Confined Solid Electrolyte Interphase Growth Space with Solid

Apr 4, 2017 - Confined Solid Electrolyte Interphase Growth Space with Solid. Polymer Electrolyte in Hollow Structured Silicon Anode for Li-Ion. Batter...
2 downloads 0 Views 7MB Size
Download PDF
Research Article www.acsami.org

Confined Solid Electrolyte Interphase Growth Space with Solid Polymer Electrolyte in Hollow Structured Silicon Anode for Li-Ion Batteries Tianyi Ma, Xiangnan Yu, Xiaolu Cheng, Huiyu Li, Wentao Zhu, and Xinping Qiu* Key Laboratory of Organic Optoelectronics and Molecular Engineering, Department of Chemistry, Tsinghua University, Beijing 100084, China S Supporting Information *

ABSTRACT: Silicon anodes for lithium-ion batteries are of much interest owing to their extremely high specific capacity but still face some challenges, especially the tremendous volume change which occurs in cycling and further leads to the disintegration of electrode structure and excessive growth of solid electrolyte interphase (SEI). Here, we designed a novel approach to confine the inward growth of SEI by filling solid polymer electrolyte (SPE) into pores of hollow silicon spheres. The as-prepared composite delivers a high specific capacity of more than 2100 mAh g−1 and a long-term cycle stability with a reversible capacity of 1350 mAh g−1 over 500 cycles. The growing behavior of SEI was investigated by electrochemical impedance spectroscopy and differential scanning calorimetry, and the results revealed that SPE occupies the major space of SEI growth and thus confines its excessive growth, which significantly improves cycle performance and Coulombic efficiency of cells embracing hollow silicon spheres. KEYWORDS: solid polymer electrolyte, hollow silicon anode, Coulombic efficiency, lithium ion battery, solid electrolyte interphase

■

graphene−silicon composites21−23 demonstrated great structural stability and cycling performance in half cells. However, porous structured materials are less successful in battery applications due to the low Coulombic efficiency. Normally, the first cycle Coulombic efficiency is generally 65−85% for silicon or silicon composite anodes, much lower than that of commercialized graphite anode (90−94%). The Coulombic efficiency of the subsequent cycles is generally in the range of 98−99.7%, lower than the needs for the full LIBs (99.9%).24 Although the prelithiated method can improve the first cycle Coulombic efficiency to 95%,25−27 the Coulombic efficiency of subsequent cycles is still difficult to improve higher than 99.9%. Continuous reconstruction of SEI on the surface of silicon is the main reason for the low Coulombic efficiency.28 Due to the low Young’s moduli (0−3 GPa),29 SEI is not so flexible to endure volume expansion and shrinking of silicon but cracks and continuously reconstructs due to the exposure of silicon with liquid electrolyte during battery cycles. To date, researchers have made great efforts to have a better understanding of the evolution of SEI on the surface of silicon. Guyomard et al. combined X-ray photoelectron spectroscopy (XPS), magic-angle spinning nuclear magnetic resonance (MAS

INTRODUCTION Lithium-ion batteries (LIBs) have been widely used as a power source for electric vehicles (EVs) because of their high energy density and long cycle life. To extend the travel distance of EVs, the development of LIBs with higher energy density has recently become more and more urgent. The energy density of LIBs is largely dependent on the capacity and potential of the anode.1,2 Silicon has been regarded as one promising anode material due to its high theoretical specific capacity (3572 mAh g−1) and low charge−discharge potential (∼0.2 V, vs Li/ Li+).3−5 However, the tremendous volume change (larger than 300%) during the lithiation/delithation reactions limits its application in LIBs. To be specific, such volume change causes pulverization of silicon particles, disintegration of the electrode, and excessive growth of solid electrolyte interface (SEI), leading to low Coulombic efficiency and capacity fading during cycling.6,7 To alleviate these problems, researchers have developed various strategies to buffer the volume expansion of silicon lithiation reaction by using porous structured silicon or its composites. Porous structured silicon such as porous silicon spheres8,9 and freestanding macroporous silicon film10 displayed higher ability for accommodation of the large volume expansion.11 Silicon/carbon composites such as porous carbon−silicon,12−14 silicon−carbon core−shell capsules,15−17 carbon coated silicon arrays,18−20 and multilayer wafer structure © 2017 American Chemical Society

Received: March 2, 2017 Accepted: April 4, 2017 Published: April 4, 2017 13247

DOI: 10.1021/acsami.7b03046 ACS Appl. Mater. Interfaces 2017, 9, 13247−13254

Research Article

ACS Applied Materials & Interfaces NMR, 7Li and 19F), and time-of-flight secondary ion mass spectrometry (TOF-SIMS) methods to analyze the electrolyte decomposition reaction and SEI growth on the surface of silicon anode. They found that the composition of SEI changes with the thickness, and the excessive growth of SEI consumes electrolyte and lithium sources from the cathode in LIBs.30 Our group also combined electrochemical impedance spectroscopy (EIS) and differential scanning calorimetry (DSC) to analyze SEI growth on lithiated silicon anode, and an exothermic peak of SEI decomposition can be detected around 50−100 °C, which is enough to quantify the growth of SEI.31 To prevent the excessive growth of SEI, construction of a coating layer on silicon to avoid the direct contact between silicon and electrolyte was suggested.32,33 However, the Coulombic efficiency of these materials is still less than 99.9%, lower than the requirements of the full LIBs. Our previous work demonstrated that hollow structured silicon can prohibit the growth of SEI, which was confirmed by results of EIS and DSC.31 In principle, SEI needs space to grow, and the occupation of this space can prohibit the excessive growth of SEI. Herein, we filled Li+ conductive solid polymer electrolyte (SPE) into a hollow structured silicon spheres (named HSSi) to occupy the space of SEI. Schematic diagrams of HSSi and SPE−HSSi samples are illustrated in Figure 1, where SPE not only serves as

100 sccm for 2 h to deposit Si on CaCO3. After deposition of silicon, the CaCO3 template was removed by washing with hydrochloric acid (5 wt %). The silicon oxide on the surface of silicon was removed by hydrofluoric acid (5 wt %). Finally, the product was washed several times with ethanol and dried in a vacuum at 60 °C for 24 h. Preparation of SPE−HSSi Composite. PEO-based SPE solutions were prepared simply by dissolving 0.88 g of PEO (20 mmol) and 0.266 g (2.5 mmol) of LiClO4 into 50 mL of deionized water under stirring for 2 h. Next, 0.2 g of HSSi material was added into the SPE solution under stirring for 12 h. Finally, the product was collected and washed by deionized water several times to remove residual SPE outside of silicon and dried in a vacuum at 60 °C for 24 h. Structural Characterization. Morphologies of the as-prepared samples were observed by transmission electron microscopy (TEM, Hitachi H-7650B operating at 80 keV) and scanning electron microscopy (SEM, Zeiss Merlin). Fourier transform infrared (FTIR) spectra were recorded on a Varian 670-IR FT-IR spectrometer with KBr pellets. The N2 sorption measurements were carried out on Quantachrome NOVA 1000e at 77.3 K, and the specific surface area and pore size distribution were calculated by the Brunauer−Emmett− Teller (BET) and Barrett−Joyner−Halenda (BJH) methods, respectively. Thermal gravimetric analysis (TGA) was performed using a METTLER TOLEDO TGA instrument with STAR system. The temperature range was set from 25 to 1000 °C with a scanning rate of 2 °C min−1. X-ray photoelectron spectroscopy (XPS) measurements were conducted on a PHI Quantera SXM spectrometer equipped with a focused and monochromatized Al Kα radiation; the binding energy scale was calibrated using the C 1s peak at 284.8 eV. TOF-SIMS was performed with the use of TOF.SIMS 5 (ION-TOF), and sputtering rate was controlled to be ∼0.3 nm s−1. Mettler-Toledo calorimeter was employed for the DSC analysis of SEI. In detail, electrode materials were scratched from the copper foil in the Ar-filled glovebox and sealed into high-pressure stainless steel crucibles (gold plated, Mettler-Toledo) for DSC tests. Temperature range was controlled from 30 to 300 °C at heating rate of 5 °C min−1. Electrochemical Measurements. The testing electrodes were prepared by coating slurry consisting of active material (60 wt %), Super P Li (TIMCAL, 20 wt %), and poly(acrylic acid) (PAA) as binder (Alfa Aesar, 20 wt %) on copper foil. 2025-type coin cells were assembled with lithium foil as counter electrode. Nonaqueous solution of 1 mol L−1 LiPF6 (Guotai-Huarong New Chemical Materials Co., Ltd.) with ethylene carbonate/dimethyl carbonate/ethylmethyl carbonate (EC/DMC/EMC) in volume ratio of 1:1:1 was adopted as electrolyte. Galvanostatic curves were collected on a Neware battery test system with constant current density between 0 and 1.2 V vs Li/ Li+. The current density was set to be 200 mA g−1 in the first 3 cycles and 600 mA g−1 in the subsequent cycles. The specific capacity was calculated based on the loading weight of silicon. EIS were collected on a PARSTAT 2273 electrochemical workstation with a threeelectrode cell (EL-CELL, ECC-REF model), where Li foil was used as both reference electrode and counter electrode. Electrochemical AC potential was controlled by applied voltage of 5.0 mV over a frequency range of 105−0.01 Hz. Before each EIS test, the electrodes were discharged to 0.01 V and stabilized at open circuit for 2 h. All batteries were assembled in an Ar-filled glovebox.

Figure 1. Schematic diagram of the lithiation and SEI growth of (a) HSSi and (b) SPE−HSSi.

conductive phase for Li+34 but also acts as a filler to occupy the space of SEI. More significantly, the flexible SPE can serve as a buffer to accommodate the volume expansion of LixSi. Benefiting from the flexible conductive SPE and inner filled strategy, this SPE−HSSi composite demonstrates electrochemical performance much better than that of HSSi material, giving a reversible capacity of 1350 mAh g−1 in 500 cycles at a constant current of 600 mA g−1. Particularly, the Coulombic efficiency of SPE−HSSi in first two cycles are 79.9 and 95.5%, respectively, and then close to 99.99% at subsequent cycles. Moreover, the SPE filling method is very simple and environmentally friendly, possibly applicable to other hollow structured anode materials which suffer from SEI excessive growth.

■

■

RESULTS AND DISCUSSION HSSi samples were fabricated with a scarification templateCVD method. The commercialized CaCO3 nanoparticles were chosen as the template because of their abundance, nontoxicity, and low cost. The SPE−HSSi composite was prepared simply by soaking PEO-based SPE solution into pores of HSSi. Figure 2 shows the morphological features of the HSSi and SPE− HSSi. The as-prepared hollow spheres are uniform (Figure 2a), with wall thickness of ca. 15 nm and outer radius of ca. 55 nm. The inner radius of spheres is close to the size of the CaCO3 template (TEM image of CaCO3 template is supplied in Figure S1). The TEM image of SPE−HSSi (Figure 2b) shows strong

EXPERIMENTAL SECTION

Materials Synthesis. Preparation of Hollow Structured Silicon Spheres. HSSi was prepared by our previous method.31 In detail, 0.4 g (4 mmol) of commercial nano-CaCO3 (ShanXi NanoMaterials Technology Co., Ltd.) was ground and used as template for chemical vapor deposition (CVD). The deposition process was carried out in a tube furnace with gradient temperature from 480 to 440 °C. A mixed gas of 5 wt % SiH4 and 95 wt % Ar was introduced with a flow rate of 13248

DOI: 10.1021/acsami.7b03046 ACS Appl. Mater. Interfaces 2017, 9, 13247−13254

Research Article

ACS Applied Materials & Interfaces

FT-IR spectra of HSSi, SPE−HSSi, and PEO powders are illustrated in Figure 3a. The stronger absorption band at the wavenumber of 878 cm−1 corresponds to the Si−Hx bond of silicon.35,36 Two weaker peaks close to 1456 and 1350 cm−1 correspond to δC−H bond of PEO. The comparison of peak intensity indicates that SPE is filled into the hollow structure of HSSi. It is worth noting that a weak peak at about 800 cm−1 can be observed, which corresponds to the bending vibration of the Si−O−Si bends, indicating the existence of silicon oxide on the surface of silicon, consistent with our previous works.31 To further investigate the pore size distribution of asprepared samples, nitrogen gas sorption measurements were performed. The nitrogen adsorption and desorption isotherms of HSSi and SPE−HSSi are presented in Figure 3b. Both of them show a sharp capillary condensation step at high relative pressures (P/P0 = 0.8−0.99), indicating the existence of mesopores.37 The pore size distribution of the two samples is given in Figure 3c. It is clearly seen that the pore size of HSSi is distributed between 20 and 100 nm, consistent with the TEM observation, while pore size in SPE−HSSi is less than 40 nm, indicating the filling of SPE. Content of silicon can be estimated from TGA curves.38 TGA curves of SPE−HSSi and HSSi are shown in Figure 3d. Weight loss of SPE−HSSi before 400 °C is ascribed to the decomposition and oxidation of SPE. Weight increase between 400 and 900 °C is owing to the oxidation of silicon to SiO2. According to TGA, silicon content in SPE− HSSi composite is calculated to be ca. 65 wt %. Due to exposure to air during preparation, the silicon surface is usually oxidized.31 Here, we combined XPS and TOF-SIMS to analyze the surface oxidation state of silicon in the SPE− HSSi sample. XPS spectra of the Si 2p orbit is shown in Figure

Figure 2. TEM images of (a) HSSi and (b) SPE−HSSi. SEM images of (c) HSSi and (d) SPE−HSSi.

contrast inside the hollow structure, confirming the fill of SPE. The thickness of silicon in SPE−HSSi is similar to that of HSSi, indicating that residual SPE on the outer surface of silicon is removed. This can also be confirmed from SEM images of HSSi and SPE−HSSi (Figures 2c and d), where both of them show a smooth surface.

Figure 3. (a) FT-IR spectra of PEO, HSSi, and SPE−HSSi. (b and c) Nitrogen sorption isotherms (77.3 K) and pore size distribution test of HSSi and SPE−HSSi. (d) TGA curve in air atmosphere of HSSi and SPE−HSSi (25−1000 °C, 2 °C/min−1). 13249

DOI: 10.1021/acsami.7b03046 ACS Appl. Mater. Interfaces 2017, 9, 13247−13254

Research Article

ACS Applied Materials & Interfaces

Figure 4. (a) XPS analysis (Si 2p) of SPE−HSSi. (b) TOF-SIMS depth profiles of O+ and SiO+ in the SPE−HSSi sample. (c) Mass peak of O+ before (blue) and after (red) the depth profile. (d) Mass peak of SiO+ before (blue) and after (red) the depth profile.

Figure 5. Electrochemical performance of HSSi and SPE−HSSi samples. (a) Charge and discharge profiles of SPE−HSSi in the first and second cycles. (b) Plots of differential capacity of SPE−HSSi in different cycles. (c) Cycling performance of HSSi and SPE−HSSi samples; inset shows active materials after 500 cycles. (d) Coulombic efficiency of HSSi and SPE−HSSi sample.

13250

DOI: 10.1021/acsami.7b03046 ACS Appl. Mater. Interfaces 2017, 9, 13247−13254

Research Article

ACS Applied Materials & Interfaces

Figure 6. (a) Nyquist diagrams of the SPE−HSSi sample by different cycles. (b) Typical DSC curve of silicon anode materials after battery cycles. (c and d) DSC curves of HSSi and SPE−HSSi in 50−100 °C, respectively.

4a. The first main 3/2−1/2 doublet (the spin−orbit splitting is 0.6 eV and the intensity ratio is 3:1), located at 99.1−99.7 eV, corresponds to Si0 (63% content); peaks of silicon oxide (mainly Si2O3 and SiO2, ∼37% content) are also detected on the surface of the composite, indicating that surface of silicon in SPE−HSSi sample is partly oxidized. Figure 4b shows the TOF-SIMS curves of SPE−HSSi. O+ and SiO+ signals were selected to characterize the SiOx in the SPE−HSSi sample in depth by TOF-SIMS. The intensity of O+ and SiO+ relates to compositional variation. As shown in Figure 4b, the gradual decrease in intensity suggests that SiOx decreased with spurring depth of SPE−HSSi. The mass peaks of O+ and SiO+ before and after the sputtering process are shown in Figures 4c and d, respectively. The main peaks of O+ (u = 16.03) and SiO+ (u = 44.025) before sputtering clearly reveal that there exists SiOx on the surface of SPE−HSSi. These peaks disappear after a sputtering depth of 18 nm, indicating that the oxidation layer on the silicon surface is very thin. The electrochemical properties of the SPE−HSSi and HSSi were characterized by galvanostatic charge−discharge tests. All of the cells performed at constant current densities of 200 mA g−1 for the first three cycles and 600 mA g−1 for subsequent cycles, and the mass loading of the anode was controlled to be 1.9−2.0 mg cm−2. The initial charge−discharge profiles of SPE−HSSi are presented in Figure 5a. The sloping potential variation shows a typical lithiation reaction of amorphous silicon.39,40 A long flat plateau observed during the first discharge is caused by irreversible reactions related to the formation of SEI and disappears in the second cycle. The specific capacity of the first discharge and charge are 2161 and 1727 mAh g−1, respectively. The irreversible capacity of the first cycle is 434 mAh g−1, corresponding to a Coulombic efficiency of 79.9%, much higher than that of HSSi (71.5%). Furthermore,

the second cycle Coulombic efficiency of SPE−HSSi reaches 95.5%, considerably higher than values reported for siliconbased anodes.3 Figure 5b shows the differential capacity curves of SPE−HSSi after 1st, 2nd, 50th, and 100th cycles, respectively. A peak at 0.2 V can be observed corresponding to the formation of Li−Si-based amorphous phase (α-LixSi) by lithiated alloy reaction.41,42 Peaks at 0.32 and 0.49 V in the first charge correspond to delithiation reaction, which is consistent with previous research works.43−45 Moreover, the almost identical peak values of 50 and 100 cycles imply little capacity loss during cycling. Figure 5c shows cyclability of SPE−HSSi, where the result of HSSi is also shown for comparison. SPE−HSSi attained a reversible discharge capacity of 1350 mA h g−1 and capacity retention of 63% after 500 cycles, which are much higher than the reversible capacity of 970 mA h g−1 and the capacity retention of 35% for HSSi. To obtain more evidence of the structural stability of SPE−HSSi, the electrodes of SPE−HSSi and HSSi after 500 cycles were removed and cleaned by dimethyl carbonate and acetic acid (2 wt %); the TEM images are displayed in the inset of Figure 5c. Nanosphere structures of SPE−HSSi and HSSi were maintained without fractures or cracks. The SPE filler (strong contrast areas) inside pores can be observed clearly, indicating stability of SPE during the long cycles. In addition to the high specific capacity and great cycle stability of SPE−HSSi, it also demonstrates better cycling Coulombic efficiencies, as shown in Figure 5d. The SPE−HSSi delivered Coulombic efficiency higher than that of the HSSi anode, and the Coulombic efficiency was close to 99.99% in the subsequent cycles. The high specific capacity and Coulombic efficiency imply that the SPE filler can not only buffer the volume expansion of the silicon alloying reaction but also effectively confine space for the SEI growth. 13251

DOI: 10.1021/acsami.7b03046 ACS Appl. Mater. Interfaces 2017, 9, 13247−13254

Research Article

ACS Applied Materials & Interfaces

The EIS and DSC results further confirmed that the internal space occupation is effective to prevent the excessive growth of SEI. Additionally, the SPE filling method is facile and easily scaled up, which has a great potential for battery applications.

To obtain more evidence of growth confinement of SEI in SPE−HSSi, EIS were collected after different cycles. A threeelectrode cell was assembled with SPE−HSSi as working electrode and lithium foil as counter and reference electrode. After being charged−discharged for specific cycles, the cell was discharged to 0.01 V and subsequently rested for 1 h to stabilize the potential before EIS. Figure 6a illustrates the Nyquist plots of SPE−HSSi at the end of lithiation process for different cycles. The semicircle in the midfrequency district corresponds to the interfacial phenomena, including charge transfer and SEI between active particles and the liquid electrolyte.46 The impedance tail in the low-frequency region can be attributed to the bulk diffusional effects. To evaluate the whole surface resistance, a quantity of RSur was adopted (equivalent circuit is shown in inset of Figure 6a), and its fitting values are plotted in Figure S2. The value of RSur increases after the initial 20 cycles and maintains at ca. 50 Ohm in subsequent cycles, indicating the stability of SEI during cycling. Compared with the Nyquist plots of electrodes of silicon nanoparticle (50 nm) of our previous work,44 whose RSur increases continuously with cycle number, it can be concluded that SEI of SPE−HSSi is stabilized due to the confinement function of SPE. To further investigate the SEI growth of SPE−HSSi during cycles, we used DSC to estimate the amount of SEI on the surface of silicon directly. In detail, batteries were discharged at a constant current density of 100 mA g−1 to the potential of 0.01 V and rested for 2 h to stabilize the potential. Then, cells were disassembled, and the electrode materials were scratched from the copper foil in the Ar-filled glovebox and sealed into high-pressure stainless steel crucibles for DSC. Figure 6b is a typical DSC curve of silicon after cycling. The exothermic peak at 50−100 °C is associated with the reaction between ROCOOLi (or ROLi) in SEI and LixSi,47 which represents the reaction between SEI and LixSi. The exothermic peaks located in 100−200 °C represent the phase transformation of amorphous LixSi, where the shape and position of the peaks are related to lithium content (x-value) .47,48 Figure 6c displays the lithiated HSSi DSC curves in the temperature range of 50−100 °C. The heat flux at 70 °C is 0.105 W g−1 after the first cycle and increases to 0.155, 0.168, and 0.178 W g−1, after 10, 30, and 50 cycles, respectively, indicating SEI growth during cycling. By comparing these with DSC tests of silicon nanoparticles in our previous work,31 where the heat flux increases to 0.285 W g−1 after 50 cycles, we can conclude that the amount of SEI on HSSi is less than that on silicon nanoparticles after 50 cycles. For comparison, DSC curves of lithiated SPE−HSSi are shown in Figure 6d. The heat fluxes of SPE−HSSi at 70 °C are 0.104, 0.108, 0.121, 0.134W g−1 after 1, 10, 30, and 50 cycles, respectively. The lower heat flux of SPE−HSSi means less SEI growth during cycles, which suggests SPE filling can effectively confine the growth of SEI. Further, the endothermic peak of PEO (around 65 °C) can be observed in the DSC curve of SPE−HSSi (DSC test of pure PEO powder is shown in the Figure S3), which further confirms that SPE was filled into the pore of hollow structured silicon.

■

ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acsami.7b03046. TEM image of the CaCO3 template, fitting results of RSur for different cycles, and DSC curve of PEO powder (PDF)

■

AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]; Tel.: +86-10-62794234; Fax: +86-10-62794234. ORCID

Huiyu Li: 0000-0003-2462-1170 Xinping Qiu: 0000-0001-5291-7943 Notes

The authors declare no competing financial interest.

■

ACKNOWLEDGMENTS The authors appreciate financial support from the National Key Project on Basic Research (Grant 2015CB251104), the Natural Science Foundation of China (Grant U1664256), and the China−United States Electric Vehicle Project (Grant S2016G9004).

■

REFERENCES

(1) Goodenough, J. B.; Kim, Y. Challenges for Rechargeable Li Batteries. Chem. Mater. 2010, 22, 587−603. (2) Arico, A. S.; Bruce, P.; Scrosati, B.; Tarascon, J. M.; Van Schalkwijk, W. Nanostructured Materials for Advanced Energy Conversion and Storage Devices. Nat. Mater. 2005, 4, 366−377. (3) Rahman, M. A.; Song, G.; Bhatt, A. I.; Wong, Y. C.; Wen, C. Nanostructured Silicon Anodes for High-Performance Lithium-Ion Batteries. Adv. Funct. Mater. 2016, 26, 647−678. (4) Su, X.; Wu, Q. L.; Li, J. C.; Xiao, X. C.; Lott, A.; Lu, W. Q.; Sheldon, B. W.; Wu, J. Silicon-Based Nanomaterials for Lithium-Ion Batteries: A Review. Adv. Energy Mater. 2014, 4, 1300882−1300904. (5) Botas, C.; Carriazo, D.; Zhang, W.; Rojo, T.; Singh, G. SiliconReduced Graphene Oxide Self-Standing Composites Suitable as Binder-Free Anodes for Lithium-Ion Batteries. ACS Appl. Mater. Interfaces 2016, 8, 28800−28808. (6) Goldman, J. L.; Long, B. R.; Gewirth, A. A.; Nuzzo, R. G. Strain Anisotropies and Self-Limiting Capacities in Single-Crystalline 3D Silicon Microstructures: Models for High Energy Density Lithium-Ion Battery Anodes. Adv. Funct. Mater. 2011, 21, 2412−2422. (7) Liu, X. H.; Zhong, L.; Huang, S.; Mao, S. X.; Zhu, T.; Huang, J. Y. Size-Dependent Fracture of Silicon Nanoparticles During Lithiation. ACS Nano 2012, 6, 1522−1531. (8) Ababtain, K.; Babu, G.; Lin, X.; Rodrigues, M.-T. F.; Gullapalli, H.; Ajayan, P. M.; Grinstaff, M. W.; Arava, L. M. R. Ionic LiquidOrganic Carbonate Electrolyte Blends to Stabilize Silicon Electrodes for Extending Lithium Ion Battery Operability to 100 Degrees C. ACS Appl. Mater. Interfaces 2016, 8, 15242−15249. (9) Ge, M.; Rong, J.; Fang, X.; Zhang, A.; Lu, Y.; Zhou, C. Scalable Preparation of Porous Silicon Nanoparticles and Their Application for Lithium-Ion Battery Anodes. Nano Res. 2013, 6, 174−181. (10) Thakur, M.; Pernites, R. B.; Nitta, N.; Isaacson, M.; Sinsabaugh, S. L.; Wong, M. S.; Biswal, S. L. Freestanding Macroporous Silicon and

■

CONCLUSION In summary, we successfully developed a new method to improve the reversible stability and Coulombic efficiency of hollow structured silicon by filling PEO-based SPE. The asprepared SPE−HSSi composite demonstrated a high initial specific capacity of >2100 mAh g−1 and cycle stability with a reversible specific capacity of >1350 mAh g−1 over 500 cycles. 13252

DOI: 10.1021/acsami.7b03046 ACS Appl. Mater. Interfaces 2017, 9, 13247−13254

Research Article

ACS Applied Materials & Interfaces Pyrolyzed Polyacrylonitrile As a Composite Anode for Lithium Ion Batteries. Chem. Mater. 2012, 24, 2998−3003. (11) Ge, M.; Fang, X.; Rong, J.; Zhou, C. Review of Porous Silicon Preparation and Its Application for Lithium-Ion Battery Anodes. Nanotechnology 2013, 24, 2001−2011. (12) Wu, L. L.; Yang, J.; Zhou, X. Y.; Zhang, M. F.; Ren, Y. P.; Nie, Y. Silicon Nanoparticles Embedded in a Porous Carbon Matrix As A High-Performance Anode For Lithium-Ion Batteries. J. Mater. Chem. A 2016, 4, 11381−11387. (13) Jeong, M.-G.; Islam, M.; Du, H. L.; Lee, Y.-S.; Sun, H.-H.; Choi, W.; Lee, J. K.; Chung, K. Y.; Jung, H.-G. Nitrogen-doped Carbon Coated Porous Silicon as High Performance Anode Material for Lithium-Ion Batteries. Electrochim. Acta 2016, 209, 299−307. (14) Roy, A. K.; Zhong, M.; Schwab, M. G.; Binder, A.; Venkataraman, S. S.; Tomovic, Z. Preparation of a Binder-Free Three-Dimensional Carbon Foam/Silicon Composite as Potential Material for Lithium Ion Battery Anodes. ACS Appl. Mater. Interfaces 2016, 8, 7343−7348. (15) Luo, J. Y.; Zhao, X.; Wu, J. S.; Jang, H. D.; Kung, H. H.; Huang, J. X. Crumpled Graphene-Encapsulated Si Nanoparticles for Lithium Ion Battery Anodes. J. Phys. Chem. Lett. 2012, 3, 1824−1829. (16) Yue, X.; Sun, W.; Zhang, J.; Wang, F.; Sun, K. Facile Synthesis of 3D Silicon/Carbon Nanotube Capsule Composites as Anodes For High-Performance Lithium-Ion Batteries. J. Power Sources 2016, 329, 422−427. (17) Zhou, J.; Qian, T.; Wang, M.; Xu, N.; Zhang, Q.; Li, Q.; Yan, C. Core-Shell Coating Silicon Anode Interfaces with Coordination Complex for Stable Lithium-Ion Batteries. ACS Appl. Mater. Interfaces 2016, 8, 5358−5365. (18) Lu, Z.; Wong, T.; Ng, T.-W.; Wang, C. Facile Synthesis of Carbon Decorated Silicon Nanotube Arrays as Anode Material for High-Performance Lithium-Ion Batteries. RSC Adv. 2014, 4, 2440− 2446. (19) Wang, W.; Kumta, P. N. Nanostructured Hybrid Silicon/ Carbon Nanotube Heterostructures: Reversible High-Capacity Lithium-Ion Anodes. ACS Nano 2010, 4, 2233−2241. (20) Pandey, G. P.; Klankowski, S. A.; Li, Y.; Sun, X. S.; Wu, J.; Rojeski, R. A.; Li, J. Effective Infiltration of Gel Polymer Electrolyte into Silicon-Coated Vertically Aligned Carbon Nanofibers as Anodes for Solid-State Lithium-Ion Batteries. ACS Appl. Mater. Interfaces 2015, 7, 20909−20918. (21) Chen, S.; Bao, P.; Huang, X.; Sun, B.; Wang, G. Hierarchical 3D Mesoporous Silicon@Graphene Nanoarchitectures for Lithium Ion Batteries With Superior Performance. Nano Res. 2014, 7, 85−94. (22) Sun, W.; Hu, R.; Zhang, M.; Liu, J.; Zhu, M. Binding of Carbon Coated Nano-Silicon in Graphene Sheets by Wet Ball-Milling and Pyrolysis as High Performance Anodes For Lithium-Ion Batteries. J. Power Sources 2016, 318, 113−120. (23) Gao, X.; Li, J.; Xie, Y.; Guan, D.; Yuan, C. A Multi Layered Silicon-Reduced Graphene Oxide Electrode for High Performance Lithium-Ion Batteries. ACS Appl. Mater. Interfaces 2015, 7, 7855− 7862. (24) Park, S.-H.; Ahn, D.; Choi, Y.-M.; Roh, K. C.; Kim, K.-B. Highcoulombic-efficiency Si-based Hybrid Microspheres Synthesized by the Combination of Graphene and IL-derived Carbon. J. Mater. Chem. A 2015, 3, 20935−20943. (25) Zhao, J.; Lu, Z.; Wang, H.; Liu, W.; Lee, H.-W.; Yan, K.; Zhuo, D.; Lin, D.; Liu, N.; Cui, Y. Artificial Solid Electrolyte InterphaseProtected LixSi Nanoparticles: An Efficient and Stable Prelithiation Reagent for Lithium-Ion Batteries. J. Am. Chem. Soc. 2015, 137, 8372− 8375. (26) Zhou, H.; Wang, X.; Chen, D. Li-Metal-Free Prelithiation of SiBased Negative Electrodes for Full Li-Ion Batteries. ChemSusChem 2015, 8, 2737−2744. (27) Kim, H. J.; Choi, S.; Lee, S. J.; Seo, M. W.; Lee, J. G.; Deniz, E.; Lee, Y. J.; Kim, E. K.; Choi, J. W. Controlled Prelithiation of Silicon Monoxide for High Performance Lithium-Ion Rechargeable Full Cells. Nano Lett. 2016, 16, 282−288.

(28) Nie, M.; Abraham, D. P.; Chen, Y.; Bose, A.; Lucht, B. L. Silicon Solid Electrolyte Interphase (SEI) of Lithium Ion Battery Characterized by Microscopy and Spectroscopy. J. Phys. Chem. C 2013, 117, 13403−13412. (29) Zheng, J.; Zheng, H.; Wang, R.; Ben, L.; Lu, W.; Chen, L.; Chen, L.; Li, H. 3D Visualization of Inhomogeneous Multi-Layered Structure and Young’s Modulus of the Solid Electrolyte Interphase (SEI) on Silicon Anodes for Lithium Ion Batteries. Phys. Chem. Chem. Phys. 2014, 16, 13229−13238. (30) Dupre, N.; Moreau, P.; De Vito, E.; Quazuguel, L.; Boniface, M.; Bordes, A.; Rudisch, C.; Bayle-Guillemaud, P.; Guyomard, D. Multiprobe Study of the Solid Electrolyte Interphase on SiliconBased Electrodes in Full-Cell Configuration. Chem. Mater. 2016, 28, 2557−2572. (31) Lv, Q.; Liu, Y.; Ma, T.; Zhu, W.; Qiu, X. Hollow Structured Silicon Anodes with Stabilized Solid Electrolyte Interphase Film for Lithium-Ion Batteries. ACS Appl. Mater. Interfaces 2015, 7, 23501− 23506. (32) Zhou, Y.; Guo, H.; Yang, Y.; Wang, Z.; Li, X.; Zhou, R.; Peng, W. Facile Synthesis of Silicon/Carbon Nanospheres Composite Anode Materials for Lithium-Ion Batteries. Mater. Lett. 2016, 168, 138−142. (33) Mi, H.; Li, F.; Xu, S.; Li, Z.; Chai, X.; He, C.; Li, Y.; Liu, J. A Tremella-Like Nanostructure of Silicon@void@graphene-Like Nanosheets Composite as an Anode for Lithium-Ion Batteries. Nanoscale Res. Lett. 2016, 11, 204−212. (34) Vignarooban, K.; Dissanayake, M. A. K. L.; Albinsson, I.; Mellander, B. E. Effect of TiO2 Nano-filler and EC Plasticizer on Electrical and Thermal Properties of Poly(ethylene oxide) (PEO) Based Solid Polymer Electrolytes. Solid State Ionics 2014, 266, 25−28. (35) Mawhinney, D. B.; Glass, J. A.; Yates, J. T. FTIR Study of the Oxidation of Porous Silicon. J. Phys. Chem. B 1997, 101, 1202−1206. (36) Sun, X. H.; Wang, S. D.; Wong, N. B.; Ma, D. D. D.; Lee, S. T.; Teo, B. K. FTIR Spectroscopic Studies of the Stabilities and Reactivities of Hydrogen-Terminated Surfaces of Silicon Nanowires. Inorg. Chem. 2003, 42, 2398−2404. (37) Li, W.; Chen, D.; Li, Z.; Shi, Y.; Wan, Y.; Wang, G.; Jiang, Z.; Zhao, D. Nitrogen-containing Carbon Spheres with Very Large Uniform Mesopores: The Superior Electrode Materials for EDLC in Organic Electrolyte. Carbon 2007, 45, 1757−1763. (38) Zhu, X. Y.; Chen, H.; Wang, Y. H.; Xia, L. H.; Tan, Q. Q.; Li, H.; Zhong, Z. Y.; Su, F. B.; Zhao, X. S. Growth of Silicon/Carbon Microrods on Graphite Microspheres as Improved Anodes for Lithium-Ion Batteries. J. Mater. Chem. A 2013, 1, 4483−4489. (39) Song, J.; Zhou, M.; Yi, R.; Xu, T.; Gordin, M. L.; Tang, D.; Yu, Z.; Regula, M.; Wang, D. Interpenetrated Gel Polymer Binder for High-Performance Silicon Anodes in Lithium-ion Batteries. Adv. Funct. Mater. 2014, 24, 5904−5910. (40) Xiaopeng, L.; Chenglin, Y.; Wang, J.; Graff, A.; Schweizer, S. L.; Sprafke, A.; Schmidt, O. G.; Wehrspohn, R. B. Stable Silicon Anodes for Lithium-Ion Batteries Using Mesoporous Metallurgical Silicon. Adv. Energy Mater. 2015, 5, 1401556−1401561. (41) Chen, X.; Gerasopoulos, K.; Guo, J.; Brown, A.; Wang, C.; Ghodssi, R.; Culver, J. N. A Patterned 3D Silicon Anode Fabricated by Electrodeposition on a Virus-Structured Current Collector. Adv. Funct. Mater. 2011, 21, 380−387. (42) Zhou, Z.-W.; Liu, Y.-T.; Xie, X.-M.; Ye, X.-Y. Constructing Novel Si@SnO2 Core-Shell Heterostructures by Facile Self-Assembly of SnO2 Nanowires on Silicon Hollow Nanospheres for Large, Reversible Lithium Storage. ACS Appl. Mater. Interfaces 2016, 8, 7092−7100. (43) Feng, W.; Lijun, W.; Key, B.; Xiao-Qing, Y.; Grey, C. P.; Yimei, Z.; GraetzGraetz, J. Electrochemical Reaction of Lithium with Nanostructured Silicon Anodes: A Study by In-Situ Synchrotron XRay Diffraction and Electron Energy-Loss Spectroscopy. Adv. Energy Mater. 2013, 3, 1324−1331. (44) Liu, Y.; Guo, X.; Li, J.; Lv, Q.; Ma, T.; Zhu, W.; Qiu, X. Improving Coulombic Efficiency by Confinement of Solid Electrolyte Interphase Film in Pores Of Silicon/Carbon Composite. J. Mater. Chem. A 2013, 1, 14075−14079. 13253

DOI: 10.1021/acsami.7b03046 ACS Appl. Mater. Interfaces 2017, 9, 13247−13254

Research Article

ACS Applied Materials & Interfaces (45) Oumellal, Y.; Delpuech, N.; Mazouzi, D.; Dupre, N.; Gaubicher, J.; Moreau, P.; Soudan, P.; Lestriez, B.; Guyomard, D. The Failure Mechanism of Nano-sized Si-based Negative Electrodes for Lithium Ion Batteries. J. Mater. Chem. 2011, 21, 6201−6208. (46) Zhang, L.; Rajagopalan, R.; Guo, H.; Hu, X.; Dou, S.; Liu, H. A Green and Facile Way to Prepare Granadilla-Like Silicon-Based Anode Materials for Li-Ion Batteries. Adv. Funct. Mater. 2016, 26, 440−446. (47) Profatilova, I. A.; Langer, T.; Badillo, J. P.; Schmitz, A.; Orthner, H.; Wiggers, H.; Passerini, S.; Winter, M. Thermally Induced Reactions between Lithiated Nano-Silicon Electrode and Electrolyte for Lithium-Ion Batteries. J. Electrochem. Soc. 2012, 159, A657−A663. (48) Wang, Y.; Dahn, J. Phase Changes in Electrochemically Lithiated Silicon at Elevated Temperature. J. Electrochem. Soc. 2006, 153, A2314−A2318.

13254

DOI: 10.1021/acsami.7b03046 ACS Appl. Mater. Interfaces 2017, 9, 13247−13254