Self-assembly of Silicon@Oxidized Mesocarbon Microbeads

Subscriber access provided by READING UNIV

Article

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 ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.7b16760 • Publication Date (Web): 16 Jan 2018 Downloaded from http://pubs.acs.org on January 16, 2018

Just Accepted “Just Accepted” manuscripts have been peer-reviewed and accepted for publication. They are posted online prior to technical editing, formatting for publication and author proofing. The American Chemical Society provides “Just Accepted” as a free service to the research community to expedite the dissemination of scientific material as soon as possible after acceptance. “Just Accepted” manuscripts appear in full in PDF format accompanied by an HTML abstract. “Just Accepted” manuscripts have been fully peer reviewed, but should not be considered the official version of record. They are accessible to all readers and citable by the Digital Object Identifier (DOI®). “Just Accepted” is an optional service offered to authors. Therefore, the “Just Accepted” Web site may not include all articles that will be published in the journal. After a manuscript is technically edited and formatted, it will be removed from the “Just Accepted” Web site and published as an ASAP article. Note that technical editing may introduce minor changes to the manuscript text and/or graphics which could affect content, and all legal disclaimers and ethical guidelines that apply to the journal pertain. ACS cannot be held responsible for errors or consequences arising from the use of information contained in these “Just Accepted” manuscripts.

ACS Applied Materials & Interfaces is published by the American Chemical Society. 1155 Sixteenth Street N.W., Washington, DC 20036 Published by American Chemical Society. Copyright © American Chemical Society. However, no copyright claim is made to original U.S. Government works, or works produced by employees of any Commonwealth realm Crown government in the course of their duties.

Page 1 of 23 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

ACS Applied Materials & Interfaces

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

a

and Li Chen*b School of Chemical Engineering and Technology, Tianjin University, Tianjin 300350, China

b

Department of Chemistry, School of Science, Tianjin University, Tianjin 300350, China

*E-mail: [email protected]

Keywords: self-assembly, electrostatic attraction, silicon@graphite/carbon, cyclic stability, anode materials

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, low-cost, simple, and easy to scale up for high performance composites.

ACS Paragon Plus Environment

ACS Applied Materials & Interfaces 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 2 of 23

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 the 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 the 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 integrity during the repeated (de)lithiation process; (2) the 3D conductive networks constructed by inner graphite and outer amorphous carbon layer could enhance the electrical conductivity of the whole electrode; (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 vapour deposition 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 0.5 C 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 with 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 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


Page 3 of 23 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


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 negatively charged O-MCMB, 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. Firstly, the raw materials of MCMB and sucrose are commercially available and relatively low cost. Secondly, the synthesis method is simple and facile which make it suitable for scalable fabrication. Thirdly, electrostatic assembly is an effective method to achieve uniform dispersion and enhance adhesion between two oppositely charged species as 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 great application potential in mass production of high-performance materials for LIBs in the future.

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 PDDA functionalized SiNPs were washed three times with water to remove excess PDDA, and collected by centrifugation at 10000 r for 20 min. The obtained functionalized SiNPs (Si-PDDA) were dried under vacuum at 70 °C for 10 h.



Page 4 of 23

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 (0.1g) was homogeneously dispersed into the mixture solvent of ethanol and water (45 mL and 5 mL, respectively) by sonicating for 1 h, and then 0.5 g O-MCMB was added into the above dispersion followed by mechanical stirring for 1 h. After that, sucrose solution (0.5 g of 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 with a heating rate of 3 °C min-1, followed by naturally cooling to room temperature. For comparison, the Si/C without O-MCMB and Si/O-MCMB without amorphous carbon were, respectively, prepared using the same process.

Materials characterizations

The structure and morphology of the product were observed using Scanning Electron Microscope (SEM, Hitachi S-4800, 10 kV). The microstructure and crystal characteristic of the samples were examined by Transmission Electron Microscope (TEM, JEOL JEM-2100F, 200 kV). X-ray diffraction (XRD) patterns were recorded with 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. The elemental contents of the composites were determined by the thermal gravimetric analyzer (TGA, Q500) with a heating rate of 10 °C min -1 from 25 °C to 1000 °C under air atmosphere.


Page 5 of 23 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


Electrochemical measurements

The electrochemical performance was evaluated using CR 2032 coin cells. The electrode slurry was prepared by dispersing active materials, superP and poly (vinylidene fluoride) (PVDF) binder with a mass ratio of 8:1:1 in N-methyl-2-pyrrolidone (NMP) solvent, then cast on a copper foil and dried at 110 °C for 9 h in a vacuum oven. 1 M LiPF6 dissolved in a mixture of ethylene carbonate (EC)/diethyl carbonate (DEC) (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 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 AC amplitude of 5 mV on IM6e electrochemical workstation. All electrochemical tests were performed at room temperature.

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 positive-charged Si-PDDA aggregates, (2) MCMB was oxidized by the mixed acid to increase the oxygen groups on the surface of MCMB, (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.



Page 6 of 23

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

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 spectra of MCMB and O-MCMB, (b) High-resolution C1s of MCMB and O-MCMB.


Page 7 of 23 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


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 SiNPs spheres in the local magnified HRTEM image, which is subsequently identified as SiOx via Fourier Transform Infrared Spectroscopy (FTIR) results (Figure S1). The characteristic peaks at about 1097 and 491 cm-1 can be ascribed to the stretching vibration of Si-OSi 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 reduced as shown in 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 and 2d. 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, while shows 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 O-MCMB are investigated with X-ray photoelectron spectroscopy (XPS) measurements. From their XPS general spectrums (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 spectra 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, respectively. O=C-O in MCMB occupies the minimum component of 1.11 at%, while 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 as compared to MCMB.18 Zeta potential measurements are conducted to shed light on the self-assembly process. The zeta potentials of Si-PDDA and O-MCMB are +27.3 and -17.0 mV, respectively, in the reaction solution. This suggests an opposite charge characteristic for Si-PDDA and O-MCMB which can lead to electrostatic attraction. The suspension stability 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



Page 8 of 23

an upper clear solution. The Si-PDDA@O-MCMB reveals a uniform black precipitate, while 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 mean of the electrostatic adsorption principle.

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

The scanning electron microscopy (SEM) images are obtained to investigate the whole structure and morphology of nano/micro-structured Si-PDDA@O-MCMB composites (Figure 4a and 4b), wherein nano-Si-PDDA clusters are stacked on the micro-O-MCMB surface. In this case, a few of these clusters may peel off from the O-MCMB surface during the long-term cycle. In addition, these SiNPs clusters are exposed directly to the electrolyte, which tend to some side effects, resulting in rapid capacity fading. In order to remedy the above mentioned challenge, a carbon shell is encapsulated onto Si-PDDA@O-MCMB composites by addition of sucrose as a carbon source. The obtained Si@O-MCMB/C sample (in Figure 4c) maintains 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)


Page 9 of 23 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


to form SiNPs-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.

Figure 5. (a) XRD patterns of Si, O-MCMB, Si/O-MCMB and Si@O-MCMB/C composites, (b) Raman spectrum for 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 and (d) pore size distribution curves of the three samples.

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@OMCMB/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, while the weak diffraction peaks at 28.4°, 47.3° and 56.1° are corresponded to (111), (220) and (311) planes of Si (JCPDS No. 27-1402). The co-existence 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



Page 10 of 23

of Si@O-MCMB/C correspond to Si. In details, the weak peak at 290 cm-1 is characteristic of 2nd 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 with 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 SiNPs surface.20 In addition, the peaks at 1346 (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 and Gband 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 thermal gravimetric analyzer (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, while the weight gain stage is attributed to the oxidation of Si to form SiOx. The calculated average weight percent of Si and C for Si@O-MCMB/C is 13.76% and 86.24%, respectively (Table S2). Besides, the DG curves of both Si@OMCMB/C and O-MCMB/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/OMCMB, 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. Based on 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 (DTG) 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 adsorption-desorption isothermal measurements. A typical type-IV isotherm is observed in Figure 5c for all the samples, indicating the presence of mesopores.25 Si/OMCMB and O-MCMB exhibit similar N2 hysteresis loop at the relative pressure P/P 0 > 0.45, which is mainly assigned to the lamella formed slit-


Page 11 of 23 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


like 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, while 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@O-MCMB/C, shows 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 (from 5.3 to 50 nm) of O-MCMB reduces, 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 expansion of SiNPs and accelerates the electrolyte infiltration.28,29

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 and (e) Long-term cycling performance of the Si@O-MCMB/C electrode at 0.3 A g-1.



Page 12 of 23

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 0.01-1.5 V (vs Li+/Li) at a scan rate of 0.1 mV s-1 are presented in Figure 6a, and these of pure SiNPs and O-MCMB under the same test condition are shown in Figure S6a and S6b, 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 with phase transition from amorphous Si to amorphous Li xSi, while 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 de-alloying of LixSi to amorphous Si, respectively. All 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 these 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 as 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 1st cycle shows differences from those 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 Si@OMCMB/C sample are 957.29 and 552.19 mAh g-1, giving rise to a Coulombic efficiency of 57.68%. The low Coulombic efficiency can be explained by 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


Page 13 of 23 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


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, 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 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 performance of Si@O-MCMB/C, SiNPs, O-MCMB and Si/O-MCMB at a current density of 0.1 A g-1 between 0.01 and 1.5 V vs. Li+/Li is evaluated (Figure 6c). As expected, the pure SiNPs 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 Si/O-MCMB 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 cycle, 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@O-MCMB/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. In order 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 Si@C sample delivers a higher initial charge capacity of 578 mAh g-1 as 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 A g-1 to 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 zero at 0.2 A g-1. After incorporating SiNPs onto O-MCMB, the Si@O-MCMB shows significantly enhanced rate capability than SiNPs. When coated with an additional carbon layer, Si@O-MCMB/C delivers the highest capacity at various current densities. A



Page 14 of 23

capacity of 540 mAh g-1 is attained at 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 faster 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 due to the combined effect of carbon coating layer and O-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 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 6th cycle. The superior electrochemical performance of the Si@O-MCMB/C may be mainly attributed to the robust structure due to the strong adhesive forces between the SiNPs and O-MCMB.

Figure 7. The electrochemical impedance plots of Si@O-MCMB/C, Si, O-MCMB and Si/O-MCMB.

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 at the right-hand side.


Page 15 of 23 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


Figure 9. (a) Charge−discharge profiles for Si@O-MCMB/C full cell at 0.2 C between 2 and 4.35 V, and (b) cycling performance of battery tested at 0.2 C for the first 5 cycles and 1 C for the later cycles.

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 these samples, the Nyquist plots consist of a depressed semicircle where a highfrequency semicircle and a medium frequency semicircle overlap each other, and an inclined line in the low frequency. The semicircle in the high

frequency can be attributed to SEI resistance (RSEI, Li ions transport in the SEI layer), and the medium frequency semicircle is correlated with the charge transfer resistance (Rct，charges transfer across the interface between electrolyte and electrode), while 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 of Zview. In comparison with 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, indicating a thinner and more stable SEI layer and faster electron transfer. Nyquist plots of the 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 of the electrode after 5, 30 and 50 cycles is 4.9, 8.6 and 15.2 Ω, respectively. The small change of RSEI during the first 50 cycles indicates formation of a stable SEI film on the Si@O-MCMB/C and therefore good cycle performance of the composite electrode. In order 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 and 8b, 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 inter-particle



Page 16 of 23

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 self-assembly 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 an initial charge and discharge capacity of 181 and 168 mAg−1, respectively, with stable cyclability. As analysis 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/cathode capacity ratio of 1.2:1, and then tested at 0.2 C for initial five cycles and 1 C for the following cycles. As shown in Figure 9, the Si@O-MCMB/C-LFP full cell demonstrates a discharge capacity of 164 mAh g−1 at the first cycle, associated with an initial 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 6th cycle, indicating the Si@O-MCMB/C composites can be regarded as a promising anode material for LIBs.

CONCLUSION

A simple self-assembly method through electrostatic attraction between positively charged SiNPs and negatively charged O-MCMB followed by carbonization of sucrose is successfully employed to fabricate Si@O-MCMB/C composites, in which SiNPs layer is deposited onto O-MCMB surface, and then completely embedded in amorphous carbon in the subsequent coating process. Herein, due to the synergistic role of carbon coating 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 (560mAh 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 self-assembling synthetic protocol is efficient to achieve uniform dispersion and enhance adhesion between two oppositely charged species as 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


Page 17 of 23 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


structure and cost efficient fabrication method, Si-based materials possess great potential for the next generation anode materials in lithium-ion batteries.

ASSOCIATED CONTENT Supporting Information FTIR spectra of SiNPs, TEM images of Si-PADDA, The photographs, SEM images of Si@MCMB/C, TGA curves of Si@O-MCMB/C, Cyclic voltammetry profiles, The initial discharge−charge voltage profiles, Cycling performance and corresponding Coulombic efficiencies (PDF)

AUTHOR INFORMATION Corresponding Author *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.

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

REFERENCES

(1) S.-J. Kim; M.-C. Kim; S.-B. Han; G.-H. Lee; H.-S. Choe; D.-H. Kwak; S.-Y. Choi; B.-G. Son; M.-S. Shin and K.-W. Park. 3D flexible Si based-composite (Si@Si 3 N 4 )/CNF electrode with enhanced cyclability and high rate capability for lithium-ion batteries. Nano Energy 2016, 27, 545-553.



Page 18 of 23

(2) X. Zhou; Y. X. Yin; A. M. Cao; L. J. Wan and Y. G. Guo. Efficient 3D conducting networks built by graphene sheets and carbon nanoparticles for high-performance silicon anode. Adv. Energy Mater. 2012, 4, 2824-2828. (3) L. Zhou; Z. Zhuang; H. Zhao; M. Lin; D. Zhao and L.Q. Mai. Intricate Hollow Structures: Controlled Synthesis and Applications in Energy Storage and Conversion. Adv. Mater. 2017, 29, 1602914-1602942. (4) X. Zhou; Y.-X. Yin; L.-J. Wan; Y.-G. Guo. Self-Assembled Nanocomposite of Silicon Nanoparticles Encapsulated in Graphene through Electrostatic Attraction for Lithium-Ion Batteries. Adv. Energy Mater. 2012, 2, 1086-1090. (5) W. Ren; Y. Wang; Q. Tan; Z. Zhong; F. Su. Novel silicon/carbon nano-branches synthesized by reacting silicon with methyl chloride: A high performing anode material in lithium ion battery. J. Power Sources 2016, 332, 88-95. (6) Q. Xiao; M. Gu; H. Yang; B. Li; C. Zhang; Y. Liu; F. Liu; F. Dai; L. Yang; Z. Liu; X. Xiao; G. Liu; P. Zhao; S. Zhang; C. Wang; Y. Lu; M. Cai. Inward lithium-ion breathing of hierarchically porous silicon anodes. Nat. Commun. 2015, 6, 8844-8852. (7) R. Xu; G. Wang; T. Zhou; Q. Zhang; H.-P. Cong; X. Sen; J. Rao; C. Zhang; Y. Liu; Z. Guo and S.-H. Yu. Rational design of Si@carbon with robust hierarchically porous custard-apple-like structure to boost lithium storage. Nano Energy 2017, 39, 253-261. (8) Fei-Hu Du; K.-X. Wang; J.-S. Chen. Strategies to succeed in improving the lithium-ion storage properties of silicon nanomaterials. J. Mater. Chem. A 2016, 4, 32-50. (9) J.-Y. Li; Q. Xu; G. Li; Y.-X. Yin; L.-J. Wan; Y.-G. Guo. Research progress regarding Si-based anode materials towards practical application in high energy density Li-ion batteries. Mater. Chem. Front. 2017, 1, 1691-1708. (10) W. Wang; P. N. Kumta. Reversible high capacity nanocomposite anodes of Si/C/SWNTs for rechargeable Li-ion batteries. J. Power Sources 2007, 172, 650-658. (11) S. Y. Kim; J. Lee; B. H. Kim; Y. J. Kim; K. S. Yang; M.S. Park. Facile Synthesis of Carbon-Coated Silicon/Graphite Spherical Composites for High-Performance Lithium-Ion Batteries. ACS Appl. Mater. Interfaces 2016, 8, 12109-12117. (12) M. Ko; S. Chae; J. Ma; N. Kim; H.-W. Lee; Y. Cui; J. Cho. Scalable synthesis of silicon-nanolayer-embedded graphite for high-energy lithiumion batteries. Nature Energy 2016, 113, 16113-16121. (13) Y.N. Jo; Y. Kim; J.S. Kim; J.H. Song; K.J. Kim; C.Y. Kwag; D.J. Lee; C.W. Park; Y.J. Kim. Si–graphite composites as anode materials for lithium secondary batteries. J. Power Sources 2010, 195, 6031-6036. (14) H. Li; C. Lu; B. Zhang. A straightforward approach towards Si@C/graphene nanocomposite and its superior lithium storage performance. Electrochim. Acta 2014, 120, 96-101. (15) W.J. Lee; T.H. Hwang; J.O. Hwang; H.W. Kim; J. Lim; H.Y. Jeong; J. Shim; T.H. Han; J.Y. Kim; J.W. Choi; S.O. Kim. N-doped graphitic self-encapsulation for high performance silicon anodes in lithium-ion batteries. Energy Environ. Sci. 2014, 7, 621-626. (16) W. Luo; Y-X. Wang; S. Chou; Y. Xu; W. Li; B. Kong; S. X. Dou; H. K. Liu and J. Yang. Critical thickness of phenolic resin-based carbon interfacial layer for improving long cycling stability of silicon nanoparticle anodes. Nano Energy 2016, 27, 255–264. (17) P.X. Han; Y.H. Yue; L.X. Zhang; H.X. Xu; Z.H. Liu; K.J. Zhang; C.J. Zhang; S.M. Dong; W. Ma; G.L. Cui. Nitrogen-doping of chemically


Page 19 of 23 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


reduced mesocarbon microbead oxide for the improved performance of lithium ion batteries. Carbon 2012, 50, 1355-1362. (18) M. Zhou; F. Pu; Z. Wang; T. Cai; H. Chen; H. Zhang; S. Guan. Facile synthesis of novel Si nanoparticles-graphene composites as highperformance anode materials for Li-ion batteries. Phys. Chem. Chem. Phys. 2013, 15, 11394-11401. (19) Y.S. Hu; R. Demir-Cakan; M.M. Titirici; J.O. Muller; R. Schlogl; M. Antonietti. Superior storage performance of a Si@SiOx/C nanocomposite as anode material for lithium-ion batteries. Angew. Chem. Int. Ed. 2008, 47, 1645-1649. (20) D. He; F. Bai; L. Li; L. Shen; H.H. Kung; N. Bao. Fabrication of Sandwich--structured Si Nanoparticles-Graphene Nanocomposites for Highperformance Lithium-ion Batteries. Electrochim. Acta 2015, 169, 409-415. (21) S. Yang; X. Feng; S. Ivanovici; K. Mullen. Fabrication of graphene-encapsulated oxide nanoparticles: towards high-performance anode materials for lithium storage. Angew. Chem. 2010, 49, 8408-8411. (22) F. M; Hassan; R. Batmaz; J. Li; X. Wang; X. Xiao; A. Yu; Z. Chen. Evidence of covalent synergy in silicon-sulfur-graphene yielding highly efficient and long-life lithium-ion batteries. Nat. Commun. 2015, 6, 8597-8608. (23) H. Sohn; D.H. Kim; R. Yi; D. Tang; S.-E. Lee; Y.S. Jung; D. Wang. Semimicro-size agglomerate structured silicon-carbon composite as an anode material for high performance lithium-ion batteries. J. Power Sources 2016, 334, 128-136. (24) J. H. Park; J. Moon; S. Han; S. Park; J. W. Lim; D. J. Yun; Y. K. Dong; K. Park and I. H. Son. Formation of Stable Solid–Electrolyte Interphase Layer on Few-Layer Graphene-Coated Silicon Nanoparticles for High-Capacity Li-Ion Battery Anodes. J. Phys. Chem. C 2017, 121, 2615526162. (25) J. Zhu; J. Yang; Z. Xu; J. Wang; Y. Nuli; X. Zhuang; X. Feng. Silicon anodes protected by a nitrogen-doped porous carbon shell for highperformance lithium-ion batteries. Nanoscale 2017, 9, 8871-8878. (26) M. Kruk; M. Jaroniec. Gas Adsorption Characterization of Ordered Organic-Inorganic Nanocomposite Materials. Chem Mater 2001, 132, 3169-3183. (27) J. Wang; D.-H. Liu; Y.-Y. Wang; B.-H. Hou; J.-P. Zhang; R.-S. Wang; X.-L. Wu. Dual-carbon enhanced silicon-based composite as superior anode material for lithium ion batteries. J. Power Sources 2016, 307, 738-745. (28) H. Park; S. Choi; S.-J. Lee; Y.-G. Cho; G. Hwang; H.-K. Song; N.-S. Choi and S. Park. Design of an ultra-durable silicon-based battery anode material with exceptional high-temperature cycling stability. Nano Energy 2016, 26, 192-199. (29) C. J. Tang; J. X. Zhu; X. J. Wei; L. Hea; K. N. Zhao; C. Xua; L. Zhou; B. Wang; J. S. Sheng and L. Q. Mai. Copper silicate nanotubes anchored on reduced graphene oxide for long-life lithium-ion battery. Energy Storage Materials, 2017, 7, 152-156. (30) D. Chen; W. Liao; Y. Yang; J. Zhao. Polyvinyl alcohol gelation: A structural locking-up agent and carbon source for Si/CNT/C composites as high energy lithium ion battery anode. J. Power Sources 2016, 315, 236-241. (31) Y.-C. Zhang; Y. You; S. Xin; Y.-X. Yin; J. Zhang; P. Wang; X.-s. Zheng; F.-F. Cao and Y.-G. Guo. Rice husk-derived hierarchical silicon/nitrogen-doped carbon/carbon nanotube spheres as low-cost and high-capacity anodes for lithium-ion batteries. Nano Energy 2016, 25, 120-127.



Page 20 of 23

(32) A. Sano; M. Kurihara; T. Abe; Z. Ogumi. Studies on Lithium-Ion Diffusion in Heat-Treated CNBs by Microelectrode Method. J. Electrochem. Soc. 2009, 156, 639-644. (33) Z. Luo; Q. Xiao; G. Lei; Z. Li; C. Tang. Si nanoparticles/graphene composite membrane for high performance silicon anode in lithium ion batteries. Carbon 2016, 98, 373-380. (34) W. Cheng; Felix Rechberger; Gabriele Ilari; H. Ma; W.-I. Lind; M. Niederberger. Amorphous cobalt silicate nanobelts@carbon composites as a stable anode material for lithium ion batteries. Chem. Sci. 2015, 6, 6908-6915. (35) N. Kim; C. Oh; J. Kim; J.-S. Kim; E.D. Jeong; J.-S. Bae; T.E. Hong; J.K. Lee. High-Performance Li-Ion Battery Anodes Based on SiliconGraphene Self-Assemblies. J. Electrochem. Soc. 2016, 164, 6075-6083. (36) Zixu Suna; Guangjin Wangb; Tingwei Caia; Hangjun Yinga; Wei-Qiang Hana. Sandwich-structured graphite-metallic silicon@C nanocomposites for Li-ion batteries. Electrochim. Acta 2016, 191, 299-306. (37) Y. Ma; H. Desta Asfaw; C. Liu; B. Wei and K. Edström. Encasing Si particles within a versatile TiO 2−x F x layer as an extremely reversible anode for high energy-density lithium-ion battery. Nano Energy 2016, 30, 745-755. (38) Y.P. Wu; E. Rahm; R. Holze. Carbon anode materials for lithium ion batteries. J. Power Sources 2003, 114, 228-236. (39) L. Wu; J; Yang; X. Zhou; M. Zhang; Y. Ren; Y. Nie. Silicon nanoparticles embedded in a porous carbon matrix as a high-performance anode for lithium-ion batteries. J. Mater. Chem. A 2016, 4, 11381-11387. (40) D. Gueon; D.-Y. Kang; J.S. Kim; T.Y. Kim; J.K. Leeb; J.H. Moon. Si nanoparticles-nested inverse opal carbon supports for highly stable lithium-ion battery anodes. J. Mater. Chem. A 2015, 3, 23684-23689. (41) N. Lin; J. Zhou; L. Wang; Y. Zhu; Y. Qian. Polyaniline-assisted synthesis of Si@C/RGO as anode material for rechargeable lithium-ion batteries. ACS Appl. Mater. Interfaces 2015, 7, 409-414. (42) Y. Yao; N. Xu; D. D. Guan; J. T. Li; Z. C. Zhuang; L, Zhou; C. W. Zhou; C. W. Shi; X. Liu and L. Q. Mai. Facet-Selective Deposition of FeOx on α-MoO3 Nanobelts for Lithium Storage. ACS Appl. Mater. Interfaces 2017, 9, 39425-39431. (43) N. Liu; Z. Lu; J. Zhao; M. T. McDowell; H. W. Lee; W. Zhao and Y. Cui. A pomegranate-inspired nanoscale design for large-volume-change lithium battery anodes. Nature Nanotech. 2014, 9, 187-192. (44) M.A. Rahman; G. Song; A.I. Bhatt; Y.C. Wong; C. Wen. Nanostructured Silicon Anodes for High-Performance Lithium-Ion Batteries. Adv. Funct. Mater. 2016, 26, 647-678. (45) Fu-Sheng Li; Yu-Shiang Wu; J. Chouc; N.-L. Wu. A dimensionally stable and fast-discharging graphite–silicon composite Li-ion battery anode enabled by electrostatically self-assembled multifunctional polymer-blend coating. Chem. Commun. 2015, 51, 8429-8431. (46) Z.L Xu; Y Gang; Mohammad Akbari Garakani; Sara Abouali; J.-Q. Huang; J.-K. Kim. Carbon-coated mesoporous silicon microsphere anodes with greatly reduced volume expansion. J. Mater. Chem. A 2016, 4, 6098-6106. (47) W. Sun; R.Z Hu; M Zhang; J.W Liu; M. Zhu. Binding of carbon coated nano-silicon in graphene sheets by wet ball-milling and pyrolysis as high performance anodes for lithium-ion batteries. J. Power Sources 2016, 318, 113-120.


Page 21 of 23 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


(48) Y. Xu; E. Swaans; S. Chen; S. Basak; P. P. R. M. L. Harks; B. Peng; H. W. Zandbergen; D. M. Borsa and F. M. Mulder. A high-performance Li-ion anode from direct deposition of Si nanoparticles. Nano Energy 2017, 38, 477-485.




Page 22 of 23

Page 23 of 23


1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


Self-assembly of Silicon@Oxidized Mesocarbon Microbeads

Recommend Documents