Small components play a big role - fillers in composite solid-state electrolytes for lithium metal batteries

Construction of lithium-ion channels of CSSEs

CSSEs are often formed by including additional inorganic fillers into the PSEs. These fillers may be split into inert fillers that insulate Li ions and active fillers that can conduct Li ions. This chapter will describe the role of each component in CSSEs and the general lithium-ion transport channels.

Lithium-ion channels in polymer electrolytes

The polymer matrix is a crucial component in CSSEs. The lithium-ion transport channel of the polymer chain should first be fully understood to facilitate a better understanding of how CSSEs work. A multitude of polar groups on the polymer chain can coordinate with lithium, facilitating the adsorption of Li ions onto the polymer chain. When the polymer chain undergoes linkage movement, Li ions will migrate from the coordination site where it was initially located to other adjacent coordination sites^[47]. Furthermore, when the temperature of the polymer matrix attains its glass transition temperature (T_g), it facilitates the movement of chain segments within the polymer, hence accelerating lithium-ion migration, which is the predominant mechanism of lithium-ion transport in polymers^[48]. The presence of space between the polymer chains facilitates the migration of Li ions, allowing them to hop across different segments of the chains. The interplay of short-range migration and long-range motion of the polymer chains facilitates the transport of Li ions in the PSEs^[49,50], as shown in Figure 2A.

Figure 2. Schematic illustrations of (A) General transport channel of Li ions in polymers^[49,50]. Reproduced from Ref.^[49] with permission. Copyright 2007, Purpose-LED Publishing. Reproduced from Ref.^[50] with permission. Copyright 2023 Wiley-VCH. (B) Fillers construct lithium-ion transport channels^[51]. Reproduced from Ref.^[51] with permission. Copyright 2023 Wiley-VCH. (C) Inert fillers reduce polymer crystallinity^[52]. Reproduced from Ref.^[52] with permission. Copyright 2017, Springer Link. (D) Lewis acid-base interaction between filler and polymer enhances cation selectivity. (E) The cation transport pathway is constructed by the Lewis acid-base interaction between filler and polymer^[53]. Reproduced from Ref.^[53] with permission. Copyright 2024, Wiley-VCH.

Lithium-ion channels in CSSEs

By adding fillers to the polymer, more lithium-ion transport channels can be constructed^[51], as shown in Figure 2B. Fillers can generally be categorized into lithium-insulated inert fillers and active fillers capable of conducting Li ions, mainly through the following ways to build lithium-ion transmission channels. The inert fillers themselves insulate lithium ions and do not directly facilitate their movement. They function as solid plasticizers inside the polymer matrix, diminishing the polymer’s crystallinity. As shown in Figure 2C, incorporating inert fillers into the polymer matrix enhances the amorphous regions of the polymer, hence augmenting lithium-ion transport channels^[52]. This behavior can be elucidated by Lewis acid-base theory, in which ether-oxygen bonds in poly(ethylene oxide) (PEO) can provide lone pair electrons. Conversely, numerous inert fillers, including metal oxides, can offer Lewis acid sites to receive lone pair electrons in PEO. The Lewis acid-base interaction facilitated coordination between the filler and the ether-oxygen bond in PEO, hence restricting the arrangement of the PEO chain segments and resulting in reduced crystallinity of PEO. In addition, the Lewis acid-base interaction between the filler and the polymer can provide fast lithium-ion transport channels, as shown in Figure 2D and E; surface treatment of inert fillers with Lewis acid surface groups can facilitate the dissociation of lithium salts, hence enhancing ionic conductivity^[53]. The inert filler does not change the transport mechanism of Li ions within the polymer, and the polymer matrix is still capable of transporting Li ions. Therefore, the concentration of the inert fillers should not be too high; otherwise, when the inert filler gathers in the polymer matrix, it will produce a continuous lithium insulation part, which limits the transport of Li ions and reduces the electrochemical properties of the CSSE^[54-56].

Unlike inert fillers, active fillers have a unique crystal structure, which makes active fillers Li ions conductors. Lithium-ion transport in active fillers mainly depends on defects in the crystal structure, among which point defects play the most crucial role in ionic transport. According to the point defects, the lithium-ion transport mechanism of active fillers can be classified into two kinds: vacancy and non-vacancy mechanisms. The defect mechanism is further classified into simple vacancy mechanism and vacancy mechanism, and the non-defect mechanism includes an interstitial mechanism and interstitial replacement mechanism^[57,58]. Various lithium-ion transport mechanisms are shown in Figure 3A.

Figure 3. (A) Schematic of Li ions in active fillers^[65]. Reproduced from Ref.^[65] with permission. Copyright 2020, Royal Society of Chemistry. (B) Interaction of active fillers with polymers and formation of space charge regions^[27]. Reproduced from Ref.^[27] with permission. Copyright 2023, Wiley-VCH. (C) Lithium-ion transport channels in polymers filled with different concentrations of active fillers^[66]. Reproduced from Ref.^[66] with permission. Copyright 2017, Elsevier.

The vacancy mechanism involves the diffusion of Li ions from their initial lattice site to an adjacent vacancy. In this process, the atomic jumps produce almost no lattice strain, so the activation energy barrier for the jump is shallow^[59]. The lithium-ion transport in the vacancy mechanism depends on the concentration of vacancies in the lattice. In addition, the class of particles in the crystal or the configuration of the doped cations affects the vacancy size to varying degrees, thus influencing the lithium-ion transport in the crystal.

Among the non-vacancy mechanisms, the interstitial mechanism is dominant and divided into direct interstitial diffusion and interstitial knockout diffusion. Direct interstitial diffusion refers to the diffusion of Li ions in the lattice interstitials to the neighboring interstitial sites, which requires the matrix atoms to be much larger than the Li ions in the interstitials, and the migration process creates a considerable lattice strain. In contrast, under the gap knockout diffusion mechanism, the Li atoms in the gap attack the matrix atoms and move into their original positions, while the displaced matrix atoms migrate to nearby gap sites. This mechanism requires that the size of the matrix atoms is not too different from the lithium-ion size and is a standard mode of Li ions diffusion in active fillers. It is divided into direct exchange and ring exchange^[60]. The former is a simultaneous movement of two atoms to complete the swap in the lattice position. In comparison, the latter is a ring position exchange by three or more atoms to migrate the atoms to the new lattice position. Compared with the vacancy mechanism, the migration barrier energy of the non-vacancy mechanism is somewhat higher.

When the active fillers are filled into the polymer matrix, Li ions are not transported only in the crystal structure of the active fillers but prefer to be transported at the polymer/active filler interface. As shown in Figure 3B, the incorporation of active fillers induces a chemical reaction between the polymer matrix and the fillers, resulting in space charge zones at the interface of the polymer and fillers^[61]. This results in the aggregation of Li ions on one side of the contact. The interconnection of space charge areas generated by each nanoparticle facilitates a continuous and quick transport channel for Li ions, resulting in a substantial enhancement of ionic conductivity. Moreover, as the quantity of active fillers increased, the lithium-ion transport channel was altered accordingly^[62]. Figure 3C shows the lithium-ion transport channels in ceramics in polymers, intermediates, and polymers in ceramics, respectively. In the presence of a limited amount of active fillers, the predominant mechanism for the transport of Li ions occurs via the polymer matrix. When the quantity of active fillers is augmented, there is a tendency for Li ions to be conveyed at the interface between the polymer and active fillers, as well as within the crystals of the active fillers^[63-66]. When the active fillers are used as the matrix phase, Li ions tend to be transported mainly in the crystal structure. Even though active fillers can provide more transport channels than inert fillers, the filling concentration still needs to be designed appropriately, and the nanoparticle agglomeration problem still reduces the ionic conductivity of CSSEs.

Key parameters in CSSEs

To better understand the performance of CSSEs, this section focuses on some critical parameters in the performance requirements of CSSEs, including ionic conductivity, lithium-ion transference number, and mechanical properties.

Metal oxide fillers

Among inert fillers, the most widely researched ones are oxide ceramics, such as Al₂O₃, SiO₂, etc. Although the mechanical properties of the electrolyte have markedly improved, there is no noticeable improvement in conductivity. The movement of Li ions is mostly influenced by the interface between the inert filler and the polymer. Consequently, the surface modification of the inert filler is essential.

Wang et al. used mesoporous SiO₂ with an ultra-high specific surface area and absorbed ethylene carbonate (EC)/propylene carbonate (PC) plasticizer in its nanopores to make SiO₂ as ionic active fillers^[115]. The electrolyte film was prepared by adding the above fillers to PEO/LiClO₄ at a ratio of 10 wt.%, and the ionic conductivity of the electrolyte could be enhanced to 1.5 × 10^-4 S cm^-1. The observed phenomenon can be attributed to the Lewis acid-base interactions between the specific surface state of active SiO₂ and the polymer chains, as well as the anions of lithium salts. Additionally, the presence of the EC/PC plasticizer in the nanopores acts synergistically to enhance the dissociation of the LiClO₄ and facilitate the rapid ionization of the Li ions, thereby promoting efficient lithium-ion transport. Wang et al. designed SiO₂ with a hydroxyl-rich surface, which was prepared in situ using hydrolysis of tetraethyl orthosilicate (TEOS) in an alkaline solution and dispersed in PEO/LiTFSI/succinonitrile (SN)^[116]. The prepared electrolyte connects PEO and SN by forming hydrogen bonds between the hydroxyl groups on the surface of SiO₂ and C, N, and F atoms. On the one hand, the addition of SiO₂ nanoparticles reduces the crystallinity of PEO, while on the other hand, it enhances the mechanical strength and the ionic conductivity of the electrolytes. Tang et al. utilized poly(propylene oxide)-poly(ethylene oxide)-poly(propylene oxide) triblock main chains and surface-modified SiO₂ nanoparticles to create a unique cross-linked nanocomposite polymer electrolyte (CNPE)^[117]. Incorporating 10 wt.% SiO₂ resulted in a room temperature conductivity of 1.32 × 10^-3 S cm^-1 and excellent mechanical stability (elongation of 700%). The enhancement is ascribed to the diminished crystallinity of PEO resulting from the incorporation of SiO₂ and the augmented hydrogen bonding facilitated by ethyl carbamate groups in CNPE, leading to improved particle conductivity and mechanical properties. Furthermore, filler surface treatment improves compatibility with the polymer matrix, creating additional migration paths for Li ions while reducing polymer matrix crystallinity and enhancing adhesion for superior mechanical properties in CSSEs.

While the incorporation of fillers can improve the characteristics of the polymer matrix, there is a critical limit for the amount of fillers utilized. It has been shown that when the inert filler content is below the threshold, lithium-ion transport is facilitated through interaction with the polymer and lithium salt. However, when the content is higher than the critical value, the fillers are subject to agglomeration problems. The agglomeration problem diminishes the effect of inert fillers on reducing polymer crystallization, and Li ions are blocked at the agglomerates as they are transported through the electrolyte^[118]. More seriously, the aggregation of inert filler particles at the anode interface leads to rapid deposition of lithium dendrites, significantly reducing the life of electrolytes. Therefore, it is crucial to reduce the concentration of the filler to obtain a homogeneous distribution inside the polymer matrix. Multiple scholarly investigations have proposed possible remedies to tackle the problem of agglomeration.

Zhan et al. prepared porous SiO₂ particles (10 wt.%) with vinyl groups, which were interconnected with the cross-linking agent polyethylene glycol diacrylate (PEGDA) to achieve an interconnected network structure in the PEO matrix, which could achieve a uniform distribution of fillers^[119]. As shown in Figure 4A and B, the uniformly distributed filler can better inhibit the crystalline behavior of PEO. It can establish chemical linkages inside the polymer matrix, therefore enhancing the compatibility at the interface between the polymer and the inorganic filler. Bao et al. deposited ZnO quantum dots into the PEO matrix using the vapor-phase infiltration (VPI) method, which can achieve the average dispersion of ZnO quantum dots in PEO, As shown in Figure 4C^[120]. The crystalline behavior of PEO is inhibited by the strong interaction of ZnO with the PEO polymer chains. In addition, certain ZnO quantum dots persist on the electrolyte surface, creating a Li-Zn alloy with the lithium anode, therefore markedly decreasing the interfacial resistance.

Figure 4. (A)Schematic illustration of uniform dispersion of p-V-SiO₂ in PEO. (B) SEM image of p-V-SiO₂^[119]. Reproduced from Ref.^[119] with permission. Copyright 2021, Licensee MDPI. (C) Schematic illustration of the VPI process^[120]. Reproduced from Ref.^[120] with permission. Copyright 2021, Elsevier. (D) BTO NF films modulate lithium dendrite growth at copper electrodes^[121]. Reproduced from Ref.^[121] with permission. Copyright 2021, American Chemical Society.

It is not difficult to see that the role of metal oxide fillers in CSSEs is very limited, and most of them need to be pre-treated or rely on advanced preparation technology to achieve the improvement of CSSE performance. Therefore, there is relatively little research on inert fillers, and materials with special functions can provide more performance improvements for CSSEs.

Ferroelectric materials fillers

In addition to the studies of some joint oxide fillers mentioned above, many scholars apply ferroelectric materials to electrolytes^[121]. Although ferroelectric materials are mainly applied as catalysts in electrocatalysis in previous studies, some scholars have found that, in addition to lowering the crystallinity of polymers, Ferroelectric materials can impede the proliferation of lithium dendrites through their spontaneous polarization phenomena^[122]. Taking the most traditional BaTiO₃ (BTO) as an example, Ti⁴⁺ with a small radius is located in the center of an octahedron consisting of six O^2-, while O^2- is squeezed in the middle of Ba²⁺ with a larger radius. When the BTO is filled into the electrolyte, the electrostatic field inside the battery causes the centers of the O^2- and Ti⁴⁺ in the batteries of the BTO to no longer coincide, which causes a spontaneous polarization from negative to positive charge^[121]. This polarization phenomenon draws Li ions from the electrolyte, facilitating accelerated transit of Li ions within the electrolyte and diminishing the concentration gradient of Li ions at the anode. In addition, the growth of lithium dendrites on the anode causes the BTO in the electrolyte to deform by accumulating stresses at the lithium dendrites. A study connecting BTO ceramic nanofibers (BTO NFs) with Cu collectors to detect the growth of lithium dendrites, as shown in Figure 4D. Positive induced charges are generated on the upper side of the deformation position to inhibit the growth of lithium dendrites at the center of deformation. In contrast, negatively generated charges on the bottom side attract lithium ions, enabling their transit perpendicular to the development direction of lithium dendrites, hence promoting uniform lithium deposition.

Singh and Chandra^[123] reported Ba_0.7Sr_0.3TiO₃ and Ba_0.88Sr_0.12TiO₃ materials; these materials significantly enhance the electrolyte conductivity by a factor of ten when 3 wt.% of them are included into PEO. Affected by the distribution of the local electric field inside the polymer matrix, the dielectric constant of the ferroelectric material affects the ion transport mode in the polymer electrolyte. The elevated dielectric constant of the ferroelectric material, in contrast to the polymer matrix, leads to an intensified electric field at the interface between the polymer and the ferroelectric material. Firstly, this phenomenon results in the facilitation of ion movement towards the interface between the polymer and filler. Secondly, as the lithium salt approaches the polymer-filler interface, its dissociation becomes more pronounced, leading to a higher concentration of Li ions in the electrolyte and thereby improving the ionic conductivity.

Carbon materials fillers

In previous studies, carbon nanotubes (CNTs) have generally been used as anode materials, and several investigations have been undertaken to include them into polymer electrolyte matrices owing to their superior mechanical strength and elevated aspect ratio^[124,125]. Li ions can be transported rapidly along the CNT/substrate interface, and the elevated electron cloud density of CNTs can significantly interact with Li ions, facilitating the dissociation of lithium salts. However, the electronic conductivity of CNTs is prone to short-circuiting in batteries, which limits their application as polymer electrolyte fillers^[126].

Tang et al. grew one-dimensional CNTs in the middle of two layers of montmorillonite (MMT) flake crystals, as shown in Figure 5A^[127]. The growth of CNTs within the sandwich resulted in the exfoliation of the MMT flake crystals to form a composite material with a high aspect ratio. The use of PEO markedly improves the mechanical characteristics of the polymer matrix and enhances ionic conductivity, and the highest ionic conductivity of 2.07 × 10^-5 S cm^-1 was reached at a filler dosage of 10 wt.%. Meanwhile, the electron transport on the CNTs is blocked by the MMT attached, hence decreasing the chance of short-circuiting in the battery. Wang et al. used a solution casting method to cast the CNTs or fullerenes (C₆₀) filled into hyperbranched star polymer HBPS-(PMMA-b-PPEGMA)₃₀ to prepare CSSEs. The addition of carbon material only decreases the T_g of the polymer matrix slightly^[128]. However, the decomposition temperature of both CSSEs prepared can reach 383 °C due to the high thermal stability of the carbon material itself. The ionic conductivity of the 0.2 wt.% CNT-filled electrolyte attains 1.06 × 10^-5 S cm^-1, while the lithium-ion transference number is 0.52, which shows high chemical stability above 5.2 V. In contrast, all the performances of the C₆₀-filled electrolyte are poorer. There are few studies on carbon materials as fillers, and there are still many questions about the detailed mechanism of ionic conductivity enhancement of CSSEs, which need further research.

Figure 5. (A) Schematics of montmorillonite clay-CNT hybrid fillers and SEM images of clay-CNT structures at various magnifications^[127]. Reproduced from Ref.^[127] with permission. Copyright 2012, American Chemical Society. (B) Migrating behavior of Li ions on PVC-s-MMT interphase and PVC electrolyte. (C) Galvanostatic cycling curves of lithium symmetrical battery utilizing the PVC, PVC-MMT, and PVC-s-MMT at the current density of 0.1 mA cm^-2. (D) Rate capacity of PVC-s-MMT at different current densities. (E) Discharge/charge profiles for Li|PVC-s-MMT|LFP battery at 1 C at room temperature^[129]. Reproduced from Ref.^[129] with permission. Copyright 2024, American Chemical Society. (F) Schematic illustration of OMMT promoting lithium-ion migration^[130]. Reproduced from Ref.^[130] with permission. Copyright 2023, Wiley-VCH.

The research on carbon materials as fillers for CSSEs is relatively limited, primarily due to the electronic conductivity of carbon materials. The advancement of interface design engineering to address the conduction issue in CSSEs filled with carbon materials at the electrode interface is anticipated to drive further progress in this area.

Clay materials fillers

MMT is a more widely studied material because of its high specific surface area and cation exchange capacity. It is a silica-aluminate with a unique layered structure, with one Al-O octahedra layer sandwiched between two Si-O tetrahedra layers and with suspended hydrated cations in the MMT layer. Li ions can easily replace these cations, and the migration energy barrier of Li ions is as low as 0.16 eV. The -OH groups are also present at the edges of the MMT to form Lewis acid sites, which can anchor the TFSI^-s and thus facilitate the transport of Li ions. Zhu et al. cured the PVC-s-MMT, which is rich in hydroxyl groups on the surface, with poly(vinyl carbonate) (PVC) by a cross-linking agent so that it can be dispersed uniformly and have good interfacial compatibility^[129]. The conductive mechanism of PVC-s-MMT is illustrated by density functional theory (DFT) calculations. As shown in Figure 5B, the coupling technique enhances interfacial channels for lithium ions, diminishes their interaction with the oxygen atom in the carbonyl group, and reduces the migration barrier for Li ions. The ionic conductivity is 7.9 × 10^-5 S cm^-1 when the dosage of MMT is 2 wt.%. As shown in Figure 5C, the Li|PVC-s-MMT|Li battery has a very low overpotential. Meanwhile, as shown in Figure 5D and E, the battery consisting of a Li anode and an LFP cathode exhibits good multiplicative performance and reversibility with a smooth platform. Zhao et al. prepared LOPPM electrolyte films by adding 20 wt.% organically modified MMT (OMMT), which is rich in Lewis acid centers, to poly(methyl methacrylate) (PMMA)/PVDF/P[VDF-hexafluoropropylene (HFP)], which possessed ionic conductivity of 1.1 × 10^-3 S cm^-1 and lithium-ion transference number of 0.54 at room temperature^[130]. This was attributed to the interaction of MMT with PVDF, which uniformly dispersed in PMMA and constructed unique multi-facial channels. As shown in Figure 5F, the negative charge on the surface of OMMT adsorbs Li ions and thus facilitates the transport of Li ions. Moreover, LOPPM exhibits excellent stability and compatibility with lithium anode; the assembled Li|LOPPM|Li symmetric battery can be cycled continuously for 200 h.

Porous materials fillers

MOFs enhance electrolyte performance by leveraging their elevated specific surface area and additional characteristics^[131,132]. The elevated porosity and specific surface area of MOF materials dictate their high surface energy, which can adsorb some side reaction products and impurities. Moreover, MOF materials have diverse structures that can be adapted to various materials so that these materials can be easily compounded with other materials.

Due to its high surface polarity, the electrochemical performance can be enhanced by modulating the Lewis acid-base interaction in the electrolyte system. Lu et al. designed a bifunctional N-CP@Sn-MOF material, an excellent precursor of lithium-philic accelerators, which can be used as a good filler in polymers^[133]. By utilizing nitrogen atoms in the host material as nucleation sites for lithium, the overpotential of lithium nucleation can be reduced, facilitating the consistent deposition and stripping of lithium metal on the host material during the charging-discharge cycle. Filling it into the PEO matrix increased the lithium-ion transference number from 0.194 to 0.367, and the electrochemical window from 3.8 to 4.5 V. Wang et al. utilized the adsorption capacity of the typical MOF HKUST-1 (Cu) to accommodate Li ion liquid (Li-IL) in the pores to form ion-conducting Li-IL@HKUST-1^[134]. The Li-IL@HKUST-1 additive, which exhibits ionic conductivity, can efficiently enhance the polymer matrix by a novel ionic conductive network. This leads to an improvement in ionic conductivity and guarantees the even distribution of lithium. The CSSE was prepared by solution casting after mixing Li-IL@HKUST-1 with PEO, and the ionic conductivity was 1.2 × 10^-4 S cm^-1 at 30 °C.

A molecular sieve is a substance characterized by homogeneous micropores, extensively utilized across various domains, and serves as a typical low-cost, environmentally sustainable filler^[135]. Molecular sieve materials have more robust Lewis acid centers in their backbones and channels, making them more effective in reducing polymer crystallization. Jiang et al. filled SBA-15 molecular sieve material into PVDF-HFP, which has stable electrochemical properties because the backbone unit of SBA-15 consists of SiO₂ and has a larger pore size, making the structure more stable^[136]. The prepared CSSE has an ionic conductivity of 5 × 10^-5 S cm^-1 at room temperature, and SBA-15 can also absorb a small number of impurities in the electrolyte, reducing the effect of impurities on the Li anode. Zeolite is a typical microporous material among molecular sieves with extremely high specific surface area. Chang et al. used commercially available zeolite to condition a typical carbonate-based electrolysis [1 mol L^-1 LiPF₆-EC/dimethyl carbonate (DMC)] without adding any additional salt or additives^[39]. The prepared electrolyte could exhibit enhanced antioxidant stability at 4.6 V and 55 °C. Long cycle life and ultra-stable Li|CSSE|NCM-811 batteries were successfully obtained with 83.2% capacity retention after more than 1,000 electrochemical cycles. Ding et al. developed an electrolyte utilizing physicochemical X-type zeolite (Li-X), incorporating 3 wt.% Li-X into PVDF-HFP^[137]. The resultant electrolyte layer demonstrated a notable ionic conductivity of 1.98 × 10^-4 S cm^-1, a substantial lithium-ion transference number of 0.55, and an extensive electrochemical window of 4.7 V. Zeolites featuring Lewis acid sites enhance the availability of free Li ions and serve to anchor the N,N-dimethylformamide (DMF) and PVDF-HFP polymer chains, thereby reducing polymer crystallinity and safeguarding the lithium metal anode.

These novel materials facilitate enhanced interaction between fillers and polymers, thereby improving the application performance of CSSEs in terms of mechanical properties, electrochemical windows, and electrode/electrolyte interface. This improvement is achieved not only through the reduction of polymer crystallinity to enhance ionic conductivity but also by providing a greater development potential than traditional metal oxide fillers.

Unlike inert fillers, active fillers contain Li ions and rely on the movement of vacancies or interstitial ions to rapidly transport Li ions. Therefore, active fillers are believed to be more efficacious in enhancing the electrochemical performance of CSSEs. The types of active fillers studied include perovskite, garnet, NASICON, and so on.

Perovskite and anti-perovskite type fillers

The general chemical formula of perovskite is ABO₃, as shown in Figure 6A, where A and B are the cations’ six and twelve oxygen coordination sites, respectively. The A site is the octahedral void, typically inhabited by cations with greater ionic radii, such as Na⁺, K⁺, Ca²⁺, Ba²⁺, La³⁺, etc. The B site is the center of the octahedron, which is occupied by cations with smaller ionic radii and forms an octahedral coordination site with the six oxygen ions, such as Sc³⁺, Al³⁺, Ti⁴⁺, Ge⁴⁺, Nb⁵⁺, and so on^[138]. The most studied perovskite fillers are namely Li_3xLa_2/3-xTiO₃. Generally speaking, LLTO is a cubic crystal system, and a large number of TiO₆ octahedra connected by co-topography form the basic skeleton of LLTO crystals, and Li ions and La³⁺ can be homogeneous in the interstices of these octahedra^[139,140]. However, LLTO will also have tetragonal phases, leading to the uneven distribution of La³⁺. The ion vacancy mechanism predominantly governs the migration of Li ions in LLTO. The vacancy of La sites is necessary for lithium-ion conductivity, and the dimensions of the anionic A and the vacancy concentration will influence the ionic conductivity of perovskite electrolytes. The LLTO-type electrolyte is stable and relatively simple to prepare, and its ionic conductivity can reach up to 10^-3 S cm^-1. Nevertheless, the assessed ionic conductivity frequently falls below 10^-4 S cm^-1. Table 2 shows the crystal synthesis method and properties of some perovskites and anti-perovskites.

Figure 6. (A) Crystal structure of cubic perovskite^[138]. Reproduced from Ref.^[138] with permission. Copyright 2021, Springer Link. (B) Octahedral framework model of cubic X₃BA anti-perovskite^[153]. Reproduced from Ref.^[153] with permission. Copyright 2022, American Chemical Society. (C) Distribution of Li ions in tetrahedral (24d) and octahedral positions (48g and 96h)^[157]. Reproduced from Ref.^[157] with permission. Copyright 2020, American Chemical Society (D) Schematic illustration of the process for preparing CD-TFSI supramolecular combination modified LLZTO. (E) Schematic illustration of lithium-ion channel in PEO/LLZTO@CD-TFSI^[176]. Reproduced from Ref.^[176] with permission. Copyright 2022, Wiley-VCH. (F) Crystal structure of LiTi₂(PO₄)₃ and M1, M2, and M3 sites of Li ions^[179]. Reproduced from Ref.^[179] with permission. Copyright 2021, Wiley-VCH. (G) Distribution of Al/Li phases in LATP and LAGP with varying Al concentrations. Yellow regions signify deficient Al/Li phases, while blue regions denote abundant Al/Li particles. Green areas represent uniform materials^[181]. Reproduced from Ref.^[181] with permission. Copyright 2014, Elsevier. (H) Variation of the CSSE films with different LAGP contents and PE skeleton at different temperatures and holding times^[196]. Reproduced from Ref.^[196] with permission. Copyright 2022, Elsevier.

LLTO electrolytes were used directly as an SSE layer to match electrodes and make solid-state batteries. However, the rigid electrolyte/electrode interface creates a tremendous interfacial impedance that limits this application. In addition, a significant problem with LLTO is that when LLTO is in contact with lithium metal, lithium can easily and rapidly insert into the LLTO structure, reducing Ti⁴⁺ to Ti³⁺, increasing the electronic conductivity, and causing an internal short circuit^[151]. Therefore, subsequent studies started to compound LLTO with PSEs, effectively solving the problem of poor interfacial contact while preventing direct contact between LLTO and Li. Romero et al. added LLTO of different sizes to PMMA, and an order of magnitude enhanced the ionic conductivity of the CSSE compared to the pure PMMA electrolyte^[152]. The CSSE had the highest room temperature ionic conductivity (1.13 × 10^-4 S cm^-1) when the average grain size of LLTO was 50 nm. In addition to LLTO, Li_3/8Sr_7/16Ta_3/4Zr_1/4O₃ (LSTZ) has attracted many studies due to its excellent stability in humid environments. Research indicates that nearly no insulating phase develops on the surface of LSTZ, leading to minimal interfacial resistance between LSTZ and PEO. Xu et al. prepared a PEO/LiTFSI/LSTZ electrolyte using LSTZ as fillers, which had ionic conductivity of 5.4 × 10^-5 S cm^-1 at 25 °C^[45]. Perovskite-type electrolytes can efficiently impede the crystallization of polymer electrolytes while offering strong ionic conductivity, but to obtain perovskite electrolyte materials with stable crystal structures, high sintering temperatures or the addition of various sintering aids are often required. Therefore, the cost of perovskite materials is generally high, and there remains a large gap between the performance of perovskite electrolytes prepared by low-temperature sintering and those prepared under high-temperature conditions.

Anti-perovskite has the same crystal structure and stoichiometric ratios as perovskite, with the general chemical formula X₃BA. However, unlike perovskite, as shown in Figure 6B, the anion (B) in anti-perovskite occupies the center of the octahedron. At the same time, the cation (X) is located at the apex of the octahedron^[153,154]. X is a monovalent cation, generally an alkali metal; B is a divalent anion, such as O^2- and S^2-, and A is a monovalent anion, generally for the halogen elements^[155]. Unlike perovskite electrolytes, the concentration of Li ions is higher in anti-perovskite. Li ions in perovskite structures are generally located in the A or B sites, with a maximum lithium-ion concentration of only 20 at%, primarily due to the necessity for cations with greater radii to occupy the A sites, hence ensuring the stability of the crystal structure. In contrast, the lithium-ion concentration in the X-site of perovskite can reach 60 at%, which has the potential to dramatically increase the ionic conductivity of the electrolyte and solve the problems caused by the reduction of high-valent metals in conventional perovskite electrolytes.

Ye et al. prepared CSSE by incorporating lithium-rich anti-perovskite Li₂OHBr as a filler into PEO^[156]. At 5 wt.% Li₂OHBr dosage; the electrolytes had a conductivity of up to 4.5 × 10^-4 S cm^-1 (60 °C). The electrolyte containing Li₂OHBr improves the cycling stability and multiplication performance in batteries with a 3D-structured lithium/copper mesh composite anode and LFP cathode. The easy processing, low cost, and environmental friendliness of the anti-perovskite electrolyte make it more readily available for application, and it has exceptional stability when in contact with the lithium metal. However, the electrochemical window of perovskite electrolytes is narrow, and they generally decompose around 3 V, which makes them less stable for most of the commonly used cathode materials.

While perovskite electrolytes offer broad options and demonstrate excellent performance, their preparation often necessitates harsh conditions. For instance, producing cubic crystalline LLTO electrolytes requires very high sintering temperatures or the addition of various sintering additives and new synthesis technologies. Additionally, the preparation process for anti-perovskite often calls for an inert atmosphere. Consequently, achieving large-scale application of CSSEs based on perovskite-type fillers presents a significant cost challenge.

Garnet type fillers

The general chemical formula of a typical garnet type is A₃B₃C₂O₁₂, with A = Ca, Mg, Y, La or rare-earth metals, B = Bi, Ge, Al, and C = Al, Fe, Ga, Ge, Mn, Ni, or V. Its A, B, and C are 8, 4, and 6 oxygen-coordinated cationic sites, respectively, which crystallize in a face-centered cubic structure. As shown in Figure 6C, Li ions occupy various vacancies within the garnet structure, namely tetrahedral vacancies (24d), octahedral vacancies (48g), and eccentric octahedral vacancies (96h)^[157-159]. The above tetrahedra and octahedra share faces to form a 3D network structure, allowing fast lithium-ion transport. The enhancement of ionic conductivity due to the incorporation of garnet electrolytes into the polymer matrix is closely associated with the space charge region. The presence of space charge regions enables the migration of Li ions from garnet lattice sites to surface sites, especially in the presence of negatively charged vacancies inside the lattice and positively charged Li ions at the surface^[160,161]. When the space charge layer regions near each nanoparticle are interconnected, a channel available for rapid lithium-ion transport is constructed. The garnet-based electrolyte that has been studied more is the LLZO-type electrolyte. Murugan et al. first proposed LLZO-type electrolytes in 2007, which can reach a room-temperature ionic conductivity of 3 × 10^-4 S cm^-1, which attracted the attention of many scholars, resulting in many studies on the modification of the LLZO-based electrolyte^[162]. Table 3 shows the crystal synthesis method and properties of some Garnets. Chen et al. used a casting method to prepare the CSSE film from PEO and 7.5 wt.% LLZO^[173]. The compliant electrolyte film had the maximum ionic conductivity (5.5 × 10^-4 S cm^-1) at 30 °C due to the synergistic effect of cubic LLZO.

LLZO structures can form both cubic and tetragonal phases, and many studies have shown that cubic-structured LLZOs have higher ionic conductivity. To obtain stable cubic-structured LLZO, the three cationic sites of LLZO are usually doped and modified using cations. Karthik et al. synthesized Al-doped Li_6.28Al_0.24La₃Zr₂O₁₂ (Al-LLZO) electrolyte using the solid-state reaction method, which was added to the PEO/LiClO₄ matrix to prepare a CSSE film^[174]. When the amount of Al-LLZO was 20 wt.%, the CSSE had an ionic conductivity of 4.4 × 10^-4 S cm^-1 at 30 °C. Nguyen et al. prepared co-doped LLZO (NAL) with Al³⁺ and Nb⁵⁺ using the sol-gel method, and X-ray diffraction (XRD) proved that NAL had a stable cubic phase^[175]. The PEO-NAL-LiTFSI-SN electrolyte film prepared with PEO as the substrate possessed an ionic conductivity of 3.09 × 10^-4 S cm^-1 at room temperature, and NAL not only facilitated the lithium-ion transport channel but also immobilized anions to enhance the solubility and transport of more Li ions. In addition, there have been more studies on Ta-doped Li_6.75La₃Zr_1.75Ta_0.25O₁₂ (LLZTO) electrolytes, and the CSSE system of LLZTO with PEO is considered to have a promising future. However, the surface wettability of PEO and LLZTO is poor, and the interfacial problem between the two has always existed. The interface between PEO and LLZTO can be efficiently modified to enhance their physical and electrochemical properties. Lu et al. designed β-cyclodextrin (CD) supramolecules containing self-assembled LiTFSI salts to modify the interface of LLZTO^[176]. The preparation process is shown in Figure 6D, where LiTFSI self-assembles within the cavity of β-CD through host-guest interactions, while β-CD associates with LLZTO via hydrogen bonding. The electrolyte film of PEO/LLZTO@CD-TFSI prepared by compositing with PEO had an ionic conductivity of 8.7 × 10^-4 S cm^-1 at 60 °C. This was ascribed to the ability of β-CDTFSI to significantly diminish the surface energy of LLZTO nanoparticles, which allowed LLZTO to be more uniformly dispersed in PEO. Moreover, β-CDTFSI with a unique cavity structure constructed a new lithium-ion channel between PEO and LLZTO, as shown in Figure 6E; the lithium-ion transport channel transitioned from the PEO matrix to LLZTO and subsequently back to PEO, resulting in a lithium-ion transference number of 0.48.

Huo et al. added 20 and 80 wt.% LLZTO to PEO, respectively, to investigate the impact of filler concentration on CSSEs^[177]. With the increase in filler concentration, the ionic conductivity of electrolyte films decreased from 1.6 × 10^-4 to 3.2 × 10^-5 S cm^-1 due to the transport channel of Li ions from migrating along the PEO-LLZTO interface to dispersed LLZTO interparticle migration. However, the tensile strength of the CSSE was elevated from 4.3 to 11.6 MPa. The influencing factors of mechanical properties mainly lie in the concentration of inorganic fillers and the degree of bonding with polymers. Generally, an increase in filler concentration enhances the strength of CSSEs; however, due to a low bonding degree between filler and polymer, a high filler content may promote polymer fluidity and consequently reduce the tensile strength of CSSEs^[178].

In addition to its use as a polymer filler, many studies have shown that the use of LLZO directly as an electrolyte layer is compatible with electrodes. Nonetheless, a significant challenge in LLZO electrolytes is the elevated resistance at the electrode/electrolyte contact, attributed to the formation of a Li₂CO₃ layer on the electrolyte’s surface. The interfacial layer obstructs the movement of Li ions at the electrode/electrolyte interface, resulting in unequal current distribution and ultimately promoting the formation of lithium dendrites.

NASICON type fillers

NASICON electrolytes were first called sodium superionic conductors, and their general chemical formula is AM₂(PO₄)₃. Alkali metal ions (Li⁺, Na⁺, K⁺) are generally used at the A-site to transport cations. According to the elements in the M-site, NASICON electrolytes can be classified into three main types: LiZr₂(PO₄)₃ (LZP), LiTi₂(PO₄)₃ (LTP), and LiGe(PO₄)₃ (LGP). NASICON consists of PO₄ tetrahedra and MO₆ octahedra, as shown in Figure 6F; the upper oxygen atoms are bonded to two MO₆ octahedra and three PO₄ tetrahedra, with Li ions located in the M1, M2, and M3 locations, respectively, which together form a three-dimensional lithium-ion channel^[179,180]. It has been found that the use of Al³⁺ with a smaller radius to replace the M-site ions can not only increase the lithium-ion concentration but also reduce the unit size of the crystal, thereby shortening the lithium-ion transport distance, which will make the electrolyte structure more compact and improve the lithium-ion transport. Therefore, the research on NASICON-type electrolytes in recent years has mainly focused on Li_1+xAl_xGe_2-x(PO₄)₃ (LAGP) and Li_1+xAl_xTi_2-x(PO₄)₃. When the Al³⁺ content is low (x < 0.2), the excess Li ions tend to be trapped near Al³⁺, which leads to higher grain boundary (GB) resistance and lower ionic conductivity^[181]. In turn, when the Al³⁺ content is too much (x > 0.7), the formation of non-Li ions conducting the AlPO₄ phase near the GBs will also be unfavorable to the ionic conductivity, and in general, the optimum Al³⁺ content is considered to be around 0.3-0.5. It has been pointed out that the ionic conductivity of LATP and LAGP is related to the inhomogeneous phase of Al in their crystals; as shown in Figure 6G, LATP and LAGP have uniformly distributed Al/Li phases and exhibit good ionic conductivity when x is 0.4 and 0.5, respectively^[182]. Table 4 shows the crystal synthesis method and properties of some NASICONs.

The LATP and LAGP electrolytes exhibit elevated room-temperature ionic conductivity and commendable chemical stability, and high oxidation voltages (6 V), among other advantages, and they are inexpensive. Compared to other electrolytes, the excellent high stability of LATP and LAGP to H₂O and CO₂ allows for stable preparation and use in air. However, side reactions occur when they come into contact with the lithium anode; Ti⁴⁺ in LATP-type electrolytes is reduced to Ti³⁺, and Ge⁴⁺ in LAGP-type electrolytes is reduced to Ge²⁺ and Ge⁰, resulting in a severe increase in the resistance at the interface until the battery fails^[193,194]. Therefore, LATP and LAGP are often not directly matched with electrode materials to be assembled into lithium-ion batteries but must be treated at the interface or combined with polymer electrolytes.

Sun et al. prepared Li_1.5Al_0.5Ge_1.5(PO₄)₃, using the lower-cost inorganic germanium (GeO₂) as a precursor by sol-gel method^[195]. The highest ionic conductivity of the LAGP samples prepared by sintering at 900 °C for 8 h reached 7.76 × 10^-4 S cm^-1. Huang et al. prepared electrolyte films by filling the LAGP with Polyphenylene Oxide (PPO) and transformed ceramic-in-polymer into polymer-in-ceramic by modulating the dosage of LAGP^[196]. The electrolyte prepared by filling 75% LAGP has an ionic conductivity of 3.46 × 10^-4 S cm^-1, lithium-ion transference number of 0.83, a voltage stability window of 4.78 V. As shown in Figure 6H, the PPO-75% LAGP electrolyte film can remain stable at high temperatures, implying that it can work stably at high temperatures.

Doping metal atoms is a common technique to enhance the ionic conductivity of NASICON electrolytes. Panda et al. prepared LATP electrolytes doped with cobalt and copper (Co-LATP and Cu-LATP), respectively^[197]. Both Cu and Co can reduce the grain size, and the resultant GBs facilitate the migration of Li ions in LATP, enhancing ionic conductivity. After the two LATP electrolytes were prepared into CSSE films with PVDF-HFP and LiTFSI, the highest ionic conductivity at room temperature was 3.18 × 10^-4 S cm^-1 (Co-LATP) and 2.99 × 10^-4 S cm^-1 (Cu-LATP), respectively.

It has been found that modulating the morphology and arrangement of LATP can improve the ionic conductivity of CSSEs. Lu et al. prepared LATP with microporous channels (p-LATP) using polyethylene-polypropylene glycol F127 as a template, and this long-range ordered tubular pore allows for the rapid transport of Li ions^[198]. p-LATP with a filling amount of 40 wt.% of the ionic conductivity of the prepared PEO-p-LATP electrolyte was 3.5 × 10^-4 S cm^-1. Moreover, the ionic conductivity may be further improved by combining p-LATP with standard dense LATP using microporous channels and dense lithium-ion conductive areas.

The primary issue constraining the utilization of LLZO and LATP electrolytes is the side reaction occurring between them and the lithium metal contact. Although the existence of a polymer matrix has alleviated this problem to a large extent, there is still no complete solution. At the same time, the morphology and arrangement of the two compounds in the polymer greatly influence the properties of CSSEs, and advanced synthesis methods are needed to obtain materials with uniform properties.

1D fillers

For inert fillers, due to the agglomeration of nanoparticles and non-ionic conductivity, the content of inert fillers should be manageable. Otherwise, it will rather reduce the ionic conductivity of the electrolyte. In this case, using 1D fillers can not only slow down the agglomeration problem, but also, 1D inert fillers have longer dimensions, providing a more continuous lithium-ion transport channel. TiO₂ is a relatively joint 1D inert filler. As shown in Figure 7A, the filling of nanowires can provide more rapid lithium-ion channels^[201]. Hua et al. filled the TiO₂ nanorods into the poly(propylene carbonate) (PPC), the utilization of rod-shaped materials rather than randomly dispersed particles to establish a continuous transport channel for Li ions, the preparation of CSSE with an electrochemical window of 4.6 V and an ionic conductivity of 1.52 × 10^-4 S cm^-1 at room temperature^[202]. Luo et al. fabricated TiO₂ micro-rods with a high concentration of oxygen vacancies^[203]. These vacancies not only decreased the crystallinity of the polymers but also facilitated the dissociation of LiTFSI by interacting with the -CF₃ group of the TFSI^- on the surface of the TiO₂ microrod. The preparation process is shown in Figure 7B. The ionic conductivity reached 1.04 × 10^-4 S cm^-1 at room temperature.

Figure 7. Schematic illustration of (A) nanowires providing fast lithium-ion channels^[201]. Reproduced from Ref.^[201] with permission. Copyright 2017, Nature. (B) The preparation of TiO₂ nanorods enriched with oxygen vacancies^[201]. Reproduced from Ref.^[203] with permission. Copyright 2022, Elsevier. (C) The preparation of Bi/HMT-MOFs nanowire and BMCPE film^[204]. Reproduced from Ref.^[204] with permission. Copyright 2022, Wiley-VCH. (D) Lithium-ion migration in VSB-5 enhanced PEO electrolytes^[205]. Reproduced from Ref.^[205] with permission. Copyright 2020, Elsevier. (E) A side view of the LiTFSI crystal structure derived from DFT calculations is adsorbed on the surface of MgO (200) and LMO (200), containing one oxygen vacancy. (F) Channels of a single Li ion severed from the LMO lattice, depicted without (left) and with (right) an oxygen vacancy^[206]. Reproduced from Ref.^[206] with permission. Copyright 2023, American Chemical Society. (G) Evaluate the performance of Li|Li-MPSE|LFP battery and Li|PSE|LFP battery at a discharge rate of 0.5 C at 30 °C^[207]. Reproduced from Ref.^[207] with permission. Copyright 2023, Elsevier.

The MOF material can additionally be synthesized as a 1D nanowire to integrate into the polymer electrolyte matrix. Xu et al. designed a colloidally dispersed nonporous MOF (Bi/HMT-MOF) nanowire with a shallow surface area using Bi³⁺ as the metal center^[204]. The preparation process is shown in Figure 7C. The 1D structure enables Bi/HMT-MOF to have better compatibility with PEO. Bi³⁺ facilitates the dissociation of lithium salts and interacts with EO to form an amorphous PEO interfacial layer on the nanowire surface, enhancing the fast transport of Li ions. The prepared CSSE has a conductivity of 5.89 × 10^-4 S cm^-1 at 60 °C.

Besides conventional nanofillers such as metal oxides, materials featuring 1D channels can also be utilized to enhance polymer electrolytes. Wu et al. used nickel phosphonate VSB-5 as a filler for PEO, which not only reduces the crystallinity of PEO, as shown in Figure 7D, but also has its large pore size of 1.1 nm 1D channels^[205]. When the amount of VSB-5 was 3 wt.%, the prepared PEO-LiTFSI-3 wt.% VSB-5 had an ionic conductivity of 4.83 × 10^-5 S cm^-1 at 30 °C. Furthermore, the incorporation of VSB-5 into the CSSE enhances its thermal stability and electrochemical stability window (4.13 V) compared to the pure PEO electrolyte. Furthermore, it enhances the inhibition of lithium dendrite development.

Similar to nanoparticles, the diffusion kinetics of Li ions can be effectively improved by rational surface defect design of 1D oxide nanofillers. Positively charged oxygen vacancies represent the predominant surface defect configuration, achievable by substituting high-valent elements with those of lower valence. This surface imperfection primarily enhances the ionic conductivity of electrolytes by facilitating the dissociation of lithium salts, hence releasing a greater quantity of Li ions. Yu et al. prepared lithium-doped MgO (LMO) nanofibers by using electrostatic spinning and calcination, and after compositing them with PEO/LiTFSI, the inorganic filler with 1D morphology firstly reduces the crystallinity of the polymer and provides a continuous Li transport channel^[206]. In addition, lithium doping can introduce many positively charged oxygen vacancies in the lattice structure, enabling LMO to adsorb more anions from Li salts and thus release more Li ions. As shown in Figure 7E, when Li is immobilized in MgO, supplementary Li-N and Li-O bonds can be established, resulting in an adsorption energy of TFSI^- at -6.63 eV, which signifies an enhancement in the obstruction of free TFSI^-. The potential barriers required for Li to spill out of the lattice are about 2.12 and 1.70 eV for LMOs with and without oxygen vacancies, respectively, as shown in Figure 7F. This indicates that flaws or compromised chemical bonds may promote the movement of Li ions adjacent to oxygen vacancies, enabling them to surmount the energy barrier and exit the lattice. In addition to utilizing 1D morphology fillers, constructing vertically oriented 1D Li ion channels can achieve the same effect. Wang et al. designed a CSSE with vertically aligned 1D lithium-ion transport channels^[207]. Lithium MMT is embedded in PVDF-HFP along a 1D vertically aligned channel to prepare a CSSE (Li-MPSE), which provides a one-dimensional efficient transport channel for Li ions. At the same time, the Li dissolution environment in the vinyl carbonate-based electrolyte was altered to establish the lithium-ion hopping pathway. This CSSE has an ionic conductivity of 1.99 × 10^-4 S cm^-1 and a lithium-ion transference number of 0.73. The rate performance test further demonstrates the superiority of Li-MPSE, as shown in Figure 7G.

When the polymer electrolyte matrix is filled with active fillers, Li ions are more inclined to be transported at the polymer/particle interface. Therefore, when the active fillers are designed as a 1D structure, the number of lithium-ion migrations between multiple filler particles can be reduced so that Li ions can be transported along the 1D active filler in a long-range and rapid manner, improving the ionic conductivity of CSSEs^[208]. In general, active fillers often require electrostatic spinning to prepare nanowires with fiber networks, followed by high-temperature sintering to obtain 1D active filler nanorods.

Zhu et al. prepared LLTO nanowires (LLTO NWs) using electrostatic spinning and high-temperature sintering, which were filled into PEO to prepare CSSEs^[209]. The ionic conductivity of the CSSE is 5.53 × 10^-5 and 3.63 × 10^-4 S cm^-1 at room temperature and 60 °C, respectively, when the amount of LLTO was 5 wt.%. Figure 8A shows the LLTO NWs prepared by electrostatic spinning, and the results show that these NWs have a diameter of about 1 μm and can be dispersed uniformly into the PEO matrix. As shown in Figure 8B, the assembled Li|CSSE|LFP battery exhibits good specific capacity and capacity retention at 0.5 C. Bi et al. took the prepared LLTO NWs and prepared PAN/SEO by using a second utilization of the PAN/SN/LLTO CSSE film prepared by electrostatic spinning^[210]. This CSSE film can exhibit an ionic conductivity of up to 2.2 × 10^-3 S cm^-1 at 30 °C and a broad electrochemical window of 5.1 V. In addition to LLTO-type electrolytes, garnet-structured LLZO electrolytes are often prepared as 1D structures to improve the performance of CSSEs. Li et al. also prepared LLZO nanofiber materials using electrostatic spinning and sintering processes, which were added to PVDF-HFP/LiTFSI to prepare a CSSE film^[211]. The CSSE containing 10 wt.% of LLZO was found to have an ionic conductivity of 9.5 × 10^-4 S cm^-1 at 20 °C. The assembled Li|PVDF-HFP/LiTFSI/LLTO|LFP battery can be stabilized with 140 mAh g^-1 capacity at 0.2 C for 120 cycles.

Figure 8. (A) Images of LLTO nanowires prepared by electrostatic spinning, surface, and cross-section images of prepared CSSE films. (B) Cycling performance of Li|PEO/5% LLTO NWs|LFP battery at 0.5 C^[209]. Reproduced from Ref.^[209] with permission. Copyright 2022, Elsevier. (C) Schematic illustration of ionic transport mechanisms in CSSEs with 10 wt.% LLZN nanowires^[212]. Reproduced from Ref.^[212] with permission. Copyright 2018, Elsevier. (D) Ionic conductivity plots of CSSEs with randomly oriented nanowires and nanowires arranged in different directions^[201]. Reproduced from Ref.^[201] with permission. Copyright 2017, Nature.

Cation doping techniques have also been applied in the design of nanowires, where the dissociation of lithium salts by active fillers can be enhanced by different cation doping, increasing the lithium-ion concentration in CSSEs. Sun et al. designed Nb-doped Li_6.75La₃Zr_1.75Nb_0.25O₁₂ (LLZNO) nanowires^[212], which were blended with PMMA and LiClO₄ to prepare a CSSE with an ionic conductivity of 2.20 × 10^-5 S cm^-1. Figure 8C illustrates the ionic transport channel within the prepared electrolyte, wherein a greater number of Li ions are dissociated through interaction with the anion. Ding et al. designed a tantalum-aluminum co-substituted LLZO nanowire electrolyte, which not only provides long-range interfaces to improve ionic conductivity but also has more lithium vacancies to further promote Lewis acid-base interaction of the electrolyte with lithium salts^[213]. The CSSE, formulated using the PEO matrix, exhibits a lithium-ion transference number of 0.79 and an ionic conductivity of 3.8 × 10^-4 S cm^-1.

Furthermore, it has been noted that well-aligned nanowires can enhance the performance of the CSSE more effectively than a randomly oriented distribution of nanowires. Liu et al. prepared CSSEs filled with randomly oriented LLTO NWs and CSSEs with oriented aligned LLTO NWs, by using PAN as the matrix^[201]. As shown in Figure 8D, the electrolyte with randomly dispersed nanowire fillers had an ionic conductivity of 5.40 × 10^-6 S cm^-1 at 3 wt.% filler and 30 °C, while the electrolyte containing well-aligned nanowires exhibited an ionic conductivity of 6.05 × 10^-5 S cm^-1. Compared to nanoparticles, nanowires provide a more continuous interface for Li ion transport, and well-aligned nanowires can transport Li ions faster than randomly aligned nanowires due to the absence of surface crossings.

2D fillers

To improve the ionic conductivity of the CSSEs, the contact area between the inert fillers and the interface may be augmented. This is because the fillers themselves are unable to transport Li ions, and instead depend on the filler/interface to create a channel for lithium-ion transport. Therefore, compared to 0D and 1D inert fillers, 2D fillers with a layered structure become a better choice^[214,215]. The design of 2D fillers significantly enhances the contact area between the polymer and the filler. Moreover, 0D and 1D fillers produce a part of unwanted volume due to their inability to transport Li ions within themselves, which is unfavorable to ionic conductivity. At the same time, the laminated 2D material reduces this part of unwanted volume very well^[216,217]. However, the disadvantages of 2D materials are also apparent, as they are prone to curling and aggregation when dispersed in a polymer matrix, which makes it challenging to take advantage of their high specific surface area. Simultaneously, the 2D material has a distribution direction that is perpendicular to the direction of lithium-ion transport. This perpendicular distribution is predicted to impede the passage of Li ions between the electrodes.

Sun et al. prepared CSSE by mixing g-C₃N₄ nanosheets with PEO and LiTFSI for the first time^[218]. The 2D configuration of g-C₃N₄ was utilized to create an expanded lithium-ion transport region, hence enhancing ionic conductivity. As the surface of g-C₃N₄ is rich in N atoms, it can better dissociate the lithium salts in the CSSE at 30 °C, the ionic conductivity of the prepared CSSE is 1.7 × 10^-5 S cm^-1, and the lithium-ion transference number reaches 0.56. Silicates are natural ore materials with a 2D layered structure. Zhao et al. used in situ polymerization of allyl acetoacetate (AAA) monomers to prepare LPAS electrolyte films containing SiO₂ nanosheets with oxygen vacancies, and the structure of LPAS is shown in Figure 9A^[51]. The SiO₂ nanosheets can dissociate LiTFSI by immobilizing the anions. The presence of the oxygen vacancies accelerates the migration of Li ions, which results in the LPAS having a 3.82 × 10^-4 S cm^-1 ionic conductivity and 0.66 lithium-ion transference number. Moreover, the 2D structure of SiO₂ forms an intercalation network in PMMA, which reduces the crystallinity while maintaining good viscoelasticity. Wu et al. incorporated 15 wt.% of 2D MoSe₂ sheets into PVDF to develop an electrolyte (PVVMS-15) in which positively charged Mo atoms and negatively charged Se atoms interact with -CF₂- and -CH₂-, respectively^[219]. This interaction enables the conversion of PVDF into the all-trans (β-phase), optimizing the solvation structure and resulting in an ionic conductivity of PVVMS-15 reaching 6.5 × 10^-4 S cm^-1.

Figure 9. (A) Schematic illustration of LPAS SPE^[51]. Reproduced from Ref.^[51] with permission. Copyright 2024, Wiley-VCH. (B) Schematic illustration of the interaction between MXene and SFP^[222]. Reproduced from Ref.^[222] with permission. Copyright 2023, Elsevier. (C) Schematic representation of interface change in a charge equalizer resulting in spherical lithium deposition for batteries^[223]. Reproduced from Ref.^[223] with permission. Copyright 2023, Wiley-VCH. (D) Schematic illustration of LFCPP, LPCPP electrolytes, and lithium deposition mechanism promoted by LFCPP^[224]. Reproduced from Ref.^[224] with permission. Copyright 2022, American Chemical Society.

MOF and covalent organic frameworks (COFs) materials can perform better when they are designed as 2D fillers. By molecular engineering, Xu et al. designed a 2D MOF sheet rich in -NH₂ on the surface. The substituents in this MOF material enhanced the electron-donating effect^[220]. When compounded with PEO and LiClO₄ to prepare an electrolyte, it could better restrict the movement of ClO^4- thus increasing the lithium-ion transference number to 0.64. This electrolyte achieved an ionic conductivity of 6.5 × 10^-5 S cm^-1 at room temperature. Saleem et al. incorporated 2D pyrazine and imine-linked COFs into a PEO electrolyte matrix^[221]. The distinctive amalgamation of electron-rich and polar functions in the COFs facilitates strong interactions with Li ions, yielding an ionic conductivity of 1.86 × 10^-3 S cm^-1 at ambient temperature. Furthermore, the prepared electrolyte film exhibits excellent flexibility and fully recovers after undergoing 360° folding, effectively mitigating the growth of lithium dendrites and interface instability.

MXene, a category of 2D transition metal carbides, is utilized in CSSEs because of its distinctive electrical and mechanical characteristics. Zhao et al. designed Li anodic interfacial layers using silk fibroin peptide (SFP), acetylene black (AB), and MXene and applied them to the LOPPM prepared in their previous work^[222]. As shown in Figure 9B, the formation of LiN₃ and LiF was successfully promoted by utilizing the -NH₂ group in SFP combined with -F in MXene. This work realized the in-situ formation of uniform SEI membranes and effectively suppressed the growth of lithium dendrites. In their other work, LiTFSI/OMMT/soybean isolate protein (SPI)/PVDF (LOSP) electrolytes were prepared, and the above interface design was applied to LOSP^[223]. As shown in Figure 9C, spherical lithium deposition was achieved at the electrolyte/electrode interface, forming a stable interfacial layer that can promote the reversible transport of Li ions. Molybdenum carbide (Mo₂C) reduces the lithium-ion transfer energy barrier and is also often used to improve electrode/electrolyte interface problems. Shan et al. designed LiTFSI/nanosheet Mo₂C (F-Mo₂C)/PVDF/PVDF-HFP electrolyte (LFCPP) and LiTFSI/nanoparticle Mo₂C (P-Mo₂C)/PVDF/PVDF-HFP (LPCPP)^[224]. This study illustrates that F-Mo₂C can create unique charge channels that enhance the uniform deposition of lithium particles, and Figure 9D shows the schematic illustration of the two electrolytes and the LFCPP to promote lithium deposition.

Compared to inert fillers, there has been less research on 2D active fillers. One possible explanation is that the current manufacturing technology cannot produce significant amounts of active nanosheets. At present, the preparation methods involve the use of template and electrostatic spinning techniques. Liu et al. prepared 2D fibrous films filled with PEO by LLTO using electrostatic spinning and then coated with a layer of PEO on both sides of the fibrous films to prevent the LLTO from direct contact with the Li negative electrode to produce a reduction reaction^[225]. In addition, the design of the PEO coating on both sides facilitates optimal interfacial contact between the electrolyte and the electrode, thereby greatly reducing interfacial impedance. The ionic conductivity of this sandwich CSSE is 1.6 × 10^-4 S cm^-1 at room temperature.

Due to the process difficulty of 2D active fillers, most current research involves preparing micron-sized 2D materials using the press-sintering process and using them directly as electrolyte-matched electrodes to prepare LMBs^[226]. Nonetheless, the contact between electrodes and electrolytes requires resolution^[227,228]. Huang et al. solved the reduction reaction of LATP on a lithium-negative electrode by constructing a flexible buffer layer of PVDF and Mg₃N₂ particles on the surface of the LATP sheet^[229]. This interfacial design constructed a multifunctional intermediate layer of Mg, LiF, and Li₃N in situ at the electrolyte/electrode interface, which inhibited the decomposition of LATP. It increased the critical current density of LATP from 0.34 to 0.76 mA cm^-2. Liu et al. prepared Ga-containing LLZO (Ga-LLZO) electrolytes, then incorporated ionic liquid into the Ga-LLZO electrolytes, and subsequently coated ionic liquid-containing PEO (IL-PEO) on the surface of Ga-LLZO electrolytes to improve the contact with the electrodes^[230]. The IL-PEO coating, on the one hand, solves the rigidity problem of Ga-LLZO, which leads to a significant electrolyte/electrode impedance problem. On the other hand, it enhances the stability of LLZO electrolytes in humid air, diminishes the formation of the Li₂CO₃ phase, and increases ionic conductivity. The electrolyte with the coating attained an ionic conductivity of 5.7 × 10^-4 S cm^-1 at ambient temperature.

Preparation method of fillers

There are many mature processes for the synthesis of fillers; the most commonly used methods are introduced here: hydrothermal reaction, solid-state reaction, and sol-gel processing. The hydrothermal method is often used to synthesize materials such as metal oxides and MOFs, and the advantage of the hydrothermal method is that it can form complex structures under relatively mild conditions. Generally, the appropriate raw material is selected according to the target product, and some auxiliary reagents are added after it is dissolved in water. It is then transferred to a high-pressure reactor, sealed, and heated in an oven between 100-300 °C, and the product can be obtained by natural cooling after the reaction^[78].

Active fillers are frequently prepared using solid-state reaction processes and sol-gel processing techniques. The solid phase reaction method is the precursor powder according to the stoichiometric ratio after mechanical grinding mixed and calcined at the appropriate temperature to obtain the material^[141]. The sol-gel processing method is the stoichiometric ratio of raw materials and metal salts dissolved in water or alcohol solvents by adding catalysts or adjusting the pH to promote a hydrolysis reaction to form a highly dispersed suspension (sol). When the particles in the sol are further cross-linked, a three-dimensional network structure (gel) is synthesized, subsequently dried, and calcined to yield the material^[142].

Preparation method of CSSEs

The most straightforward approach for preparing CSSEs is solution blending, as shown in Figure 10A, in which the polymer matrix and filler are dissolved and agitated in an organic solvent to attain uniform dispersion before being formed into an electrolyte film^[231]. Alternatively, melt blending disperses the filler in the molten polymer matrix to prepare the electrolyte film, offering the advantage of avoiding solvent residue compared to solution blending^[232]. Powder blending involves mixing polymer powder and fillers, followed by grinding and hot pressing to prepare the electrolyte film^[233]. This approach can mitigate strong shear forces during melt blending, hence maintaining the mechanical integrity of the electrolyte film to prevent polymer chain rupture. In addition to direct blending, in-situ polymerization is also utilized for CSSE preparation. The general process involves dissolving polymer monomers, initiators, fillers, and lithium salt in a suitable solvent before casting onto electrodes and solidifying using ultraviolet or thermal initiation^[234,235]. While this method effectively addresses electrode/electrolyte interface contact issues, it may result in significant solvent residues.

Figure 10. CSSEs preparation method (A) Solution casting method^[231]. Reproduced from Ref.^[231] with permission. Copyright 2022 Springer Nature. (B) Electrostatic spinning method^[95]. Reproduced from Ref.^[95] with permission. Copyright 2016, PNAS. (C) 3D printing method^[235]. Reproduced from Ref.^[237] with permission. Copyright 2019, American Chemical Society. (D) Template method^[238]. Reproduced from Ref.^[238] with permission. Copyright 2018, Royal Society of Chemistry.

In addition to the above methods, there are several techniques for building CSSEs with 3D frame structures. Examples include the electrostatic spinning method, 3D printing method, and template method. An excellent advantage of these methods is the continuous structure in CSSEs. As shown in Figure 10B, the electrostatic spinning method utilizes a DC high-voltage power supply to mix the electrolyte solution in a syringe to form nanofibers on the collector^[95]. After the solvent is removed, an electrolyte film is obtained. The 3D printing technology can efficiently investigate structure-property relationships, allowing for rapid exploration and analysis^[236]. This method uses extruded polymer adhesives to join multiple filler sheets [Figure 10C] and the prepared CSSEs have extremely high flexibility and can reduce strain during the use of the material^[237]. The template method generally adds the filler to the template first [Figure 10D], then makes the filler have a special structure by sintering and removing the template, and finally infuses the polymer matrix into the filler structure^[238]. Table 5 provides an overview of the advantages and disadvantages of general synthesis methods for CSSEs. Despite the availability of these more mature techniques, it is still a challenge to customize 3D frames with highly ordered structures accurately.

These structural designs augment the interface between fillers and polymers, facilitating continuous transport channels for lithium-ion migration while maintaining low concentration, thereby contributing to improved ionic conductivity of CSSEs. The construction of a continuous lithium-ion transport channel by active fillers within the 3D framework design further enhances ionic conductivity. However, the insulation properties of inert filler to Li ions may lead to a significant reduction in ionic conductivity if a continuous 3D structure of inert filler skeleton is constructed, resulting in a large volume ratio of lithium insulation region in the polymer matrix. As such, there are few discussions in the literature on 3D structural design for inert fillers.

Advanced characterization of CSSEs

Morphology characterization technique

Scanning electron microscopy (SEM) is a widely used technology for the characterization of microscopic morphology in materials science. In addition to imaging microscopic morphology, SEM-based energy spectrum technologies, such as energy dispersive energy spectrometers (EDS), can provide quantitative analysis of the local element composition of materials, reflecting the dispersion degree of fillers in polymer matrix to a certain extent^[239]. The atomic force microscope (AFM) enables the characterization of material morphology at nanoscale resolution and the direct characterization of insulating materials. Furthermore, AFM can be utilized for imaging in various environments. It is capable of characterizing changes in the morphology and mechanical properties of electrolytes at different temperatures; for example, Figure 11A demonstrates the characterization of the morphology, Young’s modulus, and adhesion of PEO/LiClO₄ electrolyte at 30 and 55 °C using in-situ c-AFM^[240]. When lithium deposition is observed within CSSEs by SEM, corrosion occurs during the sample transfer. This challenge can be addressed through in situ optical microscopy, enabling the observation of CSSE behavior during electrochemical reactions in a specific environment. As depicted in Figure 11B-D, cross-sectional snapshots of Li/Li₃PS₄/NCM battery charging are characterized using in situ optical microscopy^[241].

Figure 11. (A) In situ c-AFM characterization of the morphology, Young modulus and adhesion of PEO/LiClO₄ at 30 and 55 °C^[240]. Reproduced from Ref.^[240] with permission. Copyright 2021, Elsevier. Cross-sectional snapshots of Li/Li₃PS₄/NCM battery taken at 0 h (B), 6 h (C), and 18 h (D)^[241]. Reproduced from Ref.^[241] with permission. Copyright 2021, American Chemical Society (E) Cryo-HRTEM images of pristine LLZO grain and GBs. (F) Magnified cryo-HRTEM and (G) FFT images of the area within the red square in (E)^[242]. Reproduced from Ref.^[242] with permission. Copyright 2023, Wiley-VCH.

Moreover, cryopreservation high-resolution transmission electron microscopy (cryo-HRTEM) can provide detailed characterization of structural features within CSSEs, such as nanoscale structure and GB characteristics. Rapidly cooling and freezing CSSEs can help material avoid changes at room temperature. As illustrated in Figure 11E and F, cryo-HRTEM allows for the characterization of LLZO grains and the nanoscale region of GBs, with corresponding Fast Fourier transform (FFT) diagrams [Figure 11G] providing clear insights^[242].

Improve battery cycle stability

The primary factor constraining the cycle stability of ASSLMBs is the proliferation of lithium dendrites. Despite the flexible polymer’s ability to establish effective interface contact, it remains incapable of inhibiting the proliferation of lithium dendrites, hence diminishing battery longevity. The incorporation of fillers is anticipated to address this issue; fillers not only augment the mechanical strength of CSSEs to avert punctures from lithium dendrites, but numerous fillers can also establish a stable interfacial layer at the direct contact between lithium metal and electrolyte, thereby restricting the proliferation of lithium dendrites.

Li et al. added TiO₂ nanotubes to the PEO, and the prepared electrolyte formed a LiF-rich interface layer between the lithium metal and the electrolyte, achieving rapid lithium-ion transport and uniform lithium deposition^[245]. The Li|PEO-TiO₂|LFP battery was cycled more than 3,100 times at 1 C, showing excellent cycle stability. Pang et al. prepared Ga/Ta co-doped LLZO and added it to PVDF to prepare an electrolyte [PL(Ga_0.1Ta_0.3)₂₅]. PL(Ga_0.1Ta_0.3)₂₅ has a polarization voltage of only 0.08 V at a current density of 0.1 mA cm^-2[246]. The polarization voltage of pure PVDF electrolyte is higher than 0.2 V, showing uniform lithium deposition. The assembled Li|PL(Ga_0.1Ta_0.3)₂₅|LFP battery showed a capacity retention rate of 91.1% after 100 cycles at 0.2 C.

The incorporation of fillers can mitigate the formation of lithium dendrites and decrease the incidence of side reactions. For example, the DMF solvent in PVDF transfers Li ions by solvating with Li ions to form [Li(DMF)_x]⁺ complexes. However, the side reaction between DMF and lithium metal leads to low coulomb efficiency and capacity attenuation during the battery cycle. The incorporation of fillers can absorb the solvent within the polymer, facilitate uniform distribution, and establish an inorganic layer at the lithium anode to mitigate the incidence of side reactions^[247]. Li et al. added LATP to the PVDF electrolyte to prepare the electrolyte^[248]. The assembled Li|PVDF/LATP/LiTFSI|LFP battery had an initial discharge-specific capacity of 150 mAh g^-1. After 50 cycles, the specific capacity remains at 130 mAh g^-1, in contrast to the electrolyte devoid of LATP, which exhibits a specific capacity of 54 mAh g^-1 on the 50th cycle; the battery performance has markedly enhanced.

Increase battery energy density

The primary approach to enhance the energy density of a battery is to elevate its charging cut-off voltage, so it is necessary to match the high-voltage electrode material. However, the low electrochemical stability window limits the matching of PSEs to the high voltage electrode, and the filler can improve the electrochemical window of the electrolyte and reduce its resistance to oxidation and decomposition at high potential. Enhancing the intermolecular interaction forces in CSSEs is a common way to improve the electrochemical window^[249].

Wang et al. added SiO₂ to poly(vinyl ethylene carbonate) (PVEC) to prepare the electrolyte^[250]. Hydrogen bonding between the -OH on the surface of SiO₂ and the C=O group on PVEC enhanced the antioxidant capacity of the electrolyte. The electrochemical window of PVEC was elevated from 4.5 to 5.0 V. The capacity retention rate of the built Li|PVEC/SiO₂|LiCo₂ battery remained at 94% after 200 cycles at 25 °C. The capacity retention rate of the built battery lacking SiO₂ electrolyte was merely 20% after 50 cycles. Mu et al. filled directional Li_6.4La₃Zr₂Al_0.2O₁₂ (LLZAO) nanofibers into PEGDA, and the prepared electrolyte not only showed good lithium dendrite inhibition ability but also formed a cathode-electrolyte interface layer rich in LiF on the NCM811 cathode^[251]. It significantly diminishes the frequency of side reactions at elevated pressure circumstances. The assembled Li|PEGDA/LLZAO|NCM811 battery exhibits a high specific capacity of 189.6 mAh g^-1 at 0.1 C.

Conclusion

This review introduces the key parameters of CSSEs and reviews the construction of lithium-ion transport channels. Furthermore, it discusses the effects of different types of fillers on the CSSEs. In this review, a comprehensive analysis of the critical roles played by fillers in CSSEs include: (1) Fillers reduce crystallinity; (2) The interface between filler and polymer creates a lithium-ion channel. Additionally, active fillers that conduct Li ions can offer supplementary transport channels for Li ions; and (3) The size and morphology of fillers exert varying degrees of influence on the performance of CSSEs. The fillers with a higher specific surface area can offer increased lithium-ion transport channels at the filler/polymer interface, which is crucial for enhancing the ionic conductivity of CSSEs. The 3D frame structures prepared using active fillers establish a continuous network for lithium-ion conduction, thus holding promise for overcoming the bottleneck in CSSEs ionic conductivity while achieving enhanced mechanical strength. In addition, the preparation methods and characterization techniques of fillers and CSSEs were also discussed. This review focuses on the intimate interplay between filler materials and the mechanism of lithium-ion transport channels, encompassing the functional role of fillers and their underlying action principles. By synthesizing the latest research findings, an in-depth analysis is conducted to elucidate the structure-activity relationship between fillers and lithium-ion channels, thereby providing a robust foundation for the selection and dosage of fillers while ensuring optimal performance of CSSEs.

Outlook

Although this review summarizes many research works on CSSEs, for now, CSSEs research still has a high potential for exploitation. To achieve early and widespread use of CSSEs, further efforts should be made about the following challenges.