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Porous Silicon–Carbon Composite Materials Engineered by Simultaneous Alkaline Etching for High-Capacity Lithium Storage Anodes
Electrochimica Acta 196 (2016) 197–205
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Porous Silicon–Carbon Composite Materials Engineered by Simultaneous Alkaline Etching for High-Capacity Lithium Storage Anodes Myungbeom Sohna , Dae Sik Kima , Hyeong-Il Parka , Jae-Hun Kimb , Hansu Kima,* a b
Department of Energy Engineering, Hanyang University, Seoul 04763, Republic of Korea School of Advanced Materials Engineering, Kookmin University, Seoul 02707, Republic of Korea
A R T I C L E I N F O
A B S T R A C T
Article history: Received 11 December 2015 Received in revised form 11 February 2016 Accepted 16 February 2016 Available online 20 February 2016
Porous silicon–carbon (Si–C) composite materials have attracted a great deal of attention as highperformance anode materials for Li-ion batteries (LIBs), but their use suffers from the complex and limited synthetic routes for their preparation. Herein we demonstrate a scalable and nontoxic method to synthesize porous Si–C composite materials by means of simultaneous chemical etching of Si and carbon phases using alkaline solution. The resulting porous Si–C composite material showed greatly improved cycle performance, good rate capability, and high dimensional stability during cycling. Porous Si–C electrode showed an expansion of the height by about 22% after the first lithiation and only 16% after the first cycle. The material synthesis concept and scalable simultaneous etching approach presented here represent a means of improving the electrochemical properties of Si-based porous anode materials for use in commercial LIBs. ã 2016 Elsevier Ltd. All rights reserved.
Keywords: anode porous composite Si–C composite alkaline etching lithium-ion battery
1. Introduction Lithium-ion batteries (LIBs) have played important role in the development of portable electronics and have also gained a great deal of attention as power sources for emerging applications such as electric vehicles and large-scale energy storage systems [1–3]. However, in order to meet recent requirements for aforementioned emerging applications, the electrochemical performance of LIBs needs to be further advanced in terms of energy density and power density [4–6]. In an attempt to increase the energy density of LIBs, development of new electrode materials with higher capacity than currently used materials such as LiCoO2 and graphite has been pursued during the past several decades [7–9]. Silicon has been regarded as a particularly promising candidate anode material to replace graphite, because its capacity (3579 mAh g 1) is about ten times higher than that of graphite (372 mAh g 1) [10,11]. However, the practical use of Si has been limited by its extremely large volume changes during alloying and dealloying reactions with Li, leading to poor capacity retention of Si electrodes [12–14]. Various strategies have been suggested to solve these problems of Si anodes, and some of them showed improved capacity retention by managing the volume changes of Si during
* Corresponding author. Tel.: +82 2 2220 2414; fax: +82 2 2220 2489. E-mail address: [email protected] (H. Kim). http://dx.doi.org/10.1016/j.electacta.2016.02.101 0013-4686/ ã 2016 Elsevier Ltd. All rights reserved.
cycling [15]. Among these promising Si-based anode materials, Si– C composites are being regarded as highly promising anode materials because of their stable cycle performance and relatively low production cost [16–19]. Porous Si-based materials have also showed some promising behaviors, because pores in the material could effectively accommodate the mechanical strain induced by alloying and dealloying reactions of Si with Li, thereby minimizing mechanical degradation of the resulting Si-based electrodes. In this regard, many approaches to prepare porous Si-based materials have been investigated, and some examples of porous Si-based materials showed good cycle performance comparable to that of currently used graphite anode materials for LIBs [20–22]. However, most of the approaches to prepare porous Si-based anode materials require either time-consuming and expensive processing or the use of hazardous chemicals such as hydrofluoric acid and silane gas, making production process expensive [23,24]. Given that the production cost of LIBs is one of the most challenging issues for their successful use in emerging applications for electric vehicles and large-scale energy storage systems [25], further addressing this technical issues of porous Si-based anode materials is necessary to ensure their commercial success. Toward a scalable and cost-effective process to produce porous Si-based anode materials, here we demonstrate a facile route for the preparation of porous Si–C composites by means of simple ball milling followed by simultaneous one-pot chemical etching of Si
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and carbon phases in the Si–C composite using alkaline (NaOH) aqueous solution. By using this method in the present work, we could obtain porous Si–C composites with the pores inside the composite as well as at the surface of the composite, because alkaline solution has the capability to etch both carbonaceous materials and Si [26–30]. Pan et al. synthesized a york–shell Si–C nanocomposite by forming a void between the Si nanoparticle and the carbon shell using NaOH etching of the Si nanoparticle [31]. Instead of embedding Si particles in hollow carbon shell, it has been recently revealed that the structure of porous carbon surrounding nanoscale Si can accommodate the large volume
changes of Si during lithiation and delithiation, thus showing outstanding cycle performance [32,33]. However, the methods used thus far to generate pores in the carbon phase are complicated or involve toxic chemicals. Considering that both ball milling and alkaline etching are environmentally benign and already established in the industrial processes, our simple alkaline etching combined with ball-milling process would be another promising synthetic routes for porous Si–C composite anode materials for next-generation LIBs. In this work, we describe this method in detail and demonstrate its promise for the preparation of porous Si–C composite anode materials for use in next-generation LIBs.
Fig. 1. Schematic illustration of synthesizing (a) non-porous Si–C composite through annealing (heating and cooling) and (b) porous Si–C composite by etching non-porous Si–C in alkaline solution for times of T1 (105 min) and T1 + T2 (115 min). Powder SEM images of (c) ball-milled Si, (d) non-porous Si–C, and (e) porous Si–C composite (T1 + T2). Cross-sectional SEM images of (f) non-porous Si–C, (g) less porous Si–C (T1), and (h) porous Si–C (T1 + T2).
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2. Experimental 2.1. Preparation of Porous Si–C Composites To reduce particle size, bulk Si powder ( 325 mesh, 99% trace metals basis, Sigma-Aldrich) was ball-milled for 36 h at 1000 rpm using ball-to-powder ratio of 20 under an Ar atmosphere using a custom-made vibratory mill. The ball-milled Si and pitch were hand-mixed with the weight ratio of 7:3 using a mortar and pestle. The mixture was annealed in a vertical furnace under an N2 atmosphere at 800 C for 2 h with a heating rate of 5 C min 1 and then allowed to cool naturally to room temperature. Non-porous Si–C powder was obtained by grinding the annealed sample using a mortar and pestle. Porous Si–C composites were prepared by ultrasonication (Branson CPX3800H-E) of a non-porous Si–C composite together with 0.5 mol L 1 NaOH solution in a highdensity polyethylene (HDPE) bottle for either 105 or 115 min at room temperature; the etching reaction was then quenched by adding 0.2 M HCl solution.
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conductive agent (Super-P, 5 wt%), and polyacrylic acid binder (PAA, 15 wt%) in deionized water. After coating, the electrodes were dried at 120 C under vacuum conditions for 12 h and then allowed to cool naturally to room temperature. Electrochemical properties were analyzed using CR2032 coin-type half-cells assembled in a dry room. Each coin cell consisted of a porous polyethylene (PE) membrane as the separator and 1 mol L 1 LiPF6 solution dissolved in a mixed solvent of ethylene carbonate (EC) and diethyl carbonate (DEC) (3:7 v/v.) with fluoroethylene carbonate (5%) (PANAX Etec Co., Ltd.). The cycle life of cells was tested in the voltage range of 0.01 to 1.5 V vs. Li/Li+ at a current density of 100 mA g 1. The rate capability test and first cycle charge– discharge of cells for measuring electrode thickness were carried out in the voltage range 0.005 to 1.5 V vs. Li/Li+ at current densities of various values and 100 mA g 1, respectively. Electrochemical impedance spectroscopy (EIS) measurement was performed using a potentiostat (VSP-300, BioLogic) over a frequency range of 1 MHz to 10 mHz with an amplitude of 5 mV. 3. Results and Discussion
2.2. Materials Characterizations Surface and cross-sectional morphologies of the samples were observed by field emission scanning electron microscopy (FESEM, JEOL JSM-7000F). Surface areas and pore sizes were measured from Brunauer–Emmett–Teller (BET) and Barrett–Joyner–Halenda (BJH) results by using a surface analyzer (Micromeritics 3Flex). The carbon content in the non-porous and porous Si–C composites was calculated by means of thermogravimetric analysis (TGA, TA Instruments SDT Q600) under air. Raman spectroscopy (JASCO NRS-3100) was carried out to reveal the carbon phase in the Si–C composites. 3.3. Electrochemical Measurements Electrodes were prepared by coating slurries on 10 mm thickness Cu foil. The slurries were made by mixing non-porous or porous Si–C composite as the active material (80 wt%), a
A non-porous Si–C composite material was prepared by typical process, i.e., mixing ball-milled Si microparticles with pitch as a carbon precursor (the left image of Fig. 1a), followed by carbonization of pitch under a reducing atmosphere (the center and right images of Fig. 1a) [34]. As schematically depicted in Fig. 1b, pores were formed in the non-porous Si–C composite via simultaneous alkaline (NaOH) etching process. The porous Si–C composites with etching time T1 (105 min) and T1 + T2 (115 min) are denoted as less porous Si–C and porous Si–C composite material, respectively. The T2 is additional etching time (10 min) after time T1. Fig. 1c, d, and e present scanning electron microscopy (SEM) images of ball-milled Si, non-porous Si–C, and porous Si–C composite particles, respectively. Aggregation of ball-milled Si particles in carbon framework is observed in Fig. 1d, because the Si–C composite particles are much larger than the ball-milled Si particles. Pore generation inside the Si–C composite particles can be inferred from the maintenance of their dimensions on the order
Fig. 2. SEM–EDS analysis with elemental mapping of (a) non-porous, (b) less porous, and (c) porous Si–C composites.
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Fig. 3. (a) N2 adsorption–desorption isotherms and (b) pore size distributions of non-porous and porous Si–C composites.
of tens of micrometers after NaOH treatment as shown in Fig. 1e. Compared to the ball-milled Si particles, which have uneven surfaces (Fig. 1c inset), the non-porous Si–C composite particles exhibit smoother surface (Fig. 1d inset), indicating that the ballmilled Si particles are well surrounded by carbon after carbonization of the pitch. We found that the NaOH etching resulted in a roughened surface of the Si–C composite particles (Fig. 1e inset), suggesting that pores were created in the carbon phase. The nonporous Si–C composite particles have a typical microstructure in which micron-sized Si particles were dispersed in the carbon matrix (the left image of Fig. 1b and Fig. 1f). The micrometer-scale pores were formed by the chemical etching of some Si microparticles inside non-porous Si–C composite particles as shown in cross-sectional images of porous Si–C composite particles (Figs. 1g, h and 2 ). Fig. 3 shows N2 adsorption/desorption isotherms and pore size distribution of Si–C composites estimated by using Brunauer– Emmett–Teller (BET) and Barrett–Joyner–Halenda (BJH) methods. While the as-prepared non-porous Si–C composite shows typical non-porous N2 adsorption/desorption isotherms, the chemically etched Si–C composites exhibit a hysteresis loop with steep
condensation steps in the range of relative pressure, P/P0 = 0.80– 0.98, which is a typical characteristic of porous materials with disordered pore structure (Fig. 3a). Note that these isotherms show a low pressure hysteresis in which the adsorption and desorption branches have no contact points at low P/P0, which might result from the micropores in the composite materials, which have almost same size to that of adsorbates (N2) [35]. It is well known that such a narrow opening of the small pores could restrict the access of the adsorbates [36]. Considering that this hysteresis were commonly observed even at the BET curves for non-etched Si–C composite material (Fig. 3a), micropores in the composites might come from the carbon phase in the composite after carbonization of pitch at 800 C. The simultaneous chemical etching process also enlarged the surface area and created the abundant pores with the size of 50–250 nm in the composite materials (Fig. 3b). In the case of 115 min NaOH treatment, the BET surface area of the Si–C composite increased from about 76.8 to about 123.4 m2 g 1 by simultaneous chemical etching of Si microparticles and carbonaceous phase (Table 1). We also found that the pore volume of porous Si–C composite, as calculated from desorption branch using the BJH method, was dramatically increased from 2.2 10 3 to
Table 1 BET and BJH data of non-porous, less porous, and porous Si–C composites.
BET surface area BJH adsorption cumulative surface area of pores between 17 Å and 3,000 Å diameter BJH desorption cumulative surface area of pores between 17 Å and 3,000 Å diameter BJH adsorption cumulative volume of pores between 17 Å and 3,000 Å diameter BJH desorption cumulative volume of pores between 17 Å and 3,000 Å diameter BJH adsorption average pore diameter BJH desorption average pore diameter
Non-porous Si–C
Less Porous Si–C
Porous Si–C
76.7962 m2 g 1 0.108 m2 g 1 0.1149 m2 g 1 0.002372 cm3 g 0.002264 cm3 g 904.944 Å 788.469 Å
98.8482 m2 g 1 3.964 m2 g 1 5.3942 m2 g 1 0.046295 cm3 g 0.044438 cm3 g 467.208 Å 329.525 Å
123.4492 m2 g 1 9.299 m2 g 1 9.8568 m2 g 1 0.081894 cm3 g 1 0.084886 cm3 g 1 352.255 Å 344.478 Å
1 1
1 1
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Fig. 4. TGA curves of non-porous and porous Si–C composites.
80.5 10 3 cm3 g 1 as shown in Table 1. This enlarged surface area as well as increased pore volume would be helpful in accommodating the huge volume changes of Si upon lithiation and delithiation during cycling [37–39]. In order to determine the Si to carbon mass ratio in the Si–C composites, thermogravimetric analysis (TGA) was carried out. Fig. 4 shows TGA curves of the Si–C composites before and after chemical etching using NaOH solution. Carbon content in the nonporous Si–C composite is estimated to be about 22 wt%, while those of porous Si–C composites are 44 and 50 wt% after chemical etching for 105 and 115 min, respectively. These results suggest that NaOH solution is more effective in etching out Si phase than in etching the carbon matrix, which is well matched with the crosssectional observation shown in Fig. 1 h. The existence of carbon phases is revealed by D (1350 cm 1) and G (1580 cm 1) bands in Raman spectra collected at three arbitrary points in each Si–C composites (Fig. S1). Fig. 5a shows the voltage profiles of non-porous and porous Si– C composite electrodes for the first cycle. The non-porous Si–C composite electrode delivered a reversible capacity of about 2471 mAh g 1 with an initial efficiency of about 90%, while the less porous and porous Si–C composite electrodes exhibited the discharge capacities of 1438 and 1077 mAh g 1, respectively. The decrease in the reversible capacity of porous Si–C composite
electrodes is mainly attributed to the relatively lower Si contents in the composites as observed from voltage profiles (Fig. 5a) and TGA curves (Fig. 4). Note that the initial efficiencies of less porous (81%) and porous (74%) Si–C composite electrodes were also lower than that of non-porous Si–C composite electrode. This increase in the irreversible electrochemical reaction observed at porous Si–C composite electrodes is probably due to the enlarged surface area caused by chemical etching, which might increase the electrolyte decomposition reaction on the surface of the electrode during the first cycle [40,41]. However, the porous Si–C composite electrode showed a dramatic improvement in the cycle performance compared with that of non-porous Si–C composite electrode. The porous Si–C composite electrode maintained about 74% of their initial capacity even after 100 cycles, while non-porous Si–C composite electrode shows a fast capacity fading (8% of the initial capacity) within 50 cycles, as revealed in Fig. 5b. Coulombic efficiencies of Si–C composite electrodes showed the same tendency; the non-porous Si–C composite electrode showed poor Coulombic efficiencies less than 95%, while porous Si–C composite electrodes showed Coulombic efficiencies more than 98% during most cycles (Fig. 5c). In particular, the porous Si–C composite electrode obtained after chemical etching for 115 min maintained high Coulombic efficiencies without significant fading even after 100 cycles. As might be expected, the increase in both the surface
Fig. 5. Electrochemical properties of non-porous and porous Si–C composites. (a) Voltage profiles, (b) cycle performance, (c) Coulombic efficiencies of non-porous and porous Si–C composites, and (d) rate capability of porous Si–C composite only.
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Fig. 6. Cross-sectional FESEM images of (a) non-porous and (b) porous Si–C electrodes during the first cycle.
area and pore volume of Si–C composite materials by simultaneous chemical etching using NaOH solution led to much improved rate capability in the porous Si–C composite electrode (Fig. 5d) compared to that of non-porous composite electrode (Fig. S2). Although the gravimetric energy density of non-porous Si–C composite electrode is almost double that of the porous composite electrode at the current density of 0.1 A g 1 (Fig. S2), the capacity of non-porous composite electrode ( < 1000 mAh g 1) is lower than that of porous composite electrode (1100 mAh g 1) at 1 A g 1 (Fig. 5d). Porous Si–C composite prepared by simultaneous chemical etching route showed better rate capability compared with previously reported Si–C composite anode materials for lithium ion batteries. For example, yolk-shell structured Si–C nanocomposite showed a high reversible capacity of 2258 mAh g 1 at a rate of 0.1 A g 1, but this material showed a drastic capacity decrease with an increase of the current densities (about 250 mAh g 1 at a rate of 5 A g 1) [31]. Interconnected porous Si–C composites by using the Rochow reaction and hydrofluoric etching had low reversible capacity of 856.5 mAh g 1 at a rate of 50 mA g 1 and also showed poor rate capability at a high rate condition (440 mAh g 1 at a rate of 1 A g 1) [38]. However, the porous Si–C composite in this work showed a high rate performance with reversible capacity of about 1180 and 990 mAh g 1 at a rate of 0.1 and 5 A g 1, respectively. This improvement in the rate capability can be attributed to the pore created by simultaneous chemical etching of Si and carbon phase in the composite. The abundant pores formed by alkaline etching would not only allow the electrolyte to penetrate into the composite, but would also increase the interfacial area between the electrolyte and the electrode material. Besides providing high Li+ accessibility, carbon frames basically could offer the pathways for electron transport. As a result, the porous Si–C composite electrodes show much improved electrochemical kinetics compared with nonporous one. We anticipated that the pores in the Si–C composite electrode could also improve the dimensional stability of the electrode against huge volume changes of Si phase in the composite during cycling. In order to confirm this effect of the pores in the Si–C composite electrodes, the dimensional changes in the non-porous and porous composite electrodes were investigated using FESEM during the initial cycle. Fig. 6 shows cross-sectional FESEM images of non-porous Si–C composite electrode and porous Si–C composite electrode after chemical etching for 115 min after lithium insertion and removal. After fully lithiated, the non-porous Si–C composite electrode expanded up to about 310% of its pristine state (Fig. 6a and Table S1). On the other hand, the porous Si–C composite electrode showed much improved dimensional stability
during the first cycle. Even after full lithiation, the porous Si–C composite electrode expands to only 122% of its thickness of pristine electrode (Fig. 6b and Table S1). The average thickness of the porous Si–C electrode after delithiation is 40.7 mm which is slightly thicker (116%) than that of pristine electrode (35.1 mm) (Fig. S3b and Table S1). In contrast, meaningful thickness cannot be measured in the delithiated non-porous Si–C electrode because the electrode contents are sparse on the Cu foil (Fig. S3a and Table S1). This severe detachment of non-porous Si–C particles from the foil might be due to the weakened adhesion between active material and current collector by the huge volume change of Si during lithiation and delithiation. The loading of active material on the non-porous Si–C electrode (2.69 mg cm 2) is greater than that of porous Si–C electrode (1.86 mg cm 2) with similar pristine thickness (Table S1), suggesting that the density of porous Si–C material is less than that of non-porous Si–C material caused by pores generated inside the Si–C composite via NaOH etching treatment. The effects of pores are remarkable when the nonporous and porous Si–C electrodes are compared from the volumetric perspective. Although the non-porous Si–C has a gravimetric discharge capacity (2482 mAh g 1) twice that of porous Si–C (1242 mAh g 1) composite electrode, the difference between their volumetric electrode capacities based on the thickness of lithiated electrodes is only about 233 mAh cm 3 (Table S1). This indicates that the preformed pores in the Si–C composite alleviated the stresses induced by large volume changes of Si during charge and discharge. Closer observation on the cycled electrodes shows more distinctive differences between the porous and non-porous Si–C composites. Fig. 7a and b compare the cross-sectional FESEM images of the non-porous and porous Si–C composite electrode with respect to the depth of discharge. Pores observed in the pristine porous Si–C composite disappeared after full lithiation and appeared again after full delithiation, clearly showing the reversible behavior of the pores in the porous Si–C composite material (Fig. 7b and Fig. S4b). On the other hand, non-porous Si–C composite electrode showed internal cracks in the material after lithiation as shown in Fig. 7a and Fig. S4a. This difference in the morphological changes of the Si–C composite electrodes clearly demonstrates how pores in the composite play an active role in accommodating the volume changes of the Si phase during cycling. In order to further characterize the electrochemical properties of Si–C composite electrode before and after chemical etching, we carried out the EIS analysis on the electrodes after 100 cycles. The Nyquist plots of the non-porous and porous Si–C composite electrode showed two distinguishable semicircles in the high frequency and middle frequency region (Fig. 8). The oblique
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Fig. 7. High-magnification cross-sectional FESEM images of (a) non-porous and (b) porous Si–C electrodes during the first cycle.
straight lines in the low frequency region were also observed in all the Nyquist plots of the electrodes. The semicircle in the high frequency region is mainly attributed to the SEI film on the electrode, and the semicircle observed in the middle frequency region is closely related to the charge transfer resistance of active material, and the straight line in the low frequency region is related with the mass transfer, such as lithium ion diffusion in the electrode [42–44]. As shown in Fig. 8, the non-porous electrode showed a distinctive semicircle at the high frequency region which is related to the SEI formation, while both the less porous and
porous electrode showed a very small semicircle in the high frequency region after 100 cycles. This suggests that the less SEI layers were formed in the porous electrodes than in the nonporous electrode. Moreover, the porous Si–C composite electrode showed smaller semicircle in the middle frequency than that of non-porous Si–C composite electrode after 100 cycles, indicating that the porous electrode had better electronic contact between electrode components. This lower impedance could be attributed to high structural integrity in the porous electrode during cycling as the pores could alleviate the stress from the huge volume
Fig. 8. Nyquist plots of non-porous, less-porous, and porous Si–C electrodes after 100 cycles.
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change of Si. These improved electrochemical properties result in enhanced cycle performance as well as rate capability of the porous Si–C composite electrode as shown in Fig. 5b, c, and d. 4. Conclusions We demonstrated a cost-effective simultaneous chemical etching route for porous Si–C composite in which abundant pores as well as micrometer scale Si particles are incorporated into carbon matrix. This mechanically robust design of porous Si–C composites provided highly reversible Li storage over prolonged cycles and good dimensional stability upon alloying and dealloying reaction of Si phase with Li. In particular, the porous Si–C composite obtained from simultaneous etching of Si and carbon phase exhibited a high reversible capacity of more than 1000 mAh g 1 with an initial coulombic efficiency of 74% and excellent cycle performance up to 100 cycles. The highly enhanced battery performances of the proposed Si–C composite was mainly attributed to the pores in the composite, which act as a dynamic buffer phase to accommodate the volume change of the Si phase without significant degradation. This material concept and scalable simultaneous alkaline etching approach provide a means of improving the electrochemical properties of Si-based anode materials for use in commercial LIBs. Acknowledgements This work was in part supported by the Korea Evaluation Institute of Industrial Technology (KEIT), which is funded by the Ministry of Trade, Industry & Energy, Republic of Korea (No. 10046341) and in part supported by the research fund of Hanyang University (HY-2012-T). Appendix A. Supplementary data Supplementary data associated with this article can be found, in the online version, at http://dx.doi.org/10.1016/j. electacta.2016.02.101. References [1] M. Armand, J.-M. Tarascon, Building better batteries, Nature 451 (2008) 652– 657. [2] B. Dunn, H. Kamath, J.-M. Tarascon, Electrical energy storage for the grid: a battery of choices, Science 334 (2011) 928–935. [3] H. Kim, G. Jeong, Y.-U. Kim, J.-H. Kim, C.-M. Park, H.-J. Sohn, Metallic anodes for next generation secondary batteries, Chem. Soc. Rev. 42 (2013) 9011–9034. [4] J.B. Goodenough, Y. Kim, Challenges for rechargeable Li batteries, Chem. Mater. 22 (2009) 587–603. [5] J.B. Goodenough, K.-S. Park, The Li-ion rechargeable battery: a perspective, J. Am. Chem. Soc. 135 (2013) 1167–1176. [6] R. Van Noorden, A better battery, Nature 507 (2014) 26–28. [7] H. Li, Z. Wang, L. Chen, X. Huang, Research on advanced materials for Li-ion batteries, Adv. Mater. 21 (2009) 4593–4607. [8] B. Scrosati, J. Garche, Lithium batteries: status, prospects and future, J. Power Sources 195 (2010) 2419–2430. [9] S. Goriparti, E. Miele, F. De Angelis, E. Di Fabrizio, R.P. Zaccaria, C. Capiglia, Review on recent progress of nanostructured anode materials for Li-ion batteries, J. Power Sources 257 (2014) 421–443. [10] H. Liu, Z. Guo, J. Wang, K. Konstantinov, Si-based anode materials for lithium rechargeable batteries, J. Mater. Chem. 20 (2010) 10055–10057. [11] M. Obrovac, V. Chevrier, Alloy negative electrodes for Li-ion batteries, Chem. Rev. 114 (2014) 11444–11502. [12] J.H. Ryu, J.W. Kim, Y.-E. Sung, S.M. Oh, Failure modes of silicon powder negative electrode in lithium secondary batteries, Electrochem. Solid-State Lett. 7 (2004) A306–A309. [13] U. Kasavajjula, C. Wang, A.J. Appleby, Nano- and bulk- silicon-based insertion anodes for lithium-ion secondary cells, J. Power Sources 163 (2007) 1003– 1039. [14] D. Munao, J. Van Erven, M. Valvo, E. Garcia-Tamayo, E. Kelder, Role of the binder on the failure mechanism of Si nano-composite electrodes for Li-ion batteries, J. Power Sources 196 (2011) 6695–6702.
[15] H. Wu, Y. Cui, Designing aanostructured Si anodes for high energy lithium ion batteries, Nano Today 7 (2012) 414–429. [16] H. Kim, E.-J. Lee, Y.-K. Sun, Recent advances in the Si-based nanocomposite materials as high capacity anode materials for lithium ion batteries, Mater. Today 17 (2014) 285–297. [17] J. Guo, A. Sun, X. Chen, C. Wang, A. Manivannan, Cyclability study of siliconcarbon composite anodes for lithium-ion batteries using electrochemical impedance spectroscopy, Electrochim. Acta 56 (2011) 3981–3987. [18] R. Yi, F. Dai, M.L. Gordin, H. Sohn, D. Wang, Influence of silicon nanoscale building blocks size and carbon coating on the performance of micro-sized SiC composite Li-ion anodes, Adv. Energy Mater. 3 (2013) 1507–1515. [19] X. Su, Q. Wu, J. Li, X. Xiao, A. Lott, W. Lu, B.W. Sheldon, J. Wu, Silicon-based nanomaterials for lithium-ion batteries: a review, Adv. Energy Mater. 4 (2014) 1300882. [20] A. Magasinski, P. Dixon, B. Hertzberg, A. Kvit, J. Ayala, G. Yushin, Highperformance lithium-ion anodes using a hierarchical bottom-up approach, Nat. Mater. 9 (2010) 353–358. [21] D.S. Jung, T.H. Hwang, S.B. Park, J.W. Choi, Spray drying method for large-scale and high-performance silicon negative electrodes in Li-ion batteries, Nano Lett. 13 (2013) 2092–2097. [22] T. Jaumann, M. Herklotz, M. Klose, K. Pinkert, S. Oswald, J. Eckert, L. Giebeler, Tailoring hollow silicon-carbon nanocomposites as high-performance anodes in secondary lithium-based batteries through economical chemistry, Chem. Mater. 27 (2015) 37–43. [23] W. Wang, P.N. Kumta, Nanostructured hybrid silicon/carbon nanotube heterostructures: reversible high-capacity lithium-ion anodes, ACS Nano 4 (2010) 2233–2241. [24] C. Zhao, Q. Li, W. Wan, J. Li, J. Li, H. Zhou, D. Xu, Coaxial carbon-silicon-carbon nanotube arrays in porous anodic aluminum oxide templates as anodes for lithium ion batteries, J. Mater. Chem. 22 (2012) 12193–12197. [25] L. Lu, X. Han, J. Li, J. Hua, M. Ouyang, A review on the key issues for lithium-ion battery management in electric vehicles, J. Power Sources 226 (2013) 272–288. [26] S. Yumitori, Y. Nakanishi, Effect of anodic oxidation of coal tar pitch-based carbon fibre on adhesion in epoxy matrix: Part 2. comparative study of three alkaline solutions, Composites Part A 27 (1996) 1059–1066. [27] E. Raymundo-Piñero, P. Azaïs, T. Cacciaguerra, D. Cazorla-Amorós, A. LinaresSolano, F. Béguin, KOH and NaOH activation mechanisms of multiwalled carbon nanotubes with different structural organization, Carbon 43 (2005) 786–795. [28] H. Teng, T.-C. Weng, Transformation of mesophase pitch into different carbons by heat treatment and KOH etching, Micropor. Mesopor. Mater. 50 (2001) 53– 60. [29] H.-J. Ahn, J.-H. Lee, Y. Jeong, J.-H. Lee, C.-S. Chi, H.-J. Oh, Nanostructured carbon cloth electrode for desalination from aqueous solutions, Mater. Sci. Eng: A 449–451 (2007) 841–845. [30] H. Seidel, L. Csepregi, A. Heuberger, H. Baumgärtel, Anisotropic etching of crystalline silicon in alkaline solutions, J. Electrochem. Soc. 137 (1990) 3612– 3626. [31] L. Pan, H. Wang, D. Gao, S. Chen, L. Tan, L. Li, Facile synthesis of yolk–shell structured Si–C nanocomposites as anodes for lithium-ion batteries, Chem. Commun. 50 (2014) 5878–5880. [32] D. Shao, D. Tang, Y. Mai, L. Zhang, Nanostructured silicon/porous carbon spherical composite as a high capacity anode for Li-ion batteries, J. Mater. Chem. A 1 (2013) 15068–15075. [33] H.M. Jeong, S.Y. Lee, W.H. Shin, J.H. Kwon, A. Shakoor, T.H. Hwang, S.Y. Kim, B.-S. Kong, J.-S. Seo, Y.M. Lee, Silicon@porous nitrogen-doped carbon spheres through a bottom-up approach are highly robust lithium-ion battery anodes, RSC Adv. 2 (2012) 4311–4317. [34] Y.-X. Wang, S.-L. Chou, J.H. Kim, H.-K. Liu, S.-X. Dou, Nanocomposites of silicon and carbon derived from coal tar pitch: cheap anode materials for lithium-ion batteries with long cycle life and enhanced capacity, Electrochim. Acta 93 (2013) 213–221. [35] A.M. Puziy, O.I. Poddubnaya, A. Martínez-Alonso, F. Suárez-García, J.M.D. Tascón, Synthetic carbons activated with phosphoric acid II. Porous structure, Carbon 40 (2002) 1507–1519. [36] Y. Liao, J. Weber, C.F.J. Faul, Conjugated microporous polytriphenylamine networks, Chem. Commun. 50 (2014) 8002. [37] R. Zhang, Y. Du, D. Li, D. Shen, J. Yang, Z. Guo, H.K. Liu, A.A. Elzatahry, D. Zhao, Highly reversible and large lithium storage in mesoporous Si/C nanocomposite anodes with silicon nanoparticles embedded in a carbon framework, Adv. Mater. 26 (2014) 6749–6755. [38] Z. Zhang, Y. Wang, W. Ren, Q. Tan, Y. Chen, H. Li, Z. Zhong, F. Su, Scalable synthesis of interconnected porous silicon/carbon composites by the Rochow reaction as high-performance anodes of lithium ion batteries, Angew. Chem. Int. Ed. 53 (2014) 5165–5169. [39] X. Zhao, X. Rui, W. Zhou, L. Tan, Q. Yan, Z. Lu, H.H. Hng, Growth of Si nanowires in porous carbon with enhanced cycling stability for Li-ion storage, J. Power Sources 250 (2014) 160–165. [40] Z. Li, W. Wang, Z. Li, Z. Qin, J. Wang, Z. Liu, Bridging porous Si-C composites with conducting agents for improving battery cycle life, J. Power Sources 286 (2015) 534–539. [41] M.-S. Wang, W.-L. Song, J. Wang, L.-Z. Fan, Highly uniform silicon nanoparticle/ porous carbon nanofiber hybrids towards free-standing high-performance anodes for lithium-ion batteries, Carbon 82 (2015) 337–345.
M. Sohn et al. / Electrochimica Acta 196 (2016) 197–205 [42] Y. Zhou, X. Jiang, L. Chen, J. Yue, H. Xu, K. Yang, Y. Qian, Novel mesoporous silicon nanorod as an anode material for lithium ion batteries, Electrochim. Acta 127 (2014) 252–258. [43] D. Shao, S. Tang, Y. Mai, L. Zhang, Nanostructured silicon/porous carbon spherical composite as a high capacity anode for Li-ion batteries, J. Mater. Chem. A 1 (2013) 15068.
205
[44] Q. Bao, Y.-H. Huang, C.-K. Lan, B.-H. Chen, J.-G. Due, Scalable upcycling silicon from waste slicing sludge for high-performance lithium-ion battery anodes, Electrochim. Acta 173 (2015) 82–90.
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Electrochimica Acta journal homepage: www.elsevier.com/locate/electacta
Porous Silicon–Carbon Composite Materials Engineered by Simultaneous Alkaline Etching for High-Capacity Lithium Storage Anodes Myungbeom Sohna , Dae Sik Kima , Hyeong-Il Parka , Jae-Hun Kimb , Hansu Kima,* a b
Department of Energy Engineering, Hanyang University, Seoul 04763, Republic of Korea School of Advanced Materials Engineering, Kookmin University, Seoul 02707, Republic of Korea
A R T I C L E I N F O
A B S T R A C T
Article history: Received 11 December 2015 Received in revised form 11 February 2016 Accepted 16 February 2016 Available online 20 February 2016
Porous silicon–carbon (Si–C) composite materials have attracted a great deal of attention as highperformance anode materials for Li-ion batteries (LIBs), but their use suffers from the complex and limited synthetic routes for their preparation. Herein we demonstrate a scalable and nontoxic method to synthesize porous Si–C composite materials by means of simultaneous chemical etching of Si and carbon phases using alkaline solution. The resulting porous Si–C composite material showed greatly improved cycle performance, good rate capability, and high dimensional stability during cycling. Porous Si–C electrode showed an expansion of the height by about 22% after the first lithiation and only 16% after the first cycle. The material synthesis concept and scalable simultaneous etching approach presented here represent a means of improving the electrochemical properties of Si-based porous anode materials for use in commercial LIBs. ã 2016 Elsevier Ltd. All rights reserved.
Keywords: anode porous composite Si–C composite alkaline etching lithium-ion battery
1. Introduction Lithium-ion batteries (LIBs) have played important role in the development of portable electronics and have also gained a great deal of attention as power sources for emerging applications such as electric vehicles and large-scale energy storage systems [1–3]. However, in order to meet recent requirements for aforementioned emerging applications, the electrochemical performance of LIBs needs to be further advanced in terms of energy density and power density [4–6]. In an attempt to increase the energy density of LIBs, development of new electrode materials with higher capacity than currently used materials such as LiCoO2 and graphite has been pursued during the past several decades [7–9]. Silicon has been regarded as a particularly promising candidate anode material to replace graphite, because its capacity (3579 mAh g 1) is about ten times higher than that of graphite (372 mAh g 1) [10,11]. However, the practical use of Si has been limited by its extremely large volume changes during alloying and dealloying reactions with Li, leading to poor capacity retention of Si electrodes [12–14]. Various strategies have been suggested to solve these problems of Si anodes, and some of them showed improved capacity retention by managing the volume changes of Si during
* Corresponding author. Tel.: +82 2 2220 2414; fax: +82 2 2220 2489. E-mail address: [email protected] (H. Kim). http://dx.doi.org/10.1016/j.electacta.2016.02.101 0013-4686/ ã 2016 Elsevier Ltd. All rights reserved.
cycling [15]. Among these promising Si-based anode materials, Si– C composites are being regarded as highly promising anode materials because of their stable cycle performance and relatively low production cost [16–19]. Porous Si-based materials have also showed some promising behaviors, because pores in the material could effectively accommodate the mechanical strain induced by alloying and dealloying reactions of Si with Li, thereby minimizing mechanical degradation of the resulting Si-based electrodes. In this regard, many approaches to prepare porous Si-based materials have been investigated, and some examples of porous Si-based materials showed good cycle performance comparable to that of currently used graphite anode materials for LIBs [20–22]. However, most of the approaches to prepare porous Si-based anode materials require either time-consuming and expensive processing or the use of hazardous chemicals such as hydrofluoric acid and silane gas, making production process expensive [23,24]. Given that the production cost of LIBs is one of the most challenging issues for their successful use in emerging applications for electric vehicles and large-scale energy storage systems [25], further addressing this technical issues of porous Si-based anode materials is necessary to ensure their commercial success. Toward a scalable and cost-effective process to produce porous Si-based anode materials, here we demonstrate a facile route for the preparation of porous Si–C composites by means of simple ball milling followed by simultaneous one-pot chemical etching of Si
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and carbon phases in the Si–C composite using alkaline (NaOH) aqueous solution. By using this method in the present work, we could obtain porous Si–C composites with the pores inside the composite as well as at the surface of the composite, because alkaline solution has the capability to etch both carbonaceous materials and Si [26–30]. Pan et al. synthesized a york–shell Si–C nanocomposite by forming a void between the Si nanoparticle and the carbon shell using NaOH etching of the Si nanoparticle [31]. Instead of embedding Si particles in hollow carbon shell, it has been recently revealed that the structure of porous carbon surrounding nanoscale Si can accommodate the large volume
changes of Si during lithiation and delithiation, thus showing outstanding cycle performance [32,33]. However, the methods used thus far to generate pores in the carbon phase are complicated or involve toxic chemicals. Considering that both ball milling and alkaline etching are environmentally benign and already established in the industrial processes, our simple alkaline etching combined with ball-milling process would be another promising synthetic routes for porous Si–C composite anode materials for next-generation LIBs. In this work, we describe this method in detail and demonstrate its promise for the preparation of porous Si–C composite anode materials for use in next-generation LIBs.
Fig. 1. Schematic illustration of synthesizing (a) non-porous Si–C composite through annealing (heating and cooling) and (b) porous Si–C composite by etching non-porous Si–C in alkaline solution for times of T1 (105 min) and T1 + T2 (115 min). Powder SEM images of (c) ball-milled Si, (d) non-porous Si–C, and (e) porous Si–C composite (T1 + T2). Cross-sectional SEM images of (f) non-porous Si–C, (g) less porous Si–C (T1), and (h) porous Si–C (T1 + T2).
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2. Experimental 2.1. Preparation of Porous Si–C Composites To reduce particle size, bulk Si powder ( 325 mesh, 99% trace metals basis, Sigma-Aldrich) was ball-milled for 36 h at 1000 rpm using ball-to-powder ratio of 20 under an Ar atmosphere using a custom-made vibratory mill. The ball-milled Si and pitch were hand-mixed with the weight ratio of 7:3 using a mortar and pestle. The mixture was annealed in a vertical furnace under an N2 atmosphere at 800 C for 2 h with a heating rate of 5 C min 1 and then allowed to cool naturally to room temperature. Non-porous Si–C powder was obtained by grinding the annealed sample using a mortar and pestle. Porous Si–C composites were prepared by ultrasonication (Branson CPX3800H-E) of a non-porous Si–C composite together with 0.5 mol L 1 NaOH solution in a highdensity polyethylene (HDPE) bottle for either 105 or 115 min at room temperature; the etching reaction was then quenched by adding 0.2 M HCl solution.
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conductive agent (Super-P, 5 wt%), and polyacrylic acid binder (PAA, 15 wt%) in deionized water. After coating, the electrodes were dried at 120 C under vacuum conditions for 12 h and then allowed to cool naturally to room temperature. Electrochemical properties were analyzed using CR2032 coin-type half-cells assembled in a dry room. Each coin cell consisted of a porous polyethylene (PE) membrane as the separator and 1 mol L 1 LiPF6 solution dissolved in a mixed solvent of ethylene carbonate (EC) and diethyl carbonate (DEC) (3:7 v/v.) with fluoroethylene carbonate (5%) (PANAX Etec Co., Ltd.). The cycle life of cells was tested in the voltage range of 0.01 to 1.5 V vs. Li/Li+ at a current density of 100 mA g 1. The rate capability test and first cycle charge– discharge of cells for measuring electrode thickness were carried out in the voltage range 0.005 to 1.5 V vs. Li/Li+ at current densities of various values and 100 mA g 1, respectively. Electrochemical impedance spectroscopy (EIS) measurement was performed using a potentiostat (VSP-300, BioLogic) over a frequency range of 1 MHz to 10 mHz with an amplitude of 5 mV. 3. Results and Discussion
2.2. Materials Characterizations Surface and cross-sectional morphologies of the samples were observed by field emission scanning electron microscopy (FESEM, JEOL JSM-7000F). Surface areas and pore sizes were measured from Brunauer–Emmett–Teller (BET) and Barrett–Joyner–Halenda (BJH) results by using a surface analyzer (Micromeritics 3Flex). The carbon content in the non-porous and porous Si–C composites was calculated by means of thermogravimetric analysis (TGA, TA Instruments SDT Q600) under air. Raman spectroscopy (JASCO NRS-3100) was carried out to reveal the carbon phase in the Si–C composites. 3.3. Electrochemical Measurements Electrodes were prepared by coating slurries on 10 mm thickness Cu foil. The slurries were made by mixing non-porous or porous Si–C composite as the active material (80 wt%), a
A non-porous Si–C composite material was prepared by typical process, i.e., mixing ball-milled Si microparticles with pitch as a carbon precursor (the left image of Fig. 1a), followed by carbonization of pitch under a reducing atmosphere (the center and right images of Fig. 1a) [34]. As schematically depicted in Fig. 1b, pores were formed in the non-porous Si–C composite via simultaneous alkaline (NaOH) etching process. The porous Si–C composites with etching time T1 (105 min) and T1 + T2 (115 min) are denoted as less porous Si–C and porous Si–C composite material, respectively. The T2 is additional etching time (10 min) after time T1. Fig. 1c, d, and e present scanning electron microscopy (SEM) images of ball-milled Si, non-porous Si–C, and porous Si–C composite particles, respectively. Aggregation of ball-milled Si particles in carbon framework is observed in Fig. 1d, because the Si–C composite particles are much larger than the ball-milled Si particles. Pore generation inside the Si–C composite particles can be inferred from the maintenance of their dimensions on the order
Fig. 2. SEM–EDS analysis with elemental mapping of (a) non-porous, (b) less porous, and (c) porous Si–C composites.
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Fig. 3. (a) N2 adsorption–desorption isotherms and (b) pore size distributions of non-porous and porous Si–C composites.
of tens of micrometers after NaOH treatment as shown in Fig. 1e. Compared to the ball-milled Si particles, which have uneven surfaces (Fig. 1c inset), the non-porous Si–C composite particles exhibit smoother surface (Fig. 1d inset), indicating that the ballmilled Si particles are well surrounded by carbon after carbonization of the pitch. We found that the NaOH etching resulted in a roughened surface of the Si–C composite particles (Fig. 1e inset), suggesting that pores were created in the carbon phase. The nonporous Si–C composite particles have a typical microstructure in which micron-sized Si particles were dispersed in the carbon matrix (the left image of Fig. 1b and Fig. 1f). The micrometer-scale pores were formed by the chemical etching of some Si microparticles inside non-porous Si–C composite particles as shown in cross-sectional images of porous Si–C composite particles (Figs. 1g, h and 2 ). Fig. 3 shows N2 adsorption/desorption isotherms and pore size distribution of Si–C composites estimated by using Brunauer– Emmett–Teller (BET) and Barrett–Joyner–Halenda (BJH) methods. While the as-prepared non-porous Si–C composite shows typical non-porous N2 adsorption/desorption isotherms, the chemically etched Si–C composites exhibit a hysteresis loop with steep
condensation steps in the range of relative pressure, P/P0 = 0.80– 0.98, which is a typical characteristic of porous materials with disordered pore structure (Fig. 3a). Note that these isotherms show a low pressure hysteresis in which the adsorption and desorption branches have no contact points at low P/P0, which might result from the micropores in the composite materials, which have almost same size to that of adsorbates (N2) [35]. It is well known that such a narrow opening of the small pores could restrict the access of the adsorbates [36]. Considering that this hysteresis were commonly observed even at the BET curves for non-etched Si–C composite material (Fig. 3a), micropores in the composites might come from the carbon phase in the composite after carbonization of pitch at 800 C. The simultaneous chemical etching process also enlarged the surface area and created the abundant pores with the size of 50–250 nm in the composite materials (Fig. 3b). In the case of 115 min NaOH treatment, the BET surface area of the Si–C composite increased from about 76.8 to about 123.4 m2 g 1 by simultaneous chemical etching of Si microparticles and carbonaceous phase (Table 1). We also found that the pore volume of porous Si–C composite, as calculated from desorption branch using the BJH method, was dramatically increased from 2.2 10 3 to
Table 1 BET and BJH data of non-porous, less porous, and porous Si–C composites.
BET surface area BJH adsorption cumulative surface area of pores between 17 Å and 3,000 Å diameter BJH desorption cumulative surface area of pores between 17 Å and 3,000 Å diameter BJH adsorption cumulative volume of pores between 17 Å and 3,000 Å diameter BJH desorption cumulative volume of pores between 17 Å and 3,000 Å diameter BJH adsorption average pore diameter BJH desorption average pore diameter
Non-porous Si–C
Less Porous Si–C
Porous Si–C
76.7962 m2 g 1 0.108 m2 g 1 0.1149 m2 g 1 0.002372 cm3 g 0.002264 cm3 g 904.944 Å 788.469 Å
98.8482 m2 g 1 3.964 m2 g 1 5.3942 m2 g 1 0.046295 cm3 g 0.044438 cm3 g 467.208 Å 329.525 Å
123.4492 m2 g 1 9.299 m2 g 1 9.8568 m2 g 1 0.081894 cm3 g 1 0.084886 cm3 g 1 352.255 Å 344.478 Å
1 1
1 1
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Fig. 4. TGA curves of non-porous and porous Si–C composites.
80.5 10 3 cm3 g 1 as shown in Table 1. This enlarged surface area as well as increased pore volume would be helpful in accommodating the huge volume changes of Si upon lithiation and delithiation during cycling [37–39]. In order to determine the Si to carbon mass ratio in the Si–C composites, thermogravimetric analysis (TGA) was carried out. Fig. 4 shows TGA curves of the Si–C composites before and after chemical etching using NaOH solution. Carbon content in the nonporous Si–C composite is estimated to be about 22 wt%, while those of porous Si–C composites are 44 and 50 wt% after chemical etching for 105 and 115 min, respectively. These results suggest that NaOH solution is more effective in etching out Si phase than in etching the carbon matrix, which is well matched with the crosssectional observation shown in Fig. 1 h. The existence of carbon phases is revealed by D (1350 cm 1) and G (1580 cm 1) bands in Raman spectra collected at three arbitrary points in each Si–C composites (Fig. S1). Fig. 5a shows the voltage profiles of non-porous and porous Si– C composite electrodes for the first cycle. The non-porous Si–C composite electrode delivered a reversible capacity of about 2471 mAh g 1 with an initial efficiency of about 90%, while the less porous and porous Si–C composite electrodes exhibited the discharge capacities of 1438 and 1077 mAh g 1, respectively. The decrease in the reversible capacity of porous Si–C composite
electrodes is mainly attributed to the relatively lower Si contents in the composites as observed from voltage profiles (Fig. 5a) and TGA curves (Fig. 4). Note that the initial efficiencies of less porous (81%) and porous (74%) Si–C composite electrodes were also lower than that of non-porous Si–C composite electrode. This increase in the irreversible electrochemical reaction observed at porous Si–C composite electrodes is probably due to the enlarged surface area caused by chemical etching, which might increase the electrolyte decomposition reaction on the surface of the electrode during the first cycle [40,41]. However, the porous Si–C composite electrode showed a dramatic improvement in the cycle performance compared with that of non-porous Si–C composite electrode. The porous Si–C composite electrode maintained about 74% of their initial capacity even after 100 cycles, while non-porous Si–C composite electrode shows a fast capacity fading (8% of the initial capacity) within 50 cycles, as revealed in Fig. 5b. Coulombic efficiencies of Si–C composite electrodes showed the same tendency; the non-porous Si–C composite electrode showed poor Coulombic efficiencies less than 95%, while porous Si–C composite electrodes showed Coulombic efficiencies more than 98% during most cycles (Fig. 5c). In particular, the porous Si–C composite electrode obtained after chemical etching for 115 min maintained high Coulombic efficiencies without significant fading even after 100 cycles. As might be expected, the increase in both the surface
Fig. 5. Electrochemical properties of non-porous and porous Si–C composites. (a) Voltage profiles, (b) cycle performance, (c) Coulombic efficiencies of non-porous and porous Si–C composites, and (d) rate capability of porous Si–C composite only.
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Fig. 6. Cross-sectional FESEM images of (a) non-porous and (b) porous Si–C electrodes during the first cycle.
area and pore volume of Si–C composite materials by simultaneous chemical etching using NaOH solution led to much improved rate capability in the porous Si–C composite electrode (Fig. 5d) compared to that of non-porous composite electrode (Fig. S2). Although the gravimetric energy density of non-porous Si–C composite electrode is almost double that of the porous composite electrode at the current density of 0.1 A g 1 (Fig. S2), the capacity of non-porous composite electrode ( < 1000 mAh g 1) is lower than that of porous composite electrode (1100 mAh g 1) at 1 A g 1 (Fig. 5d). Porous Si–C composite prepared by simultaneous chemical etching route showed better rate capability compared with previously reported Si–C composite anode materials for lithium ion batteries. For example, yolk-shell structured Si–C nanocomposite showed a high reversible capacity of 2258 mAh g 1 at a rate of 0.1 A g 1, but this material showed a drastic capacity decrease with an increase of the current densities (about 250 mAh g 1 at a rate of 5 A g 1) [31]. Interconnected porous Si–C composites by using the Rochow reaction and hydrofluoric etching had low reversible capacity of 856.5 mAh g 1 at a rate of 50 mA g 1 and also showed poor rate capability at a high rate condition (440 mAh g 1 at a rate of 1 A g 1) [38]. However, the porous Si–C composite in this work showed a high rate performance with reversible capacity of about 1180 and 990 mAh g 1 at a rate of 0.1 and 5 A g 1, respectively. This improvement in the rate capability can be attributed to the pore created by simultaneous chemical etching of Si and carbon phase in the composite. The abundant pores formed by alkaline etching would not only allow the electrolyte to penetrate into the composite, but would also increase the interfacial area between the electrolyte and the electrode material. Besides providing high Li+ accessibility, carbon frames basically could offer the pathways for electron transport. As a result, the porous Si–C composite electrodes show much improved electrochemical kinetics compared with nonporous one. We anticipated that the pores in the Si–C composite electrode could also improve the dimensional stability of the electrode against huge volume changes of Si phase in the composite during cycling. In order to confirm this effect of the pores in the Si–C composite electrodes, the dimensional changes in the non-porous and porous composite electrodes were investigated using FESEM during the initial cycle. Fig. 6 shows cross-sectional FESEM images of non-porous Si–C composite electrode and porous Si–C composite electrode after chemical etching for 115 min after lithium insertion and removal. After fully lithiated, the non-porous Si–C composite electrode expanded up to about 310% of its pristine state (Fig. 6a and Table S1). On the other hand, the porous Si–C composite electrode showed much improved dimensional stability
during the first cycle. Even after full lithiation, the porous Si–C composite electrode expands to only 122% of its thickness of pristine electrode (Fig. 6b and Table S1). The average thickness of the porous Si–C electrode after delithiation is 40.7 mm which is slightly thicker (116%) than that of pristine electrode (35.1 mm) (Fig. S3b and Table S1). In contrast, meaningful thickness cannot be measured in the delithiated non-porous Si–C electrode because the electrode contents are sparse on the Cu foil (Fig. S3a and Table S1). This severe detachment of non-porous Si–C particles from the foil might be due to the weakened adhesion between active material and current collector by the huge volume change of Si during lithiation and delithiation. The loading of active material on the non-porous Si–C electrode (2.69 mg cm 2) is greater than that of porous Si–C electrode (1.86 mg cm 2) with similar pristine thickness (Table S1), suggesting that the density of porous Si–C material is less than that of non-porous Si–C material caused by pores generated inside the Si–C composite via NaOH etching treatment. The effects of pores are remarkable when the nonporous and porous Si–C electrodes are compared from the volumetric perspective. Although the non-porous Si–C has a gravimetric discharge capacity (2482 mAh g 1) twice that of porous Si–C (1242 mAh g 1) composite electrode, the difference between their volumetric electrode capacities based on the thickness of lithiated electrodes is only about 233 mAh cm 3 (Table S1). This indicates that the preformed pores in the Si–C composite alleviated the stresses induced by large volume changes of Si during charge and discharge. Closer observation on the cycled electrodes shows more distinctive differences between the porous and non-porous Si–C composites. Fig. 7a and b compare the cross-sectional FESEM images of the non-porous and porous Si–C composite electrode with respect to the depth of discharge. Pores observed in the pristine porous Si–C composite disappeared after full lithiation and appeared again after full delithiation, clearly showing the reversible behavior of the pores in the porous Si–C composite material (Fig. 7b and Fig. S4b). On the other hand, non-porous Si–C composite electrode showed internal cracks in the material after lithiation as shown in Fig. 7a and Fig. S4a. This difference in the morphological changes of the Si–C composite electrodes clearly demonstrates how pores in the composite play an active role in accommodating the volume changes of the Si phase during cycling. In order to further characterize the electrochemical properties of Si–C composite electrode before and after chemical etching, we carried out the EIS analysis on the electrodes after 100 cycles. The Nyquist plots of the non-porous and porous Si–C composite electrode showed two distinguishable semicircles in the high frequency and middle frequency region (Fig. 8). The oblique
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Fig. 7. High-magnification cross-sectional FESEM images of (a) non-porous and (b) porous Si–C electrodes during the first cycle.
straight lines in the low frequency region were also observed in all the Nyquist plots of the electrodes. The semicircle in the high frequency region is mainly attributed to the SEI film on the electrode, and the semicircle observed in the middle frequency region is closely related to the charge transfer resistance of active material, and the straight line in the low frequency region is related with the mass transfer, such as lithium ion diffusion in the electrode [42–44]. As shown in Fig. 8, the non-porous electrode showed a distinctive semicircle at the high frequency region which is related to the SEI formation, while both the less porous and
porous electrode showed a very small semicircle in the high frequency region after 100 cycles. This suggests that the less SEI layers were formed in the porous electrodes than in the nonporous electrode. Moreover, the porous Si–C composite electrode showed smaller semicircle in the middle frequency than that of non-porous Si–C composite electrode after 100 cycles, indicating that the porous electrode had better electronic contact between electrode components. This lower impedance could be attributed to high structural integrity in the porous electrode during cycling as the pores could alleviate the stress from the huge volume
Fig. 8. Nyquist plots of non-porous, less-porous, and porous Si–C electrodes after 100 cycles.
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change of Si. These improved electrochemical properties result in enhanced cycle performance as well as rate capability of the porous Si–C composite electrode as shown in Fig. 5b, c, and d. 4. Conclusions We demonstrated a cost-effective simultaneous chemical etching route for porous Si–C composite in which abundant pores as well as micrometer scale Si particles are incorporated into carbon matrix. This mechanically robust design of porous Si–C composites provided highly reversible Li storage over prolonged cycles and good dimensional stability upon alloying and dealloying reaction of Si phase with Li. In particular, the porous Si–C composite obtained from simultaneous etching of Si and carbon phase exhibited a high reversible capacity of more than 1000 mAh g 1 with an initial coulombic efficiency of 74% and excellent cycle performance up to 100 cycles. The highly enhanced battery performances of the proposed Si–C composite was mainly attributed to the pores in the composite, which act as a dynamic buffer phase to accommodate the volume change of the Si phase without significant degradation. This material concept and scalable simultaneous alkaline etching approach provide a means of improving the electrochemical properties of Si-based anode materials for use in commercial LIBs. Acknowledgements This work was in part supported by the Korea Evaluation Institute of Industrial Technology (KEIT), which is funded by the Ministry of Trade, Industry & Energy, Republic of Korea (No. 10046341) and in part supported by the research fund of Hanyang University (HY-2012-T). Appendix A. Supplementary data Supplementary data associated with this article can be found, in the online version, at http://dx.doi.org/10.1016/j. electacta.2016.02.101. References [1] M. Armand, J.-M. Tarascon, Building better batteries, Nature 451 (2008) 652– 657. [2] B. Dunn, H. Kamath, J.-M. Tarascon, Electrical energy storage for the grid: a battery of choices, Science 334 (2011) 928–935. [3] H. Kim, G. Jeong, Y.-U. Kim, J.-H. Kim, C.-M. Park, H.-J. Sohn, Metallic anodes for next generation secondary batteries, Chem. Soc. Rev. 42 (2013) 9011–9034. [4] J.B. Goodenough, Y. Kim, Challenges for rechargeable Li batteries, Chem. Mater. 22 (2009) 587–603. [5] J.B. Goodenough, K.-S. Park, The Li-ion rechargeable battery: a perspective, J. Am. Chem. Soc. 135 (2013) 1167–1176. [6] R. Van Noorden, A better battery, Nature 507 (2014) 26–28. [7] H. Li, Z. Wang, L. Chen, X. Huang, Research on advanced materials for Li-ion batteries, Adv. Mater. 21 (2009) 4593–4607. [8] B. Scrosati, J. Garche, Lithium batteries: status, prospects and future, J. Power Sources 195 (2010) 2419–2430. [9] S. Goriparti, E. Miele, F. De Angelis, E. Di Fabrizio, R.P. Zaccaria, C. Capiglia, Review on recent progress of nanostructured anode materials for Li-ion batteries, J. Power Sources 257 (2014) 421–443. [10] H. Liu, Z. Guo, J. Wang, K. Konstantinov, Si-based anode materials for lithium rechargeable batteries, J. Mater. Chem. 20 (2010) 10055–10057. [11] M. Obrovac, V. Chevrier, Alloy negative electrodes for Li-ion batteries, Chem. Rev. 114 (2014) 11444–11502. [12] J.H. Ryu, J.W. Kim, Y.-E. Sung, S.M. Oh, Failure modes of silicon powder negative electrode in lithium secondary batteries, Electrochem. Solid-State Lett. 7 (2004) A306–A309. [13] U. Kasavajjula, C. Wang, A.J. Appleby, Nano- and bulk- silicon-based insertion anodes for lithium-ion secondary cells, J. Power Sources 163 (2007) 1003– 1039. [14] D. Munao, J. Van Erven, M. Valvo, E. Garcia-Tamayo, E. Kelder, Role of the binder on the failure mechanism of Si nano-composite electrodes for Li-ion batteries, J. Power Sources 196 (2011) 6695–6702.
[15] H. Wu, Y. Cui, Designing aanostructured Si anodes for high energy lithium ion batteries, Nano Today 7 (2012) 414–429. [16] H. Kim, E.-J. Lee, Y.-K. Sun, Recent advances in the Si-based nanocomposite materials as high capacity anode materials for lithium ion batteries, Mater. Today 17 (2014) 285–297. [17] J. Guo, A. Sun, X. Chen, C. Wang, A. Manivannan, Cyclability study of siliconcarbon composite anodes for lithium-ion batteries using electrochemical impedance spectroscopy, Electrochim. Acta 56 (2011) 3981–3987. [18] R. Yi, F. Dai, M.L. Gordin, H. Sohn, D. Wang, Influence of silicon nanoscale building blocks size and carbon coating on the performance of micro-sized SiC composite Li-ion anodes, Adv. Energy Mater. 3 (2013) 1507–1515. [19] X. Su, Q. Wu, J. Li, X. Xiao, A. Lott, W. Lu, B.W. Sheldon, J. Wu, Silicon-based nanomaterials for lithium-ion batteries: a review, Adv. Energy Mater. 4 (2014) 1300882. [20] A. Magasinski, P. Dixon, B. Hertzberg, A. Kvit, J. Ayala, G. Yushin, Highperformance lithium-ion anodes using a hierarchical bottom-up approach, Nat. Mater. 9 (2010) 353–358. [21] D.S. Jung, T.H. Hwang, S.B. Park, J.W. Choi, Spray drying method for large-scale and high-performance silicon negative electrodes in Li-ion batteries, Nano Lett. 13 (2013) 2092–2097. [22] T. Jaumann, M. Herklotz, M. Klose, K. Pinkert, S. Oswald, J. Eckert, L. Giebeler, Tailoring hollow silicon-carbon nanocomposites as high-performance anodes in secondary lithium-based batteries through economical chemistry, Chem. Mater. 27 (2015) 37–43. [23] W. Wang, P.N. Kumta, Nanostructured hybrid silicon/carbon nanotube heterostructures: reversible high-capacity lithium-ion anodes, ACS Nano 4 (2010) 2233–2241. [24] C. Zhao, Q. Li, W. Wan, J. Li, J. Li, H. Zhou, D. Xu, Coaxial carbon-silicon-carbon nanotube arrays in porous anodic aluminum oxide templates as anodes for lithium ion batteries, J. Mater. Chem. 22 (2012) 12193–12197. [25] L. Lu, X. Han, J. Li, J. Hua, M. Ouyang, A review on the key issues for lithium-ion battery management in electric vehicles, J. Power Sources 226 (2013) 272–288. [26] S. Yumitori, Y. Nakanishi, Effect of anodic oxidation of coal tar pitch-based carbon fibre on adhesion in epoxy matrix: Part 2. comparative study of three alkaline solutions, Composites Part A 27 (1996) 1059–1066. [27] E. Raymundo-Piñero, P. Azaïs, T. Cacciaguerra, D. Cazorla-Amorós, A. LinaresSolano, F. Béguin, KOH and NaOH activation mechanisms of multiwalled carbon nanotubes with different structural organization, Carbon 43 (2005) 786–795. [28] H. Teng, T.-C. Weng, Transformation of mesophase pitch into different carbons by heat treatment and KOH etching, Micropor. Mesopor. Mater. 50 (2001) 53– 60. [29] H.-J. Ahn, J.-H. Lee, Y. Jeong, J.-H. Lee, C.-S. Chi, H.-J. Oh, Nanostructured carbon cloth electrode for desalination from aqueous solutions, Mater. Sci. Eng: A 449–451 (2007) 841–845. [30] H. Seidel, L. Csepregi, A. Heuberger, H. Baumgärtel, Anisotropic etching of crystalline silicon in alkaline solutions, J. Electrochem. Soc. 137 (1990) 3612– 3626. [31] L. Pan, H. Wang, D. Gao, S. Chen, L. Tan, L. Li, Facile synthesis of yolk–shell structured Si–C nanocomposites as anodes for lithium-ion batteries, Chem. Commun. 50 (2014) 5878–5880. [32] D. Shao, D. Tang, Y. Mai, L. Zhang, Nanostructured silicon/porous carbon spherical composite as a high capacity anode for Li-ion batteries, J. Mater. Chem. A 1 (2013) 15068–15075. [33] H.M. Jeong, S.Y. Lee, W.H. Shin, J.H. Kwon, A. Shakoor, T.H. Hwang, S.Y. Kim, B.-S. Kong, J.-S. Seo, Y.M. Lee, Silicon@porous nitrogen-doped carbon spheres through a bottom-up approach are highly robust lithium-ion battery anodes, RSC Adv. 2 (2012) 4311–4317. [34] Y.-X. Wang, S.-L. Chou, J.H. Kim, H.-K. Liu, S.-X. Dou, Nanocomposites of silicon and carbon derived from coal tar pitch: cheap anode materials for lithium-ion batteries with long cycle life and enhanced capacity, Electrochim. Acta 93 (2013) 213–221. [35] A.M. Puziy, O.I. Poddubnaya, A. Martínez-Alonso, F. Suárez-García, J.M.D. Tascón, Synthetic carbons activated with phosphoric acid II. Porous structure, Carbon 40 (2002) 1507–1519. [36] Y. Liao, J. Weber, C.F.J. Faul, Conjugated microporous polytriphenylamine networks, Chem. Commun. 50 (2014) 8002. [37] R. Zhang, Y. Du, D. Li, D. Shen, J. Yang, Z. Guo, H.K. Liu, A.A. Elzatahry, D. Zhao, Highly reversible and large lithium storage in mesoporous Si/C nanocomposite anodes with silicon nanoparticles embedded in a carbon framework, Adv. Mater. 26 (2014) 6749–6755. [38] Z. Zhang, Y. Wang, W. Ren, Q. Tan, Y. Chen, H. Li, Z. Zhong, F. Su, Scalable synthesis of interconnected porous silicon/carbon composites by the Rochow reaction as high-performance anodes of lithium ion batteries, Angew. Chem. Int. Ed. 53 (2014) 5165–5169. [39] X. Zhao, X. Rui, W. Zhou, L. Tan, Q. Yan, Z. Lu, H.H. Hng, Growth of Si nanowires in porous carbon with enhanced cycling stability for Li-ion storage, J. Power Sources 250 (2014) 160–165. [40] Z. Li, W. Wang, Z. Li, Z. Qin, J. Wang, Z. Liu, Bridging porous Si-C composites with conducting agents for improving battery cycle life, J. Power Sources 286 (2015) 534–539. [41] M.-S. Wang, W.-L. Song, J. Wang, L.-Z. Fan, Highly uniform silicon nanoparticle/ porous carbon nanofiber hybrids towards free-standing high-performance anodes for lithium-ion batteries, Carbon 82 (2015) 337–345.
M. Sohn et al. / Electrochimica Acta 196 (2016) 197–205 [42] Y. Zhou, X. Jiang, L. Chen, J. Yue, H. Xu, K. Yang, Y. Qian, Novel mesoporous silicon nanorod as an anode material for lithium ion batteries, Electrochim. Acta 127 (2014) 252–258. [43] D. Shao, S. Tang, Y. Mai, L. Zhang, Nanostructured silicon/porous carbon spherical composite as a high capacity anode for Li-ion batteries, J. Mater. Chem. A 1 (2013) 15068.
205
[44] Q. Bao, Y.-H. Huang, C.-K. Lan, B.-H. Chen, J.-G. Due, Scalable upcycling silicon from waste slicing sludge for high-performance lithium-ion battery anodes, Electrochim. Acta 173 (2015) 82–90.