Ultrasmall nanocrystalline CeO₂ fillers improving the performance of PVDF-based polymer electrolytes for lithium metal batteries

Abstract

PVDF-based polymer electrolytes have been widely studied because of their mechanical strength, easy processing, and excellent thermal/chemical stability. However, pure PVDF faces the challenges of low ionic conductivity and insufficient interface stability with the electrodes. Herein, we propose a facile strategy to fabricate Ce-MOF-derived rod-like nanocrystalline CeO₂ into PVDF-HFP polymer electrolytes for Li metal batteries. The composite polymer electrolytes (CPEs) achieved enhanced ionic conductivity and interfacial stability. Notably, the CPEs with ultrasmall nanocrystalline CeO₂ demonstrated superior Li⁺ transport kinetics (5.04 × 10⁻⁴ S cm⁻¹), lithium dendrite suppression and an extended electrochemical stability window up to 4.6 V. When applied in Li|LiMn_0.6Fe_0.4PO₄ full batteries, the quasi-solid polymer electrolyte system maintained a discharge capacity of 94 mAh g⁻¹ after 100 cycles at 0.5C, delivering good cycling stability and rate capability, and retained 99% capacity after 500 cycles at 1C. This study demonstrates that functional CeO₂ nanocrystals with tailored structures can effectively enhance the performance of PVDF-based polymer electrolytes, providing a promising strategy for the development of solid-state lithium metal batteries.

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1. Introduction

Solid electrolytes (SEs), as ion-conductive solid-state materials, represent a significant advancement in battery technology by overcoming the inherent limitations of conventional liquid electrolytes. Compared to their liquid counterparts, SEs exhibit enhanced safety, superior energy density, and extended operational temperature ranges.¹ Their implementation not only substantially improves battery energy density but also mitigates safety risks associated with liquid electrolytes.² In particular, SEs have garnered increasing attention as critical components in all-solid-state lithium metal batteries (LMBs).^3,4 Among these solid electrolytes, polymer electrolytes (PEs) have been extensively investigated due to their facile fabrication processes and excellent interfacial compatibility.³ Nevertheless, PEs face technical challenges that include low room-temperature ionic conductivity and narrow effective voltage windows, primarily attributed to their polymer chain configurations, crosslinking density, or high crystallinity in practical applications, thereby restricting their broader implementation.^5–8 Poly(vinylidene fluoride-co-hexafluoropropylene) (PVDF-HFP) has gained prominence as a polymer matrix for battery electrolytes due to its exceptional dielectric properties, robust thermal stability, and favorable mechanical strength.^9–11 Within the PVDF-HFP matrix, Li⁺ diffusion is enhanced by interactions involving residual solvent molecules and lithium salts. These solvent molecules establish dynamic pathways that enable Li⁺ migration between polymer chains.^12,13 These attributes establish PVDF-HFP as a preferred material for gel or solid polymer electrolytes in lithium-ion battery research.¹⁴ However, PVDF-based electrolytes still face critical challenges, including inefficient lithium salt dissociation, poor compatibility with lithium metal anodes, and uncontrollable dendrite growth, which significantly impede their practical applications.¹⁵

The introduction of inorganic fillers into polymer electrolytes represents an effective method for obtaining composites with concurrently improved ionic conductivity and mechanical strength.^16,17 The enhancement mechanism primarily involves crystallinity reduction of polymer matrices. Furthermore, studies have revealed that filler addition modifies the chemical environment of polymer–lithium salt complexes, facilitating lithium salt dissociation and thereby increasing Li⁺ concentration within the electrolyte.^18,19 To date, various inorganic fillers have been employed for polymer electrolyte modification, which can be categorized into two groups: non-conductive (inert) fillers and conductive (active) fillers.^17,20 The incorporation of electrochemically inactive cerium-based fillers, especially those of the oxide type, has significantly enhanced the performance of PVDF-based polymer electrolytes, including cycling stability and the electrochemical stability window.^21–23

Metal–organic frameworks (MOFs)-crystalline hybrid materials with long-range ordered porosity, exceptional surface areas, and organic–inorganic compositions have demonstrated significant potential as functional materials for enhanced ion conduction.²⁴ In polymer-based battery research, MOF-incorporated composite electrolytes have shown substantial potential in enhancing ionic conductivity and stabilizing interfacial structures.^25,26 PVDF-HFP-based polymer electrolytes can be further optimized by incorporating MOF-derived architectures, thus enhancing their structural properties and performance.²⁷ Our previous studies also showed that the anchoring of CeO₂ nanoparticles on defect-engineered Ce-MOFs significantly enhanced the performance of PVDF-based polymer electrolytes.²⁸ Currently, research on MOF-derived, phase-pure CeO₂ nanoarchitectures as functional fillers remains scarce.

Herein, we report distinct grain-sized nanocrystalline CeO₂ fillers derived from Ce-MOFs through conventional sintering and alkaline treatment, which are incorporated into a PVDF-HFP matrix to fabricate high-performance composite polymer electrolytes (CPEs). Compared with conventional inorganic fillers, the Ce-MOF-derived nanocrystalline CeO₂ fillers demonstrate superior performance by leveraging their high specific surface area and strong polymer-filler interfacial interactions, which effectively promote lithium salt dissociation and significantly enhance Li⁺ transport kinetics while improving interfacial stability with lithium metal anodes.²⁶ The assembled full-cell configuration using the high-energy-density polyanionic cathode LiMn_0.6Fe_0.4PO₄ exhibited good cycling stability and rate capability, supporting the practical potential of these CPEs for advanced lithium metal battery applications.²⁹ This work provides fundamental insights into functional filler design and contributes to the development of next-generation high-performance solid-state lithium metal batteries.

2. Experimental section

2.1 Materials

All chemical reagents were commercially obtained from commercial suppliers, including cerium(III) nitrate hexahydrate (Ce(NO₃)₃·6H₂O, 99%, Macklin), 1,3,5-benzenetricarboxylic acid (H₃BTC, 99%, Macklin), lithium bis(trifluoromethanesulfonyl)imide (LiTFSI, 99%, Macklin), poly(vinylidene fluoride-co-hexafluoropropylene) (PVDF-HFP, M_w = 600 000, Aladdin), propylene carbonate (PC, 99.9%, Aladdin), N-methyl-2-pyrrolidone (NMP, 99.9%, Sigma-Aldrich), potassium hydroxide (KOH, analytical reagent grade, Macklin), and acetone (analytical reagent grade, Aladdin).

2.2 Preparation of nanocrystalline CeO₂

Ce-BTC was synthesized as follows: Cerium(III) nitrate hexahydrate [Ce(NO₃)₃·6H₂O] (3 mmol, 1.3 g) was dissolved in 1 mL of deionized water (designated as Solution A). Then, 3 mmol (0.63 g) of H₃BTC was dissolved in 40 mL of a water–ethanol mixed solution (1 : 1 volume ratio) under stirring to obtain Solution B. Solution A was combined with Solution B under continuous magnetic stirring (10 min, room temperature), resulting in the formation of a white precipitate. The solid product was collected via filtration, sequentially washed three times each with ethanol and distilled water, and finally air-dried.³⁰ The obtained Ce-BTC material was immersed in KOH solutions of different concentrations (0.01 M, 0.1 M, 0.5 M, and 1 M, respectively) and magnetically stirred for 1 minute, followed by repeated washing with copious amounts of deionized water. Ultrasmall nanocrystalline CeO₂ (UNC-CeO₂) was ultimately obtained by drying the material at 70 °C for 12 hours. Alternatively, nanocrystalline CeO₂ (NC-CeO₂) was prepared by heating the Ce-BTC material at 600 °C in an air atmosphere for 3 hours.

2.3 Preparation of polymer electrolytes

NC-CeO₂ and UNC-CeO₂ were individually dispersed in NMP/acetone (3 : 2 v/v). PVDF-HFP (1.50 g) was then introduced to the mixture under magnetic stirring at 65 °C for 6 h, followed by the addition of LiTFSI (0.75 g) with continued agitation for 3 h. The resultant slurry was subjected to degassing for 15 min at 60 °C to produce a bubble-free viscous solution. The precursor solution was uniformly doctor-bladed onto glass substrates, subjected to primary drying (48 h, fume hood), followed by post-drying in a vacuum oven (80 °C, 12 h) to yield freestanding polymer electrolyte membranes. The fabricated PE membranes were subsequently punched into circular discs of specified diameters for comprehensive characterization. For full-cell applications, the PE membranes were precisely cut into 19 mm diameter disks, then immersed in liquid electrolyte (1 M LiTFSI in propylene carbonate) within an argon-filled glove box for 12 h. After immersion, the electrolyte content was maintained at ∼5 wt%, indicating the formation of a quasi-solid-state electrolyte (QSE). Post-saturation, the membranes were meticulously wiped with lint-free cloths prior to storage or utilization as QPEs.

2.4 Characterization

The X-ray diffraction (XRD) pattern of CeO₂ was recorded using Cu Kα radiation (λ = 0.154 nm) over the 2θ range from 5° to 90° Morphological features with different dimensions were examined using a Zeiss Sigma 300 field-emission scanning electron microscope (SEM, Germany) coupled with energy-dispersive X-ray spectroscopy (EDS). Transmission electron microscopy (TEM) observations were performed on an FEI Talos F200S instrument. Fourier-transform infrared spectroscopy (FTIR) was performed using a Thermo Scientific Nicolet iS20 spectrometer. Thermal analyses were conducted by employing a TA Instruments Q2000 DSC under N₂ purge with 5 °C min⁻¹ heating from −80 °C to 150 °C, alongside a Netzsch TG 209 F3 Tarsus TGA (Germany) instrument operating at 20 °C min⁻¹ under nitrogen flow from ambient temperature to 600 °C. Brunauer–Emmett–Teller (BET) surface areas were quantified via a Micromeritics ASAP 2460 instrument following 8 h N₂-degassing at 120 °C. Synchrotron grazing-incidence wide-angle X-ray scattering (GIWAXS) employed beamline BL17B at SSRF (China) with monochromatic radiation (λ = 1.240 Å, α_i ≈ 0.4°). The sample-detector distance was calibrated against LaB₆ standards, with data acquired under continuous N₂ purge at 1 s per frame exposure.

2.5 Electrochemical measurements

The ionic conductivity (σ) of MOFs-based polymer electrolyte films was determined by electrochemical impedance spectroscopy (EIS) using a CHI660E workstation (Chenhua Instruments, China). Symmetric blocking cells were assembled by sandwiching disc-shaped film samples between stainless steel electrodes. Impedance spectra were recorded from 0.01 Hz to 10⁶ Hz at 10 mV amplitude under open-circuit conditions. Bulk resistance values extracted from Nyquist plots were substituted into formula (1.1) for σ calculation:²⁸

σ = L/RS (1.1)

where L (cm) signifies the thickness of the MOFs-based polymer electrolyte film, S (cm²) corresponds to the electrolyte–electrode contact area, and R (Ω) represents the bulk resistance derived from electrochemical impedance spectroscopy measurements. The Li/PE/Li symmetrical batteries were assembled in a glove box (Etelux Lab 2000), and the EIS test encompassed a frequency range from 10⁶ Hz to 0.01 Hz, with the amplitude set at 0.01 V using a CHI 660E electrochemistry station. The symmetrical cell underwent chronoamperometric polarization at 10 mV. Subsequent electrochemical impedance spectroscopy (EIS) was immediately performed without disassembling the cell. The lithium-ion transference number (t_Li⁺) of the MOF-based polymer electrolyte was then determined using formula (1.2) with pre-/post-polarization resistance values:²⁸

A DC polarization potential (ΔV) of 10 mV was applied throughout chronoamperometry. The initial (I₀) and steady-state (I_s) currents were recorded for the Li|PE|Li symmetric cell, with corresponding bulk impedances designated as R₀ (pre-polarization) and R_s (post-polarization). Here, R⁰_b and R^s_b represent the initial and steady-state electrolyte resistances, respectively.³¹ Crystallite dimensions were quantified via the Scherrer equation (1.3):

D = Kλ/β cos θ (1.3)

The electrochemical stability window was evaluated by linear sweep voltammetry (LSV) on a CHI 660E electrochemical workstation using Li|PE|SS cells. Scans were performed at 1 mV s⁻¹ between 2–6 V vs. Li⁺/Li. Critical current density (CCD) was identified at the onset of a sudden voltage drop (<1 V). Lithium dissolution–deposition cycling tests employed 60-minute cycles at current densities ranging from 0.01 to 1.5 mA cm⁻².

Full-cell Li|LiMn_0.6Fe_0.4PO₄ batteries were assembled in CR2032 coin cells using a hydraulic crimper (MSK-110) within an argon-filled glovebox and tested at 25 °C (BLH-300 incubator). The cathode slurry comprised LiMn_0.6Fe_0.4PO₄ (SJ-TECHNOL), Super P, and PVDF (8 : 1 : 1 mass ratio) in N-methyl-2-pyrrolidone (NMP), blade-coated onto Al foil. After vacuum drying (80 °C, 12 h), electrodes were punched into 14 mm discs with 2.3 mg cm⁻² active mass loading. Rate performance (1C = 170 mA g⁻¹) was evaluated on a LAND CT3001A system between 2.5–4.5 V.

3. Results and discussion

3.1 Characterization of CeO₂ functional fillers derived from Ce-BTC

Fig. 1a shows NC-CeO₂ derived from Ce-BTC calcined at 600 °C and UNC-CeO₂ formed after treatment with KOH solutions of varying concentrations. The CeO₂ sample calcined at 600 °C exhibited relatively sharp crystalline diffraction peaks. In contrast, the diffraction peak intensities of KOH-treated CeO₂ samples were significantly weakened, and with increasing KOH concentration, the characteristic diffraction peaks of NC-CeO₂ gradually broadened, indicating reduced crystallinity and increased amorphous content. These results demonstrate the pronounced influence of KOH treatment on the crystal structure of CeO₂.³² The crystallite sizes of NC-CeO₂, UNC-CeO₂-0.01, UNC-CeO₂-0.1, UNC-CeO₂-0.5, and UNC-CeO₂-1 were calculated using the Scherrer equation, yielding values of 75 nm, 4 nm, 3.5 nm, 2 nm, and 3 nm, respectively. These results demonstrate that the KOH-treated CeO₂ samples all consist of ultrafine nanocrystals. Structural characterizations of CeO₂ nanocrystals with different crystallite sizes are presented in Fig. 1b, c and S1. SEM images of the UNC-CeO₂-0.5 sample (treated with 0.5 M KOH solution) reveal clustered particle distributions. At higher magnifications, these CeO₂ clusters retain the rod-like structural features of the Ce-BTC precursor. Similarly, crystalline NC-CeO₂ (calcined at 600 °C) also displays analogous structural morphology, confirming that CeO₂ inherits the precursor's structural characteristics, even after high-temperature treatment. These findings suggest that the morphologies of both NC-CeO₂ samples are significantly influenced by their precursor structure.

Fig. 1 Characterization of the Ce-MOF-derived CeO₂: (a) XRD spectra; (b and c) SEM images of NC-CeO₂ and UNC-CeO₂-0.5; (d–f) TEM images of NC-CeO₂; and (g–i) TEM images of UNC-CeO₂-0.5.

Further evidence was provided by high-resolution TEM (Fig. 1d–i).³³ Lattice fringes are clearly visible at the edges of the nanorods, while diffuse diffraction rings without distinct spots indicate amorphous regions. The observed lattice fringes suggest a degree of structural ordering, consistent with the broadened XRD peaks typically associated with nanocrystalline materials. The selected-area electron diffraction (SAED) patterns exhibit ring-like diffraction, and the XRD patterns show broad peaks without sharp reflections, indicating the possible presence of an amorphous phase. As shown in Fig. 1d–f, NC-CeO₂ nanorods retain the rod-like morphology of the precursor, with well-defined diffraction spots and dense lattice fringes, confirming high crystallinity. The dominant experimental diffraction peak corresponds to the (111) plane of CeO₂, with measured d-spacings of 0.299 nm and 0.298 nm. In contrast, Fig. 1g–i reveal that UNC-CeO₂-0.5 lacks distinct diffraction spots and well-defined lattice fringes, exhibiting structural characteristics of ultrasmall nanocrystals, in agreement with XRD and crystallite size calculations. Moreover, nitrogen adsorption–desorption measurements were performed on the two nanocrystalline samples with distinct crystallite sizes (Fig. S2). The CeO₂-600 sample exhibited a specific surface area of 29 m² g⁻¹, while UNC-CeO₂-0.5 showed a significantly higher value of 58 m² g⁻¹. These results demonstrate that the KOH-treated, ultrasmall nanocrystals possess a markedly increased specific surface area.

3.2 Electrochemical properties of CeO₂-PVDF composite polymer electrolytes

The polymer electrolytes (PE-NC-CeO₂ and PE-UNC-CeO₂-0.5) were successfully fabricated via a solution casting method, employing synthesized NC-CeO₂ (nanocrystal) and UNC-CeO₂-0.5 (ultrasmall nanocrystal) as fillers within a poly(vinylidene fluoride-co-hexafluoropropylene) (PVDF-HFP) matrix.³⁴ As shown in Fig. 2a, the composite membranes exhibit excellent mechanical flexibility, with a thickness as low as 74 μm, and are capable of bending deformation and full recovery without structural compromise. Structural analyses of the membranes were systematically conducted. The internal structure of the film was characterized by SEM, as shown in Fig. 2b, c and S3, and the smooth and uniformly distributed mesh structure on the surface of the composite electrolyte film can be seen by SEM characterization at different scales. As shown in Fig. 2d–f, energy-dispersive X-ray spectroscopy (EDS) mappings further confirm the homogeneous dispersion of fillers, as evidenced by the uniform spatial distribution of O, S, and Ce elements across the membranes.

Fig. 2 (a) Optical images of PE-UNC-CeO₂-0.5 (left), PE-PVDF-HFP (right) electrolyte films and cross-sectional SEM images of PE-UNC-CeO₂-0.5 films (middle). (b and c) SEM images of PE-NC-CeO₂ and PE-UNC-CeO₂-0.5 film surfaces, (d–f) EDS images of PE-UNC-CeO₂-0.5 films, (g–i) 2D GIWAXS images of PE-PVDF-HFP, PE-NC-CeO₂, and PE-UNC-CeO₂-0.5 films.

GIWAXS serves as a critical technique for characterizing crystallographic structures, molecular orientation,³⁵ and thin-film microstructural morphology.³⁶ For the semi-crystalline polymer PVDF-HFP, its crystalline phase can be precisely analyzed through the positions and intensities of diffraction peaks. As illustrated in Fig. 2g–i, the 2D GIWAXS patterns of PE-PVDF-HFP, PE-NC-CeO₂, and PE-UNC-CeO₂-0.5 exhibit distinct diffraction rings. Notably, composite electrolyte membranes incorporating NC-CeO₂ and UNC-CeO₂ fillers demonstrate reduced diffraction ring widths and intensities, indicating that the addition of fillers—particularly UNC-CeO₂—diminishes the crystallinity of the PVDF-HFP matrix. As shown in Fig. S4, the 1D GIWAXS profiles further reveal a characteristic β-phase diffraction peak for the PVDF-HFP matrix at *q* ≈ 1.6 Å⁻¹ (corresponding to 2θ ≈ 20°), attributed to the (200) and (110) crystallographic planes (2θ ≈ 20.6° and 18.4°, respectively). A broad amorphous halo is concurrently observed, suggesting the coexistence of crystalline and amorphous domains. For PE-NC-CeO₂ and PE-UNC-CeO₂-0.5, the intensities of these characteristic peaks are significantly attenuated and slightly shifted to *q* ≈ 1.5 Å⁻¹. This behavior implies that the introduced fillers likely suppress the crystallographic orientation of the PVDF-HFP matrix via interfacial interactions, while simultaneously optimizing the microstructural homogeneity of the composite electrolytes.

After the preliminary experiments, we investigated different ratios of fillers and conducted electrochemical impedance spectroscopy tests, as shown in Fig. S5; the optimal filler addition of 1 wt% was determined, and in the subsequent experiments, this ratio was used to prepare the composite electrolyte films. Afterwards, the ionic conductivity of the Ce-MOF-derived CeO₂ polymer electrolyte was calculated using electrochemical impedance spectroscopy and eqn (1), as shown in Fig. 3a. The amorphous CeO₂ samples derived from KOH concentrations in the range of 1 M L⁻¹ to 0.01 M L⁻¹ were systematically tested as shown in Table S1: the conductivities of PE-NC-CeO₂-1%, PE-UNC-CeO₂-0.01-1%, PE-UNC-CeO₂-0.1-1%, PE-UNC-CeO₂-0.5-1%, and PE-UNC-CeO₂-1-1% were 3.68 × 10⁻⁴ S cm⁻¹, 2.25 × 10⁻⁴ S cm⁻¹, 3.17 × 10⁻⁴ S cm⁻¹, 5.04 × 10⁻⁴ S cm⁻¹, and 4.02 × 10⁻⁴ S cm⁻¹, respectively. The analysis showed that PE-UNC-CeO₂-0.5-1% reached the highest ionic conductivity among the tested materials. To investigate the role of nanocrystalline and ultrasmall nanocrystalline CeO₂ as fillers in the polymer electrolyte system, the ionic conductivity after immersion in 1 M LiTFSI–PC solution was further examined. As shown in Fig. S6, the composite electrolyte films with UNC-CeO₂ samples as fillers exhibited lower impedance values, with the ionic conductivity of the UNC-CeO₂ composite electrolyte films after immersion reaching 5.1 × 10⁻⁴ S cm⁻¹. In contrast, the ionic conductivity of the composite electrolyte films with NC-CeO₂ as fillers after immersion was improved to 4.14 × 10⁻⁴ S cm⁻¹, but it was still lower than that of UNC-CeO₂. This result suggests that UNC-CeO₂ can more effectively promote ionic transport in the gel polymer electrolyte system, thus enhancing the ionic conductivity of the electrolyte.

Fig. 3 (a) Impedance diagrams of polymer electrolytes, (b–d) AC impedance of Li/PE/Li cells before and after polarization of PE-PVDF-HFP, PE-NC-CeO₂, and PE-UNC-CeO₂-0.5; insets: polarization curves of Li/PE/Li cells.

The lithium-ion transference number (t_Li⁺) represents a crucial performance metric for polymer electrolytes. Traditional systems exhibit dual-ion conductivity, facilitating concurrent anion/cation migration in opposing directions under applied electric fields. This ion flux establishes detrimental concentration gradients that compromise battery performance. Enhancing t_Li⁺ is therefore essential to suppress concentration polarization and enable uniform lithium deposition. Fig. 3b–d show chronoamperometry and AC impedance spectra of Li|PE|Li symmetric cells before and after polarization. The Li⁺ transference number (t_Li⁺) for MOF-based polymer electrolytes was quantified through combined electrochemical impedance spectroscopy and constant-voltage polarization, derived from formula (1.2). This approach provides a comprehensive understanding of the electrolyte's ability to facilitate Li⁺ mobility. Notably, PE-UNC-CeO₂-0.5 exhibited a Li⁺ transference number of 0.54, surpassing those of PE-NC-CeO₂ (0.47) and PE-PVDF-HFP (0.33). This disparity suggests that UNC-CeO₂ modification significantly enhances the ion transport selectivity of the polymer electrolyte. Specifically, Lewis acid sites on the surface of UNC-CeO₂ may interact with functional groups on the PVDF-HFP chains. Such interactions likely weaken the binding strength between PVDF-HFP and Li⁺, thereby facilitating rapid Li⁺ migration along the PVDF-HFP chains.

The electrochemical window serves as a critical parameter for evaluating the stability of quasi-solid-state electrolytes.³⁷ The electrochemical stability windows of composite electrolytes were evaluated via linear sweep voltammetry (LSV) in asymmetric SS|PE|Li cells.³⁸ This configuration enabled the quantification of oxidative stability for PVDF-HFP, NC-CeO₂ and UNC-CeO₂-0.5 membranes over 2–6 V vs. Li⁺/Li (Fig. 4a). The results demonstrate that the PE-PVDF-HFP electrolyte membrane exhibited an initial current surge exceeding 4.0 V, with pronounced current escalation upon further voltage elevation, indicating an electrochemical window limited to 4.2 V. In contrast, the PE-UNC-CeO₂-0.5 and PE-NC-CeO₂ systems displayed extended electrochemical windows, reaching 4.6 V and 4.7 V, respectively. These findings confirm that Ce-MOFs-derived CeO₂ fillers effectively enhance the high-voltage endurance of polymer electrolytes, providing essential compatibility with high-voltage cathode materials. To further evaluate the practical electrochemical stability, constant-current cycling (CICC) tests with incrementally increased upper cutoff voltages were conducted. This approach enables the monitoring of capacity retention and voltage profiles during cycling, with voltage curves revealing information regarding side reactions and their onset potentials. As shown in Fig. S7, the CICC results for full cells demonstrate that PE-PVDF-HFP exhibited less stability compared to PE-NC-CeO₂ and PE-UNC-CeO₂-0.5 electrolytes, which maintained more stable voltage profiles during degradation. The combined LSV and CICC test results collectively validate that the incorporation of NC-CeO₂ and UNC-CeO₂-0.5 fillers significantly improves the electrochemical window of composite electrolytes.

Fig. 4 (a) LSV curves of PE-PVDF-LiTFSI, PE-NC-CeO₂ and PE-UNC-CeO₂-0.5 electrolytes; (b) FTIR spectra, (c) TGA curves, and (d) DSC curves.

Fourier-transform infrared (FTIR) spectroscopy characterized structural features of PVDF-HFP-based electrolytes with NC-CeO₂ and UNC-CeO₂-0.5 fillers, probing their influence on LiTFSI dissociation and elucidating UNC-CeO₂-0.5's role in Li⁺ transport enhancement.³⁹Fig. 4b displays the FTIR spectra of PE-PVDF-HFP, PE-NC-CeO₂, and PE-UNC-CeO₂-0.5 in the range of 4000–500 cm⁻¹. For the PE-CeO₂-0.5 M film, the C=O peak (1785 cm⁻¹) indicated participation in Li⁺⋯C=O coordination. The infrared spectrum of PE-UNC-CeO₂-0.5 exhibited asymmetric (1785 cm⁻¹) and symmetric (1484 cm⁻¹) stretching vibration peaks of the carbonyl group (C=O), confirming that UNC-CeO₂-0.5 retains its unique structural integrity within the composite. Notably, compared to PE-PVDF-HFP, these peaks displayed significant shifts and intensity reduction, clearly indicating interactions between PE-PVDF-HFP and PE-UNC-CeO₂-0.5. Furthermore, the introduction of NC-CeO₂ and UNC-CeO₂-0.5 induced minor shifts and slight intensity attenuation in the CH₂CH₂O vibration peak (834 cm⁻¹) of PE-PVDF-HFP, further demonstrating interfacial interactions between PE-PVDF-HFP and the CeO₂ fillers. A pronounced displacement was observed for the –CF₃ asymmetric and symmetric stretching vibration peaks, originally located at 1173 cm⁻¹ and 1228 cm⁻¹ in PE-PVDF-HFP, which shifted to 1117 cm⁻¹ and 1170 cm⁻¹, respectively, in PE-UNC-CeO₂-0.5. Additionally, the –CH₂– peak in PE-UNC-CeO₂-0.5 exhibited a subtle shift relative to that in PE-PVDF-HFP. These spectral alterations highlight synergistic interactions among UNC-CeO₂-0.5, TFSI⁻ anions, and the PVDF-HFP matrix.^3,32

Thermal stability plays a pivotal role in polymer electrolytes.⁴⁰ The thermal degradation behavior of PE-PVDF-HFP, PE-NC-CeO₂, and PE-UNC-CeO₂-0.5 was systematically investigated through thermogravimetric analysis (TGA).⁴¹ As shown in Fig. 4c, although PE-NC-CeO₂ and PE-UNC-CeO₂-0.5 underwent irreversible decomposition between 300 °C and 450 °C, their thermal decomposition profiles exhibited superior stability compared to PE-PVDF-HFP. The enhanced thermal resistance can be attributed to the inherent thermal stability of NC-CeO₂ and UNC-CeO₂-0.5 fillers. Specifically, the residual mass of PE-UNC-CeO₂-0.5 (23.95 wt%) and PE-NC-CeO₂ (22.69 wt%) was significantly higher than that of PE-PVDF-HFP (21.72 wt%) at the final decomposition stage. This remarkable thermal stability ensures the safe application of PE-NC-CeO₂ and PE-UNC-CeO₂-0.5 in lithium metal batteries, even under elevated temperature conditions.

It is widely acknowledged that the amorphous phase of the PVDF-HFP matrix plays a crucial role in ionic conductivity.⁴² The phase transition behavior of PVDF-HFP was further investigated through differential scanning calorimetry (DSC),⁴³ as shown in Fig. 4d. The DSC thermograms reveal that the incorporation of NC-CeO₂ and UNC-CeO₂-0.5 fillers reduced the glass transition temperature (T_g) while increasing the melting temperature (T_m) of the PVDF-HFP matrix. Notably, the UNC-CeO₂-0.5 filler enhanced Li⁺ diffusion kinetics, yielding superior ionic conductivity. Enthalpy changes (ΔH) and crystallinity degrees for PVDF-HFP, NC-CeO₂ and UNC-CeO₂-0.5 composites, determined via DSC analysis, are compiled in Table S2. The percentage crystallinity (χ_c) of PVDF-HFP in the electrolyte was determined using eqn (1.4):

where ΔH* = 104.5 J g⁻¹ (pure PVDF-HFP melting enthalpy), ΔH = measured enthalpy change, and χ = PVDF-HFP mass fraction. The incorporation of 1 wt% NC-CeO₂ lowered χ_c from 19.9% to 15.8%, whereas 1 wt% UNC-CeO₂-0.5 further reduced it to 12.9%.⁴⁴ This crystallinity suppression arises from increased amorphous domains and enhanced polymer chain mobility induced by CeO₂ fillers.

3.3 Symmetric and full cell measurements

Fig. 5a–c present the cycling performance of Li/QSE/Li symmetric cells under varying current densities. Fig. 5a displays the voltage–time profiles of Li/Li symmetric cells tested at a low current density of 0.1 mA cm⁻² under ambient conditions. Notably, the Li/QSE-NC-CeO₂/Li and Li/QSE-UNC-CeO₂-0.5/Li symmetric cells exhibited lower polarization voltages and enhanced cycling stability compared to the Li/QSE-PVDF-HFP/Li counterpart. As shown in Fig. 5c, when the current density was increased to 0.2 mA cm⁻², the Li/QSE-UNC-CeO₂-0.5/Li cell demonstrated superior cycling stability with minimal voltage fluctuation over prolonged cycling. In contrast, the Li/QSE-NC-CeO₂/Li cell exhibited a more pronounced polarization growth under identical testing conditions. This discrepancy suggests that the UNC-CeO₂-modified polymer electrolyte (QSE-UNC-CeO₂-0.5) effectively suppresses side reactions and lithium dendrite growth at the lithium metal anode through optimized interfacial compatibility and ion transport pathways, thereby achieving more stable electrochemical cycling behavior.

Fig. 5 Performance of the Li/QSE/Li symmetric cells: (a–c) long cycle performance of the Li/QSE/Li symmetric cells, with the inset in (c) showing the locally enlarged time-voltage profile. (d) CCD curves of Li/QSE/Li symmetric cells, with the inset in (d) displaying the locally enlarged time-voltage profile.

Fig. 5b illustrates the voltage–time profiles of the Li/QSE-UNC-CeO₂-0.5/Li cell across different current densities, further confirming its robust cycling performance. These findings clearly demonstrate the exceptional capability of UNC-CeO₂-modified polymer electrolytes in inhibiting lithium dendrite proliferation, offering critical insights for designing high-safety lithium metal batteries. The critical current density (CCD), a key parameter for evaluating interfacial stability between polymer electrolytes and lithium anodes and quantifying lithium dendrite suppression efficiency, was further investigated to assess the long-term cycling stability of Ce-MOFs-derived CeO₂-based polymer electrolytes. The CCD test involved a stepwise increase in current density from 0.01 mA cm⁻² to 1.5 mA cm⁻², with detailed recording of overpotential variations in Li/QSE-NC-CeO₂/Li and Li/QSE-UNC-CeO₂-0.5/Li cells. As depicted in Fig. 5d, the Li/QSE-UNC-CeO₂-0.5/Li cell exhibited smoother polarization evolution, achieving a CCD of 0.8 mA cm⁻², whereas the Li/QSE-NC-CeO₂/Li cell reached a lower CCD of ∼0.5 mA cm⁻². These results further validate the pivotal role of UNC-CeO₂ fillers in enhancing the interfacial stability of solid-state electrolytes.

To evaluate the variation in battery resistance, in situ electrochemical impedance spectroscopy (EIS) was employed to monitor the charge/discharge cycling process, followed by analysis of the distribution of relaxation times (DRT) after cycling, as shown in Fig. 6. The mid-frequency region corresponds to the formation of the solid electrolyte interphase (SEI). It was observed that QSE-UNC-CeO₂-0.5 exhibited lower impedance during cycling compared to QSE-PVDF-HFP and QSE-NC-CeO₂, suggesting its superior capability in facilitating the formation of a stable SEI. Furthermore, the diffusion resistances of both QSE-UNC-CeO₂-0.5 and QSE-NC-CeO₂ during cycling were significantly lower than that of QSE-PVDF-HFP, indicating faster ion diffusion kinetics. To further quantify the lithium-ion transport resistance across the interface, galvanostatic intermittent titration technique (GITT) measurements were conducted (Fig. S8). The results indicated that QSE-UNC-CeO₂-0.5 possesses a higher lithium-ion diffusion coefficient than QSE-PVDF-HFP and QSE-NC-CeO₂, demonstrating enhanced ion migration kinetics; this is highly consistent with the in situ EIS data.

Fig. 6 (a–c) Distribution of relaxation times (DRT) of QSE-PVDF-HFP, QSE-NC-CeO₂ and QSE-UNC-CeO₂-0.5, with insets showing the locally enlarged mid-frequency region. (d–f) 2D intensity color map DRT profiles of QSE-PVDF-HFP, QSE-NC-CeO₂ and QSE-UNC-CeO₂-0.5 cycled between 2.5 and 4.5 V.

LiMn_xFe_1−xPO₄ (LMFP), a doped derivative of LiFePO₄ (LFP) with an olivine structure, retains the inherent safety and structural stability of its parent material while achieving enhanced operational voltage and energy density through manganese (Mn) substitution.⁴⁵ To validate the applicability of nanocrystalline and ultrasmall nanocrystalline CeO₂-filled polymer electrolytes in high-voltage cathode systems, quasi-solid-state Li|LiMn_0.6Fe_0.4PO₄ cells were assembled, and their electrochemical performance was systematically evaluated. Cycling performance and rate capability tests demonstrated the excellent compatibility between the amorphous CeO₂-based polymer electrolyte and the LMFP cathode material, as shown in Fig. 7a. The cell employing the UNC-CeO₂-0.5 electrolyte maintained a discharge capacity of approximately 94 mAh g⁻¹ after 100 cycles at a current density of 0.5C, corresponding to a high capacity retention of 94.6%. In contrast, the Li/QSE-NC-CeO₂/LMFP cell exhibited poor cycling performance, delivering an initial capacity of only 49.78 mAh g⁻¹ at 0.5C; the corresponding charge/discharge profiles are presented in Fig. 7d and e. Furthermore, under a higher current density of 1C, the UNC-CeO₂-0.5 electrolyte enabled stable cycling for over 500 cycles. Remarkably, after 500 cycles, the discharge capacity remained around 80 mAh g⁻¹, representing a capacity retention of 99%, as depicted in Fig. 7b. The rate capability of the Li/QSE-UNC-CeO₂-0.5/LMFP cell is presented in Fig. 7c. The cell delivered reversible capacities at various current densities (0.2C to 1C) as follows: 104.76 mAh g⁻¹ at 0.2C, 89.07 mAh g⁻¹ at 0.5C, and 41.85 mAh g⁻¹ at 1C. Upon reverting the current density from 1C back to 0.2C, the discharge capacity recovered to 104.51 mAh g⁻¹, with the corresponding charge/discharge profiles shown in Fig. 7f. This significant recovery highlights the outstanding rate capability and electrochemical performance of the QSE-UNC-CeO₂-0.5 electrolyte.

Fig. 7 (a) Discharge capacities of Li/QSE-UNC-CeO₂-0.5/LMFP and Li/QSE-NC-CeO₂/LMFP cells at 0.5C rate. (b) Charge/discharge capacities of the Li/QSE-UNC-CeO₂-0.5/LMFP cell at 1C rate. (c) Rate capability of the Li/QSE-UNC-CeO₂-0.5/LMFP cell at various current rates. (d and e) Charge/discharge voltage profiles of Li/QSE-UNC-CeO₂-0.5/LMFP and Li/QSE-NC-CeO₂/LMFP cells at 0.5C rate. (f) Rate-dependent charge/discharge profiles of the Li/QSE-UNC-CeO₂-0.5/LMFP cell. (g) Schematic of the reaction mechanism in the Li/QSE/LMFP cell.

As illustrated in Fig. 7g, the ultrasmall nanocrystalline CeO₂ possesses a high specific surface area and abundant active sites. These features facilitate lithium salt dissociation, enhance compatibility with PVDF, and modify its crystalline state. Consequently, interfacial compatibility with both lithium metal and the LMFP cathode is improved, enabling superior performance under the demanding conditions of reversible Li⁺ (de)intercalation reactions in the Li|LMFP system. These comparative results underscore the superior compatibility of ultrasmall nanocrystal CeO₂-modified polymer electrolytes with LiMn_0.6Fe_0.4PO₄ cathodes, effectively supporting stable cycling in all-solid-state batteries. The findings further confirm the advantages of ultrasmall nanocrystalline CeO₂ fillers in ensuring interfacial stability and long-term cyclability. This study provides experimental evidence for the application of filler-enhanced polymer electrolytes in high-energy cathode systems, highlighting their potential for use in all-solid-state batteries.

4. Conclusions

In summary, this study demonstrates an effective strategy for developing high-performance composite polymer electrolytes by incorporating Ce-MOF-derived ultrasmall nanocrystalline CeO₂ fillers into a PVDF-HFP matrix, successfully addressing the critical challenges of low ionic conductivity and lithium dendrite growth in solid-state batteries. These designed CeO₂ nanofillers, featuring a high specific surface area and abundant active sites, significantly enhance lithium salt dissociation through Lewis acid–base interactions and effectively reduce the crystallinity of PVDF electrolyte. The constructed polymer composite electrolyte achieved a superior ionic conductivity (5.04 × 10⁻⁴ S cm⁻¹) at room temperature and an expanded electrochemical stability window up to 4.6 V. The Li/QSE/LMFP full cell maintained 99% capacity retention after 500 cycles at a 1C rate. The optimized electrolyte system exhibited effective dendrite suppression capability and interfacial stability, providing guidance for the development of safe, high-energy-density lithium metal batteries. Future research should focus on the precise control of the filler morphology, optimization of scalable fabrication processes, and development of compatible cathode materials.

Conflicts of interest

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

Data availability

The data supporting this article have been included as part of the supplementary information (SI). Supplementary information is available. See DOI: https://doi.org/10.1039/d6qi00187d.

Acknowledgements

This work was supported by the Key Research and Development “LingYan” Program of Zhejiang Province (No. 2024C01189), the Funds of the Natural Science Foundation of Hangzhou (No. 2025SZRJJ2021) and the Fundamental Research Funds for the Provincial Universities of Zhejiang (No. 2024YW104). We thank the staff of the BL17B beamline (https://cstr.cn/31129.02.NFPS) at the National Facility for Protein Science in Shanghai (NFPS, https://cstr.cn/31129.02.NFPS), Shanghai Advanced Research Institute, Chinese Academy of Sciences, for their technical support in GWAXS data collection and analysis. We extend gratitude to Mr Deyou Chen from Scientific Compass (https://www.shiyanjia.com) for providing invaluable assistance with the HRTEM analysis. The authors also would like to thank Hao Wang from SCI-GO (https://www.sci-go.com) for the N₂ adsorption–desorption analysis.

References

1. S. Zou, Y. Yang, J. Wang, X. Zhou, X. Wan, M. Zhu and J. Liu, In situ polymerization of solid-state polymer electrolytes for lithium metal batteries: a review, Energy Environ. Sci., 2024, 17, 4426–4460.
2. G. Zeng, Y. Tian, M. D. Khan, K. V. Mejía-Centeno, S. Wang, N. Revaprasadu, A. Cabot and Q. Sun, Non–Flammable Electrolytes for Safe Lithium–, Sodium–, and Potassium–Ion Batteries, Electron, 2025, 3(4), e70021.
3. S. Liu, W. Liu, D. Ba, Y. Zhao, Y. Ye, Y. Li and J. Liu, Filler–Integrated Composite Polymer Electrolyte for Solid–State Lithium Batteries, Adv. Mater., 2022, 35, 211043.
4. J. Li, X. Chen, S. Muhammad, S. Roy, H. Huang, C. Yu, Z. Ullah, Z. Wang, Y. Zhang, K. Wang and B. Guo, Development of solid polymer electrolytes for solid-state lithium battery applications, Mater. Today Energy, 2024, 43, 101574.
5. Y. Fu, Z. Gu, Q. Gan and Y.-W. Mai, A review on the ionic conductivity and mechanical properties of composite polymer electrolytes (CPEs) for lithium batteries: Insights from the perspective of polymer/filler composites, Mater. Sci. Eng., R, 2024, 160, 100815.
6. Y. Dong, Y. He, B. Yuan, X. Ding, S. Zhong, J. Feng, Y. Han, Z. Liu, L. Xu, K. Feng, J. Han, H. Cheng, C. Lv and W. He, Multi–Level Regulation of Electrostatic Microenvironment With Anion Vacancies for Low–Lithium–Gradient Polymer Electrolyte, Electron, 2025, 3(3), e70010.
7. M. Zhou, K. Cui, T.-S. Wang, Z. Luo, L. Chen, Y. Zheng, B. Li, B. Shi, J. Liu, J.-J. Shao, G. Zhou, S. Yang and Y.-B. He, Bimetal Fluorides with Adjustable Vacancy Concentration Reinforcing Ion Transport in Poly(ethylene oxide) Electrolyte, ACS Nano, 2024, 18, 26986–26996.
8. Z. Luo, W. Li, C. Guo, Y. Song, M. Zhou, Y. Shi, J. Xu, L. Li, B. Shi, Q. Ouyang, J.-J. Shao and G. Zhou, Two-dimensional silica enhanced solid polymer electrolyte for lithium metal batteries, Particuology, 2024, 85, 146–154.
9. X. Liu, H. Jia and H. Li, Flame-retarding quasi-solid polymer electrolytes for high-safety lithium metal batteries, Energy Storage Mater., 2024, 67, 103263.
10. M. Lee, J. Na, S. Oh, J. Choi, K. Eun, J. Chun, J. H. Kim, I. Bae, J. Jin, T. K. Lee, S.-K. Jung, S. H. Joo and S. J. Kang, Uniaxially Aligned Relaxor Ferroelectric polymer electrolyte for high-performance solid-state lithium batteries, Energy Storage Mater., 2025, 83, 104707.
11. N. J. Kong, M. S. Kim, J. H. Park, J. Kim, J. Jin, H.-W. Lee and S. J. Kang, Promoting homogeneous lithiation of silicon anodes via the application of bifunctional PEDOT:PSS/PEG composite binders, Energy Storage Mater., 2024, 64, 103074.
12. Y.-H. Liu, F. Wang, H.-B. Wang, C.-Y. Kong, G.-X. Wang, X.-M. Liu and Y. Liu, Shining light on fillers uniform dispersion of PVDF/garnet composite electrolytes for high-performance solid-state Li batteries: fundamentals, progress and perspectives, Rare Met., 2025, 44, 5957–5979.
13. X. Yang, J. Liu, N. Pei, Z. Chen, R. Li, L. Fu, P. Zhang and J. Zhao, The Critical Role of Fillers in Composite Polymer Electrolytes for Lithium Battery, Nano-Micro Lett., 2023, 15, 74.
14. E. Carena, L. Mezzomo, N. Vallana, N. Ceribelli, G. D. Liberto, S. Mostoni, C. Ferrara, M. Mauri, R. Lorenzi, L. Giordano, R. Ruffo and P. Mustarelli, PVDF–HFP Based, Quasi–Solid Nanocomposite Electrolytes for Lithium Metal Batteries, Small, 2024, 20, 2311805.
15. Q. Yang, M. Liao, P. Ren, D. Li, N. Deng, W. Kang and G. Cai, Dual functional composite solid electrolyte constructed by double-layer sulfonated polyethersulfone (SPES)/poly(vinylidene fluoride-co-hexafluoropropylene) (PVDF-HFP) nanofiber membrane in high-performance all-solid-state lithium metal batteries, Chem. Eng. J., 2024, 495, 153448.
16. X. Zhang, S. Cheng, C. Fu, G. Yin, L. Wang, Y. Wu and H. Huo, Advancements and Challenges in Organic–Inorganic Composite Solid Electrolytes for All-Solid-State Lithium Batteries, Nano-Micro Lett., 2025, 17, 2.
17. H. Junoh, N. Awang, H. S. Zakria, N. A. S. Zainuddin, N. A. H. M. Nordin, N. S. Suhaimin, T. Enoki, T. Uno and M. Kubo, Advancements in Polyethylene Oxide (PEO)–Active Filler Composite Polymer Electrolytes for Lithium-Ion Batteries: A Comprehensive Review and Prospects, Materials, 2024, 17, 4344.
18. Z. Zhang, Y. Xu, D. Xiong, J. Yu, J. Cai, Y. Huang and Z. Yang, Improving sulfur transformation of lean electrolyte lithium–sulfur battery using nickel nanoparticles encapsulated in N–doped carbon nanotubes, Electron, 2024, 2(1), e19.
19. S. Kim, W.-G. Lim, H. Jung, Y. C. Jeong, C.-Y. Park, S. B. Yang, C. H. Lee, D. Wang, K. Sohn, J. W. Han and J. Lee, Protective catalytic layer powering activity and stability of electrocatalyst for high-energy lithium-sulfur pouch cell, Nat. Commun., 2025, 16(1), 1649.
20. Z. Zhang, X. Wang, X. Li, J. Zhao, G. Liu, W. Yu, X. Dong and J. Wang, Review on composite solid electrolytes for solid-state lithium-ion batteries, Mater. Today Sustainability, 2023, 21, 100316.
21. Y. Guo, E. Zhao, X. Zhao and S. Liu, One dimensional CeO₂ nanorods/poly(ethylene oxide) solid composite electrolyte for all-solid-state lithium-ion batteries, J. Rare Earths, 2024, 42, 570–577.
22. L. Gao, N. Wu, N. Deng, Z. Li, J. Li, Y. Che, B. Cheng, W. Kang, R. Liu and Y. Li, Optimized CeO₂ Nanowires with Rich Surface Oxygen Vacancies Enable Fast Li–Ion Conduction in Composite Polymer Electrolytes, Energy Environ. Mater., 2023, 6, e12272.
23. Z.-H. Zuo, J.-K. Hu, X.-L. Wang, S.-J. Yang, W.-Q. Mai, Y.-H. Zhu, Z. Liao, J. Liu, H. Yuan and J.-Q. Huang, Oxygen defect-mediated Li-ion transport for long-cycle solid-state lithium metal batteries, Chin. Chem. Lett., 2025, 37, 110851.
24. B. Li, C. Wang, R. Yu, J. Han, S. Jiang, C. Zhang and S. He, Recent progress on metal–organic framework/polymer composite electrolytes for solid-state lithium metal batteries: ion transport regulation and interface engineering, Energy Environ. Sci., 2024, 17, 1854–1884.
25. H. Zhu, S. Li, L. Peng, W. Zhong, Q. Wu, S. Cheng and J. Xie, Review of MOF-guided ion transport for lithium metal battery electrolytes, Nano Energy, 2024, 125, 109571.
26. S. Duan, L. Qian, Y. Zheng, Y. Zhu, X. Liu, L. Dong, W. Yan and J. Zhang, Mechanisms of the Accelerated Li(+) Conduction in MOF-Based Solid-State Polymer Electrolytes for All-Solid-State Lithium Metal Batteries, Adv. Mater., 2024, 36, e2314120.
27. H. Wang, S. Duan, Y. Zheng, L. Qian, C. Liao, L. Dong, H. Guo, C. Ma, W. Yan and J. Zhang, Solid-state electrolytes based on metal-organic frameworks for enabling high-performance lithium-metal batteries: Fundamentals, progress, and perspectives, eTransportation, 2024, 20, 100311.
28. T. Jiang, X. Meng, X. Hu, A. Tang, Z. Ruan, Q. Kang, L. Yan, Y. Zhao, N. Yu, B. Liu, M. Fan, C. Wan and T. Ma, Defective cerium-based metal-organic framework nanorod- reinforcing polymer electrolytes for lithium metal batteries, J. Power Sources, 2025, 628, 235914.
29. X. Yang, S. Wang, H. Li, J. Tseng, Z. Wu, S. Indris, H. Ehrenberg, X. Guo and W. Hua, Unveiling the correlation between structural alterations and enhanced high–voltage cyclability in Na–deficient P3−type layered cathode materials via Li incorporation, Electron, 2024, 2(1), e18.
30. S. Maiti, A. Pramanik and S. Mahanty, Extraordinarily high pseudocapacitance of metal organic framework derived nanostructured cerium oxide, Chem. Commun., 2014, 50, 11717–11720.
31. Q. Zeng, J. Wang, X. Li, Y. Ouyang, W. He, D. Li, S. Guo, Y. Xiao, H. Deng, W. Gong, Q. Zhang and S. Huang, Cross-Linked Chains of Metal–Organic Framework Afford Continuous Ion Transport in Solid Batteries, ACS Energy Lett., 2021, 6, 2434–2441.
32. H. Chen, D. Adekoya, L. Hencz, J. Ma, S. Chen, C. Yan, H. Zhao, G. Cui and S. Zhang, Stable Seamless Interfaces and Rapid Ionic Conductivity of Ca–CeO₂/LiTFSI/PEO Composite Electrolyte for High–Rate and High–Voltage All–Solid–State Battery, Adv. Energy Mater., 2020, 10, 2000049.
33. L. Wang, Y. Chen, L. Zeng, X. Huang and W. Lv, PVDF-HFP nanofiber quasi-solid electrolyte for fast charging and dendrites-free lithium batteries, J. Power Sources, 2025, 640, 236779.
34. W. Xiong, T. Huang, Y. Feng, X. Ye, X. Li, J. Liang, S. Ye, X. Ren, Y. Li, Q. Zhang and J. Liu, Rapid ionic conductivity of ternary composite electrolytes for superior solid-state batteries with high-rate performance and long cycle life operated at room temperature, J. Mater. Chem. A, 2021, 9, 18338–18348.
35. Z. Zhu, L. Lu, C. Li, Q. Xiao, T. Wu, J. Tang, Y. Gu, K. Bao, Y. Zhang, L. Jiang, Y. Liu, W. Zhang, S. Zhou and W. Qin, GIWAXS experimental methods at the NFPS-BL17B beamline at Shanghai Synchrotron Radiation Facility, J. Synchrotron Radiat., 2024, 31, 968–978.
36. M. Sanviti, D. E. Martínez-Tong, E. Rebollar, T. A. Ezquerra and M. C. García-Gutiérrez, Crystallization and phase separation in PEDOT:PSS/PEO blend thin films: Influence on mechanical and electrical properties at the nanoscale, Polymer, 2022, 262, 125475.
37. J. Guo, Z. Zhang, X. Lu, Y. Zhang, J. Wang and Z. Wei, Flame-Retardant Gel Electrolytes Reinforced by PVDF-HFP/Ceramic Nanofiber Mat for Lithium Metal Batteries, ACS Appl. Mater. Interfaces, 2025, 17, 22841–22849.
38. R. Gopal, L. Wu, Y. Lee, J. Guo and P. Bai, Transient Polarization and Dendrite Initiation Dynamics in Ceramic Electrolytes, ACS Energy Lett., 2023, 8, 2141–2149.
39. P. Zhai, Z. Yang, Y. Wei, X. Guo and Y. Gong, Two–Dimensional Fluorinated Graphene Reinforced Solid Polymer Electrolytes for High–Performance Solid–State Lithium Batteries, Adv. Energy Mater., 2022, 12, 2022967.
40. V. C. Nguyen, T. Dam, H.-B. Na and C.-J. Park, Plasticized Composite Electrolyte with Al₂O₃ Nanofiller-Reinforced PVDF-HFP for Solid-State Lithium–Metal Batteries, ACS Appl. Energy Mater., 2025, 8, 3120–3131.
41. J. Chen, H. Wang, Y. Bai, P. Pang, Z. Zheng, H. Xu, Y. Zhao, K. Jiang and G. Zhu, Dual-MOFs-cage constructed multistage-channel PVDF-HFP quasi-solid electrolytes for lithium metal battery, J. Power Sources, 2024, 614, 234975.
42. S. Wang, C. Li, Y. Ma, H. Zhang, X. Shi, L. Zhang and D. Song, Regulating Crystalline Phase/Plane of Polymer Electrolyte for Rapid Lithium Ion Transfer, Angew. Chem., Int. Ed. Engl., 2024, 64(9), e202420698.
43. H. Wang, J. Chen, P. Pang, Y. Bai, Z. Zheng, T. Huang, K. Jiang, Y. Zhao, G. Zhu and H. Xu, Multistage channel PVDF-HFP pregnant ZIF-8@SiO₂ quasi-solid electrolyte for lithium-metal batteries, J. Power Sources, 2024, 599, 234167.
44. X. Wu, Y. Zheng, W. Li, Y. Liu, Y. Zhang, Y. Li and C. Li, Solid electrolytes reinforced by infinite coordination polymer nano-network for dendrite-free lithium metal batteries, Energy Storage Mater., 2021, 41, 436–447.
45. L. Wen, Y. Shi, L.-W. Ye, C. Wang and F. Li, Progress of lithium manganese iron phosphate in blended cathode materials, Mater. Sci. Eng., R, 2025, 164, 100952.

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