Dual Carbon-Confined SnO2 Hollow Nanospheres ... - ACS Publications

Research Article Cite This: ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX

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Dual Carbon-Confined SnO2 Hollow Nanospheres Enabling High Performance for the Reversible Storage of Alkali Metal Ions Qiong Wu, Qi Shao, Qiang Li, Qian Duan, Yanhui Li, and Heng-guo Wang* School of Materials Science and Engineering, Changchun University of Science and Technology, Changchun 130022, P. R. China S Supporting Information *

ABSTRACT: To explore a universal electrode material for the highperformance electrochemical storage of Li+, Na+, and K+ ions remains a big challenge. Herein, we propose a “trinity” strategy to coat the SnO2 hollow nanospheres using the dual carbon layer from the polydopamine-derived nitrogen-doped carbon and graphene. Thereinto, hollow structures with sufficient void space could buffer the volume expansion, whereas dual carbonconfined strategy could not only elastically prevent the aggregation of nanoparticle and ensure the structural integrity but also immensely improve the conductivity and endow high rate properties. Benefiting from the effective strategy and specific structure, the dual carbon-confined SnO2 hollow nanosphere (denoted as G@C@SnO2) can serve as the universal host material for alkali metal ions and enable their rapid and reversible storage. As expected, the resulting G@C@SnO2 as a universal anode material shows reversible alkalimetal-ion storage with high performance. We believe this that strategy could pave the way for constructing other metal-oxide-based dual carbon-confined high-performance materials for the future energy storage applications. KEYWORDS: nitrogen dope, dual carbon layer, graphene, SnO2, alkali metal batteries

1. INTRODUCTION Alkali metal batteries have been or will be the efficient energy storage systems with practical or potential significance.1−4 As the best-known example, lithium-ion batteries (LIBs), which are dominant in today’s portable electronic products, are attracting increasing attention as energy sources for hybrid and electric vehicles.5−7 However, the future of LIBs is shadowed by the low natural storage of lithium resources and the increasing price of lithium year by year.8 By contrast, sodium-ion batteries (SIBs) or potassium-ion batteries (KIBs) have attracted more and more attention as promising alternatives in view of their low cost and abundant resource.9,10 Nevertheless, the larger sizes of Na+ and K+ result in much inferior capacities and problematic cycle life compared with their lithium-ion counterparts.11,12 Therefore, it is a necessary to develop the applicable electrode material that can store multiple types of alkali ions, thus speeding up the commercialization process of alkali metal batteries.13,14 Recent studies of most anodes for alkali metal batteries mainly focus on the carbon-based materials.15,16 However, larger ionic diameter limits the number of Na+ and K+ that can insert the intercalation-type electrode materials.17−19 One effective strategy to solve this obstacle is to use the alloying/ dealloying reaction, which could improve the storage capacity of alkali metal ions.20 Typically, tin-based materials have been considered as prospective electrode materials for alkali metal batteries.21,22 Of them, tin dioxide (SnO2) shows many fascinating advantages such as abundance and environmental © XXXX American Chemical Society

friendliness, especially the high theoretical capacity resulted from the alloying and pending conversion mechanism.23,24 Unfortunately, SnO2 usually suffers from the huge volume changes (∼300%) upon the charge/discharge cycling,25,26 let alone those of Na+ and K+ with a larger ionic diameter. This shortcoming could lead to the pulverization and serious aggregation of the SnO 2 , thus showing poor cycling stability.27−29 Furthermore, SnO2 shows inherent low electronic conductivity as a semiconductor, which also affects its electrochemical performance.30,31 To resolve these problems, carbon-confined SnO2 hollow nanospheres are constructed based on the stable structure as well as the high electronic conductivity of carbon materials.32,33 However, the conductive network between nanoparticles is a “plane-to-point” conductive mode.34 In contrast, two-dimensional graphene has been used as the most promising matrix to support metal-oxide nanoparticles because of the intrinsically outstanding electrical conductivity, good mechanical flexibility, and the open and porous structure. Moreover, the incorporation of graphene could construct three-dimensional (3D) hierarchical conductive network. On the one hand, this structure could effectively accommodate the mechanical stress and further improve the structural integrity, thus leading to the stable cyclability. On the other hand, it could achieve a plane-to-point conductive mode, Received: January 12, 2018 Accepted: April 17, 2018

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DOI: 10.1021/acsami.8b00605 ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX

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ACS Applied Materials & Interfaces which effectively ensures fast electron-transport pathways, thus enhancing the rate capability.35,36 However, the exposed metaloxide nanoparticles directly loaded on the surface of graphene make the occurrence of the aggregation still easy upon the cycling process, which can lead to the declined cycling performance.37,38 As a response, constructing graphene-based dual carbon-confined metal-oxide nanoparticles could be a promising electrode configuration, which could enable high performance for alkali metal batteries. Herein, the nitrogen-doped dual carbon-confined strategy was used to coat the SnO2 hollow nanosphere (denoted as G@ C@SnO2), which was employed as a universal anode material for LIBs, SIBs, and KIBs. Such a design offers a variety of advantages: (i) the hollow SnO2 nanospheres could buffer the volume change and short ions diffusion length; (ii) the nitrogen-doped carbon shell can further elastically accommodate volume variation, ensure the structural integrity, and enhance the conductivity; (iii) constructing a graphene-based 3D structure can further enhance the conductivity, prevent the aggregation of nanoparticle, and provide an open framework for the transportation of electrons and ions. As expected, the synergistic effect from this “trinity” strategy endows G@C@ SnO2 with outstanding electrochemical properties as a fascinating anode material of alkali-metal-ion batteries (LIBs, SIBs, and KIBs).

Figure 1. Schematic diagram for the synthetic strategy of G@C@ SnO2.

introduced, the GO@PDA@SnO2 composite is formed. Finally, the carbonization process is utilized to convert the PDA layer into a nitrogen-doped carbon layer and GO into graphene, thus constructing nitrogen-doped dual carbonconfined SnO2 hollow nanospheres. Scanning electron microscopy (SEM) and transmission electron microscopy (TEM) images are used to investigate the microstructures and morphologies of various samples. As shown in Figure 2a, the prepared SiO2 has discrete and uniform spherical structure with a particle size of 140−170 nm. Figure 2b clearly demonstrates that the SnO2 is monodisperse, which exhibits a highly uniform spherical shape with a particle size of 200−240 nm. It’s indicated from the red arrow that the SnO2 nanospheres are hollow. Furthermore, the G@C@SnO2 shows crumpled graphene sheets with small nanoparticles (Figure S1a,b). Figure 2c further shows that SnO2 nanospheres are tightly wrapped and connected by the wrinkled graphene nanosheets. The TEM image further demonstrates that the hollow SnO2 nanospheres are coated by the carbon layer and the graphene layer (Figure 2d). It can be found that the inner SnO2 layer is about 39 nm, and intermediate carbon layer is about 23 nm (Figure 2e). The void volume is calculated to be about 63.6% in the SnO2, which can shorten the ion diffusion length and buffer the volume change. The High-resolution TEM (HRTEM) image in Figure 2f also reveals this dual carbon-confined nanostructure, in which two kinds of lattice fringes with the space of 0.33 and 0.26 nm can be attributed to the (110) and (101) facets of SnO2, whereas outer nitrogendoped carbon is not well crystallized. To prove the crystallographic structure of SnO2 and G@C@ SnO2, X-ray diffraction (XRD) patterns are obtained. The peak at 23° confirms the formation of the amorphous silica (Figure S2a). As shown in Figure 3a, the pattern of SnO2 exhibits three characteristic diffraction peaks at 26.6°, 33.9°, 51.8°, which could be attributed to the (110), (101), and (211) planes, indicating the existence of the tetragonal rutile SnO2 phase (JCPDS no. 41-1445).36 Furthermore, XRD peaks of PDA@ SnO2 and G@C@SnO2 are identical with that of the pristine SnO2, indicating that SnO2 is successfully coated into the carbon composite. To analyze the local structure of SnO2 and G@C@SnO2, Raman spectra are carried out (Figure 3b). The peak at about 621 cm−1 confirms the presence of SnO2.39 In addition, the G@C@SnO2 displays the two broad peaks at 1353 and 1589 cm−1, corresponding to the characteristic D

2. EXPERIMENTAL SECTION 2.1. Preparation of SnO2 Hollow Nanospheres. Typically, 0.045 g of urea and 0.15 g of K2SnO3·H2O were dissolved in the suspension solution with 0.05 g of SiO2, 3 mL of deionized water, and 3 mL of ethanol. Then, the mixture solution was added in a Teflonlined stainless-steel autoclave, which was heated to 150 °C for 24 h. Subsequently, the as-prepared sample was obtained by centrifuging and washing, and then dispersed in 2 mol L−1 NaOH solution. After stirring at 50 °C for 6 h, the samples were washed to neutral and then dried under vacuum. 2.2. Preparation of G@C@SnO2. First, 0.6 g of Tris and 14.7 mL of 0.1 mol L−1 HCl were dissolved in 50 mL of deionized water. Afterward, 60 mg of SnO2 hollow nanospheres and 72 mg of polydopamine (PDA) were successively dispersed in the above solution and then stirred for 4 h, the as-prepared PDA@SnO2 was obtained by centrifuging and washing with deionized water and ethanol. Subsequently, the as-obtained PDA@SnO2 was redispersed in graphene oxide (GO) solution (1.4 mg mL−1) and sonicated for 4 h. Finally, the mixture was freeze-dried overnight and then annealed at 600 °C in N2 for 3 h using a heating rate of 5 °C min−1. For comparison, the G@SnO2 was prepared without adding PDA, C@ SnO2 was prepared without adding GO. In addition, G@C@SnO2-2, G@C@SnO2, and G@C@SnO2-8 were prepared by adding different ratios of PDA@SnO2 and reduced GO (SnO2/GO weight ratio 2:1, 4:1, and 8:1).

3. RESULTS AND DISCUSSION The overall fabrication process for nitrogen-doped dual carbonconfined SnO2 hollow nanospheres (denoted as G@C@SnO2) is schematically shown in Figure 1. First, monodisperse SiO2 nanospheres are employed as the hard template to synthesize SnO2 hollow nanospheres through a hydrothermal reaction. Then, the polymerization process is utilized to coat SnO2 hollow spheres with PDA. On the one hand, as the key nitrogen-rich polymer, PDA can be converted into nitrogendoped carbons by pyrolysis treatment; on the other hand, its amino within the polymer skeleton could promote the interaction between PDA and GO because of the hydrogen bond and van der Waals forces. Accordingly, when GO is B


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Figure 2. SEM images of SiO2 (a), hollow SnO2 (b), and G@C@SnO2 (c); TEM (d,e) and HRTEM (f) images of G@C@SnO2.

Figure 3. (a) XRD patterns and (b) Raman spectra of SnO2 and G@C@SnO2; (c) TG curve of G@C@SnO2; and (d) nitrogen adsorption− desorption isotherm and corresponding pore-size distribution (inset) of G@C@SnO2.

desorption isotherm measurement is also carried out. It can be seen clearly from Figure 3d that the measured isotherm is the type IV isotherm curve, indicating the presence of a mesoporous structure in the G@C@SnO2 composite. Meanwhile, the corresponding pore-size distribution is analyzed using the Barrett−Joyner−Halenda method,41 which reveals that the pore sizes switch from 0.65 to 8.8 nm and the average pore diameter is 4.65 nm. Especially, the G@C@SnO2 shows the Brunauer−Emmett−Teller specific surface area up to 181.166 m2 g−1. The large surface area could provide larger contact area to promote the diffusion of alkali metal ions. Meanwhile, the presence of the mesoporous structure could release the volume effect of SnO2 upon cycling. X-ray photoelectron spectroscopy (XPS) is performed to explore the surface composition and chemical states of the G@ C@SnO2. The wide-survey XPS spectrum (Figure 4a) discloses the coexistence of Sn, C, N, and O. Figure 4b shows the highresolution Sn 3d spectrum, in which two peaks located at 487.2 and 495.6 eV could be corresponded to Sn 3d5/2 and Sn 3d3/2

band and G band of disordered graphite or crystal defects and graphitic carbon, respectively.31 The Raman ID/IG ratio is calculated as 0.96, confirming the reduction from GO to graphene. To demonstrate the chemical structure of asprepared samples, Fourier transform infrared (FT-IR) spectra are studied (Figure S2b), in which the absorption peaks of the PDA at 1267, 1485, and 1640 cm−1 are ascribed to the C−O, C−H, and CN bonds.40 After the carbonization, the characteristic peaks (C−O, CN) disappear, indicating the evolution from PDA to the carbon materials. To quantify the SnO2 content in the composites, thermal gravimetric analysis is carried out (Figure 3c). The initial weight loss between 100 and 300 °C may be resulted from the removal of residue moisture. The larger weight loss occurs between 300 and 600 °C, which results from the decomposition of PDA and GO. The residual weight above 600 °C is the weight of the SnO2 in the G@C@ SnO2. The weight percentage of SnO2 in the composite could be evaluated to be 62.1%. To study the porous structure and surface area of the G@C@SnO2 composite, N2 adsorption− C


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Figure 4. (a) Survey XPS spectrum and high-resolution XPS spectra of (b) Sn 3d, (c) C 1s, and (d) N 1s in G@C@SnO2.

Figure 5. Electrochemical performance for Li+ storage: (a) CV curves of G@C@SnO2 at a scan rate of 0.1 mV s−1; (b) discharge/charge curves of G@C@SnO2 at 0.1 A g−1; (c) cycling performance of various samples at 0.1 A g−1; (d) cycling performance of G@C@SnO2 at 1 A g−1; and (e) rate capability of G@C@SnO2 at varied current densities.

spin or bit peaks of SnO2.30 Figure 4c shows the high-resolution C 1s spectrum, in which five peaks located at 283.7, 284.6, 285.1, 286.3, and 288.1 eV can be corresponded to the carbon atoms of Sn−C, C−C, C−N, Sn−O−C, and CO, respectively.42,43 Meanwhile, the peak of CO is quite weak, which indicates the reduction of GO. The high-resolution O 1s spectrum (Figure S3) is fitted with three peaks at 530.9, 531.8, and 532.7 eV, corresponding to the O 1s of SnO2, CO, COO or C−O, and H2O bonds, respectively.44 As shown in Figure 4d, high-resolution N 1s peaks are fitted to three peaks at 398.4, 400.1, and 401.2 eV, which could be ascribed to pyridinic-N,

pyrrolic-N, and graphitic-N, respectively.45 The reported results have demonstrated that nitrogen doping (especially pyridinicN) could generate favorable defect sites to reserve more Li+ ion and further increase the conductivity of carbon-based materials, thus showing the improved capacity and rate performance.46,47 The lithium storage properties of G@C@SnO2 as the anode material of LIBs is first explored. Cyclic voltammetry (CV) curves are first shown (Figure 5a). During the first cathodic scan, there is a peak located at ∼0.96 V, which is related with the formation of the solid-electrolyte interphase (SEI) layer, the decomposition of electrolyte, and the conversion from SnO2 to D


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Figure 6. Full LIBs with LiCoO2 as the cathode and G@C@SnO2 as the anode (based on anode mass). (a) Discharge/charge curves and (b) cycling performance at 0.1 A g−1, (inset) photograph of a LED powered by the full cell. (c) Schematic illustration of reversible Li+ ion storage in G@C@ SnO2.

Sn.25 Upon further charging, there is a steep reduction peak located at 0.2 V, which is related with the formation of LixSn through the lithium alloying reaction with Sn. In the following anodic scan, the oxidation peak located at 0.6 V could be attributed to the dealloying from LixSn to Sn. Meanwhile, an obvious and broad oxidation peak located at 1.2 V could be ascribed to the partial reversible formation from Sn to SnO2.30 Interestingly, the following CV curves are clearly overlapped, demonstrating good stability and reversibility of G@C@SnO2 during the Li+ insertion/extraction process. The similar results could be found for the CV curves of SnO2 and C@SnO2 (Figure S4). As shown in the representative discharge/charge curves (Figure 5b), two voltage plateaus appear during the first discharge process, which may result from the conversion from SnO2 to Sn as well as the formation of the SEI layer. During the charge process, voltage plateaus can be attributed to the Li+ insertion to form LixSn, which shows the same results with the CV curves. Moreover, the G@C@SnO2 delivers the initial discharge/charge capacities of 1363.5 and 945.9 mA h g−1, respectively. The charge capacity is near the theoretic capacity of the G@C@SnO2, which is calculated as 1068 mA h g−1 according to the content of 62.1% in the composite, indicating higher lithium storage capacity. Figure 5c shows the Coulombic efficiency of G@C@SnO2 and compares the cyclic performance of various samples. The initial Coulombic efficiency is 69.3%, which resulted from the formation of the SEI film, whereas following Coulombic efficiency increases to 95% and then remains at more than 99.8% throughout the 240 cycles. Interestingly, the G@C@SnO2 shows slightly capacity fade in its initial 25 cycles and then a high capacity of 1156.1 mA h g−1 after 240 cycles remained. Similar results have been obtained for other SnO2-based electrodes, which results from the reversible formation of a polymeric gel-like layer via electrolyte decomposition and the improvement of lithium-ion accessibility upon deep cycling.26,29 To obtain the optimum condition, various G@C@SnO2 samples with different SnO2 contents are investigated (Figure S5). It is obvious that G@C@SnO2 shows the best capacity and cycling performance among various G@ C@SnO2 samples. In addition, to show the superiority of our strategy, various samples with different components are also

compared (Figure 5c). It is obvious that the G@C@SnO2 shows higher reversible capacity and longer cycling stability than those of SnO2 and C@SnO2. To disclose the improved electrochemical properties of the G@C@SnO2, the morphology of these samples is investigated by using SEM (Figure S6). It can be seen clearly that the SnO2 and C@SnO2 are pulverized into small particles. However, the integrity of the G@C@SnO2 is generally preserved, which is further confirmed by TEM images (Figure 9c). Furthermore, the G@C@SnO2 also shows the better cycling stability than that of the G@SnO2 (Figure S7). These results demonstrate that the dual carbonconfined strategy could buffer the volume change of SnO2 during cycling, thus showing improved cycling stability. The high-rate cycling stability is also investigated at 1 A g−1 (Figure 5d). In the first five cycles, the lower current density of 0.05 A g−1 is used to generate dense SEI layers. Then, in the subsequent cycles, the high current density of 1 A g−1 is applied. It can be observed that the high reversible capacity of 487.5 mA h g−1 can still remain after 660 cycles, indicating excellent cycle performance. Moreover, the rate properties of various samples are further studied (Figure 5e). The G@C@ SnO2 exhibits the reversible capacities of 1009.7, 829.4, 767, 685.1, and 648.3 mA h g−1 at the current densities of 0.05, 0.1, 0.2, 0.5, and 1 A g−1, respectively. Interestingly, it can show a higher capacity of 614.6 mA h g−1 at the high current density of 2 A g−1. Most importantly, a specific capacity of 1054.7 mA h g−1 is regained when the current density is recovered to 0.05 A g−1. In contrast, other samples show the poor rate performance (Figure S8). All these results demonstrate that G@C@SnO2 prepared using our strategy shows outstanding cycling stability and high rate properties for LIBs. In view of the excellent performances in half cells, a full LIB with commercial LiCoO2 as the cathode and G@C@SnO2 as the anode is investigated with a potential range of 1.2−4.2 V. Figure 6a shows the charge and discharge curves with obvious discharge plateaus at 3.5 and 2.6 V, corresponding to the conversion from SnO2 to Sn.48 Figure 6b shows the discharge capacity of 345.8 mA h g−1 after 90 cycles. Moreover, the energy output of this battery could power a light-emitting diode (LED) (the inset of Figure 6b). The excellent electrochemical E


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Figure 7. Electrochemical performance for Na+ storage: (a) CV curves of G@C@SnO2 at a scan rate of 0.1 mV s−1; (b) discharge/charge curves of G@C@SnO2 at 0.025 A g−1; (c) cycling performance of G@C@SnO2 at 0.05 A g−1; and (d) rate capability of G@C@SnO2 at varied current densities.

Figure 8. Electrochemical performance for K+ storage: (a) CV curves of G@C@SnO2 at a scan rate of 0.1 mV s−1; (b) XRD patterns after first discharge and charge processes; (c) cycling performance of G@C@SnO2 at 0.05 A g−1; and (d) rate capability of G@C@SnO2 at varied current densities.

properties of G@C@SnO2 for Li+ storage result from the synergistic effects of components and their predominant structural advantages (Figure 6c). To improve the poor cycling stability, we adopt two approaches for accommodating the enormous volume changes of SnO2 during the Li+ insertion/ extraction process: on the one hand, SnO2 hollow nanospheres provide sufficient void space for internally relieving the large volume changes; on the other hand, the dual carbon-confined

strategy could not only elastically ensure the structural integrity but also effectively prevent the aggregation of nanoparticles, thus exteriorly relieving the large volume changes. To enhance the rate performance, in other words, enhance the electronic conductivity of SnO2 semiconductor, nitrogen-doped dual carbon-confined strategy is adopted. Recent reports have demonstrated that the nitrogen dope could form n-type conductive materials, which can further improve the conF


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Figure 9. (a) EIS spectra of G@C@SnO2 for LIBs, SIBs, and KIBs and the equivalent circuit model (inset); (b) dependence of Zre on the reciprocal square root of the frequency in the low-frequency region of G@C@SnO2 for LIBs, SIBs and KIBs; TEM images of the G@C@SnO2 at 0.1 A g−1 after 20 cycles in LIBs (c), SIBs (d), and KIBs (e).

rate properties with the reversible capacities of 299, 274, 248, 215, 191, and 165 mA h g−1 at the current densities of 0.05, 0.1, 0.2, 0.5, 1, and 2 A g−1, respectively. Interestingly, the discharge capacity of G@C@SnO2 could be recovered to 288 mA h g−1 when the current density is returned to 0.05 A g−1. All these results also demonstrate that G@C@SnO2 prepared using our strategy is suitable as an anode material for SIBs. The potassium storage property of G@C@SnO2 as the anode material for KIBs is explored. To investigate the potassium storage mechanism, the CV curves are investigated (Figure 8a). During the first cathodic scan, there is a reduction peak located at around 0.58, which may be ascribed to the formation of a SEI layer and conversion from SnO2 to Sn together with the formation of K2O. Furthermore, it can be observed that the broad peak is from 0.39 to 0.1 V, which could be corresponded to KxSn alloying processes. During the anodic scan, there are four peaks at 0.7, 0.83, 1.05, and 1.18 V, corresponding to the dealloying of KxSn.10 The corresponding slope and plateau can be also observed in the typical discharge/ charge curves (Figure S10a). To confirm the KxSn alloying process and the storage mechanism, XRD analysis after first discharge/charge process is performed (Figure 8b). After the first discharge to 0.01 V, it can be found that the XRD peaks of SnO2 completely disappear, which proves the KxSn alloying process. Meanwhile, there are number of diffraction peaks, which could be attributed to Sn (JCPDS card no. 04-0673), K2Sn5 (a phase is similar to a known lithium−tin alloy), and K4Sn23 (JCPDS card no. 04-004-8647).10 In the first charge to 3 V, the XRD peaks of the KxSn alloying completely disappear, which proves the KxSn dealloying process. Unfortunately, the coexistence of SnO2 and Sn phases indicates the partial oxidation from Sn to SnO2, which results in the low capacity of the G@C@SnO2 for KIBs. Figure 8c shows that the G@C@ SnO2 delivers the initial discharge/charge capacities of 708 and 195 mA h g−1 with a low Coulombic efficiency of 28% that resulted from the electrolyte decomposition. After 85 cycles, the G@C@SnO2 shows the higher reversible capacity of 147.8 mA h g−1 with high Coulombic efficiency near 98%. Interestingly, further increasing the current density, the G@

ductivity. In addition, graphene-based dual carbon-confined strategy not only immensely enhances the conductivity but also achieves a plane-to-point conductive mode, which effectively ensures fast electron-transport pathways. Furthermore, this dual carbon-coated architecture can effectively avoid the side reactions that resulted from the direct contact between the active materials with the electrolyte. Consequently, our proposed strategy of nitrogen-doped dual carbon-confined SnO2 hollow nanospheres endows the electrode of G@C@ SnO2 with extremely Li+ storage properties. The sodium storage property of G@C@SnO2 as the anode material for SIBs is investigated. First, CV curves are employed to study the sodium storage mechanism (Figure 7a). During the first cathodic scan, there is a peak at ∼1.1 V, which could be corresponded to the SEI layer formation, the electrolyte decomposition, and the reduction from SnO2 to Sn together with the formation of Na2O. Upon further charging, another broad peak between 0.65 and 0.37 V can be corresponded to the NaxSn alloying processes.25 During the anodic scan, it is obvious that the broad oxidation peak between 0.1 and 0.7 V could be possibly ascribed to the dealloying of NaxSn. Besides the first cycle, the following overlapped curves also indicate the good stability and reversibility of G@C@SnO2 during the Na+ insertion/extraction process. Figure 7b shows the typical charge/discharge profiles of G@C@SnO2 for SIBs. The initial discharge/charge capacities of 834.7 and 374.9 mA h g−1 with a Coulombic efficiency of 45% are shown. Similarly, they display an obvious voltage plateau, similar to the CV measurement. Figure 7c shows the cycling performance at 0.05 A g−1. First, it can be observed that the initial discharge/charge capacities of G@C@SnO2 are 684.4 and 298.2 mA h g−1, respectively. It can be found that the initial coulombic efficiency is 43%, it increases to 93% at the 2nd cycle, and then remains more than 98% in the subsequent cycles. The higher capacity of 277.4 mA h g−1 with the capacity retention of 93% is achieved after 100 cycles. Most importantly, upon increasing the current density to 0.1 A g−1, G@C@SnO2 could display the high reversible specific capacity of 178.6 mA h g−1 after 300 cycles with the capacity retention of 62% (Figure S9). Figure 7d displays the superior G


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ACS Applied Materials & Interfaces C@SnO2 can deliver the capacity of 122 mA h g−1 after 200 cycles (Figure S10b). Furthermore, the rate capability of G@ C@SnO2 is also tested (Figure 8d), whose average specific capacities are 320, 235, 201, 179, 150, 129, and 119 mA h g−1 at 0.01, 0.02, 0.05, 0.1, 0.2, 0.3, and 0.4 A g−1, respectively. A high reversible capacity of 241 mA h g−1 can be regained after the current density is returned to 0.01 A g−1, and then stable extended cycling is kept. From the results of the electrochemical performance for Li+, Na+, and K+ storage, it can be found that G@C@SnO2 shows much lower storage capacity in SIBs and KIBs than that in LIBs, which may be associated with the dramatic differences between sodiation/potassation and lithiation resulted from the different ionic radius of potassium (1.33 Å)/sodium (0.97 Å) compared with lithium ion (0.68 Å). For in-depth study of the different electrochemical performance of LIBs, SIBs, and KIBs, additional characteristics are further investigated and compared. It is well-known that galvanostatic intermittent titration technique (GITT) is obtained when the cell is operated under almost quasiequilibrium conditions. Therefore, GITT is used to confirm the maximum amount of Li+, Na+, and K+ ions that could reversibly insert/deinsert into the G@C@SnO2 electrode (Figure S11). First, it is found that the shape and the capacity from the GITT curves are very similar to those of the discharge/charge curves from continuous charge/discharge process, indicating that the obtained capacity under continuous charge/discharge processes is very close to the equilibrium.32,49 This result shows that the dual carbon-confined strategy endows the good ionic and electronic conductivities. The detailed mechanism and reaction kinetics of the G@C@SnO2 for LIBs, SIBs, and KIBs are further investigated using electrochemical impedance spectroscopy (EIS) (Figure 9a). The same equivalent circuit mode is used to fit all EIS plots (inset of Figure 9a), in which the inner resistance of cell (Rs) represents the Ohmic resistance and solution resistance, the charge-transfer impedance (Rct) represents the charge transfer resistance between the electrolyte interface and electrode, and the constant phase-angle element represents the double layer capacitance and the Warburg impedance (Zw) is the ion diffusion impedance in the active materials.50 In general, for a typical EIS Nyquist plot, the semicircle in medium-to-high frequency region could reflect the Rct on electrode−electrolyte interface, whereas the straight line in low-frequency region could be related with the ion-diffusion process within the electrodes (known as the Warburg behavior). Apparently, the Rct of G@C@SnO2 electrode increases gradually from 496, 1941.2 to 13 989 Ω for LIBs, SIBs, and KIBs, respectively (Figure 9a), indicating the G@C@SnO2 experiences larger interface impedance in SIBs and KIBs than that in LIBs. The small Rct of G@C@SnO2 in LIBs could be attributed to the differences in the ion radius of Li+, Na+, and K+ ions and the SEI film resistance of LIBs, SIBs, and KIBs. In fact, the formation mechanisms of SEI film for SIBs, especially for KIBs are complex; therefore, further investigation is currently in progress. In addition to the interface impedance, the ion diffusion impedance can evaluate the diffusion process of Li+, Na+, and K+ ions. According to the straight line, we could calculate the apparent ions diffusion coefficient (D) of the G@ C@SnO2 for LIBs, SIBs, and KIBs by the following eqs 1 and 251 D = (R2T 2)/(2A2 n 4F 4C 2σ 2)

Zre = R e + R ct + σω−1/2

(2)

According to eq 2, the slopes of the fitting line are 370.12, 987.65, and 1495.69 from the linear fitting for the lowfrequency region of the EIS plots (Figure 9b). Considering that the reaction mechanism of KxSn alloying cannot be determined, the average value of the K2Sn5 and K4Sn23 is used to estimate the D. According to eq 2, the D value could be calculated as 0.198 × 10−14, 0.376 × 10−15, and 0.137 × 10−15 cm2 s−1 for Li+, Na+, and K+ ions, respectively, which confirms that the transport of Na+ and K+ ions is much slower than that of Li+ ion. This impedance evolution is consistent with the electrochemical properties of the G@C@SnO2 for LIBs, SIBs, and KIBs. In addition to the difference in EIS of the G@C@SnO2 for different batteries, the detailed EIS tests of the G@C@SnO2 for LIBs, SIBs, or KIBs after different cycles are also performed (Figure S12). Apparently, with the increased cycles from 0, 10 to 20, the Rct of G@C@SnO2 for LIBs decreases gradually from 496 to 325 and 290 Ω, and the D values increase from 0.198 × 10−14 to 0.356 × 10−14 and 0.439 × 10−14 cm2 s−1, respectively. For SIBs, with the increased cycles, the Rct reduces gradually from 1941.2 to 463.8 and 320.9 Ω and the D values increase from 0.376 × 10−15 to 0.68 × 10−15 and 3.45 × 10−15 cm2 s−1, respectively. For KIBs, the Rct increases gradually from 13 989 to 10 099 and 6879 Ω, and the D values increase from 0.137 × 10−15 to 0.434 × 10−15 and 1.127 × 10−15 cm2 s−1, respectively. The decrease in Rct and D during cycling for LIBs, SIBs, and KIBs could be related with the increased active sites in the nitrogen-doped carbon shell and open framework of graphene. The dual carbon-confined SnO2 nanoparticles could achieve a plane-to-point conductive mode and an open framework, which effectively ensures fast electron-transport pathways and ion diffusion. To verify the structure stability, the morphology of G@C@SnO2 after 20 cycles in LIBs and SIBs and KIBs are also investigated (Figure 9c−e). Compared with the fresh G@C@ SnO2, the morphology has obvious changed, whereas the integrity of the dual carbon-confined SnO2 hollow nanospheres is generally preserved after cycling. For LIBs, no obvious cracking is observed, resulting in the high capacity and long cycling stability. However, for SIBs, especially of KIBs, cracking can be observed, which resulted from the destructive of the larger ionic radius of potassium/sodium. Fortunately, the broken SnO2 could be well encapsulated in the dual carbon layers, resulting in the good cycling stability to a certain extent. Certainly, the electrochemical performance of SIBs, especially KIBs still lags behind that of LIBs, therefore the reaction mechanisms need to be carefully investigated and the optimized electrolyte should be also thoroughly investigated.

4. CONCLUSION In summary, we for the first time, developed a dual carbonconfined strategy to coat the SnO2 hollow nanospheres using nitrogen-doped carbon and graphene, which could serve as the universal host material for alkali metal ions. This proposed unique structure provides rich voids and conductive carbon layers, which not only enables SnO2 with enough ability inside and out to relieve the large volume variation during the repeated cycling but also tremendously enhance the electron conductivity of SnO2. Benefitting from these structural advantages, G@C@SnO2 exhibit good Li+, Na+, and K+ storage properties including outstanding cycling and rate performance when it is served as a universal and high-performance anode material in both rechargeable alkali metal batteries (LIBs, SIBs,

(1) H


Research Article


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and KIBs). Especially, to the best of our knowledge, SnO2based anode KIB is reported for the first time. Overall, the present results demonstrate a general and convenient strategy to develop other metal oxide-based anodes for promising applications in rechargeable alkali metal batteries.

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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.8b00605. Additional results for Experimental Section, SEM images, TEM images, XRD patterns, FT-IR spectra, XPS spectra, CV curves, discharge/charge voltage profiles, GITT curves, and EIS spectra of various samples (PDF)

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. ORCID

Heng-guo Wang: 0000-0002-3704-0415 Author Contributions

The manuscript was written through contributions of all authors. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work was financially supported by the Science & Technology Department of Jilin Province (20170101177JC) and the Education Department of Jilin Province (no. 2016364).

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REFERENCES

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