Structural Engineering of Biomass‐Derived Hard Carbon With Architectured Closed Pores for Fast Sodium Storage

ABSTRACT

Hard carbons are widely used as sodium‐ion battery (SIB) anode materials owing to their abundant resources, cost‐effectiveness, and excellent electrochemical performance, but their sluggish charge transfer kinetics limit fast‐charging applications. Herein, we synthesize porous spherical nanostructured hard carbon via acid‐assisted hydrothermal precarbonization of biomass sphagnum moss, followed by high‐temperature treatment. This microstructure engineering introduces interconnected closed pores, expands the interlayer distance to 0.43 nm, and forms a uniform spherical morphology, effectively reducing ion diffusion resistance and boosting reaction kinetics. The optimized anode delivers a reversible capacity of 108 mAh g⁻¹ at 10 A g⁻¹. Assembled into a full cell with a NaNi₁/₃Fe₁/₃Mn₁/₃O₂ cathode, it retains 70% capacity after 1000 cycles at 300 mA g⁻¹, demonstrating superior cycling stability and fast‐charging capability. This work provides a feasible microstructure engineering strategy for high‐rate sodium storage in biomass‐derived carbons.

To overcome sluggish kinetics in hard carbon anodes, we engineer spherical hard carbon from biomass. Its interconnected pores and expanded interlayers enable ultrafast Na⁺ transport, achieving 108 mAh g⁻¹ at 10 A g⁻¹ and stable full‐cell cycling.

1. Introduction

To address global warming and the energy crisis, the research and development of large‐scale energy storage technologies has become a critical path for achieving renewable energy storage [1, 2]. Sodium‐ion batteries (SIBs) are regarded as strong contenders for grid energy storage due to their abundant sodium resources, low economic cost, and superior electrochemical performance [3]. Nevertheless, the large ionic radius of sodium ions (1.02 Å) and the instability of Na‐C intercalation compounds impose stringent demands on the choice of electrode materials [4, 5]. Recent advancements have been made in the development of anode materials for SIBs [6, 7, 8, 9, 10, 11, 12]. Among these, hard carbon has emerged as an ideal candidate for SIB anodes, attributed to its extensive interlayer spacing (exceeding 0.37 nm), rich pore structure, and intrinsic defects [13]. In particular, biomass‐derived hard carbon has emerged as a highly desirable commercial anode material, owing to its abundant raw materials, straightforward synthesis process, low cost, and environmental friendliness [14, 15, 16]. Despite these advantages, significant challenges such as poor cycling stability, low initial Coulombic efficiency, and inadequate rate performance still hinder its development [17, 18].

Currently, a variety of strategies have been employed to address the aforementioned obstacles. For example, Wu et al. successfully optimized the microstructure by adjusting the pyrolysis temperature of the biomass precursor, thereby achieving higher initial Coulombic efficiency (ICE) and reversible capacity [19]. Similarly, Chou et al. developed a manganese ion‐catalyzed carbonization technique, which precisely controls the degree of graphitization, eliminates defects, and maintains an efficient sodium pathway, ultimately yielding hard carbon with an ICE of approximately 92.05% [20]. Additionally, Hu et al. utilized pre‐oxidation technology to modify defects and heteroatom groups within the carbon material [21]. Furthermore, a hybrid composite design method was employed to fine‐tune the interlayer spacing and defect sites of carbon microcrystals, resulting in significant enhancements in both ICE and specific capacity [22, 23, 24]. Nevertheless, despite the performance improvements enabled by these strategies, the rate capability remains unsatisfactory, typically falling below 100 mAh g⁻¹ at current densities exceeding 5 A g⁻¹ [25, 26, 27]. This limitation, largely attributed to sluggish charge‐transfer kinetics, hinders the practical deployment of hard carbon as a viable anode material for SIBs. Consequently, there is an urgent need to develop a hard carbon structure capable of facilitating rapid sodium ion transfer across the entire voltage range [28].

In this study, we successfully synthesized a distinctive spherical hard carbon material derived from sphagnum moss, characterized by diameters ranging from 200 to 800 nm, through an acid‐mediated hydrothermal pre‐carbonization process with hydrochloric acid serving as the catalyst. The formation of spherical structures is a key factor in enhancing reaction kinetics. Beyond shortening the ion diffusion path, the ability to tune porosity and the high surface area of these structures further facilitate rapid charge transfer [29]. Moreover, the presence of moderate amounts of oxygen atoms on the surface of the spheres potentially enhances sodium storage capabilities [30]. Collectively, these factors contribute to the efficient transport of sodium ions. Specifically, the spherical‐structured HCHH‐1400 maintained an impressive specific capacity of 290.86 mAh g⁻¹ after 700 cycles at a current density of 600 mA g⁻¹, and demonstrated a remarkable reversible specific capacity of 108 mAh g⁻¹ at a current density of 10 A g⁻¹, showcasing rapid sodium ion storage. Furthermore, when integrated into a sodium‐ion full battery with a NaNi_1/3Fe_1/3Mn_1/3O₂ (NFM) cathode, it exhibited exceptional rate performance and long‐term cycling stability. This study provides a novel perspective on the rapid transport of sodium ions in biomaterials.

2. Result and Discussion

2.1. Preparation and Structural Characterization of Hard Carbons

The typical preparation process can be illustrated in Figure 1. The sphagnum moss‐derived hard carbons were synthesized by acid hydrothermal pre‐carbonization at 150°C for 5 h, followed by carbonization at 1400°C for 2 h under an Ar atmosphere, which is denoted as HCHH‐1400. In comparison, the same precursor was directly carbonized at 1400°C for 2 h and is denoted as HC‐1400. The carbonization temperature of 1400°C is chosen based on the conclusion of our initial experimental exploration, that is, hard carbon pyrolyzed at 1400°C could achieve the best balance between defects and pore structure (Figures S1–S7). Experimental details are given in the Supporting Information.

FIGURE 1.

The schematic illustration of the synthesis for hard carbon materials (blue dots: closed pore storage sites; green dots: graphitic intercalation sites; red dots: surface absorption sites).

The changes in the microscopic morphology of the prepared carbon materials have been observed through FESEM. As shown in Figure S8, the hard carbon obtained by acid hydrothermal pre‐carbonization followed by high‐temperature pyrolysis exhibits a densely packed spherical morphology ranging in size from 200 to 800 nm, whereas the directly carbonized counterpart HC‐1400 inherits the lamellar structure of sphagnum moss. This morphological change is similar to the formation of homogenous carbon microspheres via hydrothermal treatment of glucose at mild or neutral pH conditions [31]. The distinction is that the cellulose, hemicellulose, and lignin comprised of sphagnum moss cannot be completely hydrolyzed in a neutral solution [32]. Under acid catalysis, the polysaccharide components of sphagnum moss are hydrolyzed and dehydrated into 5‐hydroxymethylfurfural (HMF) and related furanic intermediates, which subsequently polymerize and self‐assemble into spherical hydrothermal‐carbon particles [33, 34]. Subsequent high‐temperature carbonization further consolidates the spherical morphology and establishes a conductive carbon network, and the compact packing density and structural stability of the resulting spherical particles are beneficial for shortening ion and electron diffusion paths and for mitigating electrode deformation during cycling. Scanning transmission electron microscopy (STEM) images further confirm the spherical and lamellar structures of HCHH‐1400 and HC‐1400, respectively (Figure 2a,d), in agreement with the SEM observations. Figure 2b and Figure S9 reveal that HCHH‐1400 is composed of many highly disordered, short‐range distorted graphite‐like layers [35], whereas HC‐1400 predominantly features extensive long‐range structured graphite‐like domains (Figure 2e; Figure S9). More importantly, the HCHH‐1400 exhibits richer interconnected closed pores surrounded by abundant short graphite‐like layers than the HC‐1400 (marked by yellow dashed lines in Figure 2b,e), which is beneficial for the fast transport of sodium ions. Furthermore, elemental mapping analysis further shows that the two carbon samples have consistent and homogenous constituent compositions (Figure 2c,f). The porous structure was further examined by small‐angle X‐ray scattering (SAXS) measurement (Figure S10a). From which one can see, HCHH exhibits stronger scattering intensity compared with HC, indicating a significantly increased number of closed pores and hierarchical pores in HCHH [36, 37]. The characteristic shoulder signal (in the range of 0.02–0.1 Å⁻¹) and the slower intensity decay in the high q region (>0.1 Å⁻¹) mechanistically correspond to the formation of well‐developed microporous structures. According to the fitting results, the average pore diameter of HCHH is calculated to be 0.94 nm, which is notably larger than that of HC (0.89 nm). Unlike isolated closed pores that are kinetically inaccessible, the closed pores in HCHH‐1400 are delimited by short and curved graphite‐like layers with an expanded interlayer spacing, which allows Na⁺ to reach the pore interior via interlayer intercalation and subsequent pore filling rather than through direct electrolyte penetration. The interconnection between adjacent pores through these distorted short graphitic layers further provides continuous pathways for Na⁺ percolation between neighboring pore cavities, thereby enabling rapid plateau‐region storage without compromising the closed‐pore character that suppresses excessive SEI formation.

FIGURE 2.

Material characterization. TEM images and elemental mapping images for HCHH‐1400 (a–c) and HC‐1400 (d–f); XRD (g) and Raman spectra (h) of HC‐1400 and HCHH‐1400; Fitted C 1s and O 1s core‐level XPS spectra of (i,j) HCHH‐1400 and (k,l) HC‐1400.

The change in defects and carbon structure of the hard carbon could be further explained with X‐ray diffraction (XRD) and a Raman spectroscope. XRD results in Figure 2g show that all samples present two amorphous peaks located at around 22° and 44°, which correspond to the typical (002) and (100) crystal planes in disordered carbons, respectively [38]. Although the materials are carbonized at the same temperature, the acid hydrothermal pre‐carbonization of the precursor causes a slight red shift in the position of the (002) peaks. Based on the Bragg equation (Table S1), the d₀₀₂ of HCHH‐1400 is 0.43 nm, which is wider than the 0.41 nm of HC‐1400, allowing for more Na ion storage due to the larger interlayer spacing. In addition, the Raman spectra of HCHH‐1400 and HC‐1400 show the typical D band dominated by disordered carbon at ∼1350 cm⁻¹ and the G band dominated by crystalline graphite at ∼1580 cm⁻¹ (Figure 2h) [39]. The HCHH‐1400 sample with acid hydrothermal pre‐carbonization has a larger peak area ratio (A_d/A_g) of 1.42 compared to 1.38 for HC‐1400. This result indicates that acid hydrothermal pre‐carbonization may induce a more disordered microstructure, leading to more reversible adsorption active sites. Furthermore, HCHH‐1400 also displays a smaller lateral width (L_a) of the carbon layer than HC‐1400, which means that the shorter or more curved carbon layer formed during the acidic hydrothermal pre‐carbonization process is conducive to the formation of a rich interconnected closed‐pore structure. (Figure S10b and Table S1).

N₂ adsorption‐desorption isotherms are performed to evaluate the porosity of carbons. According to the isotherms in Figure S11, all samples display typical IV‐type adsorption curves with an apparent hysteresis loop at high relative pressures, demonstrating the presence of mesopores and micropores [40]. Specifically, the HCHH‐1400 sample reveals a larger specific surface area of 25 m² g⁻¹ and pore volume of 0.009 cm³ g⁻¹ than HC‐1400 (10 m² g⁻¹ and 0.003 cm³ g⁻¹). Meanwhile, the pore size distribution curves (Figure S11) reveal that HCHH‐1400 samples are mainly concentrated at around 4 nm and exhibit a more uniform pore structure than HC‐1400. The reason can be attributed to the spherical microstructure formed by acid hydrothermal pretreatment, which reduces packing density between particles while increasing voids between particles, hence increasing the overall specific surface area and porosity. The relatively high specific surface area and micropore volume can increase electrode‐electrolyte interaction while also promoting sodium ion transport and adsorption [41].

To investigate the effect of the acid hydrothermal pre‐carbonization on the surface chemical environment of the hard carbon, X‐ray photoelectron spectroscopy (XPS) measurements were conducted to analyze the chemical composition and elemental states of the as‐prepared samples. As illustrated in Figure S12, the comprehensive survey spectra reveal two distinct peaks corresponding to C 1s and O 1s. Notably, the oxygen content increases slightly from 9.73 wt.% for HC‐1400 to 12.95 wt.% for HCHH‐1400, suggesting that acid hydrothermal pre‐carbonization can introduce a greater number of oxygen‐containing functional groups, thereby creating more active sites for Na⁺ [42]. Specifically, the C 1s spectra can be fitted into four peaks located at about 284.2, 284.8, 286.2, and 289.3 eV, which are assigned to C═C, C─C, C─O, and C═O, respectively (Figure 2i,k) [43]. Among them is the relative content of the C═O increase of HCHH‐1400, which is advantageous for the reversible adsorption of Na during the electrochemical process [44]. As for the O 1s fitting spectrum, it exhibits three peaks at 532.1, 533.5, and 535.2 eV, corresponding to C═O, C─O, and ─OH, respectively (Figure 2j,l) [45]. In comparison, HCHH‐1400 has higher C═O and ─OH ratios compared to HC‐1400, providing additional sloping capacity (Table S2) [46]. Overall, the characterization of the physicochemical properties of the prepared materials, including the degree of graphitization, interlayer spacing, specific surface area, porous structure, and compositional changes, reveals that acid hydrothermal pre‐carbonization can increase defect concentration, create a rich interconnected closed pore structure, and induce the formation of more oxygen‐containing functional groups, thereby providing more active sites for sodium ion storage. The electrochemical performance outcomes are likely related to the morphology and structure of the obtained materials.

2.2. Sodium Storage Performance of Half‐Cell Configuration

To understand the effect of the microstructure changes of hard carbon caused by acid hydrothermal pre‐carbonization on sodium ion storage performance, the Na storage performances were tested in a half‐cell with Na metal as the counter/reference electrode. Figure 3a illustrates the first galvanostatic charge/discharge (GCD) curves at 30 mA g⁻¹ for HCHH‐1400 and HC‐1400. The profiles are divided into a high‐potential sloping region (voltage above 0.1 V) and a low‐potential plateau region (voltage below 0.1 V). Notably, HCHH‐1400 exhibits a superior reversible charge capacity of 405.7 mAh g⁻¹ with an initial coulombic efficiency (ICE) of 75.05%, surpassing HC‐1400, which delivers 359.2 mAh g⁻¹ and an ICE of 68.78% [38, 47]. The improved ICE appears to contrast with the higher BET surface area of HCHH‐1400, which might be ascribed to the fact that the BET surface area of HCHH‐1400 remains relatively low, at 25 m² g⁻¹, with a small pore volume of 0.009 cm³ g⁻¹, indicating that it is not a highly open porous carbon. The increased surface area mainly arises from interparticle voids and moderate mesoporosity formed by the packed spherical particles, which facilitate electrolyte infiltration and Na⁺ transport but do not cause excessive SEI formation. More importantly, the internal closed pores observed by HRTEM are largely inaccessible to the electrolyte and therefore do not introduce additional SEI‐forming interfaces. These closed pores instead serve as reversible Na‐storage sites after Na⁺ intercalation into the expanded carbon layers. In addition, the spherical morphology of HCHH‐1400 promotes a homogeneous ion flux and stable SEI, reducing irreversible Na⁺ consumption. As a result, HCHH‐1400 shows a lower irreversible capacity than HC‐1400 and thus achieves a higher ICE despite its increased BET surface area. In addition, HCHH‐1400 demonstrates a greater slope and plateau capacity contribution compared to HC‐1400 at 2000 mA g⁻¹ (Figure 3b). These results demonstrate that the defects and interconnected closed pore structures from acid hydrothermal pre‐carbonization promote rapid sodium ion adsorption and filling, providing more reversible active sites for the transport of sodium ions, and exhibiting excellent specific capacity, which is consistent with the physical characterization results.

FIGURE 3.

Electrochemical performance of the half‐batteries. (a) The first GCD curves of HC‐1400 and HCHH‐1400 at 30 mA g⁻¹, (b) Capacity contribution comparison for the slope and plateau region in the discharge curve at 2000 mA g⁻¹, (c) Rate performance for HC‐1400 and HCHH‐1400, (d) Long‐cycle performance at 600 mA g⁻¹, (e) Specific capacities of HCHH‐1400 and HC‐1400 in comparison to reported carbon anodes for SIBs.

The rate capability of the hard carbon was further investigated by applying stepwise current densities ranging from 0.03 to 10 A g⁻¹ (Figure 3c). Obviously, the HCHH‐1400 exhibits excellent rate capability with specific capacities of 370, 364, 356, 350, 338, 329, 298, 263, 226, 192, and 108 mAh g⁻¹ at 0.03, 0.06, 0.15, 0.3, 0.6, 0.9, 2, 3, 4, 5, and 10 A g⁻¹, respectively, much higher than the HC‐1400 anode. The discharge capacity of HCHH‐1400 can return to its pristine value when the current density recovers to 30 mA g⁻¹, confirming its outstanding rate performance. The superior rate capacity of the HCHH‐1400 anode could be influenced by the following reasons: (1) The spherical particles formed through acid hydrothermal pre‐carbonization can uniformly distribute stress, with their dense packing enhancing the mechanical strength of the electrode material and mitigating volume changes and stress during cycling. (2) The large interlayer spacing, abundant defects, and interconnected closed pore structure of the HCHH‐1400 material provide numerous active sites for Na⁺ transport. (3) The spherical particles promote the establishment of a continuous conductive network, providing an efficient pathway for electron conduction. Furthermore, the HCHH‐1400 anode manifests great cycling stability at different current densities, further verifying its exceptional morphology and structural stability. As illustrated in Figure 3d, the HCHH‐1400 anode maintains a high reversible discharge capacity of 290.86 mAh g⁻¹ after 700 cycles at 600 mA g⁻¹, with an average CE of 99.93%, demonstrating a magnificent capacity retention of 80%. Remarkably, HCHH‐1400 also achieves an excellent reversible specific capacity of 212.7 mAh g⁻¹ after 500 cycles at a high current density of 5 A g⁻¹ (Figure S13), indicating a capacity retention of 85% and showcasing its robustness. Additionally, the high rate capability and cycling stability of HCHH‐1400 surpass the performance of most anode materials reported in the literature [44, 48, 49, 50, 51], suggesting that the microstructural and morphological strategies significantly enhance the overall electrochemical performance of the HCHH‐1400 anode (Figure 3e; Table S3). Moreover, under a high mass loading of 5.5 mg cm⁻², HCHH‐1400 still delivers a reversible capacity of 262 mAh g⁻¹ at 0.6 A g⁻¹, 1.8 times higher than HC‐1400, along with a clearly superior rate capability (Figure S23). This confirms that the kinetic merits of the spherical architecture are well preserved in practically relevant thick electrodes.

2.3. Electrochemical Kinetic Analysis

To further elucidate the influence of microstructure on electrochemical behavior, cyclic voltammetry (CV) measurements of hard carbons in the half‐cell configuration can be utilized to probe the detailed electrochemical reactions and their corresponding kinetics. The CV curves of carbon materials, recorded at 0.1 mV s⁻¹ over three cycles, are displayed in Figure S14. During the first scan, both samples show an irreversible reduction peak at approximately 0.5 V, along with a pair of redox peaks between 0.01 and 0.1 V, which correspond to the formation of the solid electrolyte interphase (SEI) during the initial discharge and the storage of Na ions in the plateau region [52]. In subsequent cycles, the irreversible reduction peak is absent, and the CV curves are nearly overlapped, indicating the high reversibility of the electrochemical reactions. Furthermore, the detailed electrochemical reactions and their corresponding kinetics can be investigated at various scan rates from 0.1 to 10 mV s⁻¹. Figure 4a,b shows that the peak current of HCHH‐1400 increases more significantly with the scan rate compared to the HC‐1400 electrode, indicating superior reaction kinetics in HCHH‐1400. Additionally, the diffusion‐control and capacitive‐dominated processes can be described by the power‐law formula:

$$i=a{v}^b$$ (1)

where i represents the peak current, v is the scan rate, and a and b are adjustable parameters. This relationship helps distinguish between diffusion‐controlled processes (b ≈ 0.5) and capacitive‐dominated processes (b ≈ 1) [53]. Through linear fitting of log v against log i, the b values for the anodic and cathodic peaks in HCHH‐1400 are found to be 0.43 and 0.37, respectively, which are higher than those of HC‐1400 (0.37 and 0.31), indicating faster reaction kinetics in HCHH‐1400 than in HC‐1400 (Figure 4c). Notably, the b value close to 0.5 implies that sodium storage is primarily dominated by diffusion processes.

FIGURE 4.

Kinetic analysis. CV profiles with different scan rates for (a) the HCHH‐1400 anode and (b) the HC‐1400 anode, (c) Linear relationship between log (i) and log (v) of HCHH‐1400 and HC‐1400, (d) The capacity‐controlled current separated from the diffusion‐controlled one at each potential in CV curves of HCHH‐1400 at 6 mV s⁻¹, (e) Calculated contributions of the capacity‐ and diffusion‐controlled proportions of HCHH‐1400 and HC‐1400, (f) AFM images of HCHH‐1400 and HC‐1400 electrodes after cycle 10.

Furthermore, to elucidate the differences observed in the CV curves, the currents contributed by capacitive‐controlled and diffusion‐controlled reactions are calculated using the following formula [54, 55]:

$$i=k_{1}v+k_2{v}^{1/2}$$ (2)

where k₁v represents the pseudocapacitive contribution and k₂v^1/2 corresponds to the diffusion‐controlled capacity. As depicted in Figure 4d and Figure S15, the response currents in the slope region are primarily capacitive‐controlled currents, while diffusion‐controlled currents dominate in the plateau region. From a kinetic perspective, the capacitive‐controlled process in the slope region is attributed to the adsorption of sodium ions on active sites such as defects, resembling a pseudocapacitive mechanism and exhibiting faster kinetics [56]. Conversely, the diffusion‐controlled reaction in the plateau region involves the slower kinetics of cation diffusion and storage within the carbon bulk phase [31]. Figure 4e illustrates the contributions of capacitive‐controlled and diffusion‐controlled processes in hard carbon materials across varying scan rates. At low scan rates, the capacity is predominantly due to diffusion‐controlled reactions in the plateau region. As the scan rate increases, the proportion of diffusion‐controlled reactions in the plateau region decreases, indicating insufficient time for these reactions to occur. Notably, HCHH‐1400 still exhibits a significant diffusion‐controlled contribution of approximately 39% even at a scan rate of 10 mV s⁻¹. These behaviors could be attributed to several factors: (1) The spherical architecture of HCHH‐1400 alters the dimensionality of Na⁺ transport relative to the layered framework of HC‐1400. In HC‐1400, Na⁺ migration is geometrically constrained by the continuous, anisotropically stacked graphitic domains inherited from sphagnum moss, forcing ions to traverse micrometer‐scale paths before reaching interior active sites. In contrast, the 200–800 nm spherical particles of HCHH‐1400 provide isotropic, 3D diffusion pathways with a characteristic length shortened by one to two orders of magnitude, which is the primary origin of its superior rate capability. Moreover, the dense packing of these spherical units establishes an interconnected 3D percolation network through abundant particle–particle contacts, simultaneously serving as continuous channels for electron conduction and electrolyte infiltration; (2) the interconnected closed pore structure formed by structural regulation provides more space for the storage of sodium ions and the higher specific surface area enhances contact between sodium ions and hard carbon, facilitating quicker penetration of sodium ions into the material; and (3) acid hydrothermal pre‐carbonization generates more active sites, enhancing both surface interactions and diffusion control capabilities, thereby promoting efficient ion and electron diffusion and exhibiting excellent electrochemical kinetics. These factors collectively elucidate why HCHH‐1400 demonstrates exceptional rate performance.

Considering that the primary distinction between the two electrodes lies in the morphology of hard carbon, as demonstrated above, it is clear that morphology directly impacts electrochemical performance. Therefore, in situ electrochemical impedance spectroscopy (EIS) and atomic force microscopy (AFM) were employed to thoroughly investigate the influence of morphology on the electrode interface. The EIS Nyquist plots in Figure S16 also reveal that spherical particles of HCHH‐1400 evenly disperse stress and offer a shorter ion transfer path [57]. In addition, the coarser surface of HC‐1400 (R_a: 149 and R_q: 181) after cycling compared to that of HCHH‐1400 (R_a: 55 and R_q: 73) is proved by AFM (Figure 4f), suggesting that the spherical morphology of HC‐1400 can homogenize the ion flux to stabilize the interface condition [58]. To further verify the role of the SEI in the enhanced kinetics of HCHH‐1400, ex situ F 1s XPS spectra of both electrodes were collected after 100 cycles (Figure S22). The deconvoluted spectra reveal that the SEI on HCHH‐1400 is dominated by inorganic NaF (58.0%) with only a minor fraction of organic P–F species (18.0%), whereas HC‐1400 shows a much lower NaF content (46.0%) and a significantly higher P–F content (35.6%). The NaF‐rich, organic‐poor SEI on HCHH‐1400 is thinner, more ionically conductive, and more mechanically robust, which simultaneously (i) accelerates Na⁺ transfer across the electrode/electrolyte interface, (ii) suppresses continuous decomposition of NaPF₆ and the solvent, and (iii) stabilizes the interface during long‐term cycling. These SEI‐level features, together with the spherical morphology, expanded interlayer spacing, and interconnected closed pores discussed above, jointly account for the fast and durable sodium storage of HCHH‐1400.

2.4. Elucidating the Sodium Storage Mechanism

To gain systematic insights into the storage mechanisms of HCHH‐1400, a series of characterization techniques, including galvanostatic intermittent titration technique (GITT), ex situ XPS, in situ Raman, and ex situ XRD, were employed to study sodium storage behavior in detail. GITT measurements were conducted to determine the diffusivity coefficient (D_Na+) of sodium ions at various voltages, elucidating the sodium storage mechanism. Figure 5a and Figure S17 illustrate two similar GITT profiles, indicating comparable sodium storage behaviors for both hard carbons. Notably, D_Na+ exhibits a characteristic rise and fall process, indicative of sodium ion adsorption on the electrodes and insertion into the pseudo‐graphite layer. Furthermore, D_Na+ increases between 0.03 and 0.01 V, aligning with a typical pore filling process, consistent with the proposed “adsorption‐insertion‐filling” mechanism [20, 59].

FIGURE 5.

Analysis of Na storage mechanism of HCHH‐1400. (a) GITT curves of the HC‐1400 and HCHH‐1400, (b) Ex situ XPS Na 1s spectra, (c) In situ Raman spectroscopy of HCHH‐1400 during the whole discharge and charge process, (d) Schematic diagram of the sodium storage mechanism of hard carbon.

Ex situ XPS Na 1s spectra were used to analyze the chemical bonding states of Na in HCHH‐1400 during different stages of sodiation and desodiation in the initial cycle. As shown in Figure 5b, a distinct Na 1s peak is observed at the binding energy of 1071.4 eV when HCHH‐1400 is discharged to 0.1 V, and this peak remains almost unchanged until 0.1 V, implying that Na is absorbed onto the electrode surface in ionic states in the high‐voltage sloping region. Notably, the peak position shifts significantly to a higher binding energy of 1072.3 eV with continuous discharge to 0.01 V. Considering that the binding energy of metal bonds is higher than that of adsorbed surface defects [39, 60], sodium is likely in a quasi‐metallic state in the low‐voltage plateau region [61]. During charging to 3 V, the binding energy of Na 1s gradually recovers to 1070.5 eV, demonstrating that sodium transitions from a quasi‐metallic state back to an ionic state in the HCHH‐1400 anode. In addition, ex situ XRD is an effective method for detecting phase evolution and interlayer spacing changes (Figure S18). The (002) peak position of the HCHH‐1400 electrode obviously shifts to lower angles when discharged to 0.1 V, indicating Na⁺ intercalation into the interlayer spaces. Furthermore, when HCHH‐1400 is discharged to 0.05 V, two distinct diffraction peaks appear, corresponding to metallic Na and NaC_x compounds. This implies that the low‐voltage plateau region is responsible for not only sodium intercalation but also sodium clustering in the pores, corroborating the ex situ XPS results.

Additionally, considering the significant impact of Na ion insertion on the intensity and position of the D and G bands, in situ Raman spectroscopy was used to investigate the structural evolution of HCHH‐1400 during discharge and charge. As displayed in Figure 5c, when the open circuit voltage is discharged to 0.45 V, both the D band and the G band reveal a slight red shift, demonstrating adsorption at the boundary of the graphite‐like domain [44]. Subsequently, the D band continued to red‐shift, and its peak intensity dramatically decreased due to Na⁺ adsorption on the surface and defect sites [62]. By analogy with previous DFT studies on alkali‐metal intercalation into carbon materials [14, 40, 63], the pronounced G‐band red shift near 0.15 V can be reasonably attributed to electron transfer from inserted Na into the π* antibonding states, which weakens and elongates the in‐plane C─C bonds. At potentials down to 0.01 V, the G band remains nearly constant, indicating that quasi‐metallic sodium fills the pores [64, 65]. Based on the above analysis, we infer the sodium storage mechanism of hard carbon as depicted in Figure 5d, following an “adsorption‐intercalation‐filling” process. This mechanism involves sodium ion adsorption at defect sites in the slope region, sodium ion insertion into the pseudo‐graphite layers, and quasi‐metallic sodium cluster filling in the pores in the plateau region.

The above sodiation mechanism further highlights that the multiple structural features of HCHH‐1400 contribute to its fast‐charging behavior in a hierarchical rather than additive manner. The spherical morphology primarily governs long‐range ion and electron transport by shortening the diffusion length and establishing a 3D percolation network. Within this framework, defects and oxygen‐containing groups account for the capacitive‐controlled adsorption in the sloping region, while the expanded interlayer spacing and interconnected closed pores enable the diffusion‐controlled intercalation and pore‐filling in the plateau region. Since these features all originate from the same acid‐assisted hydrothermal pre‐carbonization step, they are structurally coupled rather than independent, with the spherical architecture further serving as the geometric basis that renders the interior active sites kinetically accessible.

2.5. Practical Applications of Sodium‐Ion Full Batteries

The spherical morphology of HCHH‐1400 also brings a clear advantage in packing density, which is a key parameter for practical electrode fabrication. Tap‐density measurements show that HCHH‐1400 reaches ∼0.22 g cm⁻³, approximately 2 times higher than that of the lamellar HC‐1400 (∼0.12 g cm⁻³). This is because the microspheres pack much more efficiently than the anisotropic, irregularly stacked lamellae inherited from sphagnum moss. The higher packing density directly translates into a higher volumetric energy density at the electrode level, demonstrating the practical merit of the spherical architecture in addition to its kinetic benefits. ​Motivated by these intrinsic and electrode‐level advantages, we further assessed the practical application potential of HCHH‐1400 by assembling sodium‐ion full batteries. The batteries were paired with a commercial NaNi_1/3Fe_1/3Mn_1/3O₂ cathode, denoted as HCHH‐1400//NFM (Figure 6a). Figure 6b displays the GCD curves of the individual anode and cathode at the same surface current. The initial anode discharge capacity and cathode charge capacity were used to adjust the active mass of the electrodes, ensuring the capacity ratio of the negative to positive electrodes (N/P ratio) remained between 1 and 1.1, thus maximizing the stability of the full cells. Impressively, the GCD curve of full batteries (Figure S19) shows a high average operating voltage of 3.2 V and a large specific capacity of 95.9 mAh g⁻¹ at 10 mA g⁻¹ (based on the cathode mass). As the current density increases, the specific capacity of the full batteries becomes 89, 81.5, 74.9, 67.2, and 60.2 mAh g⁻¹at 20, 50, 100, 200, and 300 mA g⁻¹, respectively (Figure 6c). When the current density is returned to 10 mA g⁻¹, the capacity nearly returns to its original value, displaying superior rate capability and excellent reversibility. Furthermore, high cycling stability with a capacity retention of 70% is achieved after 1000 cycles at 300 mA g⁻¹ (Figure 6d). Additionally, the full cells were also tested at various temperatures ranging from −30°C to 60°C (Figure S20).

FIGURE 6.

Electrochemical performance of the full cells. (a) The schematic illustration of the full cell, where the HCHH‐1400 anode is coupled with the NaNi_1/3Fe_1/3Mn_1/3O₂ cathode during the charge process, (b) The initial charge‐discharge curves of the hard carbon anode and the NaNi_1/3Fe_1/3Mn_1/3O₂ cathode in the half‐cell at the same surface current, (c) Rate performance of HCHH‐1400//NFM at various current densities, (d) Long cycle performance of HCHH‐1400//NFM at 300 mA g⁻¹, (e) Long cycle performance of HCHH‐1400//NFM pouch cell at 50 mA g⁻¹, with an inserted image of a LED light powered by the pouch cell of HCHH‐1400//NFM.

As a proof‐of‐concept for scale‐up potential, single‐layer sodium ion pouch cells were assembled using HCHH‐1400 as the anode and NFM as the cathode, with increased mass loadings (anode: 3.5 mg cm⁻², cathode: 9.8 mg cm⁻²). Notably, the initial discharge/charge curve provides a distinct potential plateau around 3.2 V and a reversible capacity of approximately 102.1 mAh g⁻¹ at 20 mA g⁻¹ (Figure S21). The long cycle performance within the voltage range of 1–4 V demonstrates a capacity retention rate of 69% after 100 cycles at 50 mA g⁻¹ (Figure 6e). Furthermore, the HCHH‐1400//NFM full cell exhibits a high energy density of 203.64 Wh kg⁻¹ at a power density of 42.5 W kg⁻¹, based on the combined mass of the anode and cathode. As a practical demonstration, a single‐layer full‐cell successfully powered an LED panel (inset of Figure 6e), showcasing the promising real‐world applicability of the HCHH‐1400 materials.

3. Conclusions

In summary, this study has successfully developed hard carbon materials exhibiting high reversible capacity and exceptional charge transfer kinetics through the acid hydrothermal pre‐carbonization of biomass sphagnum moss. The optimized HCHH‐1400 demonstrates a remarkable reversible capacity of 405.7 mAh g⁻¹ at 30 mA g⁻¹, and showcases outstanding rate capability with 108 mAh g⁻¹ at 10 A g⁻¹, coupled with exceptional long‐cycle stability. Kinetic analyses have identified that the superior charge transfer capabilities are primarily attributed to the distinctive spherical structure engendered by acid hydrothermal pre‐carbonization, which not only adeptly accommodates rapid charge‐discharge cycles but also minimizes ion transmission resistance, thereby enhancing the sodium ion storage performance. Furthermore, utilizing in situ and ex situ characterization techniques, it was confirmed that the sodium storage mechanism aligns with the “adsorption‐intercalation‐filling” mechanism, underscoring its intricate and efficient storage process. Additionally, the sodium‐ion full battery, paired with a NaNi_1/3Fe_1/3Mn_1/3O₂ cathode, exhibited a commendable capacity retention of 70% after 1000 cycles at a current density of 300 mA g⁻¹. Simultaneously, single‐layer sodium ion pouch cells achieved a remarkable energy density of 203.64 Wh kg⁻¹. These findings underscore the significant potential of this hard carbon material for practical applications and validate its effectiveness as a strategy for advancing rapid sodium storage technologies.

Conflicts of Interest

The authors declare no conflicts of interest.

Supporting information

Supporting File: smll73967‐sup‐0001‐SuppMat.docx.

Data Availability Statement

The data that supports the findings of this study are available in the supplementary material of this article.