Introduction

To meet the new requirements of clean and efficient energy storage systems driven by the accelerated development of portable electronic devices and power electric vehicles, various advanced battery systems have been quickly emerging in recent decades. Sodium-ion batteries (SIBs), as a promising candidate for grid energy storage, have attracted worldwide investigation enthusiasm due to the natural abundance and low cost of sodium [1,2,3]. As an important component of sodium-ion batteries, cathode materials play a decisive role in improving electrochemical performance and reducing the cost of the whole cell. Consequently, suitable SIB cathode materials have been extensively explored, including layered metal oxides [4,5,6,7]; Prussian blue analogues (PBAs) [8, 9]; and especially polyanion-based materials, which are the most popular cathode candidates for SIBs due to their small-volume change and high structural stability, such as phosphates (e.g., Na3V2(PO4)3, Na4MnV(PO4)3, Na3MnTi(PO4)3, and Na3MnZr(PO4)3) [10,11,12,13], fluorophosphates (e.g., Na3V2(PO4)2F3) [14, 15], pyrophosphates (e.g., Na3.12Fe2.44(P2O7)2, Na3.12Mn2.44(P2O7)2) [16, 17], sulfates (e.g., Na0.97KFe(SO4)2, Na2.44Mn1.79(SO4)3) [18, 19], and mixed polyanions (e.g., Na7V4(P2O7)4(PO4)) [20].

Recently, Na2Fe2(SO4)3 has drawn considerable attention because of its lower cost, high safety, environment-friendly, and high Fe3+/Fe2+ redox potential of around 3.8 V [21,22,23,24,25]. Nevertheless, a vital problem of Na2Fe2(SO4)3 resides in its inherently poor electrical conductivity which leads to sluggish kinetics. As a result, Na2Fe2(SO4)3 experiences a fast capacity fading and poor rate capability. Tremendous studies have been committed to design rational electrodes with favorable structure. One of the most favored strategies for improving the electrochemical properties of polyanion materials is to use high conductive carbonaceous materials [26,27,28,29]. In this regard, Meng et al. synthesized an off-stoichiometric Na2 + 2xFe2-x (SO4)3/SWNT spindle-shaped composite as a cathode with improved rate capability and cycle stability for SIBs via a top-down strategy [30]. Yu et al. combined the method of electrospinning with electrospraying to fabricate freestanding Na2 + 2xFe2-x (SO4)3@porous carbon-nanofiber hybrid film [31]. Recently, we reported the successful synthesis of an off-stoichiometric Na2 + 2xFe2-x (SO4)3/rGO composite with outstanding rate performance through a facile one-stepped strategy [32].

Biomass is a renewable resource of low cost, abundant yield, rapid regeneration, and environmental friendliness. These attractive properties make them particularly appealing for the preparation of composite materials with various applications [33, 34]. As a kind of by-product of rice milling industry, rice husks (RHs) has an amazing amount of 2.9 × 107 t per year all over the globe [35]. They are mainly composed of organic lignin, cellulose, hemicellulose, and inorganic Si compound. Undergoing a series of processes, such as acidification, high temperature pyrolysis, and other processes, rice husk carbon (RHC) can be acquired from rice husks, which possesses ordered porous structure and has attracted much attention as a conductive and porous matrix of lithium/sodium-ion batteries owing to its uniform pore size, large specific surface area, interconnected porous structure, and high conductivity [36,37,38,39,40].

In this work, we report a facile strategy to prepare Na2 + 2xFe2-x(SO4)3@rice husk carbon (NFS@RHC) composite. In this NFS@RHC electrode, the NFS acts as the active material and the ordered mesoporous RHC serves as the electrode matrix carrier. The highly conductive RHC accelerates NFS electrochemistry. In comparison with bare NFS, the as-prepared NFS@RHC exhibits remarkable rate capability with specific capacity of 60 mAh g−1 at high current density of 5 C (1 C = 120 mA g−1). More importantly, the composite can maintain a high capacity of 30 mAh g−1 even at current density of 20 C. Besides the good rate capability, it also shows better cycle stability than the original sample NFS.

Materials and methods

Synthesis of alluaudite Na2 + 2xFe2-x(SO4)3

The target product of Na2 + 2xFe2-x(SO4)3 was synthesized via a classical solid-state method. Typically, in the first step, FeSO4 was prepared by heating FeSO4·7H2O at 200 °C for 12 h in an Ar/H2 (93:7, by volume ratio) atmosphere. Then, the FeSO4 precursor was mixed with off-stoichiometric Na2SO4 and ball-milled for 12 h in acetone. After being dried under vacuum at 60 °C, the mixture was then annealed in a tube furnace to 350 °C at a heating rate of 1 °C min−1 under Ar atmosphere for 24 h. Finally, Na2 + 2xFe2-x(SO4)3 material was obtained.

Synthesis of Na2 + 2xFe2-x(SO4)3@RHC

RHC was obtained from rice husks by acidification, high-temperature pyrolysis, and other processes [41]. RHC (0.0627 g) was mixed with FeSO4 (1.6 mmol, 0.2432 g) and Na2SO4 (1.2 mmol, 0.1705 g), and then ball-milled for 12 h at a rotation rate of 300 rpm in acetone, followed by drying under vacuum at 60 °C for 12 h. The collected product was further treated by heating at 350 °C for 24 h in Ar flow with the ramp rate of 1 °C min−1. The Na2 + 2xFe2-x(SO4)3@RHC was obtained when it was cooled down to room temperature.

Material characterization

Powder X-ray diffraction (XRD) was employed to identify the crystalline phases of the materials using a Bruker AXS D8 diffractometer equipped with Cu-Kα radiation (40 kV, 40 mA) in the 2θ range of 10–60°. Inductively coupled plasma-optical emission spectrometry (ICP-OES) was employed on a Thermo iCAP 7600 to quantify the elemental concentration of prepared materials. The content of RHC in the NFS@RHC material was determined by CHN elemental analysis using an Elementar Vario EL cube. Fe Mössbauer spectroscopy was collected in the transmission mode using a 57Co/Pd γ-ray source. The surface morphologies of the materials and surface structures of cycled electrodes were performed on the scanning electron microscopy (SEM; JSM-6700F). The detailed morphologies and the element distribution of NFS@RHC sample were detected by transmission electron microscopy (TEM; FEI Tecnai G2). Raman scattering data spectroscopy was evaluated by a Renishaw in Via Raman microscope with Ar-ion laser excitation (λ = 514.5 nm). X-ray photoelectron spectroscopy (XPS) was measured on an ESCALAB spectrometer using Mg-Kα light source. The particle size distribution results were performed based on a laser particle size analyzer (JL-6000). Nitrogen adsorption-desorption isotherms were performed at 77 K using a Micromeritics ASAP 2010 instrument. The specific surface area was calculated by the Brunauer-Emmett-Teller (BET) theory, and the pore-size distribution (PSD) curves were performed using the Barrett-Joyner-Halenda (BJH) algorithm. Electrical conductivity of samples were measured using four-probe test method.

Electrochemical measurements

The electrochemical performances of the obtained materials were evaluated using CR2032-type coin cells which were all assembled in Ar-filled glove box by using metallic sodium foil as the anode. In a typical experiment, the cathode electrodes were prepared by coating a slurry composed of 70 wt.% active material, 20 wt.% conductive additives (super P), and 10 wt.% polyvinylidenefluoride (PVDF) binder in N-methy-l-2-pyrrolidone onto aluminum foil. Later, the electrodes were dried in a vacuum oven at 120 °C for 12 h, pressed, and cut into the circular shapes with a diameter of about 1.0 cm. The average mass loading of the active material was controlled about 2.0 mg cm−2. The specific capacities were calculated according to the mass of Na2 + 2xFe2-x(SO4)3 excluding the carbon content. The separator was a glass fiber filter (Whatman GF/C) film. The electrolyte was 1.0 M NaClO4 dissolved in a mixed solution of ethylene carbonate (EC) and propylene carbonate (PC) (1:1, by volume ratio) with 5 wt.% fluoroethylene carbonate (FEC). Galvanostatic charge-discharge cycles were carried out using a LAND-2100 automatic battery tester in the voltage range from 2.0 to 4.3 V (vs Na+/Na). The cyclic voltammograms (CV) and electrochemical impedance spectroscopy (EIS) were performed on a Bio-Logic VSP multichannel potentiostatic-galvanostatic system. The scan rate for CV measurement was 0.1 mV s−1 and the frequency range of the EIS tests was 1 to 1 MHz with an alternative current voltage amplitude of 5 mV. Galvanostatic intermittent titration technique (GITT) was performed throughout the third discharge on the above VSP instrument. For each GITT step, the cell was discharged with a current flux of 30 mA g−1 for 0.5 h, followed by an open circuit stand for 4 h to reach the quasi-equilibrium state.

Results and discussion

The synthesis procedure of the NFS@RHC material includes ball milling and calcination which is schematically illustrated in Fig. 1. First, RHC, FeSO4, and Na2SO4 powders were mixed by ball milling in favoring the intimate contact between the different ions on the surface of RHC. Then in the second calcination step, NFS can readily form and deposit on the ordered mesoporous RHC. The operation was simple with the in situ formation of active material on the electrode matrix.

Fig. 1
figure 1

Schematic illustration of the synthesis procedure of the NFS@RHC

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The X-ray diffraction patterns of NFS and NFS@RHC are shown in Fig. 2. The Rietveld refinement of the XRD profile was carried out to investigate the structure of the as-prepared NFS. Obviously, all the observed diffraction peaks can be well fitted to the alluaudite-type structure with a monoclinic lattice and C2/c symmetry [23]. This result verifies that NFS sample is a pure alluaudite phase [42, 43]. The higher peak intensity of NFS@RHC sample manifests a greater degree of crystallinity compared with NFS. This is mainly because RHC can effectively inhibit the stacking and agglomeration of precursor during the reaction, which facilitates crystal growth in the calcination process. There is no phase transition during an in situ carbon-containing synthesis process, indicating that the addition of RHC would not change the crystal structure of the NFS material and introduce impurities. In addition, the XRD pattern of RHC clearly shows two normal characteristic peaks at 21 and 43°, corresponding to (002), (100) planes of carbon, respectively (Fig. S1). Noticeably, no diffraction peaks attributed to the carbon in the X-ray diffraction patterns of NFS@RHC can be observed because of the peak overlapping of NFS and the formation of an amorphous carbon. According to the result of ICP-OES, the exact molecular formula of Na2 + 2xFe2-x(SO4)3 can be determined as Na2.45Fe1.775(SO4)3 (x = 0.225). The content of carbon in the final product NFS@RHC was also given by CHN elemental analysis to be 12.78 wt.%.

Fig. 2
figure 2

XRD pattern of NFS@RHC (green line) and Rietveld refinement of XRD pattern of NFS showing the experimental profile (blue dots), the calculated pattern (black line), their difference (dark cyan line), and the Bragg diffraction positions (magenta ticks)

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In order to verify the chemical valence of Fe element in the samples, the detailed XPS of Fe 2p in NFS and NFS@RHC materials were carried out. As clearly observed in Fig. 3a, the two characteristic peaks at 711.4 and 724.9 eV are assigned to the Fe 2p3/2 and Fe 2p1/2, respectively, which are related to the Fe2+ species [43]. The binding energy of Fe 2p does not change after the addition of RHC, which is well consistent with XRD analyses. In addition, the XPS full spectra of NFS@RHC reveal the presence of Na, Fe, S, O, and C (Fig. S2), which is in accordance with the chemical element of the NFS@RHC. To provide sufficient evidence for the existence of RHC, Raman spectra are displayed in Fig. 3b. A peak at 1003 cm−1 in composite can be assigned to the pristine NFS. Two typical broad peaks at 1342 and 1588 cm−1 can correspond to carbon D peak and G peak [44, 45]. The D peak represents the sp3 hybridization form of carbon, which is the disordered carbon; the G peak represents the sp2 hybridization form of carbon, that is, the graphitized carbon. The ratio of the peak intensity of the D peak to the G peak (ID/IG) is often used to determine the crystallinity of the carbon material. Compared with RHC (1.022), NFS@RHC shows a lower ID/IG value of 0.955. The decrease in ID/IG ratio shows that the combination of rice husk carbon with NFS increases the degree of graphitization of carbon, which would be beneficial to enhance the electronic conductivity of NFS@RHC. As a result, electron migration of the composite can be greatly facilitated and the rate performance of NFS@RHC will be further improved. The Mössbauer spectra are shown in Fig. 3c, d. The two sets of peaks can be assigned to two distinct Fe (II) crystallographic sites, indicating the Fe (II) valence state and no Fe (III) impurities in both NFS and NFS@RHC. The fitting parameters of Mössbauer spectrum are shown in Table S1. The energy shift range of Fe (II) species is consistent with that of Mössbauer spectra of NFS in other literature, which confirms that Fe (II) is not oxidized during the synthesis process [46]. The good symmetry of Mössbauer spectrum pattern proves that RHC addition does not destroy the crystal structure of NFS.

Fig. 3
figure 3

a XPS spectra of Fe 2p for NFS and NFS@RHC composite. b Raman spectra of the RHC and NFS@RHC composite. Mössbauer spectra of c NFS and d NFS@RHC materials. The violet lines are the fit curves to the experimental data (dotted points); blue lines, LS Fe(II); green lines, HS Fe(II)

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Figure 4a–c exhibits the SEM images of the RHC, NFS, and NFS@RHC samples. As shown in Fig. 4a, the rice husk carbon keeps a lumpy and lamellar structure with thickness of approximately a few tens to hundreds of nanometers. The stacked RHC matrix forms a porous network, which makes it an ideal template for the preparation of hybrid carbon composites. The original NFS sample varies in size from 200 to 800 nm in Fig. 4b. The agglomeration phenomenon of the NFS is obvious, which leads to incomplete contact between active NFS nanoparticles and electrolyte. However, after using RHC modified, it is found in Fig. 4c that the NFS is evenly dispersed on the surface of RHC, which is beneficial to increase the contact area with electrolyte and obtain excellent electrochemical performance. Compared with the clear one-dimensional pores of the original RHC, some pores in the NFS@RHC sample have disappeared and the surface of the NFS@RHC is more rough. In order to further confirm the architecture of composite, TEM and HRTEM images of NFS@RHC are displayed in Fig. S3a, b. The distinct lattice fringes indicate the formation of NFS. Meanwhile, NFS particles are evenly deposited on the surface of RHC. To map the distribution of elements at nanoscale, TEM and the corresponding electron dispersive X-ray spectroscopy (EDS) were performed on a single NFS@RHC block (Fig. 4d–i). All the elements in NFS and C are shown to be uniformly distributed on the whole composite material block, which proves that RHC skeleton is completely covered by a continuous layer of NFS particles. The introduction of NFS into the porous carbon framework forms a good three-dimensional conductive network, which speeds up the electron and ion transfer rates and improves the electrochemical properties of the materials. To clearly demonstrate the reduced agglomeration after combination, the particle size distribution of the NFS and NFS@RHC materials is determined by the particle size analyzer (Fig. S4). Compared with NFS sample, the peak density distribution curve of NFS@RHC shifts toward smaller particle diameter. This indicates that the particle size of NFS@RHC sample is smaller than that of NFS sample. The addition of rice husk carbon can prevent particle accumulation and facilitate the acquisition of smaller and more dispersed particles.

Fig. 4
figure 4

SEM images of a RHC, b NFS, and c NFS@RHC materials. d TEM image of NFS@RHC composite and the corresponding EDS elemental mapping images for e carbon, f iron, g sodium, h oxygen, and i sulfur

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The BET surface area and the corresponding pore size distribution of samples are shown in Fig. 5. It can be seen that both of the shapes of isotherms for RHC and NFS@RHC correspond to type I adsorption-desorption curve, which is typical of microporous materials [47]. The pore size distribution of material is calculated using the BJH model (inset in Fig. 5). The pore size distribution peaks of RHC and NFS@RHC are centered at 1.08 nm and 1.14 nm (Table S2), respectively. The specific surface area and pore volume of pure rice husk carbon are calculated to be 1899 m2 g−1 and 1.02 cm3 g−1, respectively, which are higher than those of NFS@RHC (350 m2 g−1 and 0.19 cm3 g−1; Fig. 5b). After NFS loading, the specific surface area and pore volume are greatly reduced. There are two main reasons: one reason is that a large number of NFS particles are deposited on surface of RHC substrate to inevitably block its channels; the other reason is that a fraction of the pores in RHC collapse during the process of high-energy ball milling. Nevertheless, the specific surface area of the composite material is still larger than non-carbon-compounded NFS materials with specific surface area of merely 6.18 m2 g−1 (Fig. S5). The high pore volume and large specific surface area of NFS@RHC composite can offer sufficient contact area with electrolyte, which can effectively improve the electrochemical performance.

Fig. 5
figure 5

N2 adsorption–desorption isotherms of a RHC and b NFS@RHC composite. Insets are their corresponding BJH pore size distributions curves

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The electrochemical properties of the materials were evaluated in Na half-cell configuration. In order to analyze the electrochemical redox process of NFS@RHC as a cathode material for sodium-ion batteries, we tested the cyclic voltammogram of the material at different scan rates in the voltage range of 2.0–4.3 V, as shown in Fig. 6a. The two oxidation peaks appear at around 3.67 V and 4.13 V during the first charge. But during the first discharge, there are three reduction peaks at 3.91, 3.74, and 3.38 V. According to the previous literature [48], the irreversible structural changes of NFS occur during the initial cycle. In the subsequent cycles, three pairs of symmetric redox peaks can be clearly observed from the CV curve of NFS@RHC, corresponding to the reversible transformation of Fe2+/ Fe3+ redox couple [49]. The first pair of oxygen reduction peaks corresponds to the first Na+ withdrawn/embedded from the Na3 site, and the latter two of 3.87/3.74 V and 4.10/3.96 V correspond to the Na extraction/insertion at the Na2 and Na1 sites, respectively. However, for pristine NFS in Fig. S6, the redox process of CV curve is not obvious and the peak shape cannot maintain the symmetry with the increase of scan speed, which implies that the addition of RHC makes NFS have good reversibility and structural stability during the Na+ intercalation and prolapse.

Fig. 6
figure 6

a Cyclic voltammogram curves at different scanning rates of NFS@RHC. b The charge-discharge voltage profiles of NFS and NFS@RHC for the first cycle at a current density of 0.1 C. c Cycling performance at the 0.1 C and d rate capability at various current rates of NFS and NFS@RHC in the voltage range of 2.0–4.3 V

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Figure 6b shows the initial charge-discharge profiles of the bare NFS and NFS@RHC electrodes cycled between 2.0 and 4.3 V at 0.1 C. The NFS@RHC electrode possesses a voltage plateau around 3.8 V. However, the bare NFS electrode displays discharge/charge profiles around 3.7 V with larger polarization. The initial charge capacity of the NFS@RHC electrode is 113.4 mAh g−1 and discharge process delivers a reversible capacity of 96.8 mAh g−1 with an initial coulombic efficiency of 85.4%, much higher than that of the NFS (62.9 mAh g−1 and 52.5 mAh g−1, 83.5%) at 0.1 C.

To further verify the structural advantages of NFS@RHC, the long cycling performances of the NFS@RHC hybrid and pristine NFS at 0.1 C are compared in Fig. 6c. The reversible capacity of NFS@RHC sustains 81.2 mAh g−1 even after 100 cycles with a high capacity retention of 83.9% with respect to the first cycle. In contrast, the discharge capacity of the pristine material decreases to 30.3 mAh g−1 after 100 cycles, corresponding to relatively low capacity retention of 57.7%.

In addition to its good cycling performance, the composite also has excellent rate capability, as shown in Fig. 6d. The outstanding specific capacity of the composite material at various current densities of 0.1, 0.2, 0.5, 1, and 2 C can reach 95, 90, 83, 77, and 70 mAh g−1, respectively, whereas the cell without RHC exhibits a rather dramatic capacity decay. Especially when the current density continues to increase to 5 C, NFS@RHC can still show a reversible capacity of 60 mAh g−1. Nevertheless, the NFS can scarcely charge and discharge at high current density. Remarkably, the NFS@RHC exhibits a stable specific capacity of 30 mAh g−1 at 20 C (current density of 2.4 A). When the discharge current returns to 0.1 C, the specific capacity of the composite recovers to 94 mAh g−1 and remains stable.

To further understand the effect of RHC on the electrical conductivity of materials, we conducted AC impedance tests of NFS and NFS@RHC electrodes after first charge at 0.1 C. The plots are fitted by modeling AC impedance spectra based on the simplified equivalent circuit given in the inset of Fig. 7. A whole Nyquist plot consists of one semicircle at high-frequency region and one semicircle at medium-frequency area, followed by a slope line at the low-frequency zone. By knowing that the intercept of the left-hand side of the curve with the X axis represents the internal resistance (Rs) of the cell, the semicircle of the high-frequency region can be explained as the surface film resistance (Rf) in the electrode. The semicircle of the mid-frequency region is related to charge transfer resistance (Rct), while the long straight line at the low-frequency region is equivalent to the Warburg diffusion region formed by the diffusion of Na ions inside the cell [50]. The fitting results show that the Rs, Rf, and Rct of NFS@RHC material are much lower than those of the NFS material (presented in Table S3). This result confirms that the high conductivity of RHC combined with larger surface area of NFS@RHC can improve electron and ion transfer capabilities of the NFS@RHC composite.

Fig. 7
figure 7

Nyquist plots and equivalent circuit of the NFS and NFS@RHC after first charge

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Since the electrical conductivity is another key parameter to characterize the properties of materials, we evaluated the electrical conductivity of RHC, NFS, and NFS@RHC samples using four-probe test method. Particularly, RHC exhibits a comparable electrical conductivity (28.6 S m−1). The conductivity of the NFS@RHC (0.6 S m−1) is much higher than that of the bare NFS (1.5 × 10−13 S m−1). Such superiority is further validated by the enhanced electrochemical performance.

Figure 8a shows the GITT curve as a function of voltage and time during the third discharge for NFS and NFS@RHC. Each cell titration test was discharged at a current density of 60.0 mA g−1 for 0.5 h and then left open for 4 h to bring the voltage to quasi-static after this process is repeated until the cut-off voltage. Figure 8b is a plot of E vs τ1/2 for NFS and NFS@RHC. Between 60 and 300 s, the curve is close to linear, and the slope of its straight-line fit is used to calculate dE/dτ1/2. The apparent diffusion coefficient (DNa+) can be calculated according to equation, derived by Weppner and Huggins as [51]

$$ {\mathrm{D}}_{{\mathrm{Na}}^{+}}=\frac{4}{\pi \tau}{\left(\frac{m_B{V}_M}{M_BS}\right)}^2{\left(\frac{\Delta {E}_s}{\Delta {E}_{\tau }}\right)}^2 $$

where τ is the time for an applied galvanostatic current. VM is the molar volume of active material, and MB and mB are the molecular weight and the mass of active material, respectively. S is the BET surface area of the materials. ΔEs and ΔEτ are the steady-state (equilibrium) voltage and the total change of the cell voltage E during the current pulse. As shown in Fig. 8c, the Na+ diffusion coefficient of NFS@RHC at the beginning of discharge is 7.56 × 10−14 cm2 s−1 and then falls to 3.10 × 10−14 cm2 s−1. Afterward, DNa+ increased to 3.31 × 10−14 cm2 s−1, and dropped slightly to 3.39 × 10−14 cm2 s−1 again. Then, DNa+ increased to 3.47 × 10−14 cm2 s−1, and then decreased to the minimum 9.38 × 10−15 cm2 s−1. Finally, at the end of the discharge, DNa+ increased to 1.20 × 10−14 cm2 s−1. Obviously, the change of DNa+ corresponds to the peak value of CV. The minimum value of DNa+ on the discharge platform is 9.38 × 10−15 cm2 s−1, which is nearly 18 times higher than that of NFS. Thus, the presence of RHC improves the Na+ ion diffusion coefficient of NFS.

Fig. 8
figure 8

a GITT curves of NFS and NFS@RHC samples during the third discharge. b Typical E vs τ1/2 profile of the samples. c Sodium diffusion coefficients of NFS and NFS@RHC during the third discharge

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For a deeper understanding of the influence of RHC on the surface morphology of electrode after cycling, NFS@RHC and NFS electrodes were investigated by SEM after 100 cycles at 0.1 C. As shown in Fig. S7a, b, the morphology of the electrode of NFS@RHC is flat and unbroken. However, the cycling NFS electrode gives rise to a rough surface with large number of bulges and cracks. These observations indicate that RHC, as a large carbon skeleton, can prevent the positive electrode materials from being damaged.

Conclusions

In summary, NFS cathode material was well modified using biomass rice husk carbon. The NFS@RHC composite was successfully fabricated through a conventional solid-state method at low temperature. A large amount of NFS engages on the surface of RHC. The NFS@RHC composite cathode exhibits excellent long cycling stability and impressive high-rate capability. The first charge and discharge specific capacities of NFS@RHC were 113.4 mAh g−1 and 96.8 mAh g−1 at 0.1 C, respectively. About 83.9% of the initial capacity was maintained even after 100 cycles. Meanwhile, the remarkable rate performance of the composite material at the current density 5 C can reach 60 mAh g−1, which is much higher than bare NFS material. In particular, the specific capacity of about 30 mAh g−1 can still be achieved at a rate of 20 C (2.4 A). This means that RHC not only is a conductive substrate for improving the cathode conductivity but also can significantly inhibit the stacking of NFS in the electrochemical process. The present work paves the way for using low-cost, environment-friendly avenues to facilitate the development of high-performance cathodes for extensive applications in large-scale sodium-ion batteries.