Synergistic Cationic–Anionic Regulation in Ni-Doped FeSe@C Anodes with Se Vacancies for High-Efficiency Sodium Storage

Abstract

Sodium-ion batteries present an economical energy storage solution, yet their anode kinetics remain slow, impeding rate performance and cyclability. Layered FeSe anodes, characterized by metallic conductivity, hold potential, but structural decay and insufficient active sites during cycling continue to pose challenges. Herein, these challenges are addressed through the implementation of dual Ni doping and Se vacancy engineering in FeSe@C to synergistically regulate cationic/anionic configurations. The ionic substitution of larger Fe²⁺ ions (0.78 Å ionic radius) with smaller Ni²⁺ ions (0.69 Å) induces lattice distortion and generates abundant Se vacancies, enhancing electron transport, active site accessibility, and Na⁺ adsorption. These synergistic modifications effectively boost Na⁺ diffusion kinetics and electrolyte compatibility, creating a favorable electrochemical environment for fast sodium storage. Consequently, the optimized 2%Ni-FeSe@C electrode retains an exceptional discharge specific capacity of 307.67mAh g⁻¹ after 1000 cycles at an ultrahigh current density of 5 Ag⁻¹, showcasing superior rate capability and long-term cycling stability, paving the way for practical high-power SIBs.

1. Introduction

Lithium-ion batteries (LIBs) stand as a preeminent electrochemical energy storage technology, renowned for their exceptional energy density and robust cycle stability, which underpin their critical roles within the fields of portable electronics, electric vehicles, as well as grid-level energy storage systems [1]. Nevertheless, the escalating costs and scarcity of global lithium resources have driven a critical research shift toward SIBs, which utilize sodium—a naturally abundant and cost-effective alkali metal—as the charge carrier [2,3,4]. Although SIBs exhibit substantial potential for giga-scale energy storage deployment, their practical realization is constrained by the limitations of commercial anode materials: conventional graphite anodes, which are unsuitable for Na⁺ intercalation owing to the more extended ionic radius of Na⁺ (1.02 Å versus Li⁺ 0.76 Å), suffer from inadequate specific capacity and limited cycle durability [5,6]. While hard carbon, a prevalent alternative anode material for SIBs, alleviates certain intercalation challenges, its amorphous structure inherently leads to inferior electrical conductivity and substantial irreversible capacity loss, thus necessitating the investigation of high-performing anode materials with optimized architectures [7].

Transition metal dichalcogenides (TMDs), exemplified by iron selenide (FeSe), have emerged as prospective anode materials on account of their high theoretical specific capacity, abundant iron resource availability, and excellent electrical conductivity [8]. Nonetheless, FeSe-based anodes are plagued by intrinsic limitations, encompassing inefficient active material utilization, sluggish ion transport kinetics, and irreversible structural degradation during sodiation/desodiation, which collectively constrain their specific capacity and rate capability [9,10]. To tackle these challenges, researchers have investigated vacancy engineering and anion doping strategies to regulate electronic structures and boost electrochemical reactivity [11,12,13]. For instance, anion doping can create active sites and enhance alkali ion adsorption, and the abundant vacancies can induce lattice distortion and thus enhance reaction transport kinetics.

However, this strategy is primarily focused on cation or anion (e.g., S/Se) modifications, with scant research dedicated to dual doping and vacancy engineering [14]. Notably, the synergistic combination of cation doping and vacancy engineering—especially when involving metal cations with adjustable concentrations—remains largely unexplored in FeSe-based materials. Although cation doping is theoretically capable of inducing lattice distortion and electron redistribution to enhance ion storage kinetics, no universal synthesis method has been reported to achieve precise control over both cation doping ratios and vacancy formation simultaneously [3,15].

In this work, the research gap is bridged by constructing Ni-doped FeSe@C (2%Ni-FeSe@C) via a hydrothermal doping and selenization protocol. Given that the ionic size of Ni²⁺ (0.69 Å) is of a smaller magnitude than that of Fe²⁺ (0.78 Å), the Se vacancy concentration is deliberately enhanced and charge density redistribution is induced in FeSe-based anode material, thereby improving its electron mobility and enriching surface active sites [16,17]. The Ni-doped carbon layer additionally boosts electrical conductance and mitigates structural stress throughout repeated charge–discharge processes. Electrochemical characterizations reveal that the optimized 2%Ni-FeSe@C anode demonstrates a high degree of reversibility specific capacity of 427.65 mAh g⁻¹ at a current density of 1 A g⁻¹ after 100 cycles, and retains 307.67mAh g⁻¹ after 1000 cycles at 5 Ag⁻¹, demonstrating its exceptional rate capability and long-term cyclability.

2. Experimental Section

2.1. Synthesis of 2%Ni-FeOOH@C Nanomaterial

Initially, FeCl₃·6H₂O (0.2705 g), CO(NH₂)₂ (0.63 g), and Ni(NO₃)₂·6H₂O (0.00884 g) were disintegrated in 50 mL of ethylene glycol and underwent magnetic stirring for 30 min. The solution was afterward moved to a 100 mL stainless steel reactor and heated at a rate of 5 °C min⁻¹ from room temperature to 170 °C over a period of 10 h. Following this, the reactor was naturally reduced to room temperature. The resulting light green precipitate was washed once with distilled water and twice with anhydrous ethanol. The final product, 2%Ni-FeOOH@C, appeared as a green powder and was dried overnight in a vacuum oven at 60 °C.

2.2. Synthesis of 2%Ni-FeSe@C Nanomaterial

Thereafter, 100 milligrams of the manufactured 2%Ni-FeOOH was disintegrated in 100 mL of a buffer solution consisting of Tris (prepared by mixing 40 mL of water with ions eliminated and 60 mL of anhydrous ethanol with 121 mg of tris (hydroxymethyl) aminomethane) and sonicated for 30 min. Thereafter, 60 mg of dopamine hydrochloride was introduced, and the mixture was magnetically stirred for 24 h. Finally, 100 mg of 2%Ni-FeOOH@C powder was selenized with selenium powder at a mass ratio of 1:4 within a tubular furnace under a mixed atmosphere of argon (90%) and hydrogen (10%), heated at 560 °C for 9 h characterized by a heating rate of 2 °C min⁻¹, to attain the final product 2%Ni-FeSe@C. The overall synthesis yield was about 75.8%. The control samples were synthesized via the same routes except for adding different amounts of Ni(NO₃)₂·6H₂O. For more detailed information, materials, material characterization, and electrochemical measurements are provided in the Supporting Information.

3. Results and Discussion

Figure 1a is the flow chart for the preparation of 2%Ni-FeSe@C. Initially, 2%Ni-FeOOH@C was meticulously fabricated via the hydrothermal synthesis approach. Subsequently, 2%Ni-FeSe@C was successfully synthesized through the process of high-temperature selenization, which represents a crucial post-treatment step for material transformation and property enhancement.

Figure 1b presents the XRD patterns associated with the as-prepared samples. The diffraction peaks at 30.4°, 32.4°, 42.2°, 50.5°, and 55.3° are indexed to the (002), (101), (102), (110), and (103) crystal planes of hexagonal FeSe (H-FeSe, PDF#75-0608), respectively, confirming the successful synthesis of the hexagonal phase. Additionally, the weak diffraction peaks detected at 2θ = 28.7° and 47.4° are linked to the tetragonal phase of FeSe (T-FeSe, PDF#85-0735), indicating the coexistence of minor tetragonal structural domains. The layered architecture of the minor tetragonal phase, characterized by an expanded interlayer spacing of 0.59 nm, facilitates rapid sodium-ion intercalation/deintercalation, thereby enhancing cycle stability [18]. High-resolution analysis of the 30–45° diffraction region as shown in Figure 1c reveals that increasing Ni-doping ratios cause high-intensity (101) and (102) peaks to shift to higher angles. This peak shift is primarily attributed to lattice distortion induced by the substitution of Fe²⁺ with Ni²⁺, owing to the narrower ionic radius of Ni²⁺ (0.69 Å) compared to Fe²⁺ (0.78 Å). The ionic radius mismatch generates lattice strain during cation substitution, driving the observed diffraction peak displacement [19]. Raman spectroscopy (Figure 1d) reveals that both samples exhibit characteristic D-band (1346 cm⁻¹, amorphous carbon) and G-band (1585 cm⁻¹, graphitic carbon) signals, confirming the presence of carbon in the composites. The I_G/I_D intensity ratio of 2%Ni-FeSe@C increases to 1.14, compared with 0.86 for FeSe@C, reflecting a higher degree of structural disorder in the carbon layer. Such an increase ratio suggests reduced graphitization quality, which introduces more defects to facilitate electron transport and ion diffusion [20].

To characterize the microstructure and morphology of FeSe@C and 2%Ni-FeSe@C samples, SEM analysis was performed as exhibited in Figure 2a–d. The SEM images reveal that both samples exhibit spherical particles with diameters of approximately 3 μm, and distinct lamellar stacking structures are observed on their surfaces. Compared with FeSe@C, 2%Ni-FeSe@C particles exhibit a more uniform size distribution and regular morphology. The distinct structural features not only facilitate Na⁺ transport and enhance electrolyte penetration but also further reduce interfacial charge transfer resistance [21,22], thereby improving the material’s overall electrochemical performance [23,24]. HRTEM analysis (Figure 2e–h) indicates that the (102) crystal plane of 2%Ni-FeSe@C demonstrates an interplanar spacing of 0.214 nm, and the entire structure is consistently sheathed by a layer containing carbon. This uniform coating composed of carbon effectively mitigates volume expansion that occurs during charge–discharge cycles, providing structural durability for long-term electrochemical performance [21]. Elemental mapping analysis (Figure 2i,j) manifests that Fe, Ni, Se, N, and C are homogeneously distributed throughout the particles of 2%Ni-FeSe@C. This observation indicates addition of N to the carbon layer through the doping process, which alters the carbon structure and facilitates electrical conductivity. The homogeneous distribution of Ni further evidences the uniform substitution of Fe sites in FeSe by Ni²⁺, consistent with the XRD peak shift discussed earlier. Brunauer–Emmett–Teller (BET) measurements (Figure S1) demonstrate that 2%Ni-FeSe@C exhibits a high specific surface coverage of 118.3 m² g⁻¹, significantly larger than the 85.8 m² g⁻¹ observed for FeSe@C. This enhanced BET surface area increases the count of active sites situated on 2%Ni-FeSe@C, thereby boosting its electrochemical activity. The EPR characterization (Figure S4) confirms the presence of selenium vacancies in both 2%Ni-FeSe@C and FeSe@C, as evidenced by the distinct resonance signals at g = 2.003, which aligns with the well-documented signature of Se vacancies in metal selenides.

XPS was utilized to characterize the composition and chemical nature of the surface bonding states of pristine FeSe@C and 2%Ni-FeSe@C samples. The XPS profiles of FeSe@C (Figure 3a) clearly exhibit characteristic peaks for Fe, Se, N, C, and Ni. This demonstrates the successful insertion of N and Ni into the lattice during the doping process. In the Se 3d high-resolution XPS spectrum (Figure 3b), the peaks at 53.9 eV and 55.8 eV are assigned to Se 3d_5/2 and Se 3d_3/2 spin-orbit split states, respectively. A weak peak at 59 eV is ascribed to the Se-O bond, originating from the surface oxidation of selenium species [25]. The shift in Se 3d peak positions is attributed to Ni doping and Se vacancies indirectly modulating the electron cloud distribution around Se atoms [26]. The Fe 2p high-resolution XPS spectrum (Figure 3c) is deconvoluted into six subpeaks, with characteristic peaks at 710.8 eV and 724.2 eV assigned to Fe 2p_3/2 and Fe 2p_1/2 spin-orbit split states, respectively, exhibiting an energy separation (ΔE_Fe) of 13.4 eV. These peaks confirm the presence of Fe²⁺ in both FeSe@C and 2%Ni-FeSe@C samples. Additional peaks at 713.6 eV and 725.3 eV are attributed to Fe³⁺, indicating the formation of a higher iron oxidation state [27]. Notably, Ni doping induces a significant negative shift in the binding energies of both Fe²⁺ and Fe³⁺ peaks. Additionally, a peak at 706 eV, assigned to metallic Fe (Fe⁰), arises from the partial reduction during the annealing process [28]. In the C 1s high-resolution XPS spectrum (Figure 3d), three characteristic peaks at 285.0 eV, 286.7 eV, and 288.0 eV are assigned to C–C, C=O, and O–C=O groups, respectively. In the N 1s high-resolution XPS spectrum (Figure 3e), three minor peaks at 398.5 eV, 400.7 eV, and 403.9 eV are consistent with pyridinic-N, pyrrolic-N, and graphitic-N, respectively [29]. Pyrrolic-N effectively boosts electrochemical reaction activity and electrical conductivity, with such enhancement primarily attributed to N-doping-induced structural defects on the carbon surface, and Ni²⁺ substitution induces Se vacancies (Figure 3b and Figure S4), while pyrrolic-N optimizes the carbon layer’s electronic environment, collectively promoting surface-dominated pseudocapacitive Na⁺ storage (90.3% contribution at 1 mV s⁻¹) [30]. The high-resolution XPS spectrum of Ni 2p for 2%Ni-FeSe@C (Figure 3f) shows clearly discernible spin-orbit split peaks of Ni 2p_1/2 and Ni 2p_3/2 at 873.0 eV and 875.0 eV, respectively, indicating the successful incorporation of Ni into the FeSe lattice.

To elucidate the reaction mechanism of 2%Ni-FeSe@C, the cyclic voltammetry (CV) measurement was conducted as shown in Figure 4a. The first anodic scan exhibits four distinct reduction peaks at ~1.06 V, ~0.78 V, ~0.20 V, and 0.08V, respectively. The sharp peak at ~1.06 V is attributed to the formation of Na_xFeSe and a solid electrolyte interphase (SEI) via electrolyte decomposition [31]. A second cathodic peak at 0.78 V was assigned to the conversion reaction of Na_xFeSe to Fe and Na₂Se (Na_xFeSe + (2 − x) Na⁺ + (2 − x) e⁻ → Fe + Na₂Se) [32,33,34]. The subsequent charge/discharge CV curves tend to be stable, indicating the 2%Ni-FeSe@C exhibits good stability toward sodiation/desodiation [18]. In addition, there is a pair redox peaks closing to 0 V, which should be attributed to sodium-ion intercalation/deintercalation into carbon layers [21]. During the cathodic scan, oxidation peaks at ~1.52 V, ~1.80 V, and ~2.05 V are assigned to the conversion reactions of Se species in Na_xFeSe and FeSe [32,33,34,35]. Starting from the second cycle, the anodic scan exhibits a pronounced reduction peak near 0.78 V, with subsequent CV curves stabilizing, signifying the excellent stability of 2%Ni-FeSe@C during sodium insertion/extraction cycles. In pristine FeSe@C (Figure S3), the reduction peak at 1.3 V corresponds to irreversible phase decomposition (FeSe + 2Na⁺ + 2e⁻ → Fe + Na₂Se), which is fully suppressed in 2%Ni-FeSe@C due to Ni²⁺-induced lattice stabilization and strengthened Fe–Se bonding. Concurrently, the emergence of a 1.80 V oxidation peak exclusively in the doped material aligns with the Ni²⁺/Ni³⁺ redox couple activation. This Ni-mediated process not only replaces the detrimental decomposition pathway but also facilitates reversible Na⁺ storage through enhanced interfacial charge transfer [35].

The overlapping CV curves indicate the excellent reversibility, as corroborated by the galvanostatic charge–discharge (GCD) curves (Figure 4b). The active electrode was fabricated with a geometric area of 1.13 cm² and a total mass loading of about 1 mg cm⁻². At a current density of 0.1 A g⁻¹, the electrode exhibits initial discharge and charge capacities of 565.4 and 510.7 mAh g⁻¹, respectively, yielding an initial Coulombic efficiency (ICE) of 90.32%. The relatively low ICE is attributed to instability of the solid electrolyte interphase (SEI) and irreversible phase transformations [33,36,37]. The initial discharge and charge capacities of FeSe@C are measured at 551.7 and 475.1 mAh·g⁻¹, respectively, with an initial Coulombic efficiency (ICE) of 86.12% (Figure S2), markedly inferior to those of 2%Ni-FeSe@C. This disparity underscores the efficacy of Ni doping in boosting initial Coulombic efficiency. Furthermore, rate capability assessments conducted over current densities spanning 0.1 to 5 A g⁻¹ (Figure 4c) reveal the exceptional capacity retention of 2%Ni-FeSe@C. It sustains capacities of 521.2, 515.7, 503.9, 474.6, 413.1, and 355.8 mAh g⁻¹ at 0.1, 0.2, 0.5, 1, 2 and 5 Ag⁻¹, respectively, significantly outperforming other samples.

In terms of cycling stability (Figure 4d), 2%Ni-FeSe@C retains a discharge capacity of 427.65 mAh g⁻¹ after 100 cycles at 1 Ag⁻¹, surpassing FeSe@C (340.3 mAh g⁻¹), 1%Ni-FeSe@C (376.3 mAh g⁻¹), and 3%Ni-FeSe@C (366.6 mAh·g⁻¹), respectively. Even at an extremely high current density of 5 A g⁻¹ as exhibited in Figure 4e, 2%Ni-FeSe@C maintains an overall reversible capacity of 307.67 mAh g⁻¹ after 1000 cycles, whereas FeSe@C delivers only 244.1 mAh g⁻¹. This superior cycle performance is attributed to Ni doping enhancing the structural stability of FeSe. A comparison of reported SIB anodes (Figure 4f) confirms that 2%Ni-FeSe@C exhibits excellent cycling stability, which is attributed to its unique nanosheet structure, cation doping, and Se vacancies enabling fast ion/electron transport [8,32,34,38,39,40].

Figure 4. (a) CV curves of 2%Ni-FeSe@C at 0.1 mV s⁻¹; (b) charge–discharge curves of 2%Ni-FeSe@C for the first three cycles; (c) rate behavior and (d) cycling behavior of the four electrodes; (e) cycling behavior at 5 Ag⁻¹ for the FeSe@C and 2%Ni-FeSe@C electrodes; (f) contrast of the electrochemical behavior with the reported selenides anode materials [8,19,21,32,34,37,38,39].

Figure 4. (a) CV curves of 2%Ni-FeSe@C at 0.1 mV s⁻¹; (b) charge–discharge curves of 2%Ni-FeSe@C for the first three cycles; (c) rate behavior and (d) cycling behavior of the four electrodes; (e) cycling behavior at 5 Ag⁻¹ for the FeSe@C and 2%Ni-FeSe@C electrodes; (f) contrast of the electrochemical behavior with the reported selenides anode materials [8,19,21,32,34,37,38,39].

CV curves with varying scan rates were applied to evaluate the electrochemical kinetics and reversibility of FeSe@C and 2%Ni-FeSe@C, as shown in Figure 5a,b. Peak 1 and peak 2 are assigned to the conversion reactions of Se species in Na_xFeSe and FeSe [32,33,34]. Meanwhile, peak 3 is assigned to the conversion of Na_xFeSe into Fe and Na₂Se [18]. The images reveal that the curves maintain their shape with the increasing scan rates and exhibit negligible peak shifts, suggesting that the 2%Ni-FeSe@C electrode undergoes mild polarization and demonstrates excellent electrochemical reversibility [41,42,43].

The diffusion coefficient of Na⁺ can be estimated using Equation (1):

$${\mathrm{I}}_{\mathrm{P}}=2.69 \times {10}^5{\mathrm{n}}^{2/3}\mathrm{AC}{\mathrm{v}}^{1/2}{\mathrm{D}}_{{\mathrm{Na}}^{+}}^{1/2}$$

(1)

where I_p is the peak current (A) at different scan rates, n is the charge transfer number, A is the electrode surface area, C is the molar concentration of Na⁺ (mol cm⁻³), and v is the scan rate. Based on the fitting results of Figure 5c,d, the diffusion coefficients of 2%Ni-FeSe@C during the redox processes of Peaks 1, 2, and 3 are calculated to be 1.34 × 10⁻⁹, 1.86 × 10⁻⁹, and 2.19 × 10⁻⁹ cm² s⁻¹, respectively. In contrast, the Na⁺ diffusion coefficients of FeSe@C are only 1.26 × 10⁻⁹, 1.75 × 10⁻⁹, and 2.16 × 10⁻⁹ cm² s⁻¹, respectively.

Based on the relationship between the peak current (i) and the scan rate (v), the contributions of diffusion control and surface pseudocapacitive control to the Na⁺ storage performance of FeSe@C and 2%Ni-FeSe@C can be quantitatively analyzed as Equations (2)–(4):

$$\mathrm{i}=\mathrm{a}{\mathrm{v}}^{\mathrm{b}}$$

(2)

$$\log \mathrm{i}=\mathrm{b}\log \mathrm{v}+\log \mathrm{a}$$

(3)

$$\mathrm{I}={\mathrm{k}}_1\mathrm{v}+{\mathrm{k}}_2{\mathrm{v}}^{1/2}$$

(4)

where k₁v represents the current contribution from pseudocapacitive effects, and k₂v^1/2 corresponds to the current contribution from diffusion-controlled processes, with k₁ and k₂ being adjustable parameters. By fitting the values of k₁ and k₂, the total current response (I) at a specific voltage can be decomposed into two components: pseudocapacitive effects (k₁v) and diffusion effects (k₂v^1/2), thereby quantifying the proportion of pseudocapacitive contribution during Na⁺ storage.

Figure 6a–d show that the pseudocapacitive contributions of 2%Ni-FeSe@C are 71.6%, 73.2%, 76.9%, 81.3%, 85.1%, and 90.3% at scan rates of 0.1, 0.2, 0.4, 0.6, 0.8, and 1 mV s⁻¹, respectively, significantly higher than that of FeSe@C (66.9%, 72.3%, 76.3%, 80.2%, 84.4%, and 87.8%). This indicates that 2%Ni-FeSe@C exhibits superior fast charge–discharge kinetics. Quantitative analysis of the relative contributions from pseudocapacitive and diffusion-controlled mechanisms further confirms that 2%Ni-FeSe@C more effectively utilizes surface pseudocapacitive processes during fast charge–discharge, thereby significantly enhancing its rate performance.

Systematic studies of interfacial reaction resistance were conducted using EIS. As shown in Figure 7a, the Nyquist plots of both samples exhibit typical characteristics: semicircular arcs appear in the intermediate-frequency region, pertaining to charge transfer resistance (R_ct) [21]; and sloped lines appear in the low-frequency region, corresponding to Warburg impedance (W_s). Among them, the charge transfer resistance (R_ct) reflects the charge transfer process at the electrolyte/electrode interface, while the Warburg impedance (W_s) represents the diffusion behavior of Na⁺ within the electrode material—the slope of which is closely related to the diffusion coefficient of Na⁺ in the material.

The diffusion coefficient of Na⁺ (D_Na⁺) and the exchange current density (I₀) can be calculated using Equations (5) and (6) below [44].

$${\mathrm{D}}_{{\mathrm{Na}}^{+}}={\mathrm{R}}^2{\mathrm{T}}^2/2\mathrm{A}{\mathrm{n}}^4{\mathrm{F}}^4{\mathrm{c}}^2{\mathsf{\sigma }}^2$$

(5)

$${\mathrm{I}}_{0 }=\mathrm{RT}/\mathrm{nF}{\mathrm{R}}_{\mathrm{ct}}$$

(6)

Among them, R, T, A, n, F, c, and σ denote the gas constant (8.314 J K⁻¹ mol⁻¹), temperature, the area of the disk electrode, the number of electrons, the Faraday constant, the molar concentration of Na⁺, and the Warburg coefficient, respectively. According to these formulas, the value of σ can be determined by the slope of the linear relationship between Z′ and ω^−1/2 in the low-frequency region (Figure 7b). The corresponding electrochemical impedance spectroscopy (EIS) data fitted using the equivalent circuit are summarized in

Table 1. Electrochemical impedance parameters.

Samples	Rs (Ω)	Rct (Ω)	σ (Ω S−0.5)	D (cm2 S−1)	I0 (mA cm−2)
FeSe@C	3.7	117.5	244.5	5.93 × 10−9	22.18 × 10−4
2%Ni-FeSe@C	2.9	101.0	106.5	3.13 × 10−8	2.54 × 10−4

. The results show that 2%Ni-FeSe@C exhibits the lowest charge transfer resistance (101 Ω), the lowest σ value (106.5 Ω S^−1/2), and the largest Na⁺ diffusion coefficient (D_Na⁺, 3.13 × 10⁻⁸ cm² S⁻¹). These results indicate that 2%Ni-FeSe@C demonstrates a superior performance in terms of charge transfer and Na⁺ diffusion kinetics. This can be attributed to the fact that nickel doping leads to the rearrangement of charge density and additional lattice distortion in FeSe.

To clarify the kinetic differences between 2%Ni-FeSe@C and FeSe@C—particularly regarding their distinct rate capabilities—the galvanostatic intermittent titration technique (GITT) was employed. A 0.1 A g⁻¹ pulse current was applied for 30 min, followed by 60 min rest intervals, to determine the apparent Na⁺ diffusion coefficient within the electrodes of both materials.

Figure 8a shows the characteristic potential plateaus corresponding to Na⁺ desodiation and sodiation processes. Figure 8b indicates that the potential change (ΔV) in FeSe@C during the relaxation phase is 0.26 V, significantly higher than that of 2%Ni-FeSe@C (0.20 V). This can be attributed to the introduction of Se vacancies by Ni doping, which serve as fast diffusion pathways for Na⁺ and shorten the ionic migration distance. Additionally, Ni may inhibit the agglomeration, thereby reducing the contribution of diffusion impedance to ΔV.

Based on Fick’s second law of diffusion, the D_Na⁺ can be estimated using Equation (7) [45,46,47].

$${\mathrm{D}}_{{\mathrm{Na}}^{+}}={\displaystyle \frac{4}{\mathsf{\Pi }\mathsf{\tau }}}{\left({\displaystyle \frac{{\mathrm{m}}_{\mathrm{B}}{\mathrm{V}}_{\mathrm{M}}}{{\mathrm{M}}_{\mathrm{B}}\mathrm{A}}}\right)}^2{\left({\displaystyle \frac{\mathsf{\Delta }{\mathrm{E}}_{\mathrm{s}}}{\mathsf{\Delta }{\mathrm{E}}_{\mathsf{\tau }}}}\right)}^2$$

(7)

where M_B is the molar mass, V_M denotes the molar volume, m_B represents the sample mass, and A stands for the surface area of the electrode. The ΔE_S and ΔE_t values can be derived from the GITT curves. Figure 8c,d show that the Na⁺ diffusion coefficient curves of the two materials exhibit a marked similar trend. The D_Na⁺ of 2%Ni-FeSe@C is significantly higher than that in FeSe@C.

4. Conclusions

This work introduces a unique scheme for enhancing SIBs’ anode behavior through cationic Ni doping and Se vacancy engineering in FeSe@C composites. Through the substitution of smaller Ni²⁺ for Fe²⁺, lattice distortion and Se vacancies are generated, which synergistically enhance electron mobility, boost active site density, and augment Na⁺ adsorption. The optimized 2%Ni-FeSe@C electrode achieves a reversible discharge capacity of 427.65 mAh g⁻¹ over 100 cycles at 1A g⁻¹ and 307.67 mAh g⁻¹ over 1000 cycles at 5 A g⁻¹, outperforming most reported selenide anodes. Mechanistic studies reveal that 90.3% of charge storage arises from pseudocapacitive processes, attributed to surface-controlled reactions enabled by Se vacancies and Ni doping. This strategy underscores the promise of dual engineering—involving cation substitution and vacancy defects—in transition metal dichalcogenides for ultrafast alkali-metal storage, providing a scalable avenue to high-performance SIBs suited for giga-scale applications.