A Multifunctional Surface Modifier Capable of Stabilizing and Accelerating Hard Carbon Anode Kinetics via Pore Sealing and Competitive Adsorption

ABSTRACT

Hard carbon (HC), as the most suitable anode material for sodium‐ion batteries, still suffers from low Coulombic efficiency and unsatisfactory fast‐charge capability, mainly because of its numerous pores that result in severe electrolyte decomposition and sluggish Na⁺ reaction kinetics. Unlike the mainstream solid‐electrolyte interphase (SEI) strategy that solely focuses on suppressing the electrolyte decomposition or accelerating the reactions on HC anode surface, this work reveals the pore‐sealing effect of an artificial SEI through adopting a multifunctional organic small‐molecule salt (methyl blue, MB) as the surface modifier of HC and systematically elucidates its effects on the reaction kinetics of HC anode from the surface to bulk. The MB modification layer effectively seals the surface pores of commercial HC (Type‐2) and suppresses electrolyte decomposition. More importantly, theoretical calculations and empirical experiments reveal that MB accelerates the surface‐to‐bulk Na⁺ reaction kinetics of HC anode, including Na⁺ desolvation, adsorption on HC, transport across the SEI, and subsequent charge transfer and phase transition into Na clusters. Consequently, the MB‐modified commercial HC exhibits significantly enhanced Coulombic efficiency, superior rate capability and prolonged cycling life compared with unmodified HC.

An organic small‐molecule salt as the surface modifier of hard carbon effectively seals the surface pores, suppresses the electrolyte decomposition and accelerates the surface‐to‐bulk Na⁺ reaction kinetics of commercial hard carbon anode, thereby giving rise to significantly enhanced Coulombic efficiency, superior rate capability and prolonged cycling life in sodium‐ion batteries.

1. Introduction

Driven by the global pursuit of carbon neutrality and the growing demand for renewable energy storage, sodium‐ion batteries (SIBs) are regarded as a highly promising next‐generation electrochemical energy storage technology for large‐scale energy storage systems, owing to the abundance and low cost of sodium resources [1, 2, 3]. Among various anode materials, hard carbon (HC) has emerged as the most commercially promising anode for SIBs because of its unique amorphous structure, low voltage plateau versus Na/Na⁺, and high reversible capacity [4, 5, 6]. However, its practical implementation is still hindered by two critical bottlenecks. On the one hand, the abundant pores/defects and oxygen‐containing functional groups on HC surface continuously catalyze the decomposition of the electrolyte and induce irreversible Na⁺ consumption, leading to a low initial Coulombic efficiency (ICE) and deteriorated cycling stability [7, 8, 9]. On the other hand, the sluggish Na⁺ desolvation kinetics at the electrode/electrolyte interface, together with the tortuous Na⁺ diffusion paths and high transport barriers caused by the highly disordered structure of HC, result in severe electrode polarization, thereby restricting its fast charge/discharge capability [10, 11]. Particularly, the plateau capacity (<0.1 V vs. Na/Na⁺) of HC anodes often decays sharply with increasing current density, which has become a core bottleneck for the development of high‐power SIBs [12, 13, 14].

To overcome these challenges in HC anodes, tailoring the structure and chemistry of the solid‐electrolyte interphase (SEI) has been demonstrated to be highly effective [15, 16, 17, 18]. The SEI modulation can be mainly realized through two strategies. The first strategy relies on in‐situ electrochemical evolution of electrolyte components including solvents, Na salts and additives [19, 20, 21]. For instance, tris(2,2,2‐trifluoroethyl) phosphate (TFEP), as a co‐solvent in carbonate‐based electrolytes, was able to preferentially adsorb onto the HC surface and drag the salt anions, thereby forming a thin F‐rich SEI film [22]. Sodium trifluoromethanesulfonate (NaOTF) was employed as the sodium salt in ether‐based electrolytes to construct an OTF⁻ anions‐anchored solvation sheath around Na⁺ and the resultant SEI was enriched with inorganic NaF and Na₂O [23]. The fundamental mechanism of these newly developed electrolyte formulations is mainly to promote the preferential decomposition of anions, thereby constructing a low‐resistance inorganic SEI that inhibits excessive solvent decomposition at the HC surface and facilitates interfacial Na⁺ transport [24, 25, 26]. However, the electrochemical decomposition of the electrolyte components strongly depends on the active sites on HC anodes, making it difficult to regulate the uniformity of the as‐formed SEI. Consequently, continuous electrolyte consumption may occur, compromising the rate performance and long‐term cycling stability of SIBs.

The other strategy to regulate SEI is based on the pre‐modification of HC surfaces via artificial coating, which theoretically allows for the introduction of deliberately tailored SEI constituents and more uniform distribution of SEI [27, 28, 29, 30]. For instance, the surface of HC particles was grafted with 4‐(2,2,2‐trifluoroacetyl) benzoic acid, which was able to reversibly store Na⁺, promote the generation of a thin and NaF‐rich SEI, and boost the initial Coulombic efficiency to 90.0% [31]. Polymerized caffeic acid was in‐situ grafted onto the HC surface to introduce reversible carbonyl groups, which can catalyze the preferential reduction of salts, thereby forming a uniform and inorganic‐rich SEI [32]. In another report, the highly active oxygen‐containing groups of HC were converted into a Si‐O‐Si molecular layer, which was favorable for anions anchoring and contributed to the formation of an inorganic/organic hybrid SEI with reduced interfacial resistance and improved cycling stability [33]. Poly(methyl methacrylate), as an artificial organic SEI, was in‐situ coated onto a commercial HC to reconstruct the Na⁺ coordination environment at the inner Helmholtz plane and enhance the Na⁺ desolvation and interphase transport [34].

Although considerable progress has been made in modulating SEI, the influence of artificial coating on the surface pore structure of HC remains poorly explored, and the underlying mechanism governing the Na⁺ reaction kinetics from the surface to the bulk has not been clearly elucidated. Yu et al. [35]. has revealed that the closed pores smaller than 2 nm can increase the plateau capacity by facilitating the formation of stable Na clusters, while the closed pores larger than 2 nm are unfavorable for the stability of Na clusters, thereby reducing sodium storage capacity. Herein, an organic small‐molecule salt (methyl blue, MB) featuring a conjugated aromatic structure and sulfonate functional groups was deliberately selected as the surface modifier of HC. Notably, the MB modification layer can seal the pores above 2 nm on HC surface and enhance the initial Coulombic efficiency of commercial HC (type‐2). Moreover, the theoretical calculations and experimental characterizations reveal that the MB modification layer can accelerate the surface‐to‐bulk Na⁺ reaction kinetics of HC anode, including Na⁺ desolvation, Na⁺ adsorption on HC, Na⁺ transport across the SEI, and the subsequent charge transfer and phase transition into Na clusters. Consequently, the MB‐modified HC exhibits significantly improved Na‐storage performance especially at high charge/discharge rates compared with unmodified commercial HC.

2. Results and Discussion

The preparation procedure and surface structure of HC@MB are schematically shown in Figure 1a. The sulfonate (–SO₃ ⁻), imino (–NH‐) groups and conjugated aromatic rings of MB, could form ionic bonding, hydrogen bonding and π–π stacking interactions with the hydroxyl (–OH), carboxyl (–COOH) groups and localized graphite‐like layers of HC (Figure S1) [31, 33, 36]. Therefore, a layer of MB molecules is supposed to be uniformly anchored on the surface of hard carbon particles. In addition, from the perspective of modifying the performance of HC anode, the reason for adopting MB as the modifier is based on its sodium sulfonate groups and large conjugated π bonds composed of imino groups and aromatic rings. The former could not only compensate for the irreversible Na⁺ loss during charge/discharge of HC anode, but also generate a sulfonate‐based solid/electrolyte interphase (SEI) facilitating Na⁺ transport [36, 37]. The latter could promote the interfacial electron transfer and also facilitate the adsorption of Na⁺. Thus the overall charge transfer kinetics of HC@MB anode could be accelerated.

FIGURE 1.

(a) Schematic illustration of the preparation procedure and surface structure of HC@MB. (b) The adsorption energies of HC and MB towards the key electrolyte components. (c) MD simulation of 1 M NaPF₆‐TEGDME electrolyte. (d) The radial distribution function g(r) and coordination number of 1 M NaPF₆‐TEGDME electrolyte. (e) The solvation structure of Na⁺ in the electrolyte. (f) The energy barriers of the two‐step desolvation process of Na⁺[TEGDME]₂. (g) Schematic illustration of the functions of MB in modulating Na⁺ desolvation and adsorption on HC.

To evaluate the wetting behavior of the electrolyte on HC and HC@MB anodes, contact angle measurements were conducted (Figure S2). The contact angle of the electrolyte on HC@MB anode is smaller than that of unmodified HC, indicating that the MB layer improves the infiltration of the electrolyte into the electrode. To reveal the underlying mechanisms for this phenomenon, density functional theory (DFT) calculations were performed to determine the adsorption energies of HC and MB towards the key electrolyte components (Figure 1b). The adsorption energies of MB towards Na⁺, ${\mathrm{PF}}_6^{-}$ and TEGDME are generally greater the counterparts of HC, well explaining the improved infiltration of the electrolyte into HC@MB anode.

In addition, it can be seen that the adsorption energy of HC towards TEGDME molecules (–1.310 eV) is more negative than that of HC towards Na⁺ (–0.577 eV), indicating that the HC anode could preferentially adsorb the solvent molecules. The enrichment of TEGDME molecules on electrode surface would impede Na⁺ transport/adsorption at the interface. In contrast, the adsorption energy between MB and Na⁺ is as low as –3.536 eV, significantly more negative than that between MB and TEGDME (‐1.373 eV). Therefore, the HC@MB anode is supposed to preferentially adsorb Na⁺, which could effectively suppress the accumulation of TEGDME molecules at the interface and alleviate the barrier for Na⁺ migration towards the HC surface. In other words, the MB layer could contribute to accelerating the interfacial Na⁺ transport/adsorption.

By comparing the binding energy of Na⁺ with TEGDME (–1.309 eV), HC (–0.577 eV) and MB (–3.536 eV), it is expected that the MB layer could promote the desolvation of Na⁺ from the TEGDME solvation sheath at the electrode surface since MB exhibits stronger interaction with Na⁺ than TEGDME does. However, the unmodified HC could not achieve such effect because the binding energy of HC with Na⁺ is smaller than that of TEGDME.

The desolvation energy barriers of Na⁺ from the solvation sheath on HC and MB surface can be further calculated to support this deduction. Before doing this, the solvation structure of Na⁺ in the 1 M NaPF₆‐TEGDME electrolyte was analyzed through molecular dynamics (MD) simulations. The visualized structural model of the electrolyte is shown Figure 1c. The radial distribution function (RDF) (Figure 1d) shows a sharp and strong peak at r≈0.24 nm for Na–TEGDME(O), indicating a high coordination probability between Na⁺ and the O atoms of TEGDME at this distance. The Na–TEGDME(O) curve drops rapidly to the baseline at r≈0.37 nm, representing the boundary of the first solvation sheath. In contrast, a much weaker peak at r≈0.25 nm and a tiny broad peak at r≈0.45 nm are observed for Na–PF₆(F), suggesting that the distribution of ${\mathrm{PF}}_6^{-}$ around Na⁺ is far less concentrated compared to that of TEGDME. At r≈0.37 nm, the coordination number of Na⁺ with the O atoms of TEGDME and F atoms of ${\mathrm{PF}}_6^{-}$ is approximately 6.3 and 0.7, respectively, suggesting that Na⁺ is dominantly coordinated with TEGDME. Within the first solvation sheath, each Na⁺ is surrounded by two TEGDME molecules with each TEGDME providing 3–4 O atoms for coordination, while less one ${\mathrm{PF}}_6^{-}$ anion could be loosely coordinated with Na⁺ at the edge of the first sheath. Therefore, the dominant solvation structure in 1 M NaPF₆‐TEGDME is considered as Na⁺[TEGDME]₂ (Figure 1e) with Na⁺ coordinated with two TEGDME molecules.

Based on the solvation structure of the electrolyte, the desolvation process of Na⁺ was proposed and the corresponding energy barriers were calculated (Figure 1f). The Na⁺[TEGDME]₂ solvation structure undergoes two‐step stripping of TEGDME molecules at the electrode interface. The energy barriers (ΔE₁) required to strip one TEGDME molecule, forming Na⁺[TEGDME] and one free TEGDME molecule on the HC and HC@MB surfaces, are 1.687 and 0.781 eV, respectively. Subsequently, Na⁺[TEGDME] sheds the remaining solvent molecule, and the corresponding energy barriers (ΔE₂) for this step are 0.217 and −0.060 eV for HC and HC@MB, respectively. The desolvation energy barriers on HC@MB surface are obviously lower than those on pristine HC, verifying that the MB layer can effectively accelerate the desolvation of Na⁺, which is a crucial step in determining the electrochemical reaction kinetics of HC anode.

A schematic is shown in Figure 1g to illustrate the function of MB modifier on HC deduced from theoretical calculation results. On the one hand, the MB layer could effectively facilitate the desolvation of Na⁺ from the electrolyte by reducing the desolvation energy barrier. On the other hand, it favors the preferential adsorption of Na⁺ and inhibits solvent adsorption, generating an interface favorable for Na⁺ transport.

The successful coating of MB layer on the surface of HC particles are evidenced by the following various characterizations. The N₂ adsorption–desorption isotherms of HC and HC@MB are shown in Figure 2a. The introduction of MB layer significantly reduces the specific surface area from 4.7590 m² g^{− 1} of HC to 2.1296 m² g⁻¹ of HC@3%MB (Figure 2c and Table S1). With the increase of MB amount, the specific surface area further decreases, reaching 1.9524 m² g^{− 1} for HC@7%MB. The pore size distribution curves in Figure 2b show a significant reduction in the mesoporous range (2–30 nm) for the HC@MB samples. The obviously decreased total pore volume and mesopore volume (Figure 2c and Table S1) indicate that MB forms a dense coating layer and seals part of the open pores on HC surface. In this way, the MB layer could serve as a protective shield to reduce the unfavorable electrolyte decomposition at the HC surface.

FIGURE 2.

(a) N₂ adsorption/desorption isotherms, (b) pore size distribution curves, (c) BET specific surface area and pore volume of HC and HC@MB. (d) XRD patterns of HC, MB and HC@MB powders. (e) XPS survey spectrum and (f) high‐resolution S 2p core‐level spectra of HC and HC@3%MB. (g–h) TEM images of HC and HC@3%MB. (i) Zeta potentials of HC and HC@3%MB in water. (j) EDX elemental mapping of HC@3%MB.

X‐ray diffraction (XRD) patterns of HC, MB and HC@MB are compared in Figure 2d and S3. Besides the typical broad peaks of HC, HC@MB displays distinct diffraction peaks of MB at 2θ ≈ 19°, 32°, 34°, and 45°. Fourier transform infrared (FTIR) spectra of HC@MB (Figure S1) also show the characteristic absorption peaks of MB at 1500–1600 cm⁻¹, 1000–1400 cm⁻¹, 700–800 cm⁻¹ and 600–700 cm⁻¹, which are attributed to the stretching vibration of C = C bond in benzene ring, the stretching vibration of the S = O bond, the out‐of‐plane bending vibration of C–H bond in benzene ring, and the stretching vibration of S–O bond in MB molecule, respectively. X‐ray photoelectron spectroscopy (XPS) survey spectrum of HC@3%MB clearly shows the characteristic elemental peaks of MB, namely Na 1s, N 1s, and S 2p (Figure 2e) compared to that of HC. In the core‐level region of S 2p (Figure 2f), the doublet located at 168.2 and 169.5 eV corresponds to the presence of sulfonate groups.

The morphology of the MB layer on the HC surface was observed using scanning electron microscopy (SEM) and transmission electron microscopy (TEM). SEM images of HC and HC@MB are presented in Figure S4. The surface of pristine HC particles exhibits some nanoscale debris. As the MB coating amount increases, these debris become blurred gradually and the HC surface looks progressively smoother, indicating that MB successfully forms a coating on the HC surface. TEM images (Figure 2g,h) reveal that the HC@MB surface is covered by a continuous and dense MB layer with an average thickness of 14 nm. This coating layer could act as a physical barrier to mitigate the decomposition reaction of electrolyte.

Zeta potentials of HC and HC@MB (Figure 2i) in water were measured to reveal their surface charge states. Zeta potential of HC@MB is –44.5 mV, while that of HC is only –2.24 mV. The sulfonate groups of MB impart a strong negative charge to HC@MB, further suggesting the successful coating of MB layer on HC surface. The strong negative charge of HC@MB will be beneficial for its dispersion stability in the electrode slurry, thereby ensuring electrode uniformity and structural integrity. Energy‐dispersive X‐ray spectroscopy (EDS) elemental mapping (Figure 2j) shows that the distribution regions of S and O elements highly coincide with the shape of HC particle, once again verifying the continuous coating of MB layer on HC surface.

The electrochemical performance of HC and HC@MB anodes with varying coating ratios were evaluated using sodium metal as the counter electrode. The charge–discharge profiles during the first three cycles at a current density of 0.1 C are shown in Figures S5 and S6. In the first three cycles, the optimal HC@3%MB anode delivers Coulombic efficiency of 90.28%, 98.23% and 98.32%, significantly surpassing the counterparts of HC anode (86.32%, 92.96% and 93.12%). Remarkably, the MB coating layer effectively covers the open pores or defects on the HC surface, thereby suppressing the irreversible electrolyte decomposition and active Na⁺ consumption. Remarkably, the MB coating layer effectively covers the open pores or defects on the HC surface, thereby suppressing the irreversible electrolyte decomposition and active Na⁺ consumption. In addition, the sodium sulfonate groups within the MB molecule also compensate for partial active Na⁺ loss and promote rapid Na⁺ transport across interface, further increasing the efficiency for Na⁺ storage. The initial Coulombic efficiency (ICE) comparison of HC@3%MB with recently reported HC anodes (Figure 3a) further demonstrates the competitive advantage of MB modification [38, 39, 40, 41, 42, 43, 44, 45, 46].

FIGURE 3.

(a) The initial Coulombic efficiency comparison of HC@3%MB with recently reported HC anodes. (b) Rate capability of HC and HC@MB anodes. (c) The rate behavior comparison of HC@MB with recently reported HC anodes. (d) Rate capability of HC||NVP and HC@3%MB||NVP full cells. (e) The capacity evolution of HC and HC@3%MB anodes at 0.5 C under different temperature conditions. (f) Na‐storage mechanism of HC anodes. (g–h) The contributions of the slope and plateau capacity for HC and HC@3%MB anodes at different C‐rates. (i–j) CV curves and (k–l) the corresponding contour plots of HC and HC@3%MB anodes at varying scan rates.

The reversible capacities of HC and HC@MB anodes with the increase of charge‐discharge rates are compared in Figure 3b. At 0.1 C, 0.5 C, 2 C and 5 C, the capacities of HC@3%MB anode reach 295.8, 266.3, 225.1 and 94.6 mAh g⁻¹, respectively, all of which are higher than those of pristine HC (290.5, 258.7, 145.5 and 58.1 mAh g⁻¹). The capacity retention advantage of HC@MB becomes especially pronounced at high C‐rates (Figure S7). As the MB coating amount further increases (5%, 7%), the rate performance exhibits a decreasing trend. This may be attributed to the excessively thick MB layer, which prolongs the interfacial ion transport pathway and increases the diffusion resistance of Na⁺ across the artificial interface layer. The optimal coating amount of 3% MB achieves the best balance between surface protection and interfacial ion transport. The superior rate capability of HC@MB anodes indicates that the MB coating effectively enhances the reaction kinetics of HC anode and alleviates polarization under high‐current conditions.

The rate behavior comparison of HC@MB with recently reported HC anodes (Figure 3c) further demonstrates the remarkable efficacy of MB modification in improving kinetics [38, 39, 40, 41, 42, 43, 44, 45, 46]. To evaluate the practical potential of the HC@3%MB anode, full cells were assembled using Na₃V₂(PO₄)₃ (NVP) as the cathode and their performance is shown in Figures 3d and S8–S10. In terms of rate performance, the HC@3%MB||NVP full cell manifests higher capacity at all the C‐rates ranging from 0.1 C to 5 C. In contrast, the HC||NVP full cell suffers from more serious capacity decay with the increase of C‐rates.

The good rate behavior of HC@MB anodes at room temperature inspired us to study its performance at lower temperatures. The capacity evolutions of the anodes at 0.5 C under lower temperature conditions (Figures 3e and S11) show that HC@3%MB continues to exhibit higher capacity than pristine HC. At –10 °C, HC@3%MB still preserves a reversible capacity of 97.5 mAh g⁻¹, significantly higher than 50.5 mAh g⁻¹ of HC. It is evident that the MB modification layer effectively improves the electrode reaction kinetics even at low temperatures.

To reveal the origin of the enhanced capacity retention at high current density, the charge–discharge curves of HC and HC@MB anodes at different C‐rates are displayed in Figure S12. As the current density increases, the capacity of pristine HC especially the plateau capacity at the voltage near 0 V decreases substantially. In contrast, HC@3%MB exhibits a much higher capacity predominantly contributed by the plateau capacity.

Among the several mechanisms proposed for Na‐storage in HC anode, the adsorption–insertion–pore filling mechanism is widely regarded as the most plausible one [47, 48, 49, 50]. Ge et al. validated the applicability of this mechanism and revealed a dynamic evolution pathway, which is schematically illustrated in Figure 3f [51]. In the slope region of 1.5–0.15 V, Na⁺ adsorbs on the defect sites and intercalates into the graphite‐like interlayers to form NaC_x intercalation structures. In the early plateau region of 0.15–0.05 V, the adsorbed/intercalated Na⁺ ions transition to quasi‐metallic Na clusters within closed pores. In the late plateau region of 0.05–0.01 V, the progressive growth of Na clusters occurs, while the released defect/intercalation sites are re‐occupied by adsorbed/intercalated Na⁺. This “release–transition” storage pathway attributes the plateau capacity to the transition of intercalated/adsorbed Na⁺ into quasi‐metallic clusters inside the closed pores, rather than to direct pore filling. This means that the adsorption/intercalation rate of Na⁺ on HC surface in the slope voltage region of 1.5–0.15 V will directly determine the magnitude of the plateau capacity especially under fast charge‐discharge conditions.

Figure 3g,h compare the contributions of the slope and plateau capacity for the HC and HC@3%MB anodes at different current densities. It can be seen that the plateau capacity contribution of HC drops sharply at high C‐rates, indicating severely limited kinetics for the transition of Na⁺ into sodium clusters. In contrast, HC@3%MB maintains a relatively high proportion of plateau capacity even at 3 C, confirming that the MB modification effectively promotes the interfacial charge transfer kinetics by facilitating Na⁺ adsorption/intercalation and accordingly alleviating the Na⁺ transition hysteresis under fast charge‐discharge.

The optimized kinetics of MB coating was further verified by cyclic voltammetry (CV). CV curves (Figure 3i,j) and the corresponding contour plots (Figure 3k,l) at various scan rates show that the redox peaks of HC@3%MB near 0 V are sharper along with a significant enhancement in peak currents compared to those of HC, indicating faster reaction kinetics. In addition, compared to HC anode, the reduction current signal of HC@3%MB representing the Na⁺ adsorption reactions starts at obviously higher potentials as marked by the rectangular region, which can be attributed to the MB‐induced Na⁺ adsorption bringing about an earlier onset of reduction reactions.

To investigate the function of MB in modifying Na⁺ adsorption, the b values (Figure 4a) were calculated by analyzing the relationship between peak current (i) and scan rate (v) of the CV curves in the voltage range of 0.1–1.5 V (Figure S13). The b value close to 1 indicates a predominantly surface‐controlled capacitive process, while b value close to 0.5 reflects diffusion‐limited behavior. The b values of the HC@3%MB anode are 0.68 (Peak 1) and 0.75 (Peak 2), obviously higher than the corresponding values of HC anode (0.55 and 0.59), suggesting that a more prominent capacitive process in the presence of MB. The capacity contributions from capacitive‐controlled and diffusion‐controlled reactions at each scan rate (Figure S14) further reveal that the capacitive contributions of HC@3%MB surpass those of HC across all tested scan rates.

FIGURE 4.

(a) The b values of HC and HC@3%MB anodes obtained through the linear fitting of log(i_p ) vs. log(v) plots. (b) Cycling performance of HC and HC@3%MB anodes at 2C. (c) Nyquist plots and (d) DRT plots of HC and HC@3%MB anodes after 3 and 100 cycles. (e–f) Nyquist plots of HC and HC@3%MB anodes at different temperatures. (g) The calculated activation energy of the interfacial reactions in HC and HC@3%MB anodes. (h–i) In‐situ DRT plots of HC and HC@3%MB anodes. (j) The Na⁺ diffusion coefficients of HC and HC@3%MB anodes during charge. (k–l) The work function of HC and HC@3%MB anodes obtained from UPS.

The cycling stability of HC and HC@MB anodes was evaluated at a current density of 2 C (Figures 4b, S15, and S16). The HC@3%MB anode exhibits the highest capacity and optimal capacity retention. After 200 cycles, it retains a reversible capacity of 222.2 mAh g^{− 1} with a capacity retention of 98.20%, significantly outperforming the pristine HC anode (90.17%). After 500 cycles, HC@3%MB maintains a capacity retention of 96.98%, whereas that of HC is only 81.82%. The charge–discharge curves at different cycle numbers (Figure S17) further indicate that the profiles of HC@3%MB anode overlap well from the first to 500th cycle, in sharp contrast with the fast capacity decay of HC anode during cycling.

Electrochemical impedance spectroscopy (EIS) was employed to further investigate the electrode reaction kinetics. Figure 4c presents the Nyquist plots of HC and HC@3%MB anodes after 3 formation cycles and after 100 cycles. Based on the equivalent circuit model shown in Figure S18, the fitted resistances for the SEI film (R_SEI ) and charge transfer (R_CT ) are summarized in Figure S19 and Table S2. Overall, the HC@3%MB anode exhibits significantly lower R_CT than HC after both 3 and 100 cycles, indicating that the MB modification layer effectively enhances the charge transfer efficiency and maintains excellent interfacial stability during the long‐term cycling.

To further decouple the interfacial processes, distribution of relaxation times (DRT) analysis was performed on the EIS data (Figure 4d). The peak at τ ≈ 10⁻⁵ s corresponds to the ohmic resistance (R_S ). The peaks in the range of 10^−4.5∼ 10⁻³ s are assigned to ion transport within SEI (R_SEI ). The peaks in the range of 10⁻³∼ 10⁰ s correspond to the charge transfer process (R_CT ). Compared with HC, HC@MB presents a slight reduction in R_SEI and a pronounced decrease in R_CT , further validating the significant improvement in interfacial kinetics achieved by MB modification.

EIS of HC and HC@3%MB anodes at different temperatures was measured (Figure 4e,f) to quantitatively analyze the activation energy (E_a ) of the interfacial reactions through Arrhenius equation. The E_a value of HC@3%MB is only 44.16 kJ mol^{− 1}, significantly lower than the 69.01 kJ mol^{− 1} of HC (Figure 4g), indicating that the MB layer effectively reduces the energy barrier required for interfacial reactions.

In‐situ DRT analysis (Figure 4h,i) at different depth of discharge further reveals that HC@3%MB maintains a considerably lower R_CT than HC throughout the sodium storage process, providing direct evidence of its consistently optimized interfacial charge transfer kinetics. To evaluate the Na⁺ diffusion behavior in the electrode bulk, galvanostatic intermittent titration technique (GITT) measurements were performed on HC and HC@3%MB anodes. After each galvanostatic pulse, the electrode was relaxed to a steady state, and the corresponding potential relaxation curves were recorded (Figure S20a). The calculated Na⁺ diffusion coefficients (Figures 4j and  S20b) show that HC@3%MB possesses significantly faster Na⁺ transport rate than HC during the discharge‐charge.

To further elucidate the microscopic mechanism behind the reduced R_CT in HC@3%MB, ultraviolet photoelectron spectroscopy (UPS) was performed on the anodes after three formation cycles. As shown in Figure 4k,l, the work function of HC is 5.88 eV, while that of HC@3%MB decreases to 5.35 eV. The decreased work function indicates that the MB layer can modulate the electronic structure of the anode surface and enhance the electron‐donating capability. As a result, MB facilitates the charge transfer reactions across the electrode/electrolyte interface.

To investigate the interfacial morphological evolution of the HC and HC@MB anodes, TEM and SEM characterizations were performed after 100 cycles. As shown in Figure 5a,b, the HC anode exhibits an uneven and thick SEI layer with a thickness of approximately 175 nm, whereas the HC@MB anode shows a uniform and dense SEI layer with a thickness of only about 38 nm. SEM images (Figures 5c,d and S21) further reveal that the cycled HC particles are covered with obvious agglomerated decomposition products, while the surface of HC@MB appears relatively smooth with fewer deposits.

FIGURE 5.

(a–b) TEM images, (c–d) SEM images, (e–f) AFM 3D views, and (g–h) Young's modulus of HC and HC@3%MB anodes after 100 cycles at 2C. (i–l) The constant‐capacity overdischarge curves of HC and HC@3%MB anodes. (m–p) Cross‐sectional SEM images of HC and HC@3%MB anodes after constant‐capacity overdischarge.

To gain insights into the three‐dimensional morphology and mechanical properties of the electrode surface, atomic force microscopy (AFM) was employed. The 3D views (Figure 5e,f) show that the surface of HC@3%MB anode presents considerably reduced height fluctuations and decreased roughness compared with that of HC, indicating the formation of a more uniform and flat SEI. The average Young's modulus (Figure 5 g,h) of HC@3%MB (4.57 GPa) is slightly higher than that of HC (4.46 GPa), reflecting the enhanced mechanical strength of SEI on HC@3%MB.

Since the formation of Na dendrites on the anode surface will be highly detrimental to the performance and safety of SIBs, uniform sodium deposition on the anode surface is essential in the events of serious polarization or battery overcharge. So the constant‐capacity overdischarge tests (Figure 5i–l) were conducted on the HC and HC@MB anodes after three formation cycles with Na as the counter electrode. When the HC and HC@MB anodes are discharged with the capacity reaching around 309 mAh g^{− 1}, the sodium deposition reaction commences (Figure S22). The overpotential of the sodium deposition on HC@3%MB (16.1 mV) is higher than that of HC (13.6 mV), indicating the MB layer can delay sodium plating.

The top‐view SEM images of the electrodes (Figure S23) after overdischarge show that the surface of HC is covered with numerous granular sodium, whereas that of HC@3%MB seems smoother and contains less granular sodium. The cross‐sectional SEM images (Figure 5m–p) further reveal that the sodium deposited on HC surface is highly uneven and appears fluffy, whereas that on HC@3%MB seems uniform and dense. The morphological difference confirms that the MB layer significantly regulates the sodium deposition behavior and guides uniform sodium plating. The improvement in sodium deposition behavior can be attributed to the synergistic modulation of the interfacial electronic/ionic transport and SEI structure. On the one hand, the MB layer enhances the electron donating ability of the electrode surface by reducing the work function, which facilitates uniform electric field distribution at the interface and further guides Na⁺ to form a uniform ionic flux across the electrode surface. In this way, the MB layer can prevent local charge accumulation and avoid heterogeneous nucleation hotspots for sodium plating. On the other hand, the uniform structure and enhanced mechanical strength of MB‐derived SEI favors the in‐plane deposition and inhibits vertical growth of sodium.

X‐ray photoelectron spectroscopy (XPS) analysis further elucidates the regulatory effect of the MB layer on the chemical composition of SEI. Figure 6a summarizes the relative elemental contents of HC and HC@3%MB anodes after 100 cycles. The C elemental content in the SEI of HC@3%MB (71.17%) is higher than that of HC (60.95%), which can be attributed mainly to the carbon‐rich aromatic rings of MB. The combined content of O, F, Na and other elements on HC@3%MB accounts for 28.83%, significantly lower than the corresponding sum (39.05%) on HC, indicating that the MB layer effectively suppresses the electrolyte decomposition.

FIGURE 6.

(a) Atomic proportions of the key elements in the SEI of HC and HC@3%MB anodes. (b–h) XPS spectra of HC and HC@3%MB anodes after 100 cycles at 2C in the core‐level regions of S 2p, C 1s, O 1s and F 1s.

High‐resolution XPS spectra of S 2p, C 1s, O 1s, F 1s, Na 1s, and P 2p are presented in Figures 6b–h and S24. The S element existing in the form of –SO₃ ⁻ (168.4 and 169.7 eV) is detected on HC@3%MB anode, further confirming the participation of MB in SEI formation. In the C 1s core‐level region, the intensities of the peaks corresponding to C–O (285.5 eV), C = O (287.6 eV), O–C = O (289.4 eV), and CO₃ ²⁻ (290.2 eV) groups are obviously reduced for the HC@MB anode, suggesting the MB layer mitigates the decomposition of solvent. The significant decrease in the intensities of F 1s, O1s, Na 1s, and P 2p peaks in HC@3%MB further demonstrates that MB effectively inhibits the electrolyte decomposition and contributes to the formation of a thin SEI.

In addition, it is interesting to find out that fluorine on the HC@3%MB anode surface exists predominantly as NaF (685.2 eV), whereas the relative intensity of NaF on HC anode is very weak. NaF, as a common inorganic SEI component, not only possesses high mechanical strength and chemical stability but also facilitates Na⁺ conduction [52, 53, 54]. The relatively high content of NaF on HC@3%MB anode can be ascribed to the MB layer triggering the preferential adsorption of ${\mathrm{PF}}_6^{-}$ over TEGDME solvent. As shown in Figure 1b, the adsorption energy of ${\mathrm{PF}}_6^{-}$ on MB is stronger than that of TEGDME, so ${\mathrm{PF}}_6^{-}$ could approach the electrode surface and participate in the reduction reactions prior to TEGDME. These results indicate that the MB modification layer not only regulates the SEI chemistry but also helps establish a more protective and kinetically favorable interfacial layer dominated by NaF, providing favorable conditions for rapid and stable Na⁺ migration.

To gain a deeper insight into the structural changes of MB during the electrochemical cycling, MB was used as the active material and fabricated into the working electrode, with carbon black (CB) as the conductive agent. A pure CB electrode serving as a control sample was also prepared. CV curves (Figure 7a) of pure CB electrode during the first negative scan exhibit one distinct reduction peak at ∼0.7 V, which corresponds to the Na⁺ adsorption reactions on the electrode surface. The prominent reduction peaks at ∼0.05 V and oxidation peaks at ∼0.08 V are attributed to the reversible Na⁺ plating/stripping in CB electrode. Compared with pure CB electrode, besides the above redox peaks, MB electrode (Figure 7b) displays two new reduction peaks at ∼1.54 V and ∼0.12 V during the first scan, which obviously result from the electrochemical reduction reaction of MB. During the second and third cycling, the two new reduction peaks almost disappear, suggesting that the electrochemical reaction of MB in the first cycle is most probably irreversible.

FIGURE 7.

(a–b) CV curves of CB and MB electrodes during the first three cycles at 0.05 mV s⁻¹. (c–d) Ex‐situ XRD patterns and FTIR spectra of the MB electrode at different states of charge/discharge. (e) The electrochemical reaction mechanism of MB molecule. (f) Schematic illustration of the multiple functions of MB modification layer in improving the performance of HC anode.

To clarify the reaction mechanism of MB, ex‐situ XRD analysis (Figure 7c) was performed on the MB electrodes at different charge–discharge states. All the MB electrodes exhibit the characteristic diffraction peak of MB at 2θ around 32°, with no peak shift observed. Other noticeable peaks can be ascribed to Al foil. It is deduced that the overall crystalline structure of MB remains unchanged during the charge–discharge. Ex‐situ FTIR spectroscopy was used to investigate the changes in the organic functional groups of MB (Figure 7d). Upon discharge to 0.01 V, the characteristic imine (C = N) peak at ∼1600 cm^{− 1} disappears, indicating that the C = N bond in the MB molecule is reduced to C–N bond. Such a reduction can induce an electron cloud rearrangement within the triphenylmethane moiety, leading to the broadening of the benzene ring skeletal vibrations at ∼1600 cm^{− 1}. Furthermore, the reduction of C = N alters the shifts of out‐of‐plane bending vibration of aromatic C–H in the benzene ring and gives rise to a new and intense peak at ∼850 cm^{− 1}. Another newly appeared peak at ∼560 cm^{− 1} can be attributed to the formation of –N⁻–Na⁺ bond through the reduction of C = N bond.

In the pristine MB, sulfonate groups exist in two chemical environments (–SO₃ ⁻Na⁺ and –SO₃H), so the asymmetric stretching vibration of S = O splits into a doublet at ∼1350 and ∼1400 cm⁻¹. After discharge to 0.01 V, the doublet merges into a single peak at ∼1420 cm⁻¹, suggesting the unification of the chemical environment of S = O. For the MB electrode charged to 1.5 V and the one after 3 cycles, their FTIR adsorption peaks are similar to those of the electrode discharged to 0.01 V, indicating that the reduction of MB is irreversible and the structure remains stable in subsequent cycles. Based on the variation of the functional groups, the electrochemical reaction mechanism of MB is presented in Figure 7e.

When the HC@MB anode is immersed in the TEGDME‐based electrolyte, MB dissolves slightly as shown by the light blue color of the electrolyte (Figure S25). Nevertheless, it is found that the electrolyte become colorless after the electrochemical cycling. Based on the CV curves of MB electrode (Figures 7b), ex‐situ XRD patterns (Figure 7c) and FTIR spectra (Figure 7d) of MB electrodes, and XPS spectra of the cycled HC@MB electrode (Figure 6b), it is deduced that MB is reduced irreversibly at the anode during the first discharge, acts as an SEI component, and keeps structural stability after the first discharge. So in spite of its slight dissolution before cycling, the reduced MB exhibits good chemical stability and stays at the HC anode after the first discharge in the TEGDME‐based electrolyte.

Figure 7f schematically illustrates the multiple functions of MB modification layer from the perspectives of interfacial adsorption, Na⁺ desolvation surface, pore sealing and SEI optimization. On the surface of HC anode, the strong adsorption of TEGDME solvent molecules leads to solvent accumulation at the interface, making the Na⁺ desolvation kinetics sluggish. Moreover, the exposed surface pores of HC and the preferential adsorption of TEGDME onto HC result in serious electrolyte decomposition, and generate a thick and non‐uniform SEI film with high organic content and low ionic conductivity, impeding the charge‐transfer reactions. The sluggish Na⁺ kinetics at the anode surface will correspondingly lead to slow transition into Na clusters in the HC bulk. In contrast, on the HC@MB anode, the preferential adsorption of Na⁺ by the MB modification layer significantly accelerates the Na⁺ desolvation. Simultaneously, the MB layer exhibits stronger adsorption affinity toward ${\mathrm{PF}}_6^{-}$ than toward TEGDME solvent, thereby promoting the generation of inorganic components dominated by NaF. In addition, the MB layer effectively seals the surface pores of HC, suppresses the solvent decomposition, and gives rise to the formation of a thin, highly Na⁺‐conductive and structurally stable SEI. The significantly enhanced interfacial kinetics and stability accordingly accelerate the transition reactions of Na⁺ into Na clusters in the HC bulk.

3. Conclusion

Methyl blue is selected as the surface modifier and artificially coated on the surface of commercial HC particles to enhance the Na‐storage performance. The conjugated aromatic structure and sulfonate functional groups of MB enable its uniform anchoring on HC surface, which can effectively seal the surface pores of HC and suppress the electrolyte decomposition, thereby increasing the initial Coulombic efficiency of HC anode. In addition, the MB modification layer exerts a strong attraction toward Na⁺, promoting its desolvation and preferential adsorption at the electrode interface. The enriched Na⁺ further attracts ${\mathrm{PF}}_6^{-}$ anions, resulting in the formation of a thin, stable and low‐resistance SEI. As a result, the subsequent charge‐transfer reactions and the transition of Na⁺ into Na clusters within the HC bulk are also accelerated. The MB‐modified HC anode manifests more stable capacity retention and impressive rate performance compared with unmodified commercial HC. This work not only sheds lights on the critical roles of artificial SEI layers in regulating the surface pores and surface‐to‐bulk Na⁺ kinetics of HC anodes, but also provides a convenient and effective interfacial engineering strategy for fabricating high‐performance HC anodes.

Funding

National Natural Science Foundation of China (NSFC Nos. 21875156 and 22279085).

Conflicts of Interest

The authors declare no conflicts of interest.

Supporting information

Supporting File: smll74647‐sup‐0001‐SuppMat.docx.

Data Availability Statement

The datasets in the current study are available from the corresponding author on reasonable request.