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. 2026 Sep 29. Online ahead of print. doi: 10.1039/d6sc05344k

Interfacial weak interaction bridges in PVDF-HFP electrolytes via nano-ZrP enable fast charging of solid-state lithium batteries

Yanting Ma a, Songsheng Zheng a,✉, Jianyang Wu b,✉, Kai Zou c, Shujian Bai a, Zhaolin Wang a, Xiang Li d
PMCID: PMC13636258  PMID: 42835623

Abstract

Poly(vinylidene fluoride-co-hexafluoropropylene) (PVDF-HFP)-based polymer electrolytes have garnered significant interest for solid-state lithium metal batteries (SSLMBs) owing to their high ionic conductivity, robust mechanical strength, and superior thermal stability. However, undesired phase separation within the polymer matrix during film formation impedes rapid ion migration. To address this, a cost-effective zirconium phosphate (ZrP) filler featuring a two-dimensional layered architecture is incorporated into the PVDF-HFP matrix. The 2D ZrP nanosheets not only modulate polymer particle size and bridge adjacent spherulitic domains to construct continuous ion transport channels, but also leverage their abundant surface chemistry, namely strong DMF adsorption and –OH⋯TFSI− interactions, to reconfigure the Li+ solvation sheath and attenuate the Li+–TFSI− coulombic attraction. This synergistic combination of topological and chemical functionality accelerates ion migration, as evidenced by a high ionic conductivity of 1.4 × 10−3 S cm−1. Furthermore, a compact LiF-rich interphase is formed on both the anode and the cathode, effectively suppressing interfacial side reactions during high-rate cycling. Consequently, benefiting from these synergistic effects, Li‖Li symmetric cells demonstrate stable cycling at 0.3 mA cm−2, and Li‖LiFePO4 full cells deliver a high capacity of ∼96.7 mAh g−1 even after over 700 cycles at 5C. This work presents a promising filler-mediated strategy to accelerate ion migration for high-performance SSLMBs.

ZrP bridges PVDF-HFP spherulites to create fast Li+ migration pathways, while H-bonding with TFSI− weakens Li+–TFSI− attraction, enabling fast charging in solid-state Li metal batteries.graphic file with name d6sc05344k-ga.webp

1. Introduction

The escalating demand for high-energy-density energy storage systems has necessitated a critical transition from conventional lithium-ion batteries to solid-state lithium metal batteries.1,2 Lithium metal, characterized by its high theoretical specific capacity (3860 mAh g−1) and low reduction potential (−3.04 V), is regarded as a promising anode material to overcome current energy density bottlenecks.3,4 However, in conventional liquid electrolyte-based batteries, lithium metal anodes suffer from uncontrolled dendrite growth. This phenomenon not only leads to low coulombic efficiency and continuous electrolyte consumption but also poses significant safety hazards that hinder practical commercialization.5,6 Consequently, solid-state electrolytes (SSEs) have emerged as compelling alternatives, offering inherent non-flammability, superior mechanical strength, and effective physical suppression of lithium dendrites, thereby paving the way for safe and long-life lithium metal batteries. Among various SSEs, solid polymer electrolytes (SPEs) have attracted significant attention due to their lightweight nature, flexibility, excellent processability, and favorable interfacial contact with electrodes.7–11

For decades, poly(ethylene oxide) (PEO)-based electrolytes have been extensively investigated owing to their superior solvating ability for lithium salts, which facilitates ionic conduction. However, their high crystallinity results in low ionic conductivity at room temperature (typically 10−7 to 10−6 S cm−1), severely restricting their practical applications.12–14 In contrast, PVDF-HFP-based electrolytes have emerged as promising candidates due to their amorphous regions and polar functional groups, and these attributes combine high mechanical strength and excellent thermal stability with enhanced room-temperature ionic conductivity.15,16 Nevertheless, the PVDF-HFP matrix exhibits limited capability to dissociate lithium salts, necessitating the incorporation of highly polar solvents, such as N,N-dimethylformamide (DMF), to facilitate salt dissociation and achieve desirable ionic conductivity.17 Consequently, these solvent molecules inevitably participate in Li+ solvation and migrate to the electrode surfaces, where they undergo continuous reductive decomposition. This process induces repeated reconstruction of an unstable solid electrolyte interphase (SEI), accompanied by detrimental interfacial side reactions, increased impedance, and exacerbated dendrite growth, ultimately leading to rapid battery failure.18,19

To address these challenges, diverse strategies have been investigated, encompassing the optimization of lithium salt composition,20 the grafting of functional groups21 and the modulation of cycling protocols.22 Furthermore, the incorporation of inorganic fillers to fabricate composite solid electrolytes has emerged as a viable and potent approach. Notably, integrating ion-conductive active inorganic fillers, such as Li1+xAlxTi2−x(PO4)3 (LATP),16 Li1+xAlxGe2−x(PO4)3 (LAGP), and Li0.33La0.56TiO3−x (LLTO),23 enhances both ionic conductivity and interfacial stability within polymer electrolytes. However, the practical deployment of these active fillers is hindered by detrimental interfacial side reactions with the polymer matrix and involve complex, cost-prohibitive synthesis processes. Conversely, passive fillers present a compelling alternative by circumventing these limitations while bolstering performance. For example, double hydroxide nanosheets24 and metal–organic frameworks (MOFs)25 immobilize TFSI− anions to facilitate LiTFSI dissociation; palygorskite nano-wires promote ion transport via high DMF content;26 and 2D fluorinated graphene refines polymer grain size, augments mechanical strength, and improves interfacial ion conduction.13 Nevertheless, the ionic conductivity of most passive filler-based electrolytes typically remains below 10−3 S cm−1, with rate capability largely restricted to <2C, underscoring the necessity for novel approaches to satisfy the rigorous demands of high-power-density solid-state batteries.

Zeolite molecular sieves are renowned for their size-selective adsorption capabilities, derived from their well-defined pore structures (3 Å). By physically trapping DMF molecules, these sieves attenuate Li+–DMF interactions and promote an anion rich solvation sheath, thereby suppressing interfacial side reactions and enhancing cycling stability.19 Inspired by this mechanism, we investigate nano-zirconium phosphate (ZrP), a two-dimensional layered material featuring an interlayer spacing of ∼9 Å. Distinct from the three-dimensional framework of zeolites, the layered structure of ZrP offers abundant accessible Lewis-acid sites (OH) on its surface, which engage in strong interactions with polar molecules such as DMF and TFSI−. These structural attributes render ZrP a promising filler for modulating the solvation environment and facilitating fast ion conduction in polymer electrolytes. Accordingly, ZrP is incorporated into the PVDF-HFP matrix as a cost-effective, multifunctional filler. Structurally, it refines polymer particle size and bridges PVDF-HFP spherulitic domains to establish continuous interfacial ion transport channels, thereby reducing interfacial impedance. Chemically, it exhibits superior adsorption affinity for residual DMF compared to the polymer matrix, sequestering DMF from the Li+ solvation sheath and reconstructing the solvation structure. Concurrently, weak interactions between the –OH groups on ZrP and TFSI− are expected to attenuate the Li+–TFSI− electrostatic interaction. Together, these effects synergistically facilitate ion migration, as confirmed by the high ionic conductivity of 1.4 × 10−3 S cm−1. This adsorption-mediated solvation regulation fosters the formation of compact, LiF-rich interphases on both the anode and cathode, mitigating interfacial side reactions during high-rate cycling. Leveraging these synergistic effects, the Li|L-Zr3|LiFePO4 full cells deliver a high specific capacity of 120.9 mAh g−1 at 5C and maintain stable cycling over 700 cycles. This work underscores the pivotal role of two-dimensional fillers with tailored adsorption properties in solvation engineering for high-rate and long-life solid-state lithium metal batteries.

2. Results and discussion

2.1. The interfacial weak interactions in SPEs

To elucidate the interfacial interactions and their impact on solvation structures, we performed DFT calculations to obtain the electrostatic potential (ESP) distributions and adsorption energies for DMF, LiTFSI, and ZrP (Fig. 1a–c and S1, S2). ESP analysis reveals distinct negative potential regions surrounding oxygen atoms, indicative of high electron density. This suggests that in the pristine L-Zr0 electrolyte, electron-rich oxygen atoms tend to coordinate with Li+, forming ionic clusters such as Li+–O Created by potrace 1.16, written by Peter Selinger 2001-2019 C (DMF) or Li+–O Created by potrace 1.16, written by Peter Selinger 2001-2019 S Created by potrace 1.16, written by Peter Selinger 2001-2019 O (TFSI−). For ZrP, ESP analysis indicates that H, Zr, and P atoms carry positive charges, while its oxygen atoms are negatively charged. Furthermore, as illustrated in Fig. S1, ZrP possesses a layered crystalline structure; the cleaved low-energy (001) plane exposes abundant O and H atoms on interlayer surfaces, implying that Li+ may weakly interact with surface O atoms, while DMF (C Created by potrace 1.16, written by Peter Selinger 2001-2019 O) and TFSI− (S Created by potrace 1.16, written by Peter Selinger 2001-2019 O) can form weak hydrogen bonds with surface –OH groups (blue dashed lines in Fig. 1d and f).

Fig. 1. Electrostatic potential energy maps of (a) DMF, (b) LiTFSI, and (c) ZrP; density functional theory (DFT) calculations showing the interfacial weak interactions of (d) ZrP–DMF, (e) ZrP (O site)–Li+, and (f) ZrP (–OH site)–TFSI−.

Fig. 1

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Guided by these insights, DFT calculations were performed to evaluate the corresponding weak interaction energies. The results demonstrate that ZrP exhibits a significantly higher adsorption energy toward DMF (−1.5 eV) compared to the PVDF-HFP matrix (−0.178 eV) (Fig. S2). Consequently, following solvent evaporation under identical conditions, DMF coordination with Li+ in the L-Zr3 SPE is reduced, thereby altering the solvation structure and enhancing TFSI− participation in the Li+ solvation sheath. Additionally, the interaction between ZrP and LiTFSI was also investigated. As shown in Fig. 1, the interaction energy between Li+ and ZrP is 0.33 eV, suggesting that Li+ preferentially interacts with DMF in a DMF-rich environment rather than being adsorbed onto ZrP. In contrast, TFSI− exhibits a much stronger interaction with ZrP, with an adsorption energy of −2.27 eV. This pronounced difference suggests that ZrP can attenuate the coulombic attraction between Li+ and TFSI−, which is expected to lower the energy barrier for Li+ migration.

2.2. Electrolyte design and characterization

Based on the previously mentioned theoretical analysis, a nano-ZrP-mediated SPE was designed and fabricated to establish a rapid Li+ migration network. Initially, PVDF-HFP and LiTFSI were dissolved in DMF under continuous stirring to yield a clear, viscous solution. Subsequently, ZrP was introduced with thorough mixing, and the resulting homogeneous suspension was cast and heated under vacuum to fabricate the L-Zrx SPE membrane (Fig. 2a). In this architecture, PVDF-HFP serves as the polymeric matrix forming the electrolyte framework, while LiTFSI acts as the lithium source. The flake-like ZrP, characterized by a layered structure, functions as a structural bridge connecting discrete PVDF-HFP spherulites (Fig. S3). By filling voids and bridging gaps within the polymer matrix, ZrP establishes a continuous ion-conduction network that facilitates rapid ion transport across the composite electrolyte. This structural role is further validated by EDS elemental mapping (Fig. 2b) and the negligible variation in film thickness (Fig. S4): the spherical particles are predominantly composed of C and F (characteristic of PVDF-HFP), whereas the inter-spherulitic voids are enriched with Zr, O, and P. This confirms that uniformly dispersed ZrP flakes primarily occupy the voids and bridge adjacent PVDF-HFP domains.

Fig. 2. (a) Schematic illustration of the synthesis of L-Zr3 SPE. (b) EDS maps of C, F, S, P, O and Zr elements of the L-Zr3 membrane.

Fig. 2

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Notably, ZrP exhibits a BET specific surface area of 3.884 m2 g−1. Pore size distribution analysis reveals a sharp peak centered at approximately 0.9 nm within the 0.7–1.0 nm range, corresponding to the Å-scale interlayer spacing of ZrP (Fig. S5). These confined galleries are sufficiently wide to accommodate Li+ transport via a structural diffusion mechanism, wherein ions migrate through continuous dissociation/association across different solvation sheaths (<1 nm in size).27,28 Additionally, the abundant active sites on the ZrP surface interact with TFSI−, weakening the coulombic interactions between Li+ and TFSI−. This dual functionality—ion transport confined within the interlayer galleries and surface-mediated dissociation—promotes ion conduction and is anticipated to enhance ionic conductivity, thereby enabling fast battery charging. The weak interactions among ZrP, PVDF-HFP, and LiTFSI within the L-Zr3 SPE are schematically illustrated in Fig. 3, demonstrating that Li+ ions can migrate readily through the constructed electrolyte network.

Fig. 3. Schematic illustration of interfacial weak interactions accelerating fast ion migration.

Fig. 3

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For comparison, pure PVDF-HFP (denoted as L-Zr0) and a series of PVDF-HFP/ZrP composite electrolytes with varying ZrP contents (denoted as L-Zrx) were synthesized. Surface morphologies and optical images were analyzed to elucidate the influence of the filler. As shown in Fig. S6, the L-Zrx precursor solutions transitioned from transparent to white with increasing ZrP content, and ZrP was uniformly dispersed within the matrix without observable agglomeration. More significantly, as revealed in Fig. S7, the polymer particle size decreased moderately from 4.5 µm to 3.4 µm with higher ZrP loading, accompanied by a shortened ion migration path; however, no significant improvement in mechanical strength was observed (Fig. S8) upon ZrP incorporation. Instead, the decreased Young's modulus and increased elongation at break suggest that the membrane becomes more flexible and ductile. These findings suggest that ZrP physically limits polymer grain growth during solvent evaporation rather than acting as an effective nucleation site for PVDF-HFP crystallization.13 Consistent with this physical confinement, the L-Zrx series exhibited a pronounced broadening of the diffraction peak around 19° compared to the pristine L-Zr0 sample, indicating reduced crystallinity (Fig. S9). Quantitatively, the introduction of 3% ZrP (L-Zr3) reduced the melting enthalpy to −7.84 J g−1, substantially lower than that of the L-Zr0 sample (−11.76 J g−1), corresponding to a marked decrease in crystallinity (Fig. S10) as calculated using eqn (4) (SI). This implies an increased proportion of amorphous regions and enhanced segmental motion of the polymer chains. Collectively, the synergistic effects of enhanced segmental motion, reduced particle size, and filler incorporation create shorter, more continuous ion conduction pathways, laying a crucial structural foundation for the superior ionic conductivity of the L-Zr3 electrolyte. Furthermore, the L-Zr3 membrane (≈56 µm thick) demonstrates excellent flexibility, allowing it to be freely folded and curled without damage, thereby accommodating electrode deformation during repeated charge–discharge cycles (Fig. S10).

Safety concerns arising from thermal runaway hinder the practical application of lithium metal batteries. Accordingly, the thermal stability of the as-prepared electrolytes was first investigated using thermogravimetric analysis (TGA) under an N2 atmosphere. As shown in Fig. 4a, the TGA curves exhibit three distinct weight-loss stages—below 150 °C, 150–320 °C, and beyond 320 °C—corresponding to the decomposition of absorbed moisture, DMF and LiTFSI, and the PVDF-HFP matrix, respectively.29 Notably, ZrP incorporation significantly elevates the thermal decomposition temperature of the polymer matrix from 318 °C to 341 °C. This improvement is primarily attributed to the physical barrier effect of ZrP nanoplatelets and their weak interactions with DMF and TFSI−, which collectively suppress the thermal degradation of the PVDF-HFP matrix.

Fig. 4. (a) Thermogravimetric analysis (TGA) curves of L-Zr0 and L-Zr3 SPEs. (b) FTIR spectra of L-Zr0 and L-Zr3 SPEs after heat treatment at selected temperatures to identify chemical changes during the weight-loss stages. (c) FTIR spectra of L-Zrx SPEs and pure ZrP. (d) FTIR spectra of L-Zrx SPEs. (e) FTIR spectra of various precursor solutions/suspensions interacting with DMF, with an acetonitrile + LiTFSI solution included as a reference for comparison. (f) Raman spectra of L-Zr0 and L-Zr3 SPEs. (g) The fitting results of solvation structures from the Raman spectra of L-Zr0 and L-Zr3. (h) 7Li NMR spectra of L-Zr0 and L-Zr3 SPEs.

Fig. 4

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To elucidate the chemical changes occurring during weight loss, the L-Zr0 and L-Zr3 films were heat-treated at 250 °C and 320 °C and analyzed by FTIR. As shown in Fig. 4b and S12, the characteristic peaks of DMF (1633 cm−1) and TFSI− (1197 cm−1) diminish significantly, confirming their decomposition within this temperature range. Specifically, the C Created by potrace 1.16, written by Peter Selinger 2001-2019 O peak of DMF shifts to lower wavenumbers due to volatilization; this reduces DMF content and increases the relative Li+ concentration, thereby strengthening Li+-DMF coordination, polarizing the C Created by potrace 1.16, written by Peter Selinger 2001-2019 O bond, and lowering its vibrational frequency. In contrast, the TFSI− peaks exhibit a blue shift resulting from the disruption of the Li+ solvation sheath upon DMF loss, which forces Li+ to coordinate more directly with TFSI−. DMF decomposition involves volatilization and the disruption of intermolecular interactions, whereas LiTFSI decomposition likely proceeds via the cleavage of the TFSI− anion into small molecules such as imides or fluorides.30,31

Notably, after heat treatment at 320 °C, the L-Zr3 film exhibits significantly stronger PVDF-HFP peaks (880 cm−1 and 1403 cm−1) than the L-Zr0 film, although both samples show decreased peak intensities compared to their untreated counterparts. This indicates that even below the decomposition onset of L-Zr0 (320 °C), PVDF-HFP has already undergone dehydrofluorination (HF elimination) and crosslinking, which weaken its infrared absorption.32 Upon further heating above 350 °C, the resulting carbonized products thermally crack into small gaseous molecules, causing the pronounced weight loss in the TGA curves.33 Importantly, the L-Zr3 film retains stronger TFSI− characteristic peaks than L-Zr0 at 350 °C, confirming the strong anchoring effect of ZrP on TFSI−. Finally, the safety of the L-Zr3 SPE was further evaluated by heating and combustion tests. As shown in Fig. S13 and S14, the L-Zr3 SPE ceases combustion immediately upon flame removal. Even after heating above 100 °C, it exhibits significantly less shrinkage and a more stable structure than L-Zr0, confirming its flame retardancy and enhanced thermal stability afforded by the ZrP supporting framework. It should be noted that the ‘supporting effect’ of ZrP here refers to the enhanced dimensional stability and structural integrity of the membrane, rather than mechanical reinforcement.

To investigate the intermolecular interactions and solvation structures within L-Zrx solid-state electrolytes, FTIR and Raman spectroscopy were employed. As shown in Fig. 4c, compared with pure ZrP and the L-Zr0 sample, the characteristic peaks assigned to ZrP and TFSI− in the L-Zrx SPEs (x = 1, 3, 5) exhibit red and blue shifts, respectively. This confirms weak interactions between ZrP and TFSI−, a finding supported by DFT calculations. Furthermore, all samples display a broad peak in the 1633–1700 cm−1 region (Fig. 4d).

To identify the corresponding functional group, FTIR measurements were performed on a series of precursor solutions (Fig. 4e), confirming that this peak originates from the C Created by potrace 1.16, written by Peter Selinger 2001-2019 O stretching vibration of DMF. Notably, the addition of Li+ (LiTFSI) shifts this peak to lower wavenumbers, whereas the interaction between PVDF-HFP/ZrP and DMF shifts it to higher wavenumbers. Based on the deconvolution results in Fig. 4d, the C Created by potrace 1.16, written by Peter Selinger 2001-2019 O peak was resolved into two sub-bands at 1656 cm−1 and 1664 cm−1. With increasing ZrP content—particularly at 3%—the intensity of the low-wavenumber component decreases, while that of the high-wavenumber component increases. This suggests that the anchoring effect of ZrP on DMF, which competes with Li+ for interaction, alters the peak intensity distribution and profoundly influences the Li+ solvation structure.

Raman spectroscopy further elucidates this influence. The TFSI− band in L-Zr0 was deconvoluted into free anions (SSIPs: solvent-separated ion pairs, 740.5 cm−1), contact ion pairs (CIPs: where one TFSI− coordinates with one Li+, 744.7 cm−1), and aggregates (AGGs: anions coordinate with two or more Li+, 750.8 cm−1).34 Upon ZrP addition, SSIPs vanish, and the fractions of CIPs and AGGs increase from 54% to 67% and from 15% to 33%, respectively, favoring the formation of a robust anion-derived interphase (Fig. 4f and g). This conclusion is corroborated by 7Li nuclear magnetic resonance (NMR) spectroscopy, a sensitive probe for the local chemical environment of lithium ions. As shown in Fig. 4h, the 7Li NMR spectrum of L-Zr3 exhibits a downfield shift (3.11 → 3.57 ppm) relative to L-Zr0, reflecting a more deshielded lithium environment caused by increased TFSI− coordination; the electron-withdrawing effect of TFSI− lowers the local electron density around Li+.35 Collectively, these observations demonstrate that the introduction of ZrP weakens the competitive involvement of DMF in solvation, driving more TFSI− anions into the inner solvation sheath of Li+. This anion-enriched solvation environment reduces the availability of free solvent molecules, which is expected to suppress interfacial side reactions. Meanwhile, the weak –OH⋯TFSI− interactions between the ZrP surface and TFSI− render the Li+–TFSI− coordination more dynamic and prone to ligand exchange. It should be noted that the interaction between ZrP and TFSI− does not lead to significant dissociation of LiTFSI; rather, it merely attenuates the coulombic attraction within Li+–TFSI− ion pairs. This effect, together with the continuous ion-transport pathways constructed by the ZrP bridging effect, synergistically accelerates Li+ migration. DFT calculations further indicate that as the solvation sheath becomes progressively enriched with TFSI−, the LUMO level of the solvation clusters tends to shift upward, with the HOMO–LUMO gap gradually widening from 5.11 to 5.22 eV (Fig. S15 and S16). This trend is qualitatively consistent with the enhanced resistance to reductive decomposition, suggesting improved interfacial stability in the ZrP-modified electrolyte.

2.3. Physicochemical properties of electrolytes

Before evaluating the electrochemical properties, we first assessed the potential side reactions arising from residual water or surface hydroxyl groups of the ZrP filler. As shown in Fig. S17, TG-DSC analyses confirm that vacuum drying at 200 °C for 24 h effectively removes the majority of adsorbed and interlayer crystalline water from ZrP. The dehydroxylation of surface P–OH groups occurs predominantly at 480–800 °C, accompanied by a structural transformation verified by XRD.

Quantitative estimation reveals that the –OH content in the ZrP-200 filler is ∼5.91 wt%, corresponding to only ∼1.06 mg of hydroxyl groups in the composite film (containing 0.018 g ZrP-200). Given that the electrolyte membrane contains 0.6 g of LiTFSI and 0.6 g of PVDF-HFP, along with 0.018 g of ZrP-200, the total mass of these solid components is 1.218 g. Since TGA analysis indicates that residual DMF accounts for 7.78 wt% of the total membrane mass, the solid components represent 92.22 wt% of the membrane. The total membrane mass is therefore 1.218/0.9222 ≈ 1.321 g, and the residual DMF mass is ∼0.103 g. Accordingly, the mass ratios of –OH to LiTFSI and –OH to DMF are only ∼0.18 wt% and ∼1.03 wt%, respectively. These extremely low values confirm that the hydroxyl content is negligible relative to both the lithium salt and the residual solvent. Therefore, even assuming complete reaction of all hydroxyl groups with lithium metal, the resulting lithium consumption and gas evolution would be negligible and would not significantly impair the overall electrochemical performance.

The room-temperature ionic transport properties were first analyzed via electrochemical impedance spectroscopy (EIS) (Fig. 5a). As shown in Fig. 5b, room-temperature ionic conductivity exhibits a non-monotonic trend with increasing ZrP content, peaking at a loading of 3 wt% (denoted as L-Zr3). Remarkably, the L-Zr3 electrolyte delivers a maximum ionic conductivity of 1.4 × 10−3 S cm−1, markedly exceeding the values reported for most PVDF-HFP-based systems (Table S1) and nearly two times higher than that of the pristine L-Zr0 SPE (6.1 × 10−4 S cm−1). Conversely, excessive ZrP addition (e.g., 5 wt%) induces filler agglomeration, obstructing ion transport pathways and consequently diminishing ionic conductivity. To further verify the origin of the high ionic conductivity achieved in L-Zr3, two control samples (bare PVDF-HFP without ZrP and PVDF-HFP with ZrP but without LiTFSI) were prepared under identical conditions. Their ionic conductivities were only 5.75 × 10−6 and 1.24 × 10−5 S cm−1 (Fig. S18), respectively, far below that of L-Zr3 (1.4 × 10−3 S cm−1). This stark difference can be attributed to the absence of LiTFSI in these control samples: without LiTFSI, the strong interaction between Li+ and DMF (which forms solvated species such as [Li (DMF)x]+) is absent, allowing DMF to evaporate readily during the vacuum drying process, with little residual solvent.16,29 As a result, the polymer matrix becomes ionically insulating, and no effective ion-conduction pathways can be established. This contrast confirms that the coexistence of LiTFSI and DMF as well as their mutual interaction is essential for retaining sufficient solvent to facilitate ionic conduction. Owing to their insufficient ionic conductivity, these control samples could not sustain stable battery operation and were not further characterized. Beyond the salt–solvent interaction, the residual DMF content in the functional electrolytes was also quantified. Thermogravimetric analysis (TGA) indicates residual DMF contents of approximately 10.06 wt% and 7.78 wt% in L-Zr0 and L-Zr3 electrolytes, respectively (Fig. S19), implying that the superior ionic conductivity of L-Zr3 is not solely attributed to DMF retention. Temperature-dependent ionic conductivity was further analyzed using the Arrhenius equation to determine the ion transport activation energy (Ea). As depicted in Fig. 5c and S20, L-Zr3 possesses the lowest Ea (0.092 eV), signifying a lowered energy barrier for Li+ migration.36 Facilitated by efficient Li+ transport pathways, the L-Zr3 electrolyte exhibits a Li+ transference number of 0.45, surpassing those of L-Zr0 (0.30), L-Zr1 (0.34), and L-Zr5 (0.27) (Fig. 4d, e and S21). Furthermore, linear sweep voltammetry (LSV) analysis reveals an oxidation potential exceeding 4.9 V for L-Zr3 (Fig. 5f), confirming its robust electrochemical stability.37 Concurrently, the Li|L-Zr3|Li symmetric cell demonstrates a higher critical current density (CCD) of 1.3 mA cm−2 compared to 0.8 mA cm−2 for L-Zr0 (Fig. 5g).

Fig. 5. Electrochemical properties of the SPE membranes: (a) EIS plots; (b) room-temperature ionic conductivities; (c) Arrhenius plots of the ionic conductivities of SS//L-Zrx//SS cells of L-Zrx membranes. Polarization curve and corresponding impedance diagram (the inset) of the cell for (d) Li|L-Zr3|Li; (e) Li|L-Zr0|Li. (f) Linear sweep voltammetry (LSV) curves of SS//L-Zrx//Li (x = 1, 3). Lithium plating/stripping performance in Li|L-Zrx|Li symmetric cells (x = 1, 3): (g) critical current density; (h) at 0.1 mA cm−2 and 0.1 mAh cm−2; (i) at 0.3 mA cm−2 and 0.3 mAh cm−2.

Fig. 5

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Attributed to the enhanced Li+ transference number and elevated CCD, L-Zr3 promotes uniform Li+ distribution, thereby enabling fast-charging capability while effectively suppressing lithium dendrite formation. Long-term Li plating/stripping tests were conducted to further validate the inhibition of lithium dendrites and interface stability (Fig. 5h and i). The Li|L-Zr3|Li cell maintains stable cycling for over 1400 h at 0.1 mA cm−2 with a capacity of 0.1 mAh cm−2, demonstrating excellent dendrite suppression (Table S2). In contrast, the L-Zr0 cell suffers from increased polarization and short circuits within 560 h under identical conditions (Fig. 5h). Notably, L-Zr3 retains superior stability even at 0.3 mA cm−2, favorable for fast-charging applications, due to the formation of a stable and homogeneous solid electrolyte interphase (SEI) during repeated cycling (Fig. 5i). To further probe its high-rate capability, the Li‖L-Zr3‖Li symmetric cell was also tested at higher current densities of 0.5 and 1.0 mA cm−2 (Fig. S22), delivering stable operation for approximately 90 h and 10 h, respectively. The kinetic origin of this rate-dependent behavior was investigated by Tafel analysis at the electrolyte|Li interface (Fig. S23). The L-Zr3|Li interface displays a significantly higher exchange current density (0.31 mA cm−2) than L-Zr0 (0.2 mA cm−2), evidencing accelerated Li+ transfer kinetics.38

To investigate the chemical and morphological evolution at the anode/electrolyte interface, SEM, F elemental mapping, and XPS coupled with Ar+ sputtering were conducted on Li anodes retrieved from Li‖Li symmetric cells after 50 cycles at 0.1 mA cm−2 (using L-Zr0 and L-Zr3 SPEs). As depicted in Fig. 6a and b, the Li anode surface from the Li|L-Zr0|Li cell exhibits uneven and dispersed deposits (highlighted in pink) alongside a non-uniform F distribution. In contrast, the L-Zr3 cell displays a significantly more uniform and compact deposition morphology, accompanied by a homogeneous F distribution. The survey spectra (Fig. 6c and d) reveal that the SEI of L-Zr0 possesses lower F and C but higher O atomic percentages compared to L-Zr3, indicating the generation of more oxygen-containing organic products in L-Zr0.

Fig. 6. Surface characterization of lithium anodes recovered from symmetric cells after 50 cycles at 0.1 mA cm−2 and 0.1 mAh cm−2. Surface SEM images and corresponding EDS F elemental maps of Li anodes from (a) Li|L-Zr0|Li and (b) Li|L-Zr3|Li cells. Atomic percentages of elements and XPS survey spectra of the SEI at different etching depths (0, 10, and 20 nm) for (c) the Li|L-Zr0|Li cell and (d) the Li|L-Zr3|Li cell. (e and f) High-resolution C 1s XPS spectra; (g and h) high-resolution F 1s XPS spectra of the SEI from Li|L-Zrx|Li (x = 0, 3) cells.

Fig. 6

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Further analysis of the high-resolution C 1s spectra (Fig. 6e and f), which are deconvoluted into C–C/C–H, C–O, C–S, and C Created by potrace 1.16, written by Peter Selinger 2001-2019 O species (identified as DMF decomposition products),38–40 indicates that L-Zr3 exhibits a markedly higher overall fluorine content despite a higher proportion of organic carbon species. This suggests an SEI enriched with fluorine-containing components. More specifically, depth profiling reveals that at the outermost surface (0 nm), L-Zr3 contains a lower concentration of LiF than L-Zr0, indicative of an organic-rich outer layer. However, at a depth of 10–20 nm, L-Zr3 demonstrates a significantly higher F content, confirming a LiF-rich inner SEI (Fig. 6g and h). This gradient structure, characterized by an organic-rich outer layer and a LiF-rich inner layer, arises from Lewis-acid–base interactions between ZrP and TFSI−, which delay TFSI− decomposition and preserve the organic-rich outer layer. In contrast, in ZrP-free L-Zr0, TFSI− undergoes rapid decomposition, leading to premature LiF formation and uneven Li deposition.41 Additionally, the absence of characteristic peaks in the high-resolution Zr 3d and P 2p spectra confirms that ZrP does not participate in SEI formation (Fig. S24). Combined with the effective removal of residual water and the low hydroxyl content, these results demonstrate that ZrP serves as a stable and inert functional filler. Consequently, it avoids the detrimental interfacial side reactions and complex, high-cost synthesis processes typically associated with active fillers such as LATP, LAGP, and LLTO.19

2.4. Electrochemical performance of SSLMBs

The practical application potential of the L-Zr3 SPE was evaluated in ambient-temperature LiFePO4‖Li solid-state batteries, where it served concurrently as both the separator and the electrolyte. As depicted in Fig. 7a, b and S25, the cell exhibited superior rate capability, delivering average discharge capacities of 167.9, 171, 163.9, 156.3, 141.4, 130.4, 113, 90.4, and 89.4 mAh g−1 at current rates of 0.1, 0.2, 0.5, 1, 2, 3, 5, 8, and 10C, respectively. Upon reverting the current rate to 0.1C after high-rate cycling, the capacity recovered to 166.2 mAh g−1. Moreover, the Li|L-Zr3|LFP full cell exhibited lower polarization than the L-Zr0 counterpart at various current rates (Fig. 7c). These results underscore the excellent structural reversibility and electrochemical stability of the L-Zr3 cell under rapid charge–discharge conditions. In stark contrast, the L-Zr0 SPE displayed significantly diminished capacities (96.8 mAh g−1 at 5C, 72.8 mAh g−1 at 8C, and 60.9 mAh g−1 at 10C), attributable to impeded Li+ migration.

Fig. 7. Electrochemical performance of Li|L-Zr0|LFP and Li|L-Zr3|LFP full cells: (a) rate capability from 0.1C to 10C. (b) Charge–discharge profiles of the Li|L-Zr3|LFP full cell at different rates. (c) Voltage polarization bar chart of Li|L-Zr0|LFP and Li|L-Zr3|LFP full cells. (d) Long-term cycling performance at 0.5C. (e) Charge–discharge profiles of the Li|L-Zr3|LFP full cell at different cycle numbers under 0.5C cycling. (f) Long-term cycling performance at 5C.

Fig. 7

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Long-term cycling tests further highlight the advantages of the L-Zr3 modification. Specifically, at 0.5C, the L-Zr3-based battery retained ∼90.3% of its initial capacity after more than 300 cycles (decaying from 146.5 to 134.1 mAh g−1), whereas the L-Zr0 counterpart experienced a short circuit within 140 cycles (Fig. 7d and e). Notably, even under a high current rate of 5C, the L-Zr3 battery maintained a stable capacity of ∼96.7 mAh g−1 after 700 cycles with 80% retention (initial: 120.9 mAh g−1); conversely, the L-Zr0 battery exhibited a rapid capacity decay after 100 cycles (Fig. 7f).

To elucidate the contribution of ZrP to cathode interfacial stability, we analyzed LFP cathodes retrieved from Li‖L-Zr0‖LFP and Li‖L-Zr3‖LFP cells after 300 cycles at 5C; the corresponding cathodes are denoted as L-Zr0-LFP and L-Zr3-LFP, respectively. SEM images show block-like particles with visible cracks in L-Zr0-LFP, indicating structural degradation,42 whereas L-Zr3-LFP maintains well-defined particles and an intact architecture (Fig. 8a and b). EDS mapping (Fig. 8b) and SEM imaging (Fig. S26) confirm that these block-like regions are LFP, as evidenced by the uniform distribution of Fe, P and O. Consistently, XRD patterns show that the characteristic peaks of L-Zr0-LFP shift to lower angles relative to pristine LFP, indicating lattice expansion and structural degradation, while L-Zr3-LFP retains peak positions closer to those of pristine LFP (Fig. 8c). In addition, the XRD pattern of L-Zr3-LFP exhibits a more pronounced LiF characteristic peak than L-Zr0-LFP, preliminarily suggesting a higher LiF content on the surface. To further confirm the chemical composition of the CEI, XPS analysis was performed on the cycled LFP cathodes. In the F 1s spectra, L-Zr0-LFP is dominated by the C–F peak at 688.4 eV from the PVDF binder at both the surface and a depth of 10 nm, with only a weak LiF peak at 685.3 eV,42 whereas L-Zr3-LFP exhibits significantly higher LiF content at the corresponding depths (Fig. 8d and e).

Fig. 8. Surface characterization of LFP cathodes recovered from Li‖L-Zr0‖LFP and Li‖L-Zr3‖LFP full cells after 300 cycles at 5C: (a) surface SEM image of L-Zr0-LFP. (b) Surface SEM image and corresponding EDS elemental maps of L-Zr3-LFP. (c) XRD patterns of pristine LiFePO4, L-Zr0-LFP and L-Zr3-LFP after 300 cycles. High-resolution XPS spectra of the CEI (cathode electrolyte interphase) at etching depths of 0 and 10 nm: F 1s for (d) L-Zr0-LFP and (e) L-Zr3-LFP; O 1s for (f) L-Zr0-LFP and (g) L-Zr3-LFP; C 1s for (h) L-Zr0-LFP and (i) L-Zr3-LFP.

Fig. 8

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The fitted peaks (C–C, C–O and C Created by potrace 1.16, written by Peter Selinger 2001-2019 O) in the O 1s and C 1s spectra are assigned to organic decomposition products,43 and quantitative analysis indicates that the CEI on L-Zr0-LFP is more organic-rich and contains fewer inorganic species (Fig. 8f–i). These results indicate that L-Zr3 promotes the formation of a LiF-rich CEI on the LFP surface, which facilitates Li+ transport and suppresses parasitic reactions and electrolyte consumption. EIS measurements at different cycling stages (Fig. S27) further support this conclusion: the Rct of L-Zr0 increases continuously, whereas that of L-Zr3 shows only a marginal increase. The stable cathode/electrolyte interface, together with the LiF-rich SEI on the anode side, accounts for the reduced polarization and enhanced long-term cycling stability of the L-Zr3 full cell. Furthermore, representative voltage profiles (Fig. S28) revealed that the L-Zr3-based cell exhibited negligible polarization increase over cycling compared to the L-Zr0-based cell, indicating markedly enhanced interfacial stability.44 This outstanding high-rate performance is attributed to the synergistic effects of ZrP incorporation: (1) structural modulation—reducing crystallinity and constructing continuous fast-ion migration channels; and (2) interfacial engineering—inducing a uniform, dense inorganic-rich SEI that effectively suppresses lithium dendrite growth.

Beyond the moderate cathode loading used in the above tests, the practical feasibility of the L-Zr3 electrolyte was further evaluated under higher LFP loadings (13 mg cm−2, 2 mAh cm−2 and 11 mg cm−2, ∼1.72 mAh cm−2), which supported stable cycling at 0.1C and 0.2C, respectively (Fig. S29). Moreover, under a cathode loading of ∼2 mg cm−2, the cell achieved a peak discharge capacity of 113.47 mAh g−1 at 5C and retained ∼80% of its capacity after over 300 cycles (Fig. S30). The electrolyte also demonstrated high-voltage compatibility in Li‖NCM811 cells at 2C, delivering an initial discharge capacity of 146.7 mAh g−1 and retaining ∼80% of its initial capacity after 100 cycles, with an average coulombic efficiency of >98.2% (Fig. S31). Collectively, these results preliminarily validate that L-Zr3 shows promising applicability under realistic operating conditions.

3. Conclusion

In summary, a weakly coordinated L-Zr3 SPE featuring fast ion transport pathways was developed by incorporating a low-cost, multifunctional 2D layered ZrP filler for fast-charging SSLMBs. Combined theoretical and experimental analyses reveal that the introduction of nano-ZrP exerts synergistic effects via two primary mechanisms: (i) structural modification and weak interactions for fast ion migration: it reduces crystallinity, constructs continuous ion-conducting channels and weakens the coulombic trapping between Li+ and TFSI−via –OH⋯TFSI− interactions, all of which contribute to accelerated Li+ transport as confirmed by electrochemical measurements; and (ii) interfacial engineering: the filler sequesters DMF to form an anion-rich solvation structure, which, coupled with Lewis acid–base interactions, facilitates the formation of LiF-rich interphases on both the anode and cathode. This stable interface effectively suppresses dendrite growth and side reactions during high-rate cycling. Consequently, the optimized L-Zr3 electrolyte exhibits a high ionic conductivity of 1.4 × 10−3 S cm−1 and stable cycling for over 1400 h in Li‖Li symmetric cells. Notably, Li‖LiFePO4 full cells demonstrate excellent fast-charging stability, retaining 80% capacity after 700 cycles at 5C. This study highlights that adsorption-mediated filler design can effectively promote ion migration and modulate interfacial chemistry, offering new insights for developing high-rate, long-life SSLMBs.

Author contributions

Yanting Ma: writing – original draft, conceptualization, investigation, data curation, formal analysis, visualization, and validation. Songsheng Zheng: project administration, conceptualization, writing – review & editing, supervision, funding acquisition, and resources. Jianyang Wu: formal analysis, methodology (performed the DFT calculations), visualization, funding acquisition, supervision, and resources. Kai Zou: visualization. Shujian Bai: investigation. Zhaolin Wang: supervision and resources. Xiang Li: resources.

Conflicts of interest

There are no conflicts to declare.

Supplementary Material

SC-OLF-D6SC05344K-s001

Acknowledgments

This work is financially supported by Fujian Ruisen New Materials Co., Ltd (20253160A0606), Xiamen Fuhua Technology Co., Ltd (20233160A0759), and the National Natural Science Foundation of China (grant no. 12572127 and 12172314).

Data availability

The data supporting this article have been included as part of the supplementary information (SI). Supplementary information: (1) the experimental section, (2) the morphology, structure, physicochemical properties, and the related electrochemical performances of the as prepared samples, (3) the comparison of the ionic conductivity, the lithium-ion transference number, the electrochemical performance of Li/Li symmetric cells and full cells of the L-Zr3 electrolyte with previous works related to PVDF-HFP-based polymer electrolytes. (4) the supplementary references are also included in the end. See DOI: https://doi.org/10.1039/d6sc05344k.

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Associated Data

This section collects any data citations, data availability statements, or supplementary materials included in this article.

Supplementary Materials

SC-OLF-D6SC05344K-s001
SC-OLF-D6SC05344K-s001.pdf (5.3MB, pdf)

Data Availability Statement

The data supporting this article have been included as part of the supplementary information (SI). Supplementary information: (1) the experimental section, (2) the morphology, structure, physicochemical properties, and the related electrochemical performances of the as prepared samples, (3) the comparison of the ionic conductivity, the lithium-ion transference number, the electrochemical performance of Li/Li symmetric cells and full cells of the L-Zr3 electrolyte with previous works related to PVDF-HFP-based polymer electrolytes. (4) the supplementary references are also included in the end. See DOI: https://doi.org/10.1039/d6sc05344k.

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