Dual Core–Shell-Structured S@C@MnO2 Nanocomposite for Highly

Aug 17, 2017 - Herein, we first demonstrate a new synthetic strategy for dual core–shell structured S@C@MnO2 nanospheres cathode in Li–S batteries...
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Dual Core−Shell-Structured S@C@MnO2 Nanocomposite for Highly Stable Lithium−Sulfur Batteries Lubin Ni,* Gangjin Zhao, Guang Yang, Guosheng Niu, Ming Chen, and Guowang Diao* School of Chemistry and Chemical Engineering, Yangzhou University, Yangzhou 225002, Jiangsu, People’s Republic of China S Supporting Information *

ABSTRACT: Lithium−sulfur (Li−S) batteries have currently excited worldwide academic and industrial interest as a nextgeneration high-power energy storage system (EES) because of their high energy density and low cost of sulfur. However, the commercialization application is being hindered by capacity decay, mainly attributed to the polysulfide shuttle and poor conductivity of sulfur. Here, we have designed a novel dual core−shell nanostructure of S@C@MnO2 nanosphere hybrid as the sulfur host. The S@C@MnO2 nanosphere is successfully prepared using mesoporous carbon hollow spheres (MCHS) as the template and then in situ MnO2 growth on the surface of MCHS. In comparison with polar bare sulfur hosts materials, the asprepared robust S@C@MnO2 composite cathode delivers significantly improved electrochemical performances in terms of high specific capacity (1345 mAh g−1 at 0.1 C), remarkable rate capability (465 mA h g−1 at 5.0 C) and excellent cycling stability (capacity decay rate of 0.052% per cycle after 1000 cycles at 3.0 C). Such a structure as cathode in Li−S batteries can not only store sulfur via inner mesoporous carbon layer and outer MnO2 shell, which physically/chemically confine the polysulfides shuttle effect, but also ensure overall good electrical conductivity. Therefore, these synergistic effects are achieved by unique structural characteristics of S@C@MnO2 nanospheres. KEYWORDS: lithium−sulfur batteries, dual core−shell structure, mesoporous carbon, MnO2 shell, in situ redox reaction

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INTRODUCTION Lithium−sulfur (Li−S) batteries as one of the most promising candidates for next-generation energy storage system (EES) have currently attracted worldwide attention due to its high specific energy density (2600 Wh kg−1), nontoxicity, low cost, and high natural abundance of sulfur.1−3 However, the practical application of Li−S batteries is still hampered by major challenges including (1) poor electronic conductivity of sulfur, (2) dissolution of lithium polysulfides (LiPSs) and the resulting shuttle effect, and (3) severe volumetric expansion of the S electrode (80%) upon lithiation, which give rise to low sulfur utilization, rapid capacity decay and low Coulombic efficiency.4,5 To overcome these problems, an effective strategy has been focused on using specific carbon materials as sulfur hosts, (e.g., porous carbon,6−11 carbon nanofbers/nanotubes,12−15 graphene/graphitic carbon nitride16−22) to enhance the utilization of S cathode and migrate polysulfides (PSs) shuttle becuase of their high electronic conductivity, large surface area and robust © 2017 American Chemical Society

control of structures. However, the carbonaceous sulfur hosts still have obvious capacity decay upon long-term cycling owing to the weak interaction between polar LiPSs and nonpolar carbon.23 Recently, an exceptionally promising method of LiPSs chemical trapping has been developed by employing polar materials, such as metal oxides24−35/sulfides36−39/ carbides40/hydroxide41−43/nitrides44/polymers,45,46 etc. These polar hosts can offer high efficiency of chemisorption to suppress the polysulfide diffusion and then significantly improve long-term cycling stability. But they possess relatively lower electrical conductivity, which may also lead to low Coulombic efficiency, low specific capacity and a significant impedance increase compared with that of graphitic carbons. Therefore, it is vital to rationally design hosts structure in combination with polar materials and graphitic carbons, which Received: June 5, 2017 Accepted: August 17, 2017 Published: August 17, 2017 34793

DOI: 10.1021/acsami.7b07996 ACS Appl. Mater. Interfaces 2017, 9, 34793−34803

Research Article

ACS Applied Materials & Interfaces

Figure 1. (a) Schematic illustration of the synthetic process of S@C@MnO2 dual core−shell nanostructure. SEM images of (b) MCHS, (c) S@C, and (d) S@C@MnO2 composites.

Figure 2. (a) TEM images of (a) MCHS, (b) S@C composites, and (c, d) S@C@MnO2 composites (blue circle, inner carbon layer; red circle, outer MnO2 shell). (e, f) HRTEM images of δ-MnO2. (g) Dark field STEM image of S@C@MnO2 nanosphere and EDX elemental mapping on S, C, Mn, and O. (h) The TGA curves of S@C and S@C@MnO2 composites.

is expected to show unique advantages of good electrical conductivity and remarkable chemical entrapment of LiPSs for the long cycle life of the Li−S battery. In fact, carbon additives or conductive polymers as outer shell wrap around polar active materials to prepare core−shell-nanostructured ternary composites have been reported as some representative examples to date.47−49 However, the intermediate polar materials block the direct contact of sulfur with conductive carbon in these

nanostructures, thus still hindering the electronic/ionic transport. Thus, there is an urgent demand to design more efficient LiPSs-trapping structures and fundamentally solve the technical challenges of sulfur cathodes in Li−S batteries. Among known polar hosts materials, δ-MnO2 phase can be considered as a remarkable chemical inhibitor of LiPSs that effectively mediate polysulfide redox through the conversion of thiosulfate to polythionate species developed previously by Nazar and co34794

DOI: 10.1021/acsami.7b07996 ACS Appl. Mater. Interfaces 2017, 9, 34793−34803

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Figure 3. (a) XRD patterns of sublimed sulfur, S@C (85:15) and S@C@MnO2 composites. (b) XRD pattern of MnO2 shell. (c) N2 adsorption− desorption isotherms for MCHS, S@C and S@C@MnO2. (Inset: SBET for S@C and S@C@MnO2.) (d) The corresponding pore size distributions (Inset: Vtotal for S@C and S@C@MnO2).

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workers.28 Afterward, we recently reported that the surface phase evolution from γ-MnO2 to Mn3O4 and the following decomposition of Mn3O4 might give rise to the capacity degradation of MnO2−S electrode materials in Li−S batteries.31 Very Recently, other representative examples for the combination of MnO2 polar hosts and carbonaceous materials include MnO2 nanosheet/conductive polymer poly(3,4-ethylenedioxythiophene) (PEDOT) S@PEDOT/MnO250 composite or Graphene oxide cofunctionalized core−shell sulfur nanospheres S@MnO2@GO,51 and 3D porous rGO/MnO2−S aerogel composite (rGM-SA).52 Herein, we first demonstrate a new synthetic strategy for dual core−shell structured S@C@MnO2 nanospheres cathode in Li−S batteries by using mesoporous carbon hollow spheres (MCHS) as the template and then in situ δ-MnO2 growth on the surface of MCHS. The S@C@MnO2 hybrid is composed of inner highly conductive carbon layer and outer polar MnO2 shells with strong chemisorption ability toward polysulfide. In such a structure, the inner mesoporous carbon layer can keep close contact both with sulfur and outer MnO2 shell leading to the significant increase in the overall electronic conductivity, long-range electron transfer/Li+ diffusion, and structural stability of S@C@MnO2 cathode. Moreover, the outer MnO2 shell can entrap sulfur and LiPSs inside the cathode through physical confinement of dual core−shell structure and chemical interaction between MnO2 and LiPSs. The phase transformation from MnO2 to Mn3O4 occurs accompanied by the oxidation of polysulfides to intermediate polythionate via the redox reaction between LiPSs and δ-MnO2, which can further promote discharge process in the “Wackenroder reaction”. Hence S@C@MnO2 composite as cathode delivers extraordinary electrochemical performance with high specific capacity (1345 mAh g−1 at 0.1 C), very good rate capability (465 mAh g−1 at 5.0 C) and excellent cycling life over 1000 cycles (capacity decay rate of 0.052% per cycle at 3.0 C).

RESULTS AND DISCUSSION The design synthesis of dual core−shell S@C@MnO 2 composite is illustrated schematically in Figure 1a. Initially, the mesoporous monodispersed MCHS were based on the in situ generated silica primary particles in a one-pot and surfactant free approach.53 The “silica-assisted” self-assembly process to prepare MCHS involves the hydrolysis of tetrapropyl orthosilicate (TPOS) and polymerization of resorcinol−formaldehyde (RF) resin, followed by carbonization at 700 °C and selective removal of the SiO2 component. The scanning electron microscopy (SEM) images (Figure S1) of initial precursor SiO2@SiO2/RF before carbonization exhibit uniform spherical nanoparticles of ca. 300 nm in an average size. After carbonization and hydrofluoric acid (HF) etching, the MCHS not only still maintained their uniform spherical shape, but also produced outer mesoporous shell with obvious open entrance on the surface (Figure 1b). The hollow morphology and the radial porous shell with a thickness of ∼40 nm are verified from the transmission electron microscopy (TEM) (Figure 2a). The MCHS was further characterized by Raman and X-ray diffraction (XRD) spectroscopy, as shown in Figure S2.54,55 In the Raman spectrum, G-band (∼1593 cm−1) corresponds to the vibration of sp2-hybridized carbon atoms in a 2D hexagonal lattice, while D-band (∼1332 cm−1) is related to the defects in the graphite structure. Therefore, the intensity ratio of D/G bands reveals the graphitization degree of the carbon materials. The value of ID/IG for MCHS after carbonization at 700 °C is 0.84, indicating a realtive high degree of graphitization (Figure S2b).54,55 Furthermore, XRD analysis was also applied to evaluate the graphitization degree of MCHS. For MCHS smaple, a broad peak at around 24.5° and a very weak peak at about 43.7°correspond to the (002) and (101) diffractions of the graphitic materials, respectively (Figure S2a).54,55 Subsequently, sublimed sulfur was fully impregnated into presynthesized mesoporous hollow MCHS by the meltdiffusion method, to obtain core−shell S@C. Comparing 34795

DOI: 10.1021/acsami.7b07996 ACS Appl. Mater. Interfaces 2017, 9, 34793−34803

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Figure 4. XPS spectra of dual core−shell S@C@MnO2 composite. (a) S 2p, (b) Mn 2p3/2, (c) C 1s, (d) O 1s.

C@MnO2 nanocomposites both contribute to the physical and chemical entrapping of sulfur and polysulfides inside the dual core−shell structure, thus controlling the volume expansion of sulfur and shuttle effect during the lithiation. TGA reveals that 84.6 wt % of sulfur content in the S@C (85:15) was decreased to 58.2 wt % of sulfur loading in S@C@ MnO2 after in situ growth of MnO2 shell (Figure 2h). Moreover, carbon content was also reduced from ∼11 wt % in the S@C (85:15) composite to ∼8 wt % in S@C@MnO2 composite (see Experimental Section) after redox reaction. These results are in good agreement with TEM images (Figure 2a and 2d), indicating partial loss of the sulfur and carbon sources in the redox reaction with KMnO 4 solution. Furthermore, MnO2 shell can be also generated on the surface of MCHS without the existence of sulfur via the redox reaction between MCHS and KMnO4, which gives rise to the formation of the hollow C@MnO2 or MnO2 nanospheres by simply controlling the concentration of KMnO4 and reaction time (Exp. part). 56 The MnO2 shell in C@MnO2 or MnO2 nanospheres was further supported by SEM images (Figures S1e−1f), as well as XRD spectroscopy (Figure S2a). Therefore, two redox reactions between (S/C in S@C) and KMnO4 to form S@C@MnO2 composites can be described by two equations (eqs 130 and 256,57):

with MCHS, we can see that the mesoporous carbon shell of S@C become relative smooth and no clear open pores on the surface (Figure 1c), indicating that sulfur has been completely infiltrated into the mesoporous carbon shell and the void space of the MCHS. The strong sulfur signal observed in energydispersive X-ray spectroscopy (EDX) spectrum and elemental mapping of S@C both indicate a large amount of sulfur was successfully loaded into the mesoporous carbon shells (Figure S3). However, TEM image of S@C showed no sulfur aggregation in the cavity of MCHS (Figure 2b). This might be the reason that almost all of the sulfur of S@C composite was sublimated off owing to the electron beam heating effect under the ultrahigh vacuum.11 The maximum sulfur loading of about 85 wt % for mesoporous carbon shells and the inner volume of the MCHS for new synthesized S@C nanosphere composite was achieved and verified by Thermogravimetric (TGA, Figure 2h)/Elemental analysis (EA, see Exp. part). Finally, the δ-MnO2 shell is in situ generated onto the surface of S@C composites by a redox reaction between S@C and KMnO4 (Figure 1d).29,30 For S@C@MnO2 nanospheres (Figure 2c−2d), the carbon shell shrank from 40 to ca. 25 nm and an obvious self-assembled nanosheet-like MnO2 shell appeared as the new external shell of ca. 50 nm thickness. In the HRTEM image (Figure 2e−2f), two characteristic lattice spacings of 0.70 and 0.25 nm correspond to the (001) and (110) planes of birnessite δ-MnO2 phase, respectively. EDX elemental mapping confirmed that S@C@MnO2 contains S, C, Mn, and O elements (Figure 2g). Moreover, the homogeneous dispersed S overlaying with C, Mn, and O signals can further verifies this novel unique dual core−shell structure of S@C@ MnO2. XRD analysis is also used to identify the crystal phase of MnO2 and sulfur in the composites. The XRD patterns of S@C (85:15) and S@C@MnO2 confirmed the existence of elemental sublimed sulfur (Figure 3a). In line with HRTEM images, the XRD pattern of the outer MnO2 shell indicates the monoclinic birnessite (δ-MnO2, JCPDS 80-1098, Figure 3b). Therefore, the outer MnO2 shell together with inner carbon layer in S@

6KMnO4 + 3S + H 2O = 6MnO2 + K 2SO4 + K3H(SO4 )2 + KOH

(1)

4KMnO4 + 3C + H 2O = 4MnO2 + 2KHCO3 + K 2CO3 (2)

The nitrogen adsorption analysis for MCHS, S@C (85:15) and S@C@MnO2 shows type IV isotherm (Figure 3c). The total specific surface area (SBET) and pore volume (Vtotal) for the hollow structured MCHS are about 1087.1 m2/g and 1.96 cm3/ g, respectively, which are much larger than the values for S@C (14.6 m2/g and 0.1 cm3/g) and S@C@MnO2 (26.8 m2/g and 0.16 cm3/g) composites. These results indicate that S8 34796

DOI: 10.1021/acsami.7b07996 ACS Appl. Mater. Interfaces 2017, 9, 34793−34803

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Figure 5. (a) CV curves and (b) Nyquist plots of the S@C@MnO2 nanocomposite.

Figure 6. (a) Charge−discharge profiles and (b) Rate performance of S@C@MnO2 and S@C (60:40) at different rates. (c) Cycling performance and Coulombic efficiency of S@C@MnO2 and S@C at 1.0 C for 500 cycles.

The S@C@MnO2 material was further characterized by Xray photoelectron spectroscopy (XPS) techniques. The S 2p spectrum has one doublet at 163.9 and 165.1 eV assigned to elemental sulfur (SB0),58 while two small peaks located at higher energies of 168.4 and 169.6 eV could be attributed to SOx species derived from sulfur oxidation in air or covalent bonding of S to MnO2 (Figure 4a).29,59 The Mn 2p3/2 spectrum displays a characteristic multiplet with two peaks at 642.3 and 640.9 eV corresponding to the Mn(IV) and Mn(III), respectively, and two satellite peaks at 643.6, 645.2 eV (Figure 4b).29 In the C 1s spectrum (Figure 4c), the main peak at 284.7 eV is corresponding to sp2 carbon (C−C or CC). The peaks at 286.2 and 287.3 eV can be ascribed to C−O (epoxy or alkoxy) and CO (carbonyl or carboxyl) groups, respectively.58 The spectrum of O 1s (Figure 4d) could be divided into four single peaks with binding energies of 530.1, 531.6, 532.8, and 533.9 eV, which can be attributed to the oxygen bonds of (CO or M−O(latt)), C−OH, (C−O or H−O−H from adsorbed H2O), C−O−C, respectively.60,61 The XPS results also further verified dual-layer structure of carbon layer and MnO2 shell in the S@ C@MnO2 nanocomposite.

molecules penetrate into the hollow structure through pores of the carbon layer in S@C composite during the melt-diffusion process. Moreover, the maximum theoretical sulfur loading impregnated in the mesoporous carbon shell of MCHS is about 69 wt %, which can be calculated based on the density of lithium sulfide (Li2S, 1.66 g/cm3) and the total pore volume (1.96 cm3/g) of MCHS.6 Figure 3d exhibits the pore size distribution of MCHS, S@C, and S@C@MnO2 derived from the BJH method, in agreement with the SBET and Vtotal values. The major difference between S@C@MnO2 and (MCHS, or S@C) is that MCHS or S@C has only a single mesopores pattern at ∼5 nm, whereas S@C@MnO2 has both mesopores in ∼5 and ∼8 nm. The S@C@MnO2 nanocomposite exhibits increasing SBET value and new developed 8 nm mesopores owing to in situ growth of MnO2 nanoflakes onto S@C nanosphere, compared to S@C (85:15) composite. Therefore, S@C@MnO2 composite not only can confine polysulfides within their dual-mesoporous-layer, but also can construct highly efficient electron-transfer/Li+ diffusion pathways via the pore structure of inner conductive carbon layer and outer MnO2 shell. 34797

DOI: 10.1021/acsami.7b07996 ACS Appl. Mater. Interfaces 2017, 9, 34793−34803

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Figure 7. Long-term cycling performance and Coulombic efficiency of dual core−shell S@C@MnO2 composite at large C-rates of (a) 2.0 C and (b) 3.0 C.

electrodes. Therefore, dual core−shell structured S@C@MnO2 composite with outer MnO2 shell and inner mesoporous carbon layer can significantly enhance electrical conductivity of the electrode and accelerate electron transfer during the lithiation process. We next investigated the electrochemical performances of the dual core−shell S@C@MnO2 cathode in Li−S cells. The galvanostatic discharge/charge profiles of the S@C@MnO2 at different current rates from 0.1 to 5.0 C show two typical discharge plateaus at ∼2.37 and 2.10 V (Figure 6a). The voltage plateaus existing are consistent with two cathodic peaks of CV curves. The S@C@MnO2 delivers a specific capacity of ∼1345 mAh g−1 at 0.1 C without any noticeable overpotential, which has achieved more than 80% of its theoretical capacity. With increasing C-rate, the specific capacity is gradually reduced to 1237, 1204, 1121, 998, 855, 668, and 465 mAh g−1 at 0.2, 0.3, 0.5, 1.0, 2.0, 3.0, and 5.0 C, respectively, confirming the excellent electronic/ionic transport properties and reaction kinetics of the dual core−shell structure (Figure 6a). Figure 6b shows the consecutive cycling performance of the S@C@ MnO2 and S@C hybrids at different C-rates, measured for ten cycles at each current in the order of 0.1 to 5.0 C, followed by the recovery at 0.1 C. The discharge specific capacities for the S@C@MnO2 composite steadily change with C-rate increasing. The S@C@MnO2 composite still demonstrates a higher average reversible capacity of 460 mAh g−1 even after cycling at high current rate 5.0 C. When the C-rate returned back to 0.1 C, the reversible capacity of the S@C@MnO2 composite were recovered to 1259 mAh g−1, which is close to the original 0.1 C level (1345 mAh g−1). In contrast, the discharge capacities of the S@C (60:40) nanospheres decrease rapidly with the increase of charging/discharging rates. The specific capacity of 1115 mAh g−1 for S@C at 0.1C quickly drops to 257 and 98 mAh g−1 as the rate increases to 3.0 and 5.0 C, respectivrly (Figure 6b). It is quite obvious that much higher specific capacities can be obtained for all applied current densities in the case of S@C@MnO2 electrode, whereas S@C nanosphere electrode revealed very lower capacities. The results

Coin cells (2032) were fabricated using the S@C@MnO2 hybrid as the cathode and metallic lithium as the anode to evaluate the electrochemical performance. In comparison with S@C@MnO2 cathode, S@C (60:40) or S@MnO2 electrode having similar sulfur content (∼58 wt %) were also synthesized through sulfur melting diffusion method into MCHS or hollow MnO2 nanospheres (exp. part), and then their electrochemical performances were also studied. Cyclic voltammograms (CVs) of Li−S cells with S@C@MnO2 nanocomposite for the first five cycles at a scan rate of 0.1 mV s−1 in the potential range of 1.7−2.8 V are shown in Figure 5a. In the first cycle, two cathodic peaks at 2.37 and 2.00 V are assigned to (I) S8 ringopening and the conversion to soluble higher-order lithium polysulfides (Li2Sn, 4 ≤ n ≤ 8) and (II) their further reduction to short-chain lithium sulfides Li2S2/Li2S, respectively.5 Two closed anodic peaks at 2.44 and 2.48 V can be attributed to (III) the formation of higher-order LiPSs and (IV) the reverse oxidation to the neutral S8.62 The CV curves are almost identical for the first five cycles, exhibiting relatively good capacity retention (Figure 5a). Moreover, the electrochemical impedance spectroscopy (EIS) measurements of the S@C@ MnO2 were carried out at different discharged voltage states in the first discharge cycle (Figure 5b). The semicircle in the high/middle frequency region is attributed to the chargetransfer process occurring at the electrolyte-electrode interface and the straight line (Warburg impedance) in the low frequency region is associated with semi-infinite diffusion.63 The interfacial transport resistance obviously decreased the discharging process compared to that at open-circuit-voltage (OCV) state, suggesting that the transformation from insulating sulfur electrode to soluble polysulfide can improve the interfacial contacts between S species and conductive agent.31 Afterward, EIS measurements were further employed to compare the resistance among the bare S, S@MnO2, S@C (60:40), and S@C@MnO2 at OCV state (Figure S4). Noticeably, the charge transfer resistances (Rct) of the S@ C@MnO2 (89.9 Ω) and S@C (84.3 Ω) electrodes are much lower than bare S (241.9 Ω) and S@MnO2 (130.5 Ω) 34798

DOI: 10.1021/acsami.7b07996 ACS Appl. Mater. Interfaces 2017, 9, 34793−34803

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Figure 8. (a) Digital pictures of a Li2S4/DME solution after 10 min upon contact with blank, MCHS, MnO2, C@MnO2. (b) UV−vis spectra (350− 600 nm) of the Li2S4 solution after 10 min upon contact with blank, MCHS, MnO2, C@MnO2.

S@C@MnO2 hybrid nanospheres are still maintained in all the coin cells (Figure S9), indicating the improved structural stability of S@C@MnO2 hybrid. To further demonstrate the adsorption ability to the polysulfdes, the same amount of various sulfur host materials of MCHS, hollow MnO2, and C@MnO2 nanospheres was added to the lithium polysulfde solution (Figure 8a). The blank polysulfdes solution (mainly Li2S4) was prepared according to the previously reported literature12,36 and dissolved in DME solution leading to a yellow color (Exp. part, Figure 8a). As illustrated by Figure 8a, the Li2S4 solution become transparent after coming into contact with hollow MnO2 and C@MnO2 nanospheres for 10 min, indicating strong chemisorption toward the LiPSs. However, the polysulfdes solution exposed to MCHS still presents a similar deep orange color compared to the blank LiPSs solution due to the weak interaction between MCHS and LiPSs. In addition, the strong adsorption capability of MnO2 or C@MnO2 toward the LiPSs was further demonstrated by the UV−vis spectroscopy (Figure 8b).45,64 The characteristic absorption peak at 417 nm of Li2S4 in the DME solution completely vanished after adding the MnO2 or C@MnO2 composite compared to that with MCHS or in the blank solution.64 Next, we used XRD and X-ray photoelectron spectroscopy (XPS) to acquire further information about the interaction mechanism between (MnO2 or C@MnO2 composite) and LiPSs. The XRD pattern of hollow δ-MnO2 nanoshperes after Li2S4 treatment for 10 min in DME solution are presented in Figure S10. All the characteristic reflection lines of δ-MnO2 were missing, while a new broad reflection peak appeared at scanning angles (2θ) of and 36.1° corresponding to the characteristic (211) reflection line of the Mn3O4 hausmannite phase (Figure S10). Furthermore, the XRD pattern of δ-MnO2 sample exposed to sodium polysulfde (mainly, Na2S4) aqueous solution also confirms the presence of a highly crystalline Mn3O4 phase (Figure S10). Thus, the results from XRD data indicate that the δ-MnO2 is partially transformed into new Mn3O4 phase accompanied by the redox interaction between polysulfdes and δ-MnO2 host, resulting in a decrease in the oxidation state of manganese from Mn4+ to Mn3+ and Mn2+. Similar phenomena also have been observed in our previous related studies on the reaction of γ-MnO2 phase with LiPSs in Li−S cell.31 XPS analysis is currently the most widely used to identify the oxidation state and quantitative chemical composition of surface species.65 The XPS Mn 2p3/2, S 2p spectra of MnO2, and C@MnO2 composites after Li2S4

indicate the excellent reliability and stability of the S@C@ MnO2 dual core−shell structure compared to core−shell S@C nanospheres. Moreover, no significant contribution to the capacity for the δ-MnO2 shell itself as the initial discharge capacity is lower than 5 mAh g−1 at 1.0 C (Figure S5). Next, the cycling stability and Coulombic effciency of the cells based on S@C@MnO2 nanocomposite were also performed at practical level current density of 838 (0.5 C), 1675 (1.0 C), 3350 (2.0 C), and 5025 mA g−1 (3.0 C). The cycling performances of the S@C@MnO2 and S@C (60:40) hybrids at 0.5 C over 300 cycles and 1.0 C over 500 cycles are shown in Figures S6 and 6c, respectively. For S@C@MnO2, an initial discharge capacities of 1142 mAh g−1 for 0.5 C (areal capacity of 3.4 mAh cm−2, Figures S6 and S7) and 983 mAh g−1 for 1.0 C (areal capacity of 3.0 mAh cm−2, Figures 6c and S8) are obtained. The reversible capacities still remain at 712 mAh g−1 for 0.5 C (areal capacity of 2.1 mAh cm−2, Figures S6 and S7) after 300 cycles, and 550 mAh g−1 for 1.0 C after 500 cycles (areal capacity of 1.7 mAh cm−2, Figures 6c and S8) with only 0.088% capacity decay per cycle. In contrast, the initial discharge specific/areal capacities of S@C (60:40) electrode for 0.5 and 1.0 C at 800 mAh g−1/2.4 mAh cm−2 and 592 mAh g−1/1.8 mAh cm−2 fade very quickly to 219 mAh g−1/0.7 mAh cm−2 and 195 mAh g−1/0.6 mAh cm−2, respectively (Figures S6−S8). Compared with S@C, the as-synthesized dual core− shell S@C@MnO2 composites exhibited remarkably higher specific capacity and cycling stability. Furthermore, superior long-life cycling performance at a high current density of 2.0 and 3.0 C over 1000 cycles are also evaluated for S@C@MnO2 (Figure 7a). The S@C@MnO2 still can achieve good cycling stability with very low capacity fade rate of 0.060% (2.0 C) and 0.052% (3.0 C) per cycle, respectively. The Coulombic efficiency was all above 98% throughout (Figure 7), indicating that the shuttle effect has been effectively restrained by the inner mesoporous carbon layer together with outer MnO2 shell. In short, S@C@MnO2 cathode presents the high-level performance with the large discharge capacity, high-rate capability and long cyclability, which is superior to known S/ metal oxide cathodes in Li−S batteries.24−35 In addition, we further investigate the morphology and structural changes of the S@C@MnO2 cathode after cycling, to demonstrate the structural stability of the dual core−shell S@C@MnO2 electrode. The S@C@MnO2 electrode after 200 cycles at 1.0 C, 500 cycles at 2.0 C and 1000 cycles at 3.0 C in three different Li−S coin cells were disassembled and characterized by SEM (Figures S9). The morphology of the dual core−shell 34799

DOI: 10.1021/acsami.7b07996 ACS Appl. Mater. Interfaces 2017, 9, 34793−34803

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ACS Applied Materials & Interfaces

Figure 9. (a) Mn 2p3/2 XPS spectrum of hollow δ-MnO2 nanospheres. (b) Mn 2p3/2 XPS spectrum of MnO2 after Li2S4 treatment (MnO2−Li2S4). (c) Mn 2p3/2 XPS spectrum of C@MnO2 after Li2S4 treatment (C@MnO2−Li2S4). (d) S 2p XPS spectrum of Li2S4. (e) S 2p XPS spectrum of MnO2−Li2S4. (f) S 2p XPS spectrum of MnO2−Li2S4.

treatment (MnO2−Li2S4 and C@MnO2−Li2S4) are shown in Figure 9. The Mn 2p3/2 spectra of the MnO2−Li2S4 and C@ MnO2−Li2S4 samples exhibit a characteristic multiplet with a main peak at 642.3 eV corresponding to the Mn4+ and other two satellite peaks at 643.7 and 645.5 eV (Figure 9b and 9c).28 Compared to the pristine δ-MnO2 in S@C@MnO2 composite (Figure 9a), the contributions at 641.2 and 640.2 eV from Mn3+ and Mn2+ valent states obviously increases because of the partial reduction of Mn4+ to Mn3+ and Mn2+ accompanied by the oxidation of Li2S4.28,31 These results also prove that δMnO2 was reduced with LiPSs to partially evolve into lower manganese valence species Mn3O4, in accord with the conclusions about XRD patterns. Moreover, the S 2p3/2 XPS spectra of the recycled solid samples (MnO2−Li2S4 and C@ MnO2−Li2S4) can be interpreted as three doublets (Figure 9e and 9f). The relative intensities of the “bridging” (SB0, 163.7 eV) and “terminal” (ST−1, 161.9 eV) decline with respect to those of fresh Li2S4, which implies the decrease in the polysulfide species content after the redox raction between MnO2 and Li2S4 (Figure 9d−9f). The new S 2p3/2 peaks located at 169.4/168.0 eV can be mainly attributed to high valent sulfur oxysalts (SOx) polythionate.28,31 The polythionate complex is supposed to accelerate the conversion of longerchain PS into shorter-chain PS via the “Wackenroder reaction” during the discharge process.28,31 Therefore, S@C@MnO2 nanospheres as an effective Li−S host can chemically trap soluble LiPSs through the polythionate mechanism in accord with previous studies on MnO2-based sulfur host materials in Li−S cells.28−31 On the basis of these results, we can conclude that the extraordinary electrochemical performance of S@C@MnO2 is derived from the following structural advantages of S@C@ MnO2: (I) In the S@C@MnO2 hybrid, the mesoporous carbon as the inner layer maintains intimate contact with sulfur and outer MnO2 shell so as to significantly increase the overall electrical conductivity and electron/Li+ transport. (II) The mesoporous carbon wall can physically confine polysulfides within these smaller mesopores, meanwhile the weak interactions between (PSs or S) and MCHS could alleviate the shuttle effect on certain degree.66 (III) The outer MnO2 shell can further entrap polysulfides through physical structural restriction and chemical encapsulation. Because the strong redox interaction between MnO2 and LiPSs lead to the surface phase evolution from MnO2 to Mn3O4 together with the

oxidation of polysulfides to polythionate and then promoting discharge process in the “Wackenroder reaction”.28,31

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CONCLUSIONS

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EXPERIMENTAL SECTION

In summary, dual core−shell structured S@C@MnO2 nanospheres with inner mesoporous carbon layer and outer MnO2 shell were prepared from MCHS precursors via the meltdiffusion method for incorporating sulfur and then redox reaction for in situ growth of MnO2 shell. The inner carbon layer can effectively increase the overall electrical conductivity, and the outer MnO2 shell together with inner carbon layer provides physical structural restriction and chemical encapsulation for sulfur storage to restrain the shuttling of polysulfide during electrochemical reactions. Because of the synergistic effects, the S@C@MnO2 hybrid cathode exhibits high specific capacity with 1345 mAh g−1 at 0.1 C, extraordinary rate capability with 465 mA h g−1 at 5.0 C and cycling stability with capacity fade rate of 0.052% per cycle after 1000 cycles at 3.0 C in Li−S batteries. This work not only brought forward an new accessible “hands-on” strategy in the fabrication of porous carbon and metal oxide hybrid as sulfur host but also shed new light on the practical application of high performance Li−S batteries.

Synthesis of Mesoporous Carbon Hollow Sphere (MCHS). MCHS was prepared by a “silica assisted” self-assembly method according to the previous literature.53 In a typical synthesis of SiO2@ SiO2/RF, tetrapropyl orthosilicate (3.46 mL, 12 mmol) was added into the mixture of ethanol (70 mL), water (10 mL), and ammonia (3 mL, 25 wt %) under stirring at room temperature. After 15 min, 0.4 g of resorcinol and 0.56 mL of formaldehyde (37 wt %) were slowly injected into the above solution and stirred for 24 h at room temperature. The SiO2@SiO2/RF precursors were separated by centrifugation, washed several times with water and ethanol, and dried under vacuum overnight. MCHS products were obtained after calcinating SiO2@SiO2/RF at 700 °C under N2 flow for 5 h and final etching of SiO2 by HF (10 wt %). Synthesis of Mesoporous Hollow C@MnO 2 and MnO2 Nanospheres. The as-prepared MCHS (80 mg) were dispersed into 60 mL of deionized water, followed by the addition of potassium permanganate (KMnO4, 640 mg) under gradual stirring. Next, the solution was stirred continuously at 70 °C for 24 h. The C@MnO2 products were washed repeatedly with ethanol and deionized water, and dried under vacuum at 50 °C overnight. The synthetic procedure 34800

DOI: 10.1021/acsami.7b07996 ACS Appl. Mater. Interfaces 2017, 9, 34793−34803

Research Article

ACS Applied Materials & Interfaces for MnO2 nanospheres was slightly modified by using 1500 mg of KMnO4 instead of 640 mg. Synthesis of Mesoporous Core−Shell S@C and S@MnO2 Nanosphere Composites. The core−shell S@C nanosphere was obtained using a simple melt-diffusion approach. Mesoporous carbon hollow spheres (MCHS) were first mixed with sublimed sulfur at two desired weight ratios of S: C (85:15) and (60:40), respectively, and then heated at 155 °C for 20 h in an argon atmosphere. Elem anal. for S@C (85:15): C, 10.59%; S, 85.18%; for S@C (60:40): C, 39.52%; S, 58.22%. The synthetic procedure for S@MnO2 nanospheres was slightly modified by using MnO2 nanospheres instead of MCHS in the desired weight ratio of S: MnO2 (60:40) via a melt-diffusion method. Elem anal. for S@MnO2 (60:40): C, 0.47%; S, 59.25%. Synthesis of Dual Core−Shell Ternary S@C@MnO2 Nanosphere Composite. The as-prepared 54 mg of S@C (85:15) and polyvinylpyrrolidone (Mw ≈ 40 000) (20 mg) were dispersed in 80 mL of deionized water, and then the reaction mixture was kept stirring for 0.5 h. 47 mg of Potassium Permanganate (KMnO4) was added in the mixed solution by sonication for 1 h. The desired products were washed by centrifugation and dried under vacuum overnight, and the reaction yield for S@C@MnO2 was about 93% (∼70 mg). Elem anal. for S@C: C, 7.76%; S, 58.61%. Polysulfdes Adsorption Study. Li2S4 solution (0.04 M) was prepared by adding 184 mg of Li2S (4.0 mmol) and 384 mg of sublimed sulfur (12 mmol) in 100 mL of DME solution with vigorous stirring, according to previous reports.12,36 Dilute Li2S4 test solution was prepared by mixing 1 mL of 0.04 M Li2S4 in DME, 9 mL of DME. Thirty milligrams of δ-MnO2 or C@MnO2 was added into a dilute Li2S4 test solution, and the reaction mixture was stirred for an additional 10 min. All procedures were completed in an Ar-flled glovebox. The synthetic procedure for Na2S4 aqueous solution was modified by using 960 mg of Na2S·9H2O instead of 184 mg Li2S in 30 mL of water solution. Characterization. X-ray diffraction (XRD) data were conducted on a D8 advance super speed powder diffractometer (Bruker). Elemental analyses (carbon, sulfur) were performed using Vario EL cube elemental analyzer. UV/vis spectra were carried out with a PerkinElmer Lambda 650S spectrometer. Raman spectra were performed on a Renishaw via Raman spectroscopy. Scanning electron microscopy (SEM) was carried out with Zeiss Supra−55VP. Transmission Electron Microscopy (TEM) was conducted on a Philips TECNAI-12 instrument. HRTEM and HAADF-STEM were measure by a FEI Tecnai G2 F30 STWIN (USA). The TGAs were collected with a Netzsch TG209 F1 instrument under N2 flow. The Xray photoelectron spectroscopy (XPS) were carried out with a Thermo Escalab 250 system. The Barrett−Joyner−Halenda (BJH) pore-size distribution and Bruker−Emmett−Teller (BET) surface areas were recorded at 77 K on a 3H-2000PS2 apparatus after degassed at 150 °C for 12 h under vacuum. Electrochemical Measurements. Cathodes were prepared by casting of N-methyl pyrrolidone (NMP) slurry containing 70 wt % active materials (S@C@MnO2, S@C, and S@MnO2), 20 wt % conductive material Super P li (TIMCAL) and 10 wt % binder polyvinylidene fluoride (PVDF, HSV900) onto a carbon paper current collector (GDL 28 AA, SGL), followed by vacuum-dried at 60 °C overnight. The average active materials mass on each electrode disk is about ∼3.0 mg/cm2. Electrochemical measurements were performed using 2032-type coin cells with microporous membrane (Celgard 2300) as separator and lithium metal as the counter electrode. The electrolyte was 100 mL of 0.5 M lithium trifluoromethanesulfonate (LiCF3SO3) and 0.5 M LiNO3 in a mixture of 1,3-dioxolane (DOL) and dimethoxymethane (DME) (1:1 by volume). The amount of the electrolyte was strictly controlled for performance evaluation with carbon paper current collector, and the cell contained an electrolyteto-sulfur (E/S) ratio of ∼19 mL g−1. Cyclic voltammetry (CV) and electrochemical impedance spectroscopy (EIS) measurements were carried out with an electrochemical workstation (CHI660 E, Chenghua, China). The galvanostatic discharge−charge tests were conducted using a Neware Battery Measurement System (Neware, China) with a potential range of 2.8−1.7 V. The specific capacity was

calculated based on the mass of sulfur. The calculations of areal current densities and areal capacities of electrodes were based on both specific capacities and areal S mass loading of the corresponding electrode.

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acsami.7b07996. SEM images, Raman shifts, XRD patterns, TG curves, EIS plots, cycling stability curves (PDF)

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AUTHOR INFORMATION

Corresponding Authors

*E-mail: [email protected]. *E-mail: [email protected]. ORCID

Ming Chen: 0000-0002-6436-4765 Guowang Diao: 0000-0001-9224-5432 Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work was supported by the National Natural Science Foundation of China (Grant nos. 21401162, 21773203), NSFC-NRF Scientific Cooperation Program (Grant no. 21511140282), the Natural Science Foundation of the Jiangsu Higher Education Institutions of China (Grant no. 14KJB430024), the Jiangsu Provincial Postdoctoral Sustentation Fund (Grant no. 1402015B), High-Level Entrepreneurial and Innovative Talents Program of Jiangsu, and Lvyangjinfeng Talent Program of Yangzhou. Financial support from the Priority Academic Program Development of Jiangsu Higher Education Institutions and the Natural Science Foundation of Education Committee of Jiangsu Province (no. 12KJB150023) is gratefully acknowledged. We also thank the testing center of Yangzhou University for SEM, TEM and XPS measurements.

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DOI: 10.1021/acsami.7b07996 ACS Appl. Mater. Interfaces 2017, 9, 34793−34803