Scaly MoS₂/rGO Composite as an Anode Material for High-Performance Potassium-Ion Battery

Abstract

Potassium-ion batteries (PIBs) have been widely studied owing to the abundant reserves, widespread distribution, and easy extraction of potassium (K) resources. Molybdenum disulfide (MoS₂) has received a great deal of attention as a key anode material for PIBs owing to its two-dimensional diffusion channels for K⁺ ions. However, due to its poor electronic conductivity and the huge influence of embedded K⁺ ions (with a large ionic radius of 3.6 Å) on MoS₂ layer, MoS₂ anodes exhibit a poor rate performance and easily collapsed structure. To address these issues, the common strategies are enlarging the interlayer spacing to reduce the mechanical strain and increasing the electronic conductivity by adding conductive agents. However, simultaneous implementation of the above strategies by simple methods is currently still a challenge. Herein, MoS₂ anodes on reduced graphene oxide (MoS₂/rGO) composite were prepared using one-step hydrothermal methods. Owing to the presence of rGO in the synthesis process, MoS₂ possesses a unique scaled structure with large layer spacing, and the intrinsic conductivity of MoS₂ is proved. As a result, MoS₂/rGO composite anodes exhibit a larger rate performance and better cycle stability than that of anodes based on pure MoS₂, and the direct mixtures of MoS₂ and graphene oxide (MoS₂-GO). This work suggests that the composite material of MoS₂/rGO has infinite possibilities as a high-quality anode material for PIBs.

1. Introduction

Energy remains a crucial aspect in the current state of global development. Developing sustainable and clean energy is an inevitable trend, particularly given the massive consumption of traditional fossil fuels and the deteriorating global environment. Electrochemical energy storage (EES) systems are becoming increasingly popular in the field of energy technology [1,2,3]. Lithium-ion batteries (LIBs) are extensively deployed in a wide range of consumer electronics, electric vehicles, and national grids due to their high energy density and long cycle life in various EES systems. However, lithium is not an abundant resource in the Earth’s reserves, accounting for only 0.0017 wt%. Furthermore, its uneven distribution leads to high extraction and transportation costs, making it a poorly suited sustainable option for meeting the world’s long-term energy needs [4,5].

In recent years, research on batteries using Na, Mg, Al, and K has made considerable progress and received increasing attention [6,7,8,9]. Among them, potassium-ion batteries (PIBs) have been widely studied owing to the abundant K element in nature (2.09 wt%), high energy density (200–300 Wh kg⁻¹), and favorable K⁺ ions mobility in nonaqueous electrolytes (resulting from the smaller Stokes radius of K⁺ compared to Li⁺ or Na⁺) [10,11,12]. However, due to the inherent large ionic radius of K⁺ ion (3.6 Å), it encounters high diffusion barriers, and easily cause the material structure to collapse during insertion/extraction. Therefore, it is crucial to find electrode materials that can allow speed kinetics and maintain structural stability during K⁺ de-intercalation [13,14,15].

Noncarbon-based intercalation materials, in particular molybdenum disulfide (MoS₂), as typical transition metal disulfides (TMDs), have received a great deal of attention as a key anode material for PIBs [16,17]. MoS₂ is an economical and readily available mineral with a layered structure, which enables it to open up two-dimensional diffusion channels during K⁺ intercalation [18,19,20]. However, the rate performances and electrochemical stability of MoS₂ is still low due to its poor electronic conductivity (σ_RT = ~10⁻⁴ s cm⁻¹) and the huge influence of embedded K⁺ ions on MoS₂ layer [21]. Many research studies have been carried out to solve these problems. The common strategies are increasing the electrical conductivity by adding conductive agents and enlarging the interlayer spacing to reduce the mechanical strain [22,23,24]. Graphene oxide (GO) is an electrically conductive layered graphite oxide whose surface is rich in a variety of functional groups, which can provide an abundance of effective active sites [25]. However, an excess of functional groups can cause it to become too active to maintain a stable structure. Reduced graphene oxide (rGO), as a product of GO reduction, has higher electrical conductivity and structural stability while retaining functional groups. In addition, it was found that rGO can be used as a substrate material to modulate the microstructure of other materials [26]. However, simultaneous implementation of the above strategies by simple methods is currently still a challenge.

In this work, we obtained MoS₂ on reduced graphene oxide (MoS₂/rGO) composite by the one-step hydrothermal method. According to our experiments and calculations, MoS₂ in such composite has an ultrathin scale-like structure with enlarged interlayer spacing, which indirectly increases the specific surface area, provides more active sites for K⁺, and ensures structural stability. In addition, the incorporation of rGO can increase the intrinsic conductivity of MoS₂. As a result, MoS₂/rGO composite anodes exhibited a better rate performance (287.15, 266.7, 220.3, and 184.6 mAh g⁻¹, at 50, 100, 500, and 1000 mA g⁻¹, respectively) and cycle stability (99% capacity retention after 100 cycles at 500 mA g⁻¹) than that of anodes based on pure MoS₂, and the direct mixtures of MoS₂ and graphene oxide (MoS₂-GO). This work demonstrates the possibility of using MoS₂/rGO composite as anodes for PIBs.

2. Results and Discussion

The preparation of scaly MoS₂/rGO composite by a one-step reaction is illustrated in Figure 1. First, a typical modification method for GO is used [27]. The ammonium ions of cetyltrimethylammonium bromide (CTAB), which are positively charged, can attract the carboxyl group of GO, which has a negative charge, through Coulombic forces. The isolation of the lamellar GO sheets is a consequence of CTAB insertion. Next, a composite of MoS₂/rGO is synthesized in a solvothermal reaction of Na₂MoO₄·2H₂O and CH₄N₂S at 200 °C with the presence of rGO. There are several oxygen-containing groups in rGO, which are the adsorption sites for Mo⁴⁺ and K⁺, and also the nucleation sites for nanoparticles. As a result, uniform MoS₂ precursors develop on the surface of rGO. The Mo-O-C bonds located at the remaining oxygen sites can interact with van der Waals forces present among the partial rGO, ultimately leading to the binding of the ions and partial rGO. The MoS₂/rGO precursor shall undergo annealing treatment to yield MoS₂/rGO composite.

Scanning electron microscopy (SEM) was employed to examine the morphology of MoS₂ and MoS₂/rGO composite. As illustrated in Figure 2a,b, some of the MoS₂ nanosheets exhibit a tendency to agglomerate. With regard to the MoS₂/rGO composite, it can be observed that the scaly MoS₂ nanosheets are uniformly distributed on the MoS₂/rGO surface and are thinner and smaller in size. This result indicates that the presence of rGO effectively inhibits the agglomeration of MoS₂ nanosheets. The scaly MoS₂ nanosheets can enhance the specific surface area of MoS₂/rGO and facilitate the generation of more active sites for K⁺ [28]. The microstructure of MoS₂/rGO composite was investigated by field emission transmission electron microscopy (TEM). Figure 2c shows a magnified TEM image of MoS₂/rGO at 200 nm scale, which clearly demonstrates MoS₂ nanosheets embedded in translucent rGO sheets. Notably, the molybdenum disulfide nanosheets exhibit multilayered stacked lattice stripes, indicating a lattice spacing of 6.39 Å. This spacing is larger than the crystalline hexagonal 2H-MoS₂ lattice spacing (002) of 6.1554 Å, as shown in Figure 2d [29]. To further support the observations in Figure 2c, Figure 2e shows scanning transmission electron microscopy (STEM) images of MoS₂/rGO, providing additional evidence that the appearance of MoS₂ scales is a result of MoS₂ wrapped around rGO. Figure 2f–h show the total and partial elemental maps of MoS₂/rGO captured at 200 nm scale. These maps reveal the uniform distribution of carbon (C), molybdenum (Mo), and sulfur (S) in the composite. This uniform distribution further confirms the successful preparation of MoS₂/rGO composite and supports the conjecture that MoS₂ grows on rGO.

The crystal structure of the synthesized MoS₂/rGO sample was analyzed using X-ray powder diffraction (XRD) patterns. Figure 3a presents the XRD diffractograms of MoS₂/rGO and MoS₂. The four major diffraction peaks, observed at approximately 14.3°, 33.5°, 39.5°, and 58.3°, respectively, correspond to the (002), (101), (103), and (110) crystallographic planes of 2H-MoS₂ (PDF#37-1492) [30]. The disappearance of the principal diffraction peak of GO at approximately 12° and the emergence of the principal diffraction peak of rGO at approximately 26° indicate that GO was successfully reduced to rGO [31]. Notably, the peak of MoS₂/rGO at the (002) crystal plane is shifted to the left, indicating an enlarged layer spacing. The expanded layer spacing facilitates potassium insertion and extraction during the cycling process, thereby enhancing the potassium storage capacity of MoS₂/rGO.

X-ray photoelectron spectroscopy (XPS) was utilized to determine the elemental composition and chemical bonding in MoS₂/rGO. The spectrum of C 1s can be divided into four peaks, 284.7, 285.2, 268, and 286.6 (Figure 3b), which correspond to the sp2-banded of C-C, sp3-banded of C-C, C-S, and C-O, respectively [26]. As illustrated in Figure 3c, the peak at 226 in the Mo 3d spectrum is attributed to S 2s [32], while the pair of characteristic peaks at 229.7 and 232.8 are attributed to Mo 3d_5/2 and Mo 3d_3/2, respectively. This indicates that the valence state of Mo is +4 [33]. The pair of characteristic peaks at 232.2 and 236 should be attributed to Mo 3d_5/2 and Mo 3d_3/2 for Mo⁶⁺, which may be attributed to the oxidation of the MoS₂ surface, resulting in the change to Mo⁶⁺ [34]. Figure 3d depicts the spectrum of S 2p. The pair of characteristic peaks at 162.5 and 163.7 eV can be attributed to the S 2p_3/2 and S 2p_1/2 of S²⁻. The small peak at 168.3 eV is indicative of the S-O-C bond [35]. A comparison of the XPS of MoS₂ and MoS₂-GO in Figure S1 reveals that pure MoS₂ and MoS₂-GO lack the C-S bond observed in the C 1s and S 2p spectra. This suggests that rGO and MoS₂ can be more effectively bonded together in the MoS₂/rGO material, leading to enhanced stability. Furthermore, the thermogravimetric analysis (TGA), as shown in Figure S2, was used to evaluate the MoS₂ content of MoS₂/rGO. The weight loss in the temperature range 300–600 °C is due to the decomposition of rGO and oxidation reaction of MoS₂ to MoO₃. Thus, the content of MoS₂ in MoS₂/rGO can be calculated to be 83.24 wt%, which means that the content of rGO is 16.76 wt%.

In order to investigate the electrochemical properties of MoS₂/rGO and MoS₂ as anode materials, potassium-ion coin cells are fabricated. The electrochemical properties of PIBs consisting of MoS₂/rGO were determined by cyclic voltammetry (CV) at the working electrode, as shown in Figure 4a. The CV was performed linearly with a sampling frequency of 0.1 mV s⁻¹ and a voltage range of 3~0.1 V. The first cathodic scan revealed that one peaks at 0.955 V, indicating the formation and transformation reactions of K_xMoS₂ [36]. The anodic scan displayed two peaks at 1.66 V and 2.25 V, which are related to the detachment of K⁺ from the electrode material. As the electrode material interacts with the electrolytic solution during the initial charging and discharging process, a passivation layer is formed on the surface of the electrode, which is referred to as the “solid electrolyte film” or “SEI.” [37,38]. The overlap of the CV curves after the initial cycle indicates the high reversibility of MoS₂/rGO composite as PIBs during charge/discharge. In contrast, the CV curves of MoS₂ (Figure S3a) did not exhibit this phenomenon, indicating that MoS₂ has poor reversibility during the charge–discharge process.

Figure 4b illustrates the constant current charge/discharge (GCD) curve of MoS₂/rGO at a current density of 100 mA g⁻¹. The initial Coulombic efficiency (CE) of MoS₂/rGO is 44.3%, which is lower than that of MoS₂ and MoS₂-GO (Figure S3b,c). This is attributed to the formation of SEI during the initial charge/discharge process, which results in the irreversible reduction of the electrolyte and partially reversible redox reactions of MoS₂ [39]. The overlapping GCD curves after the initial cycling of MoS₂/rGO, which was not exhibited by MoS₂ and MoS₂-GO, once again demonstrates the high reversibility of MoS₂/rGO composites for PIBs. As illustrated in Figure S3d, the initial discharge capacity of MoS₂ is high at 100 mA g⁻¹ current density, yet it declines rapidly. After 50 cycles, the specific discharge capacity was only maintained at 171.95 mAh g⁻¹. MoS₂-GO not only has the lowest initial capacity (140.53 mAh g⁻¹ at 100 mA g⁻¹ current density) but also has a capacity retention of only 40.9% after 50 cycles. MoS₂/rGO demonstrated the highest initial capacity as well as best capacity retention (discharge capacity was maintained at 349.6 mAh g⁻¹ after 50 cycles). The remarkable electrochemical performance is largely attributable to the ultrathin scale structure, enlarged interlayer structure, and excellent intrinsic conductivity of MoS₂.

The rate performance of MoS₂/rGO is shown in Figure 4c, showing discharge capacities of 287.15, 266.7, 245.6, 221.6, 215.2, 197.1 and 184.6 mAh g⁻¹ at 50, 100, 200, 400, 500, 800, and 1000 mA g⁻¹, respectively, which are significantly better than those of MoS₂ and MoS₂-GO (Figure S3e,f) When the current density is restored to 50 mA g⁻¹, MoS₂/rGO can recover a specific capacity of 315.02 mAh g⁻¹, which is higher than the initial capacity, whereas MoS₂ and MoS₂-GO struggle to approach the initial capacity. These results indicate that the unique squamous structure of MoS₂/rGO can provide capable rate performance.

Figure 4d demonstrates the 100 charge–discharge cycle characteristics of MoS₂/rGO, MoS₂, and MoS₂-GO at 500 mA g⁻¹. The initial discharge specific capacity of MoS₂/rGO is 272.49 mAh g⁻¹, while the discharge specific capacity after 100 cycles is 269.5 mAh g⁻¹, with a capacity retention of 99%. In contrast, MoS₂-GO exhibited the lowest initial capacity (81.91 mAh g⁻¹) and only 36% capacity retention after 100 cycles, which was slightly higher than that of MoS₂ (35.91% capacity retention after 100 cycles). The excellent electrochemical performance at high current density is attributed to the scale-like lamellar structure that provides more K-embedding sites, and the enlarged interlayer spacing ensures that the material remains stable during the charging and discharging process [24]. At a high current density of 500 mA g⁻¹, the capacity retention of MoS₂/rGO is close to 99%, indicating that it is more suitable for operation at high current densities. In light of these findings, we sought to ascertain whether MoS₂/rGO exhibits comparable cycling stability at higher current densities. Figure 4e presents the long weekly cycle plot of MoS₂/rGO for 500 cycles at a current density of 1 A g⁻¹. It can be seen that MoS₂/rGO is still able to maintain good cycling stability compared to MoS₂ and MoS₂-GO, retaining 76% capacity after 500 cycles [40]. In addition, other MoS₂ anode electrodes are listed in

Table 1. The comparison with previous work on the MoS₂-based electrodes for PIBs.

Chemical Formula	Electrochemical Performance	Cycle Number
MoS2 [41]	102 mAh g−1 (100 mAg1)	86 mAh g−1 (200 cycles)
MoS2@C [42]	~290 mAh g−1 (500 mAg1)	241 mAh g−1 (100 cycles)
MoS2@Cnanosheets [43]	300 mAh g−1 (100 mA g−1)	164.5 mAh g−1 (350 cycles)
MoS2/MXene [44]	271.4 mAh g−1 (50 mA g−1)	206 mAh g−1 (100 cycles)
C-MoS2 [39]	~340 mAh g−1 (1000 mA g−1)	273 mAh g−1 (100 cycles)
This work	272.49 mAh g−1 (500 mA g−1)	269.5 mAh g−1 (100 cycles)

and compared to MoS₂/rGO. Of particular note is the exceptionally high cycling stability of MoS₂/rGO.

In order to further study the electrochemical stored procedure of MoS₂/rGO, CV curves at varying scan rates (ranging from 0.1 to 0.8 mV s⁻¹) were tested. It can be observed that MoS₂/rGO maintains the same shape as the initial redox peaks, whereas MoS₂ and MoS₂-GO do not exhibit the same behavior. This suggests that the MoS₂/rGO exhibit excellent electrochemical reversibility (Figure 5a and Figure S4a,b). Furthermore, the charge storage mechanism was evaluated using power-law analysis [45]:

$$i=a{\nu }^b$$

(1)

where ‘i’ represents the peak current, ‘ν’ represents the scan rate (mV s⁻¹), and ‘a’ and ‘b’ represent constant parameters. A value of ‘b’ approaching 1 indicates that capacitive behavior is the dominant phenomenon, whereas a value of ‘b’ approaching 0.5 indicates that charge storage is dependent on diffusion processes. As illustrated in Figure 5b, the b-values of the anodic and cathodic peaks of MoS₂/rGO are 0.81/0.66, indicating that the electrochemical potassium storage behavior of the material is primarily governed by the capacitance effect. The b-values of the anodic and cathodic peaks of MoS₂ and MoS₂-GO are 0.77/0.65 and 0.87/0.69, respectively, and the capacitive contributions are equally dominant (Figure S4c). The contribution of value-diffusive behavior and -capacitive behavior to the capacity can be calculated using the following equation [46]:

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

(2)

where k₁, k₂ are constants, and k₁v and k₂v^1/2 reflect the contributions of capacitive-controlled and diffusion-controlled processes, respectively. As illustrated in Figure 5c, at a sweep rate of 0.8 mV s⁻¹, approximately 85.8% of the total capacitance of MoS₂/rGO is attributed to the capacitive control process. As the sweep rate increases from 0.1 mV s⁻¹ to 0.8 mV s⁻¹, the capacitance contribution increases from 68.4% to 85.8% (Figure 5d). These results indicate that the surface diffusion behavior dominates in MoS₂/rGO. The capacitance contribution of MoS₂ and MoS₂-GO is significantly less than that of MoS₂/rGO (Figure S4d–g). This is primarily due to the unique scale-like lamellar structure of MoS₂ which provides a greater number of active sites for K⁺ and enhances charge storage. This is also the reason for the excellent rate performance of MoS₂-GO. The diffusion coefficient of K⁺ during a complete charge/discharge is calculated by the following equation [47]:

$$D=\frac{4}{\pi \tau }{\left(\frac{n_m{V}_m}{S}\right)}^2{\left(\frac{\Delta E_s}{\Delta E_t}\right)}^2$$

(3)

where τ is the relaxation time, n_m is the number of moles, V_m is the molar volume of the electrode material, S is the electrode/electrolyte contact area, the ΔE_s is the pulse-induced voltage change, and ΔE_t is the voltage change due to constant current charging (discharging). The calculated results are shown in Figure 5c, and the ion diffusion rate of MoS₂/rGO is approximately 10⁻¹¹ cm s⁻¹, which is comparable to that of other MoS₂ anodes [28]. Furthermore, the ion diffusion rates of MoS₂ and MoS₂-GO were calculated, with values of approximately 10⁻¹² cm s⁻¹ for MoS₂ and a range of 10⁻¹² to 10⁻¹¹ cm s⁻¹ for MoS₂-GO (Figure S4h,i).

The electrochemical performance of batteries is mainly assessed in terms of the charge transfer capacity and the conductivity of the electrodes [48]. To gain deeper insights into the physical characteristics and outstanding reversible capacity of the MoS₂/rGO composite electrode, a frequency range of 0.1–100,000 Hz was employed using an electrochemical workstation. Electrochemical impedance spectroscopy (EIS) was performed on a MoS₂/rGO anode to evaluate its use in PIBs. Figure 5d shows the EIS comparison images of MoS₂/rGO, MoS₂, and MoS₂-GO. The internal resistance of the cell can be analyzed in the high-frequency region, while the mid-frequency region indicates the transfer resistance of the electrode material interface. The larger the ion diffusion coefficient and the steeper the slope, the greater the diffusion capacity. In the equivalent diagram, R1 is the internal resistance, R2 is the interface transfer resistance, CPE1 is the constant phase angle element (similar to a capacitor), and W1 is the Warburg impedance. It can be seen that the semicircular diameter of MoS₂ and MoS₂-GO are larger than that of MoS₂/rGO, and the slope of the diagonal line is also smaller than that of MoS₂/rGO, which proves that the transfer of K⁺ in MoS₂ and MoS₂-GO is slow, while the interfacial transfer resistance of MoS₂/rGO is relatively small, and the reduction in the charge transfer resistance leads to the increase in the ion diffusion coefficient, and the surface of the electrode is more stabilized [49].

In order to further study the diffusion kinetics of K⁺ in MoS₂, the de/insertion process of K⁺ was calculated with first principles. All possible embedding positions of K in 2H-MoS₂, which is an AB-stacked periodic structure, have been scrutinized. K⁺ has been inserted in the AB layer at sites A and B in MoS₂ (Figure S5a,b). At site A, K⁺ is located in an octahedron of six S atoms. While at site B, K⁺ is located in a tetrahedron of four S atoms. After optimizing the structure, all K⁺ interpolation positions are located at the A site (Figure S5c,d), which indicates that site A is the stable site for K⁺ to insert. The calculation model of MoS₂/rGO is illustrated in Figure S6a, which comprises three layers of MoS₂ and three layers of graphene, with a layer spacing of 6.39 Å. The computational model for K⁺ adsorption on the MoS₂/rGO surface is shown in Figure S5b. Furthermore, the binding and adsorption energies of MoS₂/rGO were calculated to be −2.46 eV and −3.9 eV, respectively. The high absolute value of the adsorption energy indicates that K⁺ is more favored on the surface of the material, which corresponds to the capacitor behavior described above. Additionally, the negative value suggests that the structure is stabilized, which positively affects the electrochemical properties of the material.

Figure 6a illustrates the density of states of MoS₂ with a band gap of 1.2 eV. Upon the addition of rGO, the band gap of MoS₂/rGO disappears and crosses the Fermi energy level, indicating that rGO is capable of enhancing the electrical conductivity of the material. It is noteworthy that the projected density of states (PDOS) indicates that the p-orbitals of C prompt the d-orbitals of Mo to cross the Fermi energy level, thereby directly increasing the intrinsic conductivity of MoS₂ (Figure 6c,d). We considered two migration paths for K in MoS₂/rGO, one where K⁺ migrates from one octahedral position to neighboring octahedral positions and the other where K⁺ migrates to obliquely oriented octahedral positions (Figure S7). As shown in Figure 6b, the migration potential barrier of K⁺ in MoS₂/rGO for adjacent octahedral positions is 0.3 eV, which is significantly smaller than that of K⁺ in MoS₂ (0.4 eV). The results indicate that K⁺ diffuses more readily into neighboring octahedral positions.

3. Materials and Methods

3.1. Material

Sodium molybdate dihydrate (Na₂MoO₄·2H₂O), CTAB, (C₁₉H₄₂BrN), ice acetic acid (CH₃COOH), and anhydrous ethanol (C₂H₅OH) were purchased from Sinopharm Chemical Reagent Co., Ltd. (Shanghai, China). Thiourea (CH₄N₂S, 99% purity) was purchased from Shanghai Aladdin Industrial Co., Ltd. (Shanghai, China). Single-layer industrial grade graphene oxide was purchased from Suzhou Carbon Feng Technology Co., Ltd. (Suzhou, China).

3.2. Preparation of MoS₂/rGO Composite

An amount of 0.09 g (30 wt%) of GO powder was dissolved in 30 mL of deionized water (DI) and ultrasonicated for 1.5 h (a black-brown solution was generated). Next, 20 mL of N-methyl pyrrolidone was added to the solution and stirred for 10 min. An amount of ice acetic acid was added to the solution, the pH of the solution was adjusted to 6, and the solution was stirred for 5 min to ensure even mixing. CTAB (0.15 g) was added to a 10 mL DI beaker and stirred until it was a milky white solution. Finally, the two solutions were mixed and stirred for 5 h. The two solutions were stratified at the beginning of mixing (the upper layer was graphene oxide, and the lower layer was a transparent liquid), and after 5 h of magnetic stirring, they were uniform solutions with some foam (the nature of CTAB). To the above solution, 0.45 g of Na₂MoO₄·2H₂O and 0.637 g of CH₄N₂S were added and continuously stirred for 1 h. Then, the mixture was transferred into a 100 mL high-pressure reactor, kept at 200 °C for 24 h, and allowed to cool to room temperature. The black sediment was collected and rinsed 3 times with DI and anhydrous ethanol. The collection was dried overnight in a vacuum drying oven. Finally, the MoS₂/rGO composites were obtained by holding the dried black powder at 820 °C for 120 min and at a heating rate of 5 °C min⁻¹ under an argon atmosphere.

3.3. Structural Characterization

The crystalline phases of the samples were characterized by X-ray diffraction (XRD, XD3, Beijing Puxi Tongyong Co., Ltd., Beijing, China) with Cu-Kα radiation (λ = 1.54 Å). The morphologies and structures of the products were characterized by field-emission scanning electron microscopy (FE-SEM, Quanta250FEG, FEI, OR, USA) and high-resolution transmission electron microscopy (HRTEM, Talos F200, FEI, OR, USA). X-ray absorption fine structure (XAFS) spectra of the powder samples were obtained using 8C. The component and valence analysis of the synthesized samples were carried out by X-ray photoelectron spectroscopy (XPS, ESCALAB Xi+, Thermo Fisher, Morecambe, UK).

3.4. Electrochemical Characterization

Active material, acetylene black, and polyvinylidene difluoride (PVDF) at a weight ratio of 8:1:1 was mixed with N-methyl-2 pyrrolidone (NMP) solvent to prepare the anode material. The slurry mass loading on the copper collector was approximately 1 mg cm⁻². This electrode was dried in a vacuum oven at 50 °C for 12 h. Next, the active electrode was transferred into an Ar-filled glove box to assemble the cell. Coin cells were of the CR2032 type, with a fiberglass to separate the anode and K-block. The electrolyte of choice was 1 M of potassium bis (fluor sulfonyl) imide (KFSI) in ethyl methyl carbonate (EMC). Electrochemical testing of the assembled coin battery was performed at various current densities in the 0.1–3.0 V range by using the Blue Battery Test System. Galvanostatic charge/discharge tests and cyclic voltammetry (CV) measurements in the potential window of 0.1–3.0 V versus K⁺/K and electrochemical impedance spectroscopy (EIS) measurements in the frequency range 0.1 Hz to 100 kHz at open-circuit voltage were performed using an electrochemical workstation (CS350 Coster Co., Ltd., Wuhan, China). All electrochemical measurements were performed at room temperature.

3.5. Theoretical Calculation

The calculations in this work are based on the DS-PAW, as well as the Projected Augmented Wave (PAW) method for calculation [50,51]. Ultra-soft pseudopotentials described the interaction between ionic nuclei and valence electrons. This study considers the binding energy, density of states, and migration barrier of MoS₂/rGO. The exchange and correlation terms were computed using the generalized gradient approximation (GGA) of the Perdew–Burke–Ernzerh method, which was parameterized by Perdew [52,53]. Brillouin zone integrations were performed using a Gamma-centered k-point grid. MoS₂ exhibits hexagonal symmetry and belongs to the P63/MMC space group.

The calculation was performed using a cutoff energy of 480 eV and a 3 × 3 × 1 k-point grid to ensure total energy convergence. The calculations were considered converged when the maximum force on the atoms was below 0.05 eVÅ⁻¹, the maximum stress was below 0.1 GPa, and the maximum displacement between cycles was below 0.002 Å. Periodic structures are utilized in the calculations. To calculate the binding energy, density of states, and migration barriers, we extended MoS₂ to 3 × 3 × 3 and graphene to 4 × 4 × 4, and used this as the basis for constructing the MoS₂/rGO model of MoS₂-wrapped rGO. The binding energy is calculated using the following formula:

$$E_{binding}=E_{Mo{S}_2/rGO}-E_{Mo{S}_2}-E_{rGO}$$

(4)

The adsorption energy is calculated as follows:

$$E_{adsorption}=E_{Mo{S}_2/rGO}-E_{Mo{S}_2}-E_{rGO}-E_K$$

(5)

The calculation of the migration barriers took into account the distinct migration paths of K⁺ in MoS₂/rGO and MoS₂.

4. Conclusions

In summary, MoS₂/rGO composites were successfully prepared via a one-step hydrothermal method. Owing to the presence of rGO in the synthesis process, the MoS₂ anodes in such composite possess a unique scaled structure with larger layer spacing, compared with pure MoS₂. This indirectly increases the specific surface area, provides more active sites for K⁺, and ensures structural stability. In addition, after the hydrothermal reaction, the incorporation of rGO can increase the intrinsic conductivity of MoS₂, and the composite shows high electronic conductivity.

As a result, MoS₂/rGO composite anodes exhibited a better rate performance (287.15, 266.7, 220.3, and 184.6 mAh g⁻¹, at 50, 100, 500, and 1000 mA g⁻¹, respectively) and cycle stability (99% capacity retention after 100 cycles at 500 mA g⁻¹) than that of anodes based on pure MoS₂ and MoS₂-GO. This work demonstrates the significant potential of MoS₂/rGO in high-current PIBs and presents a practical strategy for the development of high-performance anode materials for PIBs.