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Facile Synthesis of Si@SiC Composite as an Anode Material for Lithium-Ion Batteries Duc Tung Ngo,†,§ Hang T. T. Le,†,‡,§ Xuan-Manh Pham,† Choong-Nyeon Park,† and Chan-Jin Park*,† †

Department of Materials Science and Engineering, Chonnam National University, 77, Yongbongro, Bukgu, Gwangju 61186, South Korea ‡ School of Chemical Engineering, Hanoi University of Science and Technology, 1 Dai Co Viet, Hai Ba Trung, Hanoi 100000, Vietnam S Supporting Information *

ABSTRACT: Here, we propose a simple method for direct synthesis of a Si@SiC composite derived from a SiO2@C precursor via a Mg thermal reduction method as an anode material for Li-ion batteries. Owing to the extremely high exothermic reaction between SiO2 and Mg, along with the presence of carbon, SiC can be spontaneously produced with the formation of Si. The synthesized Si@SiC was composed of well-mixed SiC and Si nanocrystallites. The SiC content of the Si@SiC was adjusted by tuning the carbon content of the precursor. Among the resultant Si@SiC materials, the [email protected] sample, which was produced from a precursor containing 4.37 wt % of carbon, exhibits excellent electrochemical characteristics, such as a high first discharge capacity of 1642 mAh g−1 and 53.9% capacity retention following 200 cycles at a rate of 0.1C. Even at a high rate of 10C, a high reversible capacity of 454 mAh g−1 was obtained. Surprisingly, at a fixed discharge rate of C/20, the [email protected] electrode delivered a high capacity of 989 mAh g−1 at a charge rate of 20C. In addition, a full cell fabricated by coupling a lithiated [email protected] anode and a LiCoO2 cathode exhibits excellent cyclability over 50 cycles. This outstanding electrochemical performance of [email protected] is attributed to the SiC phase, which acts as a buffer layer that stabilizes the nanostructure of the Si active phase and enhances the electrical conductivity of the electrode. KEYWORDS: Silicon anode, SiC, ultrahigh rate, cyclability, composite

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INTRODUCTION

Among various anode materials, silicon has emerged as a potential candidate to replace graphite owing to its extremely high specific capacity of 3600 mAh g−1.9−16 Theoretically, Si can store ∼10 times more Li than graphite by forming Si−Li alloy. However, Si-based materials have drawbacks such as (i) poor electrical and ionic conductivities, (ii) huge volume expansion (400%) during lithiation, and (iii) formation of an unstable solid−electrolyte interface (SEI) layer, which can lead to severe capacity fading.17−22 Therefore, it is a tremendous challenge to maintain such a high capacity in Si-based anodes during repeated charge−discharge cycles and obtain a cyclability comparable to that of current graphite anodes. To address these issues with Si-based anodes, thus far, various strategies, such as designing a Si nanostructure, embedding Si in a supporting material, and coating Si with stable buffer materials, have been proposed.23−25 In recent years, Si functionalized with carbon to form Si@C composite anodes has attracted much attention because of its simplicity and high efficiency.26−31 Several breakthroughs in the

Lithium-ion batteries (LIBs) have been known as an important component for popularizing personal electronic devices, such as laptops, mobile phones, and digital cameras, because of their unique characteristics, such as high energy and power densities, long cycle life, and low self-discharge rate, which are superior to those of earlier batteries.1,2 Further, the applications of LIBs have been extended to hybrid and electric vehicles (EVs), largescale energy storage systems, medical instruments, and military systems.3−6 However, current LIBs have not been satisfactory for application in EVs, which require an energy density more than 4 times that of conventional LIBs to match the 800 km operation range of current gasoline vehicles.7,8 To meet this requirement, a breakthrough in electrode materials is the foremost prerequisite. Specifically, a conventional LIB is composed of a graphite anode and a lithium transition-metaloxide cathode, whose capacities are limited to the theoretical capacities of 372 and ∼200 mAh g−1, respectively. Obviously, to realize LIBs with remarkably higher energy density, these traditional anode and cathode materials should both be replaced by new, more advanced materials possessing much higher specific capacities. © 2017 American Chemical Society

Received: July 20, 2017 Accepted: September 6, 2017 Published: September 6, 2017 32790

DOI: 10.1021/acsami.7b10658 ACS Appl. Mater. Interfaces 2017, 9, 32790−32800

Research Article

ACS Applied Materials & Interfaces

2, and 3 are denoted as [email protected], SiO2@C-1, SiO2@C-2, and SiO2@C-3, respectively. Next, to prepare Si@SiC with different SiC contents, the [email protected], SiO2@C-1, SiO2@C-2, and SiO2@C-3 samples containing 10 g of SiO2 were well mixed with 8.5 g of Mg powder, transferred into a stainless steel crucible, and reheated to 600 °C in Ar atmosphere for 1 h in the vertical tube furnace at a heating rate of 5 °C min−1. After the chemical reaction was complete, the furnace was cooled naturally to room temperature. The as-prepared samples were then washed with 2 M HCl and 0.1 M HF solutions in sequence to remove undesired products. Finally, the obtained Si@SiC samples were rinsed with ethanol and then dried under vacuum. The samples are denoted as [email protected], Si@SiC-1, Si@SiC-2, and Si@SiC3 after their precursor materials of [email protected], SiO2@C-1, SiO2@C2, and SiO2@C-3, respectively. For comparison, a pure Si powder sample was also synthesized using the same process as for Si@SiC, except that SiO2 was used as the Si source instead of SiO2@C. The synthesis process of Si@SiC is illustrated schematically in Scheme 1.

performance of Si-based anodes have demonstrated the success of this strategy. For instance, Pan et al. reported a facile synthesis of yolk−shell-structured Si−C nanocomposite, in which the yolk−shell-structured Si−C showed superior electrochemical performance to pure Si in all aspects of specific capacity, cyclability, and rate capability.32 Most recently, Li’s group has demonstrated that by coating with carbon the electrochemical performance of both Si@C and TiO2@C was significantly improved compared to either bare Si or TiO2.33 Nevertheless, in these studies, most of the functional carbon existed in a disordered amorphous form that had low electrical conductivity and caused undesirable side reactions with Li during the charge−discharge process. These phenomena resulted in a low rate capability and Coulombic efficiency of the anode for the first several cycles. In this regard, silicon carbide (SiC) outperforms disordered carbon owing to its chemical stability. Unfortunately, however, synthesis of SiC by conventional carbothermal reduction is not facile because the temperature for carbothermal reduction is generally quite high, 1400−2000 °C.34−37 These temperatures are even higher than the melting point of Si, which is the main active material in the anode. Accordingly, studies on the synthesis of nanostructured Si-based materials supported by SiC are rare. On the other hand, SiC can be synthesized at a low temperature of 400 °C by the sputtering method.38−40 However, this method requires expensive and sophisticated equipment systems. Moreover, the obtained product has the form of a thin SiC film, and its production yield is quite small, so it is unsuitable for application in LIBs. Recently, some research groups have reported the synthesis of SiC at a low temperature of 600 °C by magnesium thermal reduction,41−43 wherein a mixture of silica (SiO2), carbon, and magnesium is heated. As the temperature reaches the reaction temperature, the components react with each other to produce SiC and MgO. It is noteworthy that SiC is an inactive material that may serve as a supporting material in improving the stability of Si-based anodes. Thus far, few studies have reported the use of Si@SiC as an anode material for LIBs, and no study has realized mass-scalable synthesis of [email protected] The development of a simple and reliable method for massscalable synthesis of Si@SiC remains a challenging task. In this study, we propose a simple and reliable method for mass-scalable synthesis of Si@SiC, in which fine Si particles are well mixed with ultrafine SiC by a Mg thermal reduction method. In the process, a SiO2@C composite is reduced to Si@ SiC with an optimum SiC content using Mg vapor at 600 °C. Interestingly, the Si@SiC exhibits outstanding electrochemical performance characteristics, such as ultralong cyclability, excellent rate capability, and high reversible capacity. Furthermore, a Si@SiC anode is successfully coupled with a LiCoO2 cathode and employed in an LIB.

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Scheme 1. Schematic of the Synthesis of Si@SiC Derived from SiO2@C Precursors with Different Carbon Contents

Material Characterization. The Si@SiC samples were characterized using high-resolution X-ray diffraction (XRD; D/MAX Ultima III, Rigaku, Japan) with Cu Kα radiation, scanning electron microscopy (SEM) with energy-dispersive X-ray spectroscopy (S4700/EX-200, Hitachi, Japan), high-resolution transmission electron microscopy (HR-TEM; Tecnai G2, Philips, the Netherlands), X-ray photoelectron spectroscopy (XPS; Multilab 2000, VG, U.K.), Raman spectroscopy (Horiba Jobin-Yvon, France) with a 514 nm diode laser as the excitation light source, thermogravimetric analysis (TGA; TGA50, Shimadzu, Japan), and nitrogen physisorption (ASAP 2020, Micromeritics). Electrochemical Characterization. Electrodes were prepared by casting a mixture of active material ([email protected], Si@SiC-1, Si@SiC-2, or Si@SiC-3), Timcal (Super P) carbon, and lithium polyacrylate at a weight ratio of 7:2:1 onto a copper foil. Then, the foils were dried at 100 °C in a vacuum oven overnight and punched into 14 mm diameter disks. Coin cells (2032-type) were assembled in an Ar-filled glovebox; they consisted of the Si@SiC working electrode, a Li metal counter/ reference electrode, and a glass fiber (Whatman) separator containing an electrolyte of 1 M LiPF6 dissolved in ethylene carbonate and dimethyl carbonate (1:1 in vol %) with the addition of 10 wt % fluoride ethylene carbonate. The mass loading of the active material for most electrodes was controlled to be approximately 0.50−0.70 mg cm−2. Furthermore, the Si@SiC electrode with higher mass loading of 2.0 mg cm−2 was fabricated and tested for practical application. The cells were galvanostatically charged and discharged in the potential range of 0.01−1.5 V versus Li/Li+ using an automatic battery cycler (WonATech-WBCS 3000). In addition, cyclic voltammetry (CV) tests were also performed using a potentiostat (Gamry-PC 750) at a scan rate of 0.1 mV s−1 in the potential range of 0.0−1.5 V versus Li/Li+. The electrochemical impedance of the cells was measured at various frequencies ranging from 100 mHz to 100 kHz. Further, coin-

EXPERIMENTAL SECTION

Preparation of Si@SiC. Analytical-grade silica nanopowder (SiO2, 99.8%, average diameter = 20 nm), poly(vinylpyrrolidone) (PVP, average molecular weight = 40 000), and magnesium (98%, 325 mesh) were supplied by Sigma-Aldrich and used as received without further purification. To synthesize Si@SiC, SiO2@C precursors with various carbon contents were first prepared. In detail, 10 g of SiO2 was dispersed in a certain amount of 2% PVP aqueous solution to form a homogeneous suspension with a PVP/SiO2 mass ratio of 0.5, 1, 2, or 3. Then, these homogeneous suspensions of SiO2 and PVP were dried at 100 °C to obtain SiO2@PVP powders. These powders were then carbonized at 600 °C in Ar atmosphere in a vertical tube furnace for 1 h. The resultant samples starting from PVP/SiO2 mass ratios of 0.5, 1, 32791

DOI: 10.1021/acsami.7b10658 ACS Appl. Mater. Interfaces 2017, 9, 32790−32800

Research Article

ACS Applied Materials & Interfaces

Figure 1. (a) XRD patterns of as-prepared samples: (a1) Si, (a2) [email protected], (a3) Si@SiC-1, (a4) Si@SiC-2, and (a5) Si@SiC-3. (b) XRD patterns of the samples following acid leaching: (b1) Si, (b2) [email protected], (b3) Si@SiC-1, (b4) Si@SiC-2, and (b5) Si@SiC-3.

Figure 2. SEM images of the (a) pure Si, (b) [email protected], (c) Si@SiC-1, (d) Si@SiC-2, and (e) Si@SiC-3 samples. constant current/constant voltage mode. It should be noted that, prior to fabrication of the full cell, the [email protected] electrode was prelithiated by discharge−charge cycling for one cycle between 0.01 and 1.5 V

type full cells were fabricated using the [email protected] anode and a LiCoO2 cathode with a total cathode/anode capacity ratio of 0.9 and charged/discharged in the cell potential range of 2.5−4.1 V in a 32792

DOI: 10.1021/acsami.7b10658 ACS Appl. Mater. Interfaces 2017, 9, 32790−32800

Research Article

ACS Applied Materials & Interfaces Table 1. Physicochemical Properties of the Pure Si and Si@SiC Samples sample

total surface area (m2 g−1)

microporous area (m2 g−1)

mesoporous area (m2 g−1)

total pore volume (cm3 g−1)

microporous volume (cm3 g−1)

mesoporous volume (cm3 g−1)

mean pore diameter (nm)

Si [email protected] Si@SiC-1 Si@SiC-2 Si@SiC-3

40.2 90.2 94.1 123.7 105.6

1.5 1.3 0.1 4.4 12.1

38.7 88.9 94.0 119.3 93.5

0.42 0.49 0.55 0.57 0.61

0.01 0.01 0.01 0.02 0.03

0.41 0.48 0.54 0.55 0.58

32 22 23 19 23

versus Li/Li+ at the rate of C/10 using a separate half-cell. Then, the half-cell was disassembled and the prelithiated [email protected] was collected for preparing the full cell. In addition, the LiCoO2 cathode was fabricated by slurry casting method using LiCoO2 (SigmaAldrich), Super P carbon conductive additive, and PVDF binder. The LiCoO2 cathode was directly used in the full cell without any further treatment.

reaction of newly formed Si with carbon. Using Scherrer’s equation, the size of the SiC crystallites in the [email protected], Si@ SiC-1, and Si@SiC-2 samples was calculated to be 9, 7, and 6 nm, respectively. In addition, as the carbon content of the precursor varied by changing the weight ratio of PVP to SiO2 from 0.5 to 3, the carbon contents in the [email protected], SiO2@ C-1, SiO2@C-2, and SiO2@C-3 samples were measured to be about 4.37, 9.64, 18.56, and 27.06 wt %, respectively (Figure S2). In Figure 1b, the relative intensity of Si is significantly diminished for the samples with higher carbon content. This phenomenon indicates that more of the Si participated in the reaction in the samples with higher carbon content. Surprisingly, for the Si@SiC-3 sample, which was produced from the precursor (SiO2@C-3) with the highest carbon content of 27.06%, all of the Si was converted to SiC during Mg thermal reduction. Thus, no characteristic peak corresponding to the Si phase was observed in the XRD pattern of the Si@ SiC-3 sample. In addition, owing to the excess carbon, which can act as an inhibiting agent to prevent the growth of Si crystals, finer SiC crystals were obtained, as calculated from the XRD patterns; the size ranged from 9 nm for [email protected] to 6 nm for Si@SiC-3. Figure 2a shows the SEM image of the pure Si sample. This sample exhibits a porous structure composed of interconnected pore networks with an average pore diameter of 150 nm, and the interconnected Si wall thickness was approximately 100 nm. These pores appear to have formed by a selective etching process. Following acid leaching, undesirable products, such as Mg2SiO4, Mg2Si, and MgO, were removed and left empty spaces, leading to the formation of the highly porous Si structure. We recall that the pure Si sample was prepared using nanosilica with an average diameter of 10−20 nm. However, the crystalline Si grains obtained through Mg thermal reduction were much larger than the initial size of the raw silica material. This phenomenon is attributed to melting and fusion of Si particles during Mg thermal reduction. In contrast, the Si@SiC samples were composed of fine Si and SiC particles, as shown in Figure 2b−e. Obviously, owing to the presence of carbon, the SiO2 particles were enclosed by a carbon layer and reduced by Mg vapor, leading to the formation of finer Si particles. It is inferred that SiC was produced simultaneously with the reduction of SiO2. This reaction is endothermic, so it consumes energy and reduces the local temperature on the surface of the newly formed Si particles. Owing to the high thermal conductivity of SiC, it can act as an effective heat transfer agent in releasing heat energy generated during SiO2 reduction to the outside environment. Moreover, the high thermal stability of SiC is believed to contribute to preserving the fine Si particles; because of the extremely high melting point of SiC (2730 °C), it can remain stable during Mg thermal reduction. Accordingly, Si@SiC particles could not become aggregated, and the fine Si particles could be retained.

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RESULTS AND DISCUSSION Figure 1a shows the XRD patterns of the as-prepared Si@SiC samples. The as-prepared samples contained many impurity phases, such as Mg2SiO4, MgO, and Mg2Si. Although MgO is estimated to be the main product formed by the reaction of SiO2 with Mg, undesirable products, such as Mg2SiO4 and Mg2Si, were also found because of side reactions occurring during thermal reduction. These chemical reactions can be described as follows.43,45 SiO2 + 2Mg = 2MgO + Si

(1)

SiO2 + 4Mg = 2MgO + Mg 2Si

(2)

Si + Mg = Mg 2Si

(3)

SiO2 + 2MgO = Mg 2SiO4

(4)

In addition, when the precursors contain carbon, as [email protected], SiO2@C-1, SiO2@C-2, and SiO2@C-3 (Figure S1), a new phase of SiC can form as a product of the following reaction. SiO2 @C + 2Mg = SiC + 2MgO

(5)

It is known that a direct reaction between Si and C can occur at temperatures exceeding 1100 °C, which is much higher than the temperature used for Mg thermal reduction in this study. However, it is worth noting that the Mg thermal reaction releases a huge amount of energy. This energy would immediately heat the surrounding atmosphere to the temperature needed for the reaction between Si and C, leading to the formation of SiC. As shown in Figure 1a, because the peaks corresponding to the SiC phase overlap impurity phase peaks, the appearance of the crystalline peaks corresponding to SiC in the XRD patterns of the as-prepared Si@SiC samples was unclear. However, the peaks corresponding to SiC were clearly seen after the impurity phases of Mg2SiO4, Mg2Si, and MgO were removed by acid leaching, as shown in Figure 1b. Following acid leaching, the finally obtained products were pure Si or Si@SiC, depending on the precursor. In particular, for the pure Si sample, the six distinct peaks observed at 28.43, 47.30, 56.12, 69.12, 76.37, and 88.03° are in good agreement with the characteristic peaks corresponding to diamond-cubic-structured Si (JCPDS No. 27-1402). No impurity peaks were found in the XRD pattern of the Si sample. In contrast, for the [email protected], Si@SiC-1, and Si@SiC-2 samples, three additional broad peaks were observed at 35.71, 60.13, and 71.91°, corresponding to SiC formed by continuous 32793

DOI: 10.1021/acsami.7b10658 ACS Appl. Mater. Interfaces 2017, 9, 32790−32800

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ACS Applied Materials & Interfaces

Figure 3. Low- and high-magnification TEM images of (a−c) pure Si and (d, e) [email protected] samples; (f) d-spacing of SiC and Si corresponding to points 1 and 2 in (e), respectively.

characteristic absorption peaks corresponding to Si were observed at 517 and 948 cm−1. However, for the Si@SiC-3 sample, these two peaks totally disappeared. Instead, two new peaks corresponding to the disorder-induced (D) band and graphitic (G) band of carbon were found at 1353 and 1578 cm−1, respectively. This phenomenon indicates the complete conversion of Si into the SiC phase and the presence of residual carbon even after the Mg thermal reduction process was complete. The Raman analysis result is quite consistent with the above-mentioned XRD result (Figure 1) for the Si@SiC-3 sample. In contrast, the typical peaks corresponding to carbon were not observed for the [email protected], Si@SiC-1, and Si@SiC2 samples. This result suggests that all of the carbon in the samples was converted to SiC by carbothermal reaction of Si. In the presence of excess carbon, all of the newly formed Si would be converted to SiC by carbothermal reduction. In contrast, when the amount of carbon in the precursor was insufficient, the obtained samples were composed of Si and SiC. By adjusting the carbon content of the precursor material, Si@SiC products with various SiC contents could be successfully prepared as desired. The nanostructure of the pure Si and Si@SiC samples was further analyzed using HR-TEM. TEM images of the pure Si sample display dark regions corresponding to densely fused Si (Figure 3a−c), which are formed at high temperature by exothermic reaction of SiO2 with Mg. This result is in good agreement with the SEM images in Figure 2a. In contrast, TEM images of [email protected] reveal porous structure, as shown in Figure 3d. Moreover, the HR-TEM image (Figure 3e) of Si@

The porous structure of the pure Si and Si@SiC samples was further examined by N2 adsorption−desorption isotherm measurement. Figure S3a depicts the N 2 adsorption− desorption isotherm, which shows a typical type-IV isotherm having a wide range of pore sizes and the Barrett−Joyner− Halenda pore size distribution (inset) of the pure Si sample. From the N2 adsorption−desorption isotherm data, the specific area, average pore diameter, and pore volume for the pure Si sample were estimated to be about 40.2 m2 g−1, 32 nm, and 0.42 cm3 g−1, respectively. Similarly, the Si@SiC samples were also characterized by N2 adsorption−desorption isotherms, as shown in Figure S3b−e. The physicochemical properties of the samples obtained from Figure S3 are enumerated in Table 1 in detail. Overall, the total surface areas of the Si@SiC samples derived from the SiO2@C precursors with different carbon contents were higher than those of the pure Si sample from the SiO2 precursor. Specifically, the Si@SiC samples had total surface area ranging from 90.2 to 123.7 m2 g−1, which is more than twice that of the pure Si. In addition, because carbon was present in the SiO2@C precursors, the total pore volume of the Si@SiC samples was 16.7−45.2% larger than that of the pure Si. This result implies that the synthesized Si@SiC samples have superior porosity and surface area compared to the pure Si sample, which are expected to enhance the electrochemical performance of the electrode. Further, the presence of residual carbon in the Si@SiC sample synthesized from the precursor with high carbon content was verified by Raman spectra, as shown in Figure S4. For the [email protected], Si@SiC-1, and Si@SiC-2 samples, two 32794

DOI: 10.1021/acsami.7b10658 ACS Appl. Mater. Interfaces 2017, 9, 32790−32800

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ACS Applied Materials & Interfaces

Figure 4. Discharge−charge potential profiles of the (a) Si and (b) [email protected] electrodes. (c) Cyclability and (d) Coulombic efficiency of the Si and [email protected] electrodes at a rate of C/10. (e) Cyclic voltammogram of the [email protected] electrode measured at a scan rate of 0.1 mV s−1.

spectrum of the Si 2p core for the [email protected] sample was deconvoluted into three peaks centered at 99.7, 101.6, and 103.8 eV, which correspond to Si−Si, Si−C, and Si−O bonding. It is easy to understand that the Si−Si and Si−C bonds correspond to Si and SiC in the Si@SiC samples. Further, the appearance of Si−O bonding can be attributed to the surface oxidation of Si. Surprisingly, the bonding energies of Si−Si, Si−C, and Si−O shifted to lower values as the carbon content in the precursor of the Si@SiC increased (Figure S6). We note that for the [email protected] sample the intensity of Si−Si bonding in the Si 2p X-ray photoelectron spectrum was much lower than that of Si−C or Si−O bonding, although Si is the main component of the sample. This result implies that SiC was preferentially concentrated on the surface of agglomerated Si@ SiC particles. Further, in the X-ray photoelectron spectrum of Si@SiC-2, the peak corresponding to Si−Si bonding at around 90.6 eV was not detected, although the presence of Si was

SiC-0.5 exhibits many clear stripes with random lattice orientation, revealing the polycrystalline nature of the obtained Si@SiC products. To identify the nanocrystal in [email protected], the d-spacing of the nanocrystal was measured from the HRTEM image. Figure 3f (top) shows the d-spacing of the crystal corresponding to region 1 in Figure 3e; it was calculated to be about 2.5 Å, which corresponds to the (111) plane of the SiC crystal. In addition, clear lattice fringes of the (111) plane of Si with a d-spacing of 3.2 Å were also observed, as shown in Figure 3f (bottom). A similar phenomenon was observed in the TEM images of Si@SiC-1 and Si@SiC-2, as shown in Figure S5. These results demonstrate that Si and SiC were randomly mixed together in the Si@SiC samples. Further, the surface chemistry of the Si@SiC samples was carefully examined by XPS, as shown in Figure S6. Here, the binding energies were calibrated with respect to the C 1s peak at 284.8 eV. As observed in Figure S6a, the X-ray photoelectron 32795

DOI: 10.1021/acsami.7b10658 ACS Appl. Mater. Interfaces 2017, 9, 32790−32800

Research Article

ACS Applied Materials & Interfaces

is in good agreement with the report by Zhang et al.,47 who directly observed lithiation in a SiC−SiO2 core−shell using in situ TEM and found that the SiC core maintained its crystalline structure. This result suggested that SiC did not participate in the lithiation process. Fortunately, even though the specific capacity of Si@SiC was lowered by the presence of SiC, the cyclability of the [email protected] sample predominated over that of the pure Si electrode; a 76% capacity loss for the Si electrode and a 46% capacity loss for the [email protected] electrode were observed following 200 cycles. The superior cyclability of the [email protected] electrode is attributed to the presence of SiC. As verified by the SEM and TEM images in Figures 2 and 3, the particle size became finer and the pore volume was larger for the Si@SiC sample owing to the presence of SiC. The nanostructure of materials is well known to play a crucial role in suppressing pulverization by mitigating the absolute volume change of the electrode induced by lithiation/delithiation. In addition, the presence of more empty pores can significantly reduce the strain generated during lithiation. Consequently, the integrity of the structure and electrical conductivity of the electrode would be maintained continuously for a long cycling time. On the other hand, the lower proportion of active Si in [email protected] is believed to contribute to the improved cyclability of the Si@ SiC-0.5 electrode simply because the smaller volume of Si as an active material can reduce the volume expansion of the electrode. Accordingly, the cyclability of the Si@SiC electrode was better than that of the pure Si electrode. Following cycling, the morphological changes in the Si and [email protected] electrodes were observed by SEM and TEM. As shown in Figure S9, prior to cycling, the thickness of the Si electrode was ∼18 μm (Figure S9a). Following cycling, the thickness of the Si electrode changed significantly: at the fully discharged (lithiated) state, the thickness of the electrode was measured to be ∼41 μm (Figure S9b). This indicates that the thickness of the electrode increased by 127% compared with its initial value. At the full-charged (delithiation) state, the thickness was measured to be ∼34 μm corresponding to 89% increase (Figure S9c). The huge thickening of the pure Si electrode at the full lithiation state is attributed to the cracking of the electrode as well as the volume expansion of Si itself during the lithiation process. In contrast, the [email protected] electrode exhibits a smaller change in thickness. Prior to cycling, the initial thickness of the [email protected] electrode was estimated to be ∼22 μm (Figure S9d). However, following cycling, the thickness was measured to be ∼31 μm (Figure S9e) at the full lithiation state and ∼28 μm (Figure S9f) at the full delithiation state corresponding to thickness increase by 40 and 27%. Obviously, the pure Si exhibits larger volume change compared to the [email protected] electrode, due to the higher proportion of Si participating in the lithiation/delithiation process. In addition, the high specific area and porosity as well as the presence of SiC are believed to effectively mitigate the volume expansion during cell operation. The volume change of the [email protected] electrode in nanoscale following cycling was observed by TEM, as shown in Figure S9g−i. The [email protected] electrode did not exhibit any significant change in nanostructure except amorphization of crystalline Si. Obviously, following cycling, the crystalline Si converted to the amorphous Si, whereas SiC was unchanged. Prior to cycling, the SAED pattern of the Si@SiC exhibits many concentric circles (inset of Figure 3d). These circles originate from both crystalline Si and SiC. Following cycling, only two concentric circles were observed, which correspond to nano SiC only

confirmed by the XRD results, as shown in Figure 1b. This result reveals that Si was completely covered with SiC and SiO2 in this sample. Similarly, the peak assigned to Si−Si bonding definitely disappeared in the X-ray photoelectron spectrum of the Si@SiC-3 sample because Si@SiC-3 contained only SiC and residual carbon, which was confirmed by the XRD and Raman analyses (Figures 1b and S4). The quantitative compositions of [email protected], Si@SiC-1, and Si@SiC-2 were determined using TGA. From the TGA data shown in Figure S7, the Si contents of the [email protected], Si@SiC-1, and Si@SiC-2 samples were calculated to be 82.1, 60.2, and 32.8 wt %, respectively. For the Si@SiC-3 sample, the SiC content was determined to be about 70.2 wt %, as the sample contained only SiC and carbon. The electrochemical characteristics of the Si@SiC samples were investigated using a CR2032 coin-type cell. Figure 4a,b shows the potential profiles of the pure Si and [email protected] electrodes, respectively, at a charge−discharge rate of C/10 (360 mAh g−1) for 200 cycles in the potential range of 0.01− 1.5 V versus Li/Li+. In the first cycle, the pure Si electrode exhibits discharge and charge capacities of 2746 and 2231 mAh g−1, respectively (Figure 4a). This result indicates that only about 62 wt % of the Si participated in the lithiation/ delithiation process and contributed to the specific capacity, considering the theoretical capacity of Si. Here, the irreversible capacity in the first cycle for the pure Si electrode was 515 mAh g−1. The capacity loss is attributed to the formation of the SEI layer and decomposition of the native oxide.46 In contrast, the [email protected] electrode exhibits a slightly lower discharge capacity of 2134 mAh g−1 and a charge capacity of 1642 mAh g−1 in the first cycle (Figure 4b). The charge capacity of 1642 mAh g−1 corresponds to 45 wt % active Si, which is smaller than the value for the pure Si (62 wt %). As seen in Figure 4a, the first discharge potential profile exhibits a plateau below 0.1 V versus Li/Li+. This plateau is assigned to the lithiation process, which normally requires large polarization in the first cycle. In addition, another very short plateau was observed at potentials below 1.5 V versus Li/Li+. This can be attributed to the side reaction of the Si electrode, in which electrolyte decomposition and reduction of the native oxide can occur. In the next step, the electrode was charged and Li+ was released from the Li−Si alloy that had formed during the previous discharge. The main charge potential plateau was observed at around 0.46 V versus Li/Li+. Similarly, for the Si@ SiC-0.5 electrode, a plateau below 0.1 V versus Li/Li+ during the first discharge and a plateau around 0.46 V versus Li/Li+ for the first charge were also observed. From the second cycle onward, the potential profiles for both the pure Si and [email protected] electrodes exhibit two discharge plateaus centered around 0.07 and 0.26 V versus Li/Li+ and two charge plateaus located around 0.32 and 0.45 V versus Li/Li+ (Figure 4a,b). Figure 4c shows the cyclability of the Si and [email protected] electrodes measured at a rate of C/10, which corresponds to a current density of 360 mAh g−1. In the first cycle, the reversible capacities of the Si and [email protected] electrodes were estimated to be about 2231 and 1642 mAh g−1, respectively. Surprisingly, the presence of SiC in the [email protected] electrode lowered the specific capacity compared to that of the pure Si electrode. Li et al.20 reported that SiC is an active component in lithiation/ delithiation. However, in the present study, even though it occurred in ultrafine particles of size 9 nm, the SiC appears to be an inactive component that contributes very little to the specific capacity of Si@SiC, as shown in Figure S8. This result 32796

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Figure 5. (a) Rate capability of the pure Si and [email protected] electrodes under discharge−charge rates of C/20 to 10C. (b) Rate capability of the Si@ SiC-0.5 electrode under a fixed discharge rate of C/20 and varying charge rates of C/20 to 20C.

large capacity of 961 mAh g−1, corresponding to a capacity retention of 59.8%. This demonstrates that the [email protected] electrode with high mass loading can also operate stably by maintaining its high capacity. To further understand the lithiation/delithiation mechanism in the [email protected] electrode, a CV test of the [email protected] electrode was conducted using a potentiostat at a scan rate of 0.1 mV s−1 for five cycles in the potential range of 0.0−1.5 V versus Li/Li+ starting from the open-circuit potential. As shown in Figure 4e, in the first cathodic scan, a minor peak located at 1.4 V versus Li/Li+ was recorded. This peak corresponds to the side reaction occurring in the first cycle.19 Next, the cathodic current increased slightly until the electrode potential decreased to 0.12 V versus Li/Li+. Then, the cathodic current suddenly rose when the electrode potential decreased from 0.12 to 0.0 V versus Li/Li+. This step corresponds to the formation of Li−Si alloy and further formation of the SEI layer. In the first anodic scan, two delithiation peaks in the CV plot appeared at potentials of around 0.34 and 0.46 V versus Li/Li+ (Figure 4e), which are thought to correspond to the Li15Si4/amorphousLixSi and LixSi/Si + amorphous-LiySi phase transitions.19,50,51 From the second cycle onward, the cathodic curve showed a distinctly different shape from that observed in the first cycle, indicating the change in the phase -transformation process. In particular, in the fifth discharge curve, three main peaks were observed at 0.28, 0.23, and 0.06 V versus Li/Li+, as well as a tiny peak at 0.02 V versus Li/Li+. These peaks correspond to lithiation of amorphous Si in the different phases. Figure 5a shows the rate capability of the pure Si and Si@ SiC-0.5 electrodes under charge−discharge rates of C/20 to 10C for all five cycles. In the first 10 cycles under C rates lower than C/2, the pure Si electrode obviously exhibits higher specific capacity than the Si@SiC electrode because the former contained a larger proportion of the Si active material than the [email protected] electrode. As the discharge−charge rate increased to above C/2, the capacity of the pure Si electrode dropped rapidly, whereas that of the [email protected] electrode decreased negligibly. In particular, at higher discharge−charge rates, the capacity of the pure Si electrode faded severely. Specifically, the average reversible capacities of the Si electrode measured at rates of C/5, C/2, C, 2C, 5C, and 10C were 1645, 1335, 1039, 548, 144, and 5.8 mAh g−1, respectively. In contrast, the Si@ SiC-0.5 electrode illustrated superior rate capability. In particular, at the corresponding rates of C/5, C/2, C, 2C, and 5C, the average reversible capacities of 1597, 1448, 1296,

(inset of Figure S9h). Furthermore, according to HR-TEM image (Figure S9i), very fine SiC crystals were found to distribute in the amorphous Si phase. A comparison of the specific capacity of the Si@SiC electrodes reveals that the capacity declined rapidly with increasing SiC content of the electrode, as shown in Figure S8; the values were 801, 331, and 165 mAh g−1 for the Si@SiC-1, Si@SiC-2, and Si@SiC-3 electrodes, respectively, in the first cycle. This result is attributed to the presence of a relatively large amount of SiC, which can hinder lithiation of Si in the Si@SiC electrodes with higher SiC content. The lithiation process involves a series of steps: (i) Li ions approach the surface of the active material, (ii) Li passes through the electrode/electrolyte interface, and (iii) Li diffuses from the surface to the interior of the bulk in primary particles. If any step is interrupted, lithiation cannot proceed. Therefore, a high content of inactive SiC would block the route to active Si deep inside the Si@SiC bulk, leading to a lower specific capacity of the electrode with higher SiC content. Figure 4d shows the Coulombic efficiency of the Si and Si@ SiC-0.5 electrodes for 200 cycles. In the first cycles, Coulombic efficiencies of 81.2 and 76.9% were recorded for the Si and Si@ SiC-0.5 electrodes, respectively. The higher irreversible capacity of the [email protected] electrode in the first cycle can be ascribed to the higher specific area of [email protected]. As shown in Table 1, the specific surface area of the synthesized [email protected] was more than twice as high as that of the pure Si. Although SiC in the [email protected] electrode was inactive phase and SEI would not grow on its surface, SEI still grew on the surface of active Si phase present in the [email protected] electrode as normal. Herein, it is probably due to the fact that the specific surface area of distinct Si present in the [email protected] electrode was still higher than that of pure Si, resulting in the lower Coulombic efficiency of the [email protected] electrode.48,49 In the second cycle, the Coulombic efficiencies of both the pure Si and [email protected] electrodes were significantly improved to 94.2 and 94.7%, respectively. From the 10th cycle onward, both the pure Si and [email protected] electrodes exhibit high Coulombic efficiency (over 99%), demonstrating excellent reversibility. In addition, the Si@SiC electrode with higher mass loading of ∼2 mg cm−2 was fabricated and charge−discharge tested. As shown in Figure S10, even with a high mass loading of ∼2 mg cm−2, the [email protected] electrode exhibits a high specific capacity of 1607 mAh g−1 at a rate of C/10 in the first cycle. Following 100 cycles, the [email protected] electrode still delivered a relatively 32797

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Figure 6. (a) Cell potential profile and (b) cyclability of the full cell composed of [email protected] and LiCoO2.

1076, and 706 mAh g−1 were reached as for the [email protected] electrode. In general, the [email protected] electrode values showed much higher reversible capacity than that of the pure Si electrode at the same rates. Even at such a high rate of 10C (corresponding to 36 Ah g−1), the [email protected] electrode still delivered a high capacity of 438 mAh g−1. This value is much higher than the theoretical capacity of commercial graphite. In addition, when the discharge−charge rate was returned steeply to the low initial rate of C/20, the specific capacity of the Si@ SiC-0.5 electrode reached 1189 mAh g−1, which is higher than that of the pure Si electrode, 1145 mAh g−1. This superior rate capability of the [email protected] electrode is attributed to the synergistic effect of three factors: (i) nanoparticle structure of Si, which supplies a short Li diffusion path in the solid phase; (ii) high specific area of the electrode, which provides more sites for charge transfer and passage of Li though the SEI; and (iii) high chemical stability of SiC, which can hinder the formation of SEI layer on the surface of Si, resulting in better cyclability. Obviously, prior to cycling, the electrolyte resistance of both pure Si and [email protected] electrodes is ∼3 Ω and their charge transfer resistances are 66 and 48 Ω, respectively (Figure S11a,b). The charge transfer resistance of the [email protected] was found to be lower compared to the pure Si. This arises from the higher specific area of the [email protected] electrode than that of pure Si. The [email protected] electrode provided more sites for electron and Li-ion transfer across the liquid−solid interface, resulting in more favorable charge transfer process for the Si@ SiC-0.5 electrode. In subsequent cycles of charge−discharge, the charge transfer resistances of both the pure Si and [email protected] electrodes exhibit a downward trend because amorphization of Si in the first cycle makes the alloying/dealloying process easier in next cycles. From the third cycle onward, by using the model in Figure S11c, the total charge transfer resistance of the [email protected] electrode was estimated to be approximately 9−13 Ω (Figure S11b), whereas that of the pure Si increased consecutively from 68 to 83 Ω (Figure S11a). This is attributed to the growth of SEI layer on the surface of the pure Si. Herein, due to the presence of SiC film on the [email protected] electrode, the formation and growth of the SEI layer could be hampered. Accordingly, the [email protected] electrode exhibits superior cyclability and rate capability to the pure Si. The rate capability of the [email protected] electrode was further tested by fixing the discharge rate at C/20 while varying the charge rate from C/20 to 20C. Interestingly, the [email protected] electrode exhibits excellent rate capability even at the extremely high charge rate of 20C. From Figure 5b, the capacity of the

[email protected] electrode is estimated to be about 989 mAh g−1 at a rate of 20C. This result suggests that the capacity of the Si@ SiC-0.5 electrode is limited by the discharge process, which determines how much Li can be stored inside the electrode.52−54 To investigate the applicability of [email protected] in LIBs, a cointype full cell composed of a [email protected] anode and a LiCoO2 cathode was fabricated. It should be noted that the [email protected] anode was prelithiated in a specially designed cell as shown in Figure S12 before the full cell was assembled to eliminate the effect of low Coulombic efficiency in the first cycles, and the capacity ratio of the cathode and anode was controlled to be about 0.9 to prevent Li plating on the anode. The full cell was galvanostatically charge−discharge tested within the cell potential range of 2.5−4.1 V in constant current/constant voltage mode. Because the LiCoO2 was fully utilized, the capacity of the full cell was calculated on the basis of the mass loading of the cathode. Figure 6a shows the cell potential profile of the full cell. The charge−discharge curve of the full cell was close to that of a LiCoO2 half-cell, as shown in Figure S13, because the [email protected] electrode has a very low lithiation/delithiation potential, with the main plateau at around 0.3 V versus Li/Li+. The first charge capacity and discharge capacity of the full cell were about 169 and 142 mAh g−1, respectively (Figure 6a). Even though it contained the prelithiated [email protected] electrode, the full cell still exhibits a relatively high irreversible capacity of 27 mAh g−1 in the first cycle.38 This can be attributed to a further side reaction occurring in the full cell. In subsequent cycles, the capacity of the full cell gradually diminished owing to degradation of the [email protected] electrode. After undergoing 50 repeated charge− discharge cycles, the full cell retained a capacity of 115 mAh g−1, which corresponds to an average capacity loss per cycle of 0.38%. This illustrates the excellent cyclability of the coin-type full cell based on the [email protected] anode and LiCoO2 cathode.

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CONCLUSIONS We succeeded in synthesizing the composite material Si@SiC, in which fine Si active materials are well mixed with ultrafine SiC, making it an outstanding anode material for rechargeable LIBs. In particular, the [email protected] electrode, which was produced from a precursor with a PVP/SiO2 ratio of 0.5, exhibits excellent electrochemical performance characteristics, such as high specific capacity, cyclability, and rate capability. The superior electrochemical characteristics of the [email protected] electrode are attributed to the presence of nano SiC, the 32798

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ultrafine particles and high electrical conductivity of which enable fast Li diffusion and low polarization, and to the high specific area of the electrode, which facilitated the charge transfer reaction. In particular, because it contained an optimum amount of SiC, the Si@SiC electrode exhibits significantly improved cyclability. Nevertheless, the abundance of SiC degraded the specific capacity of the Si@SiC electrode because it served as an inactive buffer layer for the Si active material inside. Furthermore, a full cell fabricated by coupling the [email protected] anode with a LiCoO2 cathode exhibits excellent cyclability. These outstanding results demonstrate great promise for applicability of the [email protected] electrode as an anode material for the next generation of high-energy LIBs.

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REFERENCES

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acsami.7b10658. TEM images of the SiO2@C synthesized with various ratios of PVP/SiO2; thermogravimetric analysis of the SiO2@C samples; N2 adsorption−desorption isotherms of the Si and Si@SiC; Raman spectra of the Si@SiC; TEM images of the Si@SiC-1 and Si@SiC-2; XPS spectra of Si 2p core level of the Si@SiC; TGA data of the Si@SiC; cyclability of the Si@SiC electrodes; crosssectional-view and top-view SEM images of the pure Si before and after cycling at the fully lithiated and delithiated states; cross-sectional-view and top-view SEM images of the [email protected] before and after cycling at the fully lithiated and delithiated states; TEM images and relevant SAED pattern of the [email protected] electrode after 200 cycles of charge−discharge; cyclability of the [email protected] electrode with a high mass loading of ∼2 mg cm−2 at a charge−discharge rate of C/10 for 100 cycles; electrochemical impedance spectroscopies of the Si and [email protected] electrodes after different cycles of charge− discharge; equivalent circuit for the Si and Si@SiC electrodes after cycling; the specially designed cell for the prelithiation of the [email protected] electrode; potential profile of the LiCoO2 electrode at a rate of C/10 (PDF)

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Research Article

AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. Tel.: +82-62-530-1704. Fax: +82-62530-1699. ORCID

Chan-Jin Park: 0000-0002-3993-1908 Author Contributions §

D.T.N. and H.T.T.L. contributed equally to this work.

Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This study was supported by The Leading Human Resource Training Program of Regional Neo Industry (NRF2017H1D5A1044874), the framework of international cooperation program (NRF-2016K2A9A1A09914221), and the basic research program (NRF-2015R1D1A3A01019399) through the National Research Foundation of Korea (NRF). 32799

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ACS Applied Materials & Interfaces

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DOI: 10.1021/acsami.7b10658 ACS Appl. Mater. Interfaces 2017, 9, 32790−32800