Effect of Carbon Content on Nanocomposite Si(1-x)Cx Thin Film Anode for All-Solid-State Battery

Electrochimica Acta 147 (2014) 232–240

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Effect of Carbon Content on Nanocomposite Si(1-x)Cx Thin Film Anode for All-Solid-State Battery K.S. Lee a,b , Y.N. Lee b , Y.S. Yoon b, * a b

Department of Materials Science and Engineering, Yonsei University Shinchondong, 262 Seongsanno, Seodaemoongu, Seoul 120-749, Republic of Korea Department of Environment and energy engineering, Gachon University Seongnamdaero 1342, 461-710 Gyeonggi-do, 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 7 May 2014 Received in revised form 21 August 2014 Accepted 23 September 2014 Available online 28 September 2014

Silicon, a potential anode material for all-solid-state batteries, has the highest theoretical capacity during electrochemical reactions, but is vulnerable to structural expansion and irreversible capacity. Amorphous Si/C composite thin films have received attention as promising materials to address these obstacles. The amorphous Si/C composite thin films were synthesized by radio frequency magnetron co-sputtering, which easily fabricated uniformly dispersed nanocomposites, prevented the agglomeration of silicon particles, and more effectively resisted volume expansion than traditional techniques. Electrochemical evaluation of the Si37C63 specimen showed that it had a high 1st cycle capacity (6180 mAh cm 3) and a capacity retention of 96% from 100th to 200th cycle. These studies show that the stresses induced by volume expansion during long term cycling tests were reduced by the buffering effects of the carbon content. The homogeneously dispersed amorphous Si/C composite thin films were found to be promising anode materials for all-solid-state batteries. ã 2014 Elsevier Ltd. All rights reserved.

Keywords: Si/C thin film All-solid-state battery Anode Li ion battery

1. Introduction Lithium ion secondary batteries are generally used as power supplies for portable electronic devices [1–5]. In particular, all-solidstate (ASS) thin-film batteries have attracted considerable attention because of their many advantages, such as being organic free and having low self-discharge rates, wide operating temperature ranges, and good explosion and fire resistances [6–8]. ASS thin-film batteries are appropriate for use in the energy storage systems of micro electro mechanical system devices, smart cards, active RF cards, and medical and drug delivery devices [9–11]. ASS thin-film batteries normally use lithium metal as the anode because of its high theoretical capacity (3840 mA~h g 1) [12]. However, owing to its low melting temperature (181  C) and high moisture reactivity, fabrication of lithium containing devices is challenging and such devices have limited application windows [13,14]. Furthermore, an additional protective coating layer on the lithium negative electrode, which significantly increases cost, is necessary for the fabrication of a thinfilm battery [15]. In the pursuit of improving thin film battery performance, much attention has been paid to replace lithium metal with new active anode materials that can react with lithium ion, such as Ti [4,16], V [17,18], Si [5,19–21], and Sn [22–24] based materials. Of these potential candidates, silicon has been considered as a

* Corresponding author. Tel.: +82317505596 http://dx.doi.org/10.1016/j.electacta.2014.09.110 0013-4686/ ã 2014 Elsevier Ltd. All rights reserved.

promising anode material for ASS thin film batteries because of its high theoretical capacity (4200 mAh g 1; Li4.4Si phase) and comparatively low working potential (0.5 voltage (V)) [25]. However, silicon undergoes significant volume changes (320%; Li22Si5) and demonstrates poor capacity retention during Li ion insertion and extraction [26]. Volume expansion causes delamination from current collectors and breakdown of the electrical networks. Thus, modification of the electrode structure is considered vital to enhance the cycle stability of silicon [27]. Various approaches to enhance capacity retentions and cycle stabilities of silicon anodes have been studied. Among them, silicon nanostructures and carbon composites are considered the most promising candidates [28–30]. Carbon has many advantages over silicon, such as being relatively lightweight, having low reaction potential with lithium and reasonable electrical conductivity, and experiencing less volume expansion and higher structural stability during insertion and extraction [31,32]. The comparatively small expansion characteristics of carbon mean it can be used as a volume buffer to prevent silicon expansion. This can considerably enhance cycling performance and reduce the capacity drop induced by expansion. Hence, many efforts have been made to synthesize silicon carbon composites (Si/C) using techniques such as chemical synthesis of gels [29,33], pyrolysis [34], physical/chemical vapor deposition (PVD, CVD) [35,36], and mechanical ball milling [37]. Although, there are hundreds of papers in silicon and carbon composite, their long-term cyclabilities are still insufficient. This may be related to

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Fig. 1. Schematic illustration of co-sputtering system.

inhomogeneous dispersion of carbon onto silicon nanoparticles. In this study, to solve these problems, a homogeneously dispersed Si/C composite was prepared by sintering and used as a sputtering target. We have found that silicon expansion can be effectively prevented, even after 200 cycles, using a carbon volume buffer. Because carbon has a relatively low sputtering yield, the carbon concentrations of amorphous Si/C composite (a:Si1-xCx) thin films were controlled using a co-sputtering system with an extra carbon target, as shown in the Fig. 1. Moreover, the long term test performances (200 cycles) of a:Si1-xCx thin films with various carbon concentrations were compared. We have found that silicon expansion can be effectively prevented, even after 200 cycles, using a carbon volume buffer. 2. Experimental procedure 2.1. Material and preparation The silicon and carbon composite (Si/C) target (Sigma-Aldrich 325 mesh, 99% purity and < 45 mm, and 99.999%, respectively) was synthesized using a solid-state method. Silicon and carbon powder were mixed in a weight ratio of Si:C = 5:5. The mixture was ball milled for 12 hour with zirconia balls, forming a fine powder. The as-obtained composite powder was pressed in the form of a sputtering target, at 60 MPa for 10 min, then heated at a rate of 2  C/ min and sintered at 1000  C for 24 h in an Ar atmosphere. The extra carbon target was a commercial carbon (Pyrolytic Graphite) target (Kurt. J Lesker, 99.999% purity). Copper substrates (Nilaco Corporation 0.6 mm thick, 99.9% purity), used as current collectors, were ultrasonically cleaned with isopropyl alcohol and acetone for 20 min each. Thin films of chromium, used as glue layers, were deposited at 20 W for 15 min on the copper substrates. The base and working pressures of the chromium layer was 4.0  10 4 Pa and 6.5  10 1 Pa, respectively. To remove any surface oxides, Si/C and carbon targets were presputtered at 100 W for 30 min prior to deposition. a:Si1-xCx thin films, used as the active materials, were deposited from the Si/C and extra carbon targets using a radio frequency (r.f.) magnetron sputtering system. In order to be able to accurately compare the electrochemical properties of a:Si1-xCx thin

Fig. 2. XRD diffractogram of thin films of (a) Cu substrate, (b) Si51C49, (c) Si45C55, and (d) Si37C63.

films, a constant input power of 60 W was applied to the Si/C target and various input powers of 40, 80, and 120 W were applied to the carbon target (Table 1). The pure silicon thin film (Si100) was also deposited with input power at 60 W for 1 hour. The base and working pressures were 4  10 4 Pa and 6.7  10 1 Pa of argon (purity  99.999%), respectively. 2.2. Characterization The thicknesses of the a:Si1-xCx thin films were measured using a field emission scanning electron microscope (FE-SEM, Hitachi S-4200). X-ray diffraction (XRD, Rigaku RINT) was used to confirm the crystal structures of the specimens. The crystal structures of the thin films were also analyzed by high-resolution transmission electron microscopy (HR-TEM, JEM-2100F, JEOL). Depth profiles of the stoichiometric ratios and elemental distributions of the thin films were produce by X-ray photoelectron spectroscopy (XPS) (K-alpha, Thermo VG, U.K.) using a monochromated Al X-ray source (Al Ka line: 1486.6 eV)). 2.3. Electrochemical measurements Electrochemical characterizations were evaluated using twoelectrode coin cell (CR2032) assemblies for half-cell test. Lithium foil was used as a counter electrode and 1 M LiPF6 in EC:DEC (1:1) was

Table 1 The atomic concentration of each component of a:Si1-xCx thin films deposited with r.f magnetron sputtering from XPS depth profiles. Si input power (W) (a) 60 (b) 60 (c) 60

C input power(W)

Deposition time (hr)

Thickness (nm)

Concentration (at %)

40 80 120

1 1 1

 260  310  380

Si51C49 Si45C55 Si37C63

Fig. 3. High resolution TEM image of the as deposited Si37C63 thin film.

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used as the electrolyte. All of the coin cells were cycled with a battery cycler (WBCS3000), at room temperature (298 K), between 0.015 - 2.2 V, using a current density of 40 mA/cm2. The mass of the thin films were weighed using an analytical balance with a precision of  10 mg (Sartorius, CPA225D). Electrochemical impedance spectroscopy (EIS) analyses were carried out in a frequency range of 4 MHz to 100 mHz using a BioLogic EC-Lab (VSP-300). 3. Results and discussion 3.1. Characterization of silicon and carbon composite Fig. 2 shows XRD patterns of a:Si1-xCx thin films deposited on Cu substrate in order to clearly observe the peaks of a:Si1-xCx without interference by XRD spectra of the substrate. No obvious peaks corresponding to silicon and carbon except those of the Cu substrate (JCPDS 04-0836) can be seen in the diffraction patterns

shown in Fig. 2. The XRD peaks suggest that the Si51C49, Si45C55, and Si37C63 thin films were amorphous. HR-TEM was a more effective method for confirming the microstructures of the a:Si1-xCx thin films. Fig. 3 shows a HR-TEM image of the Si37C63 thin film. Ring type selected area electron diffraction (SAED) patterns were observed without lattice lines (inset of Fig. 3). This clearly implies that the Si37C63 thin film was amorphous and is in agreement with XRD results. It was expected that the amorphous structure would lead to improved electrochemical performance because amorphous active materials have shown homogeneous volume expansions and contractions during insertion and extraction [38]. To evaluate the volumetric specific capacities of the specimens, film thicknesses were measured by FE-SEM, as shown in Fig. 4. Specimens showed typical columnar microstructures perpendicular to the Cr/Cu substrates. Moreover, all of the specimens exhibited dense, smooth surfaces without voids or pinholes.

Fig. 4. SEM images of cross sections and surfaces of (a, b) Si51C49, (c, d) Si45C55, and (e, f) Si37C63.

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Dense, smooth surfaces facilitate fabrication of full cells and lead to lower contact resistances with electrolytes and current collectors. Film thickness increased with increasing carbon target input power. The atomic concentrations (see; Table 1) was evaluated from depth profiles of the a:Si1-xCx thin films, as shown in Fig. 5. The XPS depth profiles with argon ion bombardment show that the a:Si1xCx thin films consisted mainly of silicon, carbon, and copper substrate. Moreover, the silicon and carbon data plots of Si51C49, Si45C55, and Si37C63 without compositional fluctuations indicated that silicon and carbon were homogeneously deposited according to the depth direction as shown in Fig. 5. The XPS wide-scan spectra of the surfaces consist of O 2s, Cu 3p, Si 2p, Cu 3s, Si 2s, C 1s, O 1s, Cu LM1, Cu LM2, Cu 2p3, and Cu 2p1 spectra [39–41]. Si51C49, Si45C55, and Si37C63 have similar spectrum as shown in Fig. 5 (d). The oxygen contents generally formed native oxide during the sputtering which induces irreversible capacity corresponding to initial capacity drop as shown in Fig. 6 (a) due to formation of Li2O [42]. 3.2. Electrochemical properties Fig. 6(a) shows volumetric specific capacity (mAh cm 3) performance as a function of cycle number for bare Si100, Si51C49, Si45C55, and Si37C63 thin films, at a current density of 40 mA cm 2, in the cutoff range of 0.015–2.2 V. Si100, Si51C49, Si45C55, and Si37C63 showed initial discharge capacities of 8170, 5200, 5700, and 6180 mAh cm 3, respectively. As shown in Table 2, the 2nd cycle capacity retentions of Si100, Si51C49, Si45C55, and Si37C63 were 100, 73, 86, and 88%, respectively.

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Although Si100 shows highest retention until 4th cylcle, specific capacity of Si100 was dramatically decreased beyond the 5th cycle and it showed unstable cycle behavior. The lower capacity retention is due to the volume expansion of silicon, delamination from the current collectors and formation of cracks in the thin films [43]. Furthermore, initial capacity fading is mainly attributed to the formation of solid electrolyte interphase (SEI) layers, causing kinetic barriers for Li ion transport [44]. On the other hands, the electrochemical performances of silicon thin film normally decreases with the increase in thickness of silicon thin film, [45,46] due to low electrical conductivity of silicon and the relatively long diffusion length for Li ion. However, the Si37C63 thin film exhibited the highest 2nd cycle capacity retention of 88% although it is the thickest among three samples. Also capacity retention increased to 96% between the 100th to 200th cycles. The detailed capacity data are summarized in Table 2. Yang et al. have synthesized a:Si/C nanocomposite using spray-pyrolysis [47]. The a:Si/C nanocomposite showed high 1st and 160th cycle capacities, 1170 and 1029 mAh g 1, respectively. Sa et al. have investigated the use of Ni foam to prevent the expansion of silicon [25]. The silicon electrode on Ni foam showed 1st and 200th cycle capacities of 700 and 350 mAh g 1, respectively, which decreased rapidly thereafter with increasing cycle number. Chou et al. have synthesized a Si/graphene composite using the solvothermal method. The Si/graphene composite electrode also showed an initial capacity drop between the 1st cycle (2146 mAh g 1) and the 30th cycle (1164 mAh g 1) [48]. Wang et al. have fabricated Si/ carbon nanotubes (CNTs) using liquid injection chemical vapor deposition (CVD) [49]. Although Si/CNTs have a large initial capacity of 2552 mAh g 1, their irreversible capacities gradually

Fig. 5. XPS depth profiles of (a) Si51C49, (b) Si45C55, (c) Si37C63, and (d) XPS wide scan spectra.

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Fig. 6. Specific capacity vs. cycle number during insertion (discharge) cycled between 0.015 and 2.2 V at a current density of (a) 40 mA cm 2, and (b) 40  400 mA cm 2.

increased to 1000 mAh g 1 for the 100th cycle. All of these studies mentioned above showed 1st and end of cycle capacities, as well as long-term cycle stabilities lower than those of the Si37C63 specimen investigated in this study. Although gravimetric capacities were lower than the actual values, we observed 1st and 200th cycle capacities of 2690 and 1510 mAh g 1, respectively. The above results clearly suggest that the electrochemical performances of silicon based thin films can be improved by increasing carbon contents and making use of the buffer effect [36,50,51]. The good electrochemical properties of the a:Si1-xCx thin-film anodes reported here were due to a number of factors: (1)

homogeneously dispersed structures interrupted silicon nanoparticle aggregation; (2) volume expansion was buffered, and thereby resisted, by the carbon content; and (3) the formation of amorphous structures, as mentioned above. Because the Si37C63 thin film specimen, with the highest carbon content of the three samples, had the best electrochemical performance, we used it for further analysis. Fig. 6 (b) shows the rate capabilities of the Si37C63 thin film at different current densities, varied from 40 to 400 mA cm 2. Its reversible capacity decreased from 6700 mAh cm 3 during the 1st cycle, when using a current density of 40 mA cm 2, to 2490 mAh cm 3 during the 50th cycle, when using a current density of 400 mA cm 2. Although specific capacity decreased regularly with increasing current density, the tendency was less apparent at high current densities during the 10th to 50th cycles. When the rate returned to 40 mA cm 2 at the 50th cycle, capacity became 4810 mAh cm 3 compared with the 6700 mAh cm 3 at 1st cycle (capacity retention of 72%), indicating that the Si37C63 thin film shows a good electrochemical reversibility and structural stability at high current density. Differential capacity analyses (dQ/dV vs. potential) of the insertion (discharge) and extraction (charge (Fig. 7 (b)) were associated with phase transition and reaction mechanism of the active material [52]. The galvanostatic cycling curve of the Si37C63 thin film (voltage vs. specific capacity) was also plotted for the insertions and extractions of the 1st and 2nd cycles, as shown in Fig. 7 (a). The insertion profile of the 1st cycle shows that the Si37C63 thin film mainly reacted with the graphitic carbon content when the initial potential was decreased from 1.4 to 0.182 V, corresponding to peak potentials of 1.40 [53], 0.74 [54], and 0.18 V [55], respectively. Subsequently, the 1st insertion profile, below 0.18 V, indicates a Li13Si4 alloying process [56], amorphous LixSi [51], and Li22Si5 (full insertion phase) [55] phase transitions, corresponding to potential values of 0.15, 0.07, and 0.04 V, respectively. The differential capacity plot of the 1st extraction profile indicate a dealloying of the Si37C63 thin film from Li22Si5 at the 1st insertion to Li7Si3, then LixSi, and, finally, LixC by the presence of peaks at 0.29, 0.45 [57], and 0.92 V [53], respectively. Li13Si4 at 1st and 2nd extraction was absent due to the initial irreversible capacity. From the results obtained during 1st cycle insertion/extraction, silicon and carbon peaks are clearly observed without distortion. That means nanoscale silicon and carbon composite was homogeneously distributed and also confirmed by the XPS depth profiles (see Fig. 5). On the other hand, the 2nd cycle showed phase transitions during insertion and extraction. During 2nd cycle insertion, significant decreases in peak intensities at 1.40, 0.74, and 0.18 V were induced by irreversible capacities. However, the extraction profile of the 2nd cycle was almost identical to the 1st extraction profile, showing no decreases in peak intensities and indicating good cycling stability. The improved performance compared to pure silicon [46,58] is

Table 2 Summary of volumetric/gravimetric capacity from Fig. 6. Discharge capacityVolumetric (mAh/cm3)/gravimetric capacity (mAh/g)

Si100 Si51C49 Si45C55 Si37C63

Retention % 1stcycle

2ndcycle

100thcycle

200thcycle

8170/ 3550 5200/ 2000 5700/ 2190 6180/ 2690

8170/ 3555 4590/ 1765 4930/ 1910 5470/ 2380

–

–

3200/ 1230 3280/ 1261 3580/ 1560

3050/ 1180 3150/ 1220 3470/ 1510

1– 2ndcycle 100

100 – 200thcycle –

1–200th cycle –

73

95

58

86

96

55

88

96

56

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attributed to the carbon content, which not only buffers the silicon volume expansion during insertion/extraction but also homogeneously distributes silicon and carbon particles, avoiding agglomeration. 3.2.1. In-situ electrochemical impedance spectroscopy To better understand the capacity fading and correlations between phase transitions and Si37C63 electrode potentials, in-situ EIS measurements were performed using pre-selected potentials during 1st cycle insertion and extraction (discharge and charge). These potentials, ranging from the open circuit voltage (OCV) of 1.9 go down to 0.015 V as shown in Fig. 8. In-situ EIS spectra were fitted to a Nyquist plot using the equivalent circuit shown in the inset of Fig. 8 (b) [59]. For clarity, data points for insertion and extraction processes are separated along the horizontal axis of Fig. 8 (c). The Nyquist plot shows depressed semicircle which can be normally explained by 2 type of interpretations 1) charge transfer resistance [60] 2) interfacial and charge transfer resistance [14,22]. In our study, the depressed semicircles present two resistances, interfacial resistance (Rint: normally considered as SEI resistance [61,62]) and charge transfer resistance (Rct). However, it is difficult to obtain accurate data for individual components, Rint and Rct, because their time constants are considerably close in case of lithium ion batteries [63]. Thus, Rint and Rct were combined as the total resistance, Rtot, which consists of the ionic conductivity of the liquid electrolyte, the electrical conductivity of the active

Fig. 7. (a) charge–discharge plots of Si37C63, and (b) differential capacity plots (dQ/ dV vs. potential) of Si37C63 obtained at 1st and 2nd cycle (inset: enlarged plot of selected voltage ranges).

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materials, and the contact resistance between the active materials and the current collector. The Rtot of Si37C63 was closely connected to the cell potentials since the structure of the active material influences the lithiation state, which in turn affects lithium ion diffusion and resistance [22,63,64]. Although OCV of 1.9 V shows highest Rtot value of 5463 V, we have compared Rtot at 1.14 to 0.015 V because at OCV electrochemical reaction does not occur among lithium ion, electron, and electrolyte. In high frequency areas, from 1.14 to 0.015 V during insertion, Rtot in Fig. 8 (a) was evidently decreased from its original value of 2011 V to a minimum value of 263 V. It is expected that the improved electrical conductivity of Si37C63 was caused by the increase in amount of lithium ion inserted into the anode [65], which is similar to that of a carbon nanotube electrode [66]. In contrast, when lithium ions were extracted at 1.79 V as shown in Fig. 8 (b), Rtot reversibly increased to approximately 564 V. This indicates that the electrochemical reaction between Si37C63 and lithium ions is a stable and reversible reaction during insertion and extraction. The

Fig. 8. Typical impedance spectra of Si37C63 at different voltages during (a) insertion, (b) extraction, and (c) Rtot and exchange current density plot from (a) and (b).

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Fig. 9. Nyquist plots of Si37C63 at different cycles.

impedance data (Fig. 8 (a) and (b)) were also fitted to determine the exchange current density (io) and Rtot (Fig. 8 (c)) which characterize the total resistance and exchange current density of each individual component at various states of lithiation. Exchange current densities were calculated using Eq. (1) [65,67]. i0 = RT/nFRct (1)

(1)

Where R is the gas constant, T is the absolute temperature, n is the number of transferred electrons, F is the Faraday constant, and Rct is charge transfer resistance. The exchange current density of Si37C63 was rapidly increased during insertion from 5.11 10 2 to 3.90  10 1 mA cm 2 at 1.14 and 0.015 V, respectively. Subsequently, exchange current density behavior showed maximum values of 3.95  10 1 and 3.96  10 1 mA cm 2 at 0.3 and 0.75 V,

respectively, during extraction. This indicates that the electrochemical activity of Si37C63 at 0.015, 0.3, and 0.75 V as fully inserted lithium ions is much higher than at other potentials. From these results, it is known that exchange current density was offset by Rtot at the same potential because of the inverse relationship between Rtot and exchange current density shown in Eq.1. Moreover, exchange current density at the onset potential of lithium insertion, 1.14 V, which correspond to the differential capacity analysis shown in inset graph of Fig. 7 (b), were dramatically changed. As mentioned above, the phase transition of Si37C63, due to reaction between lithium and graphitic carbon, confirms that the enhanced electrical conductivity from in-situ EIS results was induced by lithium ions. To study the irreversible capacity of Si37C63 during cycling in more detail, impedance spectra were recorded at the end of each cycle as shown in Fig. 9. The Rtot in the pristine state was 163 V and gradually increased to 207 V (1st cycle), 234 V (3rd cycle), and 260 V (5th cycle). These values correspond to the initial capacity drop of Si37C63, seen in Fig. 6. Subsequently, the Rtot values of the 10th, 100th, and 200th cycle were 213, 205, and 209 V, respectively, which nearly overlap the semicircles. The Rtot values of Si37C63 from the 10th through 200th cycle were expected to be stable in the absence of delamination, formation of SEI layers, and volume expansion during cycling [65], thus indicating that Si37C63's cyclability was improved by carbon content. From the analyses of the two types of impedance results, we confirm that initial capacity degradation and long-term stability are clearly associated with lithiation state and Rtot. 3.2.2. Ex-situ SEM The increase of Rtot until 5th cycle and stable state of Rtot from 10th to 200th (Fig. 9) was investigated with the help of ex-situ SEM in order to analyze not only the buffer effect but also the relationship between Rtot and increasing cycles. The surface of the Si37C63 after 10th and 200th cycle (Fig. 10 (a) and (b)) was much rougher than as deposited state Fig. 4 (f) because of the volume change of silicon during insertion and extraction. In addition, 10th cycle (Fig. 10 (a)) seemed sparsely swollen surface induced by volume expansion of silicon. Although the surface of Si37C63 at after 200th cycle seemed rougher than at 10th cycle, most of the volume expansion was already caused until 10th cycle. The surface of Si37C63 at after 200th cycle shows good structural integrity without any cracks and delamination. From these results, it was confirmed that initial capacity degradation below 10th cycle (see; Fig. 6) and increasing Rtot until 5th cycle (see; Fig. 9) were related with initial volume expansion of Si37C63. In this study, we have found that a lower Si to C ratio yields a higher electrochemical performance for a:SixCy thin films which is irrespective of their thickness. Although we have found this phenomenon in our present study without controlling the Si to C ratio, we have considered it as our future work and it will contain a detailed investigation of the buffer effect depending on Si to C ratio by varying the weight ratio of sputtering targets. 4. Conclusion

Fig. 10. SEM images of the surface after (a) 5th cycle and (b) 200th cycle at a current density of 40 mA cm 2.

Thin films of a:Si1-xCx with different carbon contents were prepared by simple magnetron co-sputtering. The Si37C63 thin film showed a specific capacity of 3470 mAh cm 3 (1510 mAh g 1) even at the 200th cycle which is 4 times larger than the theoretical capacity of graphite. The capacity retentions of a:Si1-xCx thin films were found increased with carbon content. The Si37C63 demonstrated the highest capacity retention, 96% from the 100th to the 200th cycle, and good cycle stability within our rate capabilities. The behavior of the electrochemical reaction of the lithium ions at the various voltages and initial capacity degradations was

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investigated by in-situ EIS. The capacity vs. cycle number plot and EIS confirmed that increasing the carbon content effectively enhances the electrochemical performance of a:Si1-xCx. This study shows that a:Si1-xCx thin films have promise as anode materials for ASS batteries. Acknowledgements This work is supported by the National Strategic R&D Program for Industrial Technology (10043868) funded by the Ministry of Trade, Industry and Energy (MOTIE). This work was supported by the Innovations in Nuclear Power Technology of the Korea Institute of Energy Technology Evaluation and Planning (20131510400070) grant funded by the Korea government Ministry of Trade, Industry & Energy. References [1] V. Patil, D. Wook Shin, J.-W. Choi, D.-S. Paik, S.-J. Yoon, Issue and challenges facing rechargeable thin film lithium batteries, Mater. Res. Bull. 43 (2008) 1913, doi:http://dx.doi.org/10.1016/j.materresbull.2007.08.031. [2] J.M. Tarascon, M. Armand, Issues and challenges facing rechargeable lithium batteries, Nature 414 (2001) 359, doi:http://dx.doi.org/10.1038/35104644. [3] L. Baggetto, R.A.H. Niessen, F. Roozeboom, P.H.L. Notten, High Energy Density All-Solid-State Batteries: A Challenging Concept Towards 3D Integration, Adv. Funct. Mater. 18 (2008) 1057, doi:http://dx.doi.org/10.1002/adfm.200701245. [4] D. Cai, P. Lian, X. Zhu, S. Liang, W. Yang, H. Wang, High specific capacity of TiO2graphene nanocomposite as an anode material for lithium-ion batteries in an enlarged potential window, Electrochim. Acta 74 (2012) 65, doi:http://dx.doi. org/10.1016/j.electacta.2012.03.170. [5] L.B. Chen, J.Y. Xie, H.C. Yu, T.H. Wang, Si–Al thin film anode material with superior cycle performance and rate capability for lithium ion batteries, Electrochim. Acta 53 (2008) 8149, doi:http://dx.doi.org/10.1016/j.electacta.2008.06.025. [6] Y.S. Yoon, S.H. Jee, N. Kakati, J. Maiti, D.-J. Kim, S.H. Lee, H.H. Yoon, Work function effects of ZnO thin film for acetone gas detection, Ceram. Int. 38 (2012) S653, doi:http://dx.doi.org/10.1016/j.ceramint.2011.05.128. [7] Y.S. Yoon, S.H. Lee, S.B. Cho, S.C. Nam, Influence of Two-Step Heat Treatment on Sputtered Lithium Cobalt Oxide Thin Films, J. Electrochem. Soc. 158 (2011) A1313, doi:http://dx.doi.org/10.1149/2.036112jes. [8] Y.S. Yoon, S.B. Cho, S.C. Nam, A novel battery system for using in extremely low temperature conditions, Electrochim. Acta 71 (2012) 86, doi:http://dx.doi.org/ 10.1016/j.electacta.2012.03.082. [9] J.B. Bates, N.J. Dudney, B.J. Neudecker, F.X. Hart, H.P. Jun, S.A. Hackney, Preferred Orientation of Polycrystalline LiCoO2 Films, J. Electrochem. Soc. 147 (2000) 59, doi:http://dx.doi.org/10.1149/1.1393157. [10] R. Latham, Biomedical applications of batteries, Solid State Ionics 172 (2004) 7, doi:http://dx.doi.org/10.1016/j. ssi.2004.04.024. [11] A.C. Dillon, L.A. Riley, Y.S. Jung, C. Ban, D. Molina, A.H. Mahan, A.S. Cavanagh, S. M. George, S.H. Lee, HWCVD MoO3 nanoparticles and a-Si for next generation Li-ion anodes, Thin Solid Films 519 (2011) 4495, doi:http://dx.doi.org/10.1016/ j.tsf.2011.01.337. [12] T. Zhang, N. Imanishi, S. Hasegawa, A. Hirano, J. Xie, Y. Takeda, O. Yamamoto, N. Sammes, Water-Stable Lithium Anode with the Three-Layer Construction for Aqueous Lithium–Air Secondary Batteries, Electrochem. Solid-State Lett. 12 (2009) A132, doi:http://dx.doi.org/10.1149/1.3125285. [13] S.H. Jee, S.H. Lee, D.-J. Kim, J.-Y. Kwak, S.-C. Nam, Y.S. Yoon, Enhancement of Cycling Performance by Li2O–Sn Anode for All-Solid-State Batteries, Japanese Journal of Applied Physics 51 (2012) 85803, doi:http://dx.doi.org/10.1143/ jjap.51.085803. [14] S.H. Lee, S.H. Jee, K.S. Lee, S.C. Nam, Y.S. Yoon, Enhanced cycling performance in heat-treated tin-based composite oxide anode for lithium-ion batteries, Electrochim. Acta 87 (2013) 905, doi:http://dx.doi.org/10.1016/j. electacta.2012.09.002. [15] J.B. Bates, N.J. Dudney, K.A. Weatherspoon, Packaging material for thin film lithium batteries, Google Patents, 1996. [16] K.H.T. Raman, T.R. Penki, N. Munichandraiah, G.M. Rao, Titanium nitride thin film anode: chemical and microstructural evaluation during electrochemical studies, Electrochim. Acta 125 (2014) 282, doi:http://dx.doi.org/10.1016/j. electacta.2014.01.104. [17] J.M. Lee, H.-S. Hwang, W.-I. Cho, B.-W. Cho, K.Y. Kim, Effect of silver cosputtering on amorphous V2O5 thin-films for microbatteries, J. Power Sources 136 (2004) 122, doi:http://dx.doi.org/10.1016/j.jpowsour.2004.05.051. [18] S.-J. Lee, H.-K. Baik, S.-M. Lee, An all-solid-state thin film battery using LISIPON electrolyte and Si–V negative electrode films, Electrochem. Commun. 5 (2003) 32, doi:http://dx.doi.org/10.1016/s1388-2481(02) 528-3. [19] H. Li, H. Bai, Z. Tao, J. Chen, Si–Y multi-layer thin films as anode materials of high-capacity lithium-ion batteries, J. Power Sources 217 (2012) 102, doi: http://dx.doi.org/10.1016/j.jpowsour.2012.05.080.

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