Fabrication of Microporous Sulfur-Doped Carbon Microtubes for High

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Fabrication of Microporous Sulfur-Doped Carbon Microtubes for High-Performance Sodium-Ion Batteries Qiaoqiao Wang, Xufang Ge, Jingyi Xu, Yichen Du, Xin Zhao, Ling Si, and Xiaosi Zhou ACS Appl. Energy Mater., Just Accepted Manuscript • DOI: 10.1021/acsaem.8b01690 • Publication Date (Web): 22 Oct 2018 Downloaded from http://pubs.acs.org on October 27, 2018

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ACS Applied Energy Materials

Fabrication of Microporous Sulfur-Doped Carbon Microtubes for High-Performance Sodium-Ion Batteries Qiaoqiao Wang,† Xufang Ge,† Jingyi Xu, Yichen Du, Xin Zhao, Ling Si,* and Xiaosi Zhou* Jiangsu Key Laboratory of New Power Batteries, Jiangsu Collaborative Innovation Center of Biomedical Functional Materials, School of Chemistry and Materials Science, Nanjing Normal University, Nanjing 210023, China

ABSTRACT: Developing efficient anode materials for sodium-ion batteries (SIBs) is important for the storage of renewable energy. Inspired from the rapid development of biomass-derived hard carbons and heteroatom-doped carbon materials in various areas, a high-temperature sulfurizing method is exploited for the fabrication of sulfur-doped carbon microtubes (S-CMTs). Owing to high sulfur doping (10.2 wt %) and well-developed microporous structure, the as-prepared S-CMTs show a large charge capacity of 532 mAh g−1 at a current rate of 200 mA g−1, outstanding rate capability (234 mAh g−1 at 2 A g−1), and exceptional cycling stability (281 mAh g−1 after 1000 cycles at 1 A g−1), which are superior to those of biomass-derived carbons reported previously. The excellent electrochemical performance of S-CMTs in full cells paired with N, B co-doped carbon coated Na3V2(PO4)3 cathode further demonstrate the feasibility of SIBs. The simple synthesis strategy can potentially be extended to other carbon-based anode materials for sodium-ion batteries.

KEYWORDS: Microporous structure, sulfur-doping, carbon microtubes, anode, sodium-ion battery

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1. INTRODUCTION Sodium-ion batteries (SIBs) has drawn ever-growing attention as a potential alternative to lithium-ion batteries (LIBs) for implementation in large-scale electrochemical energy storage because of the low cost and widely distributed sodium.1-5 Nevertheless, since the radius of Na+ (1.02 Å) is larger than that of Li+ (0.76 Å), the diffusion kinetics of sodium ion is much tardier. Therefore, it is hard to search for suitable accommodating materials with a big enough space to hold sodium ion and to permit rapid and reversible Na+ uptake/release, especially pertaining to the negative electrodes. So far, many sodium accommodating anode materials have been reported to demonstrate appropriate specific capacity and cycle lifespan, mostly involving carbonaceous materials,6-7 titanates,8 alloys,9-12 and metal oxides/sulphides.13-18 Unfortunately, the relatively high cost of titanates and the huge volume variation of alloys and metal oxides/sulfides during sodiation/desodiation severely restrict their commercial applications.19 Hence, a great deal of studies have been continuously devoted to exploiting efficient carbon-based anode materials. As the prevailing anode material for LIBs, graphite only affords a capacity of around 35 mAh g−1 by generating NaC64, which is mainly due to the mismatch of the large sodium ions with the narrow graphite interlayer spacing (d002 = 3.35 Å). Hard carbon shows excellent sodium storage properties and diverse kinds of hard carbons, such as carbon black,20 carbon fiber,21 and biomass-derived carbons, have been surveyed.22-26 Nonetheless, their electrochemical performance is inferior to that obtained by graphite in a LIB system, with a charge capacity of only about 320 mAh g−1 at a current density as low as 30 mA g−1. Recently, plenty of studies have been conducted to enhance their specific capacity and cycling performance by fabricating novel porous carbon nanomaterials.27-34 Although remarkable enhancement has been achieved by constructing various carbon nanostructures, nitrogen, sulfur, and phosphorus doping have recently received significant interest as another effective


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measure to boost the sodium storage properties of carbon-based materials.35-40 Up to now, nitrogen has become the most extensively studied heteroatom, which could improve the electronic conductivity and electrochemical activity by creating external defects.41 Compared with nitrogen, doping with phosphorus or sulfur is relatively seldom and stands for an arising research area. Sulfur doping is especially interesting for energy storage implementations because sulfur is a high-capacity electroactive element when react with sodium.42 In particular, the integration of sulfur into the carbon materials usually offer extra deposit sites for Na+, leading to the improvement of reversible capacity. In this work, we demonstrate for the first time microporous sulfur-doped carbon microtubes (designated as S-CMTs), which are synthesized by a facile sulfurizing method, as a SIB anode material. It delivers a large charge capacity of 532 mAh g−1 at 200 mA g−1, outstanding rate property (234 mAh g−1 at 2 A g−1), and superior cycling performance (281 mAh g−1 after 1000 cycles at 1 A g−1 with a 90% capacity retention). Moreover, a full cell composed of N, B co-doped carbon coated Na3V2(PO4)3 cathode43 and the S-CMTs anode shows a large charge capacity of 370 mAh g−1 at 200 mA g−1 with a high capacity retention of >81% after 100 cycles. 2. EXPERIMENTAL SECTION Synthesis of S-CMTs In a typical procedure, cotton roll was sulfurized in a tube furnace at 700 oC for 3 h under a H2S/Ar (10:90 v/v) atmosphere with a ramp rate of 5 oC min−1 to achieve the product of S-CMTs. For comparison, cotton roll was carbonized at 700 oC by the same procedure under an argon flow to achieve the product of carbon microtubes (CMTs). Sulfur and CMTs were ground together with a weight ratio of 10.2:89.8, and maintained at 155 oC for 12 h under Ar atmosphere to obtain the product of sulfur/carbon microtubes composite (S/CMTs). Materials Characterization



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Thermogravimetric analysis (TGA) and differential thermal analysis (DTA) were conducted on a NETZSCH STA 449 F3. Field-emission scanning electron microscopy (FESEM) images were taken on a JSM-7600F (JEOL) scanning electron microscope working at 10 kV. Transmission electron microscopy (TEM) and high-resolution TEM (HRTEM) images were acquired with a JEM-2100F (JEOL) transmission electron microscope working at 200 kV. Elemental mapping analysis was performed on the JEOL JSM-7600F scanning electron microscope attached with a Thermo Fisher Scientic energydispersive X-ray spectrometer. Powder X-ray diffraction (XRD) were carried out with a Rigaku SmartLab diffractometer using Cu Kα radiation. X-ray photoelectron spectroscopy (XPS) analysis were conducted on an ESCALab 250Xi electron spectrometer. Nitrogen sorption measurements were performed on an ASAP 2050 surface area-pore size analyzer. Raman spectra were characterized with a laser wavelength of 514 nm on a Labram HR800. Fourier transform infrared (FT-IR) spectroscopy was recorded on a Thermo Nicolet Nexus 670 FT-IR spectrometer. The electronic conductivity was determined with a thickness of 1.5 and 2.2 mm for CMTs and S-CMTs under the pressure of 10 MPa by an M-3 Mini type four-probe tester, respectively. Electrochemical Characterization Electrochemical tests were performed using CR2032 coin cells. The working electrodes were fabricated by mixing 70 wt % active materials, 15 wt % carbon black (Super-P), and 15 wt % carboxymethyl cellulose sodium in water using a mortar and pestle. The obtained slurry was pasted onto a copper foil and then vacuum-dried at 40 oC for 12 h. The loading mass of the active material is approximately 1.2 mg cm−2. Thin sodium sheets were utilized as counter electrodes, and Glass fibers (GF/D) from Whatman were employed as separators. The electrolyte solution consists of 1.0 M NaClO4 in polycarbonate (PC) with 5 vol % fluoroethylene carbonate (FEC) additive. The coin cells were constructed in an argon-filled glovebox (H2O, O2 < 0.1 ppm). Galvanostatic charge/discharge analysis were conducted on a Land


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CT2001A multichannel battery tester in a voltage window of 0.01−3 V at room temperature. Cyclic voltammetry (CV) was measured by a PARSTAT 4000 electrochemical workstation. Electrochemical impedance spectroscopy (EIS) measurements were conducted in a frequency range from 100 kHz to 100 mHz. For the full cell measurement, N, B co-doped carbon coated Na3V2(PO4)3 composite (NVP@C-BN) was chosen as the positive electrode material. The NVP@C-BN composite was fabricated via a recently reported method.43 To prepare the NVP@C-BN cathode, NVP@C-BN, carbon black (Super-P), and PVDF binder with a mass ratio of 7:2:1 were blended into a slurry using a mortar and pestle. The slurry mixture was pasted onto aluminum foil and then vacuum-dried at 80 ºC overnight for further application. Prior to assembling the full cell, the NVP@C-BN cathode was initially operated within the voltage range of 2.3−3.9 V in a half-cell for activation, while the S-CMTs anode was cycled between 0.01 and 3 V to get rid of the irreversible capacity from the initial cycles. Lastly, the full battery was assembled with the cycled NVP@C-BN cathode and S-CMTs anode. To prevent Na plating on S-CMTs and make full use of the NVP@C-BN cathode, we adopted an anode-to-cathode capacity ratio of 1:0.95, corresponding to mass loading ratio of around 1:4 with the PC/FEC electrolyte. Galvanostatic charge and discharge measurements were conducted between 0.3 and 3.6 V. 3. RESULTS AND DISCUSSION



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Figure 1. Schematic representation of the synthetic process for CMTs and S-CMTs. S-CMTs was fabricated simply by annealing the natural cotton roll in a flowing atmosphere of H2S/Ar (10:90 by volume) at an optimized temperature of 700 oC for 3 h. Figure 1 presents the schematic illustration of the synthesis of S-CMTs and CMTs. For comparison, CMTs pyrolyzed from cotton roll was synthesized by the same method without H2S. S/CMTs was prepared via grinding sulfur and CMTs with a weight ratio of 10.2:89.8 and subsequent heating at 155 oC under argon protection. TGA and DTA curves of the natural cotton roll in a H2S/Ar flow show that the cotton vulcanization temperature is about 470 oC (Figure S1). As the temperature rise to 600 oC, the cotton roll has been completely carbonized and sulfurized, with excess H2S escaping freely. By using elemental analysis, the weight fractions of C, S, O, H in S-CMTs are revealed to be 73.3%, 10.2 %, 14.6%, and 1.9%, respectively (Table S1).


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Figure 2. FESEM images of (a) CMTs and (b) S-CMTs; (c) HRTEM image and (d) EDS elemental mapping images of S-CMTs. The morphologies and structures of as-synthesized S-CMTs and CMTs were investigated by FESEM and HRTEM, as shown in Figure 2a−c. CMTs are composed of distorted hollow microtubes with a diameter of around 5−11 μm (Figure 2a), while the S-CMTs sample appears as thicker hollow microtubes with a diameter of around 8−14 μm (Figure 2b). This thicker hollow microtube structure possesses many interpenetrative nanopores, which is primarily produced by releasing a lot of CS2 and SO2 during the high-temperature vulcanization process. The HRTEM image (Figure 2c) illustrates the uneven surface of the hollow microtubes with many nanopores. This unique microporous architecture can supply a large amount of interconnecting passages, permitting the whole structure to be fully infiltrated by the



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electrolyte. Elemental mapping measurement (Figure 2d) displays the uniform distribution of C, O, and S throughout the entire area of CMTs, indicating a homogeneous sulfur-doping in S-CMTs.

Figure 3. (a) XRD patterns of cotton roll, sulfur, CMTs and S-CMTs. (b) High-resolution S 2p XPS spectrum of S-CMTs. (c) N2 adsorption-desorption isotherms and (d) pore-size distribution of S-CMTs and CMTs. (e) Raman spectra of S-CMTs, CMTs, and sulfur. (f) FT-IR spectra of CMTs and S-CMTs. Figure 3a show the XRD results of cotton, sulfur, CMTs, and S-CMTs. No obvious peaks of cotton or sulfur are seen in the XRD pattern of S-CMTs, suggesting no physically adsorbed sulfur on cotton-


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derived carbon. The XRD patterns of CMTs and S-CMTs demonstrate two main wide peaks centered at around 2θ = 23.8o and 43.7o, corresponding to the (002) and (101) planes of the low crystallinity carbon materials. Based on Bragg’s equation, the interplanar distances of the (002) planes of CMTs and S-CMTs are calculated to be about 3.73 and 3.81 Å, respectively, which is consistent with the HRTEM observations (Figs. S2 and 2c). These large interplanar spacings of CMTs and S-CMTs are desirable for the insertion/extraction of Na+.32, 44 It is worth mentioning that the wider interlayer spacing of S-CMTs in comparison with CMTs indicates that the interplanar spacing of disordered carbon structure can be expanded by sulfur doping, which generally boost the transportation and Na+ uptake/release and improve the electrochemical activity of disordered carbon material. In addition, both significant sulfur diffraction peaks (Figure S3) and narrow interlayer spacing (Figure S4) in S/CMTs suggest that sulfur is physically adsorbed rather than chemically bonded. XPS survey was carried out to distinguish the chemical status of sulfur in the S-CMTs sample. The high-resolution S 2p XPS spectrum is presented in Figure 3b, which displays three peaks located at 167.7, 164.0, and 162.8 eV. The first wide peak at 167.7 eV can be ascribed to S=O and S−O groups,45 and the latter two pronounced peaks corresponds to the S 2p1/2 and S 2p3/2 states of C−S.46 These XPS peaks have lower binding energies than those of S/CMTs (Figure S5), verifying that sulfur has been covalently bonded into the carbon network of S-CMTs. Nitrogen sorption measurement was further performed to characterize the specific surface area and pore distribution of as-made S-CMTs and CMTs. As shown in Figure 3c, the Brunauer−Emmett−Teller (BET) analysis indicates a microporous structure of the S-CMTs sample, as confirmed by the I type isotherms. According to the BET characterization, S-CMTs has a specific surface area of around 307.6 m2 g−1, much larger than that of CMTs (3.3 m2 g−1). The corresponding pore size distribution data (Figure 3d) unveils the size of the majority of pores of S-CMTs locate between 0.8 nm and 1.1 nm.



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More structural information of S-CMTs and CMTs is achieved by Raman and FT-IR characterizations. In comparison with the Raman spectrum of CMTs, S-CMTs shows additional visible peaks at 388, 468, and 1709 cm−1, displayed in Figure 3e. The peak at 468 cm−1 is ascribed to the stretching vibration of the S−O bonds, and those at 388 and 1709 cm−1 are related to the stretching vibration and deformation of the C−S bonds,47 suggesting that some carbon atoms are substituted by sulfur atoms in the carbon framework. Besides, an apparent increase in D band and reduction in G band noticed in S-CMTs compared with CMTs shows the n-type doping of S-CMTs, which is able to considerably increase the electrical conductivity of S-CMTs.48 The conductivities of S-CMTs and CMTs are determined to be 35.5 and 1.96 mS cm−1, respectively, corroborating an obvious improvement of conductivity in the S-CMTs sample (Figure S6). The FT-IR adsorption peaks (Figure 3f) of S-CMTs located at 1344 and 881 cm−1 can be correlated with the stretching vibration of C−S and S−O bonds,47, 49 respectively, in accordance with the XPS and Raman analysis. Furthermore, the 1167 cm−1 band is correlated with the stretching vibration of C−O bonds, indicating that cotton is entirely carbonized to form hard carbon with many surface defects.50


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Figure 4. CV profiles of (a) CMTs and (b) S-CMTs electrodes in the voltage range of 0.01−3 V at a scan rate of 0.1 mV s−1. Charge−discharge profiles of (c) CMTs and (d) S-CMTs electrodes at 200 mA g−1. The sodium storage performances of CMTs and S-CMTs were studied by CV and galvanostatic discharge−charge testing. Figure 4a,b depict the CV profiles of the CMTs and S-CMTs electrodes for the initial five cycles. As to the CMTs electrode, a shoulder peak is found at around 0.29 V in the first cathodic scan which vanishes in the following scans. This can be attributed to the side reaction of sodium with oxygen-containing groups and the generation of solid electrolyte interface (SEI) film. The broad peaks near 0.64 V in the subsequent cathodic scans are originated from to the reactions of sodium with edge/defect sites on the CMTs surface, whereas the redox couples at 0.02/0.19 V can be ascribed to Na+ intercalation/extraction into/from turbostratic nanocrystallites. In comparison with CMTs, the CV profiles of S-CMTs (Figure 4b) demonstrate overlapped redox couples located at 1.09/1.76 V, which is able to be assigned to the redox reactions between sodium and sulfur.42 To further verify that sulfur is chemically bonded rather than physically adsorbed in the S-CMTs sample, we also studied the CV performance of S/CMTs (Figures S7 and S8). The CV curves of S/CMTs show that there are two weak redox peaks at 1.55/1.61 V and 1.96/2.16 V, and the two pairs of peaks gradually weaken to disappear during the subsequent scans. This is because the sulfur physically adsorbed on the surface of CMTs is gradually dissolved into the electrolyte as a polysulfide compound upon cycling. In contrast, a couple of stable redox peaks in the CV curve of S-CMTs further demonstrates that sulfur in S-CMTs is chemically linked to the carbon network rather than physically adsorbed. The results above suggest that the chemically bonded sulfur can provide an extra stable storage of Na,51 thus greatly improving the reversible capacity. Figure 4c,d represent the discharge and charge curves of the CMTs and S-CMTs electrodes. The CMTs electrode delivers a first-cycle charge and discharge capacity of 146 and 270 mAh g−1 at 200 mA g−1



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(Figure 4c), respectively, similar to the previously reported sodium storage properties in disordered carbons.6, 52 Apparently, the S-CMTs electrode exhibits an exceedingly large initial charge and discharge capacity of 532 and 850 mAh g−1 (Figure 4d), respectively, showing a remarkable charge capacity improvement of nearly 4 times larger than that of the CMTs electrode. Also, the initial Coulombic efficiency (CE) of the S-CMTs electrode (62.6%) is much larger than that of the CMTs electrode (54.1%), which is probably correlated with the relatively lower oxygen content of S-CMTs than that of CMTs (14.7 vs 16.9 wt %, Table S1) because oxygen-containing functional groups usually bring about many side reactions during the first cycle.48 The large charge capacity of about 530 mAh g−1 is stably preserved in the ensuing cycles for the S-CMTs anode. The voltage curve can be divided into two sloped sections: a high-potential section (0.83−3 V) and a low-potential section (lower than 0.83 V). The former is assigned to the reaction between the strongly coupled sulfur and sodium, which coincides well with the CV observations (Figure 4b). The latter resembles the voltage curve of the CMTs electrode (Figure 4c) and is related to the reversible sodium-ion storage at defects and intercalation between graphene sheets of turbostratic nanocrystallites, which devotes a discharge capacity of around 380 mAh g−1 the second cycle (Figure 4d). This value is apparently larger than that contributed by the CMTs electrode (Figure 4c) or the S/CMTs electrode (Figure S9) in the same voltage region (380 vs 153 or 308 mAh g−1), suggesting that the sodium storage property of CMTs is significantly boosted by sulfur doping. Notably, the evident capacity elevation is gained in both sections for S-CMTs, which reveals that the H2S sulfurization process not only offers extra sodium-ion storage sites via high sulfur doping, but also enhances the sodium storage performance of CMTs by enlarging the surface area, increasing the electrical conductivity, and widening the interplanar distance of the (002) plane, leading to a significant improvement of reversible capacity.


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Figure 5. High-resolution S 2p XPS spectra of S-CMTs at the (a) fully sodiated and (b) fully desodiated states. (c) Rate capability of S-CMTs, CMTs, and S/CMTs at different current rates from 0.2 to 10 A g−1. (d) Nyquist plots of the S-CMTs, CMTs, and S/CMTs electrodes after the first cycle. Inset is the simulated equivalent circuit. In order to further unravel the extra sodium storage of the chemically coupled sulfur, S 2p XPS measurements were conducted on the fully sodiated and desodiated S-CMTs electrodes. As displayed in Figure 5a,b, at the fully sodiated status of 0.01 V (Figure 5a), the binding energies of the S 2p1/2 and S 2p3/2 peaks move to lower values of 163.3 and 162.2 eV in comparison with the as-fabricated S-CMTs (164.0 and 162.8 eV, Figure 3b). The lower binding energies of sulfur in S-CMTs at the fully discharged state is resulted from the strong interaction of sodium and sulfur. When fully desodiated to 3 V (Figure 5b), the S 2p1/2 and S 2p3/2 peaks obviously increase to 163.7 and 162.6 eV (still less than the original 164.0 and 162.8 eV), respectively, suggesting the fractional irreversible reaction between Na and S in SCMTs. Moreover, three additional high binding energy peaks at 169.8, 168.6 and 167.5 eV can also be



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easily seen at the completely charged and discharged states, which may be originated from the SEI layer.53 The variation of binding energy of S 2p of S-CMTs indicates that the electrochemical reaction between S and Na includes the C−S bonds weakening and strengthening, which can be further confirmed by Raman spectra. As demonstrated in the Raman spectra in Figure S10, when totally sodiated to 0.01 V, the peaks at 339 cm−1 (C−S bond) and 438 cm−1 (S−S bond) disappear and then reemerge at the totally charged status (3 V), evidencing the weakening and strengthening of the C−S bonds. Apart from the greatly improved reversible capacity, the S-CMTs electrode also manifests excellent rate capability and superior cycling lifespan. Figure 5c shows the specific capacity of S-CMTs at various current rates from 0.2 to 10 A g−1. As the current density increases from 0.2 to 0.5, 1, 2, and 5 A g−1, the reversible capacity changes from 524 to 366, 265, 234, and 199 mAh g−1, respectively, showing a high specific capacity retention. Even at the current density of 10 A g−1, a high charge capacity of 140 mAh g−1 is still retained by S-CMTs. In contrast, the CMTs electrode and the S/CMTs electrode present inferior rate capability, and the charge capacity quickly reduces to lower than 10 mAh g−1 when the current rate is increased to 10 A g−1. After 60 cycles at different current densities, S-CMTs recovers to 491 mAh g−1 when the current rate is returned to 0.2 A g−1, showing a remarkable high-rate performance. Remarkably, the superior rate property of the S-CMTs electrode stems from the unique microtube structure with enriched micropores, which is favorable for the rapid charge transfer and Na+ diffusion by diminishing the diffusion length. EIS analysis (Figure 5d) reveals that S-CMTs has a smaller SEI layer resistance (RSEI, 231 Ω) and charge transfer resistance (Rct, 65 Ω) than those of the CMTs electrode (322 Ω and 77 Ω) or the S/CMTs electrode (242 Ω and 71 Ω) on the basis of the equivalent circuit simulation (inset in Figure 5d), respectively. This implies that the S-CMTs electrode possesses a thinner SEI layer, which facilitates fast sodium-ion uptake/release and charge transfer at the electrolyte/electrode interphase. Furthermore, sulfur doping endows S-CMTs with improved electrical conductivity, which


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may be attributed to the C−S bonds in the carbon network affording another route for the electron migration and ensuring the uninterrupted and fast electronic mobility.54 Furthermore, the cycling stability of S-CMTs was investigated at a current rate as high as 1 A g−1 displayed in Figure S11. A charge capacity of 281 mAh g−1 is achieved even after 1000 cycles with a 90% capacity retention, corresponding to a capacity attenuation of only 0.011% per cycle. The CE is close to 100% after the initial cycles, suggesting stable cycling reversibility. It is clear that the cycling property of the S-CMTs electrode is much better than that of the CMTs electrode and the S/CMTs electrode. As shown in Figure S12, the SEM observations of the S-CMTs electrode before cycling and after 1000 cycles at 1 A g−1 unveil that there is no apparent morphology variation after long cycling test at the high current density, indicating a sturdy structure of S-CMTs.

Figure 6. (a) Charge−discharge curves at 200 mA g−1, (b) rate performance, and (c) cycling stability of the NVP@C-BN//S-CMTs full cell.



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Sodium-ion full cells based on S-CMTs and NVP@C-BN (Figure S13) were fabricated to illustrate the full-battery performance of S-CMTs, as demonstrated in Figure 6. The NVP@C-BN cathode was precycled to eliminate the irreversible capacity in the initial charge/discharge processes shown in Figure S14. NVP@C-BN is revealed to be of high charge/discharge rate property (Figure S15) and stable cyclability (Figure S16), which is excellent to match with the S-CMTs anode. Figure 6a depicts the charge−discharge curves of the NVP@C-BN//S-CMTs full cell at 200 mA g−1 with cutoff voltages of 3.6 and 0.3 V. The charge and discharge capacities of the second cycle are 326 and 320 mAh g−1, respectively (normalized with the weight of S-CMTs). After 10 cycles, the full cell can retain a stable capacity of 315 mAh g−1. Apart from the large specific capacity and outstanding capacity retention, the full battery also shows superior rate performance. As demonstrated in Figure 6b, the full battery exhibits an average discharge capacity of 325, 298, 249, 157, and 45 mAh g−1 at 0.2, 0.5, 1, 2, and 5 A g−1, respectively. When the current rate is decreased to 0.2 A g−1, the reversible capacity quickly restores to 310 mAh g−1, demonstrating the remarkable tolerance to the fast Na+ uptake/release. In Figure 6c, the full cell presents an initial charge capacity of 370 mAh g−1 with a large CE of around 85.5% at 200 mA g−1, and the mean CE is larger than 99% throughout the cycling test. The high capacity retention of >81% after 100 cycles shows good cycling stability of SIBs on the basis of NVP@C-BN and S-CMTs. As shown in Table S2, such excellent electrochemical properties has rarely been obtained in previously reported carbon-based SIB anode materials. The advanced sodium storage activity of the sulfur-doped CMTs is mainly benefited from the synergistic effect of the high sulfur doping and the abundant microporous structure. First, the sulfur atoms chemically linked to the pyrolyzed CMTs are easy to store Na, resulting in high specific capacity. In addition, an expanded interplanar distance in S-CMTs warrants more efficient Na+ uptake and release, therefore boosting the sodium storage performance. Second, the unique microtube structure with an interpenetrating micropores endues the S-CMTs with additional sites


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for sodium-ion storage and favors the transportation of sodium-ions and electrons inside the electrode. Last, the electrical conductivity of CMTs is greatly enhanced via sulfur doping, leading to significantly increased sodium storage capability. 4. CONCLUSIONS In summary, advanced sulfur-doped carbon microtubes have been synthesized through hightemperature sulfurization of cotton roll. Due to the high sulfur doping and the unusual microtubes with microporous architecture, the as-synthesized sulfur-doped carbon microtubes shows a high charge capacity (532 mAh g−1 at 0.2 A g−1), superior rate performance (234 mAh g−1 at 2 A g−1), and outstanding cycling property (281 mAh g−1 at 1 A g−1 after 1000 cycles with a 90% capacity retention). A full cell with the NVP@C-BN cathode and the S-CMTs anode further verifies its excellent electrochemical performance by demonstrating high reversible capacity (370 mAh g−1 at 200 mA g−1) and long cycling stability (>81% after 100 cycles). It is expected that this facile S-doping strategy can also be used to construct other efficient biomass-derived SIB anode materials. ASSOCIATED CONTENT Supporting Information TGA curve, elemental analysis, HRTEM images, XRD patterns, S 2p XPS spectra, photographs of the measurement of resistivity, SEM images, CV curves, charge−discharge profiles, Raman spectra, long cycling performance, material characterization, rate capability, and comparison of sodium storage properties. This material is available free of charge via the Internet at http://pubs.acs.org. AUTHOR INFORMATION Corresponding Author *E-mail: [email protected].



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*E-mail: [email protected]. ORCID Xiaosi Zhou: 0000-0001-9641-7166 Author Contributions †Q.W.

and X.G. contributed equally to this work.

Notes The authors declare no competing financial interest. ACKNOWLEDGMENT This work was supported by the National Natural Science Foundation of China (Grant Nos. 51577094 and 21503112), the Natural Science Foundation of Jiangsu Province of China (BK20180086), and the 100 Talents Program of Nanjing Normal University. REFERENCES (1) Dunn, B.; Kamath, H.; Tarascon, J.-M. Electrical Energy Storage for the Grid: A Battery of Choices. Science 2011, 334, 928−935. (2) Xiang, X.; Zhang, K.; Chen, J. Recent Advances and Prospects of Cathode Materials for Sodium-Ion Batteries. Adv. Mater. 2015, 27, 5343−5364. (3) Fan, X.; Mao, J.; Zhu, Y.; Luo, C.; Suo, L.; Gao, T.; Han, F.; Liou, S.-C.; Wang, C. Superior Stable Self-Healing SnP3 Anode for Sodium-Ion Batteries. Adv. Energy Mater. 2015, 5, 1500174. (4) Yabuuchi, N.; Kubota, K.; Dahbi, M.; Komaba, S. Research Development on Sodium-Ion Batteries. Chem. Rev. 2014, 114, 11636−11682. (5) Li, Z.; Chen, Y.; Jian, Z.; Jiang, H.; Razink, J. J.; Stickle, W. F.; Neuefeind, J. C.; Ji, X. Defective Hard Carbon Anode for Na-Ion Batteries. Chem. Mater. 2018, 30, 4536−4542.


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Fabrication of Microporous Sulfur-Doped Carbon Microtubes for High

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