C Composites for High Performance Lithium Ion Battery Anode Material

Electrochimica Acta 239 (2017) 74–83

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Electrochimica Acta journal homepage: www.elsevier.com/locate/electacta

Nanoporous MoS2/C Composites for High Performance Lithium Ion Battery Anode Material Zengmei Wanga,* , Gui Weia , Kiyoshi Ozawab , Yaling Caia , Zhenxiang Chengc, Hideo Kimurab a b c

Jiangsu Key Lab. for Civil Engineering Materials, School of Materials Science and Engineering, Southeast University, Nanjing, Jiangsu 211189, PR China National Institute for Materials Science, Tsukuba, Ibaraki 305-0047, Japan Institute for Superconducting and Electronics Materials, University of Wollongong, Innovation Campus, North Wollongong, NSW 2500, Australia

A R T I C L E I N F O

Article history: Received 14 December 2016 Received in revised form 5 April 2017 Accepted 6 April 2017 Available online 8 April 2017 Keywords: MoS2 Nanoporous structure Lithium Ion Battery Anode

A B S T R A C T

MoS2 is a promising anode material due to its graphite-like structure and high theoretical capacity (
670 mAh g1). Fast capacity fading and severe structural deterioration as a result of repeated lithium insertion restrict its application, however. Recently, we successfully synthesized a new type of MoS2/C composite and found that the composite showed excellent electrochemical performance, such as by its high capacity of 982 mAhg1 after 100 cycles. The composite was synthesized by a simplified thermal reduction method in which cetyl trimethylammonium bromide (CTAB) and silica nanospheres are used as the carbon source and template material, respectively. Such favorable capacity retention could be attributed to the large specific surface area due to the effective carbon dispersion in the composite, as well as the enlargement of the interlayer spacing in the S-Mo-S structure. © 2017 Elsevier Ltd. All rights reserved.

1. Introduction As one of most important energy storage devices, lithium ion batteries (LIBs) have inspired intensive interest over the past few decades due to the deficiency of natural energy and the increasing demand for energy storage. The practical applications of LIBs largely depend on the reversible capacity and cycling stability of their electrodes. Therefore, how to improve the electrochemical performance of electrodes has become highly topical for both chemists and materials scientists [1–4]. Graphite has been used as the standard commercial anode material for LIBs for many years due to its long cycle life, low cost, and structural stability during cycling [5,6]. It suffers from a relatively low theoretical capacity (372 mAh g1), however, which greatly restricts its application in high-energy LIBs. MoS2 has an analogous structure: two sulfur atom layers sandwiching a layer of molybdenum atoms, formed one monolayer, Van der Waals bonding between the monolayer [7]. Compared with graphite, the theoretical specific capacity of MoS2 is
670 mAh g1 (based on 4 mol of Li+ ion insertion) [8], almost twice that of graphite. Therefore, MoS2 has been investigated as an anode material for lithium-ion batteries (LIBs),

* Corresponding author. E-mail address: [email protected] (Z. Wang). http://dx.doi.org/10.1016/j.electacta.2017.04.033 0013-4686/© 2017 Elsevier Ltd. All rights reserved.

based on its intercalation reaction [9]. However, there are two main problems for MoS2 as an anode material, fast capacity fading and severe structural deterioration as a result of repeated lithium insertion [10]. So, finding ways to improve the cycling stability and retain its high reversible capacity has attracted the attention of many researchers. Recently, there have been some reports on improving the capacity and cycling stability of MoS2 via morphology modification. Xiong et al. [11] present the crumpled reduced graphene oxide (RGO) decorated MoS2 nanoflowers on carbon fiber cloth, and the crumpled RGO decorated MoS2 nanoflowers anode exhibits high specific capacity (1225 mAh g1) and excellent cycling performance (680 mAh g1 after 250 cycles). In Hu et al’s. [12] work, three-dimensional (3D) hierarchical MoS2/ polyaniline (PANI) nanoflowers were successfully fabricated via a simple hydrothermal method. The obtained MoS2/C sample exhibited more excellent electrochemical performance due to its excellent electronic conductivity resulting from the close integration of MoS2 nanosheets with carbon matrix. High reversible capacity of 888.1 mAh g1 with the coulombic efficiency maintained at above 90% from the first cycle were achieved at a current density of 100 mA g1. Three-dimensional (3D) flower-like MoS2 spheres composed of nanosheets (less than 10 nm thick) was constructed by Yang et al., [13] via a simple alcohol-assisted solvothermal route. These prepared materials display high

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performance as anode materials for lithium ion batteries, which includes high initial discharge capacity (1346 mAh g1 at a current density of 100 mA g1), good coulombic efficiency (77.49% retention for the first cycle and
100% for the subsequent cycles), excellent cycling performance (947 mAh g1 at 100 mA g1 after 50 cycles), and remarkable rate performance. Hierarchical hollow nanoparticles of MoS2 nanosheets reported by Wang et al., [14] with enhanced interlayer spacing which were synthesized by a simple solvothermal reaction at a low temperature. These hollow nanoparticles exhibited a reversible capacity of 902 mAh g1 at 100 mA g1 after 80 cycles. It has been shown that MoS2/C composites with nanoporous structure and purpose-designed morphology are high performance lithium ion battery anode materials. Chen et al. [15] designed hierarchical MoS2 tubular structures internally wired by carbon nanotubes (CNTs), these porous MoS2 tubular structures are constructed from building blocks of ultrathin nanosheets, which are believed to benefit the electrochemical reactions. Benefiting from the unique structural and compositional characteristics, these CNT-wired MoS2 tubular structures deliver a very high specific capacity of
1320 mAh g1 at a current density of 0.1 A g1, exceptional rate capability, and an ultralong cycle life of up to 1000 cycles. George et al. [16] investigated the mechanistic role of carbon additive by comparing equal loading of standard Super P carbon powder and carbon nanotubes (CNTs). The latter offer a nearly 2-fold increase in capacity and a 45% reduction in resistance along with coulombic efficiency of over 90%. These insights into the phase changes during MoS2 conversion reactions and stabilization methods provide new solutions for implementing cost-effective metal sulfide electrodes, including Li-S systems in high energy-density batteries. Here, MoS2/C composite and MoS2 with different morphology were prepared by a simple thermal reduction method using ammonium tetrathiomolybdate (ATTM), cetyl trimethylammonium bromide (CTAB), and silica nanospheres as starting materials, followed by removing the silica using hydrofluoric acid. The samples were characterized by X-ray diffraction (XRD), field emission scanning electron microscopy (FE-SEM), transmission electron microscopy (TEM), thermos-gravimetric analysis (TGA), Nitrogen adsorption measurements (Brunauer-Emmett-Teller, BET), Energy Dispersive X-Ray Spectroscopy (EDX) and Raman spectrum analysis. Charge-discharge voltage profiles, cyclic voltammograms (CVs) and electrochemical impedence spectroscopy (EIS) were collected, and the cycling and rate performances of the samples were measured to evaluate their electrochemical performance. The underlying reasons for the different electrochemical performances are discussed in detail. Meanwhile, the possible growth mechanism of the samples was also probed. 2. Experiment section MoS2/C composites were prepared by a simple thermal reduction method. Firstly, SiO2 nanospheres (20 nm, >99 wt %) as the templates were dispersed in 30 ml ethanol. Then, CTAB and ATTM were successively added to the solution under ultrasonic vibration. Next, the remaining rufous powders were collected, dried in a vacuum oven at 60  C for 12 h, and then annealed at

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700  C for 4 h in an atmosphere of H2/Ar, flowing at 200 sccm (where sccm is standard cubic centimeters per minute). Finally, the silica was etched from the resultant black powder by hydrofluoric acid and dried in a vacuum oven at 80  C for 12 h. The mass of each starting material for the different samples is shown in Table 1 in detail. X-ray diffraction (XRD) was used to identify the phases of the samples. The patterns were obtained with a Rigaku Rint 220 V powder diffractometer with graphite monochromatized Cu Ka radiation (l = 0.15405 nm) in the 2u range of 10 –80 at a scanning rate of 0.15 s1. A field emission scanning electron microscope (FESEM, FEI, Sirion-200) was used to study the surface morphology of the samples. The powder samples for FE-SEM were attached to carbon conductive adhesive, and images were collected under the best operating conditions. A transmission electron microscope (TEM, FEI, G220) was used to further analyze the microstructure of the samples. A well-dispersed solution was prepared by adding a small amount of sample to absolute ethanol and sonicating for 10– 15 min. One or two drops of the dispersed solution were placed on the TEM grid for measurement. Thermal behaviours of composites were studied by thermo-gravimetric analysis (TGA, STA449 F3) from room temperature to 1000  C at a heating rate of 10  C min1 in air. The Brunauer–Emmett–Teller (BET) surface areas were measured by Nitrogen adsorption isotherms at 77 K, using an VSorb 2800P analyzer (Gold APP Instruments), and samples were degassed at 393 K for 12 h before measurements. Raman spectra of MD-S0.1-C0.05, MD-C0.05 and MD composite samples were obtained at room temperature with The Laser Raman Micro-spectroprobe from Thermo Fisher with the laser excitation of 780 nm. The electrochemical tests were carried out in CR2032 coin cells, which were assembled in an argon-filled glove box. Lithium foil was used as the counter electrode, and 1.0 M LiPF6 in ethylene carbonate/dimethyl carbonate (EC/DMC, 1:1 by volume) as the electrolyte. A polypropylene filter was used as separator. The working electrodes were prepared from a slurry of 80 wt% active materials, 10 wt% carbon black, and 10 wt% polyvinylidene fluoride (PVDF) with N-methyl-2-pyrrolidone. The slurry was spread on copper foil substrates, which were then dried at 120  C for 12 h under vacuum before fabricating the cells. 3. Results and discussion 3.1. Structure and morphology Fig. 1(a) shows the XRD patterns of the samples. All the samples display peaks which belong to MoS2 (JCPDS 37-1492). The diffraction intensity of the (002) peaks for the four samples is very different, however. Without addition of CTAB or SiO2 nanospheres, the MD sample shows a strong (002) diffraction peak, but with the addition of CTAB or SiO2 nanospheres to the starting materials, the intensity of the (002) peak decreases or even disappears (MD-S0.1-C0.05, MD-S0.1, and MD-C0.05 in Fig. 1(a)). This suggests that the presence of CTAB or SiO2 hinders the stacking of single layers of MoS2 [17]. Nevertheless, the decrease in the (002) diffraction peak of MD-S0.1 is not so obvious as MD-S0.1-C0.05 and MD-C0.05, meaning that the hindering effect of SiO2 is weaker than

Table 1 Mass of each starting material. Sample

Weight of SiO2 (g) Weight of CTAB (g) Weight of (NH4)2MoS4 (g) Physically added carbon additives Total carbon content (wt%) measured by TGA

MD-S0.1C0.05 MD-S0.1 MD-C0.05 MD

0.100

0.050

0.250

16.3%

22.3%

0.100 0 0

0 0.050 0

0.250 0.250 0.250

0 16.3% 0

0 25.7% 0

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Fig. 1. (a) XRD patterns, (b) Comparison of XRD before and after acid treatment in the MD-S0.1-C0.05 sample, (c) Composition variation before and after acid treatment in the MD-S0.1-C0.05 sample, (d) TGA curve, (e) Raman spectrum of MD-S0.1-C0.05 and MD-C0.05, (f) N2 adsorption-desorption isotherms of MD-S0.1-C0.05 samples. Inset, pore size distribution.

that of CTAB. To confirm there is no any Si-based material remained in the active electrode phase MD-S0.1-C0.05 and whether there are some effects on the active electrode phase and changes after the HF acid treatment, we compared the phase and the composition before and after acid treatment in the sample MD-S0.1-C0.05 by XRD pattern and EDX spectrum (Fig. 1(b) and (c)), it can be seen that there is no Si-based or amorphous silica phase in the XRD pattern

(the characteristic peak of SiO2 phase is in 22.5 degree), also there are only C, Mo and S three elements in the EDX spectrum of sample MD-S0.1-C0.05 after acid treatment, therefore it can be deduced that these were completely free from any Si based materials in the active electrode materials MD-S0.1-C0.05 after acid treatment. Also, from Fig. 1(b) and (c), we can only see the characteristic peak of SiO2 and Si-based element disappeared in the XRD pattern and EDX

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spectrum, there are no any negative effects of this acid treatment on the active electrode phase and changes. The remained carbon content of MD-S0.1-C0.05 and MD-C0.05 samples was determined by thermo-gravimetric analysis (TGA) shown in Fig. 1(d). The TGA curves reveal a first weight loss between 270  C to 500  C are caused by the combustion of amorphous carbon and MoS2 in air according to the following reactions [18,19]: 2MoS2 þ 7O2 ! 2MoO3 þ 4SO2

ð1Þ

The second weight loss from 650  C to 1000  C is attributed to the sublimation of MoO3. According to the weight loss of MoO3 for different samples (42.9% and 48.9% for MD-S0.1-C0.05 and MD-C0.05, respectively.), by reaction (1), we can calculate the MoS2 loss content in the first step (47.7% and 54.3% for MD-S0.1-C0.05 and MDC0.05), then deduce the carbon content for different samples (22.3% and 25.7% for MD-S0.1-C0.05 and MD-C0.05), which are a little higher than the theoretically added carbon additives (both are 16.3%), and the higher carbon weight loss in TGA can be mainly attributed to the combustion of large amount of organic adsorbed during the synthesis and the accuracy limit of TGA method, more accurate carbon content may need with the aid of more precise method. The total carbon content of all four samples were summarized in Table 1. To further explore the carbon existence, Raman spectra of the MD-S0.1-C0.05, MD-C0.05 and MD were conducted as shown in Fig. 1(e). The two sharp peaks at about 381 cm1 and 410 cm1 in Fig. 1(e) are characteristic peaks of MoS2, they are assigned to the E12g vibration mode and A1g vibration mode, respectively. The E12g mode corresponds to the in-layer displacement of Mo and S atoms, whereas, the A1g mode corresponds to the out-of-layer symmetric displacements of atoms along the c-axis [20,21]. The intense broad bands at about 1310 cm1 and 1602 cm1 in Fig. 1(c) are assigned to the disorder (D) and graphene (G) bands of residual carbon in the MD-S0.1-C0.05 and MD-C0.05 composites, respectively. The G band corresponds to one of the E2g modes, while the D band corresponds to defects and disorder in the hexagonal graphitic layers [22,23]. The relative intensity ratio between D and G bands (ID/IG) reveals the graphitic degree of the carbonaceous materials. By comparing

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the Raman spectra of MD-C0.05 and MD-S0.1-C0.05, the slight increase of the ID/IG from 1.02 to 1.04 reflects a more defective structure in MD-S0.1-C0.05 composite. It might be due to its porous architecture for MD-S0.1-C0.05 composite caused by the introduction of SiO2 template [24]. Nitrogen adsorption (Brunauer-Emmett-Teller, BET) measurements (Fig. 1(f)) were explored to investigate the porosity of MDS0.1-C0.05 samples. The porosity structure of samples was calculated from BET and BJH analyses. It reveals the pore size distributions are in range of 10 nm and 30 nm for MD-S0.1-C0.05 samples, and the specific surface area of MD-S0.1-C0.05 are 17.88 m2/ g. The porous structure will be benefit for its electrochemistry properties. The FE-SEM images of the samples are shown in Fig. 2. As shown in Fig. 2(d), the morphology of MD is that of big blocks with rough surfaces. Compared with MD, the morphology of MD-S0.1C0.05 is very different. It displays a sheet-like structure. This special morphology increases the specific surface area proved by BET results, which is beneficial to the insertion of more lithium ions into the layered structure of MoS2. The morphology of MD-S0.1 shown in Fig. 2(b) is somewhat like that of MD-S0.1-C0.05, but the sheet structure is not obvious, which gives it a more flocculent structured look. Meanwhile, bulk-like structure can also be observed in the top left corner and top right corner of Fig. 2(b), which indicates that the morphology of MD-S0.1 is inhomogeneous. MD-C0.05 also presents a bulk-like structure, but the blocks are smaller and smoother than that in the MD. Therefore, we can deduce that the SiO2 nanosphere templates play an important role in the transformation from bulk-like to sheet-like morphology of MoS2, which then increases the contact area and density of active absorption-desorption centers in the electrode for lithium ions in the electrolyte, and as a consequence, it is advantageous to the deintercalation lithium reaction. To further explore the microstructure, the samples were characterized by TEM, as shown in Fig. 3. As shown in the inset images, the samples all display layered crystal structure. The interlayer spacing of MD-S0.1-C0.05 is about 1.03 nm, while the interlayer spacing of MD-S0.1, MD-C0.05, and MD are all 0.62 nm. By comparison of the four samples, we can find that the stacking of

Fig. 2. FE-SEM images of (a) MD-S0.1-C0.05, (b) MD-S0.1, (c) MD-C0.05, (d) MD, with the insets showing higher magnification.

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Fig. 3. TEM images of (a) MD-S0.1-C0.05, (b) MD-S0.1, (c) MD-C0.05, and (d) MD. The insets are enlarged lattice resolved images of the areas marked by the rectangles. (e) HR-TEM micrographs of starting materials of silica nanoparticles.

single layers in MD-S0.1 and MD is much more ordered than that in MD-S0.1-C0.05 and MD-C0.05. These results are in good agreement with the diffraction intensity variations of the (002) peaks of the four samples in the XRD patterns shown in Fig. 1. The TEM of nanosphere-templated SiO2 was also shown in Fig. 3(e), and the diameter of silica nanoparticles was around 10–15 nm, which was consistent with the porous structure in Figs. 1(d) and 2(a). 3.2. Electrochemical performance Fig. 4 exhibits the discharge/charge potential profiles of the samples during the first, second, and third cycles at 100 mA g1 between 0.01 and 3 V. As shown in Fig. 4, the profiles of MD-S0.1C0.05 and MD-C0.05 are similar to each other, while the profiles of MD-S0.1 and MD are also similar to each other. Fig. 4(b) and (d) shows that MD-S0.1 and MD exhibit two conspicuous potential plateaus at around 1.1 and 0.6 V in the first discharge, which are

regarded as representing the formation of LixMoS2 and the conversion reaction in which Mo particles are embedded in the Li2S matrix [25], respectively. In the 2nd and 3rd discharge, a new potential plateau appears at around 1.9 V, which may correspond to the reduction of sulfur to polysulfide and then the change to Li2S [26]. In the charge process, MD-S0.1 and MD both exhibit a conspicuous plateau at about 2.2 V, which indicates that Li2S is oxidized into sulfur. As shown in Fig. 4(a) and (c), samples MD-S0.1C0.05 and MD-C0.05, on the other hand, do not exhibit any obvious plateau in the first discharge process. In the 2nd and 3rd discharge, however, a new potential plateau at around 1.9 V does appear, which is similar to those in MD-S0.1 and MD, but not as obvious as for MD-S0.1 and MD. Furthermore, in the charge process, MD-S0.1C0.05 and MD-C0.05 exhibit plateaus at about 2.2 and 1.75 V, which indicate that Li2S is oxidized into sulfur. Based on the above analysis, the samples show inconspicuous plateaus when CTAB was added into the starting materials. It is proposed that CTAB is

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Fig. 4. The initial three charge/discharge potential profiles of (a) MD-S0.1-C0.05, (b) MD-S0.1, (c) MD-C0.05, and (d) MD at 100 mA g1 between 0.01 and 3 V.

Fig. 5. Cycling performances of MD-S0.1-C0.05, MD-S0.1, MD-C0.05, and MD samples at a current density of 100 mA g1 in the potential range of 0.01-3.0 V.

transformed into amorphous carbon during the thermal reduction process, and the presence of amorphous carbon in the samples affects the discharge/charge process. Fig. 5 shows the cycling performances of the samples at a constant current density of 100 mA g1. The irreversible capacity

loss after the first cycle is mainly due to the electrolyte decomposition and formation of a solid electrolyte interphase (SEI) layer. At the end of 100 cycles, the capacities of the four samples are 982, 114, 432, and 37 mAh g1 for MD-S0.1-C0.05, MDS0.1, MD-C0.05, and MD, respectively, which is equivalent to capacity

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retention of 91%, 20%, 67%, and 4% (with respect to the second cycle). In comparison with MD, both the MD-S0.1-C0.05 and the MDC0.05 samples display higher specific capacity and better cycling stability. MD-S0.1, however, shows high specific capacity and good cycling stability in the first 40 cycles, but after that, the capacity decreases quickly. These results reveal that only the addition of both CTAB and SiO2 nanospheres (MD-S0.1-C0.05) will simultaneously improve both the capacity and the cycling stability. The electrochemical performances of MD-S0.1-C0.05 and MDC0.05 are much better than those of MD-S0.1 and MD, because the CTAB in the former samples is transformed to amorphous carbon during the thermal reduction process, increasing the electronic conductivity of the materials. While amorphous carbon itself has a certain lithium storage capacity and some flexibility during the charge/discharge process, the volume expansion is small, which leads to improve the cycling stability. The rate capability of four samples are shown in Fig. 6. After a gradually increasing current density from 100 mA g1 to 200 mA g1, 400 mA g1, 600 mA g1, and 800 mA g1, and finally to 1000 mA g1, MD-S0.1-C0.05 still maintain a high capacity of about 560 mAh g1. Furthermore, its reversible capacity can quickly recover when the current density changes from 1000 mA g1 back down to 100 mA g1, and it displays the same stable structure after repeated lithium insertion. From Figs. 5 and 6, we can find when MD-S0.1-C0.05 used as anode materials, it has the highest specific capacity and cycling performance, also most excellent rate capacities compared with other three samples MD-C0.05, MD-S0.1 and MD, which attributes to the porous structure in MD-S0.1-C0.05 samples confirmed by BET and SEM results, also improvement the electronic conductivity by the amorphous carbon. The special

porous-structure morphology in MD-S0.1-C0.05 has a large specific surface area and increases the contact area and density of active absorption-desorption centers in the electrode for lithium ions in electrolyte, so it can promote the de-intercalation lithium reaction. In the meanwhile, the extended interlayer spacing in MoS2 will relieve the volume expansion during the charge and discharge process. Fig. 7 shows the initial three CV curves of the samples MD and MD-S0.1-C0.05. In the first cathodic scan, two reduction peaks are clearly observed, namely: 0.64 and 0.18 V for the samples MD, and 0.55 and 0.20 V for samples MD-S0.1-C0.05. The plateau at 0.55– 0.64 V can be assigned to the formation of LixMoS2, leading to a phase transformation of MoS2 from trigonal prismatic (2H) to octahedral (1T) structure [27]. The other plateau at 0.18–0.20 V corresponds to the conversion reaction in which MoS2 decomposes into Mo particles that embedded in a Li2S matrix. In the subsequent lithiation cycles, the two peaks at 0.55–0.64 and 0.18–0.20 V disappear. Instead, two new peaks arise at about 1.70 and 1.00 V, indicating a multistep lithium insertion mechanism based on the following conversion reactions (reactions (2)–(4)): MoS2 + xLi+ + xe ! LixMoS2

(2)

MoS2 + 4Li+ ! Mo + 2Li2S

(3)

LixMoS2 + (4-x) e + (4-x) Li+ ! Mo + 2Li2S

(4)

For the samples MD, the shifts in the reduction potentials and the changes in the reduction peak shape after the first cycle might

Fig. 6. Rate capability of MD-S0.1-C0.05, MD-S0.1, MD-C0.05, and MD samples at different charge/discharge current densities.

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peaks at 2.34 V can still maintain, suggesting good stability and reversibility. In order to better understand the excellent electrochemical performance of MD-S0.1-C0.05 samples in comparison with other samples, the electrochemical impedance spectroscopy (EIS) was performed to study the electrochemical reaction kinetics of the MD-S0.1, MD-C0.05 and MD samples after ten cycles, as shown in Fig. 8. It can be seen that, for MD-S0.1-C0.05 samples, the diameter of the semicircle in the high/medium frequency is the smallest among three electrodes, implying that MD-S0.1-C0.05 samples possesses the lowest contact resistance and charge-transfer resistance for the insertion/extraction reactions [27,28]. It can be attributed to the incorporation of conductive carbon matrix and the unique nanostructure. Besides, the straight line in the low frequency region in the Nyquist plots is indicative of Warburg behavior arising from the diffusion of lithium ions in the electrode [29–31]. The MD-S0.1-C0.05 samples shows a steeper straight line in the low frequency region than those of MD-S0.1, MD-C0.05, and MD samples, indicating a lower diffusion impedance. The lithium ion diffusion coefficient in these electrodes can be calculated from the following formula [32]: DLiþ ¼

Fig. 7. CV profiles of (a) MD and (b) MD-S0.1-C0.05 samples at a scan rate of 0.5 mV s1.

be attributed to its poor structural stability with well-stacked layered architecture. In the first anodic scan, a pronounced peak at 2.34–2.44 V can be clearly identified, which is ascribed to the oxidation of Li2S to sulfur and lithium ions. In the subsequent cycles, the oxidation peak at 2.44 V for the samples MD becomes much weaker and exhibits a shift toward the high voltage, implying the poor structure stability of samples MD. As for samples MD-S0.1-C0.05, the prominent oxidation peaks at 2.34 V coupled with several small oxidation peaks at 1.1–1.8 V resulting from the incomplete oxidation of Mo particles are visible in the first and second anodic scan. In the successive cycles, these weak oxidation peaks at 1.1–1.8 V disappear, while its CV shape and the oxidation

   2 1 RT @v 2 An2 CF 2 @Z 0 1 2

ð5Þ

where DLi+ is the lithium ion diffusion coefficient (cm2 S1), A is the surface area of the electrode (1.13  104 m2), n is the number of the electrons per MoS2 molecule attending the electronic transfer reaction (n is 4), C is the molar concentration of Li+ in electrolyte (1.20  104 mol m3), F is the Faraday constant (96486C mol1), R is the gas constant (8.314 J mol1 K1), T is the absolute temperature (298 K), Z' is the real Warburg resistance and v is the angular frequency in low frequency region. The lithium ion diffusion coefficient DLi+ of MD-S0.1-C0.05, MD-S0.1, MD-C0.05, and MD samples calculated by Eq. (5) from the slope of @v1/2/@Z' (inset of Fig. 9) were 6.94  1016, 6.48  1017, 1.83  1016 and 1.57  1017 cm2 S1, respectively. Obviously, the MD-S0.1-C0.05 samples exhibited the largest lithium ion diffusion coefficient after 10 cycles among these four electrodes. It further proves that the sheet-like structure assembled structure of MD-S0.1-C0.05 samples not only effectively provides more active sites, shortens

Fig. 8. Nyquist plots (the inset graph of Z' plotted against v1/2 at the low frequency) of MD-S0.1-C0.05, MD-S0.1, MD-C0.05, and MD samples.

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Fig. 9. Schematic illustration of the growth mechanisms of (a) MD-C0.05; (b) MD-S0.1; and (c) MD-S0.1-C0.05.

the diffusion paths of lithium ions and electrolyte, but also facilitates the easy access of the lithium ions. As a result, MD-S0.1C0.05 samples exhibits greatly improved lithium ion kinetics in the electrode materials and enhanced rate capability when compared to MD-S0.1, MD-C0.05 and MD samples. 4. Growth mechanism Based on the above results, the possible growth mechanisms for the MD-S0.1-C0.05, MD-S0.1, and MD-C0.05 samples are proposed as illustrated in Fig. 9. As a cationic surfactant, CTAB can adsorb some of the MoS42 ions around it due to the ionic interactions. CTAB molecules are initially dispersed in the solution, as shown in Fig. 9(a). Due to stirring the solution and ethanol evaporation, the concentration of CTAB increases and gradually forms globular micelles, which then change to rod-like micelles, tabular micelles, and finally lamellar micelle [33]. During this process, the distribution of MoS42 ions is also changing accordingly. After the solvent has evaporated completely, (NH4)2MoS4 displays a bulk-like morphology, but it is smaller than for the (NH4)2MoS4 initial bulk. During the thermal reduction process, (NH4)2MoS4 reacts as follows [34], ðNH4 Þ2 MoS4 þ H2 ðgÞ ! 2NH3 ðgÞ þ 2H2 SðgÞ þ MoS2

ð6Þ

CTAB is transformed to amorphous carbon by the thermal reduction process, and (NH4)2MoS4 is reduced to MoS2, while the materials maintain their bulk-like morphology, as in Fig. 2(c). When the SiO2 nanosphere templates are added into the solution, owing to the hydroxyl groups (Si-OH) on their surfaces, the SiO2 nanospheres are similar to negatively charged micelles [35]. Following the addition of (NH4)2MoS4 into the solution, the

SiO2 negatively charged micelles will be mutually repulsive to the same negatively charged MoS42 ions, ensuring that the MoS42 ions are well dispersed, as displayed in Fig. 9(b). When the solvent evaporation is entirely completed, uniformly mixed particles consisting of SiO2 and (NH4)2MoS4 are obtained. After undergoing the thermal reduction process according to Eq. (6), then the etching away of the SiO2 templates by HF, the final product exhibits a flocculent morphology, as shown in Fig. 2(b). The crystal growth mechanism of MD-S0.1-C0.05 is demonstrated in Fig. 9(c). When CTAB and SiO2 are added to the solution simultaneously, CTAB molecules are adsorbed on the SiO2 surfaces, so as to form a thin coating layer. After the addition of (NH4)2MoS4, the MoS42 ions are surrounded by SiO2 adsorbed CTAB. After going through thermal reduction in hydrogen atmosphere as in Reaction (1), some layer-structure-disrupted MoS2/C composites are obtained, also nano-porous structure, owing to the removal of SiO2 nanosphere templates by HF, which can be proved by BET results. By SEM and TEM, we can observe an expanded layer structure in MoS2/C composites, and the intensity of the (002) diffraction peak also becomes weakened, induced by destruction of the layer structure (as indicated by XRD). Therefore, to obtain such nanoporous MoS2/C composites with purpose-designed morphology for high performance lithium ion batteries, nanosphere-templated SiO2 and CTAB play complementary crucial roles. Less-ordered and expanded layer-structure is of benefit for high lithium storage capacity with high cycling stability. 5. Conclusion In summary, MoS2/C composites and MoS2 with different morphology were prepared by a simple thermal reduction method.

Z. Wang et al. / Electrochimica Acta 239 (2017) 74–83

The characterization of microstructures by SEM and TEM indicated that the MD and MD-C0.05 samples had a bulk-like structure, MDS0.1 had a flocculent structure, and MD-S0.1-C0.05 had a sheet-like structure. The morphology observations also demonstrated that the stacking of single-layer MoS2 in MD-S0.1 and MD is much more ordered than in MD-S0.1-C0.05 and MD-C0.05. During the preparation process, CTAB and SiO2 played the critical role of preventing MoS2 from stacking in the (002) direction. Electrochemical properties measurements showed the higher specific capacity and excellent cycling stability of the MD-C0.05-S0.1 composite synthesized by adding 0.05 g CTAB and 0.1 g silica nanospheres, with a high capacity of about 982 mAh g1 after 100 cycles. The excellent electrochemical performance of MD-S0.1-C0.05 could be attributed to its special morphology and the increased interlayer spacing of MoS2 for easy lithium insertion. This study has demonstrated that the morphology plays an important role in the electrochemical performance of materials and that MD-C0.05S0.1 is a promising material for lithium ion batteries. Acknowledgements This work is financially sponsored by Natural Science Foundation of China (Grants No. 51002029). References [1] C.K. Chan, H. Peng, G. Liu, K. McIlwrath, X.F. Zhang, R.A. Huggins, Y. Cui, HighPerformance Lithium Battery Anodes Using Silicon Nanowires, Nat. Nanotechnol. 3 (2008) 31–35. [2] T.H. Han, W.J. Lee, D.H. Lee, J.E. Kim, E.Y. Choi, S.O. Kim, Peptide/Graphene Hybrid Assembly into Core/Shell Nanowires, Adv. Mater. 22 (2010) 2060–2064. [3] A.J. Manthiram, Materials Challenges and Opportunities of Lithium Ion Batteries, Phys. Chem. Lett. 2 (2011) 176–184. [4] I. Kovalenko, B. Zdyrko, A. Magasinski, B. Hertzberg, Z. Milicev, R. Burtovyy, I. Luzinov, G. Yushin, A Major Constituent of Brown Algae for Use in HighCapacity Li-ion Batteries, Sci 334 (2011) 75–79. [5] L. Lin, Z. Ji, M. Alcoutlabi, X. Zhang, Recent Developments in Nanostructured Anode Materials for Rechargeable Lithium-ion Batteries, Energy Environ. Sci. 4 (2011) 2682–2699. [6] J.M. Tarascon, M. Armand, Issues and Challenges Facing Re-chargeable Lithium Batteries, Nat. 414 (2001) 359–367. [7] H.R. Matte, A. Gomathi, A.K. Manna, D.J. Late, R. Datta, S.K. Pati, C. Rao, MoS2 and WS2 Analogues of Graphene, Angew. Chem. 122 (2010) 4153. [8] M. Wang, G. Li, H. Xu, Y. Qian, J. Yang, Enhanced Lithium Storage Performances of Hierarchical Hollow MoS2 Nanoparticles Assembled from Nanosheets, ACS Appl. Mat. Interfaces. 5 (2013) 1003–1008. [9] C. Zhang, H.B. Wu, Z. Guo, X.W. Lou, Facile Synthesis of Carbon-Coated MoS2 Nanorods with Enhanced Lithium Storage Properties, Electrochem. Commun. 20 (2012) 7–10. [10] C. Zhang, H.B. Wu, Z. Guo, X.W. Lou, Synthesis of MoS2-C One-Dimensional Nanostructures with Improved Lithium Storage Properties, ACS Appl. Mater. Interfaces 4 (2012) 3765–3768. [11] F. Xiong, Z. Cai, L. Qu, Three-Dimensional Crumpled Reduced Graphene Oxide/ MoS2 Nanoflowers: A Stable Anode for Lithium-Ion Batteries, ACS Appl. Mat. Interfaces. 7 (23) (2015) 12625–12630. [12] L. Hu, Y. Ren, H. Yang, Fabrication of 3D Hierarchical MoS2/Polyaniline and MoS2/C Architectures for Lithium-Ion Battery Applications, ACS Appl. Mat. Interfaces. 6 (16) (2014) 14644–14652.

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