Tailoring Heterogeneous Antiphase Boundaries in Multi‐Shelled Structure to Boost Na⁺ Diffusion Kinetics for High‐Rate Sodium–Ion Batteries

ABSTRACT

MoS₂ is a promising anode of sodium–ion batteries (SIBs), owing to high theoretical capacity and environmental benignity. However, its sluggish Na⁺ diffusion kinetics severely impair rate capability and cycling stability. In this study, we rationally designed MoS₂/Fe_xMo₂S₄ hollow multi‐shelled structure (HoMS). Wherein, the precise regulation of Mo/Fe molar ratio induces the formation of unique alternating antiphase boundaries (APBs) at the heterointerfaces reported for the first time arising from the lattice mismatch of (002) crystal planes between MoS₂ and Fe_xMo₂S₄. These APBs provide excess diffusion channels for Na⁺ in lattice, accompanying unique structure of HoMS to achieve 3D Na⁺ diffusion path. Additionally, the electron conductivity is promoted due to the construction of a built‐in electric field, thus facilitating fast reversible electrochemical reaction for high rate capability. Meanwhile, the void space in HoMS with APBs buffers volume change during cycling, enhancing cycling stability. Particularly, the MoS₂/Fe_xMo₂S₄ hollow double‐shelled sphere exhibits a high specific capacity of 743.6 mAh g⁻¹ at 0.2 A g⁻¹, and a capacity of 226.4 mAh g⁻¹ is still retained at 10 A g⁻¹, demonstrating excellent rate capability. This work provides an inspiration for constructing hollow multi‐shelled structure anode materials with abundant APBs lattice mismatch for fast and reversible electrochemical reactions.

The Fe_xMo₂S₄/MoS₂ HoMS around 200 nm with APBs heterojunction (MFSH‐HoMS‐D) was synthesized successfully via controlling the Mo/Fe ratio in a modified sequential templating, which brings excess diffusion channels for Na⁺ to further shorten Na⁺ diffusion path and triggers formation of a built‐in electric field to promote fast electron transfer. As anode of SIBs, the high specific capacity, excellent rate capability, and cycling stability are achieved.

1. Introduction

MoS₂ is a kind of eminent material which has been applied in energy storage [1, 2, 3], electrochemical catalysis [4, 5], drug delivery [6, 7], desalination of seawater [8, 9], and so on due to its advantages as environmentally friendly and nontoxic. Especially, it is regarded as a very promising electrode material of SIBs because of its unique 2D plane structure, which is composed of Mo (+4) and S (−2) arranged by the covalent bond as a sequence of S─Mo─S [10, 11]. The unique 2D layers stacked via Van der Waals' force are beneficial to host Na⁺ during the charging/discharging process [12, 13, 14]. However, the sluggish electron transfer rate and Na⁺ diffusion kinetics restrict further development of cycling stability and rate capability [15]. Thus, designing a delicate approach to improve electron conduction/transfer and Na⁺ diffusion in MoS₂ for excellent cycling stability and rate performance is desirable.

Hollow multi‐shelled structure (HoMS) features with a layer‐by‐layer structure with controllable shell thickness [16, 17, 18], large space between adjacent layers, and abundant pores through the layer, enabling the permeation of electrolyte from outer into inner to increase active sites for redox reaction [19, 20]. Coupled with thin shell thickness, the diffusion path of Na⁺ from electrolyte to the surface of HoMS can be shortened to improve Na⁺ diffusion kinetics [21, 22, 23]. Meanwhile, the free space between adjacent layers can alleviate volume change during the charging/discharging process to improve cycling stability [24, 25]. Constructing HoMS is an effective way to promote the diffusion kinetics of Na⁺ from electrolyte to surface and alleviate volume expansion. For example, starting with triple‐shelled WO₃ HoMS precursor, sextuple‐shelled WS_2−x heterogeneous HoMS could be obtained by precisely tuning the nanoscale Kirkendall diffusion effect in the sulfuration reaction following the chemical etching. Mechanism analyses revealed that HoMS provides good electrode–electrolyte contact and plentiful sodium storage sites as well as an effective buffer of the stress/strain during cycling [26]. However, the intrinsic electronic conductivity and Na⁺ diffusion in bulk as well as the lattice stress can't be improved just through designing HoMS.

Construction heterojunction is an effective method for improving the diffusion rate of electrons, increasing the storage site of Na⁺, and relieving lattice stress in MoS₂ to adjust rate capability and cycling stability [27, 28, 29]. Usually, two approaches, including physical and chemical paths, are applied to construct a lattice heterojunction. The physical interaction can trigger a partial irregular arrangement of the lattice face [30]. However, a lot of defects caused by the strong physical interaction will bring many side reactions in MoS₂, which is harmful for cycling stability as the anode of SIBs [31]. Another way is doping heteroatoms into lattice planes to generate a lattice heterojunction [32], which could cause the arrangement of atoms to trigger lattice distortion, thereby promoting Na⁺ diffusion rate in MoS₂ lattice and improving electron diffusion rate [33]. Therefore, constructing the MoS₂ HoMS with heterojunction structures can optimize 3D Na⁺ diffusion path from electrolyte to HoMS particles and improve Na⁺ diffusion kinetics in MoS₂ lattice, and promote electron conductivity for fast reversible electrochemical reaction, as well as alleviate volume change and lattice stress for excellent cycling stability. However, the overlong Na⁺ diffusion path in elongated MoS₂ lattice is limited by just designing MoS₂ HoMS with heterojunction structure to further maximize the Na⁺ diffusion kinetics.

In this study, a HoMS with unique alternating MoS₂/Fe_xMo₂S₄ antiphase boundaries (APBs) heterojunction lattice mismatch (MFSH‐HoMS) was successfully fabricated for the first time via a modified sequential templating approach, wherein both Fe³⁺ and MoO₄ ²⁻ were adsorbed into carbon spheres, followed by calcination in air and subsequent sulfuration in N₂ gas. During the adsorption process, the ratio of Fe³⁺ to MoO₄ ²⁻ was precisely adjusted to control the morphology and lattice structure. Notably, only when the ratio of MoO₄ ²⁻ to Fe³⁺ was 1/2, could the alternating MoS₂/Fe_xMo₂S₄ APBs heterojunction hollow double‐shelled sphere (MFSH‐HoMS‐D) be constructed. MFSH‐HoMS‐D with a larger specific surface area can provide more active sites for redox reaction, combined with thin shell to shorten the diffusion path of Na⁺ from electrolyte to HoMS particles. In addition, the voids between shells can alleviate volume expansion during the charging or discharging process. Furthermore, the original alternating MoS₂/Fe_xMo₂S₄ APBs heterojunction lattice mismatch is similar to the periodic structure of an organic polymer with two types of ligands; lattice APBs mismatch generates at the interface between MoS₂ and Fe_xMo₂S₄ due to the (002) lattice mismatch. This unique structure not only releases the Na⁺ intercalation/deintercalation induced lattice stress during cycling, but also brings excess diffusion channels for Na⁺ in MoS₂ lattice to further shorten Na⁺ diffusion path and triggers the formation of a built‐in electric field to facilitate fast electron transfer. Therefore, as anode of SIBs, the MFSH‐HoMS‐D delivers a high specific capacity of 743.6 mAh g⁻¹ at a current density of 0.2 A g⁻¹, retaining 226.4 mAh g⁻¹ at a super‐high current density of 10 A g⁻¹, demonstrating excellent rate capability. What's more, it achieves an optimized reversible capacity of 434.5 mAh g⁻¹ after 500 cycles with about 83.2% retention rate at a current density of 5 A g⁻¹.

2. Results and Discussion

The MFSH‐HoMS‐D was synthesized via a modified sequential templating approach, wherein the Fe³⁺ and MoO₄ ²⁻ both were adsorbed into carbon sphere followed by calcination in air and sulfuration in N₂ gas. The ratio of Mo/Fe during the adsorption process was adjusted, and the structure after calcination in air was explored. As shown in Figures S1 and S2, only the metal oxide particles with double‐shelled structure were generated when the ratio is 1/2. The microstructure of the final product with different Mo/Fe ratios was investigated by scanning electron microscope (SEM) and transmission electron microscope (TEM) images. As shown in Figures S3 and S4, a uniform spherical structure with the size at the range of 100–200 nm was detected at the Mo/Fe ratio of 1/2, and the TEM images demonstrated that a clear double‐shelled hollow structure with obvious interspace between the shells and huge cavities on the shells for alleviating volume change during the charging and discharging process (Figure 1a). Compared to the rod‐shaped structure of pure MoS₂ and the aggregated or broken single‐shelled hollow structure of samples with other Mo/Fe molar ratios, MFSH‐HoMS‐D exhibits excellent structure stability. Subsequently, as shown in Figure 1b, the element distribution was clarified by the element mapping image. The elements of Fe, Mo, and S are homogeneously distributed on the double‐shelled hollow sphere, from the side providing a uniformly distributed alternating APBs heterojunction. Two phases of Fe_xMo₂S₄ and MoS₂ were explored by the selected area electron diffraction (SAED) image (Figure 1c). Four distinct diffraction rings indexed to the (002), (110) planes of MoS₂ and the (−315), (040) planes of Fe_xMo₂S₄ indicate the formation of heterojunction [34]. The alternating APBs heterojunction was testified by the high‐resolution transmission electron microscopy (HRTEM) as shown in Figure 1d. The MoS₂ and Fe_xMo₂S₄, and the obvious lattice fracture due to the lattice mismatch marked by the red boxes (Figure 1d1‐d4) can be observed alternately. It corresponds to alternating APBs heterojunction with (002) lattice space of 0.613 nm in MoS₂ and (200) lattice plane of 0.593 nm in Fe_xMo₂S₄ (Figure 1e) [35], respectively, which brings excess diffusion channels for Na⁺ to further shorten Na⁺ diffusion path and triggers formation of a built‐in electric field to promote fast electron transfer [36]. In addition, the corresponding geometric phase analysis (GPA) results show that both shear and axial strains (compressive or tensile) in MFSH‐HoMS‐D are small, causing uniform and small internal stress (Figure 1f). Obviously, the MoS₂/Fe_xMo₂S₄ hollow double‐shelled sphere with alternating APBs heterojunction around 200 nm was constructed, which can mitigate volume change and improve utilization during charge–discharge cycles, as well as promote Na⁺ diffusion kinetics and electron transfer rate (Figure 1g).

FIGURE 1.

(a) The TEM image, (b) the elemental mapping images, (c) the selected area electron diffraction, and (d) the HRTEM images of MFSH‐HoMS‐D. (e) Inverse FFT lattice images of Figure 1d2. (f) Corresponding stress distributions to Figure 1d2 obtained by geometric phase analysis (GPA). (g) The advantages of HoMS and APBs heterogeneous interface in MFSH‐HoMS‐D as anode of SIBs.

The phase characterization of MFSH‐HoMS‐D at different stages during preparation was studied. As shown in Figure S5, the precursor after calcination in air corresponds to MoO₃ and Fe₂(MoO₄)₃. With the increasing ratio during adsorption process, the phase varies from MoO₃ to Fe₂(MoO₄)₃. The Fe₂(MoO₄)₃ is obtained at a Mo/Fe ratio of 1/2, even with a higher ratio than 1/2, which may be because the excess Fe ion was doped into the lattice of Fe₂(MoO₄)₃. The x‐ray diffraction (XRD) was conducted to further investigate the phase after sulfuration (Figure 2a). Two phases of MoS₂ (PDF#37‐1492) and Fe_xMo₂S₄ (PDF#71‐0379) were detected [37, 38], which changed from MoS₂ to Fe_xMo₂S₄ with the evolution of Mo/Fe ratio from 1/0 to 1/7. When the ratio of Mo/Fe is 1/2, both MoS₂ and Fe_xMo₂S₄ are explored, which is consistent with the HRTEM and SAED results as shown in Figure 1c,d. In addition, the presence of both Mo─S and S─S bonds was further detected by Fourier Transform Infrared Spectroscopy (Figure S6). With an increase in the Mo/Fe ratio, the peak position of the Mo─S bond shifts to high wavenumber gradually [39], signifying a shortened bond length due to the formation of APBs heterojunction. The electronic structure information of MFSH‐HoMS‐D and the valence states of Fe and Mo elements were analyzed by x‐ray photoelectron spectroscopy (XPS). As shown in the high‐resolution XPS spectrum of Mo 3d (Figure 2b), the valence state of Mo in MoS₂ is +4. The Mo²⁺ appears along with the Mo⁴⁺ in MFSH‐HoMS‐D, corresponding to the valence states of Mo in Fe_xMo₂S₄ [40]. And the peak intensity attributed to Mo²⁺ increases in the sample of MFSH‐HoMS‐1/7 due to the more Fe_xMo₂S₄ phase. Besides, the presence of Fe²⁺ in MFSH‐HoMS‐D was confirmed by the high‐resolution XPS spectrum of Fe 2p in Figure 2c [41].

FIGURE 2.

(a) XRD patterns of samples with different Mo/Fe ratio. (b) The Mo 3d and (c) the Fe 2p XPS high‐resolution analysis of MoS₂, MFSH‐HoMS‐D, and MFSH‐HoMS‐1/7. The Mo k‐edge (d) XANES, (e) EXAFS, and (f) the Mo EXAFS spectra in K‐space analysis, the Fe k‐edge. (g) XANES; (h) EXAFS. (i) The Fe EXAFS spectra in K‐space analysis of MFSH‐HoMS‐D. (j,k) Wavelet‐transformed (WT) Mo K‐edge EXAFS oscillations analysis of MFSH‐HoMS‐D and MoS₂. (l,m) Wavelet‐transformed (WT) Fe K‐edge EXAFS oscillations analysis of MFSH‐HoMS‐D and FeS₂.

To further analyze the chemical states of Fe and Mo, in‐depth exploration was conducted using Mo and Fe k‐edge x‐ray absorption near‐edge structure (XANES) spectra, wherein the Fe foil, Mo foil, FeS₂ and MoS₂ were taken as the references. As described in Figure 2d, compared with Mo foil and standard MoS₂, it can be determined that the total valence states of Mo in MFSH‐HoMS‐D were between 0 and +4 according to the position of the front edge line, which gave direct evidence for the mixed valence states in Mo of +2 and +4. Meanwhile, the front edge line of Fe in MFSH‐HoMS‐D was near the Fe²⁺ of standard FeS₂, indicating that the valence state of Fe was close to +2 [42]. The coordination environment can be characterized by x‐ray absorption fine structure (EXAFS) in Mo and Fe R‐space. As shown in Figure 2e, the Mo─S and Mo─Mo bonds in MFSH‐HoMS‐D were located at 1.85 and 2.96 Å [43]. Interestingly, a new Mo─Fe bond at 2.57 Å, which is shorter than Mo─Mo bond, confirmed the formation of unique Fe_xMo₂S₄, proved by the relatively low intensity of the Mo─S bond peak compared to the pure MoS₂ peak. As shown in Figure 2h, a clear peak at 1.92 Å belonging to Fe─S bond was detected [44], which is longer than the Fe─S bond in standard FeS₂ (1.8 Å) due to the electron interaction at the boundary of APBs heterojunction. Combining with other characterizations, this proved the existence of Fe_xMo₂S₄ and formation of APBs heterojunction, which was further confirmed by the coordination environment changes of MFSH‐HoMS‐D in Mo and Fe EXAFS spectra of K‐space analysis (Figure 2f,i) and the change of the maximum signal center in wavelet transform (WT) plot (Figure 2j–m).

The formation mechanism of the unique alternating MoS₂/Fe_xMo₂S₄ APB heterojunction was elucidated by tracking phase evolution throughout the sulfuration process. At the initial stage, iron molybdenum oxide was identified. After 5 min of sulfuration at 350°C, only MoS₂ was produced (Figure 3a). We hypothesized that during this early sulfuration stage, Fe species were solely doped into the MoS₂ lattice, inducing local lattice distortion (Figure 3c1) without triggering phase transformation. This observation was corroborated by high‐resolution transmission electron microscopy (HRTEM) and highlighted with a yellow rectangle in Figure 3d. Prolonging the sulfuration time—even up to 2 h—did not yield any additional phases. Upon elevating the temperature to 600°C for a further 1 h, the thermodynamically more stable Fe_xMo₂S₄ phase emerged (Figure 3b). Owing to its distinct lattice spacing relative to MoS₂, pronounced lattice mismatch was observed (Figure 3c2,e). Extending this high‐temperature treatment to 2 h led to the formation of abundant APB heterojunctions (Figure 3c3), with both MoS₂ and Fe_xMo₂S₄ phases clearly resolved in the corresponding HRTEM and fast Fourier transform (FFT) images (Figure 3f). When Fe³⁺ ions are doped into the MoS₂ lattice to form the Fe_xMo₂S₄ phase, ordered MoS₂ and Fe_xMo₂S₄ domains are generated within the crystal matrix; the misalignment in atomic arrangements between these two domains directly engenders the formation of APBs.

FIGURE 3.

The formation mechanism for the formation of MoS₂/Fe_xMo₂S₄ APBs heterojunction. (a) XRD patterns of MFSH‐HoMS‐D in different stages of sulfuration and without carbonization. (b) XRD patterns of MFSH‐HoMS‐D in different stages of carbonization and with 2.0 h sulfuration. (c) Schematic diagram of phase at different stages during the sulfuration process. HRTEM for sulfuration at (d) 0.5, (e) 1.0, and (f) 2.0 h of MFSH‐HoMS‐D with the same carbonization time. (g) HRTEM of MFSH‐HoMS‐2/1. (h) HRTEM of MFSH‐HoMS‐1/7. (i) The formation process from oxides to sulfides.

Phase compositions with Mo/Fe molar ratios of 2:1 and 1:7 were also investigated. In contrast to MFSH‐HoMS‐D, MFSH‐HoMS‐2/1 exhibited only the long‐range lattice planes of MoS₂, while MFSH‐HoMS‐1/7 displayed solely Fe_xMo₂S₄, with minimal APB‐induced lattice mismatch in both cases (Figure 3g,h). These results demonstrate that at elevated temperatures, Fe‐doped MoS₂ remains stable at lower Fe contents, whereas Fe_xMo₂S₄ dominates as the thermodynamically favorable phase at higher Fe loadings. Only when the Mo/Fe molar ratio is tuned to 1:2 can a well‐defined alternating APB structure be constructed, as this stoichiometry ensures commensurate domain sizes of MoS₂ and Fe_xMo₂S₄, thus facilitating uniform distribution of phase misalignment. At such a moderate stoichiometric ratio, a high density of intrinsic defects (e.g., vacancies and interstitial atoms) is generated [45]. To minimize the overall system energy, the crystal mitigates lattice distortion caused by these defects via the formation of APBs—low‐energy defect configurations that effectively dissipate the stress arising from stoichiometric deviations (Figure 3i).

The electrochemical performance of MFSH‐HoMS‐D as anode material for SIBs was explored by assembling a half‐cell, taking the samples with other Mo/Fe ratios as comparison. To reveal the sodium–ion storage process, cycling voltammetry (CV) tests and galvanostatic charge/discharge curves in the first circle were carried out at a voltage range of 0.01–2.8 V. Figure 4a and S7 show the CV curves of samples with different Mo/Fe ratios at a scan rate of 0.5 mV s⁻¹. There are three pairs of redox peaks corresponding to the cathode peaks of 0.46, 1.41, and 1.87 V and anode peaks of 0.32, 0.97, and 1.74 V, which are attributed to the intercalation and extraction of Na⁺ in MoS₂ and Fe_xMo₂S₄ [47]. MFSH‐HoMS‐D exhibits the smallest voltage difference between cathode and anode peaks. Correspondingly, the three platforms with the smallest voltage difference compared to other samples during the charging or discharging process in the galvanostatic charge–discharge curves indicate the lowest polarization in MFSH‐HoMS‐D due to fast Na⁺ and electron diffusion rate resulting from its unique HoMS structure and APBs heterojunction (Figure S8). The rate performance was checked by the galvanostatic charge–discharge test at different current densities. As shown in Figure 4b and S9, the MFSH‐HoMS‐D electrode exhibits higher capacities of 743.6, 635.1, 556.4, 473.4, 380.5, 334.3, and 226.4 mAh g⁻¹ at 0.2, 0.5, 1, 2, 4, 5, and 10 A g⁻¹ than that of other samples, respectively. 161.5 mAh g⁻¹ is still retained at an ultrahigh current density of 20 A g⁻¹ with an obvious charging or discharging platform (Figure S10). When the current density comes back to 0.2 A g⁻¹, the electrode can also recover the initial specific capacity, verifying the high reversibility of MFSH‐HoMS‐D due to its fast Na⁺ diffusion rate and electron transfer rate. To further evaluate the rate performance of this work, we compared the rate performance with other literature (Figure 4c) [46, 47, 48, 49, 50]. Obviously, the MFSH‐HoMS‐D exhibits higher specific capacity at different current densities, meaning an excellent rate capability. The cycling stability was explored by the galvanostatic charge–discharge test at 0.5, 1.0 and 5.0 A g⁻¹. Figure 4d and S11 reveal the cycling stability of the materials with different molar ratios of Mo/Fe at a current density of 0.5 A g⁻¹. As can be disclosed, the MFSH‐HoMS‐D displays a higher specific capacity of 510.1 mAh g⁻¹ after 100 cycles. The higher initial specific capacity of MFSH‐HoMS‐1/7 (with a Mo/Fe ratio of 1/7 and a fragmented single‐shell structure) during the first ∼30 cycles in Figure 4d can be attributed to its enhanced electrolyte accessibility and abundant exposed active sites. The long‐term cycling performance test with a high current of 1 A g⁻¹ has been designed (Figure 4e). It expressed the optimized reversible performance of 532.4 mAh g⁻¹ after 500 cycles with a 78.6% retention rate compared to the second cycle. Meanwhile, to demonstrate that MFSH‐HoMS‐D significantly improves Na⁺ diffusion kinetics, the long‐term cycling stability of the electrode was further tested (Figure 4e) at a high current density of 5 A g⁻¹. The initial discharge specific capacity of MFSH‐HoMS‐D is 522.7 mAh g⁻¹, and after 500 cycles, it exhibits a high specific capacity of 434.5 mAh g⁻¹ with a retention rate of 83.2%. In Figure 4f, even compared with previously reported iron or molybdenum‐based sulfide anode materials without conductive carbon for SIBs [51, 52, 53, 54], our work exhibited a certain of advantages due to its unique HoMS structure with APBs heterojunction.

FIGURE 4.

(a) The CV curves at a scan rate of 0.5 mV s⁻¹. (b) Rate capability test of MoS₂, MFSH‐HoMS‐D, and MFSH‐HoMS‐1/7. (c) Comparison of rate capability at various current densities [46, 47, 48, 49, 50]. (d) Cycling stability at a current density of 0.5 A g⁻¹. (e) Long cycling stability at a high current density of 1.0 A g⁻¹ and 5.0 A g⁻¹. (f) Comparison of cycling stability with previously reported MoS₂ and FeS₂ anode materials [51, 52, 53, 54]. (g) Schematic illustration of the MFSH‐HoMS‐D//NVP@C full cell. (h) GCD curves of the first cycle for MFSH‐HoMS‐D anode and NVP@C cathode in half cell. (i) The CV curves at different scan rates. (j) Rate capability test. (k) The charging and discharging curves at different current density. (l) Cycling stability at a current density of 0.1 A g⁻¹ and 5.0 A g⁻¹ of MFSH‐HoMS‐D//NVP@C full cell.

To evaluate the practical application of the material, we make a full cell with MFSH‐HoMS‐D as the anode material, and NVP@C reported previously as the cathode (Figure 4g) [55]. To match the capacity between cathode and anode materials, the specific capacity and behavior of NVP@C were explored by the galvanostatic charge–discharge test, as shown in Figure 4h and S12a. The specific capacity of NVP@C is 115 mAh g⁻¹ at a current density of 0.5 A g⁻¹. Based on the ratio of specific capacity between cathode and anode, the mass ratio of MFSH‐HoMS‐D to NVP@C is about 1/5.2. The electrochemical behavior of the MFSH‐HoMS‐D full‐cell was measured by the CV and galvanostatic charge–discharge tests. As shown in Figure S12b, two distinct charging and discharging platforms appear at a current density of 0.1 A g⁻¹. Correspondingly, at the scan rate of 0.1 mV s⁻¹, two pairs of redox peaks appear at 1.99 and 2.55 V, 1.65 and 2.05 V (Figure 4i). Meanwhile, with the increase of the scan rates from 0.1 to 0.5 mV s⁻¹, the peaks shift slowly, meaning excellent rate capability, which is confirmed by the galvanostatic charge–discharge tests at different current densities. As displayed in Figure 4j, the specific capacities of 178.3, 148.8, 137.6, and 83.7 mAh g⁻¹ are achieved in the second cycle at 0.05, 0.1, 0.2, and 0.5 A g⁻¹, and a specific capacity of 157.9 mAh g⁻¹ is still retained when the current density returns to 1.0 A g⁻¹. This demonstrates the excellent rate property of MFSH‐HoMS‐D in full cell. The cycling stability of the full‐cell was checked by the galvanostatic charge–discharge tests at 0.1 A g⁻¹ and 5.0 A g⁻¹. A specific capacity of 131.1 mAh g⁻¹ is maintained after 150 cycles at 0.1 A g⁻¹ and a specific capacity of 46.5 mAh g⁻¹ is maintained after 150 cycles at 5.0 A g⁻¹ (Figure 4l). Therefore, both the Na//MFSH‐HoMS‐D half‐cell and the MFSH‐HoMS‐D//NVP@C full‐cell exhibit an excellent rate capability and cycling stability.

The excellent electrochemical performance of MFSH‐HoMS‐D should be attributed to the fast Na⁺ diffusion kinetics and electron transfer rate due to its unique microstructure and alternating APBs heterojunction. The reason for the enhanced electrochemical performance of MoS₂/Fe_xMo₂S₄ heterojunction was clarified by the density functional theoretical (DFT) calculations. The models of MoS₂, Fe_xMo₂S₄ and MoS₂/Fe_xMo₂S₄ APBs heterojunction were constructed by the Vienna Ab‐initio Simulation Package (VASP) and the charge density difference at the interfaces was calculated on the model of MoS₂/Fe_xMo₂S₄. As shown in Figure 5a–c, the results showed that there was a significant electron depletion region (blue area) on one side of the heterojunction interface (Fe_xMo₂S₄ side), and a significant electron accumulation region (red area) on the other side of the heterojunction interface (MoS₂ side). This asymmetric and directional charge redistribution directly indicates that when the heterojunction forms, electrons spontaneously flow from the lower work function side to the higher work function side until the Fermi level reaches equilibrium. This process is the source and direct evidence of the built‐in electric field. The direction of charge accumulation and depletion clearly indicates the direction of the built‐in electric field (from the blue depletion region to the red accumulation region). The density of state (DOS) of MoS₂, Fe_xMo₂S₄ and MoS₂/Fe_xMo₂S₄ APBs heterojunction was calculated as shown in Figure 5d. The energy band gaps of MoS₂ and Fe_xMo₂S₄ were 0.475 and 0.042 V, while the DOS of MoS₂/Fe_xMo₂S₄ APBs heterojunction exhibits the crossing Fermi level, implying a relatively large effective electron transfer number near the Fermi level [56]. Furthermore, as shown in Figure 5e,f, due to the different energy levels of Mo‐d, Fe‐d and S‐p in Fe_xMo₂S₄ and MoS₂, it is further proved that the charge could transfer from Fe atom in Fe_xMo₂S₄ to Mo atom in MoS₂ at the interface to achieve electrochemical potential balance when the heterojunction formed, resulting in an internal electric field pointing from one side to the other. The corresponding migration paths of Na⁺ in Fe_xMo₂S₄ or MoS₂ of MoS₂/Fe_xMo₂S₄ heterojunction, the interface of MoS₂/Fe_xMo₂S₄ heterojunction, and the single‐phase MoS₂ and Fe_xMo₂S₄ were depicted in Figure 5g–i and S13, for detecting vivid ionic migration behaviors. Based on the above models, the Na⁺ migration energy barriers were calculated. As shown in Figure 5j, the heterojunction interface had the lowest Na⁺ migration energy barrier, and the migration energy barriers through Fe_xMo₂S₄ or MoS₂ of MoS₂/Fe_xMo₂S₄ heterojunction were lower than those of the single‐phase Fe_xMo₂S₄ or MoS₂. It proved that due to the formation of APB heterojunction inside the MFSH‐HoMS‐D, high‐rate ion transport channels are increased, thereby shortening the Na⁺ diffusion path, and the diffusion energy barrier is reduced due to the dilated lattice space. Additionally, the adsorption energy of Na⁺ for different models was calculated. As shown in Figure 5k, the interface in heterojunction has the highest adsorption energy due to the irregular arrangement of atoms resulting from a large number of dangling bonds, dislocations, and defects. And the adsorption energies in Fe_xMo₂S₄ or MoS₂ of MoS₂/Fe_xMo₂S₄ heterojunction are stronger than the single‐phase Fe_xMo₂S₄ or MoS₂, resulting in a fast sodiation process. Therefore, the built‐in electric field can be constructed at the interface of MoS₂/Fe_xMo₂S₄ APBs heterojunction, which effectively guides lattice distortion for providing excess channels of Na⁺ diffusion and redistribution of charge for fast electron transfer, reducing the energy barrier for ion diffusion and adding a “booster” to the transport of Na⁺. It was proved by the slight change of peak current with the increase of scan rate (Figure S14), the higher capacity contribution (Figure S15) originated from the larger specific surface area (Figure S16), the smaller charge transfer impedance, and the largest Na⁺ diffusion coefficient based on both Nyquist plots (Figure S17) and Galvanostatic Intermittent Titration Technique GITT test (Figure S18), respectively. Combining the electron transfer fleetly at the interface of MoS₂/Fe_xMo₂S₄ heterojunction with the fast Na⁺ diffusion kinetics, the electrochemical reactions between MoS₂/Fe_xMo₂S₄ and Na were improved.

FIGURE 5.

The density functional theory (DFT) calculations of MFSH‐HoMS‐D, MoS₂, and Fe_xMo₂S₄. (a) Average charge density and (b,c) charge density difference for MFSH‐HoMS‐D. (d) The DOS of MFSH‐HoMS‐D, MoS₂, and Fe_xMo₂S₄ samples. The PDOS of (e) MFSH‐HoMS‐D, (f) MoS₂, and Fe_xMo₂S₄ samples. Na⁺ diffusion along with the minimum energy path in (g) MFSH‐HoMS‐D, (h) MoS₂, and (i) Fe_xMo₂S₄. (j) Calculated energy barrier for Na⁺ diffuse from Fe_xMo₂S₄ to the heterojunction (heter‐Fe_xMo₂S₄), cross the heterojunction (heter–inter), and diffuse from the heterojunction to MoS₂ (heter‐MoS₂), and diffuse inside Fe_xMo₂S₄ and MoS₂. (k) Adsorption energy of MoS₂, Fe_xMo₂S₄, heter–inter, heter‐MoS₂, and heter‐Fe_xMo₂S₄.

The storage mechanism of Na⁺ for fast electrochemical reaction in MFSH‐HoMS‐D was explored by the in situ XRD measurements. As shown in Figure 6a, the MoS₂ and Fe_xMo₂S₄ were detected at the initial stage. During the discharging process, the additional diffraction peaks appearing at around 12.9° and 30.1° were attributed to the characteristic peaks of Na_xMoS₂ and Na_xFe_xMo₂S₄ originating from the insertion of Na⁺ into MoS₂ and Fe_xMo₂S₄, of which the peak intensity increased, accompanied by a decrease in that of MoS₂ and Fe_xMo₂S₄. Especially, the complicated evolution of diffraction peaks around 12.9° compared to the previous reported MoS₂ anode confirmed the coexistence of sodiation in Fe_xMo₂S₄ (As shown by the upward yellow line between 10° and 15° in Figure 6a). At the end of the discharging process, the Na₂S and Mo appeared [57], corresponding to the conversion reaction from Na_xMoS₂ or Na_xFe_xMo₂S₄ to Mo and Na₂S. During charging process, the peaks corresponding to Na₂S and Mo disappeared gradually, and the Na_xMoS₂ or Na_xFe_xMo₂S₄ appeared when the voltage was higher than 1.2 V. At the end of charging process, the MoS₂ and Fe_xMo₂S₄ phases appeared gradually with the decrease of peaks intensity corresponding to Na_xMoS₂ and Na_xFe_xMo₂S₄, meaning the reversible conversion reaction from Na₂S and Mo to MoS₂ or from Na₂S, Mo and Fe to Fe_xMo₂S₄ due to its unique APBs heterojunction for fast Na⁺ diffusion kinetic and electron transfer rate. According to the above results, we deduced the reversible electrochemical reactions of MFSH‐HoMS‐D as follows equations:

$$\mathrm{F}{\mathrm{e}}_{\mathrm{x}}\mathrm{M}{\mathrm{o}}_2{\mathrm{S}}_4+\mathrm{xN}{\mathrm{a}}^{+}+\mathrm{x}{\mathrm{e}}^{-}\leftrightarrow \mathrm{N}{\mathrm{a}}_{\mathrm{x}}\mathrm{F}{\mathrm{e}}_{\mathrm{x}}\mathrm{M}{\mathrm{o}}_2{\mathrm{S}}_4$$ (1)

$$\mathrm{N}{\mathrm{a}}_{\mathrm{x}}\mathrm{F}{\mathrm{e}}_{\mathrm{x}}\mathrm{M}{\mathrm{o}}_2{\mathrm{S}}_4+\left(8-\mathrm{x}\right)\mathrm{N}{\mathrm{a}}^{+}+\left(8-\mathrm{x}\right){\mathrm{e}}^{-}\leftrightarrow 4\mathrm{N}{\mathrm{a}}_2\mathrm{S}+2\mathrm{Mo}+\mathrm{xFe}$$ (2)

$$\mathrm{Mo}{\mathrm{S}}_2+\mathrm{xN}{\mathrm{a}}^{+}+\mathrm{x}{\mathrm{e}}^{-}\leftrightarrow \mathrm{N}{\mathrm{a}}_{\mathrm{x}}\mathrm{Mo}{\mathrm{S}}_2$$ (3)

$$\mathrm{N}{\mathrm{a}}_{\mathrm{x}}\mathrm{Mo}{\mathrm{S}}_2+\left(4-\mathrm{x}\right)\mathrm{N}{\mathrm{a}}^{+}+\left(4-\mathrm{x}\right){\mathrm{e}}^{-}\leftrightarrow 2\mathrm{N}{\mathrm{a}}_2\mathrm{S}+\mathrm{Mo}$$ (4)

FIGURE 6.

The characterization of reversible electrochemical reaction and the structure evolution during charging/discharging process of MFSH‐HoMS‐D. (a) In situ XRD analysis of MFSH‐HoMS‐D and corresponding discharge/charge curve in the first cycle. (b) In situ electrochemical impedance spectroscopy (EIS) and (c,d) in situ relaxation time distribution technique of sodium–ion batteries with MFSH‐HoMS‐D as anode. (e) Morphology evolution during discharging process characterized by in situ TEM. (f) HRTEM at the initial stage and corresponding Inverse FFT lattice images. (g) HRTEM at the final stage and corresponding Inverse FFT lattice images of MFSH‐HoMS‐D.

The reversible electrochemical reaction of MFSH‐HoMS‐D was further investigated by the in situ electrochemical impedance spectroscopy (EIS) (Figure 6b) and the corresponding distribution of relaxation times (DRT) analysis (Figure 6c,d), which were performed at various voltages during the discharging and charging process. As shown in Figure 6b,c, the semicircle at high‐frequency area representing interface charge transfer impedance (R _ct) in the Nyquist plot first increased and then decreased slowly, signifying a reversible electrochemical reaction. In Figure 6d, the EIS data were deconvoluted into six regions, labeled as τ1, τ2, τ3, τ4, τ5, and τ6 [58]. MFSH‐HoMS‐D anode demonstrated significantly reduced charge transfer resistance (R _ct) in the τ3 and τ4 timescale, indicating enhanced charge transfer kinetics. While also exhibited markedly improved interfacial Na⁺ mobility as reflected by the sodium–ion diffusion resistance in τ5 and τ6 range, collectively demonstrating the unique alternating MoS₂/Fe_xMo₂S₄ antiphase boundaries (APBs) heterojunction and hollow double‐shell structures of MFSH‐HoMS‐D achieve synergistic optimization in promoting charge transfer and ion diffusion [59]. Meanwhile, the SEI resistance, R _ct and Na⁺ diffusion impedance of MFSH‐HoMS‐D were the lowest compared to those of MoS₂ and MFSH‐HoMS‐1/7 (Figure S19). That is due to the unique HoMS structure and APBs heterojunction, which provides more active sites and transport channels for Na⁺. In situ TEM was applied to detect microstructure and lattice structure evolution during the discharging process. As shown in Figure 6e, the volume increased gradually (Figure 6e2–e5), while the HoMS structure was maintained well during the whole discharging process. The reason is that the HoMS structure could better buffer volume expansion, which is similar to a “breathable effect”. In situ TEM (Figure 6e1–e5) directly demonstrates that the MFSH‐HoMS‐D retains an intact double‐shell hollow morphology during deep sodiation, with only ∼18.4% volume expansion and no shell fracture. TEM after charge/discharge test cycles (Figure S12c) further confirms the preservation of the double‐shell structure. These results provide direct experimental evidence for the superior structural stability of the double‐shell hollow architecture over rod‐shaped or broken single‐shell structures. At the same time, the stress field and strong chemical bonding at the lattice mismatch interface can serve as a buffer zone and mechanical anchor point, effectively dissipating and redistributing the stress generated during the discharge process, preventing the generation of huge cracks and fractures in MFSH‐HoMS‐D, and adding another layer of “guarantee” for the material to maintain structural stability during cycling. The change of lattice structure was explored by the In situ HRTEM. During the discharge process, due to the insertion of Na⁺, the (002) lattice space of MoS₂ and (002) lattice plane of Fe_xMo₂S₄ in MFSH‐HoMS‐D increased from 0.615 and 0.648 nm to 0.592 and 0.609 nm, respectively, which further testified to the phase evolution in MFSH‐HoMS‐D and indeed the interface of APBs heterojunction can provide excess channels for Na⁺ diffusion.

3. Conclusions

In summary, the MoS₂/Fe_xMo₂S₄ HoMS with APBs heterojunction (MFSH‐HoMS‐D) was successfully fabricated by precisely controlling the Mo/Fe ratio in a modified sequential templating approach. MFSH‐HoMS‐D with a large specific surface area can expose abundant active sites for redox reaction; coupled with ultrathin shell, it effectively shortens the diffusion path of Na⁺ from electrolyte to HoMS particles. Besides, the interstitial voids between adjacent shells can mitigate volumetric expansion during charging‐discharging cycles. Furthermore, the formation of APBs heterojunction at the interface of MoS₂/Fe_xMo₂S₄ driven by the lattice mismatch between MoS₂ and Fe_xMo₂S₄, not only brings excess diffusion channels for Na⁺ to further shorten Na⁺ diffusion path, but also triggers the formation of a built‐in electric field to promote fast electron transfer. Consequently, as anode of SIBs, MFSH‐HoMS‐D achieves a high specific capacity, excellent rate capability, and impressive cycling stability. Specifically, it delivers a high specific capacity of 743.6 mAh g⁻¹ at a current density of 0.2 A g⁻¹, retaining 161.5 mAh g⁻¹ even at a super‐high current density of 20 A g⁻¹. Moreover, after 500 cycles at a current density of 5.0 A g⁻¹, almost 83.2% capacity was retained. This work could provide an inspiration for constructing hollow multi‐shelled structure anode materials with abundant lattice mismatch for fast reversible electrochemical reaction.

Author Contributions

Hanyue Wei: writing – original draft and data curation. Haiyan Liu: writing – original draft. Hui Zhang: writing – original draft and data curation. Jianjun Song: supervision, project administration, and writing – review and editing. Jiangyan Wang: supervision, writing – review and editing, and project administration. Shutao Gao: funding acquisition. Yongfu Tang: supervision, project administration, and writing – review and editing. Xiaoxian Zhao: supervision, writing – review and editing, project administration, funding acquisition, and data curation.

Conflicts of Interest

The authors declare no conflicts of interest.

Supporting information

Supporting File: The authors have cited additional references within the Supporting Information.

Data Availability Statement

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