Ba_0.6Sr_0.4TiO₃ ferroelectric filler-reinforced poly(vinylidene fluoride) polymer electrolytes for dendrite-free solid-state Li metal batteries

Abstract

Polyvinylidene fluoride (PVDF)-based electrolytes have attracted significant attention for their potential use in solid-state lithium batteries (SSLBs) due to their superior electrochemical performance and safety. However, their low ionic conductivity and uneven lithium deposition hinder the further application of PVDF-based electrolytes. Herein, this work focuses on incorporating Ba_0.6Sr_0.4TiO₃ (BST) ferroelectric ceramics into PVDF to form composite solid-state electrolytes (CSEs). The BST ferroelectric ceramics can create an intrinsic electric field that facilitates lithium-ion transport and enables uniform Li deposition. In addition, benefiting from the high dielectric constant of BST and dipoles generated from the asymmetric structure, PVDF–BST CSEs achieve a high ionic conductivity (1.79 × 10⁻⁴ S cm⁻¹) due to more free lithium ions, a wide electrochemical window of 4.8 V (vs. Li/Li⁺) and a high Li⁺ transference number (0.37). The assembled Li|PVDF–BST|Li symmetrical cells can steadily cycle for 1100 h at 0.1 mA cm⁻² at 25 °C. The assembled Li|PVDF–BST|LiFePO₄ cells show long-term cycling stability with a capacity retention of 85.6% after 100 cycles at 0.5C and a capacity retention of 81.4% after 200 cycles at 1C. This work provides a new strategy for improving the performance of the PVDF-based electrolytes by incorporating ferroelectric ceramics.

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1. Introduction

With the rapid advancement of technology, the demands for higher energy density and improved safety in energy storage devices are steadily increasing in modern society. Compared to conventional liquid lithium batteries, solid-state lithium batteries (SSLBs), are regarded as a key technology in the next-generation energy revolution, owing to their advantages, such as enhanced safety, higher energy density, and superior thermal stability.^1,2 The performance of solid-state electrolytes (SSEs), which are the core component of SSLBs, plays a crucial role in determining the overall efficiency and safety of these batteries.³

Significant efforts have been dedicated to the development and optimization of SSEs. SSEs are primarily divided into inorganic solid-state electrolytes (ISEs) and solid-state polymer electrolytes (SPEs).⁴ ISEs, including oxide electrolytes (garnet, NASICON, and perovskite), sulfide electrolytes (Li₁₀GeP₂S₁₂, Li₆PS₅X, and Li₂S–P₂S₅) and halide electrolytes (Li–M–X, where M = metal element and X = F, Cl, Br, I), are known for their excellent lithium conductivity, mechanical performance and safety.^5–8 However, their application are hindered by challenges in assembly processes and poor interfacial contact with electrodes. SPEs, such as PEO (polyethylene oxide), PAN (polyacrylonitrile), PMMA (poly(methyl methacrylate)), and PVDF (polyvinylidene fluoride), are considered promising due to their superior flexibility, ease of fabrication, and low cost.⁹ In particular, PVDF is regarded as a prospective candidate for SSEs because of its excellent electrochemical stability, high dielectric constant, and compatibility with the lithium metal anode. Nevertheless, PVDF SSEs are limited by low ionic conductivity, high crystallinity, and insufficient active sites.

To address these challenges, incorporating functional inert fillers into PVDF polymers to form composite solid-state electrolytes (CSEs) has been proven to be an effective strategy.^10,11 The inorganic fillers in CSEs can be categorized into MOFs (ZiF-8), 2D materials (hexagonal boron nitride [h-BN] and montmorillonite [MMT]), ferroelectric materials (BaTiO₃), etc. Among these, perovskite-type (ABO₃) ferroelectric ceramic fillers, a specific type of functional inert fillers, exhibit a high dielectric constant and promote the dissociation of lithium salts and suppress the formation of space charge layers.¹² This process generates more free lithium ions and reduces interfacial impedance between the electrode and the electrolyte.^13,14 Additionally, the displacement of the B-site cation at the center of the unit cell relative to the O atoms results in an asymmetric structure, creating numerous dipoles with opposing electronegativities on the surface of the fillers, which facilitates lithium-ion transport and improves the overall performance of the battery.^12,15

The Ba_0.6Sr_0.4TiO₃ (BST) ferroelectric material, a continuous solid solution of barium titanate (BaTiO₃) and strontium titanate (SrTiO₃), features a high dielectric constant and low dielectric loss, making it an excellent candidate for applications in dielectric energy storage, capacitors, and piezoelectric devices.¹⁶ With its high dielectric constant and intrinsic electric field induced by spontaneous polarization, BST fillers have significant potential for dissociating more free lithium ions, regulating Li⁺ flux at interfaces, and promoting uniform lithium deposition (Fig. 1). To our knowledge, functionally inert BST fillers were introduced into solid-state electrolytes for the first time, with the aim of improving the electrochemical stability and cycling performance of PVDF-based CSEs. Flexible PVDF–BST CSEs were fabricated through a simple solution-casting method, achieving a high ionic conductivity of 1.79 × 10⁻⁴ S cm⁻¹ and a lithium-ion transference number (t_Li⁺) of 0.37 at room temperature. The assembled lithium metal symmetric cell demonstrated stable cycling for over 1100 hours at a current density of 0.1 mA cm⁻². Furthermore, the Li/PVDF–BST CPEs/LiFePO₄ full cells maintained a capacity retention of up to 81.4% after 200 cycles at a 1C rate.

Fig. 1 Mechanism diagram of BST ferroelectric ceramics in PVDF-based CSE.

2. Experimental

All chemical reagents were bought from the Aladdin Company without further purification.

2.1 Synthesis of Ba_0.6Sr_0.4TiO₃ nanoparticles

Ba_0.6Sr_0.4TiO₃ (BST) ceramic nanoparticles were synthesized via a high-temperature solid-state reaction method.¹⁷ The raw materials, including BaCO₃ (AR, ≥99%), SrCO₃ (AR, ≥99%), and TiO₂ (AR, ≥99%) were precisely weighed according to the stoichiometric ratio and then mixed in a ball mill, using ethanol as the milling medium, for 12 hours. The mixture was subsequently dried in an oven and pre-sintered at 1050 °C for 3 hours. Following pre-sintering, the powder was ground and subjected to another 12-hour ball-milling process, followed by drying. The resulting powder was calcined at 1350 °C for 2 hours. Finally, the calcined powder was milled and dried again to obtain Ba_0.6Sr_0.4TiO₃ nanoparticles.

2.2 Preparation of PVDF-based CSEs

The CPE membranes were prepared using a simple solution casting method, as shown from Fig. S1.† Initially, 0.3 g of PVDF (M_w = 534 000) and 0.2 g of LiTFSI (≥99.9%) were dissolved in 5 mL of N,N-dimethylformamide (DMF) (≥99.9%) solution under stirring at 60 °C for 12 hours. Subsequently, BST nanoparticles, with varying mass fractions relative to the mass of PVDF, were added to the transparent stirred solution. The mixture was further stirred at 60 °C for 6 hours to ensure homogeneity. The solution was cast onto a horizontally placed glass slide and dried in a vacuum oven at 80 °C for 24 hours to obtain the composite electrolyte membrane. The membrane was cut into 18 mm diameter discs and transferred to a glovebox for static drying for 24 hours before use. For the PVDF electrolyte blank sample, no BST ceramic nanoparticles were added, while the other preparation steps remained the same.

2.3 Preparation of the LiFePO₄ cathode

The LiFePO₄ (LFP) cathode was prepared by mixing LiFePO₄, Super P, and PVDF binder in a weight ratio of 8 : 1 : 1 in N-methyl-2-pyrrolidone (NMP) (≥99.5%). The resulting homogeneous slurry was coated onto aluminum foil using an adjustable doctor blade, with the coating thickness controlled at approximately 16 μm. The coated foil was then vacuum dried at 80 °C for 12 hours to form the LiFePO₄ cathode. The prepared cathode sheet was cut into 10 mm diameter discs and transferred to a glovebox for storage. The active material loading of the cathode sheet was approximately 2.5 mg cm⁻².

2.4 Battery assembly

CR2025 coin cells were assembled in an argon-filled glovebox with both H₂O and O₂ levels being less than 1 ppm. The prepared LiFePO₄ was used as the cathode, and lithium metal sheets was employed as the anode. Solid-state electrolyte membranes made from PVDF and PVDF–4BST were used as the blank and modified samples, respectively.

2.5 Characterization

The crystal structure and phase composition were investigated at room temperature by X-ray diffraction (XRD, Rigaku SmartLab, Rigaku Corporation, Japan) with Cu-K_α radiation (λ = 1.5406 Å). The scan rate was set to 5° min⁻¹, and the scan range was from 10° to 80°. The morphology and elemental distribution of the CPEs were characterized using a scanning electron microscope (SEM, JSM-7610FPlus, JEOL Ltd, Japan) equipped with an energy-dispersive spectroscopy (EDS) detector. Thermogravimetric analysis (TGA, NETZSCH STA 449 F3, NETZSCH Scientific Instruments, Germany) was performed under an argon atmosphere from room temperature to 600 °C. Differential scanning calorimetry (DSC, TSC-2500, TA Instruments, USA) was conducted within a temperature range of −50 to 180 °C.

The ionic conductivity of the CPEs was measured using an electrochemical workstation (CorrTest, CS350M, China). Two stainless steel discs (diameter: 15.8 mm) were used as blocking electrodes in a symmetric cell. Electrochemical impedance spectroscopy (EIS) was performed in the frequency range of 10⁻¹ to 10⁵ Hz with an applied perturbation voltage of 10 mV, over a temperature range of 20–80 °C. The impedance spectra were analyzed using CS Analysis software. The ionic conductivity was calculated using eqn (1):¹⁸

where L is the CPE thickness, S is the contact area between stainless steel sheets and CPEs, and R is the real part of the impedance of CPEs.

The activation energy for lithium-ion conductivity (E_a) was calculated from the temperature-dependent ionic conductivity data using eqn (2):¹⁹

where A, R, and T are the pre-exponential factor, the molar gas constant, and Kelvin temperature, respectively.

The Li|CPE|Li symmetric cells were assembled to evaluate the lithium-ion transference number (t_Li⁺) using a combination of EIS and DC polarization methods. EIS measurements were performed before and after polarization to determine the impedance values of the symmetric cell under both conditions. The frequency range was set from 10⁻¹ to 10⁵ Hz, with an AC perturbation voltage of 10 mV. The DC polarization voltage and test duration were set to 10 mV and 4000 seconds, respectively. The lithium-ion transference number (t_Li⁺) was calculated using eqn (3):¹⁹

where I₀ and I_s are the current in the initial state and steady state, respectively. R₀ and R_s are the interface resistance of the initial state and steady state, respectively.

The dielectric properties of the CPEs were analyzed using an impedance analyzer (4294a, USA). Copper electrodes were deposited onto both opposing surfaces of the CPE using a high-vacuum resistive evaporation coating system. Capacitance–frequency spectra were collected at 25 °C over a frequency range of 10² to 10⁶ Hz. The dielectric constant of the electrolyte membrane was calculated using eqn (4):

where C is the capacitance, d is the thickness of CPEs, ε₀ is the vacuum permittivity, and r is the radius of the deposited copper electrode.

Cycling tests were performed at 25 °C for the Li|CPE|Li symmetric cells and the Li|CPE|LiFePO₄ full cells using a battery testing system (Land CT3002A, China). The symmetric cells underwent lithium plating/stripping cycles at different current densities, with each step lasting 1 hour. For the full cells, charge/discharge tests were conducted within a voltage range of 2.5 to 4.0 V.

3. Results and discussion

3.1 Phase and structure of PVDF-based CSEs

The prepared BST ceramic particles were characterized using XRD and SEM. As shown in Fig. 2a and S2,† the BST ceramic particles exhibit a typical perovskite structure. The diffraction peaks align well with the standard XRD pattern of Ba_0.6Sr_0.4TiO₃ (PDF#34-0411), confirming the successful synthesis of pure BST ceramic particles without the formation of secondary phases via the high-temperature solid-state method. The SEM image (Fig. S3†) reveals that the particle size is predominantly around 300 nm. Additional X-ray diffraction (XRD) analysis was performed on the prepared electrolyte membranes, as shown in Fig. 2a. Clear diffraction peaks were observed at 18.50° and 20.04°, which correspond closely to the diffraction peaks of γ-phase PVDF.²⁰ Both the β-phase and γ-phase, which exhibit polarity and similar conformations, are generally considered electroactive phases that facilitate the migration of Li⁺ ions.^20–22 The incorporation of LiTFSI and BST ceramic particles reduces the crystallinity of the PVDF polymer.²³ This decrease in crystallinity, along with the corresponding increase in amorphous phase content, enhances the performance of the PVDF-based electrolyte.

Fig. 2 (a) XRD patterns of BST particles, PVDF powder, PVDF SPE and PVDF–BST CSE. (b) SEM image of the PVDF–BST CSE with corresponding EDS elemental mapping for Ba, Sr and F. (c) FTIR spectra of PVDF–BST CPE and PVDF SPE at 2000–600 cm⁻¹. (d) The TGA curves, (e) DSC curves and (f) stress–strain curves of PVDF–BST CPE and PVDF SPE.

As shown in Fig. 2b and S4–S6,† the SEM image and EDS mapping confirm the uniform distribution of LiTFSI and BST ceramic particles in PVDF–BST CSE and PVDF SPE. LiTFSI and BST particles are uniformly attached to the surface of PVDF particles, and the obtained electrolyte membrane has a thickness of approximately 90 μm. FTIR spectra of PVDF SPE and PVDF–4BST CSE are shown in Fig. 2c. The majority of these characteristic absorption peaks match well, indicating that these absorption bands are attributable to PVDF and LiTFSI. The absorption peaks at 836 cm⁻¹, 877 cm⁻¹, and 1230 cm⁻¹ correspond to γ-phase PVDF.²⁴ The peak at 837 cm⁻¹ reflects the CF₂ symmetric stretching vibration and C–C asymmetric stretching vibration, while the peak at 1230 cm⁻¹ corresponds to CF₂ asymmetric stretching vibration, bending vibration, and C–F out-of-plane deformation vibration.^25–28 The absorption peaks at 1056 cm⁻¹, 1176 cm⁻¹, and 1351 cm⁻¹ belong to β-phase PVDF.²⁴ The absorption peak at 1331 cm⁻¹ corresponds to the characteristic chemical bond SO₂⁻ in the TFSI⁻ group of LiTFSI.²⁹ The presence of residual DMF solvent, beneficial for the electrochemical performance of the electrolyte, is confirmed by the peaks at 1391 cm⁻¹ (CH₃ bond) and 1659 cm⁻¹ (C=O vibration).^29–31 Additionally, the absorption peak at 673 cm⁻¹ corresponds to the chemical bond from [Li(DMF)_x]⁺,²³ indicating the coordination interaction between DMF molecules and Li⁺ in electrolyte membranes. Due to the interaction between PVDF chains and [Li(DMF)_x]⁺, Li⁺ could be transported among the interaction sites between [Li(DMF)_x]⁺ and the PVDF chains by the dissociation and recoordination occurring in [Li(DMF)_x]⁺ and C–F groups, further enhancing the diffusion of Li⁺ in the electrolyte.^30,32

To verify the thermal stability of PVDF SPE and PVDF–4BST CSE, thermogravimetric analysis (TGA) was performed, as shown in Fig. 2d. The mass loss below 310 °C for both PVDF SPE and PVDF–4BST CSE can be attributed to the evaporation of residual water and DMF solvent. The weight loss between 310 °C and 390 °C is primarily associated with the decomposition of LiTFSI, while the mass loss between 390 °C and 600 °C corresponds to the decomposition of PVDF. It is observed that the degradation temperature of PVDF–4BST CPE is approximately 10 °C higher than that of PVDF SPE. This may be attributed to the well-dispersed BST ceramics, which can scatter and obstruct the transmission of heat, thereby further delaying the decomposition of LiTFSI.^24,33 And the residual weight of the PVDF–4BST CPE is greater than that of PVDF SPE, indicating that PVDF–4BST CPE exhibits higher thermal stability.³⁴Fig. 2e displays the DSC curves of PVDF SPE and PVDF–4BST CSE. The degree of crystallinity (X_c) of the electrolytes can be quantified by DSC. The X_c value is calculated using the following equation:³⁵

X_c = ΔH_m/ΔH⁰_m

where ΔH⁰_m is the fusion enthalpy of full crystallized PVDF (104.7J g⁻¹) and ΔH_m is the fusion enthalpy of the electrolyte, obtained from the area under the peak at the melting temperature T_m. The X_c of PVDF SPE is approximately 34.4%, while that of PVDF–4BST CSE is 33.3%. The decrease in X_c means an increase in the amorphous regions and enhanced segmental movement of the polymer chains within the PVDF matrix, leading to improvement of the transport of Li⁺ and an increase in ionic conductivity.^13,36

The mechanical strength of CPE could influence the electrochemical stability of lithium metal batteries. As shown in Fig. 2f, the tensile fracture strength of PVDF–4BST CSE (3.02 MPa) is higher than that of PVDF SPE (2.60 MPa) due to the incorporation of BST ceramic particles. The improved mechanical properties are attributed to the intrinsic high strength of BST ceramic particles and their strong adhesion to the PVDF polymer matrix to form continuous and tensile polymer/inorganic structures.^10,31 This enhancement in mechanical strength helps to better resist the risk of lithium dendrite penetration.

3.2 Electrochemical performances

To optimize the content of BST ceramic particles in the PVDF matrix, electrochemical impedance spectroscopy (EIS) measurements were performed at 25 °C with varying ceramic contents, as shown in Fig. 3a. The bulk resistance was minimized, when the BST content reached 4% by weight. Further increases in BST content led to higher impedance, likely due to the aggregation and uneven dispersion of BST ceramics, which resulted in a decrease in ionic conductivity. Therefore, the optimal BST ceramic content was determined to be 4% by weight. Impedance measurements of PVDF–BST CPE and PVDF SPE were conducted over the temperature range of 20–80 °C, as shown in Fig. 3b and S7, S8.† The ionic conductivities of PVDF–4BST and PVDF electrolyte membranes at 25 °C were calculated to be 1.79 × 10⁻⁴ S cm⁻¹ and 7.38 × 10⁻⁵ S cm⁻¹, respectively, using eqn (1). The activation energies (E_a) were calculated using the Arrhenius equation (eqn (2)), and found to be 0.269 eV for PVDF–4BST CPE and 0.276 eV for PVDF SPE (Fig. 3c). The lower activation energy for PVDF–4BST CPE indicates an improved lithium-ion migration rate compared to PVDF SPE. This indicates that Li⁺ has a lower migration barrier in the PVDF–4BST CPE, which results in reduced energy losses during battery operation and enhanced cycling stability.

Fig. 3 (a) EIS plots of CPEs with varying BST ceramic particle contents. (b) EIS plots of PVDF–4BST CPE from 20 °C to 80 °C. (c) Activation energies, (d) LSV, (e and f) DC polarization test and related AC impedance spectra (inset) of PVDF–4BST CPE and PVDF SPE. (g) Dielectric constant and (h and i) Raman spectrum of PVDF SPE and PVDF–4BST CPE.

The electrochemical windows of these membranes are shown in Fig. 3d. The PVDF SPE exhibited an electrochemical stability window of approximately 3.7 V. In contrast, the PVDF–4BST CPE demonstrated an expanded electrochemical stability window of 4.8 V, suggesting enhanced voltage stability due to the incorporation of BST ceramics. Interestingly, DMF decomposition in PVDF SPE was observed at 3.3 V, whereas PVDF–4BST CPE effectively suppressed DMF decomposition, further enhancing its electrochemical stability. This broader electrochemical window of PVDF–4BST CPE is probably due to the strong interactions that stabilize residual DMF and suppress its electrochemical oxidation.³⁷ This enhanced electrochemical stability enables the electrolyte to effectively match with a high-voltage cathode, thereby contributing to a high energy density. As shown in Fig. 3e and f, the lithium-ion transference numbers (t_Li⁺) of PVDF SPE and PVDF–4BST CSE were measured using eqn (3). The t_Li⁺ of PVDF SPE was found to be approximately 0.18, while that of PVDF–4BST CSE increased significantly to 0.37. This enhancement is attributed to the electrostatic interactions between BST dipoles and TFSI⁻ anions.³⁸ These interactions hinder the migration of TFSI⁻ ions, leading to an increased t_Li⁺ for the PVDF–4BST CSE.

The dielectric constants of PVDF SPE and PVDF–4BST CSE were also measured in Fig. 3g. At 25 °C and 100 Hz, the dielectric constant of PVDF SPE is 6.78. In comparison, PVDF–4BST CSE exhibited an increased dielectric constant of 9.04, attributed to the high dielectric constant of the BST ferroelectric ceramic material. The increased dielectric constant can intensify the charge separation of –CH₂CF₂ dipoles and induce positional deviation of Ti atoms relative to O atoms in BST. This effect facilitates the dissociation of LiTFSI, resulting in a greater number of free lithium ions.^39,40

Fig. 3h and i show the Raman spectra of PVDF SPE and PVDF–4BST CSE. The analysis reveals that in PVDF SPE, the percentage of dissociated free TFSI⁻ ions and undissociated TFSI⁻ ions is 59.16% and 40.84%, respectively. In contrast, for the PVDF–4BST CSE, these percentages are 82.06% and 17.94%, respectively. This indicates a nearly 22% increase in the content of free TFSI⁻ ions, suggesting that the enhanced dielectric constant of PVDF–4BST CPE promotes greater dissociation of LiTFSI, which is consistent with the observed improvements in the lithium-ion transference number.

3.3 Performance and characterization of Li|PVDF-based CSEs|Li symmetric cells

To evaluate the stability of PVDF SPE and PVDF–4BST CSE with a lithium metal anode, symmetric lithium metal cells were assembled to test the critical current density (CCD). As shown in Fig. 4a, the CCD of PVDF SPE is approximately 0.3 mA cm⁻². Beyond this current density, the polarization voltage increases significantly, indicating severe interfacial degradation and increased impedance. In contrast, the CCD of PVDF–4BST CSE can reach 1.3 mA cm⁻². Gradient current density tests were conducted on both electrolyte membranes in Fig. 4b. At 0.3 mA cm⁻², the PVDF SPE caused internal short-circuiting after 10 hours of operation, whereas the PVDF–4BST CSE operated stably at 0.5 mA cm⁻². Furthermore, when the current density was reduced from 0.5 mA cm⁻² to 0.1 mA cm⁻², the cell with PVDF–4BST CSE remained stable with a low polarization voltage. This suggested that the PVDF–4BST CSE can withstand higher current densities and operate stably even at elevated current densities.

Fig. 4 (a) Critical current density (CCD) of the Li|PVDF–4BST|Li and Li|PVDF|Li cells. (b) Gradient current density (GCD) of the Li|PVDF–4BST|Li and Li|PVDF|Li cells. (c) Galvanostatic cycles of Li|PVDF–4BST|Li and Li|PVDF|Li cells at a current density of 0.1 mA cm⁻¹. (d) C 1s and (f) F 1s XPS spectra of cycled Li anodes from the Li|PVDF–4BST|Li cell. (e) C 1s and (g) F 1s XPS spectra of cycled Li anodes from the Li|PVDF|Li cell. (h and k) SEM images of the uncycled lithium. (i and l) SEM images of the cycled lithium electrode surface from the Li|PVDF–4BST|Li symmetrical cell at a current density of 0.1 mA cm⁻¹. (j and m) SEM images of the cycled lithium electrode surface from the Li|PVDF|Li symmetrical cell at a current density of 0.1 mA cm⁻¹.

With an increased number of free Li⁺ ions and multiple dipolar channels for Li⁺ transport, PVDF–4BST CPE can effectively inhibit the growth of lithium dendrites.³⁹ In contrast, the failure of PVDF SPE could be attributed to its lower t_Li⁺, which causes the accumulation of TFSI⁻ anions at the interface, forming a space charge layer and increasing interfacial impedance.^13,14,41 By contrast, BST ferroelectric ceramics, with higher dielectric constants, effectively suppress space charge layer formation, facilitating a well-defined Li⁺ transport pathway.^12,39

To evaluate the long-term lithium stability of PVDF SPE and PVDF–4BST CSE, cycling tests were conducted under conditions of 0.1 mA cm⁻² and 0.1 mA h cm⁻² (Fig. 4c). During cycling, the polarization voltage of PVDF SPE gradually increased, and the cell was able to cycle for 650 hours before short-circuiting. In contrast, PVDF–4BST CSE demonstrated stable performance for over 1100 hours, with the polarization voltage remaining stable at approximately 40 mV. This superior lithium stability of PVDF–4BST CSE is attributed to the internal electric field generated by BST ferroelectric ceramics, which promotes the uniform deposition of lithium ions and enhances the cycling performance of the cell. The lithium stability of PVDF–4BST CSE and PVDF SPE were further evaluated under conditions of 0.2 mA cm⁻² and 0.2 mA h cm⁻². Similarly, PVDF–4BST CSE exhibited better cycling stability and a lower polarization voltage. Additionally, when assessed at a higher current density of 0.5 mA cm⁻², PVDF–4BST CSE operated stably for 150 hours. These results confirm that the BST ceramic fillers can greatly boost the interfacial stability and kinetics between the Li anode and PVDF–4BST CSE. To further investigate the solid electrolyte interphase (SEI) of the cycled lithium symmetric cells, X-ray photoelectron spectroscopy (XPS) was performed. As shown in Fig. 4d–g and S12, S13,† the C–F and Li–F peaks in the spectra of F 1s were located at 688.5 and 684.7 eV, respectively. The interface of the cycled Li//PVDF–4BST CPE exhibited a higher amount of LiF compared to the cycled Li//PVDF SPE interface. This enhancement is attributed to the presence of BST near the lithium metal surface, which provides a strong dipole moment that accelerates the degradation kinetics of the C–F bond in LiTFSI, thereby forming a highly stable LiF-rich SEI film.^4,39 The abundant LiF promotes rapid Li⁺ transfer and suppresses the growth of lithium dendrites.^42,43 The SEM characterization of lithium electrodes after cycling, as shown in Fig. 4h–m, revealed that the surface morphology of the lithium electrode with PVDF–4BST CSE was similar to that of uncycled lithium. In contrast, the lithium electrode with PVDF SPE exhibited a rougher surface and severe lithium dendrite formation.

3.4 Performance of Li|PVDF-based CSEs|LiFePO₄ full cells

The assembled Li|CPE|LiFePO₄ cells were used to evaluate the compatibility of PVDF-based CSE with a LiFePO₄ electrode. As shown in Fig. 5a, the first-cycle charge capacities of PVDF SPE at various rates were 147.4, 142.4, 125.3, 105, 82.5, 58.2 and 128.3 mA h g⁻¹ at 0.1, 0.2, 0.5, 0.8, 1, 2, and 0.2C (1C = 170 mA g⁻¹), respectively. In comparison, the corresponding charge capacities for PVDF–4BST CSE were 169.6, 166.2, 157.1, 146.9, 137.9, 121 and 154.2 mA h g⁻¹. At each rate, the PVDF–4BST CSE demonstrated significantly higher initial charge capacities than the PVDF SPEs. Long-cycle tests were performed on PVDF–4BST CSE and PVDF SPE at 0.5C and 1C. As shown in Fig. 5d–e, after 100 cycles at 0.5C, both PVDF–4BST CSE and PVDF SPE exhibited high coulombic efficiencies. PVDF–4BST CSE had a higher initial capacity of 148.6 mA h g⁻¹, compared to only 112.4 mA h g⁻¹ for PVDF SPE. After 100 cycles, PVDF–4BST CSE retained a capacity of 127.6 mA h g⁻¹, while the capacity of PVDF SPEs decreased to 89.3 mA h g⁻¹. At higher cycling rates of 1C, PVDF–4BST CSE had an initial specific capacity of 133.3 mA h g⁻¹, whereas the PVDF SPE showed only 108.6 mA h g⁻¹. After 200 cycles at 1C, the PVDF–4BST CSE retained a specific capacity of 104.6 mA h g⁻¹, compared to only 77.2 mA h g⁻¹ for PVDF SPE. These results demonstrate the superior electrochemical performance of PVDF–4BST CSE compared to PVDF SPE. The improved performance can be attributed to the critical role of BST fillers in regulating ion transport at the interface. These BST fillers enhance the uniform distribution of lithium ions and reduce interfacial resistance, thereby significantly improving the cycling stability and rate capability of Li|CPEs|LiFePO₄ cells.

Fig. 5 (a) Rate capacities of LiFePO₄//Li cells with PVDF–4BST CPE and PVDF SPE. (b and c) Charge/discharge curves of Li|PVDF|LiFePO₄ and Li|PVDF–4BST|LiFePO₄ cells at different rates. (d and e) Long-term cycling performances of Li|PVDF|LiFePO₄ and Li|PVDF–4BST|LiFePO₄ cells at 0.5C and 1C.

To better understand the mechanism of BST in PVDF electrolyte, a schematic diagram of the Li plating behavior of PVDF–4BST CSE and PVDF SPE is shown in Fig. 6. In the pure PVDF SPE, the uneven lithium deposition aggravates the growth of lithium dendrites, forming a relatively fragile SEI film. The lithium dendrites penetrate the electrolyte membrane, causing a short circuit in the battery. In contrast, the large number of dipoles in BST ceramic particles spontaneously align in PVDF–4BST CPE, allowing lithium ions to deposit uniformly under the influence of the local electric field, forming a uniform and dense SEI film, enabling the battery to operate stably for a long time.

Fig. 6 The schematic diagram of the Li plating behavior of PVDF–4BST CSE and PVDF SPE.

4. Conclusion

In summary, we successfully synthesized enhanced PVDF-based electrolytes using Ba_0.6Sr_0.4TiO₃ ferroelectric ceramics. The resulting optimized PVDF–4BST CSE exhibit a wide electrochemical window of 4.8 V (vs. Li/Li⁺), an improved Li⁺ transference number (up to 0.37) and a high ionic conductivity of 1.79 × 10⁻⁴ S cm⁻¹ at 25 °C. Li symmetric cells using PVDF–4BST demonstrate stable cycling performance for over 1100 hours at 0.1 mA cm⁻² at 25 °C. When assembled with LiFePO₄, the full battery shows an impressive 85.6% capacity retention after 100 cycles and a high coulombic efficiency of 98% at 0.5C. This work presents a strategy for improving the performance the PVDF-based electrolytes by incorporating BST ceramics, which can not only prevent dendrite formation but also promote lithium-ion transport, thus advancing the practical applications of ferroelectric materials in energy storage technologies.

Data availability

Data will be made available on request.

Author contributions

Chen Yang: investigation, data curation, writing – original draw. Hongjian Zhang: methodology. Mingtao Zhu: data curation, investigation. Hao Wu: conceptualization, supervision, writing – review & editing. Qiushi Wang: writing – review & editing. Yong Zhang: conceptualization, supervision, resources, writing – review & editing.

Conflicts of interest

There are no conflicts to declare.

Acknowledgements

The research was supported by the National Natural Science Foundation of China (grant no. 52002301); the Natural Science Foundation of Hubei Province (grant nos. 2020CFB308 and 2023AFB675); the Start-up Funding of Wuhan University of Technology (grant no. 40120490); the State Key Laboratory of Advanced Fiber Materials (KF2501); the National Natural Science Foundation of China (Grant No. 22305028 and 52102211) and the Dalian Minzu University Doctoral Program (120164 and 110230).

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Footnote

† Electronic supplementary information (ESI) available. See DOI: https://doi.org/10.1039/d5se00285k

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