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MoS2 Nanosheets Hosted in Polydopamine-Derived Mesoporous Carbon Nanofibers as Lithium-Ion Battery Anodes: Enhanced MoS2 Capacity Utilization and Underlying Mechanism Junhua Kong, Chenyang Zhao, Yuefan Wei, and Xuehong Lu ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.5b07950 • Publication Date (Web): 13 Oct 2015 Downloaded from http://pubs.acs.org on October 14, 2015

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ACS Applied Materials & Interfaces

MoS2 Nanosheets Hosted in Polydopamine-Derived Mesoporous Carbon Nanofibers as Lithium-Ion Battery Anodes: Enhanced MoS2 Capacity Utilization and Underlying Mechanism Junhua Kong,a Chenyang Zhao,a Yuefan Wei,b Xuehong Lua,* a

School of Materials Science and Engineering, Nanyang Technological University, 50 Nanyang

Avenue, Singapore 639798. b

School of Mechanical and Aerospace Engineering, Nanyang Technological University, 50

Nanyang Avenue, Singapore 639798 KEYWORDS: MoS2; polydopamine; electrospinning; carbon nanofibers; lithium ion batteries.

ABSTRACT: In this work, solid, hollow and porous carbon nanofibers (SNFs, HNFs and PNFs) were used as hosts to grow MoS2 nanosheets hydrothermally. The results show that the nanosheets on the surface of SNFs and HNFs are comprised of a few grains stacked together, giving direct carbon-MoS2 contact for the first grain and indirect contact for the rest. Differently, the nanosheets inside of PNFs are of single-grain size and distributed evenly in the mesopores of

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PNFs, providing efficient MoS2-carbon contact. Furthermore, the nanosheets grown on polydopamine-derived carbon surface of HNFs and PNFs have larger interlayer spacing than that grown on polyacrylonitrile-derived carbon surface. As a result, the MoS2 nanosheets in PNFs possess the lowest charge transfer resistance, the most accessible active sites for lithiation/delithiation and can effectively buffer the volume variation of MoS2, leading to its best electrochemical performance as an lithium-ion battery anode among the three. The normalized reversible capacity of the MoS2 nanosheets in PNFs is about 1210 mAh g-1 at 100 mA g-1, showing the effective utilization of the electrochemical activity of MoS2.

INTRODUCTION Molybdenum disulfide (MoS2) is an inorganic compound with graphite-like structure, in which one layer of Mo atoms are covalently bonded to two layers of S atoms, and the S-Mo-S layers stack together by weak van der Waals forces. MoS2 has the potential for a wide range of applications such as solid lubricants,1 photoelectrodes,2-4 supercapacitors,5 and especially, electrodes of lithium-ion batteries (LIBs).6 The intercalation of lithium into MoS2 was reported decades ago.7 However, only in recent years have the electrochemical properties of MoS2 been intensively explored for LIB applications. Research shows that the structure and morphology of MoS2, such as interlayer spacing, grain size, stacking order and thickness, play crucial roles in influencing its lithium-ion intercalation capability and hence ultimate electrochemical properties as LIB electrodes. For practical applications, the difficulty in achieving desired structure and morphology of MoS2 as well as its poor electrical conductivity remains as the major obstacles for its efficient transport of charges and thus full lithiation/de-lithiation.


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To enhance the electrochemical properties of MoS2, many strategies have been formulated. The most popular ones include the fabrication of nano-sized MoS2 of various geometries,8-15 exfoliation or intercalation of MoS2 to achieve expanded interlayer spacing,16-20 and incorporation of MoS2 into carbon nanostructures.17-18, 21-42 These approaches rely on either tailoring the morphology of MoS2 or improving its charge transport capability through introducing electrically conducting agents such as carbon nanostructures43-44 to realize good electrochemical performance of MoS2. For the conducting agents, one-dimensional (1D) carbon nanostructures have been proposed as favourable hosts for MoS2. For instance, MoS2/C composite nanorods and nanotubes in powder form were obtained through preparation of MoOx/polyaniline hybrid nanostructures via multi-step chemical reactions, followed by a sulfidation reaction in H2S flow.38 It was demonstrated that the nanotubular composite possessed better cycling and rate capacity than the nanorod-supported one owing to the presence of extra interior cavities in the nanotubes, which offers larger interfacial area for interaction with active sites, facilitating lithium-ion storage.

In this article, for the first time we report the hydrothermal growth of MoS2 nanosheets in mesoporous carbon nanofibers, which allows simultaneous manipulation of both the structure and morphology of MoS2 as well as enhances the interactions between MoS2 and the conductive phase. By studying three types of 1D nanostructures, namely solid, hollow and mesoporous carbon nanofibers, as the hosts, the effects of the carbon nanofibrous hosts on the structure and morphology of the in-situ grown MoS2 are illustrated. The electrochemical properties of the three


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nanocomposites are correlated to the structure and morphology of MoS2 as well as their interactions with the 1D carbon nanostructures, revealing the underlying mechanism for the excellent electrochemical performance of the MoS2 nanosheets in mesoporous carbon nanofibers.

EXPERIMENTAL SECTION Materials. Polyacrylonitrile (PAN, Mw = 150000) was purchased from Polymer Internationals (Germany). Polystyrene (PS, Mw = 350000), dopamine hydrochloride (DOPA), tris(hydroxymethyl) aminomethane (Tris), sodium molybdate dihydrate (Na2MoO4·2H2O) and thioacetamide (TAA, CH3CSNH2) were purchased from Sigma-Aldrich Chemistry (USA). Dimethylformamide (DMF) was purchased from Tedia Company Inc (USA). Ethanol was supplied by Merck KGaA (Germany). All materials are of analytical grade and were used without further purification.

Preparation of solid, hollow and porous nanofibers. Solid PAN nanofibers, solid and porous PS nanofibers were firstly prepared via electrospinning of PAN/DMF solution (10 wt%) and PS/DMF solution (15 wt%). The applied voltage and feeding rate for the solid PAN nanofibers were 13 kv and 0.5 mL h-1, respectively. The nanofibers were collected on aluminium foils for 2 hrs. The applied voltage, feeding rate, humidity were 10 kV, 0.6 mL h-1, 25 % and 10 kV, 0.6 mL h-1, 50 % for solid and porous PS nanofibers, respectively. Ethanol was used as liquid collector 45 with collection time of 2 hrs. The solid carbon nanofibers obtained by carbonizing


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the PAN nanofibers are denoted as C-PAN SNFs. The carbonization condition is the same as our earlier report with annealing temperature of 700 °C.46 To obtain hollow and porous carbon nanofibers, the solid and porous PS nanofibers were firstly immersed in 300 g DOPA aqueous solution (0.3 mg mL-1). 0.36 g Tris was then added into the aqueous solution under vigorous shaking. The in situ polymerization of DOPA was carried out for 2 hrs under continuous shaking, forming polydopamine (PDA) coating on the nanofibers. The coating process was repeated for another 4 times using fresh DOPA solution and Tris. The PDA-coated solid and porous PS nanofibers were dried and then annealed under the following conditions to remove PS and carbonize PDA: heating from room temperature to 700 °C at heating rate of 5 °C/min in argon atmosphere, and then keeping at 700 °C for 3 hrs. The obtained hollow and porous carbon nanofibers are denoted as C-PDA HNFs and C-PDA PNFs, corresponding to that from solid and porous PS nanofibers, respectively.

Preparation of MoS2 on solid, hollow and porous nanofibers. MoS2 was introduced into the nanofibers via hydrothermal route. In brief, Na2MoO4·2H2O and TAA were firstly dissolved in de-ionized (DI) H2O to obtain precursor solution with Na2MoO4·2H2O and TAA concentration of 0.06 M and 0.13 M, respectively. The above nanofibrous mats were immersed into the precursor solution for 1 hr and then transferred into a Teflon-lined stainless steel autoclave. The hydrothermal growth was carried out at 200 °C for 24 hrs. The treated mats were washed by DI H2O several times, dried at 60 °C in vacuum for 24 hrs and then annealed at 700 °C in argon environment for 3 hrs. The samples prepared from C-PAN SNFs, C-PDA HNFs and C-PDA PNFs are denoted as MS/SNFs, MS/HNFs and MS/PNFs, respectively. For verification purpose, PDA was also coated onto PAN solid nanofibers and annealed under the same condition as that


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for C-PDA HNFs. MoS2 was then hydrothermally attached onto the above C-PDA-coated SNFs under the same condition as that for MS/SNFs. The obtained sample was denoted as MS/C-PDAcoated SNFs.

Characterization. The morphology of the samples was studied using a field-emission scanning electron microscope (FESEM, JEOL 7600F), a high-resolution transmission electron microscope (HRTEM, JEOL 2100F). For cross-section observation, the mats were embedded into epoxy, cut into ultrathin slides using ultramicrotome and attached on copper grid. Scanning transmission electron microscope-energy dispersive X-ray (STEM-EDX) elemental mapping analysis was also carried out. The structure and composition of the samples were measured using an X-ray diffractometer (XRD, Bruker D8 Discover GADDS), a thermogravimetric analyzer (TGA, Q500) and an X-ray photoelectron spectroscopy (XPS, Theta-probe, Thermo Scientic). The Brunauer-Emmett-Teller (BET) specific surface area and Barrett-Joyner-Halenda (BJH) pore size distribution of the samples were measured using a Micromeritics Tristar II-3020 nitrogen adsorption apparatus. For electrochemical characterization, the free-standing mats were used as anodes directly without using any binder and conducting agent, and assembled into 2032 coin cells in an argon-filled glove box and tested as LIB anodes. Lithium flake, Celgard 2325 membrane (USA) and 1 M LiPF6 in the mixture of ethylene carbonate and dimethyl carbonate with 1/1 volumetric ratio were used as counter electrode, separator and electrolyte, respectively. The assembled batteries were connected to a battery test system (4200, MACCOR) for electrochemical performance evaluation and an Autolab potentiostat for cyclic voltammetry (CV) and electrochemical impedance spectroscopy (EIS) test. The CV measurement was


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performed at 0.05 mV s-1 within 3.0~0.005 V. The EIS measurement was carried out in a frequency range of 106 to 10-2 Hz and at the a.c. amplitude of 5 mV.

RESULTS AND DISCUSSION Morphologies and structures of carbonaceous 1D nanostructures. In this work, three types of carbonaceous 1D nanostructures, namely, solid, hollow and porous carbon nanofibers, are used as carbon hosts for MoS2 nanosheets. PAN was electrospun into bead-less nanofibers with an average diameter of 650 nm (Figure 1a) and then carbonized into solid carbon nanofibers (CPAN SNFs, Figure 1b) via annealing in argon environment. The average diameter of C-PAN CNFs is about 550 nm. To obtain hollow and porous carbon nanofibers, solid and porous electrospun PS nanofibers, as shown in Figure 1c and 1e, were firstly obtained by adjusting environmental humidity to about 25 % and 50 %, respectively. Since water vapor is miscible with DMF, DMF can easily absorb moisture to become a nonsolvent, leading to liquid-liquid phase separation before solidification of PS nanofibers and thus creating pores in the nanofibers.47 Moreover, the mesopores in the porous PS nanofibers are interconnected, forming nanochannels (Figure 1e2). A thin PDA layer was then coated onto the solid and porous PS nanofibers via in situ polymerization. For the solid PS nanofibers, the PDA was deposited on the surface of PS, forming a continuous shell. During annealing, PS core was removed, while PDA was converted to carbonized PDA (C-PDA), forming hollow nanofibers (C-PDA HNFs) with an average diameter of around 1 µm (Figure 1d) and carbonaceous shell of about 50 nm (inset of Figure 1d2). For the porous PS nanofibers, the DOPA molecules dissolved in Tris buffer solution can diffuse into the nanochannels, triggering the in-situ coating both on the outer surface and in the nanochannels of the nanofibers. Porous C-PDA nanofibers (C-PDA PNFs) with


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interpenetrating pores are formed after annealing (Figure 1f). The diameter of the nanofibers is similar to that of C-PDA HNFs. The porous nanofibers are covered by a relatively dense shell with thickness of about 30 nm (inset of Figure 1f2) because the thickness of the PDA coating is larger than the size of the pores on the surface and hence the pores are blocked.48

Figure 1. The morphologies of (a) electrospun PAN nanofibers, (b) PAN-derived solid carbon nanofibers (C-PAN SNFs), (c1 and c2) electrospun solid PS nanofibers, (d1 and d2) C-PDA


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hollow nanofibers (C-PDA HNFs), (e1 and e2) electrospun porous PS nanofibers and (f1 and f2) C-PDA porous nanofibers (C-PDA PNFs). The insets of c2 and e2 indicate the shell thickness of the nanofibers.

Micropores, normally much smaller than 1 nm, are formed during the carbonization of PAN solid nanofibers due to the removal of heteroatoms (H, O and N).49 The presence of micropores, however, does not give large specific surface area (SSA) owing to the large diffusion resistance given by the micropores. Compared with C-PAN SNFs, the hollow geometry of C-PDA HNFs provides additional inner surface, while a large quantity of interpenetrating mesopores in C-PDA PNFs gives even larger accessible surface area. As a result, C-PDA HNFs exhibit larger SSA than C-PAN SNFs, and C-PDA PNFs show the largest SSA among the three (Figure S1a). Large SSA would offer more sites for attachment of MoS2. In addition, despite that both C-PAN and CPDA are electrically conductive, their chemical structures are different. It has been reported that C-PAN is amorphous with disordered graphitic domains,46, 50 while C-PDA is highly graphitized with a substantial amount of heteroatoms (N and O).51 The electrochemical activity of the three types of nanofibers will therefore vary owing to their differences in both morphology and structure. This will be elaborated in the following sections.

Morphologies and structures of MoS2 on the carbonaceous 1D nanostructures. MoS2 was introduced into/onto the three types of carbon nanofibers via hydrothermal route. As shown in Figure 2a1, 2b1, 2c1 and the insets, the nanofibrous morphology retains after the hydrothermal growth, whereas the nanofiber surface becomes rough, indicating the successful attachment of


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the hydrothermally grown MoS2. The formation of MoS2 is confirmed by the XRD results (Figure 3a). The peaks at around 26 °, 14 °, 33 °, 40 °, 43 °, 50 ° and 59 ° can be assigned to [002] plane of graphitic carbon, and [002], [100], [103], [006], [105], [110] of MoS2, respectively. This is further confirmed by the XPS results (Figure S2), which verify the existence of MoS2 and carbon that doped with O and/or N. The cross-sectional HRTEM images (Figure 2a2, 2b2 and 2c2) show that the surfaces of all carbon nanofibers are well covered by MoS2 nanosheets. For C-PAN SNFs, the nanosheets densely stack on the nanofiber surface with stack thickness of 20-30 nm. The nanosheets with similar thickness are observed on both the inner and outer surfaces of C-PDA HNFs, whereas the stacking is more irregular than that on C-PAN SNFs. On C-PDA PNFs, besides some large MoS2 nanosheets on the outer shell, a large amount of MoS2 nanosheets with thickness of around 5 nm are uniformly distributed in the interpenetrating mesopores (Figure 2c3). The introduction of MoS2 into the carbon nanofibers, taking C-PDA PNFs as a typical example, decreases the SSA of the hybrids while still maintains their mesoporous nature (Figure S1b1 and 1b2), which is consistent with the above morphological observation. HRTEM images (Figure 2a3, 2b3, 2c4 and 2c5) also show that the [002] interlayer d-spacings are about 0.62, 0.66 and 0.65 nm for the MoS2 nanosheets in MS/SNFs, MS/HNFs and MS/PNFs, respectively. These are close to the interlayer d-spacings calculated from XRD patterns, which are 0.62, 0.63 and 0.63 nm for MS/CNFs, MS/HNFs and MS/PNFs, respectively, and significantly larger than that of commercial MoS2 powder (0.615 nm, Figure 3a). The larger [002] d-spacing of MS/HNFs and MS/PNFs than that of MS/SNFs can be seen by a slight left shift of the diffraction peak from 2θ = 14.2° for MS/SNFs to 2θ = 14.0° for MS/HNFs and MS/PNFs (Figure 3a). This difference is caused by different structures of the two types of carbon


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(C-PAN and C-PDA), which may affect the nucleation and thus the growth of the MoS2 nanosheets. This is verified by the fact that the [002] d-spacing of MoS2 grown on C-PDAcoated SNFs is indeed larger than that directly grown on SNFs (Figure S3); the only difference between these two samples is that the outer surface of the former is composed of C-PDA, instead of C-PAN. The expansion of the interlayer spacing would relieve redox reaction-induced strain and offer more space for lithium-ion intercalation with reduced barriers.52

Figure 2. The morphologies of MoS2 nanosheets in (a1-a3) MS/SNFs, (b1-b3) MS/HNFs and (c1c5) MS/PNFs.


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Figure 3. (a) XRD patterns and (b) TGA curves (heating to 650 °C at 10 °C/min in air for the TGA measurement) of MoS2-on-carbon 1D nanostructures. XRD pattern of MoS2 bulk is also shown in (a) for comparison purpose.


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Using Scherrer equation, the average MoS2 grain sizes in the thickness direction calculated based on the full width at half maximum (FWHM) of [002] diffraction peaks (cf. Supporting Information and Table S1) of MS/SNFs, MS/HNFs and MS/PNFs are 6.7, 5.7 and 6.2 nm, respectively, which are about ten times of the [002] d-spacing, i.e., corresponding to around 10 MoS2 layers in each grain. The MoS2 nanosheet thickness in MS/SNFs and MS/HNFs is 20-30 nm, as presented above, which is 3-5 times larger than the grain size, implying that each nanosheet is comprised of parallel stacks of 3-5 single-crystalline grains (Figure 2a3 and 2b3). Taking MS/SNFs as an example, multiple grains are stacked on the surface of C-PAN (Figure 2a3), resulting in direct contact of the 1st grain with carbon while indirect contact for the rest. Distinctly, the thickness of the MoS2 nanosheets in the pores of MS/PNFs is almost identical to the average grain size, revealing that each nanosheet is a single crystal, which is directly attached on the carbon surface. Indeed, FESEM studies confirm that the thickness of most MoS2 nanosheets located in the pores as well as on the outer surface is close to 5 nm, despite that a small portion of MoS2 nanosheets attached on the outer shell possess a slightly larger thickness (Figure S4). The contents of residue from TGA tests are 60.4, 60.9 and 55.8 wt% for MS/SNFs, MS/HNFs and MS/PNFs, respectively (Figure 3b). On the assumption that MoS2 is completely oxidized to MoO3 and the carbon completely decomposes in air in the TGA tests (Figure S5),53 the MoS2 contents in MS/SNFs, MS/HNFs and MS/PNFs are 67.1, 67.7 and 62.0 wt%, respectively. Their similar composition ensures fair comparison of their electrochemical properties as LIB anodes.

Electrochemical properties. The lithiation/de-lithiation behaviours of the three samples were firstly investigated through cyclic voltammetry, and the curves for the first 3 cycles are shown in


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Figure 4. Typical irreversible alloying and reversible redox between MoS2 and lithium ions are observed.19, 26, 34 In the first cycle, the weak peak at round 1.4-1.7 V, as indicated in Figure 4a24c2, is induced by the reduction of the oxygen/nitrogen-containing functional groups of PAN and PDA-derived carbon.26 This suggests that the carbon is involved in the lithiation process. The cathodic peaks at around 0.75 V and 0.25 V are assigned to the lithiation of MoS2 that forms LixMoS2 followed by its subsequent decomposition into metallic Mo and Li2S. The initial anodic peaks at about 1.2∼1.6 V and 2.4 V correspond to the partial oxidation of Mo that produces MoS2 and the oxidation of LixS to S, respectively (Figure 4a1-4c1). In the following cycles, the cathodic peaks at about 1.8 V and 1.0 V are indicative of the lithiation of S and Mo oxidationderived MoS2, respectively, and the presence of a shoulder in the range of 0.5∼0 V verifies further decomposition of the lithiated MoS2. Comparing the CV curves of the three samples, the better overlap of the 2nd and 3rd cycle anodic curves for MS/PNFs than those for the other two samples indicates that the lithiation/de-lithiation reactions in MS/PNFs stabilize faster, which is probably due to the single-grain nature of the MoS2 nanosheets in MS/PNFs.


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Figure 4. The cyclic voltammetry (CV) curves of (a1) MS/SNFs, (b1) MS/HNFs and (c1) MS/PNFs. (a2), (b2) and (c2) show the corresponding enlarged curves.

The cycled performance of the three samples was evaluated at the current rate of 50 mA g-1 for the first 10 cycles and at 100 mA g-1 thereafter. As shown in Figure 5a, the MS/PNFs delivers much higher cycled capacity, about 950 mAh g-1, than MS/SNFs and MS/HNFs, which stabilize at about 420 mAh g-1 and 540 mAh g-1, respectively. The initial coulombic efficiencies are 79.4 %, 81.8 % and 70 %, respectively, for MS/SNFs, MS/HNFs and MS/PNFs, while all stabilize at


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around 100 % in the following cycles. The fluctuation of coulombic efficiency in MS/SNFs is most likely due to that the contact between MoS2 nanosheets and solid carbon nanofibers is insufficient to maintain the structure stability. Although the capacity of MS/SNFs is higher than that of MS/HNFs in the first 30 cycles, it drops rapidly when the current rate goes to 100 mA g-1, indicating the poor stability of MS/SNFs. By contrast, the cyclability of MS/PNFs and MS/HNFs are much better, and the capacity of MS/PNFs gradually increases, despite that there is slight fluctuation, along with the cycles due to the cycling induced activation of the Li+ pathway and expansion of defect sites that facilitates intercalation of more Li+.13, 18, 54 The above differences are mainly caused by the different structures and morphologies of MoS2 nanosheets in the three samples. As confirmed above, in MS/SNFs and MS/HNFs, thick MoS2 nanosheets with multiple grains are stacked on the surfaces of the solid and hollow carbon nanofibers. Since the hollow geometry provides higher interfacial area, as indicated by the higher SSA of C-PDA HNFs than C-PAN SNFs, MS/HNFs have more accessible reaction sites than MS/SNFs to store lithium ions and could also buffer the volume change of MoS2 during lithiation/de-lithiation, leading to its higher capacity with better stability. In MS/PNFs, most of the MoS2 nanosheets are of singlegrain size and in direct contact with the conductive C-PDA, no matter they are attached on the outer surface or located in the inner nanochannels of C-PDA porous nanofibers (Figure S4). The good structural integration of MS/PNFs gives rise to further enhancement of the cycling stability.


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Figure 5. Cycled discharge and charge capacity and coulombic efficiency of (a) MS/SNFs, MS/HNFs and MS/PNFs and (b) C-PAN SNFs, C-PDA HNFs and C-PDA PNFs. (c) Normalized cycled capacity of MoS2 in MS/SNFs, MS/HNFs and MS/PNFs. (d) Rate capacity of MS/PNFs, C-PDA PNFs and normalized rate capacity of MoS2 nanosheets in MS/PNFs.

In order to analyze the contribution of MoS2 and carbon phases to the overall capacity, the normalized capacity of MoS2 nanosheets was calculated by assuming that the anode capacity is the sum of those of MoS2 and carbon in each cycle. As shown in Figure 5b, C-PAN SNFs, CPDA HNFs and C-PDA PNFs deliver stable capacity of about 360, 420 and 530 mAh g-1,


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respectively, at 100 mA g-1. The differences are due to their different morphologies, graphitic structures as mentioned earlier, as well as N-doping and doping-induced defects that enhance the reversible lithium-ion storage.45, 51, 55 Taking these values as the capacities of the carbon phases, the normalized capacities of MoS2 nanosheets in MS/SNFs, MS/HNFs and MS/PNFs are around 450, 600 and 1210 mAh g-1, respectively, at 100 mA g-1 (Figure 5c). Considering that the theoretical capacity of MoS2 is only 670 mAh g-1 (calculated by taking the weight of both Mo and S into account),6 the above analysis implies that the MoS2 in MS/PNFs has been effectively utilized, while additional physical or chemical absorption may also contribute to lithium-ion storage. By contrast, the utilization rate is only about 90 % and 67 %, respectively, for the MoS2 nanosheets in MS/SNFs and MS/HNFs. It is also worth noting that the normalized capacity of MoS2 nanosheets in MS/PNFs at 100 mA g-1 is comparable to the highest reported value for MoS2-based anodes,37 1290 mAh g-1 at 50 mA g-1, which is the normalized capacity of MoS2 in a MoS2-graphene hybrid without taking the electrochemical activity of carbon black (conducting agent used) into account (there is no carbon black in MS/PNFs anode).

The rate performance of MS/PNFs, C-PDA PNFs and the corresponding normalized rate capacity of MoS2 nanosheets are shown in Figure 5d. At a high current rate of 2000 and 5000 mA g-1, MS/PNFs still possess capacities of around 380 and 170 mAh g-1, respectively, which correspond to normalized capacities of around 500 and 230 mAh g-1 for MoS2 nanosheets. A full rebound is observed upon the current rate is reduced back to 100 mA g-1. This indicates the excellent rate performance of MS/PNFs.


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Mechanism studies. To further investigate the mechanisms behind the excellent electrochemical performance of MS/PNFs, EIS tests were conducted. For all the three samples, a semicircle in high-frequency region (1), two semicircles in moderate-frequency region (2 and 3) and an inclined line in low-frequency region (4) are observed (Figure 6). This is different from the reported EIS results which typically show only a high-frequency semicircle corresponding to resistance Rf and constant phase element of SEI film (CPE1), a moderate-frequency semicircle correlated to charge transfer resistance Rct and constant phase element of electrode-electrolyte interface (CPE2), and a low-frequency inclined line for Warburg impedance arising from the diffusion of lithium ions in the electrode.19, 35-36 It is believed that the two moderate-frequency semicircles in this case are likely to be associated with the charge transfer resistances (Rct1 and Rct2) and corresponding constant phase elements (CPE2 and CPE3) of multiple interfaces, including that between the electrolyte and carbon (Rct1) as well as the one between the electrolyte and MoS2 (Rct2) because both MoS2 and carbon are involved in the lithiation/de-lithiation process, as proven above. The 1st semicircles of the three samples are in the same range, indicating their similar SEI film resistance Rf. The 2nd semicircle of MS/SNFs is much larger than those of MS/HNFs and MS/PNFs, which are alike to each other, suggesting that the charge transfer resistance Rct1 (electrolyte/carbon interface) of C-PAN SNFs is higher than those of CPDA HNFs and C-PDA PNFs. This is verified by the EIS results directly measured from C-PDA PNFs and C-PAN SNFs (Figure S6). The 3rd semicircles of the three samples are in a decreasing order from MS/SNFs, MS/HNFs to MS/PNFs. The charge transfer resistance Rct2 (electrolyte/MoS2 interface) of MS/PNFs is obviously much lower than those of the other two samples, which can be attributed to the efficient contact between the two phases in MS/PNFs. The stacking of multiple grains in MS/SNFs and MS/HNFs results in indirect contact of most


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MoS2 grains with carbon and hence large resistance for charge conduction, while the effective contact between MoS2 single grains and carbon in MS/PNFs significantly reduces the charge conduction resistance during lithiation/de-lithiation. The larger [002] interlayer d-spacing of the MoS2 nanosheets in PNFs and HNFs also gives rise to lower diffusion resistance of lithium ions in the electrodes, as indicated by the smaller circle 4 for MS/HNFs and MS/PNFs than that for MS/SNFs. In addition, the single-grain nature of most MoS2 nanosheets in MS/PNFs ensures complete and smooth lithium insertion/extraction, and large quantity of internal cavities in CPDA PNFs and their high SSA offer abundant reaction sites for lithium-ion uptake as well as efficient volume buffering function. These lead to the best cycling performance and the highest capacity of MS/PNFs.

Figure 6. (a) Nyquist plots obtained from the impedance spectroscopy of MS/SNFs, MS/HNFs and MS/PNFs. (b) Enlarged Nyquist plots in high and moderate frequency region.


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To support the above claims, the morphologies of the three samples were examined after charge/discharge tests for 100 cycles. The results show that the layered structure of the MoS2 nanosheets disappears in all cases due to the predominant irreversible conversion from crystalline MoS2 to amorphous S in the initial cycles (Figure 7). Compared with the lithiated/delithiated compounds in MS/SNFs and MS/HNFs (Figure 7a and 7b), which tend to aggregate on the surfaces and be embedded in the SEI film, those on the outer surface of MS/PNFs are still well separated probably because the amount of MoS2 initially on the outer surface is smaller. Although the aggregation is also observed inside MS/PNFs, nevertheless the active materials are still constrained within the pores of the PNFs as verified by the HRTEM image and corresponding STEM-EDX elemental mapping results (Figure 7c). This confirms the best integration of the two phases in MS/PNFs.


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Figure 7. HRTEM images of (a1-a3) MS/SNFs, (b1-b3) MS/HNFs and (c1-c3) MS/PNFs after charge/discharge for 100 cycles. (C4) the STEM-EDX elemental mapping results (SE, SK, and MoK represents bright field image, mapping of sulfur and mapping of molybdenum, respectively) of MS/PNFs after charge/discharge for 100 cycles.

CONCLUSIONS


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MoS2 nanosheets were successfully introduced into three different types of 1D carbon nanostructures, namely solid, hollow and porous nanofibers, via in-situ hydrothermal route. The morphology and structure of the MoS2 nanosheets varies with the carbon hosts. In the solid and hollow nanofibers, the MoS2 nanosheets are comprised of a few grains that stacked parallel on the surface. Consequently, only the first grain is in direct contact with the carbon host. In porous nanofibers, most MoS2 nanosheets are of single-grain size and in contact with the carbon phase directly. The interlayer spacing of MoS2 in all samples is larger than that of commercial MoS2 bulk powder, while it is relatively larger in hollow and porous nanofibers because PDA-derived carbon surface may affect the nucleation and growth of MoS2 nanosheets. The above morphological and structural features lead to the most efficient contact between MoS2 and carbon, the lowest charge conduction resistance, the most active sites for lithiation/de-lithiation and the best buffering effect for MoS2 volume variation, and thus the best electrochemical performance of MS/PNFs among the three samples. The electrochemical activity of the MoS2 is very effectively utilized in MS/PNFs, and additional physical or chemical absorption may also contribute to the lithium-ion storage. The normalized capacity of MoS2 nanosheets in PNFs is comparable to the highest reported value for MoS2-based anodes.

ASSOCIATED CONTENT Supporting Information. The isotherm curves of PAN-derived solid carbon nanofibers (C-PAN SNFs), PDA-derived hollow carbon nanofibers (C-PDA HNFs), porous carbon nanofibers (CPDA PNFs) and MoS2 nanosheets-in-C-PDA porous nanofibers, the XPS results of MS/SNFs, MS/HNFs and MS/PNFs, the XRD patterns of MS-SNFs and MS-C-PDA-coated C-PAN SNFs,


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HRTEM images of MoS2 nanosheets in C-PDA porous nanofibers, the TGA curves of C-PAN SNFs, C-PDA HNFs, C-PDA PNFs and pure MoS2, Nyquist plots obtained from the impedance spectroscopy of C-PAN SNFs and C-PDA PNFs, calculation of grain size from Scherrer equation. This material is available free of charge via the Internet at http://pubs.acs.org. AUTHOR INFORMATION Corresponding Author *

E-mail: [email protected]

ACKNOWLEDGMENT This work was supported by Science and Engineering Research Council of the Agency for Science, Technology and Research (A*Star) and Ministry of National Development, Singapore under Grants 132 176 0013 and 132 176 0011.

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