Amorphous carbon-coated prickle-like silicon of micro and nano hybrid anode materials for lithium-ion batteries

Solid State Ionics 260 (2014) 36–42

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Amorphous carbon-coated prickle-like silicon of micro and nano hybrid anode materials for lithium-ion batteries Jung Sub Kim a,b, Martin Halim a,c, Dongjin Byun b, Joong Kee Lee a,c,⁎ a b c

Center for Energy Convergence, Korea Institute of Science and Technology, Hwarangno 14-gil 5, Seongbuk-gu, Seoul 136-791, Republic of Korea Department of Material Science & Engineering, Korea University, Seoul 136-713, Republic of Korea Department of Energy and Environmental Engineering, University of Science and Technology, Gwahangno, Yuseong-gu, Daejeon, 305-333, Republic of Korea

a r t i c l e

i n f o

Article history: Received 18 November 2013 Received in revised form 4 March 2014 Accepted 11 March 2014 Available online 28 March 2014 Keywords: Prickle-like silicon Carbon coating Polypropylene Thermal chemical vapor deposition

a b s t r a c t Carbon coated prickle-like Si particles (PS@C) are prepared by metal-assisted chemical etching and subsequent coating with an amorphous carbon film carried out by thermal chemical vapor deposition (CVD). The electrochemical characteristics of PS@C employed as anode material for lithium-ion batteries are investigated in order to find a relationship between interfacial properties and electrochemical performance. The unique morphology of prickle-like Si (PS) having empty space can accommodate volume expansion during the lithiation and delithiation. Additionally, an amorphous carbon coating layer with a thickness of 10–15 nm deposited onto the PS prepared by thermal CVD is investigated as an effective way to enhance the cycle stability and rate capability of the PS electrode due to improved interfacial characteristics. The micro and nano hybrid structure of the PS material combined with the 12 wt.% amorphous carbon layer plays an important role in enhancing the electrochemical performance. © 2014 Elsevier B.V. All rights reserved.

1. Introduction As a prospective anode material for lithium-ion batteries, Si has been considered as an alternative to carbon-based anodes for next generation lithium-ion batteries because of its natural abundance, availability, environmental friendliness, and most importantly its low discharge potential and high theoretical capacity (4200 mA h g−1 in Li4.4Si) [1]. However, the practical application of Si anodes has so far been mainly hindered by its low electrical conductivity, low lithium diffusion rate and enormous volume change (300–400%) occurring during the discharge (lithiation) and charge (delithiation) process [2]. The volume change causes bulk Si to be pulverized and lose electrical contact with the conductive additive or current collector, and can also lead to the instability of the solid electrolyte interphase (SEI) [3]. The latter issue encourages continuous consumption of Li-ions for reformation of the SEI layer caused by the breakage of the silicon surface with the progress of the cycle, which leads to an increase in irreversible capacity. In order to improve the cycling stability of Si anodes, great efforts have been made to mitigate the pulverization of Si and improve the stability of the SEI layer. These efforts include the development of Si materials composed of nanostructures, porous structures, or nano-composites,

⁎ Corresponding author. E-mail address: [email protected] (J.K. Lee).

http://dx.doi.org/10.1016/j.ssi.2014.03.013 0167-2738/© 2014 Elsevier B.V. All rights reserved.

the addition of coating layers, and the application of electrolyte additives and novel binders [4–8]. These studies investigate the fact that morphology, volume change, conductivity and surface characteristics of the Si-based electrode materials play a key role in producing high performance active material as anode in lithium-ion batteries. In particular, carbon coating the surface of the active material is very useful due to its nature of forming a stable SEI, structural integrity, and high electric conductivity [9,10]. Accordingly, the well-defined uniformity of carbon layer is of high importance for the electrode performance. Since the nanostructured materials have a small size and large surface area, they are expected to inherently provide short diffusion length for lithium-ions, large interfacial area between the electrodes and the electrolyte, and large alleviation of lattice stress during the operation in lithium-ion batteries. Studies of Si as an anode material are leaning toward almost the whole nanostructure area [11–13]. The nanostructuring of electrodes has been demonstrated to result in lithiumion batteries with superior electrochemical performance, i.e. higher storage capacity, higher rate capability, and excellent cycling performance. In contrast, the high surface area of nanostructured materials may result in undesirable side reactions with the electrolyte under certain conditions, leading to a significant fading of Li storage capacity and, therefore, poor cycling performance. To avoid these undesired problems, various surface modification or morphologies have been investigated such as hollow, core–shell, nanotube, and nest structures [14–17]. On the other hand, when micro-sized Si (10 μm) is fabricated as an anode to be made into a cell, it has a discharging capacity of only

J.S. Kim et al. / Solid State Ionics 260 (2014) 36–42

about 3260 mA h g−1, and a charging capacity of 1170 mA h g−1, indicating an initial coulombic efficiency of 35%. When the cell is continuously charged and discharged over 5 cycles, its discharge capacity rapidly decreases to about 300 mA h g−1, which results in poor cycling performance due to internal resistance caused by volume change [18]. Therefore, micro and nano hybrid materials enable us to use the advantages of each component and to make up for the imperfection. In another approach, Yoo et al. [19] have synthesized Si nanotubes via surface sol–gel reaction and magnesium reduction. The empty space in the cores of nanotube accommodates the large volume changes which results in an improved cycle performance. Chen et al. [15] have also shown that Si core-hollow carbon shell nano-composites with tunable buffer voids tolerate physical stress during charge and discharge of cells. Metal-assisted chemical etching of bulk Si powders is rarely reported and attractive for its unique nanostructures [20–25]. In this work, prickle-like Si (PS) of micro and nano hybrid material is prepared via metal-assisted chemical etching and used as an anode material for lithium secondary batteries. Their electrochemical characteristics are investigated in comparison to pristine bulk Si and amorphous carbon coated prickle-like Si (PS@C). 2. Experimental PS obtained from bulk Si was synthesized via metal-assisted chemical etching. Therefore, commercially available Si powder (Kojundo co., 99.9%) was dispersed in etching solution that consists of 49 wt.% hydrofluoric acid (HF, 4.6 mol L− 1), 60 wt.% silver nitrate (AgNO3, 0.04 mol L− 1) and de-ionized water. The etching solution is stirred mildly for 1 h in order to prevent the detachment of silver. After the etching process, residual silver nanoparticles were eliminated from the PS using a 50 wt.% HNO3 aqueous solution, and the etched silicon powders are filtered out from the etching solution and dried in a vacuum oven at 80 °C for 5 h. The as-synthesized sample (PS) was transferred to a thermal chemical vapor deposition (CVD) reactor. The CVD was heated to 700 °C at a rate of 4.3 °C min−1, and then stream of diluted gases of propylene and argon (10% C3H6, 90% Ar) at a flow rate of 100 cc min− 1 was applied for 10 h to get a uniform carbon coating. After coating, the samples were cooled down to room temperature. The morphologies of pristine Si, PS, and PS@C were observed by scanning electron microscopy (SEM, NOVA NanoSEM200, FEI Corp.) and transmission electron microscopy (TEM, Tecnai, FEI Corp.) as well as chemically analyzed with energy-dispersive X-ray spectroscopy (EDX). The Brunauer–Emmett–Teller (BET) method using an N2 sorption isotherm plot on a Micromeritics apparatus (ASAP 2010) was used to measure surface area, average pore diameter, and the change of total pore volume before and after etching and carbon coating. An investigation of the crystal structure was carried out using an X-ray diffractometer (XRD, Rigaku) equipped by Cu-K α (λ = 1.5418 Å) radiation in the 2 θ range from 10 to 90°. After the etching process, the chemical bonding of the remaining Ag was analyzed by X-ray photoelectron spectroscopy (XPS, PHI 5000 VersaProbe). The structures of pristine Si, PS, and PS@C were obtained by Raman spectroscopy (Nicolet Almega XR dispersive Raman, Thermo electron co.) using a 632 nm excitation laser. A 50× microscope objective and an exposure time of 4 s combined with low laser power was used to avoid heating effects. The spectra were recorded at 4 cm−1 resolution between 4000 and 90 cm−1. The laser spot has a diameter of about 2 μm on the sample surface. Elemental analysis (ES) by the CHN method was carried out to measure the carbon content. Therefore, the samples were burned in excess oxygen and then combustion products (carbon dioxide, water, and nitric oxide) were collected by various traps. Furthermore, the collected combustion products were weighed to calculate the unknown sample. Electrochemical analysis were carried out using two-electrode cells with pristine Si, PS, and PS@C as the working electrode and lithium metal as counter electrode. The base electrolyte used was 1 M LiPF6 dissolved in a 1:1:1 (v/v) mixture of ethylene carbonate, ethylmethyl and

37

dimethyl carbonate. As reductive additive 30% vol fluoroethylene carbonate (FooSung co.) were added into base electrolyte. The working electrodes were fabricated by mixing 60% active materials (pristine Si, PS, and PS@C), 20% conductors (Denka black) and 20% binder (poly acrylic acid) onto a copper foil. The weight of the active materials in each electrode was around 1 mg. The coated electrodes were dried in a vacuum oven at 80 °C for 24 h. The CR2032 coin-type test cells were assembled in a dry room. The galvanostatic discharge–charge cycling was carried out by MACCOR cycle tester in the potential range of 0.05 to 2.0 V at various current densities of 100, 200, 400, 800 and 4000 mA g−1. The electric current was calculated based on the amount of the active material, not including the weight of the additives in the electrode. Electrochemical impedance spectroscopy (EIS) experiments with the fully charged cells (up to 0.0 V) were performed using an impedance/gain-phase analyzer (Solartron SI 1260) equipped with an electrochemical interface (Solartron SI 1286). The AC amplitude was 5 mV over a frequency range of 1 mHz to 100 kHz. 3. Results and discussion The synthesis process of PS@C is schematically presented in Fig. 1. Commercially available Si with an average particle size of 5 μm was used as raw material. The pristine Si for synthesis of micro and nano hybrid material was then added into the etching solution consisting of HF and AgNO3. In a galvanic displacement process, the reduction of oxidizing metal ions (Ag) to metallic species, and the local dissolution of bulk Si powder occurred spontaneously on a silicon surface. The Si atoms act as a reducing agent for the metal ions in an aqueous HF solution. After the etching process, the color of pristine Si changed from dark gray to brown due to the formation of nanostructures (PS) with lower refractive index than that pristine Si. Carbon coating of the PS obtained from metal-assisted chemical etching is carried out via thermal CVD using gaseous polypropylene at 700 °C for 10 h. After the coating process the color changed from brown to black due to the carbon layer. The synthesis based on metal-assisted chemical etching is quick and easy, making it a cost efficient process to manufacture nanostructures with large surface areas at low temperature. Moreover, the produced quantities can be easily industrialized and are ultimately limited only by the dimensions of the vessel. Furthermore, carbon layers are well known for the formation of stable SEI layers on coated surfaces resulting in enhanced electrochemical performance [26,27]. Fig. 2 shows the SEM images of pristine bulk Si, surface-modified PS, and PS@C. Pristine bulk Si powders feature particle sizes of 5 μm and

Fig. 1. A schematic illustration showing the synthetic process of PS@C.

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Fig. 2. (a) SEM images of (a) pristine bulk Si, (b) PS (low magnification), (c) PS (high magnification), (d) PS@C (low magnification), and (e) PS@C (high magnification).

smooth surfaces of polygonal Si (Fig. 2a). After the etching process, morphologies of PS are quite different compared with bulk Si. It exhibits a clearly porous Si nanostructure with a prickle shape with empty space (Fig. 2b and c). As shown in Fig. 2d and e, the voids are slightly decreased after the carbon coating on the PS. In the TEM images depicted in Fig. 3a and b of PS and PS@C the different structures with and without carbon coating are visible. TEM analysis indicates that carbon is uniformly formed along the surface of PS with thickness of 5–10 nm (Fig. 3d). Additionally, line scanning analysis confirmed that the outer layer mainly consists of carbon decomposed from C3H6 gas. These results demonstrate that the thermal CVD process is an effective route for uniform coating of the PS surface on a nanometer scale. Fig. 4 shows nitrogen adsorption–desorption isotherms and BarrettJoyner–Halenda (BJH) pore size distributions of pristine Si, PS, and PS@ C. The adsorption–desorption isotherms of PS and PS@C exhibit typical IV shapes, indicating their mesoporous characteristics. Pore size distribution (PSD) of PS and PS@C with mesoporous characteristics is concentrated on a value of 30–50 nm. The BET surface areas and total pore volume of pristine Si, PS, and PS@C are calculated to values of 1.2, 14.5, 10.9 m2 g−1, and 0.001, 0.038, 0.026 cm3 g−1, respectively. The BET surface area and total pore volume of PS produced by metalassisted chemical etching is increased compared to pristine Si due to the empty volume. Furthermore, the pore size distribution (7 ± 2.8 nm),

BET surface area and total pore volume of PS@C are significantly decreased, which indicates that those pores are partially filled with a carbon layer. Fig. 5a shows XRD results in accordance with the synthesis process of PS@C. The intensities of pristine Si are clearly observed at scattering angles (2θ) of about 28.4, 47.4, 56.2, 69.2, 76.5, 88.2° which are ascribed to the (111), (220), (311), (400), (331), and (422) planes of Si crystallites, respectively. Although any metal should be removed by HNO3 etching, the still remaining Ag is oxidized resulting in Ag2O formation. Schön et al. [28] have shown that Ag2O is converted to metallic Ag due to variation of the oxygen state at temperatures above 400 °C. During the carbon coating process the temperature is increased to 700 °C; thus, the remaining impurity converts from Ag2O to metallic Ag (Fig. 5c). More specifically, five crystalline peaks are observed in the diffraction patterns at 2θ° of 38.1, 44.3, 64.5, 77.5, and 81.7° which can be attributed to the (111), (200), (220), (311), and (222) crystallographic planes of face-centered cubic silver crystals, respectively. However, no peaks corresponding to carbon could be detected, which suggests that the carbon obtained by thermal CVD is amorphous. To identify the chemical state of Ag, XPS measurements were performed. Fig. 5b shows that high resolution XPS spectra in the Ag3d-region from PS@C. Two signals can be observed at 372.5 and 366.5 eV, originating from Ag 3d3/2 and 3d5/2, respectively. In the literature, the XPS spectrum includes Ag\Ag and Ag\Si bonds at low

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Fig. 3. TEM images of (a) PS and (b) PS@C, (c) STEM-HAADF image and (d) corresponding EDX line scanning profiles of Si (purple), O (blue), and C (orange) obtained from BSNR@C along the red line.

concentration of AgNO3. However, an Ag\O bond is observed with increasing AgNO3 concentration due to the high Ag layer thickness [29]. In the PS@C spectrum, no Ag\O peaks are observed, as shown in Fig. 5b. This demonstrated that the amount of Ag on the Si surface is very small (atomic concentration: 0.82%, relative Ag/Si ratio: 0.018), as shown in Table 1. Moreover, an F impurity of 1.67% is observed in the PS@C due to contamination caused by the HF solution. The residual Ag metal in the PS@C particle is important because it definitely reacts with Li+ ion. However, it is therefore assumed that the low content of Ag metal has no significant effect on the electrochemical reaction. To identify the actual carbon contents, EA analysis was carried out by the CHN method. The amounts of carbon in pristine Si, PS, and PS@C are measured to about 0.03, 0.12, and 12 wt.%, respectively. To further investigate the carbon structure in the PS@C samples, Raman spectroscopy with using a laser wavelength of 632 nm was performed on pristine Si, PS, and PS@C samples, as shown in Fig. 4c and d. The well-known Si\Si vibration appearing at ~510 (first-ordered peak) and ~950 cm−1 (second-ordered peak) was observed [30,31]. After etching and carbon coating, the Si peak

intensity is decreased by reduced scattering due to nano-structuring and masking effects. Both curves in the PS@C sample exhibit two broad bands at 1350 and 1580 cm−1 corresponding to the D and G band peaks for carbon, respectively. The G-band is related to the vibration of sp2-bonded carbon atoms in a two-dimensional hexagonal lattice. Whereas, the D-band either results from vibrations of carbon atoms with dangling bonds in crystal lattice plane terminations of disordered carbon, or from defects in the curved graphene sheets. The intensity ratio of the D-band to the G-band, ID/IG, gives an approximate evaluation of graphitic quality of carbon materials. For the carbon coating of the PS@C material the ratio of the D band to the G band was measured to be 1.5, indicating an amorphous carbon structure. This result is also in good accordance with the XRD measurements. Fig. 6a shows the voltage profiles of pristine Si, PS, and PS@C at a current density of 400 mA g−1 in the potential range from 0.05 to 2 V. The different behavior of the first and second cycles at all samples is originated from the phase change from crystalline to amorphous during Liion insertion and extraction. The first and second discharge–charge

Fig. 4. (a) Nitrogen adsorption–desorption isotherms, (b) BJH pore size distributions of pristine Si, PS, and PS@C. The solid and open symbols are adsorption and desorption curve, respectively.

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Fig. 5. Characterization of pristine Si, PS, and PS@C. (a) XRD of pristine Si, PS, and PS@C, (b) high resolution XPS spectra in the region of Ag3d obtained from PS@C, (c) Raman spectra of Si, PS, and PS@C in the wide spectra range, and (d) Raman spectrum showing the D and G band obtained carbon layer. The spectrum was fitted by 95% Gaussian–5% Lorenz function.

capacities of pristine Si are 2748, 2523, and 2583, 2504 mA h g−1, exhibiting coulombic efficiencies of 91 and 96%, respectively. In comparison with pristine Si, discharge–charge capacities of PS@C electrodes show a lower capacity due to the 12 wt.% carbon layer. However, the charge capacity of pristine bulk Si rapidly diminished after the second cycle. After 20 cycles, no significant capacity was detectable for the pristine Si electrode (Fig. 6b). This can be attributed to the cracking of bulk Si because of physical stress during insertion and extraction of Li+ ions. At the same time the coulombic efficiencies of the pristine Si electrode declined to 82% (Fig. 6d). Solid electrolyte interphase (SEI) formation for pure Si generally occurs in a wide voltage range, from 0.8 to 0.4 V. When cracks occur during cycling bare silicon surfaces are exposed to the electrolyte leading to more SEI formation. In this case, the irreversible reaction consuming active lithium and solvents of electrolyte occurs continuously during cycling due to the unstable SEI layer and is not limited to the first cycle. The slopes of sharp decrease and increase in the coulombic efficiencies and accumulated irreversible capacities suggested that the structure has been completely broken after 13 cycles (Fig. 6d). More interestingly, the charge capacities (1028 mA h g− 1 after 100 cycles) of PS electrode show relatively good cycle stability in comparison with bulk Si, because the prickle-like nanostructures with empty space prepared by metal-assisted chemical etching can mitigate the volume expansion. The normalized capacity retention (based on initial charge capacity) of PS@C (65%) electrode is higher than that of pristine Si (2%) and PS (47%) electrodes after 100 cycles (Fig. 6c). Table 1 Atomic concentration of various elements in PS@C. C1s (%)

N1s (%)

O1s (%)

F1s (%)

Si2P (%)

Ag3d (%)

Ag/Si ratio

13.88

0

40.05

1.67

43.58

0.82

0.018

As shown in Fig. 7, the rate capabilities of pristine, PS, and PS@C are investigated as a function of the cycle number at various current densities of 100, 200, 400, 800, 4000 and 100 mA g−1. The specific capacities of pristine Si rapidly decreased after few cycles at a low current density of 100 mA g−1 due to large volume change. However, the rate capability of PS is better than that of pristine Si because the porous structures play an important role in determining the transport of Li-ions and the electrode/electrolyte interface directly affects the electrochemical performance. The specific capacities of PS electrode without carbon layer are decreased to 300 mA h g− 1 at high current densities of 4 A g−1 and the capacities did not recover after decreasing the current density back to 100 mA g − 1. More interestingly, the data reveals that the PS@C electrode with conformal carbon layer retains a specific capacity over 1330 mA h g−1 after 25 cycles at a high current density of 4 A g−1, and the specific capacity of the PS@C electrode is recovered up to 1980 mA h g−1 at a current density of 100 mA g−1. Impedance is considered to be one of the key factors which reflected cell performance because it causes the polarization of cells in operation and has a vigorous effect on capacity. Impedance spectroscopy technique was employed in order to identify the resistance, and the resulting Nyquist plots were investigated. The imaginary part of the impedance is plotted as a function of the real part over a wide range of frequencies, as shown in Fig. 8. Impedance spectra were obtained from pristine Si, PS, and PS@C after rate cycling in a frequency range of 1 mHz to 100 kHz, and the equivalent circuits utilized in the impedance analyses are given in the inset of Fig. 8; where RE is the resistance of the bulk electrolyte, CPESEI is the dispersed capacitance of the SEI layer, and RSEI is the resistance for Li+ conduction in the SEI layer. In addition, Rct, CPEdl, and ZW represent the charge transfer resistance, the dispersed capacitance of double layer at the electrode surface, and the Warburg impedance, respectively. It is generally known that

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Fig. 6. (a) Voltage profiles (solid line: 1st cycle, dot line: 2nd cycle), (b) charge capacities, (c) normalized capacity retention (initial charge capacity is equal to 100%), and (d) accumulated irreversible capacities of pristine Si, PS, and PS@C at a current density of 400 mA g−1.

the semicircles formed in the high- and middle-frequency region reflect the interfacial characteristics of the electrodes, whereas the straight line in the low-frequency region represents the Warburg diffusion of lithium ions in the electrochemical cells containing silicon electrodes. As shown in Fig. 8, it is clear that the semicircle diameter of PS@C corresponding to 115 Ω is much smaller than those of the pristine Si and PS corresponding to 385 and 295 Ω. The increased charge transfer resistance of pristine Si and PS indicates a reduction of the active site area over lithium insertion and extraction cycling process. This decrease in the active site area is probably due to the reduced boundaries among

Fig. 7. Rate performances of pristine Si, PS, and PS@C at various current densities of 100, 200, 400, 800, 4000, and 100 mA g−1 (cut-off voltage: 0.05 to 2 V).

the active material, current collector and electrolyte. Metal-assisted etching and carbon coating improve charge transfer at the current collector/electrode/electrolyte interface. It is suggested that combination of metal-assisted etching and amorphous carbon coating is an effective method to improve the cycle stability and rate capability of Si anode for lithium-ion batteries. 4. Conclusions Micro and nano hybrid material of conformal carbon coated pricklelike Si powders were successfully synthesized via metal-assisted chemical etching and thermal chemical vapor deposition and then employed

Fig. 8. Nyquist plots of pristine Si, PS and PS@C cells after rate cycling.

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as an anode material for lithium-ion batteries. In comparison to bulk Si, the prickle-like structure with empty voids prepared by chemical etching is effective in accommodating the large volume changes, resulting in improved electrochemical performance, i.e. cycle stability and rate capability. Surprisingly, the amorphous carbon coating layer with an ID/IG ratio of 1.5 on the unique prickle-like Si powders acts as a buffer layer to suppress the volume change and prevents the Si surface from direct contact with the electrolyte and thus increases the structural stability during the insertion and extraction of Li+. In particular, the carbon layer improves charge transfer resistance, as well as the kinetics of lithiation and delithiation resulting in superior performance at a high current densities of 4 A g−1. The results obtained in this study suggest that the introduction of metal-assisted etching and conformal carbon coating on the Si interface could be an effective way to enhance the cycle stability and rate capability of prickle-like Si. Acknowledgment This work was supported by KIST institutional program and research grants by the National Research Foundation under Ministry of Science, ICT & Future, Korea (NRF-2012M1A2A2671792). References [1] B.A. Boukamp, G.C. Lesh, R.A. Huggins, J. Electrochem. Soc. 128 (1981) 725–729. [2] D. Larcher, S. Beattie, M. Morcrette, K. Edström, J.-C. Jumas, J.-M. Tarascon, J. Mater. Chem. 17 (2007) 3759–3772. [3] N.-S. Choi, K.H. Yew, K.Y. Lee, M. Sung, H. Kim, S.-S. Kim, J. Power Sources 161 (2006) 1254–1259. [4] H. Yoo, J.-I. Lee, H. Kim, J.-P. Lee, J. Cho, S. Park, Nano Lett. 11 (2011) 4324–4328. [5] A. Magasinki, P. Dixon, B. Hertzberg, A. Kvit, J. Ayala, G. Yushin, Nat. Mater. 9 (2010) 353–358. [6] Y.-M. Lin, K.C. Klavetter, P.R. Abel, N.C. Davy, J.L. Snider, A. Heller, C.B. Mullins, Chem. Commun. 48 (2012) 7268–7270.

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