multiple non‐ionic N‐MOFs, UiO‐66@LiTFSI/DME‐IL, [171], HKUST‐1‐LiClO₄ [149], PEO@quasi‐ZIF‐67@LiTFSI [172] with OMSs deliver the high Li⁺ conductivity up to ∼0.5–4 × 10⁻³ S cm⁻¹. The Li⁺ conductivity of these pseudo ionic N‐MOFs has reached the level of intrinsic ionic N‐MOFs. Meanwhile, the synthesis and stability of these N‐MOFs are relatively good. More importantly, almost all non‐ionic N‐MOFs can be transformed into pseudo ionic conductors using this method, which cannot be achieved by other strategies.

In short, the prospect of N‐MOFs capable of conducting alkali metal ions depends on the strategic trade‐off between synthesis accessibility and final performance. The chemical modification of grafted ionic N‐MOFs is flexible, but it is limited by the strong ion pair effects and requires the use of guest solvents to exert its effect, placing their conductivity in the 10⁻⁴ S cm⁻¹ range. Pseudo ionic N‐MOFs have higher conductivity (∼10⁻³ S cm⁻¹) via a non‐covalent anion capture mechanism, and this method is simpler to synthesize and have a wider range of applications. But its weakness is that unsaturated metal sites are easily poisoned by strong coordinating species. To achieve practical applications, it is necessary to simultaneously optimize the ion functional groups (e.g., type and density), pore structure (e.g., pore size, tortuosity, and connectivity), and particle morphology. In addition, computational screening should be combined to evaluate the dissociation energy of ionic groups and the binding strength of unsaturated metal sites to anions, thereby accelerating the search for the optimal N‐MOFs.

3.3.2. Regulation of the N‐MOFs

Similar to bulk MOFs, the organic ligands of N‐MOFs are highly tunable. If the derived organic ligands with high similarity to the parent counterparts are selected, then the intrinsic properties (e.g., pore size, functional group types, functional group numbers, ion types and ion numbers) in the oriented nanochannels can be optimized without changing the morphology of N‐MOFs [32]. Besides, the porosity of N‐MOFs guarantees their nanochannels to accommodate some species with special properties, which can alter the physicochemical properties of their directional nanochannels [32]. Moreover, changing the crystalline state of N‐MOFs also could change the physicochemical properties within their nanochannels owing to amorphous alters the ion conduction pathways [173]. Therefore, regulating the ion conduction behaviors of N‐MOFs can be differentiated into the following categories: (1) regulating the functional group/anion numbers in N‐MOFs. (2) adjusting the nanochannel size for N‐MOFs. (3) immersing the guest substances into N‐MOFs. (4) transforming the crystal phase of N‐MOFs. Herein, it is important to note that the regulation of N‐MOFs is not a separate regulation process, but a dynamic and comprehensive process. The specific modification of N‐MOFs only means which regulation plays a leading role in the ion transport process.

The functional group/anion number regulations, on the one hand, promote the formation of proton/Li⁺ transport networks and ion hopping sites, and accelerate the dissociation of other alkali metal ions on the other side. Thus, they are significantly crucial for ion transport. The impact of ‐SO₃ ⁻ density is most clearly illustrated by the Cr‐MIL‐88B derivatives, Cr‐MIL‐88B‐PSA with more ‐SO₃H units delivers the high proton conductivity than the Cr‐MIL‐88B‐PESA at T = 100°C and RH = 85% (Figure 13a), (e.g., 1.58 × 10⁻¹ S cm⁻¹ for Cr‐MIL‐88B‐PSA and 4.50 × 10⁻² S cm⁻¹ for Cr‐MIL‐88B‐PESA) [145]. In principle, more ‐SO₃ ⁻ should provide more proton donors and a more extensive hydrogen bonding network, enabling a significant increase in conductivity. In practice, the functional group density of the ionic N‐MOFs increased considerably, but the conductivity improved only by 3.5 folds. This indicates a nonlinear relationship between functional group density and conductivity, which may be caused by the pore crowding and counterion condensation of high‐density grafting. Among the conductive MOFs (2D cMOFs), namely M₃(HIR₃‐TAT)₂ (M = Ni, Cu; R = nBu, 1EG, 2EG), Ni₃(HI1EG₃‐TAT)₂ with 1EG side chain exhibits high Li⁺ conductivity than Ni₃(HInBu‐TAT)₂ with aliphatic nBu chain in 1 m LiClO₄ aqueous electrolyte (1.1 × 10⁻⁴ S cm⁻¹ for Ni₃(HI1EG₃‐TAT)₂ and 4.18 × 10⁻⁵ S cm⁻¹ for Ni₃(HInBu‐TAT)₂) (Figure 13b), demonstrating that group side chains in nanochannels would influence the hopping site density and H‐bonding interactions between neighboring −NH and −O‐ group [174]. More importantly, increasing the number of ‐COO⁻ can enhance the O sites for Li⁺ hopping. Therefore, UiO‐2COOLi@PC has higher ionic conductivity than that of UiO‐COOLi@PC and UiO‐66@PC [166]. This indicates that the additional ‐COO⁻ indeed provides more hooping sites and reduces the migration energy barrier. However, the ion pairing effect between ‐COO⁻ and Li⁺ is too strong, and the ionic conductivity cannot be significantly improved. This once again proves that for cationic conductive anionic N‐MOFs, anionic chemical properties are more important than site density.

FIGURE 13.

Regulation of the N‐MOFs. (a) Cr‐MIL‐88B derived Cr‐MIL‐88B‐PSA and Cr‐MIL‐88B‐PESA. (i) The structure and conductivity of Cr‐MIL‐88B‐PESA. (ii) The structure and conductivity of Cr‐MIL‐88B‐PSA. Reproduced with permission [145]. Copyright 2020, American Chemical Society. (b) The structure of M₃(HIR₃‐TAT)₂. Reproduced with permission [174]. Copyright 2025, American Chemical Society. (c) The surface electrostatic potentials of MIL‐53‐SO₃H‐X. Reproduced with permission [176]. Copyright 2025, Wiley‐VCH. (d) IM‐UiO‐66‐AS. (i) the preparation of the IM‐UiO‐66‐AS. (ii) The proton conductivity of IM‐UiO‐66‐AS. Reproduced with permission [177]. Copyright 2019, the Royal society of chemistry. (e) AHF‐MOFs [178]. (i) The structure of AHF‐PDC and AHF‐BPDC. (ii) The simulation structure of PDC@PH. (iii) The simulation structure of BPDC@PH. (iv) The conductivity of PDC@PH and BPDC@PH. Reproduced with permission [178]. Copyright 2025, Wiley‐VCH. (f) UIO‐66 in UIO‐67. (i) The structure of UIO‐66 in UIO‐67. (ii) The ionic conductivity of UIO‐66 in UIO‐67. Reproduced with permission [148]. Copyright 2021, Wiley‐VCH. (g) FU@MOF‐808‐SO₃H [179]. (i) The structure of FU@MOF‐808‐SO₃H. (ii) The proton conductivity of FU@MOF‐808‐SO₃H. Reproduced with permission [179]. Copyright 2025, Wiley‐VCH. (h) glassy ZIF‐62 [180]. (i) Na⁺ conduction in crystalline ZIF‐62. (ii) Na⁺ conduction in glassy ZIF‐60. Reproduced with permission [180]. Copyright 2025, Elsevier.

Slightly different from the ion number‐regulated ion transport, increasing the functional group numbers or types in N‐MOFs mainly strengthens or enhances the ion hopping sites. For strengthening the ion hopping site, the introduction of strongly electronegative –F into UiO‐66‐(SO₃H)₂ [175] and MIL‐53‐SO₃H‐X [176] enhances the acidity of −SO₃H (Figure 13c), which increases the ordered hydrogen bonding networks, incurring UiO‐66‐F₂(SO₃H)₂ and MIL‐53‐SO₃H‐F to feature high proton conductivity than their non functionalized counterparts. Such as, the proton conductivity of 298.3 mS cm⁻¹ for MIL‐53‐SO₃H‐F at 80°C and 98% RH [176]. The advantage of this strategy is to improve the efficiency of existing proton sources rather than increasing the number of proton sources. It not only avoids the problem of pore crowding caused by high‐density grafting, but also improves the continuity of hydrogen bonding networks. For enhancing the ion hopping site, taking UiO‐66 derivatives as an example [177], when the introduction of ‐NH₂ into UiO‐66 with zero proton sources and one hopping site, UiO‐66‐NH₂ will increase a proton hopping site, resulting UiO‐66‐NH₂ delivers the higher proton conductivity (3 × 10⁻⁶ S cm⁻¹) than UiO‐66 at T = 80°C and RH = 98%, Meanwhile, while the addition of an imidazole functional group with a proton donor and a proton hopping site into ‐NH₂ of UiO‐66‐NH₂, UiO‐66‐NH₂‐IM will produce a proton source and three proton hopping sites, delivering the high conductivity of 1.59 × 10⁻⁴ S cm⁻¹. Furthermore, when the imidazole functional group is introduced into UIO‐66‐NH₂‐SO₃H. M‐UiO‐66‐NH₂‐SO₃H can have both two proton sources and four hopping sites, leading the super high proton conductivity (1.54×10⁻¹ S cm⁻¹) at the same condition (Figure 13d‐i,d‐ii) [177]. These results suggest that the density and chemical properties of proton sources and hopping sites are the determining factors of proton conductivity in N‐MOFs, which are more important than pore structure, specific surface area, and even water absorption.

As one of the unique advantages of N‐MOFs, their porosity significantly influences the ion transport network and nanoconfinement effect in directional channels. However, the relationship between pore size and ionic conductivity may be not monotonic, but depends on the specific conditions. Because larger porosity will both weaken the confinement effect and enhance the ability to accommodate guest molecules. These two effects may either promote or conflict with each other. For instance, the large pore of UiO‐67 [149] can reduce the confinement effect and allow more efficient solvated Li⁺ transport compared with UiO‐66, resulting LPC@UiO‐67 have high ionic conductivity than their counterparts (0.65 mS cm⁻¹ for LPC@UiO‐67). In addition, although AHF‐BPDC (AHF: azole hybrid framework, BPDC: 4, 4‐biphenyl dicarboxylic acid) with large pore weakens the salt confinement effect, it increases the entry of TFSI⁻ into the pore and interact with the pore wall sites as well as accelerate the Na⁺ diffusion (Figure 13e), enabling BPDC@PH electrolyte with high Li⁺ conductivity of 0.74 mS cm⁻¹ [178]. More importantly, MOF‐in‐MOF (UIO‐66 in UIO‐67) structures can fully utilize the hybridized MOFs channels, in which UIO‐67 that have large pore and a high specific surface area, could adsorb a large amount of ILs electrolyte, while UIO‐66 with small pores can restrict the transport of large‐sized ions (Figure 13f‐i). Such special structural design alleviates the slow Li⁺ migration behavior on the one hand, and also avoid the defects of low electrolyte loading on the other hand. In view of this, Li‐IL@UIO‐66@67 delivers the high conductivity of 2.1 × 10⁻³ S cm⁻¹ (Figure 13f‐ii) [148]. Notably, this high conductivity surpasses those of the individual N‐MOFs loaded with the same electrolyte, confirming that pore‐architecture engineering, rather than simply maximizing the pore volume, is the more impactful strategy for optimizing the conductivity.

The flexible pore structure of N‐MOFs allows them to be inserted by some guest nano species, which can alter the physicochemical properties within the nanochannels. Although the experimental operation of this method is simple, the difference in effectiveness can reach several orders of magnitude, and the key lies in the chemical properties of the inserted species [175, 179, 181] (e.g., protein, zwitterion, 5‐fluorouracil and amine molecules, etc). The immobilization of guest 5‐fluorouracil (FU) and in MOF‐808‐SO₃H [179] and zwitterion sulfobetaine methacrylate (SBMA) in ZIF‐8 [181] forms ordered ionic channels and hydrogen bonding network with water to rapidly transport proton (Figure 13g‐i), inducing FU@MOF‐808‐SO₃H and SBMA‐MOF outputs the higher proton conductivity than their counterparts. Such as the proton conductivity of 0.072 S cm⁻¹ for FU@MOF‐808‐SO₃H at T = 353 K and RH = 95% (Figure 13g‐ii). This strategy is one of the largest increases in ionic conductivity reported in the guest immersion method, indicating that the density and continuity of the hydrogen bonding network constructed by the guest are the true determinants of ionic conductivity, not just the presence or absence of proton source groups in N‐MOFs.

Amorphous N‐MOFs, which is obtained by partially or entirely transforming the crystallized N‐MOFs, are a special state of N‐MOFs, which weakens the interaction in directional nanochannels and shortens the ion transport pathway (Figure 13h‐i,h‐ii). Consequently, crystal phase regulation of N‐MOFs enhances the ionic conductivity of N‐MOFs. For example, amorphous ZIF‐62 [180] weakens the interaction between components and guests and provides abundant ion transport channels. Disordering ZIF‐4 [173] shorten the Li⁺ transport pathway by no grain boundaries and isotropy, resulting in NaTFSI/PC@ZIF‐62 and glass ZIF‐4 exhibit the high ionic conductivity up to 1–2 × 10⁻⁴ S cm⁻¹ at 30°C [180]. The ionic conductivity caused by amorphous N‐MOFs is not particularly prominent, but it achieves a qualitative change in the ion conduction mechanism. More importantly, amorphization is a universal processing strategy, which is applicable to almost all N‐MOFs. However, due to the loss of long‐range order in N‐MOFs, the final ionic conductivity achieved by this method is limited in the absence of regular 1D nanochannels.

Among the various regulation strategies for N‐MOFs mentioned earlier (Table 1), the most effective and convenient strategy is to immerse guest nano species with specific properties into N‐MOFs. Therefore, the immersion of guest nano species should be preferred when considering the modification cost and time. On top of that, grafting ions and functional groups into N‐MOFs is the most essential approach to improve their conductivity, and more attempts should be made in this regard in the future. Finally, considering the critical effect of intrinsic characteristics on the ion transport of N‐MOFs, the amorphous state, the matching relationship between pores and ions, and the influence of nanometer dimension and size on ion transport should be systematically studied.

TABLE 1.

Summary of the key regulation strategies for nano ion conductors.

Regulating Strategy	Material Examples	Mechanism	Performance (S cm− 1)	Advantages	Limitations	Refs.
Ion type regulation	LiCON‐2 (−COOH) vs. LiCON‐3 (−SO3H)	Stronger acidity increases dissociation	4.36 × 10−6 3.21 × 10−6	Enhancing carrier concentration	Limited by acid–base dissociation	[37]
	CF3‐Li‐ImCOF vs. CH3‐Li‐ImCOF	Electron‐withdrawing delocalizes anion charge	7.2 × 10−3 8.0 × 10−5	Electronic modulation is powerful	Requiring co‐solvent	[61]
Ion number regulation	LiO3S‐COF1 vs. LiO3S‐COF2	More ‐SO3 − sites for Li+ hopping	2.67 × 10−5 4.46 × 10−5	Scalable improvement	Not significant	[123]
	TpPa‐SO3H vs.TpBd‐SO3H	Optimized ‐SO3H spacing for SHB network	0.54 1.389	Qualitative leap in hydrogen bonding network	Requiring precise spatial control	[41, 117]
Functional group type regulation	UiO‐66, UiO‐66‐NH2 vs. M‐UiO‐66‐NH2‐SO3H	Enhanced proton sources/hopping sites	3 × 10−6 1.59 × 10−4 1.54 × 10−1	Five‐order‐of‐magnitude enhancement	Not significant	[177]
	MIL‐53‐SO3H‐H vs. MIL‐53‐SO3H‐F	−F enhances ‐SO3H acidity	6.65 × 10−2 2.93 × 10−1	Enhancing the electron cloud density	Requiring halogenated monomers	[176]
Guest molecule intercalation	COF vs. PEG@COF LiClO4	Enriched pathways and promoted dissociation	10−8 1.45 × 10−6	Dramatic enhancement Synthetically simple	Not significant	[130]
Porosity/ pore size regulation	UiO‐66 vs. UiO‐67	Larger pores reduce confinement for solvated ions	0.17 0.65	Improved guest loading and mobility	Larger pores may reduce selectivity	[149]
	UiO‐66 and UiO‐67 vs. MOF‐in‐MOF	Hierarchical pores: large for loading, small for sieving	4.2 × 10−4 5.8 × 10−4 2.1 × 10−3	Synergistic enhancement	Complex synthesis	[148]
Amorphization	ZIF‐4 crystal vs. ZIF‐4 glass	Isotropic pathways, weakened interaction	8.21 × 10−5 1.61 × 10−4	Universal strategy; eliminating grain boundaries	Loss of long‐range order	[173]
OMS creation	Pseudo ionic N‐MOFs	OMSs adsorb guest anions	0.5–4 × 10−3	Universal applicability	OMSs susceptible to poisoning	[171, 172]
Flexible framework design	Rigid sulfonated MOFs vs. BUT‐8(Cr)	Maintained H‐bond network at low RH	∼0 (at low RH) 6.32 × 10−3 (65% RH)	Low humidity conductivity	Difficult design	[164]

4. Theoretical Calculations for Nano Ion Conductors

Theoretical calculations have emerged as an indispensable tool for elucidating the microscopic ion transport mechanisms of nano ion conductors. When the dimensions of ion conductors are reduced to the nanometer scale, interfacial, confinement, and quantum size effects become tightly intertwined, which renders the classical bulk ion transport theories inadequate. Concurrently, conventional experimental characterization of nano ion conductors faces critical bottlenecks due to the insufficient spatial resolution and the difficulty of performing in situ observations. Theoretical calculations such as density functional theory (DFT), ab initio molecular dynamics (AIMD), and classical molecular dynamics (MD) can capture the key information at the atomic and molecular scales [120, 122], including ion migration kinetics, ion solvation structures, and electronic properties. Therefore, these approaches serve as a critical bridge between the practical performance of nano ion conductors and their underlying microscopic structures.

4.1. Ion Migration Pathway Calculation

The Arrhenius behavior of ionic conductivity is dominated by the activation energy. Therefore, the migration barrier and local energy landscape, which are calculated by DFT, serve as the primary theoretical indicators for evaluating ion transport pathways and transport performance in nano ion conductors. As a classic example of inorganic nano ion conductors, bulk AgCrS₂ exhibits a calculated Ag⁺ migration barrier of 0.48 eV, which agrees closely with the experimental value (0.47 eV). Upon exfoliation to a monolayer, this barrier drops dramatically to 0.19 eV, demonstrating that the room‐temperature ionic conductivity increases by three orders of magnitude (Figure 14a‐i) [46]. This dimensionality‐driven improvement, which is quantified by DFT, establishes that reducing the layer thickness is a powerful strategy for boosting ion transport. Nevertheless, static barrier calculations alone cannot adequately describe the cooperative ion migration in multi‐ion systems. DFT simulations of Li_xAg_1‐xCrS₂ reveal a more intriguing microscopic phenomenon, large‐radius Ag⁺ acts as interlayer pillars that rigidify the layered AgCrS₂ framework and weaken the binding of S²⁻ to Li⁺, enabling Li⁺ to occupy high‐energy tetrahedral sites (Figure 14a‐ii). The resulting Coulomb repulsion lowers the calculated activation energy to 0.18 eV, significantly enhancing the interlayer Li⁺ transport (Figure 14a‐iii) [45]. This multi‐ion synergistic mechanism, which does not exist in pure AgCrS₂ or LiCrS₂, illustrates how theory can reveal ion transport paradigms that cannot be achieved solely through experiments.

FIGURE 14.

Theoretical calculations for nano ion conductors. (a) AgCrS₂ and Li_xAg_1‐xCrS₂ based nano ion conductors. (i) Migration path of Ag⁺ in AgCrS₂. Reproduced with permission [46]. Copyright 2021, Springer Nature. (ii) DFT calculation of ionic migration in Li_xAg_1‐xCrS₂. Reproduced with permission [45]. Copyright 2020, American Chemical Society. (b) Theoretical calculations for TpPa‐SO₃Li and Li‐CON. (i) Theoretical calculations and elucidation of Li⁺ migration in TpPa‐SO₃Li. Reproduced with permission [122]. Copyright 2019, American Chemical Society. (ii) Energy barriers of Li⁺ diffusion in Li‐CON. Reproduced with permission [182]. Copyright 2020, Elsevier. (c) Theoretical calculations for LDH‐L and LDH‐H and TpPa‐SO₃H. (i) The trajectory map of OH⁻ and H₂O for LDH‐H and LDH‐L. Reproduced with permission [76]. Copyright 2025, American Chemical Society. (ii) The fast proton dissociation calculation for TpPa‐SO₃H. Reproduced with permission [41]. Copyright 2022, Springer Nature. (d) Snapshots from AlMD simulation for Li@Zn‐MOF‐74, Li@HKUST‐1 and Li@MOF‐5. Reproduced with permission [160]. Copyright 2023, Wiley‐VCH. (e) The coordination structures of NaTFSI‐TPDBD and NaTFSI‐TPDBD‐COO⁻. Reproduced with permission [120]. Copyright 2023, Springer Nature.

In N‐MOFs, DFT calculations have elucidated how the anisotropy of ion transport pathways determines macroscopic ionic conductivity. In TpPa‐SO₃Li, the calculated axial migration barrier is only 7.6 kcal·mol⁻¹, whereas the in‐plane barrier reaches 31.6 kcal·mol⁻¹. This pronounced anisotropy originates from the short spacing and relay‐like spatial arrangement of ketonic oxygen atoms along the axial direction (Figure 14b‐i) [122], establishing that the arrangement of ion hopping sites along the pore axis is a universal design principle.

In N‐COFs, theoretical calculations on the Li‐CON system further reveal a lithiation effect. The migration barrier of CON decreases drastically from 14.8 to 0.53 eV after lithiation (Figure 14b‐ii) [182], demonstrating that the introduction of charge carriers can reconstruct the energy landscape. These theoretical studies may offer a fundamentally new activation strategy for nano ion conductors.

4.2. Ion Migration Mechanism Simulation

AIMD and MD simulations can identify the dynamic ion transport mechanisms in nano ion conductors by tracking the atomic trajectories or analyzing the mean‐square displacement (MSD). In 2D ion conductors, AIMD simulations of the confined interlayer space in LDHs reveal that the OH⁻ undergoes long‐range transport through the proton reverse hopping migration (Grotthuss mechanism) within the hydrogen bonding network. This mechanism has been confirmed through the simulation snapshots and MSD analysis [76]. More intriguingly, simulations show that LDH‐L with smaller interlayer spacing exhibits faster in‐plane transport, owing to the combination of confinement and z‐axis/in‐plane motion velocities (Figure 14c‐i) [76]. This theoretical finding overturns the assumption that large interlayer spacing always favors ion transport, and emphasizes that the hydrogen bonding network structure is the determining factor. MD simulations on 2D BiOI further quantify the relative contributions of the Grotthuss and vehicular mechanisms with increasing OH⁻ substitution, the contribution from the Grotthuss mechanism progressively increases, while the vehicular contribution remains nearly unchanged [102]. This calculation result directly explains the rapid increase in conductivity observed in the experiment, and determines that the expansion of the continuous hydrogen bonding network is the mechanism of the enhanced OH⁻ transport in these systems. AIMD simulations on N‐MOF systems disclose an even more intricate phenomenon. In Zn‐MOF‐74, simulations indicate that Li⁺ is co‐solvated by TFSI⁻ and DME and follows a vehicular mechanism [160], providing direction for regulating the guest electrolyte.

In proton transport systems, theoretical simulations demonstrate a richer mechanistic landscape. AIMD simulations of TpPa‐SO₃H reveal that the electrostatic attraction exerted by ‐SO₃ ⁻ induces the dissociated protons to undergo rapid hopping rather than transport within the confined hydrogen bonding network (Figure 14c‐ii) [41]. When the distance between hydrophilic groups is below 1.0 nm, simulations show that the confined water regions overlap, leading to proton delocalization, which strongly supports the Grotthuss mechanism. MD simulations of LiCl@UiO‐66‐F₂(SO₃H)₂ demonstrate a humidity‐driven mechanistic transition. This indicates that immobilized hydrated Li⁺ forms a charged layer, then the charge repulsion simultaneously reduces the kinetic barriers for both the ion hopping and vehicular mechanisms [175].

4.3. Ion Dissociation and Solvation Calculation

Fast ion transport in nano ion conductors requires not only low migration barriers but also the generation of sufficient charge carriers. Consequently, DFT and MD simulations have played a pivotal role in quantifying host‐guest interactions and elucidating the driving forces for salt dissociation. In N‐MOFs, DFT calculations show a quantitative ranking of the binding energies: −1.31 eV for TFSI⁻/Zn‐MOF‐74 with OMSs, −1.16 eV for TFSI⁻/HKUST‐1 with OMSs, and −0.83 eV for TFSI⁻/MOF‐5 [160] (Figure 14d), revealing that the enhanced anion anchoring and promoted lithium salt dissociation originate from the OMSs. AIMD simulations further disclose the immobilizing effect of OMSs on TFSI⁻. These calculations establish that OMSs are the structural determinant of anion anchoring and salt dissociation in N‐MOFs, providing a direct guideline for optimizing carrier concentration through metal site engineering.

In N‐COF systems, theoretical calculations reveal an even more exquisite design landscape for regulating ion solvation. DFT calculations on ziCOFNs disclose a dissociation/enhancement bifunctional mechanism. Specifically, the ‐COO⁻ group binds Li⁺ with a strong adsorption energy of ‐5.56 eV, surpassing all carbonate solvents. MD simulations confirm that the Li⁺ solvation sheath in ziCOFNs is the loosest and exhibits the lowest coordination number [140]. This indicates that the zwitterionic framework simultaneously achieves Li⁺ enrichment and anion immobilization. MD simulations of TPDBD‐CNa show that the negatively charged ‐COO⁻ pore walls enrich Na⁺ and repel TFSI⁻, forming a bioinspired sub‐nanometer transport zone (Figure 14e). In addition, the calculated diffusion coefficient increases from 0.085 to 0.239 × 10⁸ Ų·ns⁻¹, quantitatively confirming the critical role of pore‐wall charge modification in constructing fast ion transport pathways. These results provide a theoretical foundation for the design of single‐ion conducting COFs [120]. Furthermore, the ion‐sieving effect within the nanochannels of N‐MOFs has been quantified through theoretical calculations. In HKUST‐1, DFT calculations show that the migration barrier of TFSI⁻ crossing the pore aperture is highly anisotropic (0.05 eV along the longitudinal direction vs. 0.68 eV along the transverse direction). DFT‐MD further confirms that the confinement of N‐MOFs renders TFSI⁻ mobility far lower than that of Li⁺ [144], highlighting the necessity of optimizing pore size from a theoretical perspective.

Collectively, the theoretical studies summarized above yield the following design principles. First, reshape the energy landscape of 2D inorganic ion conductors through dimension reduction, structural disorder, or ion doping to construct continuous and interconnected ion transport pathways. Second, exploit electron‐withdrawing effects to modulate the ion‐framework binding strength. Third, utilize pore/nanochannel confinement and charged interfaces to realize ion sieving and selective transport within nano ion conductors. Looking ahead, the deep integration of DFT screening and machine learning will accelerate the discovery and optimization of novel nano ion conductors, while the integrated development of multiscale simulation methods has the potential to bridge the gap between atomic‐scale mechanistic understanding and macroscopic device performance.

5. The Major Challenges of Current Nano Ion Conductors

Despite the significant advances achieved in the design of nano ion conductors (Table 2), several fundamental and technological challenges must be addressed before their practical applications.

TABLE 2.

Comparative analysis of the ionic conductivity for representative nano ion conductors.

Material class	Representative material	Ion type	Ionic Conductivity (S cm− 1)	Measurement conditions	Conduction mechanism	Potential applications	Refs.
2D Inorganic	AgCrS2	Ag+	33.2×10−3	RT anhydrous	Multi‐ion effect	Batteries	[46]
	LixAg1‐xCrS2	Li+	19.6×10−3	RT	Multi‐ion effect		[45]
	BiOI1‐x(OH)x	OH−	168×10−3	90°C	Grotthuss	Fuel cells	[102]
	H3Sb3P2O14	H+	1.02	90°C 100% RH	Grotthuss	Batteries Fuel cells	[97]
	HSbP2O8	H+	1.18		Grotthuss		[97]
	Cd0.85PS3Li0.15H0.15	H+	0.95	90°C 98% RH	Grotthuss		[40]
	TiO2	Li+	0.62	90°C 100% RH	Water‐ mediated	Batteries	[75]
	LDH (Co–Al)	OH−	1×10−1	RT	Grotthuss	Fuel cells	[76]
Anionic COFs	TpPa‐SO3H	H+	0.54	80°C pure water	Grotthuss	Batteries Fuel cells	[117]
	TpBd‐SO3H	H+	1.389	90°C 100% RH	Grotthuss (SHB)		[41]
	IPC‐COF	H+	0.38	80°C	Grotthuss (SHB)		[116]
	TpPa‐SO3H	H+	0.167	30°C 98% RH	Grotthuss		[115]
	SDT‐COF	H+	0.095	98% RH	Grotthuss		[118]
	TpMbh‐PO3Li2	Li+	1.7 × 10−3	RT	Grotthuss	Batteries	[119]
	ziCOFNs@LE	Li+	1.5 × 10−4	35°C	Lewis acid–base		[140]
	LiO3S‐COF2	Li+	4.46 × 10−5	RT	Solvent‐assisted		[123]
	PC@Li‐ImCOF	Li+	5.3 × 10−3	RT	Lewis acid–base		[61]
Cationic COFs	QACOFMs (@H3PO4)	H+	0.380	200°C, anhydrous	Anion‐decoupled	Batteries Fuel cells	[127]
	COF‐3OH	OH−	>0.25	80°C 100% RH	Grotthuss	Fuel cells	[126]
	COF‐SDQA	OH−	0.329	80°C 100% RH	Grotthuss		[137]
	Li‐TpTGTFSI‐COF	Li+	2.09 × 10−4	70°C	Anion‐trapping	Batteries	[129]
Non‐ionic COFs	CD‐TpAzo@H3PO4	H+	0.78	150°C anhydrous	Guest‐confinement	Batteries Fuel cells	[135]
	H3PO4@TPB‐DMeTP‐COF	H+	∼10−1	High T anhydrous	Guest‐confinement		[132]
	PA@PyTTA‐BMTP‐COF	H+	0.023	140°C	Guest‐confinement	Batteries Fuel cells	[133]
	COF‐H (pristine)	H+	8×10−8	80°C 85% RH	—	Batteries Fuel cells	[53]
Anionic MOFs	H2SO4@MIL‐101‐SO3H	H+	1.82	70°C 90% RH	Grotthuss+Vehicle	Batteries Fuel cells	[146]
	MIL‐53‐SO3H‐F	H+	0.298	80°C 98% RH	Grotthuss		[176]
	BUT‐8(Cr)	H+	0.127	80°C 100% RH	Flexible‐framework		[164]
	BUT‐8(Cr)	H+	6.32 × 10−3	80°C 65% RH	Flexible‐framework		[164]
	Cr‐MIL‐88B‐PSA	H+	0.158	100°C 85% RH	Grotthuss+Vehicle		[145]
	UiO‐66‐(SO3H)2	H+	0.084	80°C 90% RH	Grotthuss		[177]
	M‐UiO‐66‐NH2‐SO3H	H+	0.154	80°C 98% RH	Grotthuss		[177]
	FU@MOF‐808‐SO3H	H+	0.072	353 K 95% RH	Additional H‐bond		[179]
	MIL‐101(Cr)‐NH2‐SO3H	H+	∼10−1	70–100°C 85%–98% RH	Grotthuss+Vehicle		[146]
Cationic MOFs	D‐UiO‐66‐NH2+PEG	Li+	1.6 × 10−4	80°C	Anion‐trapping	Batteries	[165]
Non‐ionic MOFs	EIMS‐HTFSA @MIL‐101	H+	2 × 10−4	140°C anhydrous	acid‐confinement	Batteries Fuel cells	[153]
	Li‐PC@ CuBTC‐PSS	Li+	4.47 × 10−3	25°C	OMS coordination	Batteries	[62]
	LPC@UiO‐67	Li+	6.5 × 10−4	RT	Porous enrichment		[149]
	LiTFSI@UiO‐66‐D4	Li+	1.42 × 10−3	∼30°C	Defect and coordination		[155]
	ILs+Na salts @MIL‐101	Na+	1.32 × 10−2	50°C	Pore effect		[63]
	Li‐IL@UIO‐66@67	Li+	2.1 × 10−3	RT	MOF‐in‐MOF		[148]
Pseudo ionic N‐MOFs	UiO‐66@Li TFSI/DME‐IL	Li+	0.5–4 × 10−3	RT	OMS anion‐trapping	Batteries	[171]
	HKUST‐1‐ClO4 −	Li+	0.5–4 × 10−3	RT	OMS ‐trapping		[149]
	PEO@ZIF‐67@LiTFSI	Li+	0.5–4 × 10−3	RT	OMS ‐trapping		[172]
Amorphous MOFs	G‐ZIF‐4	Na+/Li+	1–2 × 10−4	30°C	Isotropic transport	Batteries	[173]
Hybrid composite	SPEEK/MOF‐808	H+	0.077	High T high RH	Polymer synergy	Batteries Fuel cells	[151]
	PIL0.5 @mTpPa‐SO3H	H+	0.102	90°C 100% RH	hydration layer		[141]

Resolving ion transport mechanisms through advanced operando characterization. The current mainstream research on the proton transport mechanisms still relies on empirical calculation formulas or some simple characterization techniques. Few reports have described the specific arrangement and structure of hydrogen bonding networks in nanopores. Similarly, for the transport of alkali metal ions, only some simple characterization techniques are used to study their ion transport state within the nanochannels, and then the proposed ion transport mechanism is verified through theoretical calculations. Therefore, it is vital to develop and utilize advanced characterization techniques to study the specific ion transport process within the nanopores.

Expanding the operational temperature window toward extreme environments. The developed nano ion conductors have achieved desirable ion transport performance in the temperature range of ‐20‐100°C, but they cannot operate under extremely harsh environments, which conflicts with the application of a large number of devices in extremely high or low temperatures. The core reason is that ion transport strongly depends on the phase behavior of the medium, and high‐temperature dehydration or low‐temperature solidification of the medium all can lead to a sudden drop in ionic conductivity. Future research needs to shift toward alternative media and design flexible frameworks to adapt to thermal expansion and contraction without damaging ion conduction pathways.

Diversifying ion conduction beyond H⁺ and Li⁺. The developed nano ion conductors are mainly used to achieve the rapid conduction of H⁺ and Li⁺, but in this era of charged devices as daily necessities, the demand for various ion conduction‐based electrical devices is also increasing. However, the design principles established for H⁺ and Li⁺ cannot be directly applied to multivalent ions or monovalent ions with larger radius due to the strong coulomb interactions, large ionic radius, and high steric hindrance.

Overcoming the aqueous exfoliation bottleneck for 2D nano ion conductors. Among nano ion conductors, the exfoliation of 2D nano ion conductors is mainly achieved in aqueous systems, and they are prone to aggregation in non‐aqueous systems, which significantly limits their application fields. Hence, it is necessary to develop non‐aqueous exfoliating technology and an interlayer expansion process, which is a key frontier for significantly expanding the application of 2D nano ion conductors.

Designing ion conductors with both nanoshape and ion conduction function. Simultaneously achieving controllable nano morphology and high ionic conductivity in the same material is extremely challenging. Ion conductors with high performance are mostly micron sized grains with high interparticle resistance, reducing the size will damage the crystallinity. Therefore, the essential contradiction of this challenge lies in the fact that enhancing ion conduction disrupts morphology control, while strengthening nanostructures damages ion conduction function.

Understanding the ion distribution at heterogeneous interfaces. In practical devices, the ionic conductivity of nano ion conductors is often limited by the interfacial resistance (such as electrode, separator, and binder interfaces) rather than the bulk properties of nano ion conductors. At present, the understanding of the ion distribution, ion morphology, and ion dynamics remains lack on the interface. In addition, the ion concentration, solvation, and migration rate at the interface differ significantly from those of nanopores/nanochannels due to space charge, specific adsorption, and dielectric discontinuity. Therefore, if the interfacial ion distribution is designed properly, it can promote charge transfer, otherwise it will become a limiting step of ion transport.

6. Perspective

Nanomaterial‐based ion conductors exhibit a combination of unique performance: high conductivity, good interfacial contact, and high transference number. These advantages compensate for the significant shortcomings faced by traditional SSICs when they are used, namely, achieving fast ion transport without thoroughly comprising other properties of the corresponding bulk ion conductor. The research focus of nano ion conductors spans from pure organic/inorganic materials to organic‐inorganic hybrid materials based on the diversity of constituent elements. It is certain that the glamour of nano ion conductors will become more attractive because their unique fast ion transport makes the application of ion conductors more widely. Looking to the future, many aspects still require optimization, and the mechanisms need to be clearly elucidated before the large‐scale application of nano ion conductors (Figure 15).

1. Clarifying the specific ion transport mechanisms within directional channels and the distribution of ions at the interface. Although many characterizations and calculations have provided preliminary insight into the ion transport in directional channels and the ion distribution at the surface, they struggle to elucidate the specific existence mechanism of ion transport. In this scenario, remarkable in situ characterization and simulation techniques should be used to obtain the precise ion transport mechanism in nanochannels and/or at the surface. For example, in situ NMR can determine the distribution of ions inside and outside the pores in different charged states, in situ infrared (IR) and in situ impedance (EIS) can analyze the surface distribution of ions in different charged states, in situ AFM is able to monitor the volume expansion rate of nanopores to infer the change of ions in the pores. In addition, finite element simulation should be developed to simulate the distribution of ions in pores and at the surfaces.

2. Achieving ultra‐high ionic conductivity at high and low temperatures. Due to the unique ion transport effect in nanochannels, nano ion conductors can work well between −20°C and 100°C, which limits their work in extreme environments, such as the operating temperature above 200°C and below −50°C. Considering that eutectic electrolytes have good ion transport performance at low temperatures, improving the low‐temperature performance for nano ion conductors should start from the combination of eutectic electrolytes. In addition, ILs feature high thermal stability, and they can avoid adverse effects such as thermal volatilization at high temperatures. Therefore, perhaps the realization of nano ion conductors with good high‐temperature properties should start from the combination of ILs.

3. Developing nano ion conductors that conduct other alkali metal ions. The current conduction of ions for nano ion conductors is mainly concentrated on H⁺ and Li⁺, and there are few studies about the conduction of other alkali metal ions. In theory, Li⁺ in the nano ion conductors can be replaced with other alkali metal ions through ion exchange. However, in reality, alkali metal ions (e.g., Na⁺ and K⁺) have larger radius and may be trapped by small pores, while large pores can weaken the confinement effect. Meanwhile, their weak acidity affects ion dissociation and ion solvation. Therefore, Computational screening should combine with experimentation to systematically study the matching of pore size/nanochannel, charge density, and ion characteristics, enabling the rapidly development of nano ion conductors.

4. Enriching the types of nano ion conductors through predictive discovery. Compared to bulk ion conductors, there are still fewer types of nano ion conductors, which is mainly restricted by the difficulty of synthesizing nanomaterials with ion conduction as well as the limited monomers. Therefore, the breakthrough path should shift from experience trial and error method to predictive discovery. In addition to classical DFT and MD calculations, three computational techniques are crucial. First, the interatomic potential based on machine learning can enable nanosecond‐to‐microsecond simulations with near DFT accuracy, which captures the rare events that traditional DFT based dynamics cannot achieve, such as proton hopping and cation desolvation. Second, the enhanced sampling methods systematically map a multidimensional free energy landscape, which can reveal the competitive pathways and metastable intermediates of ion transport beyond the minimum energy pathway. Third, high‐throughput screening, which crosses MOF, COF, and 2D material databases, can determine the optimal pore size, functional group density, and framework flexibility. Then, integrating these methods into a unified material gene library, in which high‐throughput screening identifies the candidate ion conductors, enhanced sampling maps energy landscapes, and machine learning of interatomic potential function validates the dynamics under realistic conditions. Therefore, the constructed material gene library will greatly accelerate the discovery of conductors with different ion transport.

5. Developing non‐aqueous exfoliating technology. 2D nanosheet based ion conductors currently occupy a significant share in nano ion conductors. However, the exfoliation of 2D materials is mainly performed in an aqueous solution, which limits their ion conduction in non‐aqueous environments. Therefore, future research should attempt to use some strong polar solvents (e.g., NMP, EA, and EtOH) to exfoliate 2D ion transport materials on the one hand, and try to exfoliate them with a mixture of strong polar solvents and water on the other hand. In addition, attention should be paid to the issues, such as residual coordination of high boiling point solvents, and solvents need to be screened through simulation and experiments. Furthermore, the long‑term dispersion stability of nanosheets in organic electrolytes must be addressed, which will promote the application of 2D materials in non‐aqueous devices.

6. Emphasizing the mass transport in nanochannels. The mass transport in nanochannels plays a key role in the ion transport of nano ion conductors. Specifically, it affects the concentration of internal ionophores and determines the formation of the ion transport pathway. However, few studies have focused on this aspect so far. Given that EQCM technology is an effective tool to research the microscopic mass change in materials. its application to nano ion conductors will certainly clarify the mass transport of ions in nanochannels.

7. Thoroughly studying the relationship between the internal structure and ion conduction of nano ion conductors. Current studies have already explained that the crystalline state of nanomaterials, the size of nanochannels, and the functional groups in nanochannels present significant effects on ion transport. However, few studies have elucidated the specific relationship between these structures and ion conduction performance. Therefore, when enriching the types of nano ion conductors, more research should utilize characterization and testing techniques to explore the specific influence of material structure on ion conduction.

8. The design concept of nano ion conductors can be applied to other ion conduction systems (e.g., aqueous electrolyte, organic electrolyte, and hybrid electrolyte) and material design (e.g., the construction of fast ion transport channels within electrode materials). In addition, while improving the ionic conductivity of nano ion conductors, it is necessary to study the influence of the change of conductivity on other physicochemical properties of nano ion conductors, which is very critical to the comprehensive performance of actual devices.

FIGURE 15.

Illustration of strategies to optimize nano ion conductors.

Overall, although nano ion conductors address the most pressing problem of low conductivity, their development and utilization still have a long way to go. We believe that, starting from the current composition, the advanced design of nano ion conductors will provide shimmers for the new generation of fast ion conductors, which promotes the broad application of ion conductors in ion electronics, especially in fuel cells and energy storage devices as well as transistors and diodes.

Conflicts of Interest

The authors declare no conflicts of interest.

Data Availability Statement

The data that support the findings of this study are available from the corresponding author upon reasonable request.