Molybdenum-doped few-layered SnS₂ architectures with enhanced electrochemical supercapacitive performance

Abstract

Molybdenum doped few-layered SnS₂ (m-SnS₂) architectures have been fabricated via a facile hydrothermal route. The results indicate that the resultant m-SnS₂ is composed of numerous few-layered nanosheets and possesses a hierarchical architecture with mesoporous features. When evaluated as an electrode material for supercapacitors, the m-SnS₂ architecture exhibits obviously enhanced electrochemical performance with high specific capacitance, good cycling stability and rate capability in comparison with the bare SnS₂ sample. It is believed that the superior electrochemical properties can be attributed to the robust mesoporous architectures with large surface area, the few-layered nature as well as other positive factors such as expanded interlayer space and rich dislocations which originate from the molybdenum doping.

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1. Introduction

To date, the urgent demand for clean and sustainable energy has driven considerable effort toward the development of energy storage and conversion systems.¹ Supercapacitors are an important power source and attract much attention due to their fast charge and discharge process, high power performance, long cycle life, and relatively low cost.^2,3 According to their energy storage mechanism, supercapacitors are classified into electrical double layer capacitors (EDLCs) and pseudocapacitors (PCs). The capacitance of the EDLCs comes from the charge accumulated at the electrode/electrolyte interface. Carbon-based materials such as carbon nanotubes, mesoporous carbon as well as graphene have been intensively used for EDLCs.^4,5 While the capacitance of the pseudocapacitors results from fast faradic redox reactions in the electrode material. Pseudocapacitors are being widely investigated because of their large specific capacitance and high-energy properties. Since pseudocapacitance arises from the redox reaction of electroactive materials, conductive polymers such as polypyrrole,^6,7 transition metal oxides^8–12 and sulfides^13–15 with several oxidation states are considered promising electrode materials for pseudocapacitors. Whether EDLCs or pseudocapacitors, the capacitor performance is largely dependent on the electrode material. Therefore, considerable effort has been exerted to explore the new electrode materials with good capacitive performance.

The last decades have witnessed the rise and the great success of graphene, which have encouraged the exploration of other kinds of 2D atomic-layer crystals.^16–18 Among them, layered transition metal dichalcogenides (LTMDs) with strong covalent bonds in plane and relatively weak van-der-Waals-like forces between the layers have received significant attention for a wide spectrum of applications in catalysis, nano-electronics, optoelectronics, Li-ion batteries and supercapacitors.^19–21 Moreover, the layered structure analogue to graphite not only makes it a suitable intercalation host for foreign ions such as alkali metal ions storage, but also is favourable to be exfoliated into graphene-like monolayer and few layers. Compared to their bulk layered compounds, exfoliated single or few-layered 2D crystals usually exhibit extraordinary properties such as large specific surface area and high carrier mobility and have found potential applications in energy storage and conversion.^22–24 For example, graphene-like LTMDs can not only exhibit better capacitive properties in their electrochemical double layers, but also can give rise to a faradaic capacitance due to the diffusion of the ions such as Li⁺, Na⁺ and K⁺ ions into the LTMDs layers, which plays an important role in enhancing charge storage capabilities.^25–28 Recently, graphene-like MoS₂ with specific morphologies has been intensively designed and investigated as an electrode material for electrochemical capacitors.^29–38

Besides MoS₂, as a typical LTMD, SnS₂ has a layered hexagonal CdI₂-type structure and has been widely investigated as a possible catalyst,³⁹ gas sensor⁴⁰ and alternative anode in lithium and sodium ions batteries.^41,42 To improve its properties, fabricating composite and doping with metal elements have been proved to be effective methods. For instance, Ce doped SnS₂ material demonstrates obviously enhanced Li storage performance due to the fact that Ce doping can stabilize the crystal lattice during charge/discharge and improved photocatalytic effect for the modified electronic structure of SnS₂.^43,44 Doping SnS₂ with copper can increase the carrier density and the optical absorption and makes the photocatalytic process more efficient.⁴⁵ However, to our knowledge, the investigation of SnS₂ nanostructures as a supercapacitor electrode is seldom and the electrochemical properties of SnS₂ for pseudocapacitors still need to be further investigated.

In view of the similar layered microstructure between SnS₂ and MoS₂ as well as the possible resultant positive effect on electrochemical property by introduction of molybdenum into the lattice of SnS₂, in this work, we introduce a facile hydrothermal strategy to fabricate molybdenum doped few-layered SnS₂ architectures and investigate their electrochemical supercapacitive performance. The samples are well characterized by XRD, SEM, HRTEM, EDX, XPS, nitrogen adsorption–desorption, Raman spectra and electrochemical impendence spectroscopy (EIS) techniques.

2. Experimental section

2.1 Synthesis of molybdenum-doped few-layered SnS₂ architectures

In a typical synthesis, 0.447 g of SnCl₄·5H₂O and 1.10 g of L-cysteine (L-cys) were successively dissolved into 50 mL of deionized water under constant stirring. Then 10 mL of sodium molybdate aqueous solution (22.5 mM) was slowly dropwise added into the above solution under vigorous stirring. The laurel-green precipitates immediately appeared. After stirring for another 1 h, the resultant mixture was transferred into the 100 mL Teflon-lined stainless steel autoclave, which was heated and maintained at 200 °C for 24 h. After cooling naturally, the black solid product was collected by centrifugation and washed several times with deionized water and ethanol, and dried in the vacuum oven at 80 °C for 12 h to obtain Mo-doped SnS₂ sample termed as m-SnS₂. For comparison, the bare SnS₂ sample was also prepared by a similar synthetic route.

2.2 Characterizations

XRD patterns were recorded with a Thermo X'TRA X-ray diffractometer with Cu Kα radiation (λ = 0.154056 nm). The morphologies of the samples were observed by using a SIRION-100 field emission scanning electron microscopy (FE-SEM). The elemental compositions were analyzed by energy dispersive X-ray spectroscopy (EDX, GENENIS-4000). X-ray photoelectron spectra (XPS) were processed on a Perkin-Elmer PHI5000c XPS, using C 1s (B. E. 284.6 eV) as a reference. TEM and HRTEM characterizations were performed on a JEOL JEM-2010 TEM operating at 200 kV, for which each sample was prepared by dispersing it in acetone and drop-casting onto a 200 mesh copper grid coated with holey carbon. The nitrogen adsorption–desorption was conducted at 77 K on a Quantachrome NOVA 2000e sorption analyzer (USA), and the Brunauer–Emmett–Teller (BET) surface area was estimated from the adsorption data. Raman spectra were recorded with 514 nm Ar ion laser at 6 mW for 50 s employing a Jobin Yvon LabRam HR spectrometer. Electrochemical impendence spectroscopy (EIS) was obtained by applying a sine wave with amplitude of 0.5 mV in the frequency range 200 kHz to 0.01 Hz on a PARSTAR 2273.

2.3 Electrochemical measurements

Electrochemical measurements were performed with a CHI660D electrochemical analyzer. And the electrochemical tests were taken in 3.5 M KOH electrolyte using a three electrode cell comprising a saturated calomel electrode (SCE) electrode as reference, Pt wire as counter and the as-prepared samples-based electrode as the working electrode. Working electrodes were prepared by mixing electroactive material, carbon black and polytetrafluoroethylene (PTFE) binder in a mass ratio of 80 : 10 : 10 to obtain a slurry. Then the slurry was pressed onto the nickle foam current collector (1 cm²) and dried at 60 °C for 12 h. The specific capacitance can be calculated from the galvanostatic discharge curve according to eqn (1):

C = IΔt/ΔVm (1)

where C is the specific capacitance, I is the current, Δt is the discharge time, ΔV is the potential window, and m is the loading mass of the electroactive materials.

3. Results and discussions

3.1 Microstructure, morphology and component of the Mo-doped few-layered SnS₂ architectures

Fig. 1 depicts the XRD patterns of the as-prepared bare SnS₂ and molybdenum-doped SnS₂ (m-SnS₂) samples. It can be seen from Fig. 1a that the diffraction peaks of the bare SnS₂ can agree well with the hexagonal SnS₂ phase (JCPDS no. 23-0677). High and sharp diffractions are indicative of good crystallinity. Especially, the high (001) peak situated at 15.19° suggests the well-developed and stacked SnS₂ layers along c axis direction. According to Bragg's equation, the d spacing of (001) plane is calculated as 0.588 nm, which is in agreement with that (0.589 nm) of hexagonal SnS₂. Compared with the bare SnS₂, the diffractions of m-SnS₂ can also be well assigned to the hexagonal SnS₂ phase (JCPDS no. 23-0677). No characteristic diffraction of MoS₂ phases can be detected in the XRD pattern, indicating that a small amount of molybdenum has been doped into SnS₂ crystal lattice. Note that the obviously weak and broadened (001) diffraction peak located at 14.68° of m-SnS₂ indicate that (001) plane growth of the m-SnS₂ crystals is greatly inhibited and a few-layered SnS₂ is formed, which was also observed in Zn or Ce doped SnS₂.^43,44,46 The slight shift of the (001) diffraction peak of m-SnS₂ to a lower angle can be attributed to Mo-doping and means enlarged d-spacing. The d spacing of (001) plane of m-SnS₂ is calculated as 0.604 nm, which is slightly larger than that of the bare SnS₂. The decreased layers and enlarged interlayer spacing of the m-SnS₂ could be attributed to the generated structural strain and dislocation due to the expansion of the crystal lattice after substitution of Mo with larger radius for smaller Sn atoms.^43–46 The microstrain (ε) is calculated using the following relation:⁴⁴

ε = β cos θ/4 (2)

where β is the full width at half maximum of the peak; θ is Bragg's angle.

Fig. 1 XRD patterns of the as-prepared (a) bare SnS₂ and (b) m-SnS₂.

Additionally, the dislocation density (δ) is also calculated using the formula (3):⁴⁴

δ = 1/D² (3)

in which D is the crystallite size of the sample estimated using Scherrer's formula.

The lattice parameters such as microstrain, dislocation density and unit cell volume are calculated and summarized in Table 1.

Table 1 The microstructural parameters of the bare SnS₂ and m-SnS₂

Sample	Lattice parameters		Microstrain ε × 10^3	Dislocation density δ × 10^14 lines per m^2	Unit cell volume nm^3
	a	c
SnS2	3.6264	5.8847	4.58	77.45	67.02
m-SnS2	3.6448	6.1601	7.80	225.21	69.78

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It can be noticed that the values of lattice parameters and thus the unit cell volume with Mo doping are higher than those of bare SnS₂, which confirms that the substitution of Sn by larger Mo can cause the lattice distortion and lead to the few-layered structure and expanded interlayer distance of the m-SnS₂. The few-layered feature can effectively reduce interlayered resistance and promote electron fast transport from the active sites to the electrodes.^19–22 Moreover, the expanded interlayer space can decrease the energy barrier of ions insertion/extraction and accommodate more ions. In addition, the additional dislocations or defects can act as new sites for ions insertion and extraction. It is envisioned that the as-obtained m-SnS₂ can exhibit higher capacitance and superior rate capability in comparison with the bare SnS₂.

EDX is also adopted to identify the elemental compositions of the m-SnS₂ sample. The results are summarized in Table 2. The Table 2 shows that the m-SnS₂ sample contains Sn, Mo and S element. The calculated atomic ratio of S to (Sn + Mo) element is about 2.02, approaching the theoretical stoichiometric value. The mass fraction of the doped Mo in the m-SnS₂ sample can be ascertained to be around 7.94%.

Table 2 Elemental composition of the SnS₂ and m-SnS₂ sample. S: (Sn + Mo) represents the atomic ratio of the bare SnS₂ and m-SnS₂

Sample	Sn/wt%	Mo/wt%	S/wt%	S: (Sn + Mo)
SnS2	64.76	0.00	35.24	2.01
m-SnS2	56.20	7.94	35.86	2.02

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The samples were further investigated by Raman spectra as shown in Fig. 2. The spectra of both samples illustrate the strong characteristic peaks at 314.89, 314.16 cm⁻¹, respectively for the bare SnS₂ and m-SnS₂ sample, which could be assigned to the A_1g mode of SnS₂ according to the group theory analysis given by Lucovsky et al.⁴⁷ It is noticed that the intensity of A_{1 g} peak of the m-SnS₂ sample are visibly lower than that of the bare SnS₂. The weakened A_1g modes could be due to the reduced interlayer interactions and dielectric screening of the long-range Coulomb forces with decrease of the layer numbers.^48,49 Moreover, it is observed that m-SnS₂ sample has a down-shift of A_1g peak in comparison with bare SnS₂. The variation of A_1g peaks can be evidence of the few-layered SnS₂, which can also be observed in other few-layered transition metal sulfides.^50,51

Fig. 2 Raman spectra of the as-prepared (a) bare SnS₂ and (b) m-SnS₂ sample.

XPS is further used to study the chemical states of the m-SnS₂ sample. The two strong peaks at around 486.64 and 495.05 eV displayed in Fig. 3a can be attributed to Sn 3d_3/2 and 3d_5/2 respectively, which agrees well with the result of Sn⁴⁺ in SnS₂.^39,41–43 The Mo 3d orbit of the m-SnS₂ exhibits two peaks located at 229.00 and 232.41 eV, which are ascribed to the doublet of Mo 3d_5/2 and Mo 3d_3/2, respectively, suggesting the presence of Mo(IV) in the prepared Mo-doped SnS₂ product.^{32–36,52,53} In addition, it can be seen that the S 2p core-level XPS spectrum can be divided into two peaks, which indicates that there exist two chemical environments. The peaks at 161.57 and 162.79 eV could be attributed to the binding energies of S 2p_3/2 and S 2p_1/2,^41–43 indicating the existence of other binding signals, such as bridging S₂²⁻or apical S²⁻.^{41–43,52,53}

Fig. 3 XPS spectra of the as-prepared m-SnS₂ sample. (a) Sn 3d; (b) Mo 3d and S 2s and (c) S 2p.

The morphologies and microstructures of the samples are further inspected by SEM and HRTEM techniques. Fig. 4a shows that the bare SnS₂ exhibits a leaf-like nanosheet morphology and the SnS₂ nanosheets are inclined to be curved and interlaced together to diminish the surface energy. Furthermore, the detailed microstructures of the SnS₂ samples can be carefully examined by HRTEM. Fig. 4b distinctly displays the straightly arranged lattice fringes of the bare SnS₂ nanosheets with an interlayer spacing of 0.589 nm, corresponding to the (001) plane of SnS₂. The layer number can also be estimated about 18 layers. By contrast, the m-SnS₂ sample demonstrates a quasi-3D architecture with interconnected porous structures as shown in Fig. 4c. It can be observed that this hierarchical architecture is comprised of numerous nanosheets. Fig. 4d manifests about 10 slightly curved layers with a larger d spacing of 0.608 nm for the (001) plane of the m-SnS₂. Fig. 4e and f display that some discontinuous or twisted lattice fringes appear in the few-layered m-SnS₂, indicating rich defects of the sample. HRTEM results confirm the formation of few-layered m-SnS₂ and are well in line with above XRD analysis.

Fig. 4 SEM and HRTEM images of (a and b) bare SnS₂ nanosheets; (c–f) m-SnS₂ architectures.

Nitrogen adsorption–desorption isotherm measurements were carried out to determine the surface area of the as-prepared samples. Fig. 5 shows the corresponding N₂ adsorption–desorption isotherms and Barrett–Joyner–Halenda pore-size distribution plots for both the SnS₂ nanosheets and m-SnS₂ architectures. Fig. 5a and b show a typical type IV with a H3 hysteresis loop (at P/P₀ > 0.8), indicating the presence of a mesoporous structure for both samples. The Brunauer–Emmett–Teller (BET) specific surface area of the m-SnS₂ was measured as about 31.4 m² g⁻¹, which was larger than that of the bare SnS₂ (17.5 m² g⁻¹).

Fig. 5 Nitrogen adsorption–desorption isotherm of (a) SnS₂ nanosheets and (b) m-SnS₂ architectures.

Time-dependent experiment was also carried out to investigate the formation process of the as-prepared m-SnS₂ hierarchical architecture. Prior to hydrothermal reaction, the laurel-green tin molybdenum oxide precipitates were collected and characterized by SEM technique. Fig. 6a shows that the tin molybdenum oxide precursors are large-sized irregular particles. When hydrothermal time is 1 h, agglomerate nanoparticles appear as shown in Fig. 6b. As hydrothermal reaction proceeds for 4 h, the profiles of the nanoparticles become more clear from Fig. 6c. Moreover, the surfaces of the nanoparticles are covered with nanosheets as seen in the inserted figure in Fig. 6c. With the hydrothermal time reaches to 8 h, it can be observed in Fig. 6d that numerous underdeveloped nanosheets are crosslinked together to form hierarchical architectures.

Fig. 6 SEM images of the m-SnS₂ samples prepared with different hydrothermal time. (a) 0 h; (b) 1 h; (c) 4 h; (d) 8 h.

Based on the above experimental results, the formation of the hierarchical m-SnS₂ architectures in our synthetic route could be illustrated in Fig. 7. As MoO₄²⁻ anions were introduced to the Sn⁴⁺-contained solution before hydrothermal treatment, a laurel-green precipitate was observed, suggesting the formation of probable tin molybdenum oxide. During the following hydrothermal process, the resultant tin molybdenum oxide could act as a precursor and was gradually vulcanized to form Mo-doped tin disulfide nanosheets by H₂S arising from the decomposition of L-cys and further assemble into architectures. In view of the larger radius of Mo than Sn, Mo doping caused the structural dislocation and lattice distortion, which could inhabit the growth of SnS₂ layers along c axis and result in the few-layered structure with expanded interlayer distance.

Fig. 7 Schematic illustration of the formation process of the Mo doped-SnS₂ architectures.

3.2 Electrochemical performance of the samples

To evaluate the electrochemical performances of the as-prepared samples as supercapacitor electrode materials, we carried out cyclic voltammetry (CV) and galvanostatic charging/discharging measurements in a three-electrode system. Fig. 8a shows typical CV curves of the bare nickel foam, the bare SnS₂ nanosheets and the m-SnS₂ hierarchical architectures within the potential range from −0.1 to 0.4 V at a scan rate of 10 mV s⁻¹. It can be seen that the CV curves of all the sample electrodes are distinct from the ideal rectangular curves originating from the electric double-layer capacitance and can be indicative of the pseudocapacitance characteristics.^2–4 The obvious cathodic peak situated at around 0.27 V for both bare SnS₂ and m-SnS₂ could be attributed to the electrochemical insertion of K⁺ ions in the interlayer of layered SnS₂.^29–38 Additionally, the emerging reduction peak at about 0.21 V could result from the intercalation of K⁺ ions into the S–Mo–S layers.^34,54 The nearly symmetrical redox peaks of all the sample electrodes indicate the reversibility of the redox processes. In addition, it is well-known that the specific capacitance is proportional to the area of the CV curves.^2–4 Therefore, the capacitance contributed by nickel foam can be ignored for its negligible CV area. Moreover, it can be reckoned that the m-SnS₂ electrode possesses a higher capacitance than that of bare SnS₂ electrode. Fig. 8b and c depict the CV curves of the sample electrodes at different scan rates. It can be observed that even at a high scan rate, a similar CV shape and symmetric characteristics of the current response on voltage for all the sample electrodes can be maintained, which indicates good capacitive behaviour and reversibility as well as the fast diffusion of ions into the sample electrodes.

Fig. 8 Cyclic voltammogram (CV) curves of (a) the sample electrodes at 10 mV s⁻¹; (b) the bare SnS₂ electrode and (c) the m-SnS₂ electrode at different scanning rate.

Galvanostatic charge–discharge is a complementary method for measuring the specific capacitance of electrochemical capacitors at constant current. The galvanostatic charge–discharge curves of the bare SnS₂ nanosheets and m-SnS₂ architectures electrodes within the potential range from −0.1 to 0.4 V at different current densities are plotted in Fig. 9. It can be seen from Fig. 9a that the bare SnS₂ electrode demonstrates poor electrochemical performance with low specific capacitance. The discharge capacitance values of the bare SnS₂ at a current density of 1 A g⁻¹, 2 A g⁻¹, 4 A g⁻¹ and 10 A g⁻¹ are calculated about 89.4, 71.3, 64.9 and 50.6 F g⁻¹, respectively. In sharp contrast, the m-SnS₂ delivers larger specific capacitance of 213.2, 184.2, 152.6 and 78.7 F g⁻¹ at a current density of 1 A g⁻¹, 2 A g⁻¹, 4 A g⁻¹ and 10 A g⁻¹, respectively (Fig. 9b). The specific capacitance variations at different current densities for the sample electrodes are also described in Fig. 9c.

Fig. 9 Galvanostatic charge–discharge curves of (a) bare SnS₂ and (b) m-SnS₂ sample electrodes at different current densities and (c) capacitances versus current densities of the sample electrodes.

Because a long cycling performance is among the most important criteria for supercapacitors, an endurance test is conducted using galvanostatic charge–discharge cycles at 4 A g⁻¹. Fig. 10a and b exhibit the specific capacitance of the bare SnS₂ and the m-SnS₂ electrodes as a function of cycle number for 1000 cycles. The bare SnS₂ electrode displays an obvious capacity decay during cycles. In contrast, as for the m-SnS₂ electrode, still about 89% of the initial capacitance can be maintained after 1000 cycles without noticeable decrease. The superior electrochemical performance of the m-SnS₂ architecture could be attributed to following factors. Firstly, the few-layered characteristic of m-SnS₂ is advantageous to decreasing interlayer resistance and enhancing the kinetics of electrode process. Moreover, rich dislocations or defects could act as new sites for ions insertion. Secondly, the robust mesoporous architectures can effectively cusion the volume change, guarantee the stability of microstructures during cycling and provide more exposed surfaces for access of electrolyte and more channels for insertion/extraction of ions as well as shorter paths for rapid diffusion of charge carriers. Thirdly, the enlarged interlayer space can decrease the energy barrier of ions insertion/extraction and accommodate more ions. Finally, the additional capacitance contribution could be arising from the incorporated Mo of m-SnS₂.

Fig. 10 Cycling performance of (a) the bare SnS₂ and (b) m-SnS₂ electrodes at a current density of 4 A g⁻¹.

Fig. 11 shows Nyquist plots of the EIS data obtained for the bare SnS₂ and m-SnS₂ electrode. Both impedance spectra are composed of semicircles in the high frequency region followed by a linear slope in the low frequency region, which reveals that the capacitive behavior is due to the non-faradaic charge storage mechanism. It is known that a small semi-circle for the electrode suggests a low interfacial charge-transfer resistance.^30–36 It is clearly observed that the diameter of the Nyquist circles of m-SnS₂ is smaller than that of bare SnS₂, which also confirms the superior electrochemical performance of m-SnS₂.

Fig. 11 EIS spectra of the as-prepared bare SnS₂ and m-SnS₂ sample.

4. Conclusions

In conclusion, Mo doped SnS₂ architectures composed of few-layered nanosheets with mesoporous structure have been prepared through a facile hydrothermal route and investigated as an electrode material for supercapacitor. The electrochemical tests indicate that compared with the bare SnS₂, the Mo doped SnS₂ architectures exhibit prominently improved supercapacitive properties including high specific capacitance, good cycling stability and rate capability, which could be ascribed to their structural merits such as robust architectures with large surface area, the few-layered features, expanded interlayer space and rich dislocations originated from the Mo doping.

Acknowledgements

This work is financially supported by the Guangdong Natural Science Foundation for Joint Training Innovative Talents in East-West-North Guangdong (2014A030307030, 2015A030310195), the Science and Technology Planning Project of Guangdong Province (2014A010106032), the Stong Innovation School of Engineering Program of of Department of Education of Guangdong Province (Distinctive Innovation Project No. 2014KTSCX157), Natural Science Foundation of Lingnan Normal University (LZL1502, LZL1402), Distinguished Young Talents Foundation in Higher Education of Guangdong (2013LYM_0054).

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