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Rational design of Si@SiO2/C composite using sustainable cellulose as carbon resource for anode in lithium-ion batteries Dazhi Shen, Chaofan Huang, Lihui Gan, Jian Liu, Zhengliang Gong, and Minnan Long ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.7b16724 • Publication Date (Web): 09 Feb 2018 Downloaded from http://pubs.acs.org on February 10, 2018

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ACS Applied Materials & Interfaces

Rational design of Si@SiO2/C composite using sustainable cellulose as carbon resource for anode in lithium-ion batteries Dazhi Shen, Chaofan Huang, Lihui Gan, Jian Liu*, Zhengliang Gong*, Minnan Long* College of Energy, Xiamen University, Xiamen, 36105, China * Corresponding author. Tel.:+86-592-5952787; Fax:+86-5922188053

ABSTRACT In this work, we propose a novel and facile route for rational design of Si@SiO2/C anode material by using sustainable and environment-friendly cellulose as carbon resource. In order to simultaneously obtain SiO2 layer and carbon scaffold, a special designed homogeneous cellulose solution and commercial Si nanopowder are used as the start materials, and cellulose/Si composite is direct assembled by an in-situ regenerating method. Subsequently, Si@SiO2/C composite is obtained after carbonization. As expected, Si@SiO2 is homogenously encapsulated in cellulose-derived carbon network. The obtained Si@SiO2/C composite shows a high reversible capacity of 1071 mA h g−1 at the current density of 420 mA g−1 and 70 % capacity retention after 200 cycles. This novel, sustainable and effective design is a promising approach to obtain high performance and cost-effective composite anodes for the practical applications. Keywords: Cellulose, Silicon anode, Li-ion batteries, Carbon, Electrochemical performance

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1 Introduction Lithium-ion batteries have the advantages of high energy density, good rate capability, long cycle life, low self-discharge rate

1-3

. Therefore, in the past several decades, they are widely used in many

aspects, including electric vehicles, large scale energy storage, and power tools

4-5

. It is well known that

the commercial graphite anode has been widely used as the anode material in lithium-ion batteries. When Li ions are inserted into natural graphite, lithium-intercalated graphitic compounds are reversibly formed at about 0.1-0.2 V offering a theoretical specific capacity of 372 mAh g-1

6

. Low volume

expansion/contraction of 9% during the lithiation-delithiation is the major advantage of graphite as the anode material. However, graphite anodes used in the current commercial Li-ion batteries can not meet high energy density challenges due to its low theoretical capacity 3. Among the many anode candidates, silicon (Si) has been considered as one of the most promising anode material to replace or complement graphite in lithium-ion batteries 7. Silicon has a high theoretical capacity at 4200 mAh g-1 for forming Li22Si5 and a very low discharge potential of around 0.2 V with respect to lithium metal 8. Moreover, Silicon is abundant in source, environmental friendly and low cost 1. However, the severe volume expansion/contraction during the lithiation/delithiation process leads to the pulverization of the electrode materials which will reduce capacity during the cycling process and thus causes series of challenges to the practical use of Si as an anode material. Therefore, various silicon nanostructures have been designed to tackle the aforementioned problems associated with Si anode

9-16

.

Recently, the synthesis of Si/SiOx/C structure composites are investigated as an effective method to improve the electrical performance

17-21

. SiOx (0 < x ≤ 2) shells not only alleviate the formation of SEI

film and the volume expansion 22, but also can strengthen the growth of the carbon layer on Si particles

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and provide a good electron-transfer network for low-conductivity Si

23-24

. Therefore, developing the

Si@SiO2/C composite is an effective method to improve the electrochemical performance of Si-based materials. Nevertheless, the common carbon resources of Si/C composite are mainly from phenolic resin 15, 25-26

, PVA14, 27, toluene 28-29, etc. These carbon precursors are generally from non-renewable petroleum

resource. Moreover, in order to obtained the SiO2 layer on the surface of the Si nanoparticles, raw Si material is often treated through additional calcination in air

24

or alkali treatment 19, which hinder their

practical applications. Therefore, it is still a challenge to develop a simple, green and cost-effective method for preparing Si/SiO2/C anode materials by using renewable raw as carbon resource. Cellulose is one of the most abundant biopolymer on the earth, which is sustainable, biodegradable and inexpensive resources

30-31

. Compared with petrochemical-derived polymers, cellulose is a kind of

renewable, environment-friendly, and cost competitive nature material. Therefore, we design a novel and sustainable route for rational synthesis of Si@SiO2/C composite anode by using cellulose as carbon precursor. First, Si/cellulose precursor was prepared by dispersing nano-Si in homogeneous solution of cellulose/NaOH/urea solvent and then coagulated and regenerated in acid solution. Owing to the existing of NaOH in the precursor solution, the SiO2 layer can be generated through the reaction between Si, H2O and NaOH in the blending process. Therefore, Si@SiO2/C composite was obtained after carbonization of the Si/cellulose precursor. In this work, NaOH performs a dual function. First, it serves as an alkaline resource for the dissolution of the cellulose. Secondly, it can react with Si to generate SiO2 layer coating on the surface of the Si nanoparticles. Furthermore, the resulting Si@SiO2/C composite shows a high reversible capacity due to the reliable buffering of SiO2 layer and mechanical support of carbon scaffold during cycles.

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2 Experimental section 2.1 Preparation of Si@SiO2/C composite The cellulose solution was prepared by a typical method with a slight modification 32. In brief, 7.0 g NaOH, 11.0 g urea and 2.0 g microcrystalline cellulose were dissolved in 80 mL deionized water and kept at -12 °C for 5 h to obtain 2 wt% transparent cellulose solution. Commercial Si nanopowder (Si NPs) (~100 nm, Naiou, China) and cellulose solution were used as starting materials. Fig. 1 illustrates the preparation process for Si@SiO2/C using cellulose as carbon resource. Typically, 0.5 g silicon powder was dispersed in 50 g cellulose solution (2 wt% cellulose) with a vigorous agitation at -12 °C for 5 min. Subsequently, the obtained sample was added into the 5 wt.% HCl for the coagulation and regeneration of the cellulose to obtain cellulose/Si composites. The resulting composites were washed with deionized water and then freeze-dried. After that, the as-obtained Si/cellulose composite was carbonized at 700 °C for 3 h with a heating rate of 5 °C min-1 in an argon atmosphere to produce the Si@SiO2/C composite.

Fig. 1 Schematic illustration of the synthesis of Si@SiO2/C composite.

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2.2 Characterization The X-ray powder diffraction (XRD) patterns were recorded on an Ultimate IV (Rigaku, Japan) spectrometer. FT-IR spectra were recorded on a Nicolet iS5 spectrometer (Thermo fisher, USA). Raman spectra were measured on a DXR (American Thermo Electron) Raman spectrometer with a He-Ne laser (633 nm) as the light source. Thermal analysis curves were obtained by using a STA449F5 TGA analyzer from 40 °C to 700 °C with a heating rate of 10 °C min-1 under air atmosphere. The microstructures were studied by SEM (Zeiss Sigma) and TEM (Zeiss Libra200FE). XPS experiments were performed with an EscaLab 250Xi system (Thermo fisher, USA). The carbon content of the samples was determined by element analysis (Element Vario EL III). The surface area of the samples were measured on a Brunauer-Emmett-Teller (BET) surface area analyzer (Micrometric, ASAP 2010) in a static measurement mode. 2.3 Electrochemical measurements The working electrodes were comprised of active material, acetylene black and sodium alginate at the weight ratio of 8:1:1. A solution of 1 M LiPF6 in a 1:1 (v/v) of ethylene carbonate (EC) and dimethylcarbonate (DMC) containing 10 wt% fluoroethylene carbonate (FEC) was used as the electrolyte. The obtained slurry was coated onto pure copper foil and dried in a vacuum oven at 80 °C for 12 h. The mass loading of active material (Si@SiO2/C) on the electrodes was about 0.5 mg cm-2. The anode was assembled in an Ar-filled glove box. The microporous Celgard 2400 membrane and Li foil were used as the separator and the counter electrode, respectively. In the galvanostatic battery tests, one pre-cycle was performed at 100 mA g-1 to stabilize the SEI layers for all the electrodes, and subsequent cycles were performed at different current densities in the voltage range between 0.01 and 1.5 V versus Li / Li+ using

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a battery tester (LAND CT2001A, Wuhan, China) at room temperature. The cyclic voltammograms (CV) and electrochemical impendence spectroscopy (EIS) were conducted on CHI 660 (Chenhua, Shanghai, China) electrochemical workstation. The electrochemical impedance spectra of the cells measured at a frequency range of between 0.01 and 10000 Hz and an amplitude ratio of 5 mV. The current density and specific capacity were calculated based on the mass of Si@SiO2/C composite. 3 Results and discussion The morphology and structure of the products have been investigated by scanning electron microscopy (SEM) and transmission electron microscopy (TEM). As shown in Fig. 2 a and c, the individual Si NPs with a diameter of about 100 nm were encapsulated and connected by the cellulose network. No conglomeration of Si NPs was found on the surface of the material, indicating that most nanosilicons have been uniformly distributed in the cellulose matrix. Similar morphology is also observed in Si@SiO2/C composite (Fig 2 b and d), which indicates that the structure is preserved well during the carbonization process. As shown in Fig. 2e, Nano-Si shows the lattice spacing value (0.31 nm) which correspond to the (311) plane of Si. The amorphous SiO2 layer (~4 nm) is coated on the surface of crystalline Si NPs. Fig. 2f shows the TEM-energy dispersive spectrometer (TEM-EDS) line-scanning of the composite, which further intuitively confirms the SiO2.layer on the Si surface.

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a

c

b

Si NPs

d

Carbon

Cellulose

Si NPs

Fig. 2 SEM image of (a) Si/cellulose composite and (b) Si@SiO2/C. TEM image of (c) Si/cellulose composite and (d) Si@SiO2/C, (e) HRTEM of Si@SiO2/C, and (f) the corresponding TEM-EDS line-scanning.

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Fig. 3. (a) XRD pattern, (b) FTIR spectra, (c) Raman spectrum, (d) TGA curve of Si@SiO2/C The XRD patterns obtained from Si NPs, Si/Celluose and Si@SiO2/C materials are depicted in Fig. 3 a. The peaks of all patterns show the sharp reflection at 29°, 47°, 56°, 69°, and 77°, which are ascribed to the (111), (220), (311), (400) and (331) lattice planes of Si. These peaks are the typical reflections of crystalline face-centered cubic Si (JCPDS Card No.27-1402) 33. The result suggests that Si NPs are well retained in the composites. In the pattern of Si/cellulose composites, the diffraction peaks at around 20° and 22° are assigned to the (110) and (020) lattice planes of the crystalline cellulose, respectively

34

,

indicating that the cellulose has been generated and encapsulated Si NPs in the prepared process. For the pattern of Si@SiO2/C, the characteristic peaks of the cellulose vanished and a broad (002) peak of amorphous carbon at around 23° were observed 35, indicating that cellulose was carbonized after heating to 700 °C in an argon atmosphere.

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In order to understand the chemical formation evolution of the Si@SiO2/C during the synthesis process, fourier transform infrared spectroscopy (FTIR) characterization was performed to investigate the chemical structure of the Si, Cellulose, Si/cellulose and Si@SiO2/C, respectively. As shown in Fig. 3 b, the spectra of Si NPs showing a broad band at approximate 3500 cm-1 can be assigned to the symmetric stretching of the surface hydroxyl group and absorbed water. The pristine cellulose shows type adsorptions of C–O–C stretching within the pyranose ring skeletal (at 1050 cm-1), O–H stretching (3400 cm-1), and C–H stretching (2900 cm-1) 36. As expected, these characteristic absorption peaks also appear in the Si/cellulose composite which suggests the Si/cellulose composite has been successfully synthesized by using our method. This result is also in good agreement with the XRD result. For the spectra of the Si@SiO2/C, the broad bands at around 1080 cm−1, 810 cm−1 and 470 cm−1 can be assigned to asymmetric stretching vibrations, symmetric stretching vibrations and bending vibrations of Si-O-Si bonds, respectively

37

. These characteristic peaks are also found in the Si/Cellulose composites. The results

indicate that amorphous SiO2 has been detected in the sample. Based on the above results, the formation mechanism of Si@SiO2/C nanocomposite can be ascribed to the reaction of the Si, NaOH and H2O during the synthesis process. The formation process of SiO2 layer can be interpreted by the following equation 38. Si+2NaOH+H2O→Na2SiO3+2H2↑ (1) Na2SiO3+2HCl→SiO2+2NaCl+H2O (2) When Si NPs were added into the cellulose solution with a vigorous agitation, the NaOH existing in the cellulose solution will react with a fraction of Si to produce Na2SiO3 and release hydrogen. Subsequently, the product of Na2SiO3 will react with the hydrochloric acid to produce amorphous SiO2 during the regeneration process of mixture solution of cellulose/Si. Thus, Si@SiO2/C composite was

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obtained after calcining cellulose/Si precursor in an argon atmosphere. In this work, NaOH performs a dual function during the synthesis process. First, it serves as an alkaline for the dissolution of the cellulose 39

. Secondly, it can react with Si NPs to generate SiO2 layer coating on the surface of the Si NPs in the

process of preparation. Fig. 3 c shows the Raman spectra of the pure Si nanoparticles and the Si@SiO2/C composite. A strong peak at 510 cm

-1

and two broad peaks around 300 cm-1 and 935 cm-1 are ascribed to the

characteristic peaks of Si NPs

9, 33

. The result indicates that Si NPs are well retained in the composites.

Two wide peaks at around 1360 cm-1 and 1590 cm-1 are assigned to the D band originated from disordered carbon and G band originated from sp2 graphitic carbon, respectively intensity ratio of the peaks ID/IG, is used as the index of the graphitization degree

40-41

42-43

. The relative

. The lower value

correlates with a higher degree of graphitization and superior electrical conductivity. The intensity ratio of D-band and G-band (ID/IG) of sample is 0.78, indicating that the carbon in the composite shows the relatively high graphitization degree, which can provide superior electrical conductivity carbon matrix. Furthermore, the presence of the SiO2 layer on the Si NPs surface will promote the growth of carbon upon heating 23, 44, and thus strengthen a good electron-transfer network for low-conductivity Si. Fig. 3 d shows the TGA curve of Si@SiO2/C composite anode measured from 40 to 700 °C at a heating rate of 10 °C min-1 under air atmosphere. In the TGA curve, the burning temperature of the sample was 569 °C and the sample showed a weight loss of 22 wt% at 700 °C. The weight loss can be mainly attributed to the oxidation of amorphous carbon in air. In comparison, the carbon content from the element analysis is 24%, a slightly higher than the result of TGA. It can be ascribed to the oxidation of Si NPs, which contributes to a weight gain of the sample 45.

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Fig. 4. XPS spectra of the samples. (a) Survey; (b) O 1s; (c) C1s deconvolution and (d) Si 2p deconvolution. In order to get additional information of the surface chemical composition and chemical states of the samples, the Si@SiO2/C composite was investigated by XPS. As shown in the Fig 4, the sample is composed of O, Si and C. The binding energy of 533.1 eV is the signal of O 1s 35. The chemical states of C and Si have been distinguished from the positions of the high-resolution by the curve-fitting procedure. The high resolution of C1s spectrum can be fitted to three peaks. The C1s peak located at 284.8 eV is assigned to the free carbon in the sample

45

. The two weak peaks located at 286.2 and 288.8 eV are

attributed to C-O and C=O bonds, respectively

46

. For the high-resolution spectrum of Si2p, the peak

binding at 103.5 eV is assigned to Si (4+) corresponding to SiO2 and the peak binding at around 99.5 eV is assigned to Si(0). The two band peaks between 103.5 and 99.5 eV can be ascribed to SiOx (0 < x < 2) 33, 47

. The presence of Si (4+) indicates the successful encapsulation of SiO2 on the Si NPs, which is well

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consisted with the results of TEM and FTIR. The specific surface areas of samples were measured by a Brunauer-Emmett-Teller (BET) surface area analyzer. As shown in Tab. S1, the nano-Si exhibits relatively low specific surface area of 42.6 m2 g-1 and pore volume of 0.12 cm3 g-1, while Si@SiO2/C shows higher specific surface area of 74.8 m2 g-1 and pore volume of 0.18 cm3 g-1. The reason can be ascribed to the carbon matrix, which provides high specific surface area. The relatively high specific surface area of Si@SiO2/C composite can facilitate the infiltration of electrolyte into electrode and promote the charge transfer process for the composite electrode 48.

Fig. 5 Electrochemical measurement of pure Si and Si@SiO2/C anode: (a) the cycling performance of the pure Si anode and Si@SiO2/C anode at 420 mA g-1; (b) the rate performance of the Si@SiO2/C composite

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at various current densities. (c) CV curves of Si@SiO2/C anode; (d) Nyquist plots of Si@SiO2/C after the 5th, 15th and 100th charging-discharging cycles. The

long-term

cycling

stability

of

Si@SiO2/C

was

evaluated

using

galvanostatic

charging/discharging in the voltage interval of 0.05−1.5 V at a constant current density of 100 mA g−1 for the first cycle and 420 mA g−1 for the following cycles. For comparison, we also tested LIBs by utilizing pure Si as anode. As shown in Fig. 5 a, in the first cycle, the Si@SiO2/C composite exhibits an initial discharge/charge capacity of 2506/1525 mA h g−1. The initial coulombic efficiency (ICE) is 61% and the average CE from 2nd to the 200th is up to 98.4%. The irreversible capacity of the initial cycle can be ascribed to reduction of the electrolyte, resulting in the formation of a solid electrolyte interphase (SEI) on the surface of the active particles and/or from irreversible lithium insertion into nanocomposites. 18, 49. Furthermore, the composite electrode can still exhibit a relatively high reversible capacity of 1071 mA h g−1 after 200 cycles at a current density of 420 mA h g−1. In contrast, pure Si anode exhibits a very high initial discharge/charge capacity of 2777/2188 mA h g−1 and high ICE of 78 % in the first cycle, but its capacity decreased rapidly to 390 mA h g-1 after 140 cycles. The significant improvement of the lithium-storage properties for Si@SiO2/C nanocomposite electrode is mainly attributed to the SiO2 layer and the carbon network which provided reliable buffer layer and good mechanical support for the Si NPs without breaking during charging/discharging. Therefore, Si@SiO2/C composite preventing the continuous formation of extra SEI responsible for active material loss and capacity fading, and thus, the cycle performance of the Si@SiO2/C anode is improved. Moreover, The cycle performance of Si@SiO2/C obtained by this method is also comparable to those of recently reported Si/SiOx/C (0 < x ≤ 2) anode materials (Tab. 1) , which has shown a competitive performance.

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Tab. 1 Comparison of electrochemical performance for various Si-based anode materials Materials

Cycling performance Current density

References

Reversible specific capacity

−1

100 mA g

786 mAh g−1 after 100 cycles

24

400 mA g−1

830 mA h g-1 after 100 cycles

27

1000 mA g−1


29

500 mA g−1


50

Double Core-Shell Si@C@SiO2

200 mA g−1


18

Si/SiO2@C composite

50 mA g−1


25

100 mA g−1


7

Si@SiOx@C nanocomposite

150 mA g−1


49

Si@SiO2/C composite

420 mA g−1


Double-walled core-shell structured Si@SiO2@C nanocomposite Three-dimensional network structure Si/C anode Si-C nanocomposites with yolk-shell structure Silicon/nitrogen-dopedcarbon/carbon nanotubespheres

Colloidal

routes

to

synthesize

silicon@carbon composites This work

The rate capability was evaluated by conducting the charge/discharge test at various current densities from 420 to 8400 mA g−1. As shown in Fig. 5 b, after the first activation cycles at 100 mA g-1, the current density gradually increases to 420, 840, 2100, 4200 and 8400 mA g-1. Accordingly, the discharge capacity of Si@SiO2/C changed from 1472 to 1283, 1069, 850 and 586 mA h g-1. Moreover, through multiple rate cycles, the reversible capacity is able to return to 1390 mA h g−1 when current density returned back to 420mAg-1. This result suggests that the Si@SiO2/C composite remain stable even under extremely high current density. The reason can be ascribed to the SiO2 layer and carbon matrix coated on the Si NPs surface, which contribute to a good cushion and significantly enhanced electrical conductivity at high current density. Furthermore the carbon matrix can not only absorb Li-ions but also act as the fast transport channel for electrons and Li-ions

21, 51

, which leads to the good electrical

conductivity and Li-ion diffusion during the discharge/charge cycles. Fig. 5 c shows the CV curves of Si@SiO2/C composite electrode at a scan rate of 1 mV s-1 between

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0.01 and 3 V. The CV curve of the first cycle is apparently different from the subsequent cycles. In the first cathodic sweep, there are obvious reduction peaks at about 1.2V, which can be attributed to the decomposition of electrolyte on the composite surface and the irreversible formation of solid electrolyte interphase (SEI) film on the electrode surface 25, 52. The reduction peak at about 0.33 V can be ascribed to the reduction of SiOx and the formation of elemental silicon and a series of silicates 53. After first sweep, these reduction peaks disappear in the subsequent curves indicates the formation of stable SEI

10

. In the

rest cathodic scan, when the potential gets lower to 0.1 V, a sharp reduction peak is observed, which corresponds to the lithiation of Si NPs to form a LixSi alloy

54-55

. The charging branch shows two

oxidation peaks at 0.37 and 0.55 V, which are related to the extraction of ions from LixSi. Obviously, the two oxidation peaks exhibit an enhanced intensity during anodic sweep, which can be attributed mainly to the gradual activation process of anode 45. For comparison, the pure Si anode was also tested. As shown in Fig. S1, the similarly peaks were also appeared but absence of the reduction peak at 0.33 V. This can be ascribed to the absence of SiO2 structure in pure Si anode, which is in good agreement with the FTIR results. To investigate the effect of SiO2/C structure on electrical resistance of electrodes, The electrochemical impedance spectra of pure Si and Si@SiO2/C anode at specified cycles were measured at a frequency range of between 0.01 and 10000 Hz and an amplitude ratio of 5 mV. As shown in Fig. 5d, all the curves consist of a depressed semicircle in the high frequency range and a slopping line in the low frequency range. The diameter of the depressed semicircle corresponds to the charge transfer resistance at the electrode/electrolyte interface, while the sloping line reflects the diffusion resistance of ions into the active material

56

. Compared with the semi circle of the electrode before cycling of Si@SiO2/C anode

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(Fig. 5d), a decreasing trend is observed for the subsequent curves. This reduction of the resistance can be ascribed to the activation of the composites electrodes as well as the stable structure of SiO2 layer and carbon scaffold provided a good buffering and electronic/ionic conductivity during the cycles 35. The pure Si anode was also tested as a control. As shown in Fig. S2, the impedance of the pure Si electrode significantly increased after 100 cycles. The result is mainly from the electrode pulverization after cycling caused by the large volume changes of Si during cycles process

35

. Therefore, the rational structure of

Si@SiO2/C composite is beneficial to accommodate the volume change and retaining the integral morphology during the cycles. Moreover, the @SiO2/C structure contributes to a sufficient contact between the activated materials, carbon matrix and the contact can be enhanced during the cycles, and thus the electrical resistance of electrodes was reduced during the cycles. To further investigate the morphology changes of Si@SiO2/C composite after cycling, TEM characterization was carried out on delithiated composite electrode. The cycled electrode was disassembled and washed by acetonitrile for three times before observation. As shown in Fig. 6(a), the nanostructure of composite still maintains integrity and shows no pulverization after cycles. Meanwhile, the lattice of crystalline Si is not detected in HRTEM image (Fig. 6 b). The reason can be ascribed to the conversion of crystalline Si to amorphous Si after lithiation/delithiation process. Moreover, to investigate the distribution of Si in the composite, the recycled sample was characterized by TEM-EDS. As shown in Fig 6(c), Si element is mainly distributed in the center of composite. The result further indicates that Si is well encapsulated in the composite after cycles. The above results confirm the stability of composite, which significantly ensures the good cycling stability.

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Fig. 6 (a) TEM, (b) HRTEM and (c) TEM-EDS line scanning images of Si@SiO2/C composite after 200 cycles 4 Conclusions In summary, Si@SiO2/C composite for Li-ion battery anode was successfully prepared using the renewable and low-cost cellulose as carbon resource by a novel method for the first time. In this unique architecture, SiO2 layer and carbon scaffold synergistically provide a good buffering for the volume change and a highly conductive network for electrons. The resulting Si@SiO2/C composite exhibits excellent electrochemical performance in LIBs, The specific capacity reached 1071 mA h g−1 at the current density of 420 mA g−1 after 200 cycles. More importantly, the renewable cellulose utilization and the rational and integrated preparation method make the Si@SiO2/C composite cost-effective. Therefore, the composite synthesized with presenting method have good prospects for the industrial application in

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high-energy lithium-ion batteries. ASSOCIATED CONTENT Supporting Information. Supporting information associated with this article is available free of charge on the ACS Publications website at DOI: Additional figures about the CV curves and the impedance curve for the pure Si electrode. AUTHOR INFORMATION Corresponding Author* [email protected] [email protected] [email protected] Notes The authors declare no competing financial interest. ACKNOWLEDGMENTS This work was supported by the President Fund of Xiamen University (No. 20720150090), the research fund from the Xiamen Southern Oceanographic Center (No. 4GZP59HJ29), Fujian Provincial Department of Ocean and Fisheries (No. 2015-27). References 1.

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127x81mm (300 x 300 DPI)


C Composites Using Sustainable

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