Tuning S-doping in Porous Carbon for Enhanced Initial Coulombic Efficiency and Rate Performance for Sodium/Potassium-Ion Batteries

Amorphous carbon has emerged as a promising anode material for sodium-ion (SIBs) and potassium-ion batteries (PIBs) due to its high specific capacity, abundant defects, and low production cost. Nevertheless, its practical application remains hindered by suboptimal rate capability and low initial Coulombic efficiency (ICE). Heteroatom doping, particularly with sulfur, has proven to be an effective strategy for addressing these limitations. In this study, sulfur was successfully incorporated in a controlled amount into the amorphous porous carbon framework via a combination of low-temperature carbonization and a facile fumigation process. The introduction of sulfur and the low-temperature treatment synergistically induced a higher density of structural defects, which contributed to significantly enhanced rate performance of the sulfur-carbon (S/C) composite. Furthermore, controlled sulfur doping facilitated the formation of a thinner and more stable solid electrolyte interphase (SEI) during the initial cycling, thereby minimizing irreversible sodium consumption and leading to a substantial improvement in ICE. As an anode for SIBs, the optimized S/C composite delivers a high reversible capacity of 480 mAh g-1 at 0.1 A g-1 and maintains a capacity of 232.6 mAh g-1 even at 5.0 A g-1. Notably, an initial Coulombic efficiency of 83.1% is achieved under a current density of 0.1 A g-1. When applied as the anode for PIBs, the composite exhibits a reversible capacity of 439.2 mAh g-1 and an ICE of 61.2% at the same current density. This work provides a viable and scalable approach to simultaneously enhance the rate capability and ICE of amorphous carbon. 1. Introduction In recent years, SIBs and PIBs have garnered significant attention due to the abundance of resources and their cost-effectiveness. The similar physicochemical properties between Na + /K + and Li + enable analogous electrochemical reactions, positioning SIBs/PIBs as promising alternatives to lithium-ion batteries (LIBs) [1,2]. However, the larger ionic radius of Na + /K + compared to Li + imposes more stringent requirements on the selection and structural stability of anode materials [3]. As a result, commercial graphite, which is widely used in LIBs anodes, is no longer suitable for SIBs and PIBs [1,2]. Currently, the most commonly explored anode materials for SIBs/PIBs include carbon-based materials [4], alloy-type [5], titanates [6] and conversion-type materials [7]. Nevertheless, alloy-type materials and conversion-type materials suffer from severe volume expansion during cycling [5,7], whereas titanates are hindered by low electrical conductivity and low energy density-both factors limiting their practical application [6]. Therefore, carbon-based materials are considered the most promising candidates for SIB/PIB anodes [1,2]. Among them, amorphous carbon has received considerable interest due to its high specific capacity and abundant defects. Despite these advantages, the low ICE and suboptimal rate performance of amorphous carbon continue to hinder its practical application and commercialization in SIBs/PIBs [8]. Thus, it is critical to develop amorphous carbon materials with both high rate performance and high ICE. To address the aforementioned challenges, researchers have proposed various optimization strategies, including heteroatom doping [9], modulation of the structural disorder in carbon frameworks [10], and the construction of nanostructures [11]. Li et al. prepared amorphous carbon anode materials using activated carbon as a precursor through high-temperature annealing and surface coating techniques [10]. The resulting material exhibited a unique microstructure characterized by high internal porosity and a dense surface carbon layer. This architecture effectively prevents electrolyte infiltration and maximizes the utilization of internal storage sites, thereby enhancing the ICE of the material. When applied as an anode for SIBs, the material demonstrated an ICE of 80% at a current density of 0.4 C. However, the employed high-temperature annealing and coating methods did not improve its rate capability, retaining only a reversible capacity of 65 mAh g -1 at a high current density of 8.0 C. Zhong et al. developed a nitrogen-doped hollow carbon nanotube network using MoO 3 as a template and a nitrogen-containing pyrrole-based carbon precursor via pyrolysis [11]. The interconnected network facilitated the formation of an efficient electron transport pathway, thereby enhancing the transport kinetics of both electrons and Na + . As an SIBs anode material, it delivered a discharge capacity of 208 mAh g -1 at 100 mA g -1 . Even after a 50-fold increase in current density, the material maintained 52.4% of its original capacity. However, the ICE at 100 mA g -1 remained below 40%, indicating significant initial irreversible capacity loss. Li et al. also synthesized a sulfur-rich amorphous carbon material via a simple calcination process using 1,4,5,8-naphthalenetetracarboxylic dianhydride (NTCDA) and elemental sulfur as precursors [9]. The material exhibited a high reversible capacity of 360 mAh g -1 at 0.05 A g -1 and retained 158 mAh g -1 even at 4 A g -1, demonstrating excellent rate performance. The substantial sulfur incorporation significantly enhanced the rate capability. However, the ICE remained low at only 63.2% under a current density of 0.02 A g -1 . The above strategies have proven effective in improving the internal porosity and electrical conductivity of amorphous carbon, thereby enhancing either the ICE or rate performance while improving cycling stability. Nevertheless, these approaches have yet simultaneously improving both the ICE and the rate capability of amorphous carbon materials. Therefore, the development of amorphous carbon anodes with both high ICE and excellent rate performance remains a critical challenge in the field of SIBs/PIBs research. In this study, we developed a S/C composite by combining low-temperature carbonization with a simple vapor-phase S infusion technique, using lignite-derived amorphous porous carbon as the carbon framework (Figure 1a). The introduction of S not only increases the defect density and active sites but also enhances the electronic conductivity of S/C composite. Additionally, the controlled introduction of S helps modulate the specific surface area of the S/C composite and induced the formation of a thinner, stable SEI film during initial charge-discharge cycles. These combined effects significantly enhance both the ICE and rate capability of S/C composite. As an anode material for SIBs, the S/C composite delivered an outstanding reversible capacity of 480 mAh g -1 at 0.1 A g -1, high ICE of 83.12% at 0.1 A g -1, and maintained 232.6 mAh g -1 at 5.0 A g -1 . When applied in PIBs, it also demonstrated a high reversible capacity of 439.2 mAh g -1 at 0.1 A g -1 . This study proposes a facile S doping strategy to simultaneously address the key challenges of suboptimal rate performance and low initial coulombic efficiency typically associated with amorphous carbon anodes. 2. Results and discussion To systematically investigate the microstructure and structural properties of C and S/C composite materials, we characterized them using scanning electron microscopy (SEM) and transmission electron microscopy (TEM). Figures 1b and 1c show that pure C materials exhibit a honeycomb-like porous structure. The morphology of the S/C-16% composite differs significantly from that of pure C. High-magnification SEM images (Figures 1e, 1f) reveal that the surface of the S/C-16% composite is covered with numerous nanoparticles, increasing surface roughness. SEM images of S/C composites with varying S contents (Figure S1) show that as S content increases from 14% to 18%, both the number and size of nanoparticles on the surface of the carbon sheet layers increase. These results suggest that S-doping in the carbon matrix directly influences the surface morphology and roughness of the composites [12]. However, when the S content increases to 18%, the size of the nanocrystals significantly increases, gradually filling the carbon pores and leading to the collapse of some porous structures. HRTEM images of the porous carbon material (Figure 1d) show no clear lattice fringes or diffraction spots, indicating its amorphous nature [13]. HRTEM images (Figure 1g) show that the S/C-16% composite also lacks clear lattice fringes and distinct diffraction spots. Energy Dispersive X-ray Spectroscopy (EDS) element distribution maps (Figure 1h) illustrate the spatial distribution of C (red) and S (yellow) elements, indicating that carbon forms a continuous porous network with S element uniformly dispersed throughout [12]. The two elements are highly interwoven and uniformly distributed, reflecting the strong interaction between the C and S phases. The intact porous carbon framework provides rapid electronic and ionic transport pathways, mitigating sulfur’s volumetric expansion and thereby enhancing the long-cycle stability of this electrode material [14]. All samples display typical type IV nitrogen adsorption-desorption isotherms (Figure S2a) with H3-type hysteresis loops [15], confirming their mesoporous nature. As the S content in the S/C composites increases, their nitrogen adsorption capacity decreases progressively. This trend suggests that the introduction of S significantly reduces the specific surface area and pore volume of the composites. The corresponding pore size distribution curves (Figure S2b) show that the pore sizes of all samples are primarily concentrated in the 0-20 nm range, indicating that both pure carbon and S/C composites consist mainly of micropores (< 2 nm), with some mesopores (2-50 nm). However, as the S content increases from 0% to 18%, the proportion of micropores gradually decreases, suggesting that S incorporation leads to the collapse of the microporous structure, and a reduction in the number of micropores [16]. As shown in Table S1, with increasing S content, the pore structure of the composites gradually changes, with both specific surface area and pore volume decreasing progressively and tending to stabilize at higher S contents, indicating that the pores are approaching S filling saturation. Figure 1. (a) Illustration of the synthesis strategy of S/C composite material via low-temperature carbonization and vapor-phase S infusion technique; SEM images of pure C at (b) low magnification and (c) high magnification, and (d) HRTEM and SEAD image; SEM images of S/C-16% composite at (e) low magnification and (f) high magnification, (g) HRTEM image of S/C-16% composite, and (h) EDS elemental mapping of S/C-16% composite. Elemental analysis (EA) indicates that as the S content in the S/C composites increases from 14% to 18%, the carbon content correspondingly decreases (Figure 2a). The small amount of oxygen is likely derived from the organic precursor used in the synthesis. Figure S3 show the TGA curves of C and all S/C samples under Ar atmospheres. Based on the results of thermogravimetry and elemental analysis, we can calculate that the contents of elemental S in the S/C-14%, S/C-16% and S/C-18% composite materials are 1.3%, 3.4% and 5.1% respectively, while the valence-bond S in them are 12.4%, 12.5% and 12.8% respectively (Figure 2a). Ultraviolet (UV) spectroscopy was used to characterize the optical properties of C and all S/C samples (Table S2). Ultraviolet analysis indicated elemental S contents of 0%, 1.32%, 3.38%, and 5.13%, respectively, consistent with the thermogravimetric analysis results. The X-ray Diffraction (XRD) patterns (Figure 2b and S4) show two broad peaks at 23° and 43°, corresponding to the (002) and (100) planes of amorphous carbon [13]. As S content increases, the intensity of the diffraction peaks decreases slightly. Meanwhile, the two diffraction peaks at 23° and 43° gradually shift towards lower angles and become increasingly prominent. However, the peak shapes remain almost unchanged, suggesting that S incorporation does not affect the microstructure of the carbon matrix [17]. The defect levels of C and all S/C composite samples were further characterized via Raman spectroscopy. The D and G bands were deconvoluted into four distinct sub-peaks located at approximately 1200, 1350, 1500, and 1580 cm -1, corresponding to the D4, D1, D3, and G bands, respectively. These sub-peaks are attributed to the vibrations associated with sp2-sp3 hybridized structures (D4), disordered graphitic lattices (D1), short-range sp3 carbon (D3), and ideal graphitic domains (G) [18,19]. Among them, the intensity ratio of the D1band to the G band (I D1 /I G ) serves as an indicator of the overall structural disorder [20]. The I D1 /I G values for the C, S/C-14%, S/C-16%, and S/C-18% samples are 1.57, 1.60, 1.88, and 1.75, respectively (Figure 2c and Figure S5). As the S content increases, the I D1 /I G ratio first rises and then falls, suggesting a non-linear evolution of defect density associated with S incorporation. This trend suggests that an optimal level of S doping (S/C-16%) effectively enhances the structural defects and disorder in amorphous carbon, thereby increasing the active sites available for sodium ion storage [21]. However, when the S content exceeds a certain threshold (S/C-18%), the I D1 /I G ratio decreases, likely due to the aggregation of excess S atoms and S-S interactions, which suppress defect formation in the carbon backbone and ultimately degrade the electrochemical performance of the electrode material [21]. X-ray Photoelectron Spectroscopy (XPS) spectra (Figure 2d) show only C and O in the C sample, while the S/C-16% composite displays distinct S peaks, consistent with the results of elemental analysis and FTIR (Figure S7), further confirming the successful integration of S into the porous C. As S content increases from 14% to 18% (Figure S6), the S peak intensity increases progressively, consistent with the added S content. In the XPS C 1s spectrum (Figure 2e and S8a-c), the five main peaks correspond to the following species: sp2 (284.4 eV), C-H (284.9 eV), sp3 (285.5 eV), C-C (286.2 eV), O=C-O bonds (287.5 eV) [22]. As the S content increases, the sp2 peak intensity decreases markedly. This suggests that the defect density in the S/C electrode materials increases progressively with higher S incorporation. [23]. For all S/C samples, the S 2p spectra (Figure 2f and Figure S8d-e) display double peaks at 162.9 eV and 164.3 eV, corresponding to S 2p 3/2 and S 2p 1/2, associated with elemental S, while the peaks at 166.5 eV and 168.1 eV correspond to C-SO x -C bonds [17]. As S content increases, the intensity of the S peaks and C-SO x -C signals rises, suggesting that the S in the S/C samples mainly consists of elemental S and valence bond S. The appearance of C-SO x -C indicates partial oxidation of S ether during synthesis or storage [9]. For pure C (Figure S8f), the O 1s spectrum is fitted with four peaks: C-O (530.4 eV), C=O (531.2 eV), C-OH (532.1 eV), and adsorbed -OH (533.2 eV). After S incorporation, the O 1s spectrum exhibits significant changes [24]. For S/C-14%, S/C-16%, and S/C-18%, the O 1s spectra display four peaks, with a new S-O peak appearing at 530.4 eV (Figure S8g-i). The formation of S–O bonds can be attributed to the presence of C-SO x -C bonds in the S/C samples. As S content increases, the relative intensity of the S-O peak gradually increases, confirming the formation of S-O bonds and further supporting the presence of the C-SO x -C group, in line with the S 2p results. Figure 2 . (a) EA of C and S/C composites; (b) XRD patterns, (c) Raman spectra, and (d) XPS survey spectra of C and S/C-16% composites; High-resolution XPS spectra of (e) C 1s and (f) S 2p. To investigate the sodium storage mechanism of the electrode materials, CV measurements were carried out on the C and S/C-16% electrodes at a scan rate of 0.1 mV s -1 within the voltage range of 0.01-3.0 V (Figure 3a, b). As observed in Figure 3a, a broad cathodic peak appears in the potential range of 0.75-0.01 V, which corresponds to the insertion of Na + into the C framework. A distinct anodic peak near 3.0 V is associated with the extraction of Na + from the C matrix [25]. No additional redox peaks are observed throughout the entire voltage window. In contrast, the CV profile of the S/C-16% electrode exhibits a reduction peak around 2.2 V, which can be attributed to the electrochemical reaction between Na + and S, forming sodium polysulfides such as Na 2 S 3, Na 2 S 2, and Na 2 S [26]. The cathodic peak at approximately 1.1 V corresponds to the further reduction of sodium polysulfides to Na 2 S [27]. The broad reduction feature in the 0.75-0.01 V range is associated with the formation of a SEI film [25]. Additionally, the anodic peak near 1.8 V is related to the desodiation process and the decomposition of Na 2 S into Na + and elemental S [27]. Figure 3c and Figure S9a present a comparative evaluation of the rate capabilities of the C electrode and S/C electrodes with varying S contents. Under identical current densities, all S/C electrodes demonstrate higher reversible specific capacities than the undoped C electrode. Notably, the reversible capacity of S/C electrodes exhibits a non-monotonic dependence on S content, initially increasing and subsequently decreasing as the S content increases, suggesting the existence of an optimal S doping level. Specifically, the S/C-16% electrode delivers reversible capacities of 480, 436.6, 386.7, 346.6, 312.5, 282.6, 251.3, 240.9, and 232.6 mAh g -1 at current densities of 0.1, 0.2, 0.4, 0.6, 0.8, 1.0, 2.0, 3.0, and 5.0 A g -1, respectively. In contrast, the C electrode exhibits significantly lower capacities under the same conditions, with values of 194.1, 171.1, 142.9, 116.9, 101.3, 60.0, and 22.3 mAh g -1 . Moreover, the SIBs employing the C electrode undergoes structural failure at a current density of 3.0 A g -1, indicating limited high-rate tolerance. Remarkably, the S/C-16% electrode exhibits excellent rate reversibility, with its capacity recovering to 427.1 mAh g -1 when the current density is returned from 5.0 to 0.1 A g -1 . These results underscore that S incorporation into the carbon framework introduces additional defect sites, enhances capacitive contributions to the overall charge storage [24], and S doping is likely to improve the electrical conductivity and facilitate Na + transport kinetics within the electrode matrix [28]. The synergistic effect of these structural and electrochemical modifications leads to a pronounced enhancement in rate performance. When benchmarked against other reported heteroatom-doped porous C materials (Figure 3d), the S/C-16% electrode exhibits a clear performance advantage, especially at high current densities [9, 11, 29-40]. The cycling stability of the electrodes at a current density of 0.1 A g -1 is shown in Figure 3e and Figure S9b. The S/C-16% electrode maintains a substantially higher reversible capacity than the other electrodes throughout the cycling test, retaining 577.5 mAh g -1 after 100 cycles. By contrast, the C electrode only delivers 124.2 mAh g -1 under identical conditions. These results confirm that moderate S doping substantially enhances the long-term cycling stability of C electrodes. However, a further increase in S content results in a decline in reversible capacity (Figure S9b). In addition to improved capacity and cycling performance, S doping also significantly enhances the ICE of the porous C material. The S/C-16% electrode achieves an ICE of 83.12% at 0.1 A g -1, significantly higher than that of the C electrode (29.98%) and other S/C composites (Figure 3f). This result also compares favorably with those of previously reported heteroatom-doped carbon anodes for SIBs (Table S3). The galvanostatic charge-discharge profiles of the S/C-16% electrode (Figure 3g) exhibit highly overlapping curves after the initial cycle, indicating excellent electrochemical reversibility and structural stability. The S/C-16% electrode delivers an initial discharge and charge capacity of 747.1 and 621.0 mAh g -1, respectively (Figure S9c), which are significantly higher than those of all other samples. From the third cycle onward, the Coulombic efficiency remains nearly 100% over 100 consecutive cycles (Figure S10), further attesting to its robust reversibility. To further assess the long-term cycling durability of the S/C electrodes, extended galvanostatic cycling was conducted at 1.0 A g -1 (Figure 3h and Figure S9d). The S/C-16% electrode delivers a substantially higher reversible capacity than both the C electrode and other S/C variants, maintaining 349.2 mAh g -1 even after 1000 cycles. These findings highlight the excellent structural integrity and electrochemical resilience of the S/C-16% electrode, underscoring its potential as a high-performance anode material for sodium-ion batteries. Figure 3. Electrochemical performance of sodium -ion battery anodes: (a,b) CV curves of C and S/C-16% electrodes; (c) Rate performance of C and S/C-16% electrodes; (d) Comparison of rate performance from different references; (e) Cycling of C and S/C-16% at 0.1 A g -1 ; (f) ICE of S/C with different S contents; (g) Charge/discharge of S/C-16% at 0.1 A g -1 ; (h) Long-term cycling stability of C and S/C-16% at 1.0 A g -1 . To elucidate the underlying mechanism by which the S/C-16% electrode achieves a favorable balance between substantial capacity and high ICE, the evolution of the SEI film on both S/C-16% and S/C-18% electrodes was systematically investigated by HR-TEM after various cycling intervals (Figure 4a and b). For the S/C-16% electrode, the SEI thickness was determined to be approximately 5.1 nm after 5 cycles and increased moderately to 7.0 nm after 20 cycles. The SEI film formed on S/C-16% exhibits a relatively uniform morphology with minimal fluctuation in thickness, indicative of stable interfacial properties and efficient Na + transport [41]. In stark contrast, the SEI on the S/C-18% electrode exhibited a more pronounced increase in thickness, from 4.8 nm after 5 cycles to 9.94 nm after 20 cycles, revealing a significantly higher growth rate. Moreover, the SEI formed on the S/C-18% electrode was substantially less uniform, displaying localized regions of excessive thickness. This inhomogeneous SEI formation is likely attributable to the non-uniform distribution of S within the electrode, stemming from the excessive S content [42]. Such compositional heterogeneity can induce spatially localized and intensified electrochemical reactions during the initial cycles, thereby promoting the formation of a thicker and more heterogeneous SEI [43]. It is worth noting that the formation of a thicker SEI inevitably consumes a larger quantity of Na + during the initial cycling stages, leading to increased irreversible capacity loss and reduced ICE [25]. Consequently, the S/C-16% electrode, owing to its more controlled SEI growth and uniform interfacial characteristics, delivers a significantly higher ICE relative to the S/C-18% electrode. Figures 4c and 4d depict the in situ electrochemical impedance spectroscopy (In-situ EIS) measurements conducted during the first charge-discharge cycle for the S/C-16% and S/C-18% electrodes, respectively. In the high-frequency region, the semicircular arc corresponds to the resistance of the SEI ( R SEI ), whereas the medium-frequency semicircle is associated with the charge-transfer resistance ( Rct ). In the low-frequency region, a linear Warburg-type impedance is observed, indicative of Na + diffusion within the electrode matrix [44]. Detailed analysis of these impedance components across the voltage profile enables a comprehensive understanding of the interfacial reaction pathways and associated electrochemical kinetics. For the S/C-16% electrode (Figure 4c), a relatively high impedance is observed in the high-voltage range (3.0-2.5 V), primarily attributed to the initial formation and growth of the SEI film. As the potential decreases to an intermediate voltage window (approximately 2.0-1.0 V), a pronounced reduction in impedance is evident, signifying the stabilization of the SEI and progressive activation of the electrode material, which facilitates the formation of additional Na + storage sites. Upon further discharge to the low-voltage region (< 0.5 V), the impedance reaches its minimum, reflecting enhanced electronic and ionic conductivity due to extensive Na + intercalation [45]. Importantly, during the subsequent charging process, the impedance variation remains relatively smooth and continuous, suggesting a highly reversible interfacial evolution and robust structural stability. In comparison, the S/C-18% electrode (Figure 4d) exhibits a qualitatively similar impedance evolution trend; however, the absolute impedance values are markedly higher throughout the entire voltage range. Notably, in the high-voltage region (3.0-2.0 V), R SEI is significantly elevated, indicating more pronounced SEI formation. Furthermore, the impedance response of S/C-18% is characterized by sharper fluctuations, with distinct local maxima observed at specific voltages (2.5 V and 1.5 V), implying heterogeneous interfacial reactions [46]. During the subsequent charge process, a substantial increase in total impedance is observed, reflecting deteriorated interfacial reversibility and less stable SEI behavior, likely due to the non-uniform S distribution and associated side reactions. Figures 4e and 4f present the quantitative fitting results of the Nyquist plots based on the equivalent circuit model (Figure S11) in bar graph form, highlighting the comparative evolution of R SEI and Rct for the two electrodes. While both systems exhibit similar overall trends, the S/C-18% electrode consistently displays higher resistance values and greater fluctuations across both interfacial components. These findings provide direct evidence that excessive S content leads to interfacial instability and compromised reaction kinetics. In contrast, the S/C-16% electrode achieves a lower and more stable interfacial resistance profile, indicating a more favorable balance between electrochemical activity and interfacial integrity, which is essential for the development of high-performance sodium-ion battery anode materials [41]. Figures 4g and S12a present TOF-SIMS analyses, which elucidate the spatial distribution of various ionic species within the SEI film of the S/C-16% and S/C-18% electrodes, respectively. Each dataset includes both surface projection maps and 3D volume-rendered images, providing intuitive visualization of key fragment ions, including NaS -, ClO -, SO 3 -, CHO 2 -, NaCO 3 -, and NaCl 2 -, at the electrode–electrolyte interface. For the S/C-16% electrode (Figure 4g), the surface mapping reveals a relatively homogeneous distribution of NaS - signals, with moderate-intensity red features, indicating a controlled degree of interaction between sodium ions and S species [47]. The signals for ClO - and SO 3 - are weaker and more sparsely distributed, suggesting that the formation of oxygen and sulfur-oxygen species is relatively limited during cycling. CHO 2 - and NaCO 3 - signals exhibit medium intensity and uniform spatial distribution, implying moderate formation of carbonate-based SEI components [48]. The NaCl 2 - signal appears weak and localized to a few confined regions. These observations are further corroborated by the 3D volumetric reconstructions, which confirm the overall homogeneity and compositional moderation of the SEI formed on the S/C-16% electrode. In contrast, the S/C-18% electrode (Figure S12a) exhibits substantially different ion distribution characteristics, with significantly enhanced signal intensities and pronounced spatial heterogeneity for all detected species. These features indicate the formation of a thicker, more compositionally complex SEI film, which likely impedes the formation of a compact and stable interface, ultimately contributing to the observed deterioration in ICE and overall electrochemical performance. Specifically, the NaS - signal in the S/C-18% electrode is markedly intensified, as evidenced by the prevalence of high-brightness regions in the surface map and the increased density of red markers in the 3D volume image [47]. This directly reflects a higher degree of Na-S compound formation induced by the elevated S content. An excessive accumulation of NaS - implies substantial S leaching from the active S/C matrix during cycling, leading to undesirable side reactions that compromise the electrochemical reversibility and ICE [49]. Moreover, the ClO - signal also displays a significant increase in both intensity and coverage area, indicating a higher concentration of chlorinated species within the SEI [50]. Similarly, the SO 3 - signal exhibits stronger intensity and a broader spatial spread, revealing pronounced S oxidation. Such oxidative degradation not only consumes electrochemically active material but may also lead to the accumulation of electronically insulating byproducts, thereby further degrading electrode performance [49]. The CHO 2 - signal in the S/C-18% electrode is not only stronger in magnitude but also markedly uneven in distribution, with several localized regions of high intensity. The NaCO 3 - signal exhibits similar aggregation behavior, suggesting the formation of abundant but unevenly distributed carbonate species-consistent with the previously observed SEI inhomogeneity from HR-TEM analyses [48]. Furthermore, NaCl 2 - signals are significantly enhanced, with multiple high-intensity domains appearing in both the surface and volumetric maps [50]. Overall, the SEI film formed on the S/C-18% electrode contains a more diverse and concentrated array of ionic species compared to that of S/C-16%. The universally higher intensities of all detected ions in the S/C-18% electrode strongly indicate a more complex and aggressive interfacial reaction environment [51]. While such reactions may be driven by the elevated S content, their severity ultimately undermines SEI uniformity and structural integrity, thereby exerting a detrimental effect on the electrochemical performance of the S/C anode. To gain a deeper understanding of the electrochemical reaction mechanisms and structural evolution of the C electrode and the S/C-16% electrode during the charge-discharge process, in-situ Raman spectroscopy was conducted on both electrodes (Figure 4h and Figure S12b). As the discharge process proceeds, the intensity of the D-band in the S/C-16% electrode gradually decreases without significant shift in its position. This phenomenon can be attributed to the interaction between S and Na +, as well as the adsorption of Na + at defect sites, pores, and the edges of the carbon surface within the composite [52]. These results indicate that a sodium storage mechanism dominated by adsorption is operative in the charging and discharging process. Concurrently, the G band intensity also decreases progressively during discharge, which can be ascribed to the continuous intercalation of Na +, resulting in electron occupation of the π* antibonding orbitals and the consequent weakening of C-C bonds [19,53]. Notably, during the subsequent charging process, both the D and G bands gradually recover to their initial states, indicating good reversibility of the S/C-16% electrode during cycling. In contrast, the in-situ Raman response of the C electrode shows markedly different behavior. During discharge, both the D and G bands of the C electrode completely vanish, whereas in the S/C-16% electrode, these bands merely diminish in intensity. This difference can be attributed to the introduction of S, an appropriate amount of S enhancing the structural stability of the porous carbon matrix and preserving the Raman activity of the S/C-16% composite throughout the charge/discharge process [52]. As a result, the Raman signals are retained during cycling. This also explains why the introduction of S can substantially enhance the rate capability of porous carbon-based electrode materials. Figure 4. HR-TEM images of (a) S/C-16%, and (b) S/C-18% electrodes showing SEI thickness; In-situ EIS spectra of (c) S/C-16% and (d) S/C-18% electrodes at the first cycle; (e) R SEI and (f) R ct of S/C-16% and S/C-18% electrodes at different voltages; (g) TOF-SIMS 3D views of the sputtered surface and sputtered volume for S/C-16% electrodes; (h) In-situ Raman of S/C-16% electrode. To better understand the effect of S doping on electrode interfacial dynamics, we compared and analyzed the EIS of C and S/C composite electrodes. From the Nyquist plot (Figure 5a), the S/C-16% electrode shows a significantly smaller semi-circle in the high-frequency region, indicating lower R SEI and R ct [44]. Further analysis shows that as S doping increases, the R SEI and R ct values of the S/C electrodes first decrease and then increase, with the S/C-16% electrode exhibiting the lowest interfacial impedance. This suggests that an optimal S doping level enhances the electrode/electrolyte interface structure, promoting rapid interfacial charge transfer. Notably, in the low-frequency region, the S/C-16% electrode displays a steeper slope, indicating smaller Warburg impedance and improved Na + diffusion kinetics within the electrode material [45]. As shown in Figure 5b, Galvanostatic Intermittent Titration Technique (GITT) was employed to calculate the sodium ion diffusion coefficients (D Na+ ) of different samples. Based on Fick’s second law, the D Na+ of the S/C-16% electrode during charging and discharging ranged from 5.4×10 -9 to 5.2×10 -9 cm 2 s -1, which is significantly higher than those of the C electrode (2.7×10 -9 to 2.5×10-9 cm 2 s -1 ) (Figure 5c) [54]. This significant improvement in D Na+ is attributed to S doping, which optimizes the transport channels of the electrode material and effectively reduces the diffusion resistance and path length of Na + . The CV curves of the C and S/C-16% electrodes were measured at scan rates of 0.1, 0.2, 0.4, 0.6, 0.8, and 1.0 mV s -1, as shown in Figure S13a and Figure 5d. For electrode materials, the current and scan rate relationship is expected to follow the theoretical equation: \begin{equation} i=av^{b}\nonumber \\ \end{equation}\begin{equation} \log\left(i\right)=blog\left(v\right)+log(a)\nonumber \\ \end{equation} In this context, i represents the peak current, ν is the scan rate, and a and b are constants obtained from the intercept and slope of the logarithmic plot of scan rate versus peak current [55]. As shown in Figure 5e, the b values for peak 1 and peak 2 of the S/C-16% electrode are 0.95 and 0.64, respectively. In contrast, the b values for peak 1 and peak 2 of the C electrode are 0.76 and 0.84, respectively (Figure S13b). This suggests that the electrochemical process of the S/C-16% electrode is influenced by both capacitance and diffusion. Additionally, the contributions of the capacitive-controlled process ( k 1 v ) and the diffusion-controlled process ( k 2 v 1/2 ) to the total current can be quantified using the following equation [56]: \begin{equation} i\left(v\right)=k_{1}v+k_{2}v^{1/2}\nonumber \\ \end{equation} As shown in Figure 5f, the capacitive contributions of the S/C-16% electrode increased from 46.1% to 73% as the scan rates increased, which are higher than those of the C electrode (Figure S13c). At a scan rate of 0.6 mV s -1, the capacitive contribution of the S/C-16% electrode reaches 64% (Figure S14), higher than the C electrode’s 58.3% (Figure S13d), which is consistent with the results in Figure 5f. To further assess the potential of S/C-16% as a promising anode material for practical applications, we paired it with a Na 3 V 2 (PO 4 ) 3 (NVP) cathode to assemble a sodium-ion full battery. The full battery operates within a voltage window of 0.5 to 3.8 V. After 200 cycles, the battery retained a reversible specific capacity of 197.4 mAh g -1 after 65 charge-discharge cycles at 1.0 A g -1, demonstrating good cycle stability (Figure S15). As presented in Figure 5g, after activation, the first three cycles of S/C-16%//NVP full cell at 0.1 mV s -1 almost completely overlapped, indicating its excellent reversibility. The rate performance of the full battery is shown in Figure 5h. At a low current of 0.05 A g -1, the initial discharge capacity was about 470.8 mAh g -1 . As the current density increased, the capacity gradually decreased, but the decline was mild, indicating good electrochemical kinetics. Importantly, even at 2.0 A g -1, the battery maintained a capacity of about 107.8 mAh g -1, demonstrating that the S/C-16% anode material has excellent rate capability, crucial for high-power density applications. Figure 5i shows the charge-discharge curves of the full battery at various current densities. All the charge-discharge curves exhibit distinct voltage plateaus. As current density increases, the electrochemical plateaus gradually shrink, accompanied by a decrease in capacity and an increase in polarization. This phenomenon is primarily due to electrochemical kinetic limitations at high current densities, including increased charge transfer resistance and limited Na + diffusion within the electrode material [57]. Figure 5. (a) Nyquist plots of C and all S/C electrodes; (b) GITT curves of C and S/C-16% electrodes; (c) Na + diffusion coefficients calculated from GITT data of C electrode and S/16%C electrode; (d) CV curves of S/C-16% electrode at different scan rates; (e) log(i) vs. log(ν) plot derived from (d); (f) Capacity contribution ratios of S/C-16% electrode at different scan rates; (g) CV curves, (h) Rate performance and (i) Charge and discharge curves at different current densities of sodium-ion full battery Figure 6a and Figure S16a present the rate capabilities of C and S/C composite anodes for PIBs. It is evident that all S/C electrodes exhibit markedly enhanced reversible specific capacities compared to the C electrode. Moreover, the reversible capacity of the S/C composites displays a non-monotonic dependence on S content: it initially increases and then decreases as the S content rises. Specifically, the S/C-16% electrode delivers reversible capacities of 523.4, 492.0, 449.4, 415.6, 373.5, 306.4, and 268.3 mAh g -1 at current densities of 0.1, 0.2, 0.4, 0.6, 0.8, 1.0, and 2.0 A g -1, respectively. In contrast, the C electrode only exhibits capacities of 159.9, 112.5, 80.5, 49.5, 35.0, 22.9, and 16.4 mAh g -1 under the same conditions. Importantly, when returning to 0.1 A g -1, the S/C-16% electrode recovers a high reversible capacity of 526.0 mAh g -1, which closely approximating its initial value, whereas the C electrode only regains 145.8 mAh g -1 . These observations provide compelling evidence that appropriate S incorporation introduces additional defect sites within the porous C framework, and moderate S doping is inferred to improve the electronic conductivity and interfacial ion kinetics of the C host, synergistically contributing to the observed improvements in rate performance [28]. Figure 6b further illustrates that S content exerts a significant influence not only on rate capability but also on the ICE of the electrodes. The ICE of the S/C-16% electrode reaches 61.26%, substantially higher than that of the pristine C electrode (28.74%) and other S/C variants. As depicted in Figure 6c and Figure S16b, the cycling stability of various electrodes was evaluated at 0.1 A g -1 over 100 cycles. The S/C-16% electrode delivered a reversible capacity of 439.2 mAh g -1, corresponding to a capacity retention of 74.89%, which is significantly higher than those of C electrode and other S/C composites. Additionally, a consistent trend of initially increasing and then decreasing capacity with higher S content is observed across the S/C series, reaffirming that an optimal S content is critical for achieving balanced electrochemical performance. To further elucidate the electrochemical reversibility of the S/C-16% electrode, Figure 6d presents its charge-discharge profiles at various cycles under 0.1 A g -1 . With the exception of the initial formation cycle, the subsequent charge-discharge curves exhibit excellent overlap, indicative of high cycle-to-cycle stability and consistent electrochemical behavior, in agreement with the long-term cycling results shown in Figure 6c. Furthermore, the Coulombic efficiency of the S/C-16% electrode stabilizes near 100% after the sixth cycle (Figure S16c). The long-term cycling performance of the S/C-16% electrode under high-rate conditions was assessed at a current density of 1.0 A g -1, as shown in Figure 6e. Remarkably, the electrode retains a reversible capacity of 189.9 mAh g -1 after 1000 cycles, significantly exceeding that of the pristine C electrode under identical conditions. Figure 6. Electrochemical Performance of Potassium-Ion Battery Anodes: (a) Rate performance of C and S/C-16% electrodes; (b) Initial Coulombic efficiency of S/C electrodes with different S contents; (c) Cycling performance of C and S/C-16% electrodes at 0.1 A g -1 ; (d) Charge/discharge curves of S/C-16% electrode at 0.1 A g -1 ; (e) Long-term cycling stability of C and S/C-16% electrodes at 1.0 A g -1 3. Conclusion In summary, this work presents a facile strategy for incorporating S into lignite-derived porous C via a combination of low-temperature carbonization and vapor-phase infiltration. The aim is to enhance the ICE, cycling stability, and rate performance of the material as an anode for sodium/potassium-ion batteries. Specifically, the introduction of S effectively increases the defect density within the porous C framework, thereby enhancing the capacity of the S/C composite anode and improving its rate capability. Moderate S incorporation also facilitates the formation of a thinner and more stable SEI, while concurrently improving the electronic conductivity and ionic reactivity of the porous carbon matrix. These synergistic effects result in a favorable balance between high ICE, enhanced rate performance, and excellent cycling stability. Electrochemical evaluations reveal that, as an anode material for SIBs, the S/C composite delivered an outstanding reversible capacity of 480 mAh g -1 at 0.1 A g -1, high ICE of 83.1% at 0.1 A g -1, and maintained 232.6 mAh g -1 at 5.0 A g -1 . When applied as a potassium-ion battery anode, the S/C-16% electrode exhibits a reversible capacity of 439.2 mAh g -1 and an ICE of 61.2% at 0.1 A g -1 . This study provides a novel and effective design approach for developing advanced porous C anodes for sodium/potassium -ion batteries. 4. Experimental Section

are available in the Supporting Information.

We would like to express our sincere gratitude to the Ordos City Industrial Innovation and Entrepreneurship Talent Team Project (Innovation Team for Green and High-Value Utilization of Coal Chemical Waste) for their generous financial support. This work was also supported by the National Natural Science Foundation of China (Nos. 51962027, and 21968022), Program for Young Talents of Science and Technology in Universities of Inner Mongolia Autonomous Region (No. NJYT24002), and Central Guidance Fund for Local Scientific and Technological Development (No. 2024ZY0012). Conflicts of Interest The authors declare no conflicts of interest. Data Availability Statement The data that support the findings of this study are available upon request from the author. Supporting Information Additional supporting information can be found online in the Supporting Information section. Received: ((will be filled in by the editorial staff)) Revised: ((will be filled in by the editorial staff)) Published online: ((will be filled in by the editorial staff))

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Li* Tuning S-doping in Porous Carbon for Enhanced Initial Coulombic Efficiency and Rate Performance for Sodium/ Potassium-Ion Batteries Information & Authors Information Version history Peer review timeline Published ENERGY & ENVIRONMENTAL MATERIALS Version of Record16 Feb 2026Published Copyright This work is licensed under a Non Exclusive No Reuse License. Collection

Authors Metrics & Citations Metrics Article Usage 228views 153downloads Citations Download citation Jiahao Zhao, Lankun Shi, Mengting Jia, et al. Tuning S-doping in Porous Carbon for Enhanced Initial Coulombic Efficiency and Rate Performance for Sodium/Potassium-Ion Batteries. Authorea. 11 August 2025. DOI: https://doi.org/10.22541/au.175494810.02181197/v1 DOI: https://doi.org/10.22541/au.175494810.02181197/v1 If you have the appropriate software installed, you can download article citation data to the citation manager of your choice. Simply select your manager software from the list below and click Download. For more information or tips please see 'Downloading to a citation manager' in the Help menu.