Construction of an ultrathin multi-functional polymer electrolyte for safe and stable all-solid-state batteries

Abstract

The ever-increasing demand for safe and high-energy-density batteries urges the exploration of ultrathin, lightweight solid electrolytes with high ionic conductivity. Solid polymer electrolytes (SPEs) with high flexibility, reduced interfacial resistance and excellent processability have been attracting significant attentions. However, reducing the thickness of SPEs to be comparable with that of commercial separators increases the risk of short-circuiting. Herein, an ultrathin (≈7 μm), flexible and mechanical robust SPE was constructed from a rationally designed multi-functional polymer network, i.e., poly[2,2,2-trifluoroethyl methacrylate-r-(2-ethylhexyl acrylate)-r-methyl methacrylate-r-1,4-bis(acryloyloxy)butane] (PTEM) and porous polyethylene (PE). The resultant PTEM@PE electrolyte possesses a high tensile strength of 128.0 MPa with extensibility up to 34.8%, which could effectively prevent short-circuiting and minimize the interfacial resistance of cells. The obtained all-solid-state Li|PTEM@PE|LiFePO₄ cell exhibited stable cycling performance over 1500 cycles at 0.5 C with a capacity retention of 74.4%. With high-voltage NCM811 as the cathode, the cell fabricated with PTEM@PE showed a remarkable capacity retention of 84.2% over 500 cycles. Even with the high-mass loading (≈3 mA h cm⁻²) NCM811 cathode, the cell could be operated at ambient temperature, demonstrating superior ion-migration kinetics. The current design provides a promising strategy to develop ultrathin and multifunctional solid electrolytes for safe, long-cycling and high-energy-density all-solid-state batteries.

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New concepts

Given the great demand for safe energy storage devices with high energy density, it is crucial to rationally design polymer electrolytes with integrated desired properties that enable the stable operation of all-solid-state lithium metal batteries. However, poor lithium-ion (Li⁺) migration kinetics inhibit their large-scale applications. To date, most research has been limited to enhancing the ionic conductivity of polymer electrolytes, including copolymer molecular structure design, lithium salt exploitation and additive engineering. In this manuscript, in contrast to conventional approaches, we demonstrate an ultrathin, flexible and mechanically robust solid polymer electrolyte (SPE) based on a rationally designed multi-functional polymer network. By greatly reducing the thickness of the polymer electrolyte, the Li⁺ diffusion distance and time were efficiently shortened. The constructed SPE possesses superior ionic conductance, wide electrochemical window and high mechanical strength, enabling the operation of all-solid-state cells at ambient temperature. The current design of ultrathin multi-functional SPE construction provides guidance on breaking the bottleneck of poor ionic conductivity of SPEs towards building rechargeable batteries with high energy density and safety features.

Introduction

With the rapid development of electric vehicles and portable devices, currently dominant energy storage devices, i.e., lithium (Li)-ion batteries (LIBs) based on graphite anodes, suffer from unsatisfactory energy density.^1–4 Given its extremely high theoretical specific capacity (3860 mA h g⁻¹), low density (0.59 g cm⁻³), and lowest redox potential (−3.04 V vs. H⁺/H₂), metallic Li has been attracting great attention as a replacement for conventional graphite anodes in LIBs.^4–8 Unfortunately, coupling highly reactive Li with conventional liquid electrolytes causes side reactions and safety concerns. The resulting Li dendrites can penetrate the separator, leading to short-circuiting and even thermal runaway of batteries.^7–10 Solid-state electrolytes (SSEs) can suppress side reactions and mechanically block Li dendrite penetration, thus significantly improving safety. Therefore, replacing combustible liquid electrolytes with SSEs to construct all-solid-state Li metal batteries is considered a promising solution to address the dilemma between high energy density and safety.^11–14

Generally, an ideal SSE should simultaneously possess high ionic conductivity, superior mechanical robustness, excellent processability and low electrode/electrolyte interface resistance.^15–18 Up to now, numerous studies on SSEs have reported significant achievements, including improving ionic conductivity, decreasing interfacial resistance and enhancing mechanical strength.^19–24 However, the electrolyte thickness, which is also significant for increasing energy density and improving electrochemical performance, has surprisingly received less attention. According to the computational studies by Sun and co-workers,¹⁹ the energy density of a pouch cell assembled with Li₁₀GePS₁₂ and a Li-rich cathode (mass loading of 4 mA h cm⁻², N/P ratio at 2) was only ≈100 W h kg⁻¹ when the SSE was 500-μm thick, while it achieved ≈300 W h kg⁻¹ on reducing SSE thickness to 100 μm. Regarding the electrochemical performance, according to the equation t = l²/D, the Li⁺ diffusion time (t) is proportional to the square of electrolyte thickness (l), and D is the diffusion constant. A thin electrolyte will shorten the ion diffusion distance/time, alleviate Ohmic resistance, and afford enhanced electrochemical performance.²⁵ Therefore, it is of great significance to develop ultra-thin SSEs for all-solid-state batteries with high energy density and superior electrochemical performance.

SSEs can be classified into inorganic solid electrolytes (ISEs) and solid polymer electrolytes (SPEs).^26–28 ISEs have high ionic conductivity, mechanical robustness and high cationic transference numbers (close to 1). Nevertheless, to preserve their intact structure during battery manufacturing and prevent potential short-circuiting caused by the defects in ISEs, thick ISEs are always utilized (>200 μm), which dramatically impedes the energy density.^29–33 In contrast, SPEs with high flexibility,³⁴ lightweight, low interfacial resistance and excellent processability can be easily manufactured with low thickness. However, the decreased thickness of SPEs inevitably deteriorates the overall mechanical robustness and interfacial compatibility with electrodes, leading to a high risk of short-circuiting and unsatisfactory cycling performance, respectively.^35–41

Herein, an ultrathin SPE (≈7 μm) was constructed using a rationally designed multi-functional polymer network of poly[2,2,2-trifluoroethyl methacrylate-r-(2-ethylhexyl acrylate)-r-methyl methacrylate-r-1,4-bis(acryloyloxy)butane] (PTEM, as shown in Scheme 1(a)) and porous polyethylene (PE). Due to the ultrathin thickness (efficiently shortened Li⁺ transport distance and time), multi-functional polymer network (high ionic conductivity, superior electrochemical stability and spontaneous adhesion) and porous matrix (mechanical robustness), the resultant SPE possessed excellent ionic conductance of 165 mS at 25 °C, a wide electrochemical stability window of 4.7 V, and high tensile strength of 128.0 MPa with elongation up to 34.8%. The assembled Li|PTEM@PE|LiFePO₄ (LFP) cell exhibited outstanding long-term cycling performance over 1500 cycles with a capacity retention of 74.4%. With high-voltage LiNi_0.8Mn_0.1Co_0.1O₂ (NCM811) as the cathode, excellent cycling stability with capability retention of 84.2% over 500 cycles could be achieved. To achieve high energy density, a Li|PTEM@PE|NCM811 cell with an areal capacity of 3.0 mA h cm⁻² was also assembled, which delivered a high discharge specific capacity of ∼170.0 mA h g⁻¹ at ambient conditions. The current design of ultrathin SSE with high-voltage stability, good transport kinetics and superior interface compatibility provides a promising strategy for constructing next-generation rechargeable batteries with high energy density and safety.

Results and discussion

Synthesis and characterization of PTEM@PE

Typically, to assemble safe high-energy-density Li-metal batteries, polymers with high ionic conductivity, superior electrochemical stability and high mechanical strength are preferred as SPEs.^42,43 To avoid the trade-off between ionic conductivity and mechanical strength of the polymer electrolyte, a fibre-reinforced quasi-solid polymer electrolyte has been reported.⁴² By decreasing electrolyte thickness and thereby improving the energy density, we constructed a 5-μm-thick ultra-thin polymer electrolyte film earlier,⁴⁴ but such electrolytes cannot be coupled with high-voltage cathodes. In this work, for practical battery applications that simultaneously demand higher energy density and safety, a rationally designed multi-functional polymer network PTEM was chemically synthesized via radical copolymerization of 2,2,2-trifluoroethyl methacrylate (TFEMA), 2-ethylhexyl acrylate (2-EHA) and methyl methacrylate (MMA) using 1,4-bis(acryloyloxy)butane (BAOB) as the crosslinker, as shown in Scheme 1(a). TFEMA can enhance oxidation resistance and hence afford high voltage stability due to its abundant fluorine atoms.⁴⁵ 2-EHA, which is widely employed in the fabrication of pressure-sensitive adhesives,⁴⁶ was utilized to form intimate interfacial contacts with the electrodes. Moreover, during cathode fabrication, succinonitrile (SN) was also added to the cathode slurry to suppress the interfacial side reactions due to the shielding effect after combining with the active sites, i.e., nickel/cobalt/manganese ions. Moreover, SN can promote intimate interfacial contact between the electrodes and electrolyte, which would endow the cell with superior electrochemical performance. According to Li and co-workers, poly-MMA can assist in Li deposition during the repeated charging/discharging process,⁴⁷ Therefore, MMA monomers were also used for the fabrication of the polymer network with homogenous Li deposition. By casting the polymer solution into a Teflon dish and then removing the solvent, a freestanding, optically clear, flexible PTEM film was obtained, as displayed in Fig. S1 (ESI†).

Scheme 1 (a) Synthesis of the PTEM multi-functional network. (b) Schematic of the fabrication process of the PTEM@PE electrolyte.

The polymer structure of the as-prepared multi-functional PTEM was confirmed by the ¹H NMR and FT-IR spectra. The peaks at 2950 and 1720 cm⁻¹ could be assigned to the stretching vibration of the C–H units and –C=O bonds, respectively (Fig. 1(a)). Compared with the FT-IR spectra of the monomers, the absence of C=C signals at 1630 cm⁻¹ in the spectrum of PTEM suggests successful polymerization, which was also confirmed by the ¹H NMR spectra. Characteristic peaks between 5–6.5 ppm corresponding to the C=C units of TFEMA, 2-EHA and MMA could be observed clearly, while these peaks were invisible in PTEM spectrum (Fig. 1(b) and Fig. S2, ESI†).⁴² The differential scanning calorimetry (DSC) results (Fig. 1(c)) identified a low glass transportation temperature (T_g) of the polymer electrolyte, i.e., −72.6 °C, contributing to fast segmental dynamics and thus high ionic conductivity. The rheological behaviour of the copolymers was evaluated by oscillatory frequency sweeping, and the master curves of storage moduli (G′) and loss moduli (G′′) were obtained based on the time-temperature superposition during frequency sweeping at different temperatures, as displayed in Fig. 1(d). At high temperatures (low frequency), PTEM exhibited liquid-like behavior with G′ < G′′. With decreasing temperature (increasing frequency), G′ approached and surpassed G′′, demonstrating the transition of PTEM from a liquid-like state to a mechanically robust solid state.^42,48

Fig. 1 (a) FT-IR spectra of TFEMA, 2-EHA, MMA monomers and resulting PTEM. (b) ¹H NMR spectra of the prepared PTEM. (c) DSC curve of the PTEM polymer. (d) Frequency sweep of storage moduli (G′) and loss moduli (G′′) of PTEM based on time-temperature superposition at a reference temperature of 25 °C. (e) Stress–strain curves of PTEM, PE and PTEM@PE. (f) Interfacial adhesion test of Cu foils with PTEM as the binder (the inset is the image of the 180° peeling test). (g) Cross-sectional and (h) top-surface SEM images of the PE film. (i) Cross-sectional and (j) top-surface SEM images of PTEM@PE.

The ultra-thin, robust and flexible PTEM@PE electrolyte was prepared by allowing PTEM and lithium bis(trifluoromethane)sulfonimide (LiTFSI) to infiltrate the porous PE films. As shown in Fig. 1(g)–(j), the thickness of the porous structured PE film identified from the cross-sectional SEM image was only 4 μm, and after PTEM and LiTFSI infiltration, the thickness of PTEM@PE was slightly increased to around 7 μm, which is much thinner than most of the solid electrolytes reported previously.^19,49 As shown in Fig. S3 (ESI†), PTEM@PE demonstrated good thermal stability with a degradation onset temperature of ≈300 °C, as revealed by simultaneous thermal analysis in an N₂ atmosphere. As illustrated by the stress–strain curves of PE, PTEM and PTEM@PE in Fig. 1(e), PTEM had a fracture strain of 870% with an ultimate tensile strength of only 1.05 MPa, while the porous PE film presented an ultimate tensile strength of 168.3 MPa. After loading PTEM into the porous PE film, the constructed PTEM@PE film delivered a decent tensile strength of ≈128.0 MPa along with a relatively high extensibility of 34.8%. Seamless contact between the electrolyte and electrode is always preferred as it can dramatically decrease the interfacial resistance.⁵⁰ As illustrated in Fig. 1(f), in the 180° peeling test, the as-prepared polymer electrolyte delivered strong adhesion with a peeling force of 61.4 N m⁻¹ (adhesion energy of 61.4 J m⁻²), which is much higher than the baseline for an intimate interface, i.e., >5 J m⁻².⁵¹

Electrochemical performance of PTEM@PE

Electrostatic potential (ESP) calculations were performed to investigate the coordination of PTEM with Li⁺, which is closely related to non-covalent interactions between the macromolecules and Li⁺.⁵² Herein, different atoms exhibited a variety of electrostatic potentials (Fig. 2(a)), promoting the coordination of electronegative and electropositive atoms. Especially, the iso-potential surface of PTEM exhibited a strong electron density in the C=O group (red colour), which could form nucleophilic coordination with positive Li⁺ (Fig. S4, ESI†). To investigate the effect of electrolyte thickness on ion-transport kinetics and thus the electrochemical performance, PTEM coupled with a 22 μm-thick PE skeleton was also prepared (PTEM@PE-22 in Fig. S5, ESI†). To confirm PTEM loading in the electrolytes with different thicknesses, the PTEM loading was calculated by measuring the weight of the stainless-steel symmetric cells before and after PTEM loading. As shown in Table S1 (ESI†), PTEM loading in PTEM@PE was 2.13 wt%, which increased to around 3.40 wt% in PTEM@PE-22. Although PTEM@PE-22 had more porous channels to uptake the electrolyte, the thinner PTEM@PE film could efficiently shorten the Li⁺ diffusion distance, resulting in dramatically improved ionic conductance. The ionic conductivity of the prepared SPEs were tested by electrochemical impedance spectroscopy (EIS) measurements in assembled cells at temperatures ranging from 25 to 60 °C (Fig. S6, ESI†). As displayed in Fig. 2(b), PTEM@PE (if not specified, PTEM@PE refers to the electrolyte with a thickness of ≈7 μm) exhibited an ionic conductivity of 5.3 × 10⁻⁵ S cm⁻¹ at 25 °C, which was higher than that of PTEM@PE-22 (4.3× 10⁻⁵ S cm⁻¹), suggesting accelerated ion transport dynamics in the ultrathin electrolyte.

Fig. 2 (a) ESP of PTEM. (b) Vogel–Fulcher–Tammann (VFT) fitting results of the ionic conductivity of PTEM@PE and PTEM@PE-22 electrolytes. (c) Temperature-dependent ionic conductance of PTEM@PE and PTEM@PE-22 electrolytes. (d) Impedance spectra of the Li|PTEM@PE|Li and Li|PTEM@PE-22|Li cells after different storage times. (e) LSV plot of PTEM@PE at 2.0–5.2 V and 25 °C. (f) DFT analysis of the HOMO and LUMO of different polymer units. (g) Cycling performance of a symmetrical cell at 0.1 mA h cm⁻²; the insets show the voltage profiles of the Li|PTEM@PE|Li cell at 900 h.

Notably, the ionic conductivity of 5.3 × 10⁻⁵ S cm⁻¹ is much lower than the commonly set threshold of 10⁻⁴ S cm⁻¹. However, the ionic conductance, which is closely related to the thickness of the electrolyte, can better reflect ion transport in the electrolyte and determine the electrochemical performance of the cell.^{25,33,38,44,53–57} As shown in Fig. 2(c), PTEM@PE showed an ionic conductance of 165 mS at ambient temperature, which is over seven times higher than that of PTEM@PE-22. The activation energy reduced from 0.14 to 0.12 eV with lower film thickness, implying a lower energy barrier for ion transport in PTEM@PE than in PTEM@PE-22.⁵⁸ Moreover, as shown in Fig. 2(d) and Fig. S7 (ESI†), after the rest period of 480 h, the symmetric cell with the PTEM@PE electrolyte had lower interfacial resistance (166 Ω) than that of the cell with PTEM@PE-22 (280 Ω). It is believed that the ultrathin electrolyte facilitates superior interfacial compatibility with the Li anode, contributing to the superior electrochemical performance of the cell.

High electrochemical stability is essential for the application of SPEs in high-energy-density batteries. As shown in Fig. 2(e), the decomposition-onset potential of the PTEM@PE electrolyte was 4.7 V, implying that the PTEM@PE can be used with high-voltage cathodes, such as NCM811. To get deep insights into the oxidation resistance of PTEM@PE, density-functional theory (DFT) calculations were carried out in different monomer and polymer systems, as shown in Fig. 2(f). PEO, MMA and TFEMA showed comparable lowest unoccupied molecular orbital (LUMO) energies. The highest occupied molecular orbital (HOMO) energy for MMA and TFEMA were calculated to be −7.50 eV and −7.75 eV, respectively, which are obviously lower than that of widely studied PEO. This suggests the improved anti-oxidation capability of the MMA- and TFEMA-based polymer network because of the strong electron-withdrawing effect of the fluorine atoms. Polymer electrolytes with a high cationic transport number exhibit great potential in reducing the buildup of ion concentration gradients.^43,59,60 As shown in Fig. S8 (ESI†), the Li⁺ transport number (t_Li⁺) for PTEM@PE obtained by the Bruce–Vincent method was 0.57, which can reduce ion polarization and effectively suppress dendritic Li growth according to numerous computational studies.^61,62 Moreover, the critical current density was measured to evaluate the capability of PTEM@PE to suppress Li dendrite growth (Fig. S9, ESI†). Due to elevated Li⁺ mobility and enhanced interfacial stability, PTEM@PE showed a critical current density of 0.5 mA cm⁻². With the repeated Li plating/stripping process, as illustrated in Fig. 2(g), the Li|PTEM@PE|Li symmetric cell showed superior cycling performance over 1600 h at a current density of 0.1 mA cm⁻² (polarization voltage lower than 41 mV). Moreover, at the current density of 0.2 mA cm⁻², the cell showed good cycling stability over 400 h, demonstrating the superior compatibility of the as-prepared SPE with the highly reactive Li electrode (Fig. S10, ESI†).

Full-cell evaluation of PTEM@PE

To evaluate the electrochemical performance of PTEM@PE in all-solid-state batteries, the Li|PTEM@PE|LFP full-cell was assembled. As displayed in Fig. 3(a) and Fig. S11 (ESI†), at a current density of 0.5C and 25 °C, the Li|PTEM@PE|LFP cell delivered stable cycling performance over 1500 cycles with capacity retention of 74.4% and an average Coulombic efficiency (CE) of 99.5%. Notably, CEs over 100% could be observed in some cycles. This can be explained by incomplete delithiation and excess lithiation of the cathode during the activation process, as well as possible reactions during the charging/discharging process, which may lead to such phenomena.⁶³ At a current density of 1C, the Li|PTEM@PE|LFP cell delivered outstanding cycling stability with capacity retention of 76.6% over 1000 cycles (Fig. S12, ESI†). Moreover, the rate performance also indicated that even at a current density of 2C, the cell could deliver a high discharge capacity of ≈136 mA h g⁻¹ (Fig. S13, ESI†). When reverted to 0.5C, good capacity reversibility was observed, indicating the superior electrochemical stability of the electrolyte at high current densities.³⁸ The superior rate capability can be explained by the shortened ion transport distance/time in the ultrathin electrolyte, and the rationally designed polymer network leads to good interfacial compatibility.

Fig. 3 (a) Cycling performance of the Li|PTEM@PE|LFP full-cell at a current density of 0.5C. (b) Cycling performance of the Li|PTEM@PE|LFP pouch cell at a current density of 0.5C (current collector with 25 mm Li). (c) Cycling performance of the Li|PTEM@PE|LFP cell with industrial-standard LFP as the cathode (96.8% active material and areal capability = 2.5 mA h cm⁻²). (d) Schematic of the SPE in the cell. (e) Top-surface and (f) cross-sectional SEM images of the cycled Li electrode using the PTEM@PE electrolyte. (g) Top-surface and (h) cross-sectional SEM images of the cycled Li electrode using the MMA@PE electrolyte.

One advantage of all-solid-state batteries based on SPEs is their mechanical flexibility, which allows their application in flexible devices with improved safety. As shown in Fig. 3(b), a Li|PTEM@PE|LFP 5.3 mAh pouch cell was assembled with a 25-μm Li anode (detailed parameters in Table S2, ESI†). At a current density of 0.5C and 25 °C, the pouch cell was reused for 70 cycles, and after the abuse test (folding), a light-emitting diode (LED) could still be powdered, revealing their superior safety. Evaluating the ion transport dynamics and interfacial resistance of SPEs with high-areal-capacity cathodes is essential for practical applications. To demonstrate the applicability of the as-prepared PTEM@PE solid electrolyte, a full-cell with industrial standard LFP as the cathode (96.8 wt% active material, 2.5 mA h cm⁻² loading based on one side of the active material) was also assembled. As shown in Fig. 3(c), at a current density of 0.2C and 25 °C, the cell exhibited stable cycling performance over 95 cycles with a capacity retention of 90%, which was much better than that of the cell with a liquid electrolyte (27 mA h g⁻¹ with a capacity retention of only 20% after 100 cycles). The fluctuation of discharge specific capacity can be explained by the electrochemical activation of the cathode in the initial cycles and the loss of cyclable lithium in the subsequent cycles.⁶⁴ The cyclic voltammetry (CV) curves of Li|PTEM@PE|LFP exhibited two distinct peaks located at 3.64 and 3.22 V (Fig. S14, ESI†). After the first cycle, the two peaks shifted to 3.62 and 3.24 V, respectively, probably due to the formation of an SEI layer. With increasing cycle numbers, the peak intensities showed only a slight change, confirming stable discharge/charge efficiency and electrochemical activity.

The excellent electrochemical performance of the cell based on the PTEM@PE electrolyte can be explained by the following reasons. Firstly, with reduced thickness and superior interfacial compatibility of the solid electrolyte, the Li⁺ diffusion distance and time can be dramatically shortened, providing sufficient ion conductance in the cell operated at room temperature. As shown in Fig. 2(c), the ultra-thin PTEM@PE electrolyte exhibited an ionic conductance of 165 mS with an activation energy of 0.12 eV, suggesting accelerated ion transport dynamics. Moreover, as illustrated by the electrolyte/electrode interfacial resistance measured in Fig. S15 (ESI†), the cell exhibited gradually increasing interfacial resistance in the initial 20 cycles followed by a decrease in resistance, suggesting the construction of stable ionic conducting channels. Secondly, the rationally designed polymer network with unique functional groups enables homogeneous current distribution in PTEM@PE, facilitating the uniform deposition of Li⁺. As illustrated by the surface and cross-sectional images of the Li electrode from the cycled Li|PTEM@PE|LFP cell in Fig. 3(e) and (f), with PTEM@PE as the solid electrolyte, a dense and smooth surface was observed, implying uniform ion deposition and suppressed Li dendrite formation. In contrast, with MMA@PE as the solid electrolyte (prepared via the same method as PTEM@PE), fractured and porous structured Li was observed (Fig. 3(g) and (h)). To further confirm the role of the as-prepared SPE in battery performance, XPS was utilized to analyze the cycled Li electrode from the Li|PTEM@PE|LFP cell, as shown in Fig. S16(a)–(d) (ESI†). In the C1s spectra, the characteristic peaks assigned to C=O (287.5 eV) and CF₃ (292.9 eV) were observed, corresponding to the PTEM chain and LiTFSI from the SPE. As described in the F1s spectrum, the peak assigned to LiF (689.1 eV) agrees well with the peak detected in the Li1s spectrum, and LiF has good electronic insulation and high shear modulus, contributing to the formation of a stable SEI layer.

To build safe high-energy-density batteries, the SPEs should possess high electrochemical stability to be used in all-solid-state batteries with high-voltage cathodes, such as NCM811. Considering the decomposition onset of the PTEM@PE electrolyte at 4.7 V, as confirmed by Fig. 2(e) and (f), an Li|PTEM@PE|NCM811 cell was also assembled. As shown in Fig. 4(a), at a current density of 0.5C and 25 °C, the Li|PTEM@PE|NCM811 cell delivered outstanding cycling stability with capacity retention of 84.2% over 500 cycles. Moreover, even with a mass loading of 1.2 mA h cm⁻², the Li|PTEM@PE|NCM811 cell could exhibit an initial discharge capacity of 206 mA h g⁻¹ with capacity retention of 78.8% over 100 cycles (Fig. 4(c) and (d)). As shown in Fig. S17 (ESI†), the CV curves of the Li|PTEM@PE|NCM811 cell were collected at 0.1 mV s⁻¹. The CV curves exhibited three pairs of anodic and cathodic peaks, corresponding to the transitions of the cathode phases. The potential of the anodic peaks decreased slightly from the 1st cycle to the 2nd, which is mainly due to irreversible reactions. The comparable potentials of the 2nd and 3rd cycles suggest that the cathode was stable after the 1st charging/discharging process.⁶⁵ Even with a high mass loading of 3.0 mA h cm⁻², the cells were operational with an initial discharge capacity of 167.9 mA h g⁻¹ (Fig. S18, ESI†). The rate capability of the Li|PTEM@PE|NCM811 cell was also evaluated. As displayed in Fig. 4(e), at a current density of 0.2C, the cell delivered a high discharge capacity of 196 mA h g⁻¹, and even at a high current density of 2C, a high discharge capacity of ≈100 mA h g⁻¹ was achieved, implying fast Li⁺ conductance and superior interface compatibility. Notably, the Li|PTEM@PE|NCM811 full cell exhibited superior electrochemical performance with a long cycle life and relatively high areal capacity compared to previous reports,^{41,44,53–60} as summarized in Fig. 4(f) and Table S3 (ESI†).

Fig. 4 (a) Cycling performance of the Li|PTEM@PE|NCM811 cell at a current density of 0.5C. (b) Schematic of the SPE in the cell. (c) Cycling performance of the NCM811 full-cell with a cathode areal capacity of 1.2 mA h cm⁻². (d) Charge–discharge profiles of the NCM811 full-cell with a cathode-areal capacity of 1.2 mA h cm⁻² at different cycles. (e) Rate capability of the Li|PTEM@PE|NCM811 cell. (f) Comparison of the areal capability and cycle number of solid-state Li batteries with typical polymer- or composite-based electrolytes.

Apart from the reduced thickness and rationally designed polymer network, to further investigate the enhanced electrochemical performance of the cell based on the PTEM@PE electrolyte, the relationship between the electrolyte and degradation of the cathode was also studied (Fig. 5(a)). Till now, three degradation mechanisms have been identified for the NCM cathode: the layered-to-spinel/rock salt phase transformation, the side reactions between the cathode and electrolyte, and the corrosion/dissolution of the cathode materials in liquid electrolytes.^66–70 In this study, the soft and flexible polymer could form a homogeneous and compatible surface with low charge transfer resistance, accommodating drastic volume changes of the cathode and inhibiting the side reactions. As shown in Fig. 5(b) and (c), the cathode particles could maintain structure integrity after repeated charging/discharging cycles, consistent with its high cycling stability. On the contrary, with MMA@PE as the electrode, obvious cracks could be observed in the cathode particles (Fig. 5(d) and (e)). To further confirm the role of the coating layer in improving the cycling performance, the chemical compositions of the pristine and cycled NCM811 cathode particles were analyzed by XPS measurements, as shown in Fig. 5(f)–(i). In the C1s spectra, C–O (286.5 eV) and C–C/C–H (284.8 eV) could be detected on both pristine and cycled cathode surfaces. In the cycled NCM electrodes, new characteristic peaks assigned to CF₃ (292.9 eV, contributed by LiTFSI) and the O–C=O group (288.8 eV, contributed by PTEM) were observed, suggesting the formation of a coating layer on the cathode surface.^71–74 In the O1s spectra, the lattice O (529.5 eV) peak was only present in the pristine NCM electrode. With increasing cycle numbers, the disappearance of lattice O peak indicates that the NCM811 particles were covered by the cathode electrolyte interphase (CEI) layer. During the cycling process, the formed flexible and elastic CEI layer can inhibit the side reactions and stabilize the NCM cathode, contributing to stable cycling performance.

Fig. 5 (a) Schematic of the detailed mechanism of SPE underlying the enhanced performance of the cell. (b) and (c) SEM images of the cycled NCM811 cathode particles from the Li|PTEM@PE|NCM811 cell. (d) and (e) SEM images of the cycled NCM811 cathode particles from the Li|MMA@PE|NCM811 cell. (f)–(i) C1s and O1s XPS spectra of the pristine and cycled NCM811 cathodes.

Conclusions

In summary, a rationally designed polymer network was synthesized in this work from three types of monomers and one crosslinker, with each of them furnishing different functionalities, such as oxidation resistance, intimate interface contact, efficient ionic conducting and mechanical robustness. Derived from the multi-functional and porous structured PE film, the ultrathin (≈7 μm) and flexible solid polymer electrolyte (SPE) PTEM@PE exhibits superior ionic conductance (165 mS at 25 °C), high electrochemical stability (≈4.7 V) and high tensile strength (128.0 MPa). The assembled Li|PTEM@PE|LFP cell displayed excellent cycling performance over 1500 cycles with a capacity retention of 74.4%. With the high-voltage cathode NCM811, the Li|PTEM@PE|NCM811 cell exhibited an initial discharge capacity of 178.6 mA h g⁻¹ with capacity retention of 84.2% over 500 cycles. Even with a high-mass-loading NCM811 cathode (3 mA h cm⁻²), the cells could still operate at room temperature. The detailed mechanisms underlying the electrochemical performance enhancement were thoroughly investigated; the results show that the ultra-thin PTEM@PE electrolyte with abundant functional groups can effectively shorten the ion diffusion distance/time, facilitate uniform deposition of Li⁺ and inhibit the degradation of the cathode. The design principle demonstrated here provides guiding principles for designing high-performance SPEs with functional polymers towards achieving Li metal batteries with both high energy density and safety features.

Author contributions

S. Gao, P.-F. Cao and H. Yang conceived the research and designed the experiments. Y. Zhang, T. Cheng, H. Ding, Z. Li, L. Li and D. Yang performed the material fabrication and characterization. H. Yang, S. Gao and P.-F. Cao revised the grammar in use and property characterizations. The manuscript was written through contributions of all authors. All authors have given approval to the final version of the manuscript.

Data availability

The data supporting this article have been included as part of the ESI.†

Conflicts of interest

There are no conflicts to declare.

Acknowledgements

This work was financially supported by the National Natural Science Foundation of China (22379073, 52373275), the Natural Science Foundation of Tianjin, China (18JCZDJC31400), the MOE Innovation Team (IRT13022). P.-F. Cao acknowledges financial support by Fundamental Research Funds for the Central Universities (buctrc202222).

Notes and references

1. Y. Guo, Q. Wu, L. Liu, G. Li, L. Yang, X. Wang, Y. Ma and Z. Hu, Adv. Sci., 2022, 9, 2200411.
2. J. Tao, D. Wang, Y. Yang, J. Li, Z. Huang, S. Mathur, Z. Hong and Y. Lin, Adv. Sci., 2021, 9, 2103786.
3. Z. Ju, J. Nai, Y. Wang, T. Liu, J. Zheng, H. Yuan, O. Sheng, C. Jin, W. Zhang, Z. Jin, H. Tian, Y. Liu and X. Tao, Nat. Commun., 2020, 11, 488–497.
4. S. Gao, Z. Li, N. Liu, G. Liu, H. Yang and P. F. Cao, Adv. Funct. Mater., 2022, 32, 2202013.
5. P. Albertus, V. Anandan, C. Ban, N. Balsara, I. Belharouak, J. Buettner-Garrett, Z. Chen, C. Daniel, M. Doeff, N. J. Dudney, B. Dunn, S. J. Harris, S. Herle, E. Herbert, S. Kalnaus, J. A. Libera, D. Lu, S. Martin, B. D. McCloskey, M. T. McDowell, Y. S. Meng, J. Nanda, J. Sakamoto, E. C. Self, S. Tepavcevic, E. Wachsman, C. Wang, A. S. Westover, J. Xiao and T. Yersak, ACS Energy Lett., 2021, 6, 1399–1404.
6. L. Zhou, T.-T. Zuo, C. Y. Kwok, S. Y. Kim, A. Assoud, Q. Zhang, J. Janek and L. F. Nazar, Nat. Energy, 2022, 7, 83–93.
7. S. L. Gao, A. Cannon, F. Y. Sun, Y. Y. Pan, D. D. Yang, S. R. Ge, N. Liu, A. P. Sokolov, E. Ryan, H. B. Yang and P. F. Cao, Cell Rep. Phys. Sci., 2021, 2, 100534.
8. S. Gao, F. Sun, N. Liu, H. Yang and P.-F. Cao, Mater. Today, 2020, 40, 140–159.
9. T. Zizheng, W. Shu-Bo, L. Yu-Kai, H. Shu-Fen and L. Ru-Shi, ACS Appl. Mater. Interfaces, 2020, 12, 47181.
10. S. Liu, W. Liu, D. Ba, Y. Zhao, Y. Ye, Y. Li and J. Liu, Adv. Mater., 2023, 35, e2110423.
11. J.-H. Kim, J. Sun, J.-S. Jang, D.-H. Park, S.-Y. Ahn, W.-C. Kim, K. Min and K.-W. Park, Energy Storage Mater., 2024, 64, 103080.
12. Z. Li, L. Sun, L. Zhai, K. Oh, J. Seo, C. Li, D. Han, J. Baek and S. Lee, Angew. Chem., Int. Ed., 2023, 62, e202307459.
13. C. Zhifu, P. Quan, A. Zhitao, T. Yiting, Z. Qingfeng and X. Shuhong, Chem. Commun., 2023, 59, 10416.
14. D. Zhang, X. Meng, W. Hou, W. Hu, J. Mo, T. Yang, W. Zhang, Q. Fan, L. Liu, B. Jiang, L. Chu and M. Li, Nano Res. Energy, 2023, 2, e9120050.
15. W.-Q. Ding, F. Lv, N. Xu, M.-T. Wu, J. Liu and X.-P. Gao, ACS Appl. Energy Mater., 2021, 4, 4581–4601.
16. D. Wu, L. Chen, H. Li and F. Wu, Appl. Phys. Lett., 2022, 121, 120502.
17. H. J. Yoon, J. W. Cho, H. T. Kang, M. Lee and R. G. Yi, Macromol. Res., 2021, 29, 509–518.
18. N. Meng, F. Lian and G. Cui, Small, 2021, 17, e2005762.
19. X. Yang, K. R. Adair, X. Gao and X. Sun, Energy Environ. Sci., 2021, 14, 643–671.
20. Y. Yuan, B. Wang, K. Xue, Y. Ma, X. Liu, X. Peng, M. Liu and H. Lu, ACS Appl. Mater. Interfaces, 2023, 15, 17144.
21. S. Qi, M. Li, Y. Gao, W. Zhang, S. Liu, J. Zhao and L. Du, Adv. Mater., 2023, 35, e2304951.
22. M. Wu, J. Song, J. Lei and H. Tang, Nano Res., 2023, 17, 1482.
23. Y. He, C. Wang, P. Zou, R. Lin, E. Hu and H. L. Xin, Angew. Chem., Int. Ed., 2023, 62, e202308309.
24. P.-M. Cristina, S. Preston, M. Isabella, K. Aristotelis, M. Worarin, W. Christoph, S. Ullrich and G. Ilja, J. Mater. Chem. A, 2023, 11, 9521.
25. J. Y. Wu, Z. X. Rao, Z. X. Cheng, L. X. Yuan, Z. Li and Y. H. Huang, Adv. Energy Mater., 2019, 9, 1902767.
26. J. Zheng, C. Sun, Z. Wang, S. Liu, B. An, Z. Sun and F. Li, Angew. Chem., Int. Ed., 2021, 60, 18448–18453.
27. Q. Zhao, S. Stalin, C.-Z. Zhao and L. A. Archer, Nat. Rev. Mater., 2020, 5, 229–252.
28. Y. Lu, C.-Z. Zhao, H. Yuan, J.-K. Hu, J.-Q. Huang and Q. Zhang, Matter, 2022, 5, 876–898.
29. M. B. Dixit, N. Singh, J. P. Horwath, P. D. Shevchenko, M. Jones, E. A. Stach, T. S. Arthur and K. B. Hatzell, Matter, 2020, 3, 2138–2159.
30. J. Su, M. Pasta, Z. Ning, X. Gao, P. G. Bruce and C. R. M. Grovenor, Energy Environ. Sci., 2022, 15, 3805–3814.
31. Z. Ning, D. S. Jolly, G. Li, R. De Meyere, S. D. Pu, Y. Chen, J. Kasemchainan, J. Ihli, C. Gong, B. Liu, D. L. R. Melvin, A. Bonnin, O. Magdysyuk, P. Adamson, G. O. Hartley, C. W. Monroe, T. J. Marrow and P. G. Bruce, Nat. Mater., 2021, 20, 1121–1129.
32. R. J. Y. Park, C. M. Eschler, C. D. Fincher, A. F. Badel, P. Guan, M. Pharr, B. W. Sheldon, W. C. Carter, V. Viswanathan and Y.-M. Chiang, Nat. Energy, 2021, 6, 314–322.
33. J. Wan, J. Xie, X. Kong, Z. Liu, K. Liu, F. Shi, A. Pei, H. Chen, W. Chen, J. Chen, X. Zhang, L. Zong, J. Wang, L. Q. Chen, J. Qin and Y. Cui, Nat. Nanotechnol., 2019, 14, 705–711.
34. S. Chen, K. Zhang, Z. Li, Y. Wu, B. Zhu and J. Zhu, Supramol. Mater., 2023, 2, 100032.
35. Y. Lin, M. Wu, J. Sun, L. Zhang, Q. Jian and T. Zhao, Adv. Energy Mater., 2021, 11, 2101612.
36. X. Lin, J. Yu, M. B. Effat, G. Zhou, M. J. Robson, S. C. T. Kwok, H. Li, S. Zhan, Y. Shang and F. Ciucci, Adv. Funct. Mater., 2021, 31, 2010261.
37. X. Yang, M. Jiang, X. Gao, D. Bao, Q. Sun, N. Holmes, H. Duan, S. Mukherjee, K. Adair, C. Zhao, J. Liang, W. Li, J. Li, Y. Liu, H. Huang, L. Zhang, S. Lu, Q. Lu, R. Li, C. V. Singh and X. Sun, Energy Environ. Sci., 2020, 13, 1318–1325.
38. Z. Wang, L. Shen, S. Deng, P. Cui and X. Yao, Adv. Mater., 2021, 33, e2100353.
39. J. Liu, L. Zhang, Y. Cao, Z. Wang, X. Xia, J. Zhou, X. Shen, X. Zhou, T. Qian and C. Yan, Appl. Phys. Lett., 2022, 121, 153905.
40. A. G. M. Mathieson, W. M. Dose, H.-G. Steinrück, C. J. Takacs, S. Feldmann, R. Pandya, A. J. Merryweather, D. Mackanic, A. Rao, F. Deschler and M. De Volder, Energy Environ. Sci., 2022, 15, 4323–4337.
41. H. Wang, Y. Li, Y. Li, Y. Liu, D. Lin, C. Zhu, G. Chen, A. Yang, K. Yan, H. Chen, Y. Zhu, J. Li, J. Xie, J. Xu, Z. Zhang, R. Vilá, A. Pei, K. Wang and Y. Cui, Nano Lett., 2019, 19, 1326–1335.
42. S. Gao, Y. Zhang, M. Ma, Z. Li, Z. Sun, M. Tian, H. Yang and P.-F. Cao, Mater. Adv., 2023, 4, 3452.
43. S. Gao, M. Zhang, C. Gainaru, A. P. Sokolov, H. B. Yang and P.-F. Cao, Matter, 2022, 5, 2457–2460.
44. S. Gao, M. Ma, Y. Zhang, L. Li, S. Zhu, Y. He, D. Yang, H. Yang and P.-F. Cao, J. Mater. Chem. A, 2024, 12, 9469–9477.
45. L. Tang, B. Chen, Z. Zhang, C. Ma, J. Chen, Y. Huang, F. Zhang, Q. Dong, G. Xue, D. Chen, C. Hu, S. Li, Z. Liu, Y. Shen, Q. Chen and L. Chen, Nat. Commun., 2023, 14, 2301.
46. S. Gao, Y. Pan, B. R. Li, M. A. Rahman, M. Tian, H. B. Yang and P.-F. Cao, Adv. Funct. Mater., 2022, 33, 2210543.
47. Y. Guo, Y. Ouyang, D. Li, Y. Wei, T. Zhai and H. Li, Energy Storage Mater., 2019, 16, 203–211.
48. S. G. Y. Pan, Z. Rashid, S. Gao, A. Erwin, V. Tsukruk, K. D. Vogiatzis, A. P. Sokolov, H. Yang and P.-F. Cao, ACS Appl. Energy Mater., 2020, 3, 3387–3396.
49. J. Y. Wu, L. X. Yuan, W. X. Zhang, Z. Li, X. L. Xie and Y. H. Huang, Energy Environ. Sci., 2021, 14, 12–36.
50. B. L. Watson, N. Rolston, A. D. Printz and R. H. Dauskardt, Energy Environ. Sci., 2017, 10, 2500–2508.
51. M. J. Lee, J. Han, K. Lee, Y. J. Lee, B. G. Kim, K. N. Jung, B. J. Kim and S. W. Lee, Nature, 2022, 601, 217–222.
52. X. Huang, S. Huang, T. Wang, L. Zhong, D. Han, M. Xiao, S. Wang and Y. Meng, Adv. Funct. Mater., 2023, 33, 2300683.
53. D. Zhou, Y.-B. He, R. Liu, M. Liu, H. Du, B. Li, Q. Cai, Q.-H. Yang and F. Kang, Adv. Energy Mater., 2015, 5, 1500353.
54. J. M. Whiteley, P. Taynton, W. Zhang and S.-H. Lee, Adv. Mater., 2015, 27, 6922–6927.
55. Y. X. Ma, J. Y. Wan, Y. F. Yang, Y. S. Ye, X. Xiao, D. T. Boyle, W. Burke, Z. J. Huang, H. Chen, Y. Cui, Z. A. Yu and S. T. Oyakhire, Adv. Mater., 2021, 33, e2105329.
56. F. He, W. Tang, X. Zhang, L. Deng and J. Luo, Adv. Mater., 2021, 33, e2105329.
57. L. Gao, H. R. Liang, J. X. Li, B. W. Cheng, N. P. Deng and W. M. Kang, J. Power Sources, 2021, 515, 230622.
58. Z. Lin, X. Guo and H. Yu, Nano Energy, 2017, 41, 646–653.
59. S. Gao, Z. Li, Z. Zhang, B. Li, X. C. Chen, G. Yang, T. Saito, M. Tian, H. Yang and P.-F. Cao, Energy Storage Mater., 2023, 55, 214–224.
60. S. Zhao, Y. Zhang, H. Pham, J.-M. Y. Carrillo, B. G. Sumpter, J. Nanda, N. J. Dudney, T. Saito, A. P. Sokolov and P.-F. Cao, ACS Appl. Energy Mater., 2020, 3, 12540–12548.
61. H. O. Ford, B. Park, J. Jiang, M. E. Seidler and J. L. Schaefer, ACS Mater. Lett., 2020, 2, 272–279.
62. Z. Chen, D. Steinle, H.-D. Nguyen, J.-K. Kim, A. Mayer, J. Shi, E. Paillard, C. Iojoiu, S. Passerini and D. Bresser, Nano Energy, 2020, 77, 105129.
63. T. Lin, T. Seaby, Y. Hu, S. Ding, Y. Liu, B. Luo and L. Wang, Electrochem. Energy Rev., 2022, 5, 26–86.
64. L. Ding, R. Leones, T. Schmeida, K. Nielsch and D. Mikhailova, J. Power Sources, 2022, 521, 230955.
65. M. Yao, Q. Ruan, S. Pan, H. Zhang and S. Zhang, Adv. Energy Mater., 2023, 13, 2203640.
66. U. Nisar, N. Muralidharan, R. Essehli, R. Amin and I. Belharouak, Energy Storage Mater., 2021, 38, 309.
67. P. Yan, J. Zheng, M. Gu, J. Xiao, J.-G. Zhang and C.-M. Wang, Nat. Commun., 2017, 8, 14101.
68. F. Lin, I. M. Markus, D. Nordlund, T.-C. Weng, M. D. Asta, H. L. Xin and M. M. Doeff, Nat. Commun., 2014, 5, 3529.
69. B. Xu, C. R. Fell, M. Chi and Y. S. Meng, Energy Environ. Sci., 2011, 4, 2223.
70. S. Malunavar, L. Porcarelli, P. C. Howlett, D. Mecerreyes and M. Forsyth, RSC Appl. Polym., 2024, 2, 384–394.
71. H. Lyu, Y. Li, C. J. Jafta, C. A. Bridges, H. M. Meyer, A. Borisevich, M. P. Paranthaman, S. Dai and X.-G. Sun, J. Power Sources, 2019, 412, 527–535.
72. M. Yoon, Y. Dong, J. Hwang, J. Sung, H. Cha, K. Ahn, Y. Huang, S. J. Kang, J. Li and J. Cho, Nat. Energy, 2021, 6, 362–371.
73. S. Jiao, X. Ren, R. Cao, M. H. Engelhard, Y. Liu, D. Hu, D. Mei, J. Zheng, W. Zhao, Q. Li, N. Liu, B. D. Adams, C. Ma, J. Liu, J.-G. Zhang and W. Xu, Nat. Energy, 2018, 3, 739–746.
74. D. Ensling, M. Stjerndahl, A. Nyten, T. Gustafsson and J. O. Thomas, J. Mater. Chem. A, 2009, 19, 82–88.

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Footnote

† Electronic supplementary information (ESI) available. See DOI: https://doi.org/10.1039/d4mh01037j

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