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. 2026 Sep 15. Online ahead of print. doi: 10.1039/d6sc05957k

Local electrostatic field achieved by charge redistribution for stable solid-state lithium metal batteries

Panpan Mao a, Qi An a,✉, Xin Wang a, Nan Tang a, Cuiping Luo a, Wenwei Li a, Fanyu Xie a, Yufeng Fan a, Hongjia Liu a, Guiquan Zhao a, Yunchun Zha b, Genfu Zhao a, Qing Liu a, Xiaomin Yang c,✉, Hong Guo a,b,✉
PMCID: PMC13618588  PMID: 42808011

Abstract

Poly(vinylidene fluoride-co-hexafluoropropylene) (PVDF-HFP)-based solid-state lithium batteries (SSLBs) possess excellent safety and high energy density, but they suffer from sluggish Li+ transport and unstable interphase. A fluorinated covalent organic framework (F-COF) is fabricated and integrated into PVDF-HFP-based solid polymer electrolytes to tackle these issues. The strongly electron-withdrawing fluorine groups can realize the charge redistribution, thereby modulating the local electrostatic environment. This modulation promotes lithium salt dissociation, effectively enabling synergistic dual-ion regulation and suppressing concentration polarization. Meanwhile, F-COF induces the transition of PVDF-HFP from the α-phase to the polar β-phase to reduce polymer crystallinity and build fast ion-transport channels. Furthermore, F-COF injects electrons into TFSI− through anion-π interactions, thereby facilitating the breakage of C–F bonds to obtain a stable LiF-rich SEI film. Consequently, a high ionic conductivity of 1.34 × 10−3 S cm−1 is obtained for the F-COF@PVDF-HFP electrolyte. Li|F-COF@PVDF-HFP|Li cells demonstrate ultra-stable cycling for over 4400 h. LFP|F-COF@PVDF-HFP|Li delivers stable cycling performance for more than 400 cycles, while the LCO|F-COF@PVDF-HFP|Li maintains stable cycling beyond 300 cycles. This study develops a molecular design strategy to regulate the local electrostatic environment for SSLBs.

Fluorine-substitution induced skeleton charge redistribution constructs push–pull local electrostatic fields that accelerate Li+ transport and form a stable LiF-rich SEI.graphic file with name d6sc05957k-ga.webp

Introduction

Lithium metal stands out as the most promising anode material for constructing next-generation high-energy-density batteries.1–5 Nevertheless, traditional liquid lithium metal batteries (LMBs) suffer from safety hazards such as flammability and explosion, as well as interfacial side reactions.6–11 Solid polymer electrolytes (SPEs) represent an attractive approach to address these challenges, offering a combination of superior safety, favorable interfacial contact and ease of processing.12–14 The development of high-performance SPEs requires rational material selection and structural design. It is worth noting that poly(vinylidene fluoride-co-hexafluoropropylene) (PVDF-HFP) is recognized as an optimal choice due to its prominent thermal stability, great mechanical rigidity and favorable electrochemical stability.15–18 Nevertheless, a high level of intermolecular crystallinity in the α-phase typically leads to poor ionic conductivity at ambient temperature. Furthermore, conventional PVDF-HFP electrolytes lack selective chemical regulation of ion transport, thereby failing to reconcile the kinetic mismatch between rapid interfacial electrochemical reactions and sluggish bulk ion migration.19,20 Under an applied electric field, mobile anions migrate freely and accumulate near the anode surface. Continuous depletion of Li+ at the electrode–electrolyte interface induces pronounced Li+ concentration polarization and interfacial transport instability, ultimately leading to lithium dendrite growth.21,22 Therefore, achieving long-term stable operation of LMBs requires the active regulation of interfacial ion transport kinetics to suppress dendrite initiation.23

According to Chazalviel's space-charge theory, dendrite growth can be suppressed by slowing free anion transport and reducing their interfacial accumulation.24,25 Early studies have demonstrated that immobilizing anionic species on polymer backbones to form single-ion conducting electrolytes increases the Li+ transference number and mitigates concentration polarization. However, elevated local anionic charge density strengthens Li+–anion interactions, thereby hindering ion dissociation and restricting Li+ transport, which results in reduced ionic conductivity and poor rate performance.26,27 Hence, regulating either Li+ or anion transport alone is insufficient to suppress dendrite formation, highlighting the need for synergistic dual-ion regulation to enhance electrochemical performance. The electronic structure of the electrolyte governs charge distribution and electrostatic potential, which in turn tunes Li+–framework interactions, facilitates ion-pair dissociation, and adjusts the relative contributions of cation and anion transport.28–31 Building on this regulation principle, covalent organic frameworks (COFs), featuring precisely engineered functional groups and well-defined nanoporous channels, provide a structural platform for directional ion transport control.32–35 Tailoring framework electronegativity creates a push–pull electrostatic microenvironment that promotes Li+ transport while repelling anions, enabling synergistic dual-ion regulation and suppressing concentration polarization. Meanwhile, regulating the local charge distribution around lithium salt anions can promote their reductive decomposition. This favors the formation of a chemically stable LiF-rich SEI layer, which effectively suppresses interfacial side reactions and improves battery cycling stability. Nevertheless, it remains a significant challenge to rationally design framework electronegativity at the molecular level for regulating the local electrostatic environment of PVDF-HFP-based SPEs, achieving synergistic bulk-interface optimization, and establishing an electronic structure–transport property relationship.

Herein, we prepare a novel F-COF-modified PVDF-HFP-based SPE for LMBs. Scheme 1 illustrates the mechanism of F-COF for promoting Li+ transport and improving interfacial chemistry through the regulation of the local electrostatic environment. First, the strongly electron-withdrawing F-group can induce the charge redistribution to increase the framework electronegativity, establishing a local electrostatic field that promotes Li+ affinity and anion repulsion. Second, polarized PVDF-HFP chains are induced to transform into the β-phase by the chemical complexation of C–F bonds in COF, which leads to reduced crystallinity and generates more amorphous regions in the polymer, so fast and stable Li+ transport channels are constructed to enhance Li+ transport dynamics. Third, the electron-rich sites in F-COF form anion-π interactions with TFSI− to inject electrons into the TFSI−, consequently weakening C–F bond polarity to decrease the cleavage energy barrier and further forming a dense LiF-rich SEI layer. This interphase layer effectively inhibits the uncontrolled propagation of lithium dendrites, greatly boosting the cycle life and safety performance of batteries.36,37 The long-range electrostatic repulsion of TFSI− and its strong short-range interaction with the framework may appear contradictory, but they operate at distinct spatial scales.38,39 Far from the interface, the negative electrostatic potential of the fluorinated framework repels TFSI− and attracts Li+. When TFSI− approaches the electron-deficient aromatic ring to an equilibrium distance of approximately 3 Å, however, the anion-π interaction takes over as the dominant force. Driven by electrostatics and ion-induced polarization, this interaction brings TFSI− into close contact with the π system and enables charge transfer, including electron injection into the sigma antibonding orbital of the C–F bond.40–42 Both effects share an electrostatic origin and dominate at different distances, so they are not mutually contradictory. Based on this design, at ambient temperature, the F-COF@PVDF-HFP composite electrolyte reaches an ionic conductivity of 1.34 × 10−3 S cm−1. Cells assembled with this electrolyte deliver desirable cycling stability for LFP cathodes, retaining 95% of the initial capacity after 400 cycles. Meanwhile, a high capacity retention of 93.7% can be maintained for LCO cells following 300 cycles. Notably, the Li|F-COF@PVDF-HFP|Li cell achieves excellent long-term cycling stability over 4400 h at 0.1 mA cm−2. This work proposes a molecular modification strategy for optimizing the local ionic environment to accelerate Li+ migration and regulate interfacial chemistry, providing new insights for improving the performance of SPEs.

Scheme 1. The mechanism by which F-COF promotes Li+ transport and the formation of LiF-rich SEI.

Scheme 1

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Results and discussion

Powder X-ray diffraction (PXRD) measurement is carried out to characterize the crystallinity of the two COFs. Illustrated in Fig. 1a, three obvious diffraction signals appear in the PXRD pattern of TPPA-COF at 4.6°, 12.1° and 26.9°, which can be assigned to the (100), (210) and (001) lattice planes separately. For F-COF, diffraction peaks appear at 4.7°, 7.8°, 9.5°, 12.1° and 27.2°, corresponding to the (100), (110), (200), (210) and (001) planes, respectively (Fig. 1b).43 In TPPA-COF, the peak observed at 7.98° is ascribed to the merging of the (110) and (200) diffraction peaks, a phenomenon that is supported by the existing literature.44–46 The simulated PXRD patterns of both COFs after geometric optimization match well with the experimental results, both corresponding to an AA stacking mode (Fig. S1). Both COFs possess identical pore topologies, which decouple geometric pore effects from intrinsic electronic effects. Fluorine substitution alone triggers charge redistribution over the backbone to tune framework electronegativity and generate distinct local electrostatic microfields. To explore the structural features of COFs, Fourier-transform infrared (FT-IR) spectroscopy is carried out. As depicted in Fig. 1c, the relevant spectra of 2,3,5,6-tetrafluoro-1,4-phenylenediamine (F-DA) and 1,3,5-triformylphloroglucinol (TFP) show obvious characteristic peaks of –NH2 (3430 and 3324 cm−1) and –CHO (2892 and 1664 cm−1), respectively. Once the reaction procedure is accomplished, the characteristic peaks attributed to –NH2 and –CHO nearly vanish. Meanwhile, F-COF displays the stretching vibration of C Created by potrace 1.16, written by Peter Selinger 2001-2019 C at 1595 cm−1 and C–N stretching bands at 1241 cm−1, similar results are obtained for TPPA-COF (Fig. S2). The above evidence verifies the effective occurrence of the Schiff-base reaction and the feasible fabrication of target COFs. Solid-state 13C nuclear magnetic resonance (NMR) spectroscopy further confirms the chemical constitution of the COFs (Fig. 1d and S3). Characteristic peaks of TPPA-COF and F-COF at 185, 150, 140, 110 and 105 ppm can be attributed to ketone R(CO)R, alkene C Created by potrace 1.16, written by Peter Selinger 2001-2019 C, aromatic ring C Created by potrace 1.16, written by Peter Selinger 2001-2019 C, and two types of R2C Created by potrace 1.16, written by Peter Selinger 2001-2019 CR alkene structures, respectively.47–49 Moreover, F-COF exhibits an intense signal at 125 ppm assigned to fluoroalkene R Created by potrace 1.16, written by Peter Selinger 2001-2019 (CF)R, whereas TPPA-COF presents split peaks from the aromatic carbon. Fluorinated olefin moieties act as core sites for charge redistribution. Strong electron-withdrawing fluorine atoms shift electron clouds of conjugated skeletons to raise overall framework electronegativity and construct push–pull local electrostatic fields within pore channels. The chemical composition of TPPA-COF and F-COF is further explored via X-ray photoelectron spectroscopy (XPS) (Fig. S4 and S5). The survey XPS spectrum clearly confirms the presence of fluorine in F-COF, while no fluorine signal is observed in TPPA-COF. For the F 1s spectrum of F-COF, a characteristic peak corresponding to semi-ionic C–F bonds emerges at 687.4 eV. This signal stems from the coordination interaction between fluorine atoms and neighboring sp2-hybridized carbon atoms (Fig. 1e).50 As shown in Fig. 1f, the O 1s spectrum shows peaks for C–O and C Created by potrace 1.16, written by Peter Selinger 2001-2019 O. Semi-ionic C–F bonds serve as vital structures triggering global charge redistribution across the skeleton, which directly alters electrostatic potentials on pore walls and creates local electrostatic fields capable of selective cation–anion sieving. Subsequently, the porous structures of TPPA-COF and F-COF are evaluated by low-temperature (77 K) N2 adsorption–desorption tests. The Brunauer–Emmett–Teller (BET) technique determines the specific surface areas to be 202.9 m2 g−1 for F-COF (Fig. 1g) and 1080.8 m2 g−1 for TPPA-COF (Fig. S6). Meanwhile, F-COF exhibits a confined pore size distribution centered around 1.9 nm. Regular confined pores supply stable spatial environments for local electrostatic fields derived from charge redistribution, enabling full interactions between ions and electrostatic fields to realize synergized regulation of Li+ attraction and TFSI− repulsion. Moreover, to identify the morphological characteristics of the two COF samples, scanning electron microscopy (SEM) and transmission electron microscopy (TEM) are employed for structural observation. Both F-COF and TPPA-COF exhibit a flake-like morphology (Fig. S7 and S8). For F-COF, elements including C, N, O and F present homogeneous distribution across the material (Fig. 1h). These findings confirm the successful fabrication of F-COF and TPPA-COF. Homogeneous fluorine doping ensures uniform charge redistribution throughout the entire skeleton and yields continuous, defect-free distribution of local electrostatic fields inside pores.

Fig. 1. (a) XRD patterns of TPPA-COF. (b) XRD patterns of F-COF. (c) FTIR spectra of F-COF. (d) 13C NMR spectra of F-COF. (e) XPS F1s spectra of F-COF. (f) XPS O 1s spectra of F-COF. (g) Nitrogen adsorption desorption isotherms of F-COF (inset: pore size distribution). (h) TEM image of F-COF and EDS mapping of C, N, O, F element.

Fig. 1

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The solution casting technique is employed to prepare PVDF-HFP-based SPEs. SEM is conducted to observe the morphology of PVDF-HFP-based SPE films. As presented in Fig. 2a and S9, the C, N, O, F, and S elements all display homogeneous dispersion, demonstrating that COF fillers uniformly distribute within the PVDF-HFP-based SPEs without obvious agglomeration. Uniform dispersion of F-COF generates continuous and evenly distributed charge redistribution sites inside polymers and builds interconnected channels of local electrostatic fields throughout the electrolyte matrix. The F-COF@PVDF-HFP has a thickness of ≈43 µm (Fig. 2b), while the thickness of the TPPA-COF@PVDF-HFP is ≈38 µm (Fig. S10). Furthermore, the photograph of the F-COF@PVDF-HFP electrolyte reveals a uniform surface, and it remains crease-free even after repeated folding, indicating the excellent toughness of the film (Fig. 2c). A similar situation is also observed for the PVDF-HFP and TPPA-COF@PVDF-HFP electrolytes (Fig. S11). As shown in Fig. 2d and e, confocal laser scanning microscope (CLSM) identifies surface roughness and morphological differences between the two electrolytes. The PVDF-HFP features an extremely rugged surface with a height gap of approximately 18.03 µm. On the contrary, F-COF@PVDF-HFP owns a level and smooth surface with a height deviation of merely ≈6.90 µm. Smooth interfaces eliminate anion accumulation arising from local regions lacking electrostatic fields and mitigate localized concentration polarization at electrode–electrolyte boundaries. Furthermore, as illustrated in Fig. 2f, thermogravimetric (TG) analysis is performed to determine the composition of the two COF-doped electrolytes, which primarily consists of solvent molecules, lithium salt, and the PVDF-HFP substrate. The crystalline structure of PVDF-HFP-based SPEs is explored via PXRD measurement. As illustrated in Fig. 2g, the diffraction peak located at 18.5° matches the (020) plane of the α-phase, while the diffraction peak at 20.1° is attributed to the (110/200) planes of the β-phase.51,52 Notably, compared with TPPA-COF@PVDF-HFP (37.1%) and PVDF-HFP (46.1%), F-COF@PVDF-HFP shows a lower degree of crystallinity (23.8%), indicating that the incorporation of F-COF effectively lowers the crystallinity of PVDF-HFP (Table S1). Additionally, FT-IR results reveal that intense β-phase absorption bands appear at 838 and 881 cm−1 for F-COF@PVDF-HFP, while faint α-phase signals are observed at 759 and 790 cm−1 (Fig. 2h). The incorporation of F-COF induces the transition of PVDF-HFP from the α-phase to the polar β-phase, driven by the polarization traction of local electrostatic fields on polymer molecular dipoles. These local electrostatic fields are constructed by fluorine-triggered skeleton charge redistribution. This phase transition reduces the polymer crystallinity and generates more amorphous regions, thereby constructing fast and stable Li+ transport channels. These findings reveal that the F-COF@PVDF-HFP electrolyte can markedly facilitate Li+ transport. Meanwhile, the stress–strain curves show that the F-COF@PVDF-HFP membrane delivers a higher tensile strength of 12.77 MPa, compared with 7.04 MPa for pure PVDF-HFP (Fig. 2i). As depicted in Fig. 2j, the maximum piercing force of F-COF@PVDF-HFP reaches 1.37 N, that of TPPA-COF@PVDF-HFP is 1.14 N, and that of pure PVDF-HFP is 0.61 N. The above results confirm that the F-COF@PVDF-HFP can efficiently hinder the generation of lithium dendrites. Combined with anion repulsion effects induced by charge redistribution, a dual dendrite suppression mechanism integrating physical blocking and electrostatic sieving is established. The chemical shift of Li shifted from −0.57 ppm for PVDF-HFP to −0.47 ppm for TPPA-COF@PVDF-HFP, and finally to −0.36 ppm for F-COF@PVDF-HFP (Fig. 2k), indicating the highest concentration of free Li+ in the F-COF@PVDF-HFP electrolyte, which matches well with the Raman results. Essentially, local electrostatic fields formed through F-COF charge redistribution competitively trap Li+, weaken Li+-TFSI− association, and liberate abundant free lithium cations. As observed in Fig. 2l and S12, for the PVDF-HFP electrolyte, TFSI− almost exclusively exists in the form of ion clusters (747–750 cm−1), indicating that PVDF-HFP is ineffective in promoting lithium-salt dissociation. In contrast, for F-COF@PVDF-HFP, the proportion of TFSI− present in the dissociated state (740–744 cm−1) is 32.8%, which is higher than those of PVDF-HFP (7.7%) and TPPA-COF@PVDF-HFP (18.6%).53,54 This demonstrates that the incorporation of F-COF can regulate the charge distribution around TFSI−, weaken the interaction between TFSI− and Li+, and effectively promote lithium salt dissociation. This provides direct microscopic evidence that local electrostatic fields achieve synergized cation–anion regulation and suppress concentration polarization.

Fig. 2. (a) SEM images of the top-view of F-COF@PVDF-HFP and energy-dispersive X-ray maps of F-COF@PVDF-HFP. (b) Cross section image of F-COF@PVDF-HFP. (c) Optical photographs of F-COF@PVDF-HFP. Confocal laser scanning microscopy (CLSM) images of (d) PVDF-HFP and (e) F-COF@PVDF-HFP. (f) TGA curves of TPPA-COF@PVDF-HFP and F-COF@PVDF-HFP. (g) XRD patterns of PVDF-HFP, TPPA-COF@PVDF-HFP, and F-COF@PVDF-HFP. (h) FTIR spectra of PVDF-HFP, TPPA-COF@PVDF-HFP, and F-COF@PVDF-HFP. (i) Stress–strain curves of PVDF-HFP-based SPEs. (j) Puncture curves of PVDF-HFP-based SPEs. (k) 7Li NMR spectra of PVDF-HFP-based SPEs. (l) Raman spectra of PVDF-HFP and F-COF@PVDF-HFP.

Fig. 2

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Stainless steel symmetric cells are adopted to determine the Li+ conductivity of PVDF-HFP- based SPEs from 25 °C to 80 °C for evaluating their ion migration capability. Notably, F-COF@PVDF-HFP possesses optimal ionic conductivity throughout the measured temperature interval. At 25 °C, the lithium-ion conductivity of F-COF@PVDF-HFP reaches 1.34 × 10−3 S cm−1, which far exceeds that of PVDF-HFP (2.8 × 10−4 S cm−1) and TPPA-COF@PVDF-HFP (5.3 × 10−4 S cm−1) (Fig. 3a and S13). As depicted in Fig. 3b, F-COF@PVDF-HFP exhibited a reduced activation energy (0.124 eV) relative to PVDF-HFP (0.203 eV) and TPPA-COF@PVDF-HFP (0.154 eV). Local electrostatic fields constructed via charge redistribution form oriented electrostatic potential wells for Li+ transport, reduce migration barriers of lithium ions, and greatly boost Li+ transport kinetics. The Li+ transference number of F-COF@PVDF-HFP is 0.68 (Fig. 3c and Table S2), surpassing that of pristine and TPPA-COF modified electrolytes (0.34 and 0.47, respectively) (Fig. S14). These observations confirm that the incorporation of F-COF into PVDF-HFP can efficiently facilitate Li+ migration. High-voltage electrochemical stability serves as a crucial criterion for evaluating the compatibility between electrolytes and high-voltage cathode materials. Therefore, linear sweep voltammetry (LSV) is employed to characterize the electrochemical stability window of the prepared electrolytes. The F-COF@PVDF-HFP electrolyte presents an extended electrochemical working window of 4.82 V, exceeding those of TPPA-COF@PVDF-HFP (4.7 V) and PVDF-HFP (4.56 V). This result demonstrates that F-COF@PVDF-HFP is well suited for application with high-voltage cathode systems (Fig. 3d). Uniform local electrostatic fields stabilize reduction pathways of anions and restrain electrolyte oxidative degradation triggered by unregulated anion decomposition. To further assess the interfacial contact properties of SPEs toward lithium metal anodes (LMAs), the Tafel measurements are conducted on lithium symmetric cells (Fig. 3e). The exchange current density (j0) of F-COF@PVDF-HFP reaches 0.32 mA cm−2, outperforming PVDF-HFP (0.15 mA cm−2) and TPPA-COF@PVDF-HFP (0.21 mA cm−2). This difference confirms that F-COF facilitates the generation of inorganic-rich SEI film and optimizes interfacial chemistry. The dendrite suppression capability of F-COF@PVDF-HFP is further verified by systematically testing the lithium deposition/dissolution behaviors of lithium symmetric cells. As illustrated in Fig. 3f, critical current density (CCD) tests are performed on these symmetric cells. The Li|F-COF@PVDF-HFP|Li cell exhibits an outstanding CCD value (2.0 mA cm−2), outperforming the other two counterparts. Furthermore, the long-cycle performance of these lithium symmetric cells is assessed. As revealed in Fig. 3g, under 0.1 mA cm−2, the Li|F-COF@PVDF-HFP|Li symmetric cell realized a stable ultralong cycling life exceeding 4400 h, which is much longer than those of the Li|PVDF-HFP|Li cell (450 h) and the Li|TPPA-COF@PVDF-HFP|Li cell (800 h). Meanwhile, with the increase in current density, F-COF@PVDF-HFP maintains stable voltage fluctuation under a high current density of 0.8 mA cm−2, demonstrating excellent rate capability (Fig. 3h). Nevertheless, the PVDF-HFP electrolyte shows a rapid increase in polarization voltage after 0.3 mA cm−2, leading to a short circuit and preventing stable cycling for long periods of time. These results demonstrate that the introduction of F-COF can effectively facilitate Li+ migration and induce the generation of a stable SEI layer, thus greatly improving the battery cycling stability. Steady polarization over long cycles directly verifies that local electrostatic fields formed by charge redistribution durably suppress concentration polarization and prevent dendrite-induced short circuits caused by persistent ion depletion.

Fig. 3. (a) The Nyquist curves of F-COF@PVDF-HFP. (b) Arrhenius plots of the ionic conductivities of PVDF-HFP, TPPA-COF@PVDF-HFP, and F-COF@PVDF-HFP. (c) DC polarization curve (the inset shows the Nyquist plots) of Li symmetric cell with F-COF@PVDF-HFP. (d) LSV curves of PVDF-HFP, TPPA-COF@PVDF-HFP, and F-COF@PVDF-HFP. (e) Tafel plots of the Li‖Li symmetric cells. (f) Critical current density (CCD) of the Li‖Li symmetric cells. (g) Galvanostatic cycling curves of the Li‖Li symmetric cells with PVDF-HFP, TPPA-COF@PVDF-HFP, and F-COF@PVDF-HFP at 0.1 mA cm−2. (h) Different rates of cycling performance of symmetric cells.

Fig. 3

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The Li plating behavior on the anode in different Li symmetric cells is visually monitored in real time using in situ optical microscopy. As depicted in Fig. 4a, cross-sectional images of the lithium symmetric cell are captured at four different time points during cycling at 2 mA cm−2. Before deposition, the lithium foils in both cells exhibit smooth surfaces. After 5 min of Li deposition, the battery with PVDF-HFP gradually generates lithium dendrites. Without modulation from local electrostatic fields generated by charge redistribution, freely migrating anions accumulate to trigger interfacial concentration polarization, localized Li+ depletion and subsequent dendrite growth. As the deposition time increases, an increasing number of dendrites are observed, and after 30 minutes, large lithium dendrites nearly occupy the entire cell interior, readily triggering battery short-circuit failure and posing significant safety hazards. On the contrary, Li|F-COF@PVDF-HFP|Li cell inhibits dendrite growth and achieves homogeneous Li deposition. The thickness of the Li-plated layer progressively increases over time. These observations confirm that F-COF@PVDF-HFP effectively suppresses the nucleation of lithium dendrites and promotes uniform lithium deposition. The fluorine groups in F-COF induce skeleton charge redistribution to construct local electrostatic fields that repel anions and homogenize the Li+ flux, while the anion-π interactions between F-COF and TFSI− promote the formation of a dense LiF-rich SEI. These effects jointly enable uniform Li deposition and effectively suppress dendrite growth. As presented in Fig. S15, the polarization resistance of lithium symmetric cells after 100 cycles is analyzed by distribution of relaxation times (DRT). The DRT spectrum can be divided into three segments, representing the resistance of the SEI layer (RSEI), charge transfer resistance (RCT), and ion diffusion resistance (RID), respectively.55 Compared with the Li‖Li cell using PVDF-HFP electrolyte, the lithium symmetric cell with F-COF@PVDF-HFP electrolyte delivers lower SEI film resistance. Moreover, this work tracks the impedance evolution of lithium symmetric cells throughout different cycling periods (Fig. 4b–e). The comparison shows that F-COF@PVDF-HFP based symmetric cell possesses lower impedance and superior Li+ conduction kinetics than pure PVDF-HFP counterpart. This phenomenon is contributed by the special framework of F-COF. Local electrostatic fields constructed through skeleton charge redistribution directionally alter electron cloud distribution of anions, accelerate anion decomposition and in situ generate LiF-rich SEI layers. To clarify the internal mechanism by which F-COF@PVDF-HFP suppresses Li dendrite formation and improves interfacial stability, XPS is utilized to investigate the lithium metal anode after 50 cycles (0.1 mA cm−2, 0.1 mA h cm−2) (Fig. 4f, g and S16–S18). According to C 1s and O 1s spectral analysis, functional groups including C–C, C–O, C Created by potrace 1.16, written by Peter Selinger 2001-2019 O, C-O-C, and CF3 stem from the breakdown of solvents and lithium salt during lithium deposition and dissolution cycles.56 The F 1s spectrum displays two characteristic peaks corresponding to LiF (685 eV) and CF3 (689 eV), indicating that the SEI film formed with the F-COF@PVDF-HFP electrolyte delivers an elevated LiF proportion. Local electrostatic fields continuously block massive anions from reaching anode surfaces, reduce disordered byproducts and form homogeneous LiF-enriched interphases. Similarly, the Li 1s spectrum verifies that the SEI film of the F-COF@PVDF-HFP cell possesses an obviously increased LiF proportion.The changes in the lithium salt within the electrolyte under the influence of electrochemical potential are monitored via in situ FT-IR spectroscopy. Fig. 4h displays the in situ FT-IR battery model established in this work. As shown in Fig. 4i and j, the absorption peaks corresponding to S–N–S (1048 cm−1), S Created by potrace 1.16, written by Peter Selinger 2001-2019 O (1117 cm−1), and C–F (1180 cm−1) in LiTFSI gradually diminish during cyclic charge/discharge operation. This indicates that the F-functional groups effectively boost the dissociation of LiTFSI, thus elevating the concentration of free Li+ and facilitating the formation of LiF-rich SEI.

Fig. 4. (a) Optical microscopy observation of the Li plating in Li/PVDF-HFP/Li and Li/F-COF@PVDF-HFP/Li. EIS spectra of Li|PVDF-HFP|Li (b) and Li|F-COF@PVDF-HFP|Li (d) at different cycles. DRT analysis of the EIS data of Li‖Li symmetric cells at different cycle for PVDF-HFP (c) and F-COF@PVDF-HFP (e). XPS characterization of the electrodes for F-COF@PVDF-HFP (f). (g) The percentage of each integrated peak area in the F 1s spectrum. (h) Analog cell for in situ FT-IR. (i) In situ FT-IR spectra and (j) contour maps of the F-COF@PVDF-HFP electrode/electrolyte interface at different discharge and charge steps.

Fig. 4

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Cyclic voltammetry (CV) is performed to study the kinetic reaction behavior and Li+ transport characteristics in Li|LiFePO4 batteries. With the scan rate rising from 0.1 to 0.5 mV s−1, the oxidation peak voltage gap (ΔV) of F-COF@PVDF-HFP is 0.213 V, which is smaller than those for TPPA-COF@PVDF-HFP (0.272 V) and PVDF-HFP (0.338 V) (Fig. S19). This finding indicates that F-COF@PVDF-HFP exhibits superior redox reaction reversibility and a lower polarization voltage (Fig. 5a). To study the practical feasibility of the electrolytes, half-cells are assembled by pairing the PVDF-HFP-based SPEs with different cathode materials (LiCoO2 (LCO) and LiFePO4 (LFP)). The electrochemical properties of these cells are further evaluated. Fig. 5b reveals that the LCO battery based on F-COF@PVDF-HFP possesses outstanding rate capability, accompanied by discharge capacities of 163.3, 155.2, 147.1, 139.4, 130.0, and 123.2 mA h g−1 under current densities of 0.1C, 0.2C, 0.5C, 1C, 2C, and 3C, respectively. On the contrary, the LCO|TPPA-COF@PVDF-HFP|Li and LCO|PVDF-HFP|Li cells exhibit lower discharge capacities of 163.2, 146.7, 135.5, 125.9, 111.1, 103.3 mA h g−1 and 154.6, 120.0, 115.4, 111.7, 106.4, 101.8 mA h g−1, respectively, under consistent operating conditions. Furthermore, the long-cycle results at 1C (Fig. 5c) show that the LCO|F-COF@PVDF-HFP|Li cell still retains 93.7% of its initial capacity over 300 cycles. By contrast, the LCO|TPPA-COF@PVDF-HFP|Li cell exhibits inferior cycling performance, with the coulombic efficiency and discharge capacity fluctuating abruptly after 190 cycles, and the LCO|PVDF-HFP|Li cell only cycles stably for 110 cycles, the result originates from internal short circuits triggered by uncontrolled lithium dendrite growth. At 0.1C, 0.2C, 0.5C, 1C, 2C, and 3C, the discharge capacities of the LFP|F-COF@PVDF-HFP|Li cell reach 165.9, 165.0, 161.8, 155.6, 142.8, and 130.7 mA h g−1, respectively. In comparison, the LFP|PVDF-HFP|Li and LFP|TPPA-COF@PVDF-HFP|Li cells exhibit lower discharge capacities at various rates (Fig. 5d). Meanwhile, The long-cycle performance of cells matched with LFP cathode is tested at 1C (Fig. 5e). The LFP|F-COF@PVDF-HFP|Li cell delivers superior cycling durability, keeping 95% of its initial capacity over 400 cycles. In comparison, both the LFP|PVDF-HFP|Li and LFP|TPPA-COF@PVDF-HFP|Li cells exhibit greatly deteriorated cycling stability. As shown in Fig. 5f, Li|SS cells are fabricated and subjected to CV measurements to explore the redox behavior of the electrolyte towards LiTFSI. From the CV curves of the three electrolytes, the highest redox peak current is observed for F-COF@PVDF-HFP. The higher peak intensity indicates a more vigorous redox reaction of TFSI−, signifying faster kinetics of TFSI− in the F-COF@PVDF-HFP electrolyte. Moreover, the interfacial kinetics of the cathode during charge–discharge cycling are explored by means of in situ DRT characterization (Fig. 5g and k). Distinct electrochemical processes are represented by different relaxation time constants (τ). The DRT spectra The DRT spectra are divided into three components: CEI/SEI film impedance, interfacial charge transfer impedance, and ionic diffusion impedance. During charging, the LFP|F-COF@PVDF-HFP|Li cell exhibits lower impedance than the LFP|PVDF-HFP|Li cell, demonstrating faster Li+ transport kinetics (Fig. 5h, i, l and m). Simultaneously, during discharging, the Li|F-COF@PVDF-HFP|LFP cell also shows a lower resistance (Fig. S20). The obvious impedance reduction benefits from the homogeneous ion migration enabled by F-COF@PVDF-HFP, which further forms a stable electrode–electrolyte interface.57 Subsequently, in situ differential electrochemical mass spectrometry (DEMS) is used to track gas release throughout the cycling process. The LFP|Li cell employing the F-COF@PVDF-HFP electrolyte generates barely any gas, even when cycled at an extremely high voltage of 4.5 V (Fig. 5n). Comparatively, the PVDF-HFP electrolyte based cell produces a substantial amount of H2, O2, and CO2 (Fig. 5j).

Fig. 5. (a) CV curves at different scan rates obtained from Li-LFP cells based on F-COF@PVDF-HFP. (b) Rate capability of LCO cells at different current densities. (c) Long-term cycling performances of LCO at 1C. (d) Rate capability of LFP cells at different current densities. (e) Long-term cycling performances of LFP at 1C. (f) The CV results of Li‖stainless steel (SS) cells in three electrolytes at 25 °C. In situ EIS spectra of LFP|PVDF-HFP|Li (g) and LFP|F-COF@PVDF-HFP|Li (k) during discharging and charging process. DRT analysis of the EIS data of LFP cells at different voltages using PVDF-HFP (h and i) and F-COF@PVDF-HFP (l and m). Voltage profile and corresponding in situ DEMS results of various mass signals for the LFP‖Li battery with PVDF-HFP (j) and F-COF@PVDF-HFP (n).

Fig. 5

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Electronic density of states and electrostatic potential calculations indicate that, compared to TPPA-COF, F-COF achieves enhanced electron delocalization, strongly demonstrating that the introduction of fluorine-containing groups promotes electron transfer (Fig. 6a–d). Such electron displacement essentially originates from fluorine-induced charge redistribution over the skeleton, which directly forms push–pull local electrostatic fields within pore channels. In addition, the charge density difference calculations further visualize this redistribution (Fig. S21). The green isosurfaces correspond to charge accumulation regions and the blue isosurfaces to charge depletion regions at an isovalue of 0.06e Å−3. Compared with TPPA-COF, F-COF shows significantly more charge accumulation and depletion regions over the skeleton, especially around the fluorinated aromatic rings, indicating that fluorine substitution greatly enhances charge separation and electron delocalization within the framework. Therefore, incorporating F-COF into PVDF-HFP-based SPEs optimizes the local coordination environment of Li+, thereby facilitating Li+ migration and enabling the construction of LiF-rich SEI, effectively suppressing lithium dendrite propagation and improving cycling performance. To reveal the intrinsic mechanism by which F-COF@PVDF-HFP enhances Li+ migration kinetics, density functional theory (DFT) calculations are performed. To begin with, DFT simulations are adopted to systematically examine the interactions between Li+ and each component. Different species exhibit distinct binding energies with Li+, ranked in the order of Li–F-COF (−2.807 eV), Li-TPPA-COF (−2.397 eV), Li-PVDF-HFP (−2.247 eV), and Li-TFSI− (−2.110 eV) from high to low (Fig. 6e). The stronger binding energy between F-COF and Li+ indicates that local electrostatic attraction derived from high framework electronegativity competitively captures Li+, breaks Li+-TFSI ion pairs and facilitates dissociation of lithium salts. Furthermore, by simulating the Li+ transport process, it is found that the migration rate of lithium ions in TPPA-COF@PVDF-HFP is remarkably lower than that in F-COF@PVDF-HFP (Fig. 6f–h). This indicates that incorporating F-COF into PVDF-HFP efficiently modulates the local coordination environment of lithium ions, thereby accelerating Li+ transport. In addition, mean square displacement (MSD) analyses demonstrate that F-COF@PVDF-HFP delivers a greater Li+ flux compared to TPPA-COF@PVDF-HFP (Fig. 6i). As shown in Fig. 6j, the dissociation energy of the lithium salt itself is 5.915 eV. After the addition of TPPA-COF, the dissociation energy decreases to 4.744 eV, and with the incorporation of F-COF, it reaches the lowest value of 4.105 eV. This indicates that the incorporation of COFs can accelerate the dissociation of lithium salts. To further clarify the underlying mechanism by which F-COF modulates Li+ migration, molecular dynamics (MD) simulations are utilized to explore the local coordination environment of Li+ (Fig. 6k and m). The radial distribution functions (RDFs) for Li–N (Li+/TFSI−), Li–F (Li+/TFSI−), Li–O (Li+/TFSI−), and Li–S (Li+/TFSI−) are analyzed. In contrast to the TPPA-COF@PVDF-HFP system (Fig. 6l), the g(r) peak intensities of the above Li+-TFSI− pairs are distinctly diminished in F-COF@PVDF-HFP (Fig. 6n). This phenomenon reveals that the introduction of F-COF weakens the binding interaction of Li+ and TFSI−. The results intuitively prove that local electrostatic fields generated by charge redistribution weaken coulombic association between cations and anions, achieve dual-ion coordinated regulation and suppress concentration polarization.

Fig. 6. Localized orbital locator of TPPA-COF (a) and F-COF (b). The optimized geometric configurations and electrostatic potential of TPPA-COF (c) and F-COF (d). (e) Adsorption energies of PVDF-HFP, TPPA-COF, F-COF, and TFSI− with Li+. Theoretical elucidation of Li-ion migration behaviors inside the TPPA-COF@PVDF-HFP (f) and F-COF@PVDF-HFP (g) with corresponding energy diagrams (h). (i) Calculated mean square displacement (MSD) of Li+ as a function of the simulation time. (j) The dissociation energy of pristine LiTFSI and the dissociation energies in the presence of TPPA-COF and F-COF. Theoretical calculation model of TPPA-COF@PVDF-HFP (k) and F-COF@PVDF-HFP (m). Li+ RDFs of TPPA-COF@PVDF-HFP (l) and F-COF@PVDF-HFP (n).

Fig. 6

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Conclusion

In summary, we successfully synthesize F-COF and integrate it into PVDF-HFP-based SPEs. Using a molecular-level regulation strategy enabled by fluorine-substitution induced skeleton charge redistribution, we construct push–pull local electrostatic fields that synergistically accelerate Li+ transport and form a highly stable LiF-rich SEI layer. The strong electron-withdrawing fluorine sites anchored on F-COF skeletons induce global skeleton charge redistribution and remodel pore electrostatic potentials to construct push–pull local electrostatic microenvironments. This effect facilitates lithium salt dissociation and releases more free Li+. Meanwhile, F-COF induces PVDF-HFP phase transition and constructs low-energy-barrier ion transport pathways, thus enhancing Li+ transport dynamics. In addition, F-COF donates electrons to TFSI− through anion-π interactions, facilitating the breakage of C–F bonds to generate dense and stable LiF-rich SEI, which effectively homogenizes Li+ flux. As a result, both Li dendrites and interfacial parasitic reactions undergo effective suppression. The electrolyte's ionic conductivity reaches as high as 1.34 × 10−3 S cm−1. Lithium symmetric cells maintain stable operation for more than 4400 h. LFP and LCO full cells show superior cycling stability. This work provides theoretical guidance for the molecular-level design of fluorinated COF materials.

Experimental section

Detailed experimental methods are provided in the SI.

Author contributions

Panpan Mao conceived the research idea, carried out material preparation, experimental characterizations and data analysis, and drafted the original manuscript. Qi An provided guidance on experimental testing and manuscript writing. Xin Wang and Nan Tang took part in the fabrication of solid polymer electrolytes. Cuiping Luo and Wenwei Li participated in experimental measurements. Fanyu Xie performed in situ FTIR characterization. Yufeng Fan completed XRD testing. Hongjia Liu was responsible for battery assembly. Guiquan Zhao conducted in situ mass spectrometry measurements. Yunchun Zha carried out Raman characterization. Genfu Zhao contributed to the computational calculations of this work. Qing Liu, Xiaomin Yang and Hong Guo participated in manuscript writing. Hong Guo also contributed to funding acquisition and conceptualization.

Conflicts of interest

There are no conflicts to declare.

Supplementary Material

SC-OLF-D6SC05957K-s001

Acknowledgments

The authors would like to acknowledge financial support provided by National Natural Science Foundation of China (52372232), Yunnan Fundamental Research Project (202601BC070002), Undergraduate Innovation Training Program (202410673001), the Major Science and Technology Projects of Yunnan Province (202403AA080019-3), Science and Technology Project for Universities of Yunnan Province to Serve Key Inaustries-Doctoral Students Project for Cultivating Industrial Scientific Research Innovatior (FWCY-BSPY2024030) and the Laboratory of Solid-State Ions for Green Energy of Yunnan University, and the Electron Microscope Center of Yunnan University for the support of this work. The authors extend their gratitude to Ms. Fu Zhen (from Scientific Compass https://www.shiyanjia.com) for providing invaluable assistance with the XPS analysis.

Data availability

All the data supporting this article have been included in the main text and the supplementary information (SI). Supplementary information is available. See DOI: https://doi.org/10.1039/d6sc05957k.

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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-D6SC05957K-s001
SC-OLF-D6SC05957K-s001.pdf (2.8MB, pdf)

Data Availability Statement

All the data supporting this article have been included in the main text and the supplementary information (SI). Supplementary information is available. See DOI: https://doi.org/10.1039/d6sc05957k.

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