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Tunable Synthesis of Yolk-shell Porous-silicon@Carbon for Optimizing Si/C-based Anode of Lithium-Ion Batteries Sichang Guo, Xiang Hu, Yang Hou, and Zhenhai Wen ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.7b13035 • Publication Date (Web): 09 Nov 2017 Downloaded from http://pubs.acs.org on November 10, 2017

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ACS Applied Materials & Interfaces

Tunable Synthesis of Yolk-shell Porous-silicon@Carbon for Optimizing Si/C-based Anode of Lithium-Ion Batteries Sichang Guoa, Xiang Hua, Yang Houb* and Zhenhai Wena* a Key Laboratory of Design and Assembly of Functional Nanostructures, Fujian Provincial Key Laboratory of Nanomaterials, Fujian Institute of Research on the Structure of Matter, Chinese Academy of Sciences, Fuzhou 350002, China b College of Chemical and Biological Engineering, Zhejiang University, Hangzhou 310027, China E-mail: [email protected], [email protected]

ABSTRACT Significant “breathing effect” calls for exploring efficient strategies to address the intrinsic issues of silicon anode of lithium ion batteries (LIBs). We here report a controllable synthetic route to fabricate the silicon-carbon hybrids, in which porous silicon nanoparticles (p-SiNPs) are loaded in void carbon spheres with forming the yolk-shell p-SiNPs@HC nanohybrids tunable. A set of controlled experiments accompanying with systematic characterizations demonstrate that the void space and mass loading of Si can be adjusted in an effective way, so that the nanostructure can be optimized with achieving improved electrochemical performance as anode of Lithium ion batteries (LIBs). The optimized p-SiNPs@HC nanohybrids show excellent performance as anode for Li-ion battery, delivering a capacity of more than 1,400 mA h g

−1

after 100 cycles at 0.2 A g−1 and 720 mA h g−1 at a high current

density of 4 A g-1. The present work may provide us with an attractive and promising strategy for advancing Si-based anode materials due to advantages of tunable structure of silicon-carbon nanohybrids for optimizing electrochemical performance.

Keywords: Silicon anode, yolk-shell, tunable structure, magnesiothermic reduction, lithium ion batteries

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Introduction Rechargeable lithium ion batteries (LIBs) have evoked a plethora of research to meet the ever-growing demands of portable electronics and energy storages thanks to their high energy density, long cycle life, low self-discharge, and high operating voltage.1-3 Although the initial commercialization of LIBs has been developed for more than two decades, exploring new electrode materials with high specific capacity are still highly desirable to satisfy the widespread energy storage applications.4 Among the various anode candidates, silicon (Si) based materials have aroused great interests owing to the advantages of high capacity, low charge potential, and abundant natural resources.5 Si can deliver a theoretical capacity with more than decuple increase compared to the current commercialized carbon anodes, i.e., 4,200 mA h g−1 for the highest lithiated stoichiometry of Li22Si5 vs. 372 mA h g−1 for the graphite.6-7 However, there exist at least three intrinsic issues at Si-based anode hindering their practical application: 1) Low intrinsic electrical conductivity limits the transmission of Li-ion through the Si material;8 2) Li-ion does not intercalate into crystal Si topotactically but proceed via solid-state alloying reaction, which cause a drastic volume change (≈300%) with severely cracking down the electrode,9-10 leading to increasing battery impedance and decreased specific capacity due to the electrical contact loss between the active material and current collector and continuous growth of unstable solid electrolyte interphase (SEI); 3) There still remains daunting challenges in implementing tunable synthesis of favorable Si-based anode with achieving large-scale production.11-12 So far, immense efforts and numerous strategies, including nanostructured Si (nanoparticles, nanotubes, and nanowires), stress-relief buffer matrix, and physical compartment, have been developed to mitigate or tackle these critical issues aforementioned.13-15 Extensive studies have shown that Si-anode materials with particle size below 150 nm are beneficial to prevent fast mechanical fracture while offering high specific surface area, fast electron transfer and ionic diffusion.16 Accordingly, various synthetic methods, such as chemical vapor deposition (CVD),

2


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sol-gel, chemical-assisted thermal techniques, electrospin, and template synthesis have been applied to produce a variety of Si-based nanomaterials.17-21 Moreover, buffer matrix and physical compartment can prevent the Si from reuniting and withstand stresses upon volume changing, thus keeping the structural integrity. Among them, Si-carbon composites are highly attractive because of their outstanding electric conductivity and mechanical-strength. Recently, a large number of Si-carbon hybrids have been developed for improving the long cycling stability of Si-based anodes.22-25 Porous silicon (p-Si) nanomaterials with the advantages of effectively releasing the stress that are created by the huge volume expansion of Si also trigger a lot of research. In industry, p-Si is normally produced by etching bulk Si, chemical de-alloying, and template-assisted synthesis methods. However, these synthetic approaches need a significant amount of toxic chemicals, high temperature, high pressure

or

expensive

precursors.26-27

Compared

with

these

approaches,

magnesiothermic reduction (MR) technology with low-energy consumption can convert silica directly into Si while basically maintaining the structure and shape. Recent studies about the synthesis of porous Si through the MR method have been reported.28-30 Nevertheless, we must strive to develop reliable synthesis method to realize low-cost and scalable preparation of Si-based anode with goal to achieve a satisfied electrochemical performance for putting forward the commercialization of Si-based anode. In this work, we describe a deliberate design for the tunable synthesis of porous Si nanoparticles (p-SiNPs) loaded in hollow carbon nanosphere (HC) hybrids, as denoted as p-SiNPs@HC. As shown in Fig. 1a, the p-SiNPs@HC hybrid demonstrates the structure advantages coping with the volume variation during lithiation/delithiation process. Fig. 1b illustrates the entire fabrication process of a series of Si-based materials. Specifically, the preparation starts from synthesis of SiO2 nanospheres using a minor modified Stöber’s route associated with coating a layer of resorcinol-formaldehyde (RF) polymer on SiO2 nanospheres, and subsequently carbonization treatment.31 After the tunable water etching of SiO2, the resulting SiO2 was further evolved into p-Si that consists of a large number of SiNPs during the MR. 3



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The inner SiO2 nanoparticles in SiO2@C were controllably etched via hydrothermal method with producing SiO2@void@C, during which SiO2 will transformed into Si(OH)4 that are soluble under high temperature and high pressure.32 In this way, the void size between yolk and shell could be efficiently tuned through adjusting etching reaction time, temperature, and solution concentration. We found that the etching time is a key factor to impact the morphology of SiO2@void@C: The etching dissolution of SiO2 is inefficient in the beginning of etching process, and looks like a distinct void can be formed with a hydrothermal time of 15 hours. Six samples, i.e., SiO2@C, SiO2@C-0, SiO2@C-1, SiO2@C-2, SiO2@C-3, and SiO2@C-4, were prepared by fixing the hydrothermal (180 ºC) etching time of 0, 15, 20, 25, 30, and 35 h, respectively. The hybrids of SiO2@C-0 has the similar morphology of SiO2@C (Fig. 5a,b and Fig. S7a-c), providing not enough space for volume variation of Si, while the SiO2@C-3 and SiO2@C-4 have the too large void space (Fig. S7d-i), indicating most of SiO2 were removed, making against the capacity of hybrids. Therefore, SiO2@C, SiO2@C-1, and SiO2@C-2, were chosen as the objects for study. Thermogravimetric analysis (TGA) was used to investigate the mass variation of the samples (Fig. S1) and to illustrate the influence of etching time. Since the MR process is a well-known exothermic reaction (Eq. 1), the local temperature in the core area is usually higher than the setting temperature, resulting in the generation of SiC (Eq. 2).33 Given that SiC is a semiconductor with poor electrical conductivity for potentially impairing the performance of Si-anode materials, NaCl was used as heat scavenger to digest the excessive heat to impair the formation of SiC. X-ray powder diffraction (XRD) patterns show that SiC is the major product without assistance of NaCl, while a rather pure Si phase can be produced upon adding NaCl (Fig. S2b-2c). This result suggests that the presence of NaCl does assist to restrain the formation of SiC during the MR step. It is worth mentioning that the gaseous magnesium reacts with silica in the MR process,34 thus the experiments were developed in sealed quartz tubes (Fig. S3). 2Mg(g) + SiO2 (s) → 2MgO(s) + Si(s) Si(s) + C(s) → SiC(s) 4


(1) (2)

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Results and discussion XRD was carried out to further investigate the composition and crystalline structure of the MR products. As expected, the sharp peaks of all samples can be indexed to Si with a cubic structure (JCPDS, no.27-1402), suggesting that SiO2 has been successfully converted into crystalline Si through the MR process (Fig. 2). Five peaks located at 28.4, 47.2, 56.1, 69.1, and 76.3° are attributed to the crystal planes of (111), (220), (311), (400), and (331) respectively. Besides, the weak broad peak around at 23.0° belongs to the infinitesimal silica or amorphous carbon, which are beneficial to restrict the stress variation during lithiation/delithiation process. It can be clearly seen that the Si peaks of p-SiNPs@HC (Fig. 2a), p-SiNPs@HC-1 (Fig. 2b), and p-SiNPs@HC-2 (Fig. 2c) are gradually weakened with increasing time of etching, indicating that Si content in hybrids decreased associated with the sacrifice of SiO2. The result was further supported by TGA. As presented in Fig. 3a, the samples have a large mass loss between 350 and 650 °C, corresponding to the loss of carbon in the samples. The gradually increased mass at above 650 °C is attributed to the oxidation of Si because it is stable below 600 °C. The mass ratio of Si in the p-SiNPs@HC, p-SiNPs@HC-1, and p-SiNPs@HC-2 are calculated to be 78%, 65%, and 48%, respectively. The digital photos of all samples before and after TGA measurements were displayed in Fig S4, in which colour change from black to earthy yellow is due to the oxidation reaction of products at high temperature. Raman spectroscopy of the samples is shown in Fig. 3b. The peak at about 497 cm−1 corresponding to the Si gradually weaken from p-SiNPs@HC, p-SiNPs@HC-1, to p-SiNPs@HC-2, further implying the decrease of mass loading of Si in the composites. Compared with the pure SiNP (500 cm-1), there is a blue shift for all samples, which may be attributed to the phonon confinement effect or masking effect of carbon coating on p-Si.35 Two characteristic peaks at 1,344 and 1,592 cm-1 for all composites match well with the D-band (disordered band) and G-band (graphite band) of graphite, respectively. X-ray photoelectron spectroscopy (XPS) measurements were applied to analyze the surface elemental composition and valence states of the samples. Fig. 3c shows the survey spectra of p-SiNPs@HC-1, in which several 5



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distinct peaks at around 100.5, 152.5, 283.8, and 533.5 eV, correspond with Si2p, Si2s, C1s, and O1s, respectively.36 The O1s peak may be ascribed to tape substrate, adsorbed oxygen, unreacted silica, and carbon-shell.37 The high-resolution XPS spectrum of the Si2p in Fig. 3d can be deconvoluted into three positions: 99.9 eV for Si-Si, 101.7 eV for Si-C, and 102.9 eV for Si-O.38-39 XPS measurements were also carried out for the other two samples. The relatively intensity of O1s peak (Fig. S5a) is decreasing from p-SiNPs@HC, p-SiNPs@HC-1, to p-SiNPs@HC-2, indicating that the SiO2 has a close contact with carbon shells will be not transformed into Si easily. The signal of SiC (Fig. S5a) has a similar tendency with O1s peak, indicating the void space have an effect on the formation of SiC during the MR process by prevent the direct contact between carbon and silica. The SEM images for SiO2@RF (Fig. S6) and SiO2@C (Fig. 4a) clearly illustrate the monodispersed nanoparticles with uniform diameter of approximately 180 nm. The SiO2@C-1 (Fig. 4b) and the SiO2@C-2 (Fig. 4c) samples well maintain the original

nanosphere

structure

of

the

SiO2@C,

demonstrating

a

good

mechanical-strength of carbon shells. The SEM images of p-SiNPs@HC (Fig. 4d) and p-SiNPs@HC-1 (Fig. 4e) show that the original sphere morphology are basically maintained, indicating that the MR process can successfully convert silica into Si without structure failure. The appearance of pores on the carbon shells in p-SiNPs@HC-2 (Fig. 4f) is attributed to the large void space, which leads to the unstable structure integrity in the MR procedure. Transmission electron microscopy (TEM) was applied to uncover the structure evolution of the samples during the synthetic process. As shown in Fig. 5 and Fig. S7, the void space in these samples increases with the hydrothermal time. Typical TEM images of SiO2@C-1 (Fig. 5c, d) and SiO2@C-2 (Fig. 5e, f) reveal that the carbon layer has a thickness of ~10 nm. The p-SiNPs@HC-1 (Fig. 6a) displays a similar morphology with SiO2@C-1 (Fig. 5c), indicating that there was almost no change in morphology before and after the MR, while a close observation on SiNPs implies the appearance of nanopores within them. The high resolution TEM (HRTEM) image (Fig. 6b) shows that the pores within p-Si have a diameter about 6.9 nm with a distinct 6


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lattice fringe space of 0.31 nm, corresponding to the (111) plane of Si with diameter of less than 7 nm. Fig. 6c displays the corresponding selected area electron diffraction (SAED) pattern. There are three circles in the SAED patterns that agree well with Si (111), (220) and (311) planes. The EDS mapping measurements were also performed to determine the elemental distribution in p-SiNPs@HC-1. Fig. 6d shows the enrichment of disconnected Si signals in the cores and homogeneous distribution of C signals in the shells, further confirming the yolk-shell configuration. Additionally, the O signal was also detected, which may due to the adsorbed oxygen, carbon shells, and the residue SiO2 phase. EDS tests were also applied to p-SiNPs@HC (Fig. S8) and p-SiNPs@HC-2 (Fig. S9). One can observe that only p-SiNPs@HC has the signal of Si on the carbon shells, which may be due to the formation of SiC or unreacted SiO2. Fig.7a shows the first four cyclic voltammetry (CV) curves of p-SiNPs@HC-1 at a scan rate of 0.1 mV s−1 between 0.01 and 1.5 V (vs. Li/Li+), which exhibits typical electrochemical characteristics of Si-based materials. In the cathodic scan, there is a sharp reductive peak below 0.2 V, corresponding to the transformation of Si from crystalline phase to amorphous phase. This phenomenon has been described as electrochemically-driven solid-state amorphization.40 In the anodic scan, two oxidation peaks are observed at around 0.36 V and 0.58 V in the first scan and retained nearly the same positions in the following scans. Moreover, the current increase with cycling, implying more Si-substance can be further activated upon continuous scan. This “activation” phenomenon should be mainly attributed to gradual breakdown of cubic Si structure, which depends greatly on the diffusion rate of Li+ into p-SiNPs and the formation rate of amorphous Si–Li phase.41-43 Fig. 7b shows the first discharge-charge profiles of p-SiNPs@HC, p-SiNPs@HC-1, and p-SiNPs@HC-2 electrodes at 100 mA g-1. The p-SiNPs@HC-2 displays the first discharge and charge capacities of 2,657 and 1,029 mA h g-1, respectively, corresponding to an initial coulombic efficiency (CE) of 38.7%. The p-SiNPs@HC shows a higher CE of 49.3 % with discharge and charge capacities of 3,809 and 1,876 mA h g-1, respectively. As expected, the p-SiNPs@HC-1 attained the highest CE of 55.0 % with 3,143 and 1,730 mA h g-1 in the discharge and charge process. Regardless 7



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of the lower initial coulombic efficiency of p-SiNPs@HC-1, prelithiation is an effective way for these hybrids in real batteries. It is well known that the low initial CE of anode material is mainly due to the formation of SEI layers, as well as the consumption of Li ions and electrolyte. The p-SiNPs@HC-2 has the lowest CE possible caused by the highest surface area (Fig. 8c), which potentially leads to consumption of irreversible Li+ and electrolyte due to side reaction. The p-SiNPs@HC-1 has the highest CE in all electrodes, which may be attributed to the unique structure with appropriate void space. The carbon shells favor the formation of a stable SEI because the void space can alleviate the volume change issues of Si anodes, preventing the breakdown of structure. Due to the lack of sufficient void space, the structure of p-SiNPs@HC will be destroyed during the discharge process, resulting in the low CE. Cycling performance of all samples were performed at 100 mA g-1 in the initial three cycles and then for later cycles at 200 mA g-1 (Fig. 7c). The p-SiNPs@HC shows an initial delithiation capacity of 1,876 mA h g-1, and drops dramatically in the later cycle. In contrast, the reversible capacity of p-SiNPs@HC-1 still reaches 1,436 mA h g-1 even after 100 cycles, demonstrating excellent long life stability. The CE of the other two samples (Fig. S10) shows the similar phenomenon. The p-SiNPs@HC-2 has a similar cycling curve with p-SiNPs@HC-1 but with a lower specific capacity, because of the low mass loading of Si in the hybrid. Furthermore, the rate performance was further investigated at varied current densities. As displayed in Fig. 7d, the p-SiNPs@HC-1 display a discharge capacities of

1,756, 1,442, 1,271, 1,004,

853, 777, and 720 mA h g-1 at the current densities of 0.1, 0.2, 0.5, 1, 2, 3, and 4 A g-1, respectively, manifesting a good rate capability. When the current density was switched back to 100 mA g-1, the capacity recovers to 1,529 mA h g-1, maintaining around 87 % of its initial capacity. It should be noted that, the capacity is slightly dropped during the later cycling when go back to the low current density. As presented in Fig. 7e, the galvanostatic discharge-charge profiles of p-SiNPs@HC-1 tends to be a horizon line with the increasing current density. Fig. 7f also shows the long cycle stability of p-SiNPs@HC-1 at a current density of 1 A g-1, delivering a 8


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significant cycling stability with a high reversible capacity of 600 mA g-1 after 600 cycles. The excellent cycling stability and outstanding rate capacity of p-SiNPs@HC-1 can be attributed to the yolk-shell structure with appropriate void space, mitigating both mechanical and chemical degradation during discharge/charge process. The TEM images of p-SiNPs@HC-1 (Fig. S11) after the discharge-charge cycling further demonstrates the structural integrity of p-SiNPs@HC-1. The electrochemical impedance spectroscopy (EIS) measurements were tested to better understand the improved electrochemical performance in p-SiNPs@HC-1. Fig. 8a shows the Nyquist plots recorded for the p-SiNPs@HC, p-SiNPs@HC-1, and p-SiNPs@HC-2, respectively. All samples consist of a semicircle in the high frequency region, which represents the interfacial electrode characteristics. The p-SiNPs@HC-1 shows the smallest semicircle, suggesting the lowest charge transfer resistance. According to the previous investigation, the deep lithiation causes the concomitant huge volume expansion, and the later inelastic deformation of Li22Si5 generates nanopores.16,44 The p-SiNPs@HC-1 has a suitable void space to accommodate the volume variation of p-Si without breakdown of yolk-shell structure. The shorter distance between carbon shells and p-Si yolk can guarantee higher efficiency of charge and Li-ion transfer. Fig. S12 shows the Z′-ω−1/2 (ω = 2πf) curves in the low frequency region, and the lowest slope of the p-SiNPs@HC-1 electrode indicates its better lithium ion kinetics in the internal of electrode. For p-SiNPs@HC-2, the large void space hinders the charge transfer in the hybrid, leading to the low electrical conductivity. In terms of p-SiNPs@HC, the volume expansion of p-Si during alloy process destroyed the structure of carbon shells, resulting contact loss between active materials and current collector. Fig. 8b shows the EIS plots for p-SiNPs@HC-1 after 5, 100, and 200 cycles, respectively. Only a slight variation in these impedance spectroscopies confirms that the p-SiNPs@HC-1 can maintain high conductivity during cycling. Fig 8c presents the isotherm N2 adsorption-desorption curve for all three samples, the distinguishable hysteresis loop at high relative pressure

region

demonstrate

the

presence

of

mesoporous

structure.

The

Brunauer-Emmett-Teller (BET) surface area is 237.6, 311.1, 359.7 m2 g-1 for 9



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p-SiNPs@HC, p-SiNPs@HC-1, and p-SiNPs@HC-2, respectively. The average pore size based on the Barrett–Joyner–Halenda (BJH) method (Fig 8d) is 7.25 nm for p-SiNPs@HC, 10.76 nm for p-SiNPs@HC-1, and 12.2 nm p-SiNPs@HC-2, which could be produced during the MR process. These pores can facilitate the transmission of Li+ and electronic, and alleviation of the stress, thus promoting the electrochemical performance. The excellent cycling stability and outstanding rate capacity of p-SiNPs@HC-1 can be attributed to the unique yolk-shell structure in several respects. Firstly, the mesoporous carbon shell can offer high electric conductivity for electron transfer and path for Li ion transport, thus enhancing the rate performance. Secondly, the carbon-shell can prevent the direct contact between Si active materials and the electrolytes, in this way, SEI film will be formed on the surface of carbon rather than on internal p-SiNPs, in favor of maintaining the internal void space and stabilizing the CE performance. Thirdly, the self-supporting and mechanically strong carbon shells with the void space allow the SiNP to expand freely without breaking the structural integrity during the lithiation/delithiation process. Fourthly, the SiNPs produced in the MR process with a large amount nanopores has several advantages with the size less than 150 nm and hoisting the efficiency of active materials, and relieve the volume change; Finally, the raw materials are very cheap without any toxic chemicals and the production is quite simple and straightforward, which make it compatible for scale-up production and commercial process. Conclusion In summary, we have developed a tunable and reliable methodology to fabricate the Si-C nanohybrids with “yolk-shell” structure. The unique void space and nanopores structure are beneficial for the formation of stable SEI film, structure integrity, and increment of efficient channels for fast transport of both electrons and Li ions during the

cycling.

The

resulting

p-SiNPs@HC-1

hybrid

exhibits

an

improved

electrochemical performance compared with that of other Si-based nanohybrids materials. This work may open up new opportunities to meet the ever increased requirements of energy storage applications due to the advantages of non-toxic, 10


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low-cost, and facile scalability of synthetic route.

Supporting Information Description of the materials; Experimental section, digital pictures, EIS patterns, X-ray

photoelectron

spectra,

XRD

patterns,

SEM

and

TEM

images,

thermogravimetric analysis curve, galvanostatic discharge−charge curves, cycling performance and Coulombic efficiency.

Acknowledgments We would like to thank 1000 Plan Professorship for Young Talents in China, Hundred Talents Program of FuJian Province, and the Fujian Science and Technology Key Project (Item Number. 2016H0043) for financial support.

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Figure. 1 (a) Schematic of the lithiation/delithiation reaction in the p-SiNPs@HC anode for LIBs. (b) Schematic illustration of the p-SiNPs@HC design.

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Figure. 2 XRD patterns of (a) p-SiNPs@HC, (b) p-SiNPs@HC-1, and (c) p-SiNPs@HC-2.

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Figure. 3 Structural characterization of the products. (a) TGA and (b) Raman spectra curves of p-SiNPs@HC, p-SiNPs@HC-1, and p-SiNPs@HC-2. (c) XPS spectrum of p-SiNPs@HC-1 and (d) high-resolution XPS spectra of Si2p in p-SiNPs@HC-1.

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Figure. 4 SEM images of (a) SiO2@C, (b) SiO2@C-1, (c) SiO2@C-2, (d) p-SiNPs@HC, (e) p-SiNPs@HC-1, and (f) p-SiNPs@HC-2.

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Figure. 5 TEM images of (a,b) SiO2@C，(c,d) SiO2@C-1, and (e,f) SiO2@C-2.

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Figure. 6 (a) TEM image, (b) HRTEM pattern, (c) SAED and (d) HAADF-STEM mapping images of p-SiNPs@HC-1.

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Figure. 7 (a) CV curve of the p-SiNPs@HC-1 fabric electrode obtained at the first four cycles.

(b) Voltage

profiles

of

p-SiNPs@HC,

p-SiNPs@HC-1,

and

p-SiNPs@HC-2 plotted for the first cycle. All electrodes were first cycled at 0.1 A g-1 for first three cycles and then for later cycles.（c）Delithiation-lithiation capacity performance of all hybrids and coulombic efficiency of p-SiNPs@HC-1. (d) Rate performance at current densities of 0.1, 0.2, 0.5, 1, 2, 3, 4 and 0.1 A g-1, respectively, of all hybrids. (e) Galvanostatic charge-discharge profiles of p-SiNPs@HC-1 at current density of 0.2, 0.5, 1, and 2 A g-1, respectively. (f) Cycle performance of p-SiNPs@HC-1 at a current density of 1 A g-1.

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Figure. 8 (a) Nyquist curves of p-SiNPs@HC (black), p-SiNPs@HC-1 (blue), and p-SiNPs@HC-2 (red). (b) Evolution of Nyquist curves of p-SiNPs@HC-1 during cycling. (c) Nitrogen adsorption-desorption isotherm curves and (d) pore size distribution

plot

of

p-SiNPs@HC

(black),

p-SiNPs@HC-1

p-SiNPs@HC-2 (red).

24


(blue),

and

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