2-hydroxyethyl acrylate (> 98.0%), 2-hydroxyethyl acrylamide (> 98.0%), succinonitrile (SN, 98%), 2,2′-azobis(2-methylpropionitrile) (AIBN), dichloromethane (99.5%, water ≤ 50 ppm), anhydrous dichloromethane (97.0%), pyridine (99.5%), magnesium sulfate (MgSO₄, 98%), hydrochloric acid (37.0%), petroleum ether (99.0%), ethyl acetate (99.0%) and deuterated chloroform (99.8%) were purchased from Energy Chemical Technology (Shanghai) Co., Ltd. Lithium bis(trifluoromethanesulphonyl)imide (LiTFSI) and lithium difluoro(oxalato)borate (LiDFOB) were acquired from Suzhou DodoChem Technology Co., Ltd. Malonylchloride and pimeloylchloride were sourced from Shanghai Macklin Biochemical Co., Ltd. Polyvinylidene difluoride (PVDF), Super P, N-methyl pyrrolidone (NMP) and LiNi_0.8Co_0.1Mn_0.1O₂ (NCM811) were provided by Canrd Technology Co., Ltd. All chemicals were used as received without further purification.

A Nicolet iS50 spectrometer (Thermo Scientific, USA) was utilized to acquire Fourier transform infrared (FTIR) spectra. For surface analysis, a Thermo K-Alpha XPS system (Thermo Scientific, USA) with a monochromatic Al-Kα X-ray source was employed to record X-ray photoelectron spectroscopy (XPS) data. Scanning electron microscopy (SEM, SU8220, Hitachi, Japan) was used to characterize the morphology. Thermal stability was evaluated using a Q50 thermogravimetric analyzer (TA Instruments, USA); TGA measurements were conducted at a heating rate of 5 °C·min^-1 under an argon atmosphere. Bruker AVANCE III HD 400 instrumentation (Bruker, Germany) was used to analyze Nuclear magnetic resonance (NMR) spectra.

All Density Functional Theory (DFT) calculations were performed using the CP2K package^[25]. A hybrid Gaussian and plane-wave (GPW) approach was employed, wherein the Kohn-Sham orbitals were expanded using Gaussian-type basis sets, while the electron density was represented by an auxiliary plane-wave basis. Specifically, molecularly optimized triple-zeta valence doubly polarized (TZV2P-MOLOPT) basis sets^[26] were utilized for all elements (S, N, O, F, Li, C, and H), coupled with a plane-wave energy cutoff of 500 Ry. The exchange-correlation interactions were described by the Perdew-Burke-Ernzerhof (PBE) functional^[27]. To account for weak interactions, Grimme’s D3 dispersion correction was applied throughout^[28]. Wave function optimization was achieved via matrix diagonalization, and Fermi smearing at an electronic temperature of 300 K was used to aid self-consistent field (SCF) convergence. Geometry optimizations were conducted using the Broyden-Fletcher-Goldfarb-Shanno (BFGS) algorithm. Furthermore, the basis-set superposition error (BSSE) was evaluated for all adsorbate systems and found to be negligible (< 0.05 eV)^[27].

The DAH and Am-DAH monomers are synthesized via esterification and amidation reactions, respectively [Supplementary Figure 1]. In terms of molecular structure, DAH contains ester linkages connecting the polymerizable vinyl groups, whereas Am-DAH features amide functionalities. This substitution of oxygen with nitrogen (-CONH- vs. -COO-) endows Am-DAH with the capability to form hydrogen bonding networks. The molecular structures of the synthesized DAH and Am-DAH monomers are first verified via ¹H nuclear magnetic resonance [Supplementary Figure 2]. Subsequently, the p(DAH-co-Am-DAH) solid polymer electrolyte (SPE) is fabricated by solution casting and thermal curing in the presence of SN plasticizers, lithium bis(trifluoromethanesulphonyl)imide (LiTFSI) and lithium difluoro(oxalato)borate^[29]. The optimal Am-DAH dosage is determined by measuring the ionic conductivity of the copolymers with varying Am-DAH contents [Supplementary Figure 3]. Despite the relatively high conductivity of pristine p(DAH), the copolymer exhibits a non-monotonic trend, where conductivity initially increases but subsequently decreases with higher Am-DAH ratios. This drastic reduction is attributed to the dense hydrogen-bonding network introduced by the amide groups, which severely constrains polymer segmental motion and impedes Li⁺ transport^[30,31]. By precisely regulating this composition, the copolymerization approach achieves a critical synergy. The optimized p(DAH-co-Am-DAH) effectively utilizes the hydrogen-bonding network to anchor molecules and enhance stability, while simultaneously avoiding excessive chain restriction to ensure unimpeded Li⁺ conduction. Therefore, the subsequent discussion primarily focuses on the comparative analysis of p(DAH) and the optimized p(DAH-co-Am-DAH) (DAH:Am-DAH = 4:1, by weight) to elucidate the effect of the amide groups.

With the optimal composition established, the polymerization process and structural evolution are examined. As illustrated in Figure 1A, the precursor solution transforms from a flowable liquid into a transparent, gel-like solid following free radical thermal polymerization. Figure 1B presents an optical photograph of the composite electrolyte supported on a glass fiber membrane, exhibiting a milky-white and semi-transparent appearance. The surface morphology is analyzed via scanning electron microscopy (SEM). The p(DAH-co-Am-DAH) surface appears smooth and uniform, devoid of significant protrusions [Figure 1C]. This high degree of surface flatness ensures intimate contact between the electrolyte and the electrode materials, thereby effectively reducing interfacial impedance and maintaining low internal resistance within the battery. FTIR spectroscopy is employed to monitor the evolution of functional groups [Figure 1D]. Notably, the characteristic absorption peak at 1,670 cm^-1, attributed to the terminal C=C bonds of the precursors, completely disappears in the cured electrolyte, indicating a high degree of free radical polymerization. Furthermore, any trace unreacted monomers potentially remaining below the FTIR detection limit are effectively eliminated during the subsequent rigorous purification and vacuum drying steps, thereby ensuring the high purity of the final copolymer.

To further evaluate the electrolyte’s suitability for practical applications, thermogravimetric analysis (TGA) and tensile testing are employed to characterize the fundamental physicochemical properties. Figure 1E displays the TGA profiles, where both electrolytes exhibit two distinct thermal decomposition stages. The first stage at 119.30 °C corresponds approximately to the flash point of SN, while the second at 325.40 °C represents the decomposition of the polymer backbone. The thermal degradation curves of p(DAH-co-Am-DAH) and p(DAH) are highly coincident, indicating that the introduction of the amide modifying group does not significantly alter the thermal decomposition behavior. With backbone decomposition temperatures exceeding 300 °C, both electrolytes demonstrate good flame resistance. To evaluate the intrinsic mechanical properties of the polymer networks, tensile testing is performed on free-standing pure polymer films prepared without glass-fiber support [Figure 1F]. The p(DAH) exhibits an elongation at break of 16.52% and a tensile strength of 2.01 MPa. In comparison, p(DAH-co-Am-DAH) shows superior mechanical performance, with an elongation at break of 21.18% and a tensile strength of 2.99 MPa. This improvement is attributed to the role of Am-DAH as an amide hydrogen bond donor, which forms a hydrogen bond network with acceptors such as SN, TFSI^- and segments of the polymer backbone. This network strengthens the interactions between electrolyte components, thereby enhancing both tensile strength and elongation. Consequently, the electrolyte not only offers improved safety against thermal runaway but also possesses sufficient mechanical strength to suppress lithium dendrite growth^[32,33].

Density functional theory (DFT) calculations and FTIR spectroscopy are employed to characterize and verify the hydrogen bonding interactions between the polymer backbone and components such as SN, Li⁺ and TFSI^- in the p(DAH-co-Am-DAH) electrolyte. As shown in Figure 2A, DFT results show that Am-DAH exhibits a relatively lower binding energy with Li⁺ (-0.59 eV) compared to DAH (-0.95 eV). This suggests that the introduction of Am-DAH effectively reduces the overall affinity of the polymer backbone for Li⁺, thereby facilitating Li⁺ dissociation^[34]. In contrast to this weak Li⁺ affinity, Am-DAH demonstrates significantly stronger interactions with TFSI^- and SN compared to DAH, owing to the presence of the amide groups. Specifically, the binding energies of Am-DAH with TFSI^- and SN increase to -1.82 and -1.70 eV, respectively, markedly surpassing the values of -1.41 and -1.21 eV observed for DAH. This significant enhancement confirms the formation of a robust hydrogen bond network.

Subsequently, FTIR spectroscopy is employed to investigate the interactions of the terminal nitrile group (-C≡N) of SN with the p(DAH) and p(DAH-co-Am-DAH) electrolytes. As shown in Figure 2B, in the p(DAH) electrolyte, a peak appears at 2,275.1 cm^-1, corresponding to the Lewis acid-base interaction between the nitrile group and Li⁺. However, in the p(DAH-co-Am-DAH) system, the introduction of the amide hydrogen bond donor induces a blue shift of this peak to 2,277.6 cm^-1 This shift indicates that, in addition to interacting with Li⁺, the SN molecules also form hydrogen bonds with the amide groups on the polymer backbone. To corroborate these hydrogen bonding interactions, the response of the secondary amide (-NH-) groups on the polymer backbone is further analyzed. Upon the incorporation of SN and lithium salts into the p(DAH-co-Am-DAH) matrix, the bending peak shifts from 1,541.2 to 1,550.8 cm^-1 [Figure 2C], and the stretching peak shifts from 3,257.4 to 3,375.5 cm^-1 [Figure 2D]. These significant blue shifts provide compelling evidence of the strong interaction network formed between the polymer backbone, SN, and TFSI^-. In contrast, the p(DAH) spectrum naturally lacks these amide-specific features. Consequently, these spectral results confirm that the strong intermolecular forces in p(DAH-co-Am-DAH) effectively anchors the SN and TFSI^- components, thereby ensuring the stable localization of these species within the polymer matrix.

Raman spectroscopy is employed to investigate the variations in ion solvation structure induced by the Am-DAH hydrogen bond donor. First, the coordination states of SN are examined [Figure 3A and B]. SN exists primarily in two forms, identified as bound SN (interacting with the backbone or ions) at 2,254.0 cm^-1 and free SN at 2,278.0 cm^-1[35,36]. In the p(DAH) electrolyte, SN exists predominantly as free solvent molecules (90.1%), with only 9.9% in the bound state. In comparison, the p(DAH-co-Am-DAH) system exhibits a significantly higher proportion of bound SN, which reaches 30.4%. This increase occurs because the introduction of Am-DAH facilitates hydrogen bonding between the polymer backbone and SN molecules. Simultaneously, the weakened interaction between the backbone and Li⁺ releases more Li⁺, which further promotes SN coordination. Subsequently, the solvation states of the TFSI^- are analyzed [Figure 3C and D]. The anion exists in three forms, identified as free TFSI^- at 738.0 cm^-1, contact ion pairs (CIPs) at 741.6 cm^-1, and aggregates (AGGs) at 751.2 cm^-1[37,38]. In the p(DAH) electrolyte, TFSI^- exists primarily as free anions (44.5%), with CIPs and AGGs accounting for 31.2% and 24.3%, respectively. In comparison, the p(DAH-co-Am-DAH) system exhibits an aggregate-rich solvation structure, characterized by a lower fraction of free TFSI^- (40.9%) and a higher proportion of AGGs (27.3%). Similarly, the Am-DAH backbone anchors TFSI^- while releasing Li⁺, which collectively drives the formation of ion aggregates. These distinct solvation structures have profound implications for interfacial chemistry. The abundance of free SN in p(DAH) tends to undergo parasitic decomposition, leading to the formation of an unstable, organic-rich solid electrolyte interphase (SEI). Conversely, the enrichment of CIPs and AGGs in p(DAH-co-Am-DAH) promotes anion-derived reduction, fostering a robust, inorganic-rich SEI that enhances electrochemical stability^[39].

To evaluate the impact of the Am-DAH donor on electrochemical performance, room-temperature ionic conductivity and exchange current density are measured. As shown in Figure 4A, the p(DAH-co-Am-DAH) electrolyte exhibits a higher ionic conductivity (0.92 mS·cm^-1) than the p(DAH) (0.79 mS·cm^-1). This enhancement is initially attributed to the weakened interaction between the polymer backbone and Li⁺, which facilitates Li⁺ dissociation. To validate this mechanism, the ion migration activation energy is determined by fitting the temperature-dependent ionic conductivity [Supplementary Figure 4] to the Arrhenius equation [Figure 4B]. The p(DAH-co-Am-DAH) electrolyte exhibits a lower ion transport activation energy of 0.258 eV compared to 0.278 eV for p(DAH). This reduction confirms that the Am-DAH modification weakens the interaction between the polymer backbone and Li⁺, thereby lowering the energy barrier for Li⁺ dissociation into the SN/TFSI^- mobile phase. Beyond bulk ion transport, the interfacial Li⁺ plating/stripping kinetics are also assessed using Tafel plots [Figure 4C]. The p(DAH-co-Am-DAH) electrolyte exhibits a higher exchange current density of 0.162 mA·cm^-2 compared to 0.147 mA·cm^-2 for p(DAH), indicating that the optimized solvation structure facilitates faster interfacial charge transfer^[33]. These superior kinetic parameters are crucial for reducing internal resistance, improving rate capability and mitigating heat-induced degradation.

While high conductivity and fast kinetics are essential, a high Li⁺ transference number is equally critical to ensure efficient Li⁺ transport. The Li⁺ transference number is determined using the Evans method [Supplementary Figure 5, combining AC impedance and DC polarization]. The p(DAH-co-Am-DAH) electrolyte achieves a remarkably high Li⁺ transference number of 0.82, significantly surpassing the 0.52 of p(DAH) [Figure 4D]. This improvement stems from a unique structural configuration where SN and TFSI^- are effectively anchored to the polymer backbone via hydrogen bonds, thereby restricting their mobility. Consequently, the ionic current is predominantly carried by highly mobile Li⁺. A high Li⁺ transference number enhances energy efficiency, reduces concentration polarization, and promotes uniform Li deposition, effectively suppressing dendrite growth and extending cycle life^[40]. Furthermore, the intrinsic electrochemical stability windows of p(DAH) and p(DAH-co-Am-DAH) are evaluated via linear sweep voltammetry. Both electrolytes demonstrate impressive oxidation resistance, with onset potentials of 5.53 and 4.81 V, respectively [Supplementary Figure 6]. Such wide electrochemical windows ensure full compatibility with commercial cathodes operating at elevated voltages, confirming their suitability for high-energy-density battery systems^[41].

To evaluate the long-term compatibility of the electrolytes with lithium metal anodes, Li||Li symmetric cells are assembled using p(DAH) and p(DAH-co-Am-DAH) electrolytes and subjected to galvanostatic cycling tests [Figure 5A]. The p(DAH-co-Am-DAH) cell exhibits superior stability, operating for 1,700 h at a current density of 0.5 mA·cm^-2 with a stable overpotential. In comparison, the p(DAH) succumbs to an internal short circuit after approximately 1,600 h, accompanied by a noticeably higher overpotential. This electrochemical durability is further corroborated by post-cycling morphological analysis via SEM [Figure 5B and C]. The lithium surface cycled with p(DAH-co-Am-DAH) retains a smooth and dense texture free of obvious dendrites, whereas the p(DAH)-cycled surface appears significantly rougher, confirming that the modified electrolyte promotes more uniform lithium deposition.

To elucidate the chemical origin of this enhanced stability, the SEI composition is probed using XPS. As a highly surface-sensitive technique (typically probing the top 5-10 nm), XPS specifically reveals the chemical composition and robust nature of the SEI formed on the electrode surface, rather than the bulk electrolyte. Accordingly, Figure 5D-G displays the high-resolution spectra of C 1s, O 1s, F 1s, and N 1s for p(DAH) and p(DAH-co-Am-DAH), respectively. The C 1s spectra are deconvoluted into four peaks at 284.8 eV (C-C), 286.8 eV (C-O), 289.4 eV (-C≡N) and 293.1 eV (-CF₃)^[42,43]. The C-C and C-O signals originate from the decomposition of the polymer backbone, while the -C≡N and -CF₃ are derived from the decomposition of SN and TFSI^-, respectively^[44]. The O 1s spectra further confirm the partial decomposition of the polymer skeleton with peaks at 528.1 eV (Li₂O), 530.8 eV (C=O), 531.7 eV (Li₂CO₃) and 532.5 eV (C-O)^[44]. Crucially, the F 1s and N 1s spectra reveal peaks at 684.8 eV (LiF), 687.1/689.0 eV (-C-F/-CF₃), 398.4 eV (Li₃N), and 399.5 eV (-C≡N)^[44-46]. Notably, compared to p(DAH), the p(DAH-co-Am-DAH) exhibits higher relative intensities for LiF and Li₃N peaks, indicating that a greater proportion of TFSI^- anions participated in the SEI formation^[47]. This high-quality interface effectively suppresses dendrite growth and minimizes side reactions, thereby ensuring the superior long-term cycling stability and compatibility of the p(DAH-co-Am-DAH) electrolyte.

To further validate the practical applicability of the electrolytes beyond symmetric cells, full cells are assembled using a nickel-rich LiNi_0.8Co_0.1Mn_0.1O₂ (NCM811) cathode. NCM811 is regarded as one of the most promising cathode materials for high-energy-density lithium-ion batteries due to its high specific capacity and high operating voltage. The Li|p(DAH-co-Am-DAH)|NCM811 and Li|p(DAH)|NCM811 coin cells are assembled and subjected to rate capability and long-term cycling tests within a voltage range of 3.0-4.3 V. Figure 6A and B displays the rate capability of the cells at currents of 0.1, 0.2, 0.5, 1, and 2 C (1 C = 200.0 mAh·g^-1). The Li|p(DAH-co-Am-DAH)|NCM811 cell delivers discharge capacities of 220.8, 212.8, 200.5, 187.3, and 152.6 mAh·g^-1, respectively. In comparison, the Li|p(DAH)|NCM811 cell yields lower capacities of 200.6, 193.4, 181.1, 167.4, and 137.1 mAh·g^-1, respectively. This superior rate performance is attributed to the higher ionic conductivity and higher Li⁺ transference number of p(DAH-co-Am-DAH) compared to p(DAH), which enables faster charge/discharge kinetics. The robustness of the system is further confirmed by the rate recovery test [Figure 6C]. After undergoing the variable-rate testing protocol, the Li|p(DAH-co-Am-DAH)|NCM811 cell recovers a discharge capacity of 214.4 mAh·g^-1 at 0.1 C, retaining 97.1% of its initial capacity.

Finally, the long-term cycling stability at 1 C is evaluated, as shown in Figure 6D. The p(DAH-co-Am-DAH) system demonstrates exceptional durability, with capacity retentions of 99.9%, 92.9% and 85.3% at the 200th, 350th, and 500th cycles, respectively. In comparison, the p(DAH) system suffers a sudden capacity decay after 413 cycles and fails to recover, accompanied by severe fluctuations in Coulombic efficiency. To confirm that the enhanced performance originates from intrinsic material properties rather than experimental variations, a stringent high-voltage cycling test (2.8-4.6 V) is conducted [Supplementary Figure 7]. Under these harsh conditions, the performance gap becomes highly pronounced. The p(DAH-co-Am-DAH) SPE maintains stable operation for 200 cycles with 92.2% capacity retention. In contrast, the p(DAH) cell experiences a sudden capacity drop at the 114th cycle, followed by rapid degradation. This contrast demonstrates that the p(DAH-co-Am-DAH) exhibits superior interfacial compatibility and structural stability with the NCM811 cathode, ensuring a prolonged lifespan under severe electrochemical stress. Post-cycling SEM analysis on the NCM811 cathodes with p(DAH-co-Am-DAH) and p(DAH) [Supplementary Figure 8] provides direct morphological evidence to support this conclusion. This long-term stability significantly outperforms that of representative electrolytes reported in recent literature [Figure 6E and Supplementary Table 1], further demonstrating the immense potential of the p(DAH-co-Am-DAH) polymer electrolyte for high-voltage cathode applications.