Ultra‐Uniform Lithium‐Ion Transport Enabled by Supramolecular Polymeric Networks as Artificial Solid Electrolyte Interphase Layers for Highly Stable Lithium‐Ion Battery Anodes

ABSTRACT

Polymer‐based artificial solid electrolyte interphase (SEI) layers have emerged as a promising solution to address the inherent limitations of silicon‐carbon nanocomposite (SCN) anodes. However, their practical implementation remains hindered by the inherent trade‐off between achieving complete surface coverage and maintaining a thin, uniform coating. This trade‐off often compromises either the electrolyte‐blocking capability or the Li‐ion transport efficiency. To overcome these challenges, we aim to enhance the ionic conductivity of the artificial SEI layer to levels comparable to liquid electrolytes, while simultaneously improving Li‐ion dissociation properties. To this end, we developed a polymer‐based supramolecular artificial SEI layer incorporating p‐phenylenediamine (pPD) as a bridging agent. The supramolecular network formed via pPD introduces robust hydrogen bonding and facilitates the formation of Li‐ion hopping channels through its benzenoid–quinoid transition. As a result, the incorporation of pPD significantly increases the ionic conductivity of PEO and PMMA polymers to 0.215 and 0.106 mS cm⁻¹, respectively. Furthermore, SCN anodes coated with this supramolecular SEI exhibited over fourfold improvement in cycling stability under ultra‐lean electrolyte conditions, closely mimicking commercial operating environments, compared to uncoated SCN in full‐cell configurations. This study offers a robust platform for the design of advanced artificial SEI layers tailored for high‐performance anode materials.

Incorporation of p‐phenylenediamine (pPD) into the polymer–pPD supramolecular artificial SEI reduces the activation energy for Li‐ion desolvation at the interface and promotes Li‐ion transport through the layer, thereby ensuring uniform delivery of Li‐ion into the SCN and stable cycling performance regardless of the inherently uneven thickness of the polymer coating.

1. Introduction

Employing polymeric artificial solid electrolyte interphase (SEI) layer on silicon–carbon nanocomposite (SCN) anodes provide a rational and effective strategy to address the persistent interfacial instabilities that limit the practical implementation of SCN in lithium‐ion batteries [1]. On the surface of SCN, Li‐ion migrate through the electrolyte in a solvated state, solvated Li‐ion inevitably carries its solvent molecules to the SCN surface, where both Li‐ion and solvent participate in interfacial reduction reactions [2, 3]. Such parasitic processes consume electrolyte, generate unstable SEI components, and reduce Coulombic efficiency, ultimately compromising long‐term cycling stability [4]. To fundamentally suppress these side reactions, it is essential to introduce a mechanism that enforces de‐solvation of Li‐ion before it reaches the SCN surface [5]. Polymeric coatings can serve this role by functioning as artificial SEI layers that not only tolerate repeated volume changes through their elasticity but also regulate ion transport at the molecular level [6]. The polar functional groups within the polymer coating facilitate the removal of solvent molecules, thereby ensuring that only bare Li⁺ migrates across the interface, while solvent penetration is effectively blocked [7]. This selective ion transport pathway minimizes continuous electrolyte decomposition and stabilizes SEI evolution, enabling improved Coulombic efficiency and durable cycling. Consequently, polymer‐coated SCN anodes integrate the high capacity, the conductivity of carbon, and the interfacial stability of an engineered SEI.

To effectively block electrolytes and inhibit parasitic side reactions, it is crucial that the polymer artificial SEI layer fully encompasses the surface of the SCN particles. At the same time, this polymer layer, which has a relatively lower ionic conductivity than liquid electrolytes, needs to be applied in a thin and uniform manner to ensure efficient Li‐ion transport to the SCN surface [8, 9]. However, achieving consistent and comprehensive coverage on the complex surfaces of micron‐sized SCN particles is challenging when using cost‐effective and scalable techniques such as solution mixing coating, spray coating, and slurry coating [10]. These methods, while practical for mass production, often struggle to maintain thin and uniformity, leading to areas that may remain uncoated [11, 12]. This inconsistency in the application of the SEI layer compromises its protective capabilities, leaving portions of the SCN surface directly exposed to the electrolyte. Conversely, attempts to ensure complete coverage can result in non‐uniform layer thicknesses, with certain regions of the SCN particles being excessively coated. These overly thick polymer layers hinder Li‐ion transport due to their inherently low ionic conductivity (10⁻⁶ – 10⁻⁷ S cm⁻¹), which is several orders of magnitude lower than that of liquid electrolytes (∼10⁻² S cm⁻¹). As a result, such transport limitations can lead to non‐uniform lithiation and delithiation behavior, ultimately deteriorating electrochemical performance and inducing mechanical stress within the SCN during cycling [13].

Additionally, the stability of the polymer‐based artificial SEI layer through multiple cycles requires robust mechanical strength [14]. Consequently, the strategy of enhancing ionic conductivity by promoting chain mobility with low‐molecular‐weight polymers is impractical. Given these limitations, traditional methods of applying polymer‐based artificial SEI layers face significant challenges in fully realizing their intended benefits. Therefore, to fundamentally improve the performance of SCN anodes, polymeric artificial SEI layers must not only serve as passive protective coatings but also simultaneously and indispensably fulfill three critical requirements: (i) desolvation of lithium ions at the electrode interface, (ii) effective electrolyte‐blocking capability through sufficiently robust coatings, and (iii) consistent and efficient Li‐ion transport across the interphase [15]. These three requirements are not optional or independent design elements; rather, they constitute essential and interdependent conditions that must be integrated within a single polymeric SEI architecture to ensure stable and efficient operation. However, recent studies on artificial SEI layers for SCN anodes have often focused solely on surface stability or exclusively on ionic conductivity, while neglecting to address these considerations in an integrated manner [6, 16, 17].

Based on this rationale, our objective is to concurrently achieve efficient Li‐ion desolvation, high ionic conductivity, and sufficient mechanical robustness by incorporating functional additives into the polymer matrix. This integrated approach enables highly uniform ion transport behavior alongside strong electrolyte‐blocking capability within the polymer‐based artificial SEI layer. As a result, even when the practical solution‐based mix‐coating process leads to inconsistencies in coating thickness, the combination of efficient Li‐ion desolvation and enhanced ionic conductivity ensures uniform lithiation–delithiation of the SCN anode, regardless of local variations in layer thickness [18, 19]. To address this challenge, we have developed a polymer‐based supramolecular artificial SEI layer using p‐phenylenediamine (pPD) as a molecular bridging agent, applied through a scalable solution‐based mix‐coating method. The interaction between pPD and the anionic functional groups of the polymer induces strong hydrogen bonding, facilitating the formation of a supramolecular network structure [20]. Within this architecture, the conjugated structure of pPD promotes the desolvation of electrolytes and Li ions at the SEI surface, while simultaneously serving as ion‐transport pathways through the polymer matrix via the benzenoid–quinoid transition, enabling rapid Li‐ion hopping. Additionally, the introduction of pPD improves the packing density and mechanical integrity of the polymer coatings, significantly enhancing their electrolyte‐blocking performance and operational stability [21].

In experiments employing pPD as a bridging agent within poly(ethylene oxide) (PEO) and poly(methyl methacrylate) (PMMA) polymers, which is most, both of which are readily applicable for polymer based artificial SEI layers, optimal ionic conductivity was achieved. The most favorable performance was observed at a monomer‐to‐pPD ratio of 6:1, where the polymer layers exhibited ionic conductivities exceeding the practical threshold of 0.1 mS cm⁻¹. Specifically, the conductivities increased to 0.215 mS cm⁻¹ for PEO and 0.106 mS cm⁻¹ for PMMA, in contrast to their baseline values of 0.031 and 0.004 mS cm⁻¹, respectively. Additionally, computational analyses based on density functional theory (DFT) and molecular dynamics (MD) simulation revealed that the incorporation of pPD promotes Li‐ion desolvation and accelerates Li‐ion transport within the pPD–polymer supramolecular framework. As a result, the enhanced Li‐ion desolvation at the surface together with the improved ionic conductivity minimizes the transport disparity relative to liquid electrolytes, thereby enabling uniform Li‐ion migration through the artificial SEI layer regardless of the coating thickness. In addition, the mechanical strength of the modified polymer coatings was enhanced by a factor of 2.5 compared to the unmodified counterparts.

Therefore, following the successful application of a pPD‐based supramolecular artificial SEI layer on the SCN, we achieved enhanced cyclic stability, with initial capacities of 1869.4 and 1845.1 mAh g⁻¹, and a retention of 72.9% and 73.3% of the reversible capacity after 100 cycles at 0.5 C in half‐cell tests, showing improved cyclic stability compared to both bare SCN and purely polymer‐coated SCN anodes. Additionally, full‐cell tests conducted under extremely lean electrolyte conditions (∼2 µL mAh⁻¹), representative of practical electrolyte‐to‐capacity ratios and simulating industrial‐scale cell operation, demonstrated that modified SCN‐based electrodes exhibited more than fourfold greater stability in electrolyte consumption compared to bare SCN based electrodes [22]. We believe that the presented approach establishes a fundamental platform for designing artificial SEI layers and developing advanced anode materials, with optimized management of electrolyte side reactions on the active material surface.

2. Results and Discussion

2.1. Coating of Polymer‐Based Supramolecular Artificial SEI Layer

As depicted in the schematic illustration of Figure 1a, the formation of the artificial SEI layer on the surface of SCN began with the application of either PEO or PMMA polymer layers via a solution‐based mix‐coating process, a method that is readily scalable and compatible with commercial implementation. The SCN used in this study, as shown in Figure S1, had an average size of about 3 µm, was composed of 60% Si, and demonstrated a capacity of approximately 1800 mAh g⁻¹. The primary objective at this stage was to achieve complete coverage of the SCN's entire surface. Prioritizing full coverage often leads to a polymer coating with inherently uneven and randomly varying thicknesses. This variation in thickness can cause non‐uniform Li‐ion transport to the SCN, as illustrated in Figure 1b, placing significant stress on the SCN during repeated lithiation‐delithiation cycles and potentially accelerating its structural degradation. In contrast, the polymer‐pPD supramolecular layer, formed through hydrogen bonding, facilitates highly uniform Li‐ion transport due to the incorporation of pPD, which promotes efficient Li‐ion desolvation and facilitates rapid Li‐ion hopping. Remarkably, the activation energy required for Li‐ion transport in pPD is five times lower than that in PEO, a commonly used polymer in solid polymer electrolytes, demonstrating its superior ion transport capabilities. Additionally, the semi‐crosslinked structure offers superior electrolyte‐blocking and mechanical properties compared to a single polymer coating, thereby improving interfacial stability. Consequently, the SCN coated with the polymer‐pPD‐bridged supramolecular artificial SEI layer supports selective and efficient lithiation and delithiation at the SCN surface, enabling the artificial SEI layer to fulfill its protective role effectively.

FIGURE 1.

(a) Schematic images illustrating the process of coating SCN surfaces with polymer‐pPD supramolecular artificial SEI layer. Li‐ion transport scheme in a random thickness artificial SEI layer: (b) polymer‐only layer (left) and polymer‐pPD supramolecular layer (right). TEM and EDS mapping images of (c) EP‐SCN and (d) MP‐SCN. (e) ¹H NMR data of EP‐SCN and MP‐SCN. XPS data of (f) pPD, EP‐SCN, and MP‐SCN and (g) deconvolution data for the N 1s peak of EP‐SCN and MP‐SCN.

For convenience of discussion, PEO and PMMA complexed with pPD are referred to as PEO/pPD and PMMA/pPD, respectively. The SCN coated with PEO/pPD and PMMA/pPD are designated EP‐SCN and MP‐SCN, respectively, while the control samples, where SCN is coated with PEO and PMMA alone, are named E‐SCN and M‐SCN, respectively. The coated PEO/pPD and PMMA/pPD layers are visible through transmission electron microscope (TEM) imaging and energy‐dispersive X‐ray spectroscopy (EDS) mapping. As presented in Figure 1c,d and Figures S2 and S3, the polymer layers coated on SCN in all samples exhibit a randomly distributed coating thickness, ranging approximately from 5 to 40 nm. Furthermore, unlike the EDS mapping images of E‐SCN and M‐SCN shown in Figure S4, the EDS mapping images of both EP‐SCN and MP‐SCN samples reveal the presence of nitrogen elements, suggesting that the artificial SEI layer containing pPD has been successfully applied. Additionally, Thermogravimetric analysis (TGA) data shown in Figure S5 confirms that the coating amount is exceptionally small, with PEO/pPD on EP‐SCN at 1.1 wt% and PMMA/pPD on MP‐SCN at 0.8 wt%. Such a minimal thickness of the artificial SEI layer effectively minimizes additional weight, thereby reducing the gravimetric capacity loss of the electrode attributable to additives. As shown in Figure S6, the coated artificial SEI layer forms an electronic barrier on the SCN surface after coating. This electronic barrier effectively suppresses electrolyte‐induced side reactions in regions of the surface that are not in direct contact with the conductive additive.

The hydrogen bonding‐utilized molecular bridging of pPD in the PEO/pPD and PMMA/pPD coatings on EP‐SCN and MP‐SCN was confirmed through ¹H Nuclear Magnetic Resonance (NMR) and X‐ray photoelectron spectroscopy (XPS) analyses. In Figure 1e, the peaks of pPD observed around 7 ppm in ¹H NMR appear in a split form, rather than as a single peak, indicating the presence of hydrogen bonding. The comprehensive XPS data shown in Figure S7, particularly the N peak observed around 400 eV, further confirm the presence of pPD on the SCN surface. Additionally, in Figure 1f, the N 1s peak of pPD, typically at 398.9 eV, shifts to 399.5 eV in EP‐SCN and 400 eV in MP‐SCN. In these hydrogen bonds, the N atom, acting as a proton donor, transfers electron density to the H atom to from stable hydrogen bonding with an O atom [23, 24]. This results in a relative reduction in electron density around the nitrogen, observed as an increase in the binding energy of N in XPS analysis following the formation of hydrogen bonds. As illustrated in Figure 1g, deconvolution of the N 1s peak for EP‐SCN and MP‐SCN reveals that the amine peak of pure pPD (398.9 eV) is barely detectable. Instead, the amine peaks corresponding to hydrogen bonding with PEO (399.6 eV) and PMMA (399.9 eV) are predominantly observed, indicating specific hydrogen bonding interactions with each respective polymer.

2.2. Characterization of Polymeric Supramolecular Network Structures

Hydrogen‐bonded pPD serves not only as a bridging agent but also facilitates Li‐ion transport within polymers (Figure 2a). The mechanism through which pPD supports Li‐ion transport involves a benzenoid‐quinoid transition, illustrated in Figure 2b. This structural shift enables Li ions hopping across the polymer network. When Li‐ion is introduced to pPD, the adjacent amine group transforms into a quinoid‐based imine structure, facilitating charge transfer across the pPD molecule. This leads to Li‐ion binding to the aromatic ring and subsequently hopping to the oppositely placed N atom. In this process, the activation energies for Li‐ion hopping in these transitions were determined to be 5.2 kcal mol⁻¹ for the amine‐to‐benzene and 4.8 kcal mol⁻¹ for the benzene‐to‐amine transition [21]. These activation energies are significantly lower than those for ionic conduction in PEO (25.1 kcal mol⁻¹), highlighting the efficiency of pPD as a low‐barrier pathway for Li‐ion transport. To confirm the benzenoid‐quinoid transition, Fourier transform infrared (FTIR) and NMR analyses were conducted on the film‐cast PEO/pPD and PMMA/pPD with lithium salt, providing insights into ion transport behavior. For clarity, the explanation was based on the positioning of benzenoid and quinoid states as shown in Figure 2b. In the FTIR spectrum, depicted in Figure 2c, the amine peak of pPD at 1513 cm⁻¹ exhibits a blue shift to 1519 cm⁻¹ due to hydrogen bonding with PEO or PMMA and lithium ion doping (2‐position and b‐position) [25]. Additionally, a new peak appears at 1330 cm⁻¹, which corresponds to the C═N bond formed in the quinoid structure of pPD (d‐position) [26].

FIGURE 2.

(a) Mechanism of increased ionic conductivity by pPD molecules. (b) Schematic illustration of the cross‐liked pPD and its Li‐ion doped phase. (c) FTIR spectra of pPD, PEO, PMMA, PEOpPD, and PMMApPD. ¹H‐NMR of pPD molecules and (d) PEO and PEOpPD and (e) PMMA and PMMApPD. ¹³C‐NMR of pPD molecules and (f) PEO and PEOpPD and (g) PMMA and PMMApPD. (h) Ionic conductivity, (i) Arrhenius plot, and (j) strain‐stress curve of PEO, PMMA, PEOpPD, and PMMApPD. All the polymer, except PEOpure and PMMApure, were blended with bis(trifluoromethane sulfonyl)imide Li salt (LiTF₂N) powders prior to analysis.

Figure 2d–g presents the results of ¹H NMR and ¹³C NMR analyses, which confirm the benzenoid‐quinoid transition of pPD following hydrogen bonding with PEO and PMMA. The complexation behavior observed in PEO was analyzed in D₂O, while the complexation in PMMA was assessed in acetone‐d₆. In ¹H‐NMR, pure pPD exhibits peaks at 6.6 ppm (1‐positon) and 4.7 ppm (2‐ position) in D₂O, and at 6.6 ppm (1‐positon) and 3 ppm (2‐ position) in acetone‐d₆. Upon interaction with PEO and PMMA, the peak at the 1‐position distinctly splits into multiple peaks (a‐, b‐, and c‐positions), clearly indicating the presence of hydrogen bonding interactions. In ¹³C‐NMR, the C‐H (1‐position) and C‐N (2‐position) peaks of pure pPD are recorded at 118 and 138 ppm in D₂O, and at 120 and 140 ppm in acetone‐d₆. When combined with PEO and PMMA, these peaks shift to 120 and 142 ppm for PEO/pPD, and to 121 and 147 ppm for PMMA/pPD, respectively. Furthermore, due to the quinoid transition, additional peaks arising from the d‐position and c‐position emerge near the a‐position. These observations further substantiate the structural transformation of pPD in response to its interactions with PEO and PMMA, demonstrating the benzenoid‐quinoid transition that occurs upon hydrogen bonding with the polymers and doping with lithium salt, thus facilitating rapid Li‐ion transport.

The hydrogen bonding of pPD introduces additional Li‐ion transport channels within the polymer matrix. As pPD forms hydrogen bonds with adjacent surrounding polymers, it facilitates pathways that enhance Li‐ion conductivity. The resulting changes in ionic conductivity, as influenced by these newly formed transport channels, are depicted in Figure 2h. The ionic conductivity of pure PEO and pure PMMA was initially measured at 3.12 × 10⁻⁵ and 3.90 × 10⁻⁶ S cm⁻¹, respectively. With the integration of pPD, the ionic conductivity of PEO/pPD and PMMA/pPD increased significantly, reaching 2.15 × 10⁻⁴ and 1.06 × 10⁻⁴ S cm⁻¹, representing enhancements by approximately 7‐fold and 25‐fold, respectively. The Arrhenius plot presented in Figure 2i and Figure S8 exhibits a decrease in the activation energy required for Li‐ion transport. Considering that the glass transition temperature (T _g) of PEO is ∼60°C, the plot was segmented into temperature regions above and below T _g. The activation energies exhibit a reduction from 0.71 eV and 0.45 eV to 0.46 eV and 0.43 eV for the PEO polymer, and from 0.66 eV to 0.40 eV for the PMMA polymer, indicating more efficient ion transport. It is crucial to note that this enhancement in ionic conductivity can be ascribed to the formation of additional Li‐ion hopping channels, which significantly improve Li‐ion transport from one polymer chain to another. It should be noted that, although the incorporation of pPD enhances the overall ionic conductivity, the local ion dissociation behavior can vary depending on the functional groups of the polymer backbone, and the chain mobility of the polymer also contributes to ionic conductivity in polymer‐based systems.

Additionally, hydrogen‐bonded pPD acts as a bridging agent (i.e. semi‐crosslinking agent) that connects polymer chains, enabling PEO/pPD and PMMA/pPD to form a supramolecular structure that enhances mechanical strength. Figure 2j assesses the impact of pPD incorporation on the mechanical fracture behavior of PEO, PMMA, PEO/pPD, and PMMA/pPD samples through tensile testing (ASTM D882). The PEO and PMMA films initially exhibited tensile stresses of 2.6 MPa at 0.98% strain and 19.4 MPa at 1.95% strain, respectively. In contrast, the introduction of bridging through pPD molecules resulted in enhanced mechanical properties, with PEO/pPD achieving an improved tensile stress of 6.6 MPa at a higher strain of 2.18%, and PMMA/pPD demonstrating a tensile stress of 56.3 MPa at 3.21% strain. These improvements in strain limits indicate that molecular bridging by pPD facilitates the formation of a supramolecular network structure through hydrogen bonding‐based interlinking, enhancing the overall elasticity of the complexed materials.

2.3. Computational Evaluation of Li‐Ion Desolvation and Transport in the Supramolecular Artificial SEI Layer

To elucidate the mechanism of Li‐ion transport from the electrolyte (ethylene carbonate, EC) into polymer‐based supramolecular artificial SEI layers, we conducted DFT calculations. Specifically, we evaluated the migration energy barriers associated with Li‐ion desolvation from Li‐EC complexes into polymer hosts, namely PEO or PMMA, in both the presence and absence of pPD. The positions of Li ions along the reaction coordinate are shown in Figure S9. In the case of the PEO‐based SEI (Figure 3a), the migration barrier was found to be 2.13 eV without pPD, indicative of a significant energetic hindrance to Li‐ion transfer across the electrolyte/SEI interface. Notably, the incorporation of pPD introduces bridging channels that lower this barrier to 1.37 eV, thereby enhancing both Li‐ion desolvation and interfacial transport. A similar trend was observed in the PMMA‐based SEI (Figure 3b), where the energy barrier decreased from 2.10 eV to 1.60 eV upon the incorporation of pPD. These results clearly demonstrate that pPD effectively facilitates Li‐ion desolvation across diverse polymeric matrices. The reduction in energy barriers is anticipated to increase the population of mobile Li‐ions at the interface, thereby enabling more efficient transport into the SCN and promoting uniform lithiation/delithiation behavior. Furthermore, Raman spectroscopy was employed to verify the desolvation of Li ions from Li–EC complexes in relation to the reduced energy barrier for Li ions. As shown in Figure S10, the characteristic Raman peak of EC at 743 cm⁻¹ exhibits a red shift to 741 cm⁻¹ upon incorporation of pPD. This red shift signifies a weakened coordination environment between Li⁺ ions and EC molecules, indicating that partial desolvation of Li⁺ ions occurs as a result of pPD incorporation.

FIGURE 3.

DFT‐calculated formation energies of Li‐ion transferred from the electrolyte into (a) PEO‐based and (b) PMMA‐based artificial SEI layers. MD simulation models of Li⁺ transport (c) within PEO (left) and PEO/pPD (right), and (d) PMMA (left) and PMMA/pPD (right). MSD data of (e) PEO vs. PEO/pPD and (f) PMMA vs. PMMA/pPD, respectively, with the corresponding diffusion coefficients shown on the right‐hand side of each panel.

Achieving uniform and rapid Li‐ion transport within SCNs requires not only effective desolvation of Li‐ions from the electrolyte but also their efficient migration across the polymer‐based artificial SEI layer. To investigate this, therefore, MD simulations were also performed to investigate Li‐ion transport through the supramolecular artificial SEI layers, as illustrated in Figure 3c,d. All MD simulations were performed using the Forcite module within the Materials Studio 2024 software package under three‐dimensional periodic boundary conditions [27, 28]. The simulations were carried out on the nanosecond timescale, with all calculations extended up to 2 ns. The results reveal that the incorporation of pPD substantially enhances Li‐ion mobility in both polymer systems. This enhancement is quantitatively corroborated by the mean square displacement (MSD) profiles and the corresponding diffusion coefficients shown in Figure 3e,f. In the PEO‐based system, the diffusion coefficient increased from 2.25 × 10⁻⁷ to 4.11 × 10⁻⁷ cm² s⁻¹ upon pPD incorporation. Likewise, in the PMMA‐based system, the diffusion coefficient rose significantly from 1.48 × 10⁻⁸ to 1.00 × 10⁻⁷ cm² s⁻¹. These findings confirm that pPD plays a critical role in facilitating Li‐ion transport through the artificial SEI layer. Notably, such an accelerated and ultra‐uniform Li‐ion transport alleviates local lithiation/delithiation heterogeneity, which is an inherent limitation of conventional non‐uniform polymer coatings. As a result, it facilitates not only enhanced electrochemical performance but also improved operational stability.

2.4. Evaluation of Half‐Cell Performance for the Introduction of Supramolecular Artificial SEI Layer

In the case of conventional polymer coating layers with insufficient Li‐ion conductivity (below 0.1 mS cm⁻¹), the Li‐ion transport characteristics at the SCN surface are high sensitivity by the coating thickness [26]. Therefore, uneven Li‐ion transport can occur due to an unavoidable non‐uniform coating layer with low ionic conductivity, causing localized lithiation‐delithiation of SCN that can diminish the lifespan of batteries. Conversely, when the polymer‐based artificial SEI layer possesses high Li‐ion conductivity (above 0.1 mS cm⁻¹), Li‐ion transport characteristics at the SCN surface are little affected by the coating thickness. This allows for the formation of a fully covered coating layer that minimizes performance degradation associated with lithiation‐delithiation processes. As illustrated in Figure 2, the incorporation of pPD into the coating layer enhances its ionic conductivity, effectively resolving these issues. Additionally, the improved mechanical properties of the coating significantly enhance the stability of the SCN active material.

Compared to SCN with either a bare or conventional polymer‐based artificial SEI layer, the pPD‐based supramolecular artificial SEI layer facilitates fast Li‐ion transport at the SCN anode surface and offers superior electrolyte‐blocking capabilities by acting as a bridging agent. This enhanced ion transport aids in the activation of the SCN and mitigates localized electron accumulation, thereby improving cycle stability. Figure 4a,b presents the half‐cell test results for each anode material. The half‐cell test was performed in constant current‐constant voltage (CC‐CV) mode, a more rigorous condition compared to conventionally employed constant current (CC) mode. This is because the extent of Li‐ion lithiation in CV mode is much greater, which increases the likelihood of side reactions and lithium deposition. With standard SCN, a marked decline in performance due to electrode degradation is observed after 50 cycles. While simple polymers like PEO or PMMA when used as coatings alleviate performance degradation to some degree, a gradual decline remains evident. In contrast, SCN coated with the pPD‐based supramolecular artificial SEI layer demonstrates stable operation for up to 100 cycles.

FIGURE 4.

Cell performances were evaluated with coin cells. The coin cells were cycled in CC‐CV mode. Comparison of cycle performances of SCN, (a) E‐SCN, EP‐SCN, (b) M‐SCN, and MP‐SCN anodes. Galvanostatic charge‐discharge profiles of (c) EP‐SCN and (d) MP‐SCN anodes. (e) Rate capability test of SCN, E‐SCN, EP‐SCN, M‐SCN, and MP‐SCN anodes.

As depicted in Figure 4c,d, and Figure S11, the initial specific capacity of SCN is measured at 1799.6 mAh g⁻¹. SCN samples coated with PEO and PMMA (referred to as E‐SCN and M‐SCN, respectively) exhibit initial specific capacities of 1798.6 and 1787.5 mAh g⁻¹. Conversely, EP‐SCN and MP‐SCN display significantly enhanced specific capacities of 1869.4 and 1845.1 mAh g⁻¹. This indicates that while polymer coatings improve the electrolyte wettability of the SCN surface, they often lack sufficient ionic conductivity, which hinders rapid ion transport and thus limits the achievable capacity of SCN. However, the integration of pPD into the coating layer achieves improved ionic conductivity, facilitating faster ion transport and resulting in a higher capacity for SCN. Additionally, with the incorporation of pPD, the density of the coating layer is improved, resulting in an initial efficiency increase from 83.67% (bare SCN) to 84.95% (E‐SCN) and 84.13% (M‐SCN), while EP‐SCN and MP‐SCN reach higher initial efficiencies of 86.51% and 86.21%, respectively. Furthermore, the average Coulombic efficiency from cycles 3 to 40 rose from 99.28%, 99.45%, and 99.31% for SCN, E‐SCN, and M‐SCN, to 99.69% and 99.62% for EP‐SCN and MP‐SCN. These half‐cell cycle performance results demonstrate that, while PEO and PMMA coatings alone did not fully enhance the stability of the cell, the inclusion of pPD markedly improved the stability of SCN without any performance degradation.

As shown in Figure S12, the extent of electrode volume change during lithiation clearly depends on the presence of pPD. The SCN, E‐SCN, and M‐SCN half cell electrodes exhibited approximately twofold volume expansion, accompanied by pronounced cracking after lithiation. In contrast, the pPD‐containing EP‐SCN and MP‐SCN half‐cell electrodes showed substantially reduced expansion of approximately 1.5‐fold, with no observable crack formation. These results suggest that the incorporation of pPD not only suppresses the volume expansion of SCN by forming a robust artificial SEI layer, but also enhances the mechanical integrity of the electrode through hydrogen bonding interactions with the incorporated PAA polymer binder [21].

This is further substantiated by the capacity retention of the half‐cell under varying rate conditions. As shown in Figure 4e and Figure S13, rate capability tests for each anode material demonstrated significantly higher capacity retention across all C‐rates for EP‐SCN and MP‐SCN. Notably, at relatively higher rates of 1.0 C and 2.0 C, E‐SCN and M‐SCN exhibited lower capacity retention than the bare SCN anode. These results emphasize the critical role of a polymer‐based artificial SEI layer with high ionic conductivity in achieving improved performance. Additionally, upon returning to 0.2 C after cycling at 2.0 C, the rate retention of EP‐SCN and MP‐SCN reached improved values of 96.0% and 95.7%, respectively, in contrast to the 89.2%, 89.8%, and 91.5% retention observed for SCN, E‐SCN, and M‐SCN. This enhanced stability suggests that the pPD‐based supramolecular artificial SEI layer not only enhances ionic conductivity but also forms a highly dense layer, which contributes to the sustained integrity of the artificial SEI layer and improved electrolyte‐blocking capabilities on the SCN surface. Such properties serve to lithiaiton–delithiation processes and reinforce electrode stability over extended cycling.

2.5. Evaluation of Enhanced Li‐Ion Transport Characteristics Through the pPD Incorporation

To evaluate the uniform lithiation of the anode active material, surface potential measurements were carried out using atomic force microscopy (AFM), and the results were visualized through surface potential mapping (Figure 5a; Figures S14 and S15). As Li‐ion lithiation in SCN progresses, the work function of electrons is reduced, resulting in an increased surface potential [29, 30]. Using this principle, the surface potential mapping of samples lithiated to 10 mV at a 0.5 C rate in CC mode after the formation process, as depicted in Figure 5a and Figure S15, shows that SCN, E‐SCN, and M‐SCN exhibit insufficient and non‐uniform lithiation. In contrast, EP‐SCN and MP‐SCN display both sufficient and uniform lithiation. Moreover, the average surface potentials measured for SCN, E‐SCN, and M‐SCN were 2.19, 2.07, and 2.06 V, respectively, while EP‐SCN and MP‐SCN showed higher averages of 2.32 V and 2.39 V, respectively, indicating a higher degree of lithiation. This enhancement is attributed to the incorporation of pPD, which facilitates rapid and uniform Li ion transport to the active material by improving surface ion transport kinetics.

FIGURE 5.

(a) Surface potential mapping image of E‐SCN, M‐SCN, EP‐SCN, and MP‐SCN electrodes after lithiation under 0.5C constant current (CC) mode. Cyclic voltammetry of (b) SCN, (c) EP‐SCN, and (d) MP‐SCN electrodes. (e) Linear plots of maximum currents at corresponding square roots of scan rate based on the cyclic voltammetry curves of each electrode. Differential capacity analysis curves of SCN, (f) E‐SCN, EP‐SCN, (g) M‐SCN, and MP‐SCN after formation cycle.

Additionally, the diffusion characteristics of Li‐ions were investigated and measured at various scan rates using cyclic voltammetry (CV) (Figure 5b–d; Figure S16). From these analyses, the current released at approximately 0.05 V during lithiation and around 0.5 V during delithiation was compared across electrodes with the same loading levels under cathodic and anodic potential sweeps, with results presented in Figure 5e. As shown in Figure S17, at the same scan rate, EP‐SCN exhibited approximately 5.3 times the current of SCN and 2.4 times that of E‐SCN, while MP‐SCN showed about 6.3 times the current of SCN and 2.2 times that of M‐SCN. Furthermore, using the measured CV profiles, the Randles‐Sevcik equation was employed to evaluate Li‐ion diffusion performance during lithiation and delithiation phases. As illustrated in Figure 5g, EP‐SCN demonstrated slopes of ‐0.396 (lithiation) and 0.223 (delithiation), and MP‐SCN exhibited slopes of ‐0.428 (lithiation) and 0.223 (delithiation), both showing significantly higher values compared to SCN, E‐SCN, and M‐SCN. This enhanced Li‐ion diffusion behavior, facilitated by the incorporation of pPD which bridges Li‐ion hopping channels, enables faster Li‐ion transport within SCN, thereby contributing to improved activation and stability throughout extended cycling.

The voltages at which lithiation‐delithiation reactions reach their maximum were compared for each electrode using the differential capacity analysis (dQ/dV) plot. Improved Li‐ion diffusion reduces the overpotential during operation, leading to an increase in the lithiation reaction voltage and a decrease in the delithiation reaction voltage. This phenomenon reflects the improved kinetics and reduced resistance within the electrode, contributing to enhance electrochemical performance. As depicted in Figure 5f,g, during the second cycle immediately following the formation stage, the SCN electrode demonstrated the highest lithiation reaction peak at 0.025 V and the highest delithiation reaction peak at 0.475 V. For the E‐SCN electrode, the corresponding peaks were observed at 0.049 V for lithiation and 0.473 V for delithiation, while the M‐SCN electrode exhibited its lithiation and delithiation peaks at 0.052 V and 0.472 V, respectively. Notably, the EP‐SCN and MP‐SCN electrodes, which demonstrated the highest Li‐ion diffusion performance, displayed lithiation peaks at 0.060 V and 0.058 V, respectively, and delithiation peaks at 0.456 V and 0.469 V, respectively. These results indicate that the addition of pPD, acting as Li‐ion hopping channels, enables enhanced electrochemical kinetics and significantly improved ionic transport properties. This improved ionic conductivity facilitates a more uniform distribution of Li‐ion flux across the SCN anode, thereby promoting consistent lithiation–delithiation behavior.

2.6. Post‐Cycling Evaluation With Incorporating Supramolecular Artificial SEI Layer

To investigate the inhibitory effect of the artificial SEI layer formed by pPD addition on side reactions at the SCN surface, post‐cycling evaluation was conducted using EIS, XPS and XRD analyses. The resistance associated with the lithiation‐delithiation reactions on the electrode surface was measured via EIS test at 0 cycle, after 1 cycle, and after 100 cycles (Figure 6a–c). Initially, the resistance values were 187.1 Ω for SCN, 52.3 Ω for E‐SCN, 78.5 Ω for M‐SCN, 45.1 Ω for EP‐SCN, and 54.1 Ω for MP‐SCN. After the formation step at first cycle, when the active material was activated, overall resistance values decreased, with measurements of 28.3 Ω for SCN, 18.7 Ω for E‐SCN, 29.7 Ω for M‐SCN, 17.7 Ω for EP‐SCN, and 17.1 Ω for MP‐SCN, respectively. The substantially low resistance in EP‐SCN and MP‐SCN initially or after 1 cycle, when minimal electrolyte side reactions have occurred, can be attributed to the enhanced ionic conductivity facilitated by the addition of pPD. After 100 cycles, SCN, E‐SCN, and M‐SCN exhibited substantial increases in resistance, over twice the initial post‐formation values, reaching 59.8 Ω, 40.6 Ω, and 36.7 Ω, respectively. In particular, for SCN without a coating layer, the well‐defined emergence of two distinct semicircles in the EIS plot indicates the presence of two separate resistive layers. This suggests that on the SCN surface, Li‐ions undergo rapid desolvation and transport, while continuous surface SEI formation reactions induce the growth of a thick resistive layer, thereby increasing the interfacial resistance. In contrast, EP‐SCN and MP‐SCN exhibited stable resistance values, lacking the presence of two distinct semicircles, with recorded resistance values of 21.4 Ω and 20.3 Ω, respectively. The consistently lower resistance observed after 100 cyles for EP‐SCN and MP‐SCN can be attributed not only to the fast Li‐ion transport facilitated by pPD on the SCN surface but also to the enhanced density and mechanical strength of the artificial SEI layer. These improvements contribute to superior electrolyte‐blocking capabilities and structural stability, ensuring the durability of the coating layer throughout extended cycling.

FIGURE 6.

Comparison of EIS spectra obtained from the electrodes completing (a) before cycle, (b) 1 cycle, and (c) 100 cycles of cell operations. Comparison of XPS spectra obtained from the SCN, EP‐SCN, and MP‐SCN electrodes completed 50 cycles of cell operations in (d) C 1s and (e) F 1s regions. (f) The proportion of C containing group in the SEI layer and F containing group in the SEI layer. (g) Crystallographic changes of SCN, Ep‐SCN, and Mp‐SCN anodes at 50 cycles.

In the XPS analysis, the presence of the N 1s peak form pPD even after 50 cycles provides evidence that the PEO/pPD and PMMA/pPD artificial SEI layers remain on the SCN surface (Figure S18). As shown in Figure 6d,e and Figure S19, the deconvolution of the C 1s and F 1s peaks after 50 cycles for each electrode revealed a Li₂CO₃ peak at 289.8 eV, attributed to side reactions of the electrolyte, and a LiPO_xF_y peak at 686.8 eV, resulting from side reactions of the lithium salt anions [31]. In the absence of a protective coating layer, the SCN surface is continuously exposed to the electrolyte throughout cycling, leading to the repeated formation and breakdown of the SEI. This recurring process amplifies the incidence of side reactions, as evidenced by the pronounced Li₂CO₃ peaks in the C 1s spectra and LiPO_xF_y peaks in the F 1s spectra, highlighting the instability and inefficiency of bare SCN surfaces in mitigating undesirable interfacial reactions. In stark contrast, E‐SCN and M‐SCN, which are equipped with polymeric coating layers, exhibit a notable suppression of such side reactions. The presence of the polymeric layer acts as an effective barrier, restricting direct contact between the SCN surface, the electrolyte, and lithium salt anions during cycling. This stabilization effect is reflected in the significant reduction of both Li₂CO₃ and LiPO_xF_y peaks. Further advancements are achieved in EP‐SCN and MP‐SCN, where the incorporation of pPD leads to the development of a highly robust and densified artificial SEI layer. This enhanced interfacial coating layer further minimizes side reactions, as indicated by the markedly diminished intensities of the Li₂CO₃ and LiPO_xF_y peaks. The superior performance of EP‐SCN and MP‐SCN is quantitatively demonstrated in Figure 6f, where the area ratios of the deconvoluted Li₂CO₃ and LiPO_xF_y peaks, alongside other related peaks, distinctly indicate a significant reduction in side reaction activity. It should be noted that, the suppression of Li₂CO₃ and LiPO_xFᵧ formation is attributed to the combined effects of the barrier function of the pPD‐based artificial SEI layer and the enhanced Li⁺ dissociation behavior at the interface.

Furthermore, XRD analysis was conducted to monitor the crystalline growth of Li₂CO₃, allowing for a comparative evaluation of SEI layer formation resulting from electrolyte side reactions. In Figure 6g, prior to cycling, the electrodes from all samples displayed peaks at 26.5° and 28.5°, aligning with the theoretical peaks for graphite and Si, as shown in Figure S20. After 50 cycles, additional peaks corresponding to Li₂CO₃ were identified at 21.3°, 30.5°, 31.8°, and 33.8° in the SCN. Additionally, peaks at the same positions were observed in the E‐SCN and M‐SCN electrodes, indicating that simple polymer coatings such as PEO or PMMA are insufficient to provide an effective electrolyte‐blocking effect. In contrast, these peaks were absent in the EP‐SCN and MP‐SCN electrodes, indicating that the growth of Li₂CO₃ resulting from parasitic side reactions with the electrolyte was effectively suppressed on the SCN surface. These post‐cycling results demonstrate that pPD‐based modifications provide a sustained electrolyte‐blocking effect, effectively mitigating the persistent issue of increasing cell impedance in SCN electrodes during prolonged cycling.

2.7. Evaluation of Full‐Cell Performance Under Extremely Lean Electrolyte Conditions

The implementation of extremely lean electrolyte conditions (∼2 µL mAh⁻¹) establishes a stringent framework for evaluating cycle life performance of battery cells, with a particular focus on electrolyte consumption and the associated loss of available lithium ions during cell operation. Under these restrictive conditions, the continuous depletion of the electrolyte due to side reactions inevitably leads to an insufficient amount of electrolyte, hindering the effective transport of lithium ions to active materials. This issue manifests a significant increase in cell impedance, causing severe degradation in battery performance and an abrupt capacity fading over time. By monitoring the onset of rapid capacity decline under lean electrolyte conditions, it is possible to quantitatively assess the extent of side reactions at the electrode‐electrolyte interface. This approach offers insights into the direct impact of these reactions on overall cell performance. Importantly, incorporating protective coating layers on electrodes can mitigate electrolyte consumption by suppressing parasitic side reactions, thereby enhancing the long‐term stability of the cell. To validate the effectiveness of such coating layers, full‐cell tests were conducted under lean electrolyte conditions. These tests not only highlight the role of the coatings in minimizing electrolyte depletion during cycling but also establish their practical applicability for real‐world cell configurations [32].

The cathode utilized in this study was made from nickel cobalt aluminum oxide (NCA) with a total capacity of 5.83 mAh and a high loading level of 16.5 mg cm⁻². The anode was a blended electrode consisting of 90% graphite and 10% SCN, with an loading level of approximately 6.3 mg cm⁻². The anode was designed with a negative electrode total capacity of 6.41 mAh and an electrode compaction density of 1.5 g cm^{− 3}, resulting in an N/P ratio of 1.1 with respect to the positive electrode. During the initial formation cycles, the cells were charged to above 95% SOC at 0.1 C and to over 85% SOC at 0.2 C. Subsequently, during cycling at 1 C, the cells were operated after being charged to more than 65% SOC. The electrolyte was introduced in a total volume of 12 µL (2.05 µL mAh⁻¹) to match the overall capacity of the full cell, reflecting the stringent conditions of extremely lean electrolyte usage.

As shown in Figure 7a,b, the cycling performance analysis revealed that SCN‐based graphite‐blended full‐cells exhibited a rapid capacity drop around 300 cycles. In comparison, E‐SCN and M‐SCN cells showed capacity degradation between 500 and 550 cycles. Notably, cells with EP‐SCN and MP‐SCN demonstrated stable cycling performance for over 1000 cycles, with significant capacity fading observed only after approximately 1100 cycles. The observed capacity drop behavior in SCN‐based graphite‐blended full cells is primarily attributed to continuous electrolyte consumption driven by parasitic side reactions, which predominantly affect the SCN component, as opposed to the graphite which typically forms a stable SEI. Therefore, the absence of a coating layer in SCN‐based full cells led to the earliest occurrence of abrupt capacity drop. In contrast, EP‐SCN and MP‐SCN cells benefit from a robust pPD‐based supramolecular artificial SEI layer, which effectively blocks electrolyte access to the SCN surface. This protective layer significantly delays the occurrence of abrupt capacity drop by suppressing electrolyte decomposition reactions. Furthermore, the pPD‐involved SEI facilitates fast Li‐ion transport to the SCN, enabling both EP‐SCN and MP‐SCN based full cells to achieve higher capacity compared to other samples. The marked improvement in capacity and cycle stability under extremely lean electrolyte conditions is primarily due to the supramolecular artificial SEI layer formed by the addition of pPD. This layer not only enhances electrolyte‐blocking capabilities but also promotes fast and uniform Li‐ion transport characteristics, contributing to the overall enhanced performance of the cells.

FIGURE 7.

Cycle performance of SCN‐based graphite‐blended coin full‐cells: (a) E‐SCN and EP‐SCN, and (b) M‐SCN and MP‐SCN, plotted at 5‐cycle intervals to compare the electrochemical stability and cycling behavior of each sample under lean electrolyte conditions. Voltage profiles after the (c) first, (d) third, and (e) 100th cycles are also presented, illustrating the electrochemical behavior and performance differences across the samples under lean electrolyte conditions.

Additionally, the voltage profiles at the first, third, and 100th cycles reveal differences in capacity and overpotential during lithiation and delithiation, reflecting the influence of the specific SCN sample type utilized (Figure 7c–e). In the initial cycle, the initial Coulombic efficiencies of SCN, E‐SCN, M‐SCN, EP‐SCN, and MP‐SCN were 81.4%, 83.6%, 84.09%, 87.45%, and 88.66%, respectively, clearly indicating that the incorporation of pPD leads to a pronounced improvement in the initial Coulombic efficiency. Operating at a high loading level (6.3 mg cm⁻²) and a fast rate of 1C accentuates the differences in expressed capacity, which are primarily governed by the Li‐ion transport properties to each sample. Consequently, the capacities exhibited across all cycles show pronounced variation, with EP‐SCN and MP‐SCN demonstrating the highest capacities. This superior performance is attributed to the pPD‐based coating layer on the SCN surface, which facilitates Li‐ion transport. Differences in Li‐ion transport characteristics are also evident in the overpotential observed during lithiation and delithiation in each cycle. Inefficient lithiation and delithiation of SCN lead to increased polarization, manifesting as higher overpotential. This is clearly reflected in the voltage profiles, with SCN displaying the highest overpotential, while EP‐SCN and MP‐SCN exhibit significantly lower values. Furthermore, in the third and 100th cycles after the formation step, the overpotential variation observed during the 5‐min voltage relaxation after discharge is significantly lower for EP‐SCN and MP‐SCN. This indicates that rapid and uniform Li‐ion transport occurs on the SCN surface, effectively minimizing polarization.

3. Conclusions

To summarize, this study addresses the fundamental challenges inherent in conventional polymer‐based artificial SEI layers used in SCN‐based anodes, specifically the conflicting objectives of attaining complete surface coverage and securing a coating layer that promotes uniform and rapid Li‐ion transport while effectively blocking electrolyte penetration. In designing the coating layer, our strategy prioritized full coverage over SCN particles while enhancing the ionic conductivity of the coating to minimize the dependence of Li‐ion transport performance on polymer layer thickness as much as possible. To accomplish this, we incorporated pPD molecules into the polymer coating to develop a supramolecular artificial SEI layer. Notably, pPD acts as a bridging agent that simultaneously promotes Li‐ion desolvation, enables ultra‐fast Li‐ion transport, and improve mechanical properties. The pPD‐based supramolecular artificial SEI layer formed on the SCN surface integrates with coated PEO or PMMA to form an interconnected structure via extensive hydrogen bonding. This configuration promotes rapid Li‐ion transport at the SCN interface, enhances electrode capacity, and enables smoother lithiation‐delithiation processes. Furthermore, it effectively mitigates polarization and ensures the stable, efficient operation of SCN‐based electrodes, addressing key performance limitations inherently encountered in conventional approaches.

Furthermore, pPD, serving as a molecular bridging agent, not only effectively blocks electrolyte penetration but also enhances the elasticity of the complexed polymeric layer, making it well‐suited to accommodate the volume expansion of SCN. Consequently, the incorporation of the pPD‐based supramolecular artificial SEI layer substantially improved the cycling performance of full cells under extremely lean electrolyte conditions (2.05 µL mAh⁻¹). This sophisticated coating enabled the cells to achieve an impressive initial capacity of 6.3 mAh with the lowest overpotential, while maintaining sufficient electrolyte to ensure efficient Li‐ion transport, thereby supporting stable operation for over 1000 cycles. The superior performance highlights the dual functionality of the coating layer in both mitigating electrolyte depletion and facilitating efficient Li‐ion transport. Therefore, the introduction of this supramolecular network coating layer on the SCN surface effectively addresses critical challenges encountered at the anode interface. This innovative approach provides valuable insights for the future design of anode surfaces and has the potential to serve as a foundational platform for advancing next‐generation battery technology.

4. Experimental Section/Methods

4.1. Materials

p‐Phenylenediamine (pPD, M_W = 108.14 g mol⁻¹, ≥99.0%, SigM‐Aldrich), Poly(ethylene oxide) (PEO, M_V ≈ 100,000, SigM‐Aldrich), Poly(methyl methacrylate) (PMMA, M_V ∼98,500, SigM‐Aldrich), Bis(trifluoromethane sulfonyl)imide lithium salt (LiTFSI, 99.95%, SigM‐Aldrich), 1‐Methyl‐2‐pyrrolidone (NMP, 99.5%, SigM‐Aldrich), Si (CN vision, diameter ≈ 80 nm), Pitch (Dong‐Suh Chemical), Poly(acrylic acid) (PAA, M_W ≈ 250000, 35wt% in H2O, SigM‐Aldrich), Hexadecyltrimethylammonium bromide (CTAB, SigM‐Aldrich) Super P (Wellcos), 1.3 M LiPF6 in EC/DMC 30/70 v/v% + 5% FEC (Panaxetec), natural graphite (MTI Korea), Carboxymethyl cellulose sodium salt (CMC, SigM‐Aldrich, Mw ≈ 1 000 000), and Styrene‐butadiene rubber (SBR, MTI Korea)

4.2. Synthesis of Silicon Carbon Nanocomposite as Anode Materials

To prepare the silicon‐carbon nanocomposite (SCN), silicon nanoparticles are mixed with coal‐tar pitch, serving as the carbon source, at a weight ratio of 1:1 via ball‐milling. The resulting mixture undergoes thermal carbonization at 900°C with a heating rate of 5°C min^{− 1}. Following carbonization, the material is ball‐milled for 20 min and then sieved using a 635‐mesh screen. The final product is referred to as SCN.

4.3. Synthesis of SCN With Polymer‐Based Artificial SEI Layer

To coat SCN with PEO, 1 g of SCN was dispersed in 100 mL of deionized (DI) water along with 5 mg of CTAB. The mixture was subjected to bath sonication for 10 min. Subsequently, 100 mg of PEO was added to the solution, and the mixture was stirred for 24 h. After stirring, the solution was filtered using vacuum filtration and thoroughly washed with DI water. The resulting material was dried in a vacuum oven at 80°C to yield PE‐coated SCN, referred to as E‐SCN.

To coat SCN with PMMA, 1 g of SCN was dispersed in 120 mL of acetone, 20 mL of DI water, and 5 mg of CTAB. The mixture was subjected to bath sonication for 10 min. Subsequently, 100 mg of PMMA was added to the solution, and the mixture was stirred for 24 h. After stirring, the solution was filtered using vacuum filtration and thoroughly washed with acetone. The resulting material was dried in a vacuum oven at 80°C to yield PMMA‐coated SCN, referred to as M‐SCN.

For the preparation of EP‐SCN or MP‐SCN, SCNs were coated with a supramolecular complex of polymer‐pPD. During the polymer addition step, pPD was mixed with the polymer (PEO or PMMA) at a molar ratio of 12:1 (polymer:pPD) before being incorporated into the solution. The subsequent steps followed the respective procedures for E‐SCN or M‐SCN to obtain EP‐SCN or MP‐SCN.

4.4. Synthesis of Polymer and pPD‐Polymer Supramolecular film

500 mg of PEO or PEOpPD (mixed at a molar ratio of 6:1) was dissolved in 5 mL of deionized (DI) water and stirred for 24 h. After stirring, LiFSI was added to the solution to achieve a molar ratio of 14:1 (EO:Li). The resulting mixture was drop‐cast onto a PTFE plate and dried to form the electrolytic film.

500 mg of PMMA or PMMApPD (mixed at a molar ratio of 6:1) was dissolved in 5 mL of acetone and stirred for 24 h. After stirring, LiFSI was added to the solution to achieve a molar ratio of 14:1 (EO:Li). The resulting mixture was drop‐cast onto a PTFE plate and dried to form the electrolytic film.

4.5. Physicochemical Property Characterization

Structural characteristics of MOF filler was examined by Field‐Emission Scanning Electron Microscopy (FE‐SEM, JSM‐IT800, JEOL) with energy‐dispersive X‐ray spectroscopy (EDS). The crystalline analysis of the samples was measured by using an X‐ray diffractometer (D8 ADVANCE, Bruker, Germany) with Cu K radiation. Thermogravimetric analysis (TGA, TG/DTA7300, Seico inst.) was conducted at a rate of 10°C min⁻¹ under air atmosphere. To analyze the molecular structure, 700 MHz of 1H and 13C nuclear magnetic resonance equipment (NMR, AVANCEIII700, Bruker Corporation). To analyze of binding energy, X‐ray photoelectron spectroscopic (XPS, Thermo Fisher ESCALAB 250Xi) was conducted. The tensile test was performed on polymer films prepared via the solution casting method. The film used in the tensile test had a width of 25 mm and a length of 150 mm, prepared in accordance with the ASTM D882 standard. The samples were clamped at both ends, leaving a gauge length of 10 mm where the tensile force was applied. The test was conducted at a constant elongation rate of 10 mm min^{− 1} using a Universal Testing Machine (UTM, model 3343, Instron), and the results were recorded. The surface potential was measured using an Atomic Force Microscope (AFM, NX10) in Kelvin Probe Force Microscopy (KPFM) mode in dry room. The scan rate was set to 0.3 Hz to mapping of the surface potential distribution. To obtain cross‐sectional images of the electrodes, an ion beam cross‐section polisher (IB‐19530CP, JEOL) was employed for sample preparation. The polished cross‐sections were subsequently examined using scanning electron microscopy (EVO10, ZEISS). As the electrodes were in a lithiated state during characterization, all sample preparation and SEM measurements were conducted in a dry room (dew point < −50°C, temperature: 20°C) environment to prevent moisture‐induced side reactions.

4.6. Electrochemical Characterization

The ionic conductivity of the electrolytic film was evaluated using electrochemical impedance spectroscopy (EIS). For this measurement, coin cells were assembled by sandwiching the CSE membranes between two electrodes in an argon‐purged glovebox. The bulk resistances of electrolytic film was determined using EIS with two stainless‐steel electrodes. The EIS measurements were conducted at a 0.5 mV AC oscillation amplitude over a frequency range from 10⁵ to 10⁻¹ Hz. The ionic conductivity (𝜎) was calculated using the following equation:

$$\begin{array}{c}\sigma =\frac{L}{R\times S},\end{array}$$ (1)

where 𝜎 is the ionic conductivity, L is the thickness, S is the area of stainless‐steel electrodes, and R represents the resistance of the solid electrolyte. The activation energy was subsequently determined based on the ionic conductivity using the Arrhenius equation:

$$\begin{array}{c}{\sigma }_T={\sigma }_0\exp \left(-\frac{E_a}{RT}\right),\end{array}$$ (2)

where 𝜎_T is the ionic conductivity of the measured solid electrolyte, 𝜎₀ is the pre‐exponential factor, T is the absolute temperature, and R represents the gas constant.

4.7. Preparation and Measurement of SCN‐Based Anode for Half Cell Test

The anode slurry for half‐cell test, consisting of 70 wt% active material, 15 wt% Super P, and 15 wt% PAA, was coated onto a copper (Cu) foil substrate. The coated electrodes were dried in a vacuum oven at 120°C for 24 h. The active material loading was 1.5 mg cm⁻², and the electrode density was maintained at 0.5 g cm⁻³. The half‐cell tests were conducted using coin cells (CR2032 type). The anodes were punched into disks with a diameter of 8 mm and lithium metal was used as the counter electrode, and a microporous polypropylene (PP) film (Celgard 2400, Celgard) served as the separator. The electrolyte consisted of a 1.3 M solution of LiPF₆ in a 3:7 (v/v) mixture of ethylene carbonate (EC) and dimethyl carbonate (DMC), with fluoroethylene carbonate (FEC) added as an additive. The cell assembly process was performed inside a glovebox under an argon atmosphere with humidity maintained at ≤1 ppm. After resting at room temperature for 10 h to stabilize the cell, galvanostatic measurements were carried out using a multichannel potentiostat/galvanostat (WMPG 1000, WonATECH). The galvanostatic charge‐discharge testing was performed in the voltage range of 0.01‐1.5 V (vs. Li/Li⁺). The first cycle employed the constant current‐constant voltage (CC‐CV) mode at a current density of 0.1 C, followed by subsequent cycles at a current density of 0.5 C. Cyclic voltammetry (CV) was performed using half‐cells at various scan rates ranging from 0.1 mV s⁻¹ to 0.7 mV s⁻¹ in the potential range of 0.01‐1.5 V (vs. Li/Li⁺). The lithium diffusion coefficient was calculated based on the obtained CV data using the Randles‐Sevcik equation:

$$\begin{array}{c}{i}_p=0.4463nFAC\left(\frac{nF\nu D}{RT}\right)=\alpha D^{0.5}{\nu }^{0.5},(\alpha =0.4463nFAC{\left(\frac{nF}{RT}\right)}^{0.5}\end{array}$$ (3)

where i_p is the maximum current, n is the number of electrons involved in the redox reaction, A is the electrode area, F is the Faraday constant (C mol⁻¹), D is the diffusion coefficient, C is the concentration of lithium ions in electrolyte, ν is the scan rate (V s⁻¹), R is the gas constant (J K⁻¹ mol⁻¹), and T is the temperature (K). The galvanostatic intermittent titration technique (GITT) was conducted using half‐cells. To ensure stable cell conditions, the first two cycles were carried out at a current density of 0.1 C in the potential range of 0.01‐1.5 V (vs. Li/Li⁺). Following this stabilization, a current pulse of 0.1 C was applied for 10 min, followed by a relaxation period of 1 h. This sequence was repeated until the voltage reached the defined potential limits of 0.01 and 1.5 V (vs. Li/Li⁺). The diffusion coefficient form GITT data was calculated using the following Equation:

$$\begin{array}{c}{D}_{Li}=\frac{4}{\pi \tau }{\left(\frac{n_m{V}_m}{A}\right)}^2{\left(\frac{E_s}{E_t}\right)}^2\end{array}$$ (4)

where τ is the current applying time, n_m is the number of moles, V_m is the molar volume of the electrode, A is the area of used electrode, E_s is the steady‐state voltage change at a relaxation stage, and E_t is the voltage change while applying the constant current.

4.8. Preparation and Measurement of Graphite/SCN‐Blended Anode for Full Cell With Extremely Lean Electrolyte Testing (ELET) Conditions

In the case of graphite/SCN‐blended electrode, carboxymethyl cellulose (CMC) and styrene butadiene rubber (SBR) were used as binder. The composition of active materials for electrodes was 90:10 of a weight ratio for graphite:SCN. And the slurry was prepared with 96.5:1:1:1.5 of a weight ratio for active materials:Super P:CMC:SBR and the slurry was cast on a copper foil with doctor blade and dried at 120°C vacuum oven. The mass loading of electrode was controlled to be 6.3 mg cm⁻² and electrode density was set to be 1.5 g cm⁻³ for full‐cell tests. For cathode, commercialized lithium cobalt aluminum (NCA) was used. For coin‐type full‐cell tests, the electrodes were prepared by punching the anodes into 16 mm diameter disks and the cathodes into 15 mm diameter disks. The galvanostatic charge‐discharge tests were conducted within a potential range of 2.5–4.2 V (vs. Li/Li⁺) for the first and second cycles, with current densities of 0.1 C and 0.2 C, respectively. Starting from the third cycle, the potential range was adjusted to 2.5–4.0 V (vs. Li/Li⁺), and the current density was increased to 1 C.

4.9. DFT and MD Computational Details

All migration energy barrier calculations were performed based on ab initio calculations using the Vienna Ab initio Simulation Package (VASP 5.4.4) [33, 34, 35, 36]. The projector‐augmented wave (PAW) method was employed, and the exchange‐correlation energy was described using the Perdew–Burke–Ernzerhof (PBE) functional within the generalized gradient approximation (GGA) [35, 36, 37, 38, 39]. To accurately account for non‐bonding van der Waals interactions, Grimme's DFT‐D3 dispersion correction was incorporated [39, 40]. A plane‐wave basis set with energy cutoff of 450 eV was used for all calculations. To avoid interactions between periodic images, a vacuum spacing of at least 25 Å was applied in all directions. Due to the large size of the supercell, the Brillouin zone was sampled solely at the Gamma point. The atomic positions for the initial and final state structures were fully relaxed until the residual forces on each atom were less than 0.04 eV/Å, and the electronic self‐consistency loop was converged to 10⁻⁴ eV. The minimum energy path and the corresponding diffusion barrier were determined using the Climbing Image Nudged Elastic Band (CI‐NEB) methodology [41].

All Molecular Dynamics (MD) simulations were conducted using the Forcite module within the Materials Studio 2024 software package as mentioned in results and discussion part. The detailed composition of each simulated system is provided in Table 1. The simulation parameters, describing both bonded and non‐bonded interactions, were taken from the Condensed‐phase Optimized Molecular Potentials for Atomistic Simulation Studies (COMPASSIII) force field [42]. To maintain charge neutrality, a partial charge of +0.7e was assigned to each Li‐ion, with a corresponding ‐0.7e charge on each PF₆ ⁻ anion; this charge scaling was employed to better reproduce experimental observations [43, 44, 45].

TABLE 1.

Detailed composition of the simulation system.

	Q a	M b	N c	Repeat unit/Chain	Total number of Repeat unit
PEO	1.56 × 10−6	9253.15	3	210	630
PMMA	−5.11 × 10−6	21026.60	3	210	630
pPD	−1.00 × 10−7	108.14	105
Li+	+0.70	6.94	45
PF6 −	−0.70	144.96	45

^a Charge (e).

^b Molecular weight of molecule (g/mol).

^c The number of molecules loaded in the system.

Initial cells were constructed using the Amorphous Cell module [46, 47, 48]. The resulting amorphous cells were subsequently stabilized through a multi‐step protocol comprising geometry optimization, thermal annealing, and isothermal‐isobaric (NPT) equilibration. In the geometry optimization step, both the lattice constants and internal atomic positions were fully relaxed until the residual forces converged to below 0.001 kcal·mol⁻¹·Å⁻¹. The annealing procedure involved cyclic temperature variation between 300 and 500 K to overcome local energy minima and promote structural relaxation. Thereafter, the systems were equilibrated in the NPT ensemble for 1 ns at 353 K and 1 atm (1.013 × 10⁻⁴ GPa) to ensure convergence of both density and cell dimensions.

Following equilibration, a 10 ns production run was performed in the canonical (NVT) ensemble to investigate the transport behavior of Li‐ions. Understanding the Li‐ion diffusion mechanism at room temperature is computationally challenging, as it requires simulation times extending beyond the nanosecond scale. Therefore, to accelerate the diffusion process and gain clearer insight into the transport kinetics within a feasible timeframe, all simulations were conducted at an elevated temperature of 398 K. A time step of 1 fs was used, and the system temperature was controlled by a Nosé–Hoover thermostat [49]. Long‐range electrostatic interactions were treated using the Particle‐Particle Particle‐Mesh (PPPM) method, whereas van der Waals interactions were calculated with an atom‐based scheme and a real‐space cutoff distance of 18.5 Å. The diffusion coefficient (D) of Li⁺ ions was obtained from the mean squared displacement (MSD) according to the Einstein relation.

Conflicts of Interest

The authors declare no conflicts of interest.

Supporting information

Supporting File: smll73093‐sup‐0001‐SuppMat.docx.

Data Availability Statement

The data that support the findings of this study are available from the corresponding author upon reasonable request.