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Micro-sized porous SiOx@C composites synthesized through aluminothermic reduction from rice husks and used as anode for lithium-ion batteries Jinlong Cui, Yongfu Cui, Shaohui Li, Hongliang Sun, Zhongsheng Wen, and Juncai Sun ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.6b10260 • Publication Date (Web): 20 Oct 2016 Downloaded from http://pubs.acs.org on October 24, 2016

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ACS Applied Materials & Interfaces

Micro-sized porous SiOx@C composites synthesized through aluminothermic reduction from rice husks and used as anode for lithium-ion batteries Jinlong Cuia, Yongfu Cuia, Shaohui Lib, Hongliang Suna, Zhongsheng Wena, Juncai Sun*a a

Institute of Materials and Technology, Dalian Maritime University, Dalian 116026, P. R. China

b

School of Materials Science and Engineering, Nanyang Technological University, Singapore

639798, Singapore KEYWORDS: lithium-ion batteries, anode, porous SiOx@C, rice husks, aluminothermic reduction ABSTRACT: Micro-sized porous SiOx@C composites used as anode for lithium-ion batteries (LIBs) are synthesized from rice husks (RHs) through low-temperature (700 °C) aluminothermic reduction. The resulting SiOx@C composite shows mesoporous irregular particle morphology with a high specific surface area of 597.06 m2/g under the optimized reduction time. This porous SiOx@C composite is constructed by SiOx nanoparticles uniformly dispersed in the C matrix. When tested as anode material for LIBs, it displays considerable specific capacity (1230 mAh/g at a current density of 0.1 A/g) and excellent cyclic stability with capacity fading of less than 0.5% after 200 cycles at 0.8 A/g. The dramatic volume change for the Si anode during lithium-

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ion (Li+) insertion and extraction can be successfully buffered because of the formation of Li2O and Li4SiO4 during initial lithiation process and carbon coating layer on the surface of SiOx. The porous structure could also mitigate the volume change and mechanical strains and shorten Li+ diffusion path length. These characteristics improve the cyclic stability of the electrode. This low-cost and environment-friendly SiOx@C composite anode material exhibits great potential as an alternative for traditional graphite anodes. 1. INTRODUCTION Petroleum fuel shortage, environmental crisis, and escalating consumer demands have impelled exploration of cleaner and efficient electrochemical energy systems.1 LIBs, as a new type of power source, have gained increased attention because of their high energy density, outstanding cycling performance and low cost. LIBs could satisfy the requirements for development of electric vehicles (EVs) and plug-in hybrid EVs.2 Current commercial graphite anodes for LIBs have a theoretical capacity of 372 mAh/g, which limits the extensive applications of these materials.3 Si-based anode materials have drawn much interest because of their relatively low working potential and high theoretical capacity of 4200 mAh/g, as well as the natural abundance of Si.4,5 However, the remarkable change in the volume of Si during alloying/dealloying leads to internal cracks, loss of electrical contact, and increased impedance. These characteristics seriously affect the cycling performance of Si electrodes.6,7 Numerous approaches have been explored to surmount these defects and boost the electrochemical properties of Si-based anodes. C-coated silicon oxide (SiOx) composites have obvious advantages over pure Si in terms of cycling performance; Li2O and Li4SiO4 generated during first lithiation process could serve as buffer components to enhance the cyclic stability of the electrode.8−10 The C matrix coated on the surface of SiOx particle could also contribute to


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controlling changes in the volume of the electrode during cycling, improving electronic conductivity, and preventing agglomeration of SiOx.3,11,12 Therefore, the preparation of SiOx@C anodes has been extensively investigated using ethyltriethoxysilanes [EtSi(OEt)3], polysiloxane, or tetraethyl orthosilicate (TEOS) as Si source and resorcinol/formaldehyde (RF), epoxy resin, cetyltrimethyl ammonium bromide (CTAB), cellulose fibers, polyvinylpyrrolidone (PVP), citric acid, polyvinylidene fluoride (PVdF), or glucose as C source.3,6,8,13−19 Excellent electrochemical performance has been achieved. Rice husk (RH), as a renewable source, is a widely available and inexpensive agricultural residue from rice-producing countries. Approximately 600 million tons of rice is produced annually worldwide, generating more than 100 million tons of RHs.20,21 In china, only a small portion of the generated RHs is used to compost organic fertilizers or cultivate edible fungi. Most of them are usually used as a low-value energy resource, burned in the field or abandoned, which are detrimental to the environment.22 RHs are mainly comprised of organic lignin, cellulose, hemi-cellulose, and inorganic Si compound (Table S1). Thus, RHs can be used to prepare lowcost electrode materials, such as carbonaceous and silicon-based anode materials. Several studies have explored the synthesis of anode materials by using C23−26 and nanostructured Si27−29 with RHs as raw materials. Moreover, silica, which naturally exists as nanoparticles, is assimilated from soil in the form of silicic acid during the growth of rice.27,30 Silica is uniformly deposited on the cell wall material as a polymer of hydrated and amorphous silica (Scheme 1). Therefore, silica is homogeneously coated by the disordered C generated during the carbonization of RHs under Ar atmosphere without any complex C coating technology (Figure S1). In addition, silica remains uniformly distributed when the carbonized RHs are activated using ZnCl2 as activation agent to obtain a porous structure.31 Thus, the RHs exhibit great potential as the only raw


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material suitable for preparing porous SiOx@C composites. The homogeneous distribution of silica in the RHs carbon matrix could promote the electrical contact between SiOx and electrolyte as well as prevent the desquamation and pulverization of the electrode materials. These characteristics greatly contribute to the remarkable cyclic performance of the SiOx@C electrode.3,6,8,32 Recently, Ju et al. have prepared SiOx/C composite using RHs as raw material by a two-step carbonization process. After an initial carbonization process, the carbonized RHs powder was further heated to 900 oC to obtain SiOx/C composite in the Ar/H2 atmosphere. The resulting SiOx/C composite delivered a reversible capacity of about 600 mAh/g at a current density of 0.1 A/g.32 In this study, we synthesized porous SiOx@C composite using RHs as the raw material of both Si and C source through a low-temperature aluminothermic reduction process. The synthesis of SiOx@C samples is described in Scheme 1. The RHs powder was carbonized under Ar atmosphere after pretreatment. Then, the carbonized RHs powder was activated with ZnCl2 to obtain porous carbonized RHs powder (PCRH) under Ar atmosphere. Finally, PCRH was directly reduced using Al powder to receive porous SiOx@C samples (Scheme 1). The specific capacity of the resulting porous SiOx@C electrode material can achieve as large as 1230 mAh/g at a current density of 0.1 A/g, which is significantly higher than that prepared using chemical reagent as raw material (840 mAh/g at 0.1 A/g).3 The overall method is controllable and easy to scale-up, and does not require expensive reagents. This work provides a method for possible industrial applications to transform RHs into commercial anode materials.


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Scheme 1. Synthesis of SiOx@C composites from rice husks. 2. EXPERIMENAL SECTION RHs were collected from a rice mill of Wuchang City in China. The chemical components of RHs are listed in Table S1. All chemical reagents used in the synthesis process were provided by Tianjin Kermel Co. Ltd (analytic grade). RHs were rinsed drastically with deionized water to eliminate adhering dust and soil. The RHs were then soaked in 2 M HCl for 12 h at room temperature to eliminate the metal impurities. The cleaned RHs were collected by filtration, rinsed to neutrality with deionized water, and dried at 100 °C for 12 h. Then, the cleaned RHs were crushed to 60 mesh and stored in double ziplock bags for subsequent experimental use. The RHs powder was carbonized under Ar atmosphere in a tube furnace at 550 °C for 1 h. The carbonized RHs powder was impregnated with ZnCl2 solution (weight ratio of ZnCl2-tocarbonized RHs powder was 2:1) for 12 h and then dried overnight at 100 °C. The impregnated carbonized RHs powder was placed in a quartz boat and heated from room temperature to 550 °C


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at a heating rate of 10 °C/min under Ar environment in the tube furnace. The sample was sintered at 550 °C for 1.5 h and then naturally cooled to ambient temperature. The resulting product (PCRH) was washed with 2 M HCl and rinsed with deionized water. The PCRH was dried again in air at 100 °C for overnight. Then, 0.5 g of PCRH was thoroughly ground with 0.25 mg of Al powder (200 mesh) and sealed in 430 stainless steel reactor, which was designed by our group, inside an Ar-filled glove box. The reactor was heated to 700 °C at a heating rate of 10 °C/min inside the tube furnace under Ar atmosphere at reduction times of 1, 2, 3, and 4 h. The corresponding products were denoted as S1, S2, S3, and S4, respectively. Then, the reactor was removed from the furnace after the temperature was reduced to below 50 °C. The obtained powders were immersed in 2 M HCl for 8 h to eliminate Al2O3 and excessive Al. The powders were dried under vacuum at 50 °C overnight, yielding the SiOx@C samples. As a control experiment, the sample labeled as S0 was also synthesized via direct reduction of carbonized RHs powder using the synthesis process of S3. To investigate the contribution of free carbon and SiOx to the observed capacity of S3, some S3 powder was soaked in 5 M HF for 6 h to receive free carbon without any SiOx. XRD spectra were received on a RigakuD/MAX-3A X-ray diffractometer (XRD) using a Cu Kα radiation. Raman spectra were obtained via a T64000 triple Raman system with a 532 nm Ar-ion laser. The O and C contents of SiOx@C composites were measured using a LECO/ONH836 O/N/H analyzer and a LECO/CS230 C/S analyzer, respectively. The amount of Si was calculated as a difference to 100%, because there was less than 1 wt% of other elements in the material. The microstructures of samples were determined using a field-emission scanning electron microscope (FE−SEM), and a high-resolution transmission electron microscope (HRTEM). The elemental distribution of samples was measured via an


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energy-dispersive X-ray spectrometer (EDS) attached to the SEM. The lattice structure of the typical sample was analyzed by selected area electron diffraction (SAED) before and after cycling. The specific surface area (SBET) and pore size of PCRH and SiOx@C composites were analyzed via the Brunauer–Emmett–Teller (BET) method. The particle size distributions of SiOx@C samples, PCRH, and carbonized RHs were obtained by a laser particle-size analyzer. Fourier-transform infrared reflection (FTIR) spectra of SiOx@C composites were recorded with an FTIR instrument using KBr pellet. The slurry, which contained 70 wt% of the SiOx@C samples, 15 wt% of C black, and 15 wt% of PVdF, was uniformly smeared onto a thin copper foil and dried to prepare electrodes. The electrochemical performance of the samples were tested in 2025 coin cells, with SiOx@C electrodes as the working electrode and Li foil as the counter and reference electrode. Mass loading of active substance on the copper foil is 1.8-2.3 mg/cm2. Briefly, 1 M LiPF6/EC + DMC (1:1 in volume) solution was applied as electrolyte. Galvanostatic charge/discharge measurements were performed using a Land CT2100 battery test system from 0.01 to 3.0 V (vs. Li/Li+). Cyclic voltammetry (CV) and electrochemical impedance spectroscopy (EIS) were conducted at ambient temperature with CHI660D electrochemical station (CHI Instrument). The scan rate for CV was 0.05 mV/s with a voltage range of 0.01−3.0 V. EIS was carried out at open circuit potential from 100 kHz to 0.01 Hz with an amplitude of 10 mV. The cells in the EIS tests with different SiOx@C samples as working electrode had the same mass of active substance. 3. RESULTS AND DISCUSSION


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Almost all SiOx@C composites are produced using the carbothermic reduction of silica. This process not only requires high temperatures, but also has a relatively weak reducing capacity.3,6,8,13−19 The reaction temperature when Al powder is used as reducing agent for silica in PCRH could decrease to 700 °C, following a relatively strong reducing power. In addition, contrary to Mg powder as reducing agent, the excess Al powder could not react with the obtained Si leading to the loss of Si during reduction.27,33,34 The possible chemical reactions between SiO2, Al, and C at 700 °C are shown in Reactions (1) and (2) 33,35−38, as follows: 3SiO 2 ( s ) +4Al ( l ) → 3Si ( s ) +2Al2 O3 ( s )

(1)

2Al 2O3 ( s ) +9C ( s ) → Al4 C3 ( s ) +6CO ( g )

(2)

Al2O3, Al4C3, and unreacted Al are removed by HCl treatment, yielding SiOx@C composites and some SiC. Al melt at 700 °C could not react with C but has certain permeability that allows Al to pass through the C matrix to react with SiO2.39 The atomic ratio of O to Si of sample S0 is obviously higher than that of S3 (Table S3), although these samples were synthesized similarly. The only difference in the synthesis of S0 and S3 is the SBET of their precursors. Carbonized RHs used as the precursor of sample S0 has a SBET of 117.39 m2/g, whereas PCRH, as the precursor of sample S3, has a SBET of 362.19 m2/g (Table S4). Greater SBET of the precursor of SiOx@C composites indicates better contact of the aluminum melt with the silica in the precursor, resulting in higher reduction of silica. The O/Si ratio decreases with increasing reduction time (Table S3), demonstrating that the degree of the aluminothermic reduction of PCRH increases with reduction time. The particle size of the SiOx@C composites reduces with increasing reduction time (Figure S2a), and the particle size of these composites is smaller than that of their precursors (Figure S2b). In addition, the average pore diameter of SiOx@C composites increases with prolonged


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reduction time (Table S4). The possible reason for these phenomena is that the Al2O3 byproduct near the C matrix is further converted to Al4C3 [Reaction (2)]. Then, Al4C3 could react with HCl releasing CH4 gas during washing process (Al4C3+12HCl→4AlCl3+3CH4↑), which causes a certain loss of C.37 The C loss ratio of PCRH during the reduction process increases with reduction time (Table S3). This ratio was calculated and shown in the Supplementary Information. This result further proves that the Al2O3 byproduct near the C matrix is further converted to Al4C3. The phases and crystallinities of the four SiOx@C samples were examined by XRD. As shown in Figure 1a, the characteristic diffraction peaks at 2θ=28.4°, 47.2°, 56.1°, 69.2°, 79.5°, and 88.0° can be assigned to the (111), (220), (311), (400), (331), and (422) crystal planes of Si crystallites (JCPDS No. 27−1402), respectively.13,15,40 Three additional diffraction peaks corresponding to the (111), (220), and (311) planes of SiC (JCPDS No. 29−1129)13 are observed in S4. No obvious diffraction peaks from SiC are found in the XRD patterns of other three samples. In addition, the wide peak at approximately 2θ=23° can belong to the overlap between the amorphous C and amorphous SiOx in each spectrum of SiOx@C samples.13,15 The Raman spectra of the four SiOx@C samples (Figure 1b) show three characteristic peaks at approximately 517, 1343, and 1598 cm−1 which are in excellent accordance with the Si crystal and the D-band and G-band of C, respectively.41,42 The relative dimensional intensities from the D-band to the Gband of the SiOx@C samples are all estimated to be close to 1.0, demonstrating a low degree of graphitization and the existence of disordered C.40,43 This finding is completely in conformity with the XRD result. A broad peak at 796 cm−1, which appears on the Raman spectrum of S4, is related to SiC.44 This result further testifies the existence of the SiC phase in S4. This characteristic is probably due to the fact that the obtained silicon could react with C to form SiC


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(Si+C→SiC) with aluminum as catalyst when the heating time is more than 4 h at 700 °C.38,45 The electrochemical behavior of the obtained SiOx@C samples was investigated as anode materials of LIBs. Figure 1c exhibits the cyclic performance of the obtained SiOx@C samples from the second cycle at a current density of 0.1 A/g. The specific capacity rises with the increase in reduction time of samples, except for S4, which displays a relatively lower specific capacity compared with S3 (Figure 1c). In addition, all SiOx@C electrodes exhibit excellent cycling performance without any tendency of degradation (Figure 1c) and low initial coulombic efficiency because of the irreversible reduction of SiOx and decomposition of electrolyte (Figure S3). S3 shows the largest specific capacity of 1230 mAh/g. The high Si content is very favorable for Li+ storage of SiOx phase. Therefore, the specific capacity of SiOx@C samples increases with the Si content in the resulting samples. Moreover, the Si−C bonds are inactive for Li+ insertion and thus reduce the specific capacity of SiOx@C anode material.13 S4 exhibits a relatively lower reversible specific capacity because of its Si−C bonds, although S4 has the smallest O/Si ratio (Table S3). EIS is a useful technique to characterize the electrochemical behavior of electrode materials. Figure 1d presents the Nyquist plots of the synthesized SiOx@C electrodes after 25 cycles from 10 mHz to 100 kHz. The plots in Figure 1d are composed of one depressed semicircle in the high-frequency range corresponding to charge transfer resistance and the solid electrolyte interface (SEI) film resistance and an inclined line in the low-frequency region associated with the Li+ diffusion within electrodes.13,18 The electrochemical impedance spectra of SiOx@C electrodes decrease with increasing reduction time, except for S4. Moreover, the diameter of the semicircle for S3 electrode is much smaller than those of other SiOx@C electrodes, manifesting that S3 electrode exhibits smaller electrode reaction resistance.


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Figure 1. (a) XRD patterns and (b) Raman spectra of SiOx@C samples. (c) Cyclic performance of SiOx@C electrodes from the second cycle at 0.1 A/g. (d) EIS of SiOx@C electrodes after 25 cycles. As shown in Figure 2a, S3 shows irregular particles morphology with a wide-size distribution and there are many pores on the surface of it, which is in conformity with the BET results (Figure S4 and Table S4). The sample with larger SBET and more pores can reduce Li+ diffusion path length to improve the electrochemical properties, such as cycling ability and rate performance.14,46 Elemental mapping images from SEM show the relatively uniform distribution of O, C and Si in S3 (Figure 2b). In order to describe the detailed internal structure of S3, TEM characterizations are applied in this work. Figure 2c obviously displays the stark contrast between SiOx and C in which the black core corresponds to the SiOx nanoparticles. The SiOx


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nanoparticles are uniformly dispersed in the C matrix and no evident aggregation of SiOx nanoparticles is observed, which is consistent with the result of EDS mapping analysis. These results indicate the good contact between the SiOx particle and C matrix. Therefore, the aggregation of SiOx particles could be efficiently prevented during Li+ insertion and extraction.36 The HRTEM image of S3 (Figure 2d) also shows that the SiOx particles are coated by a layer of disordered C with varying thickness, which could effectively alleviate the remarkable change in the volume of Si during cycling. Lattice fringes can be observed in the ordered part, and the measured d-spacing values of the composite are approximately 0.313 and 0.192 nm, which can agree with the crystalline planes of Si (111) and (220), respectively.42 The SAED pattern of S3 (inset Figure 2d) illustrates that the diffraction spots of S3 can be placed in four concentric rings, and the lattice plane spacing calculated from these rings are in line with Si (111), (220), (311), and (511).16,17 These phenomena are completely in conformity with the XRD results.


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Figure 2. (a) SEM pictures of S3, (b) EDS mappings of C, O, and Si in S3. (c) TEM images and (d) HRTEM image and SAED patterns (inset) of S3. Figure S5a shows the cyclic voltammograms (CV) for S3 electrode measured at a scan rate of 0.05 mV/s in the voltage range of 0.01–3.0 V after the 1st, 2nd, 10th, 22nd, and 130th cycle. A broad cathodic peak is observed at approximately 0.7 V for the first cathodic half-cycle (Li+ insertion). This characteristic is owing to the electrolyte decomposition and accompanying SEI film formation.14,40,47,48 Then, the peak at 0.7 V disappears at subsequent cycles, suggesting the formation of stable SEI layer in the first discharge process. The broad anodic peaks at 0.2–


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1.5 V correspond to the dealloying processes of LixSi, which is in conformity with other researcher’s observation.14 Figure 3a presents the discharge/charge profiles of S3 electrode at a current density of 0.1 A/g for the 1st, 2nd, 10th, 22nd, and 130th cycle. In the first cycle, the voltage drops quickly and forms one short plateau at 0.7 V, which is in agreement with the first CV curve. Otherwise, the curves of the 22nd and 130th cycles overlap, confirming that the S3 electrode has excellent cycling stability. The cyclic performance of S3 electrode was tested at 0.1 A/g with voltage range of 0.01–3 V, and the result is plotted in Figure 3b. The S3 electrode shows high initial specific capacity of 3096.6 mAh/g, but the initial coulombic efficiency is only 45.04%. The high irreversible capacity is mainly ascribed to the generation of inactive Li4SiO4 and Li2O. In addition, the SEI film formation could also consume some Li+ in the electrolyte irreversibly.13,14,36 Nevertheless, the S3 electrode exhibits excellent cycling stability after 22 cycles and high reversible capacity, approximately 1230 mAh/g, which is nearly 4 times higher than the theoretical specific capacity of a traditional graphite anode (LiC6, 370 mAh/g). Although the cyclic performance curve is slightly fluctuating, this value is much larger than those of the SiOx@C samples synthesized using chemical reagent as raw materials. The obtained value is also close to those of the Si@SiOx@C electrode (1640 mAh/g at 0.1 A/g),41 Si/C electrode (1630 mAh/g at 0.4 A/g),49 and SiOx/C@RGO electrode (1284 mAh/g at 0.1 A/g).50 Table 1 lists the comparison of specific capacities of SiOx@C samples synthesized using chemical reagent as raw materials. The coulombic efficiency of S3 is gradually stabilized at almost 100% in the subsequent cycles. The S3 electrode also exhibits good rate capability. The electrode displays a stable capacity at each current density when the current density is gradually increased from 0.1 to 1.6 A/g. The specific capacity can be recovered perfectly to the original value every time when the current density is returned to 0.2 A/g (Figure 3c), indicating that the


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electrode is not destroyed during the high rate test. Figure S5b shows that the resistance of S3 electrode reduces because of activation and ultimately keeps almost 95 Ω. The diameter of the depressed semicircle shows no significant increase after 22 cycles, indicating that S3 electrode could well adapt to the volume expansion of SiOx. The electrochemical performance was also investigated under high discharge rate. The cycling performance of S3 electrodes were tested at current densities of 0.2, 0.4, and 0.8 A/g (Figure S5c and d). The electrode exhibits good stability with discharge capacities of approximately 946.7, 751.7, and 707.1 mAh/g at 0.2, 0.4, and 0.8 A/g, respectively. Otherwise, in order to investigate the contribution of free carbon and SiOx to the obtained capacity of S3, the cycling performance of the free carbon electrode was also studied in this work. Its discharge specific capacity keeps at about 410 mAh/g at 0.1 A/g (Fig. S6). This means that the specific capacity of SiOx in S3 is 1990 mAh/g (410 mAh/g×0.481 (for free carbon) + 1990 mAh/g×0.519 (for SiOx) = 1230 mAh/g (for the composite)). Therefore, nearly 84% of the S3’ specific capacity comes from the SiOx nanoparticles. Table 1 Comparison of specific capacities of the SiOx@C composites. Materials

Specific capacity

Current density

Cycling capability

Ref.

(mAh/g)

(A/g)

TEOS and epoxy resin

800

0.1

93.4% retention after 50 cycles

[3]

EtSi(OEt)3, RF and CTAB

720

0.1

100% retention after 350 cycles

[6]

TEOS and citric acid

840

0.1


[13]

EtSi(OEt)3 and RF

630

0.05


[14]

TEOS and glucose

674.8

0.1


[15]

TEOS and CTAB

1050

0.15


[16]

TEOS and cellulose fibers

940

0.1


[17]

TEOS and sucrose

820

0.1


[18]

TEOS, CTAB and sucrose

780

0.1


[19]

Rice husks

1230

0.1


This work


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Figure 3. (a) Charge/discharge curves at 0.1 A/g for the 1st, 2nd, 10th, 22nd, and 130th cycles, (b) Cyclic performance at 0.1 A/g and (c) Rate performance of S3 electrode. The satisfactory electrochemical properties and lasting structural stability of obtained SiOx@C composites could be mainly attributed to several factors. First, samples with larger SBET and more pores can shorten Li+ pathway, provide large interface to facilitate electrochemical reactions, and supply sufficient void space, which successfully mitigates the mechanical stress caused by volume change.17 This hypothesis is testified using S0 as contrast sample. S0 was


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synthesized via direct reduction of carbonized RHs powder using the synthesis process of S3. As shown in Figure S7, there is almost no pore on the surface of S0. The SBET of S0 (155.88 m2/g) is much smaller than that of S3 (597.06 m2/g) (Table S4). Meanwhile, its specific capacity (approximately 543.4 mAh/g) is much lower than that of S3 (approximately 1230 mAh/g) (Figure 4a), and its resistance (approximately 250 Ω) is much larger than that of S3 (approximately 95 Ω) (Figure 4b). Second, Figure S8 shows three obvious absorption peaks at 1080, 810, and 460 cm−1 which are associated with the Si–O bond of SiOx. Therefore, the inert Li4SiO4 and Li2O would be generated during the first lithiation process. The HRTEM image in Figure 5a is used to verify the irreversibility of the crystallized Li4SiO4. After the electrode is charged to 3.0 V, the Li4SiO4 phase is still observed except Si nanocrystallites, indicating that Li4SiO4 phase is irreversible during cycling.12,15,51 The result could be further proved by the SAED pattern of S3 electrode charged to 3.0 V (Figure S9). Although Li2O is not observed in the HRTEM image as its amorphous form, the reaction mechanism could still be summed up as follows:3,15,17,52 SiO x +2xLi + +2xe − → xLi 2 O+Si

(3)

SiO x +xLi + +xe − → ( x/4 ) Li 4SiO4 + (1-x/4 ) Si

(4)

Reactions (3) and (4) only take place in the initial discharge process, which is an important factor for the large irreversible capacity loss at the first cycle. However, the Li4SiO4 and Li2O can act as buffer components to mitigate the remarkable change in volume caused by the lithiation of active Si to a certain degree. Third, C, Si, and O present a relatively homogeneous distribution in the resulting samples, suggesting the good electrical contact between SiOx particle and C matrix. Aggregation of SiOx particles can be efficiently prevented during Li+ insertion and extraction, thereby remarkably contributing to the superb cyclic performance of the obtained


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SiOx@C electrode. Finally, the disordered C coating layer on the surface of SiOx particle not only provides good electronic conductivity of the active material but also prevents the considerable change in volume caused by the lithiation of active Si.53,54 The SEM pictures of S3 electrode before and after cycling further confirm this characteristic (Figure 5b and c). In addition to a thick SEI film, the desquamation and pulverization of electrode materials, which frequently exist in conventional Si electrodes, are not seen in this study. These characteristics could guarantee high reversible capacity and superb cyclic stability of the resulting SiOx@C composites.

Figure 4. (a) Cyclic performance of S0 electrode at 0.1 A/g. (b) EIS of S0 and S3 electrodes after 25 cycles.


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Figure 5. (a) HRTEM image S3 electrode charged to 3.0 V. SEM pictures of S3 electrode (b) before cycling and (c) after 200 cycles. 4. CONCLUSION We have successfully synthesized a micro-sized porous SiOx@C composite with RHs as raw material using a low-temperature (700 °C) aluminothermic reduction process. The obtained composite shows an impressive electrochemical performance with a specific capacity of approximately 1230 mAh/g at 0.1 A/g in the voltage range of 0.01–3 V. The electrode displays excellent cyclic stability with less than 3% degradation after 200 cycles at 0.1 A/g. Therefore, the obtained SiOx@C composite is a promising alternative for graphite as a new anode material for LIBs in the future. This work also provides a new effective method for harnessing RHs. ASSOCIATED CONTENT Supporting Information Four tables of chemical compositions and porous structure for rice husks, carbonized RHs, PCRH and each SiOx@C samples. First cycle voltage profiles, FTIR spectra, particle size distribution curves and N2 adsorption/desorption isotherms of each SiOx@C samples and PCRH. SAED pattern, cyclic voltammetry, EIS and cycling performance of S3. Cycle performance of


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the free carbon. SEM pictures of S0 and carbonized RHs. This material is available free of charge via the Internet at http://pubs.acs.org. AUTHOR INFORMATION Corresponding Author * Tel: +86-411-84727959. Fax: +86 411 84725960. E-mail address: [email protected]. Author Contributions The manuscript was written through contributions of all authors. We have given approval to the final version of the manuscript. The author Jinlong Cui was responsible for the implementation of the experiment, analysis of tests and data results, and manuscript writing. Yongfu Cui and Hongliang Sun were responsible for the collection of raw materials. Shaohui Li and Zhongsheng Wen were responsible for the revision of English expression. The author Juncai Sun was responsible for the communication and coordination of laboratory equipment, guidance on experiments and proofread the manuscript carefully. Funding Sources This work was supported by the National Natural Science Foundation of China (Grants 51479019 and 21476035) and Fundamental Research Funds for the Central Universities (3132016341). Notes The authors declare no competing financial interest.


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ACKNOWLEDGMENT We are very grateful to Professor Feng Zhou and Professor Yingtao Yu (Environmental Science and Engineering, Dalian Maritime University) for help of BET and Raman spectrum measurements, respectively. We also thank the Analytical and Testing Centre of the Department of Materials Science and Engineering for performing characterization of various samples. Financial support from the National Natural Science Foundation of China (No. 51479019 and 21476035) and the Fundamental Research Funds for the Central Universities (3132016341) are gratefully acknowledged.

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