Self-Assembly of Silicon@Oxidized Mesocarbon Microbeads

Jan 16, 2018 - After that, sucrose solution (0.5 g of sucrose (Jiang Tian Chemical Company) dissolved in 1 mL of water) was injected into the above sy...
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Cite This: ACS Appl. Mater. Interfaces 2018, 10, 4715−4725

Self-Assembly of Silicon@Oxidized Mesocarbon Microbeads Encapsulated in Carbon as Anode Material for Lithium-Ion Batteries Huitian Liu,† Zhongqiang Shan,† Wenlong Huang,† Dongdong Wang,† Zejing Lin,† Zongjie Cao,† Peng Chen,† Shuxian Meng,† and Li Chen*,‡ †

School of Chemical Engineering and Technology and ‡Department of Chemistry, School of Science, Tianjin University, Tianjin 300350, China S Supporting Information *

ABSTRACT: The utilization of silicon/carbon composites as anode materials to replace the commercial graphite is hampered by their tendency to huge volumetric expansion, costly raw materials, and complex synthesis processes in lithium-ion batteries. Herein, self-assembly method is successfully applied to prepare hierarchical silicon nanoparticles@oxidized mesocarbon microbeads/carbon (Si@O-MCMB/C) composites for the first time, in which O-MCMB core and low-cost sucrose-derived carbon shell not only effectively enhance the electrical conductivity of the anode, but also mediate the dramatic volume change of silicon during cycles. At the same time, the carbon can act as “adhesive”, which is crucial in enhancing the adhesive force between Si and O-MCMB in the composites. The as-obtained Si@O-MCMB/C delivers an initial reversible capacity of 560 mAh g−1 at 0.1 A g−1, an outstanding cyclic retention of 92.8% after 200 cycles, and respectable rate capability. Furthermore, the synthetic route presented here is efficient, less expensive, simple, and easy to scale up for high-performance composites. KEYWORDS: self-assembly, electrostatic attraction, silicon@graphite/carbon, cyclic stability, anode materials

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INTRODUCTION Lithium-ion batteries (LIBs), one of the highest-performing energy-storage systems, are widely used in portable electronics, emerging electric vehicles, and other industries.1−3 Silicon (Si) is deemed as one of the most promising alternatives to the commercial graphite for the next-generation anode materials in LIBs due to its high theoretical capacity of 4200 mAh g−1 (with the formation of Li4.4Si).4 However, its practical application is severely hindered by several factors: (1) significant volume change (>300%) inevitably induces the fracture and pulverization of electrode materials during charge/discharge process, resulting in the formation of new solid electrolyte interphase (SEI) layer on Si surface and hence leading to a rapid capacity decay and poor rate performance.5−7 (2) Si suffers from inherent low electrical conductivity, which leads to inferior rate capability and inadequate electroactive material utilization.8 Recombining Si and carbon is an appealing approach to overcome these drawbacks. Graphite has shown to be an excellent carbon framework due to its low cost, excellent electrical conductivity, and high Coulombic efficiency.9 Nevertheless, it has limited impact on mitigating the volume variation of Si.10,11 Besides, the binding strength between the Si and the graphite needs to be enhanced. Given these, the rational design should consist of a graphite core and Si embedded in an amorphous carbon matrix shell, and such design has multiple virtues: (1) the hierarchical structure could buffer the large volume change of Si and consequently maintain the structural © 2018 American Chemical Society

integrity during the repeated (de)lithiation process; (2) the three-dimensional conductive networks constructed by inner graphite and outer amorphous carbon layer could enhance the electrical conductivity of the whole electrode; and (3) the carbon layer anchored onto the Si/graphite surface could have effect on contributing to a stable SEI and the structural integrity. Many strategies have been adopted to synthesize Si/ graphite/carbon composites. By chemical vapor deposition of Si and carbon on the surface of graphite with the use of silane (SiH4) gas and high-purity acetylene, Si-nanolayer-embedded graphite/carbon hybrids have been successfully prepared.12 This structure delivers a capacity of ∼500 mAh g−1 after 100 cycles at a 0.5C rate with 96% capacity retention. However, the use of expensive equipment and toxic precursor, such as SiH4, is not suitable for mass production. In this case, simple ball mixing or ultrasonic stirring approach followed by carbonization has been applied to prepare the Si/graphite/carbon composites, which deliver enhanced electrochemical performance compared to pure Si.11,13 However, silicon nanoparticles (SiNPs) are difficult to be coated evenly on the graphite surface due to serious agglomeration. Besides, the obtained materials are heterogeneous, resulting from incompatibility between the Received: November 3, 2017 Accepted: January 16, 2018 Published: January 16, 2018 4715

DOI: 10.1021/acsami.7b16760 ACS Appl. Mater. Interfaces 2018, 10, 4715−4725

Research Article

ACS Applied Materials & Interfaces

Figure 1. Schematic of the fabrication process for Si@O-MCMB/C composites.

irregular particle sizes of SiNPs and graphite.12 Therefore, it is highly desirable to explore an advanced synthetic method to prepare Si/graphite/carbon composites with the homogeneous distribution of SiNPs on the graphite surface. Herein, electrostatic self-assembly is successfully employed to synthesize hierarchical silicon nanoparticles@oxidized mesocarbon microbeads/carbon (Si@O-MCMB/C) composites for the first time. Due to the attractive electrostatic forces between the positively charged SiNPs and the negatively charged OMCMB, Si-PDDA@O-MCMB aggregations containing uniformly distributed Si-PDDA on O-MCMB are achieved. Subsequently, they are coated with carbon, which served as a “binder” to further enhance the cohesion between SiNPs layer and O-MCMB. Different from the established Si/graphite/ carbon composites, our work has several significant advantages. First, the raw materials of MCMB and sucrose are commercially available and relatively less expensive. Second, the synthesis method is simple and facile, which make it suitable for scalable fabrication. Third, electrostatic assembly is an effective method to achieve uniform dispersion and enhance adhesion between two oppositely charged species compared to mechanical blend. The achieved results indicate that Si@O-MCMB/C composites exhibit enhanced cycling stability and excellent rate capability. Hence, this route has a great application potential in the mass production of high-performance materials for LIBs in the future.

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sucrose (Jiang Tian Chemical Company) dissolved in 1 mL of water) was injected into the above system. The solution was further evaporated at 65 °C under stirring, and the obtained solid composites were heated at 700 °C for 2 h under argon atmosphere at a heating rate of 3 °C min−1, followed by natural cooling to room temperature. For comparison, the Si/C without O-MCMB and Si/O-MCMB without amorphous carbon were prepared using the same process. Materials Characterizations. The structure and morphology of the product were observed using a scanning electron microscope (Hitachi S-4800, 10 kV). The microstructure and crystal characteristics of the samples were examined by a transmission electron microscope (JEOL JEM-2100F, 200 kV). X-ray diffraction (XRD) patterns were recorded using a X-ray diffractometer (Rigaku D/max 2550 P, Cu Kα). Raman spectroscopy was carried out with a wavelength of 632.8 nm. X-ray photoelectron spectroscopy (XPS) was conducted on Kratos Analytical Axis Ultra XPS device. The elemental contents of the composites were determined by a thermogravimetric analyzer (Q500) at a heating rate of 10 °C min−1 from 25 to 1000 °C in air atmosphere. Electrochemical Measurements. The electrochemical performance was evaluated using CR 2032 coin cells. The electrode slurry was prepared by dispersing active materials, super-P, and poly(vinylidene fluoride) (PVDF) binder with a mass ratio of 8:1:1 in N-methyl-2pyrrolidone solvent, then cast on a copper foil, and dried at 110 °C for 9 h in a vacuum oven. LiPF6 (1 M) dissolved in a mixture of ethylene carbonate/diethyl carbonate (volume ratio of 1:1) and fluoroethylene carbonate (FEC) (5 wt %) and Celgard 2400 microporous polypropylene membrane were used as the electrolyte and separator, respectively. In the half-cell, Li metal was used as the counter electrode and galvanostatic charge/discharge measurements were performed at a cutoff voltage range of 0.01−1.5 V with a constant current density of 0.1 or 0.3 A g−1 for cycling test and different current densities for rate test. In the full Li-ion cell, the Si@O-MCMB/C electrode was prelithiated by direct contact with lithium foils in the electrolyte for 1 h. The cathode electrode was fabricated with LiFePO4 (LFP, Taiwan Likai Power Technology Co., Ltd.), super-P containing 1% graphene, and PVDF in mass ratio of 8:1:1. The capacity ratio of the anode to cathode (N/P ratio) was 1.2:1. The full-cells were measured galvanostatically in the voltage window of 2−4.35 V at the current density of 1C (1C = 170 mA g−1). Cyclic voltammetry (CV) tests were conducted at a scanning rate of 0.1 mV s−1, using a Chen Hua CHI660E electrochemical workstation. The electrochemical impedance spectra (EIS) curves were obtained in a frequency range of 100 kHz to 0.1 Hz with an alternating current amplitude of 5 mV on IM6e electrochemical workstation. All electrochemical tests were performed at room temperature.

EXPERIMENTAL SECTION

Synthesis of Positively Charged SiNPs. SiNPs (0.5 g, 50−80 nm in diameter, Guangzhou Hongwu Materials Technology Co., Ltd.) were uniformly dispersed in deionized water (100 mL) by ultrasonic vibration; then, poly(acrylamide-co-diallyldimethyl ammonium chloride) (PDDA) (5.0 g, 10 wt %, Macklin) was poured into the above solution and the mixture was further sonicated. Later on, the PDDAfunctionalized SiNPs were washed three times with water to remove excess PDDA and collected by centrifugation at 10 000 rpm for 20 min. The obtained functionalized SiNPs (Si-PDDA) were dried under vacuum at 70 °C for 10 h. Preparation of Negatively Charged O-MCMB. MCMB (12−15 μm in diameter, Rongtan Technology Co., Ltd.) was dispersed in a mixed acid (sulfuric acid (98%)/nitric acid (65%) = 3:1 by volume), and the mixture was vigorously stirred at 70 °C for 10 h. Then, the obtained product (O-MCMB) was washed with distilled water for several times until the pH of the solution reached a constant value of about 7.0 and further dried at 70 °C under vacuum condition to gain functionalized O-MCMB. Preparation of Si@O-MCMB/C Composites. The Si-PDDA (0.1 g) was homogeneously dispersed into the mixture solvent of ethanol and water (45 and 5 mL, respectively) by sonicating for 1 h and then 0.5 g of O-MCMB was added into the above dispersion, followed by mechanical stirring for 1 h. After that, sucrose solution (0.5 g of

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RESULTS AND DISCUSSION Si@O-MCMB/C composites were synthesized via a simple self-assembly of SiNPs and O-MCMB, followed by carbonization of sucrose, as shown in Figure 1. It involves the following steps: (1) SiNPs were modified by PDDA to obtain positively charged Si-PDDA aggregates; (2) MCMB was oxidized by the mixed acid to increase the oxygen groups on 4716

DOI: 10.1021/acsami.7b16760 ACS Appl. Mater. Interfaces 2018, 10, 4715−4725

Research Article

ACS Applied Materials & Interfaces

Figure 2. TEM images of (a) pure Si (inset is its high-resolution TEM image) and (b) Si-PDDA. SEM images of (c) MCMB and (d) O-MCMB at different magnifications.

Figure 3. (a) XPS survey images of MCMB and O-MCMB; (b) high-resolution C 1s XPS images of MCMB and O-MCMB.

transform infrared spectroscopy (FTIR) results (Figure S1). The characteristic peaks at about 1097 and 491 cm−1 can be ascribed to the stretching vibrations of Si−O−Si and Si−OH, respectively. The study shows that SiOx is prone to be negatively charged; thus, it can electrostatically attract the positively charged PDDA.15,16 As a result, a polymer layer is clearly detected on SiNPs for Si-PDDA particles (Figure 2b). However, SiNPs cannot be completely dispersed as single nanoparticle and form a few agglomerated Si-PDDA clusters. By tuning sonication duration, the dispersion state of Si-PDDA can be improved, and Si-PDDA clusters can be significantly

the surface of MCMB; and (3) Si@O-MCMB/C composites were synthesized via electrostatic attraction between positively charged Si-PDDA and negatively charged O-MCMB, followed by uniformly coating pyrolytic carbon. As shown in Figure 2a, pristine SiNPs possess uniform spherical morphology with the size ranging from 50 to 80 nm, and the nanoparticles tend to agglomerate to some degree due to their high surface energy.14 Moreover, an obvious shell with a thickness of 5−6 nm can be observed on the surface of SiNP spheres in the local magnified high-resolution TEM (HRTEM) image, which is subsequently identified as SiOx via Fourier 4717

DOI: 10.1021/acsami.7b16760 ACS Appl. Mater. Interfaces 2018, 10, 4715−4725

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Figure 4. SEM images of (a, b) Si-PDDA@O-MCMB and (c, d) Si@O-MCMB/C at different magnifications. (e) TEM and (f) HRTEM images of Si@O-MCMB/C.

reduced as shown in the TEM image of Figure S2, when ultrasonication duration is prolonged to 5 h. The SEM images of MCMB and O-MCMB are shown in Figure 2c,d, respectively. Both MCMB and O-MCMB exhibit near-spherical structure with the diameter of 10−15 μm. After mixed acid oxidation treatment, the O-MCMB retains the spherical shape without structural collapse, whereas it shows a relatively rough surface with some cavities about 200 nm in diameter, which can facilitate the penetration of electrolyte.17 In addition, the surface chemical compositions of MCMB and OMCMB are investigated by X-ray photoelectron spectroscopy (XPS) measurements. From their XPS general spectra (Figure 3a) and the surface composition data (Table S1), it can be seen that the atomic percentage of O on the surface increases from 1.4% for MCMB to 9.08% for O-MCMB after acid treatment. C 1s high-resolution XPS images in Figure 3b can be resolved into four individual peaks: C−C at 284.5 eV, C−O at 285.4 eV, CO at 286.4 eV, and OC−O at 289.0 eV. OC−O in MCMB occupies the minimum component of 1.11 at %, whereas the content of OC−O in O-MCMB sharply increases to 2.52 at %. The increase in OC−O content can be ascribed to the oxidation of the partial hydroxyl and carboxide. Furthermore, the content of C−O raises from 9.63 at % for MCMB to 11.84 at % for O-MCMB. As a result, the

ionization of carbonyl and phenolic hydroxyl groups leads to a more negatively charged O-MCMB compared to MCMB.18 ζ-Potential measurements are conducted to shed light on the self-assembly process. The ζ-potentials of Si-PDDA and OMCMB are +27.3 and −17.0 mV, respectively, in the reaction solution. This suggests an opposite charge characteristic for SiPDDA and O-MCMB, which can lead to electrostatic attraction. The suspension stabilities of Si-PDDA, O-MCMB, Si-PDDA@O-MCMB, and the other contrast samples are further compared in Figure S3. After mechanical agitation and several hours standing, almost all samples settle down to the bottom of bottles and leave an upper clear solution. The SiPDDA@O-MCMB reveals a uniform black precipitate, whereas the other three composition samples (Si@MCMB, Si-PDDA@ MCMB, Si@O-MCMB) show obvious stratification phenomenon owing to the different density between SiNPs (Si-PDDA) and MCMB (O-MCMB). The comparison directly attests that Si-PDDA and O-MCMB can more efficiently combine to form aggregates by means of the electrostatic adsorption principle. The scanning electron microscopy (SEM) images are obtained to investigate the whole structure and morphology of nano-/microstructured Si-PDDA@O-MCMB composites (Figure 4a,b), wherein nano-Si-PDDA clusters are stacked on the micro-O-MCMB surface. In this case, a few of these clusters 4718

DOI: 10.1021/acsami.7b16760 ACS Appl. Mater. Interfaces 2018, 10, 4715−4725

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ACS Applied Materials & Interfaces

Figure 5. (a) XRD patterns of Si, O-MCMB, Si/O-MCMB, and Si@O-MCMB/C composites. (b) Raman spectra of Si, Si/O-MCMB, and Si@OMCMB/C. (c) N2 adsorption/desorption isotherms of O-MCMB, Si/O-MCMB, and Si@O-MCMB/C materials. (d) Pore size distribution curves of the three samples.

may peel off from the O-MCMB surface during the long-term cycle. In addition, these SiNP clusters are exposed directly to the electrolyte, which tend to some side effects, resulting in rapid capacity fading. To remedy the above-mentioned challenge, a carbon shell is encapsulated onto Si-PDDA@OMCMB composites by addition of sucrose as a carbon source. The obtained Si@O-MCMB/C sample (in Figure 4c) maintains a similar spherical shape to Si-PDDA@O-MCMB after carbon coating. The corresponding high-magnification SEM image (Figure 4d) reveals that a thin-film layer wraps SiNPs onto O-MCMB, which plays a vital role in sticking SiNPs and O-MCMB together. As a result, the structural stability of Si-PDDA@O-MCMB is improved due to the introduction of outer carbon shell. The details about the structure of Si@O-MCMB/C composites are further magnified in Figure 4e, in which irregular amorphous carbon is cladded on SiNPs (indicated by arrows) to form SiNP-embedded amorphous carbon layer with tight junction between neighboring SiNPs. In addition, the HRTEM image in Figure 4f (combining with Figure 2a) clearly shows a multilayer structure, in which the SiNPs with a SiOx layer are well encapsulated into the amorphous carbon matrix. As shown in Figure S4, mixing SiNPs and MCMB in sucrose solution (without using self-assembly method) results in inhomogeneous SiNPs on the MCMB. The X-ray diffraction (XRD) patterns of Si@O-MCMB/C, SiNPs, O-MCMB, and Si/O-MCMB are shown in Figure 5a. In the case of Si@O-MCMB/C, the distinct diffraction peaks at 26.4, 42.2, 54.5, and 77.3° are assigned to (002), (100), (004), and (110) lattice planes of graphite (JCPDS no. 41-1487), respectively, whereas the weak diffraction peaks at 28.4, 47.3, and 56.1° correspond to (111), (220), and (311) planes of Si (JCPDS no. 27-1402). The coexistence of graphite and Si XRD peaks indicates the co-presence of O-MCMB and SiNPs in the composites, which do not change their respective crystalline

structure after sonication and pyrolysis. No distinctive diffraction peaks of carbon are detected, meaning that the sucrose-derived carbon is amorphous in nature.19 Raman spectroscopy (Figure 5b) is used to further investigate the sample structure. The peaks at 290, 513, and 929 cm −1 in the Raman spectra of Si@O-MCMB/C correspond to Si. In detail, the weak peak at 290 cm−1 is characteristic of the second acoustic phonon mode of Si. The strong peak located at 513 cm−1, which can be used to confirm the crystal structure of Si, shows an obvious blue shift toward higher frequency in comparison to that in pure Si (around 495 cm−1), which can be attributed to the pyrolytic carbon coating.14 Meanwhile, the peak at 929 cm−1 indicates the immobilization of SiO2 on SiNP surface.20 In addition, the peaks at 1346 cm−1 (D-band) and 1590 cm−1 (G-band) are ascribed to the vibration of defective carbon atoms and the stretching vibration of sp2-hybridized carbon atom, respectively.21,22 Generally, the area ratio R of D-band to G-band reflects the degree of graphitization, and its lower value indicates the higher graphitization degree of the composites.23 Apparently, the R value of Si@O-MCMB/C composites (1.73) is significantly larger than that of Si/O-MCMB (0.99), which suggests a lower degree of graphitization for Si@O-MCMB/C due to the incorporation of pyrolyzed carbon. In this case, a uniform and compact SEI film can be formed on the carbon surface of Si@O-MCMB/C composites, thereby effectively preventing capacity fading, resulting from continuous electrolyte degradation during cycles.24 The contents of Si, O-MCMB/C, and Si@O-MCMB/C composites are determined by thermogravimetric analysis (TGA) curves in Figure S5a. Two distinct weight change stages can be observed in the TGA curves: the weight loss stage between 500 and 800 °C is ascribed to the complete combustion of O-MCMB and pyrolytic carbon, whereas the weight gain stage is attributed to the oxidation of Si to form 4719

DOI: 10.1021/acsami.7b16760 ACS Appl. Mater. Interfaces 2018, 10, 4715−4725

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ACS Applied Materials & Interfaces

Figure 6. (a) Cyclic voltammetry profiles of Si@O-MCMB/C. (b) Voltage profiles of the Si@O-MCMB/C composites for the 1st, 2nd, 50th, 100th, and 200th cycle at a current density of 0.1 A g−1. (c) Cycling performance of Si@O-MCMB/C, Si, O-MCMB, and Si/O-MCMB at 0.1 A g−1. (d) Rate performance of the four samples. (e) Long-term cycling performance of the Si@O-MCMB/C electrode at 0.3 A g−1.

adsorption−desorption isothermal measurements. A typical type-IV isotherm is observed in Figure 5c for all of the samples, indicating the presence of mesopores.25 Si/O-MCMB and OMCMB exhibit similar N2 hysteresis loops at the relative pressure P/P0 > 0.45, which is mainly assigned to the lamellaformed slitlike pores around 2.07 nm in O-MCMB (Figure 5d).26 This illustrates that deposited SiNPs do not damage the original surface state of O-MCMB, whereas the surface area increases from 16.86 m2 g−1 for O-MCMB to 21.04 m2 g−1 for Si/O-MCMB due to the introduction of SiNPs with high specific surface area. However, the final product, Si@OMCMB/C, shows a significantly different N2 hysteresis hoop, which is characteristic of uniform channel-like pores.26 The different pore sizes for the three samples may be attributed to the insertion of sucrose molecules into the pores of O-MCMB during the coating process.27 After the carbon-coating process, the original pores (5.3−50 nm) of O-MCMB reduce, but abundant well-ordered small pores (centered at 1.7 nm) in the pyrolytic carbon form due to the pyrolysis of sucrose. The porous structure effectively accommodates the volume

SiOx. The calculated average weight percents of Si and C for Si@O-MCMB/C are 13.76 and 86.24%, respectively (Table S2). Besides, the TGA curves of both Si@O-MCMB/C and OMCMB/C are similar to each other but distinctly different from that of Si/O-MCMB (without amorphous carbon) between 500 and 800 °C, which shows only one weight loss process (combustion of O-MCMB). In contrast to the single combustion of O-MCMB for Si/O-MCMB, the first rapid weight loss for Si@O-MCMB/C and O-MCMB/C is attributed to the decomposition of amorphous carbon at the lower temperature between 500 and 620 °C, and the second is ascribed to the combustion of O-MBMC occurring between 620 and 800 °C. On the basis of these different reactions in weight loss stage, the content of amorphous carbon (mainly relates to the carbonization of sucrose) is further determined according to the two derivative peaks of the differential thermogravimetry profiles in Figure S5b. The calculated contents of amorphous carbon and O-MCMB are about 27.8 and 58.4% in Si@O-MCMB/C, respectively. The porous structures of Si@O-MCMB/C, Si/O-MCMB, and O-MCMB are further characterized with nitrogen 4720

DOI: 10.1021/acsami.7b16760 ACS Appl. Mater. Interfaces 2018, 10, 4715−4725

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ACS Applied Materials & Interfaces

implying high structure stability and hence leading to excellent electrochemical performance.44 Figure S7 compares the initial discharge/charge curves of MCMB and O-MCMB. The initial discharge and charge capacities are 781.25 and 354.21 mAh g−1, respectively, for O-MCMB, which are much higher than those for MCMB (394.5 and 212.83 mAh g−1). The high specific capacity of O-MCMB is probably related to the more cavities in O-MCMB, which provide more active sites for Li+ to intercalation/deintercalation and facilitate the diffusion of Li+.45 The cycle performances of Si@O-MCMB/C, SiNPs, OMCMB, and Si/O-MCMB at a current density of 0.1 A g−1 between 0.01 and 1.5 V vs Li+/Li are evaluated (Figure 6c). As expected, the pure SiNP electrode displays the highest initial charge capacity of 2016.4 mAh g−1, but decays rapidly to less than 20 mAh g−1 after 20 cycles, which can be ascribed to the poor connection between active materials and copper foil induced by the huge volume expansion of SiNPs.46 By comparison, Si/O-MCMB shows a slower capacity fading than SiNPs, and its capacity increases slightly within 50 cycles due to the repeated expansion and contraction of Si layer, which allows gradual infiltration of electrolyte into the Si/OMCMB electrode.47 However, the average reversible capacity of Si/O-MCMB is still as low as 429.26 mAh g−1 within 200 cycles due to the lower Si utilization. After carbon coating, the charging capacity of Si@O-MCMB/C rapidly increases to 552.2 and 556.4 mAh g−1 in the 1st and 100th cycles and only decreases slightly to 510.8 mAh g−1 after 200 cycles, showing a high capacity retention of 92.8%. The cyclability of Si@OMCMB/C is significantly improved by carbon coating, which not only offers more effective paths for the electron transmission but also stabilizes the secondary structure of Si@O-MCMB/C by alleviating the volume dilation of Si and strengthening the connection between SiNPs and O-MCMB during cycles. To illustrate the merit of Si@O-MCMB/C hierarchical configuration, the cycling performance and corresponding Coulombic efficiencies of Si@O-MCMB/C and Si/C are compared in Figure S8. Although the Si@C sample delivers a higher initial charge capacity of 578 mAh g−1 compared to Si@O-MCMB/C, its capacity drops rapidly to 148 mAh g−1 after 200 cycles, demonstrating a poor capacity retention of 25.6%. Therefore, both O-MCMB and amorphous carbon are of great importance in achieving superior cycle performance in Si@O-MCMB/C composites. The rate capabilities of SiNPs, O-MCMB, Si/O-MCMB, and Si@O-MCMB/C are investigated by testing at current densities between 0.1 and 1.6 A g−1, followed by stepwise decreasing in current densities to 0.1 A g−1. As shown in Figure 6d, the capacity of SiNPs declines sharply at 0.1 A g−1 and then decreases to 0 at 0.2 A g−1. After incorporating SiNPs onto OMCMB, the Si@O-MCMB shows significantly enhanced rate capability compared to SiNPs. When coated with an additional carbon layer, Si@O-MCMB/C delivers the highest capacity at various current densities. A capacity of 540 mAh g−1 is attained at a current density of 0.1 A g−1, and a reversible capacity of above 230 mAh g−1 is achieved at a higher current density of 1.6 A g−1. For all samples, the capacities decrease with increasing current density due to the limited lithium-ion migration at high rates. However, Si@O-MCMB/C exhibits higher lithium-ion migration rate than Si and Si/O-MCMB. In addition, when the current density backs to 0.1 from 1.6 A g−1, the specific capacity of Si@O-MCMB/C recovers 99.8% of the original capacity. Thus, the excellent rate performance may be due to the combined effect of carbon-coating layer and O-

expansion of SiNPs and accelerates the electrolyte infiltration.28,29 Half-coin cells are utilized to examine the electrochemical performance of Si@O-MCMB/C composites. The cyclic voltammetry (CV) curves of Si@O-MCMB/C for the initial five cycles between 0.01 and 1.5 V (vs Li+/Li) at a scan rate of 0.1 mV s−1 are presented in Figure 6a, and those of pure SiNPs and O-MCMB under the same test condition are shown in Figure S6a,b, respectively. For Si@O-MCMB/C, the first cathodic scan (lithiation process) reveals a weak peak at 0.95 V corresponding to the decomposition of electrolytic additive of FEC and a broad peak between 0.85 and 0.5 V derived from the formation of SEI layer.30,31 In addition, the sharp cathodic peak below 0.15 V can be attributed to the insertion reactions of Li into both crystal Si and O-MCMB.32 After that, two new peaks have emerged as the cycling process continued. The former located at 0.1 V is correlated to phase transition from amorphous Si to amorphous LixSi, whereas the latter at 0.17 V is assigned to the graphite intercalation compound (LixC), and it changes with the lithium-ion insertion degree.33 The anodic scan curves exhibit a sharp peak at 0.24 V and a gradually obvious hump at around 0.5 V with cycling, which correspond to Li extraction out of O-MCMB and the dealloying of LixSi to amorphous Si, respectively. All of the above-mentioned peaks can also be found in the CV curves for either bare Si or O-MCMB, revealing that both Si and OMCMB are electrochemically active materials for lithium storage and not affected by each other.21,34 The cathodic current in the first cycle is larger than those of the second and third cycles due largely to the formation of SEI layer in the first discharge process.35 Nevertheless, both the cathodic and anodic peak intensities gradually increase from the second cycle to fifth cycle, indicating an activation process.36,37 Notably, after the first cycle, the CV curves of Si@O-MCMB/C generally overlap reasonably with subsequent cycles. This suggests a higher reversibility in lithium-ion intercalation/deintercalation for Si@ O-MCMB/C compared to Si electrode.15 The charge−discharge profile curves of Si@O-MCMB/C at a current density of 0.1 A g−1 and a potential range of 0.01−1.5 V (vs Li+/Li) are shown in Figure 6b. In the initial discharge curve, a slope ranging from 1.2 to 0.2 V corresponds to the formation of SEI layer, and a distinct plateau below 0.2 V indicates the insertion of lithium ions in both crystalline SiNPs and O-MCMB.38 The lithiation plateau region (below 0.1 V) of the first cycle shows differences from that of the subsequent cycles (below 0.4 V), mainly due to the transformation of Si from crystalline to amorphous phase during the first cycle.39 The result coincides with the above CV result. The initial discharge and charge capacities of the Si@O-MCMB/C sample are 957.29 and 552.19 mAh g−1, respectively, giving rise to a Coulombic efficiency of 57.68%. The low Coulombic efficiency can be explained by the formation of SEI on the electrode surface, the secondary reaction of Li+ with the oxygen functional groups of O-MCMB, and the high density of lithium-trapping sites in amorphous carbon.40−43 Moreover, the cycling performance curve of Si@O-MCMB/C in Figure 6c also indicates that its Coulombic efficiency increases with cycling, achieving above 99% for 20th cycle and 99.5% for 100th cycle. This can be interpreted by the gradually enhanced stability of SEI layer on Si@O-MCMB/C surface with cycling.22 Most significantly, the charging voltage plateau at 100th cycle is still visible and the voltage profiles almost overlap upon further cycling except the first discharge−charge cycle, 4721

DOI: 10.1021/acsami.7b16760 ACS Appl. Mater. Interfaces 2018, 10, 4715−4725

Research Article

ACS Applied Materials & Interfaces

charges transfer across the interface between electrolyte and electrode), whereas the line represents the Warburg impedance (Zw, lithium ions diffuse through active materials).27,48 The fitted equivalent circuit is given in Figure S9; then, RSEI and Rct are estimated using simulation software ZView. In comparison to Si/O-MCMB (10.5 and 20.2 Ω) and Si (50.1 and 42.3 Ω), Si@O-MCMB/C possesses the smallest RSEI and Rct of 5.2 and 15.3 Ω, respectively, indicating a thinner and more stable SEI layer and faster electron transfer. Nyquist plots of Si@OMCMB/C before and after different numbers of cycles are also shown in Figure S10. The fresh cell exhibits only one semicircle and a line, corresponding to Rct and Zw, due to without SEI films. The RSEI values of the electrode after 5, 30, and 50 cycles are 4.9, 8.6, and 15.2 Ω, respectively. The small change of RSEI during the first 50 cycles indicates the formation of a stable SEI film on the Si@O-MCMB/C and therefore good cycle performance of the composite electrode. To further investigate the reason behind the excellent performance of Si@O-MCMB/C, the electrode thickness and surface morphology are compared before and after 200 cycles in lithiation state. As shown in Figure 8a,b, the electrode thickness increases from about 28 μm for the pristine electrode to 33.77 μm after 200 cycles, giving rise to a small volume expansion of 21%. In addition, its high-magnification images reveal no cracks nor morphology changes on the surface of the cycled Si@O-MCMB/C, which suggests good interparticle connectivity during the lithiation/delithiation process. To address the issue regarding volume expansion during lithiation/delithiation process, the structural integrity of the anode materials has to be considered. By utilizing a simple selfassembly approach and subsequent carbon coating, a structurally stable Si@O-MCMB/C can be achieved as a result of strong adhesive force between positively charged SiNPs and negatively charged O-MCMB, which demonstrates outstanding performance. To evaluate the possibility for practical application in LIBs, a full cell is fabricated with Si@O-MCMB/C anode and a commercially available LiFePO4 (LFP) cathode. The half-cell performance of LFP is shown in Figure S11, and it delivers initial charge and discharge capacities of 181 and 168 mAh g−1, respectively, with stable cyclability. As analyzed above, the anode exhibits an initial irreversible loss of 42.32%, which will lead to a lessened total energy density. Herein, to eliminate the first irreversible capacity loss, prelithiation is achieved by direct contact with metal Li in electrolyte solution. The full cell is designed to cathode limited with anode to cathode capacity ratio of 1.2:1 and then tested at 0.2C for initial five cycles and 1C for the following cycles. As shown in Figure 9, the Si@OMCMB/C-LFP full cell demonstrates a discharge capacity of 164 mAh g−1 at the first cycle, associated with an initial

MCMB, which provide mechanical backbone with excellent electronic network. As a result, most SiNPs become electrochemically active by means of the good electrical connection with O-MCMB and hence leading to high capacity due to the high utilization of SiNPs. The long-term cycling stability of the Si@O-MCMB/C is further evaluated at a relatively high current density of 0.3 A g−1 (Figure 6e). A reversible capacity of 458 mAh g−1 is obtained after 500 cycles, which translates to a good capacity retention of 86.5% to the capacity of the sixth cycle. The superior electrochemical performance of Si@O-MCMB/C may be mainly attributed to the robust structure due to the strong adhesive forces between the SiNPs and O-MCMB. The electrochemical impedance spectroscopy (EIS) measurements of SiNPs, O-MCMB, Si/O-MCMB, and Si@O-MCMB/ C are performed in lithiation state after five cycles. As shown in Figure 7, for all of these samples, the Nyquist plots consist of a

Figure 7. Electrochemical impedance plots of Si@O-MCMB/C, Si, OMCMB, and Si/O-MCMB.

depressed semicircle, where a high-frequency semicircle and a medium-frequency semicircle overlap each other, and an inclined line in the low frequency. The high-frequency semicircle can be attributed to SEI resistance (RSEI, Li ions transport in the SEI layer), and the medium-frequency semicircle is correlated to the charge-transfer resistance (Rct,

Figure 8. SEM images of electrode cross section of Si@O-MCMB/C (a) before cycling and (b) after 200 cycles at a current density of 0.1 A g−1. Top views of these electrodes are shown on the right-hand side. 4722

DOI: 10.1021/acsami.7b16760 ACS Appl. Mater. Interfaces 2018, 10, 4715−4725

Research Article

ACS Applied Materials & Interfaces

Figure 9. (a) Charge−discharge profiles for Si@O-MCMB/C full cell at 0.2C between 2 and 4.35 V, and (b) cycling performance of battery tested at 0.2C for the first five cycles and 1C for the later cycles.

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Coulombic efficiency of 90.08%, which is considerably higher than that for the half-cell due to the prelithiation process. In addition, the reversible capacity maintains at 135 mAh g−1 after 100 cycles, with a capacity retention ratio of 96.1% to the capacity of the sixth cycle, indicating that the Si@O-MCMB/C composites can be regarded as a promising anode material for LIBs.

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*E-mail: [email protected]. Tel: (+86)22 27403475. Fax: (+86)22 27403475. ORCID

Li Chen: 0000-0001-9617-5224 Notes

The authors declare no competing financial interest.

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CONCLUSIONS

ACKNOWLEDGMENTS This work was supported by the National Key Research and Development Program of China (2016YFB0100511) and the State Scholarship Fund of China Scholarship Council.

A simple self-assembly method through electrostatic attraction between positively charged SiNPs and negatively charged OMCMB, followed by carbonization of sucrose is successfully employed to fabricate Si@O-MCMB/C composites, in which SiNP layer is deposited onto the O-MCMB surface and then completely embedded in amorphous carbon in the subsequent coating process. Herein, due to the synergistic role of carboncoating layer and O-MCMB, the structural stability and electrical conductivity of the Si@O-MCMB/C composites are enhanced. As a result, the as-obtained Si@O-MCMB/C delivers a high initial reversible capacity (560 mAh g−1), outstanding cyclic stability (92.8% capacity retention at 0.1 A g−1 after 200 cycles, and 86.5% capacity retention even at high current rate of 0.3 A g−1 after 500 cycles), high Coulombic efficiency, and good rate capability. Furthermore, this selfassembling synthetic protocol is efficient to achieve uniform dispersion and enhance adhesion between two oppositely charged species compared to simple blend. Meanwhile, it is more suitable for industrial production compared to conventional approaches that require toxic precursors or solvent, expensive equipment, and harsh condition. Hence, by adopting reasonable hierarchical electrode structure and cost-efficient fabrication method, Si-based materials possess great potential for the next-generation anode materials in lithium-ion batteries.

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AUTHOR INFORMATION

Corresponding Author

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REFERENCES

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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.7b16760. FTIR spectra of SiNPs, TEM images of Si-PADDA, photographs, SEM images of Si@MCMB/C, TGA curves of Si@O-MCMB/C, cyclic voltammetry profiles, initial discharge−charge voltage profiles, cycling performance, and the corresponding Coulombic efficiencies (PDF) 4723

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