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Dual Core−Shell Structured Si@SiOx@C Nanocomposite Synthesized via a One-Step Pyrolysis Method as a Highly Stable Anode Material for Lithium-Ion Batteries Bolun Jiang,† Shi Zeng,† Hui Wang,† Daotan Liu,‡ Jiangfeng Qian,† Yuliang Cao,† Hanxi Yang,† and Xinping Ai*,† †
Hubei Key Lab. of Electrochemical Power Sources, College of Chemistry and Molecular Science, Wuhan University, Wuhan 430072, China ‡ China Electric Power Research Institute, Beijing, 100192, China S Supporting Information *
ABSTRACT: Silicon (Si) has been regarded as a promising high-capacity anode material for developing advanced lithiumion batteries (LIBs), but the practical application of Si anodes is still unsuccessful mainly due to the insufficient cyclability. To deal with this issue, we propose a new route to construct a dual core−shell structured Si@SiOx@C nanocomposite by direct pyrolysis of poly(methyl methacrylate) (PMMA) polymer on the surface of Si nanoparticles. Since the PMMA polymers can be chemically bonded on the nano-Si surface through the interaction between ester group and Si surface group, and thermally decomposed in the subsequent pyrolysis process with their alkyl chains converted to carbon and the residue oxygen recombining with Si to form SiOx, the dual core−shell structure can be conveniently formed in a one-step procedure. Benefiting from the strong buffering effect of the SiOx interlayer and the efficient blocking action of dense outer carbon layer in preventing electrolyte permeation, the obtained nanocomposite demonstrates a high capacity of 1972 mA h g−1, a stable cycling performance with a capacity retention of >1030 mA h g−1 over 500 cycles, and particularly a superiorly high Coulombic efficiency of >99.5% upon extended cycling, exhibiting a great promise for practical uses. More importantly, the synthetic method proposed in this work is facile and low cost, making it more suitable for large-scale production of high capacity anode for advanced LIBs. KEYWORDS: Li-ion battery, Si anode, dual core−shell, poly(methyl methacrylate) (PMMA), pyrolytic synthesis
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INTRODUCTION Developing advanced lithium ion batteries (LIBs) with considerably improved energy density and cyclability are in urgent demands for the light-weighting of portable electronics and the range-extending of electric vehicles. Under such a technological background, Si anode material has received increasing attention because of its relatively low delithiation potential (∼0.4 V vs Li+/Li) and extra-high theoretical capacity of 4200 mAh g−1, which is at least more than ten times higher than that of the commercial graphite anodes currently used in LIBs.1 As revealed by recent literatures, the energy density of LIBs can be remarkably improved by at least 30% on the original basis even if the graphite anode is replaced by a Si anode with capacity of only 1000 mAh g−1, and the future LIBs based on Si anodes are expected to achieve a considerably high energy density of >400 Wh kg1− on the cell level.2,3 Nevertheless, most Si anodes reported so far still suffer from poor cycle life originating from the detrimental volume expansion/contraction (∼300%) during lithiation/delithiation cycles.4 Such a huge volume change can cause pulverization of Si particles, electrical disconnection of active components in electrode framework, as well as repeated destructions/ © XXXX American Chemical Society
reconstructions of the surface solid electrolyte interphase (SEI) film, thus leading to rapid capacity fading upon cycling.5 To address these problems, many attempts have been made to alleviate the volume change of Si anodes during lithiation by nanosizing of Si particles,6,7 fabricating porous nanoarchitectures,8,9 using Si-based alloys as active phase10,11 and embedding nano-Si in less-active/inactive buffering matrices,5,12 so as to minimize the mechanical stress and improve the structure stability on particle and electrode levels. Besides, building a protective carbon coating on the Si nanoarchitectures has become a generally adopted method to extend the cycling life of Si anodes. On the one hand, carbon coating can prevent aggregation and improve the electrical conductivity of nano-Si particles; On the other hand, carbon coating can act as both buffering and barrier layer to release the mechanical stress and prevent the electrolyte from reaching the inner Si core, thus stabilizing the electrode/electrolyte interface and depressing the reductive decomposition of organic electrolyte on Si surface. Received: August 4, 2016 Accepted: November 1, 2016
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DOI: 10.1021/acsami.6b09775 ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX
Research Article
ACS Applied Materials & Interfaces
core−shelled Si@SiOx@C nanocomposite. For comparison, other two Si@C nanocomposites were also synthesized in a similar way using polyacrylonitrile (PAN) and sucrose as carbon precursors, respectively. The weight ratios of Si to carbon precursor were designed as 1:2.5. Structural Characterization. The morphologies and structure features were analyzed by scanning-electron microscope (SEM, Zeiss, MERLIN Compact) and transmission electron microscopy (TEM, JEOL, JEM-2010-FEF). The FT-IR spectra of Si@PMMA composite and PMMA polymer were obtained on a Continuum FT-IR spectrometer using KBr pellets. Thermalgravimetric analyses (TG) of Si@PMMA composite and PMMA polymer were carried out using a TGA Q500 under N2, while the TG curves of Si@SiOx@C composite and Si particles were recorded under air. The Si@SiOx@C composite and Si particles were also characterized by X-ray photoelectron spectroscopy (XPS, Thermo Fisher, ESCALAB 250Xi). The BET (Brunauer−Emmett−Teller) specific surface area of the Si/C nanocomposites was derived from the nitrogen adsorption−desorption isotherms, which were measured on ASAP 2020. Electrochemical Measurements. Coin-type half cells (2016type) were assembled in an argon filled glovebox to evaluate the electrochemical properties of the as-prepared Si@SiOx@C nanocomposite electrodes. For preparing working electrodes, a mixture of the active material, carbon black (Super P) and poly(acrylic acid) (PAA) binder with a weight ratio of 6:2:2 was added in distilled water to form a slurry, which was then pasted onto a 20 mm thick copper foil and dried at 80 ◦C under a vacuum for overnight to remove the water. A Li disk and a Celgard 2400 microporous membrane were used as counter electrode and separator, respectively. The electrolyte was 1 M LiPF6 dissolving in a mixture solvent of ethylene carbonate (EC), dimethyl carbonate (DMC) and ethyl methyl carbonate (EMC) (1:1:1, by volume, Shinestar Battery Materials Co., Ltd., China). The mass loading of the Si@SiOx@C nanocomposites in the electrode was ∼1.5 mg cm−2. The PAA binder used in this study was purchase from Alfa Asear with an average molecule weight of 240 000. The electrochemical performances were conducted in a voltage interval of 0.01 to 1.5 V on Land Battery Testing System (Wuhan Kingnuo Electronics Co., Ltd., China) at 25 °C. Except as specifically noted, the Si electrodes were first discharged and charged at 100 mA g−1 for one cycle and then cycled at 500 mA g−1 for the following cycles.
Along with these strategies, various core−shell Si/C nanocomposites were reported to exhibit remarkably high capacity and rather stable cycling performance in lithium-half cells.13−16 However, the structure stability of these core−shell nanocomposites is still problematic for their battery applications. At the repeated action of volume expansions/contractions, the carbon shells only with a weak interface bonding are easily cracked and detached from the Si surface. This situation allows the electrolyte to permeate into and direct contact with the inner core, and then undergo a successive reductive decomposition on the constantly changed Si surface during cycling, giving rise to low Coulombic efficiencies and ceaseless consumption of Li ions. Since the cathodes in full batteries can not provide an enough amount of Li ions to compensate for such an irreversible consumption like Li counter electrode in half cells, the Si anodes often cause a fast capacity fading in practical batteries. To enhance the structure stability, Si/C composites with dual core−shell structure have been elaborately designed and fabricated in different ways.17−19 For instance, Zhang et al. prepared a double-walled Si@SiO2@C nanocomposite by calcination of silicon nanoparticles in air and subsequent carbon coating, which demonstrates a stable capacity of about 786 mAh g−1 over 100 cycles.20 Fan et al. synthesized Si@SiOx/ C nanoporous spheres, which have a stable capacity of 913 mA h g −1 and a good capacity retention of 97% after 60 cycles, through magnesiothermic reduction of mesoporous silica followed by carbonization of carbon precursor.21 This improved Li-storage behavior should be attributed to the double-shelled structure, in which the SiOx interlayer serves not only as a rigid buffering layer to confine the volume expansion and release the mechanical strain, but also as a bonding layer to enhance the interfacial adhesion between carbon shell and silicon core, thus maintaining the integrity of core−shell structures at cycling.22 However, the double shells in these nanocomposites are generally formed through multiple synthetic steps with rather complicated procedures. Herein, we propose a new, facile and large-scale method for preparation of dual core−shelled Si@SiOx@C nanocomposite in a one-step procedure by a polymer pyrolysis method using poly(methyl methacrylate) (PMMA) polymer as both carbon and oxygen source, and nanosized Si as Si source. Benefiting from the strong buffering effect of the SiOx interlayer and the efficient blocking action of dense outer carbon layer to electrolyte permeation, the obtained nanocomposite demonstrates a high capacity of 1972 mA h g−1 with an initial Coulombic efficiency of 77%, a long-term cyclability with a retained capacity of 1030 mA h g−1 over 500 cycles. Particularly, a superiorly high Coulombic efficiency of >99.5% is also obtained upon extended cycling, showing a great promise for battery application.
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RESULTS AND DISCUSSION In this work, we chose PMMA polymer as precursor to construct dual shells for Si nanoparticles, mainly because of its unique thermal-degradation behavior on the Si surface. In general, PMMA can be completely degraded into volatile molecular species with a weight loss of 100% under inert atmosphere at temperature above 300 ◦C, and therefore was widely used as pore former and organic template in material processing.23,24 But on the α-Al2O3 ceramics substrate, it was found that PMMA can be pyrolysized to yield a small amount of residual carbon, due to the bonding interaction between ester group on PMMA and Al2O3 surface hydroxyl groups, which immobilizes a part of carbon to avoid the complete escaping of thermally degraded species from the substrate surface.25 Considering the analogous surface chemistry of Si with α-Al2O3, we expected that the ester group can also be chemically bound on the surface of Si and thermally decomposed to carbon with the simultaneous formation of SiOx through the combination of oxygen with Si in the subsequent pyrolysis process. Based on this consideration, a new strategy for fabricating a double shelled nano-Si composite was proposed directly by heat treatment of the PMMA-coated nano-Si particles in a one-step procedure, the feasibility of which has been evidenced by SEM and TEM characterizations together with the thermalgravimetric analysis (TGA) results.
EXPERIMENTAL SECTION
Preparation of the Si@SiOx/C Nanocomposite. To reduce the costs, millimeter-scale silicon powders were selected as starting material to prepare the Si@SiOx/C nanocomposite in this work. Silicon powders (325 mesh, Sigma-Aldrich) were first ball-milled in a high-energy planetary ball mill (QM-3A, Nanjing, China) for 10 h, during which the particle size was reduced to nanometer level. After then, the as-prepared nano-Si particles (0.1 g) were mixed with PMMA (0.25 g) in N,N-dimethylformamide (10 mL) solvent at a weight ratio of 1:2.5 and stirred for 2 h. The mixture was dried in a vacuum oven at 100 °C for 2 h to get Si@PMMA composite, which was further calcinated under N2 flow at 900 °C for 2 h to obtain dual B
DOI: 10.1021/acsami.6b09775 ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX
Research Article
ACS Applied Materials & Interfaces
Figure 1. SEM image (a, b), TEM image (c) and (d) HRTEM image of Si@SiOx@C composite. (e) TG curves of pure PMMA polymer and Si@ PMMA precursor composite under Ar atmosphere, nano-Si powders and Si@SiOx@C composite under air atmosphere.
Figure 2. (a) FTIR spectra of pure PMMA polymer and Si@PMMA precursor composite. (b) Si 2p XPS spectra of the pristine nano-Si and Si@ SiOx@C composite. (c) Schematic illustration for the formation mechanism of dual core−shelled Si@SiOx@C composite.
Figure 1 gives the SEM and TEM images of thus-obtained Si@SiOx@C composite. As observed in Figure 1a, the Si@ SiOx@C particles are irregular with their size at 200−500 nm. The magnified SEM image in Figure 1b exhibits that each particle is an aggregate of many small crystalline with a mean size of ∼50 nm. Figure 1c shows the TEM image of the Si@ SiOx@C nanocomposite. The dual core−shell structure can be clearly observed from the TEM image of a single particle. As can be observed from the locally magnified image of the edge of a particle (HRTEM image, Figure 1d), the Si core is completely surrounded by a uniform SiOx interlayer with a thickness of about 15 nm and a carbon outer layer with varying thickness. TGA results demonstrate the existence of residual carbon and oxygen in the pyrolytic product of Si@PMMA. As shown in Figure 1e, the pure PMMA exhibits a continuous weight loss in the temperature range of 170 °C − 400 °C under Ar atmosphere without any residue. But on the Si surface, the weight loss of PMMA is significantly decreased. According to
the designed weight ratio of Si/PMMA = 1/2.5, the theoretical residue weight of Si@PMMA precursor after pyrolysis should be 28.6%, whereas in fact it exhibits an actual residue weight of 37.1%, indicating a totally different thermal degradation behavior. The content of silicon in the resulting composite can be easily calculated from the ratio between theoretical and actual residue weight of Si@PMMA precursor composite, while the carbon content can be estimated from the weight loss of the resulting composite of Si@SiOx@C relative to nano-Si powders under Air atmosphere in TGA measurements. Based on the TGA data reflected in Figure 1e, the mass contents of silicon and carbon in the resulting Si@SiOx@C composite are calculated to be 77.1 and 12.5 wt %, respectively, while the remaining mass fraction of 10.4 wt % should be corresponding to the oxygen content. The formation mechanism of the dual core−shelled Si@ SiOx@C composite was investigated by Fourier transform infrared spectroscopy (FTIR) and X-ray photoelectron specC
DOI: 10.1021/acsami.6b09775 ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX
Research Article
ACS Applied Materials & Interfaces
Figure 3. (a) Charge/discharge profiles of the Si@SiOx@C composite electrode between 0.01 and 1.5 V at 100 mA g−1 for the first cycle, and 500 mA g−1 for the subsequent cycles. (b) High-rate capability of the Si@SiOx@C composite electrode at various currents from 100 mA g−1 to 4000 mA g−1. (c) Cycling behavior of the Si@SiOx@C composite electrode. The electrode was cycled at 100 mA g−1 for the 1st cycle, 500 mA g−1 for the 2nd cycle, and 1000 mA g−1 for the subsequent cycles.
According to this mechanism, the oxygen content in the resulting composite would not be affected by the weight ratio of PMMA to nano-Si in the precursor composite only if the amount of PMMA is sufficient to form a complete coverage on the Si surface, because the surface bonded ester groups, as the sole oxygen source, can only exist in the form of monomolecular layer. To confirm this mechanism, we selected 1.5:1, 2.5:1 and 3.5:1 as the starting weight ratios of PMMA to Si to prepare the composite examples using a polymer pyrolysis method, and determined the contents of residue carbon and oxygen in the pyrolytic products by TGA analyses. The TGA curves are given in Figure S1, and the calculation results are listed in Table S1. As can be found, the three composite examples have almost the same oxygen content, further confirming the rationality of the proposed formation mechanism. The Li-storage behaviors of the Si@SiOx@C composite electrode were evaluated by galvanostatic charge and discharge of 2016 coin-type half cells. Figure 3a displays the charge− discharge profiles of the coin cells in the voltage interval of 0.01−1.5 V at a constant current density of 100 mA g−1 for the first cycle and 500 mA g−1 for the following cycles. In the first cycle, the Si@SiOx@C electrode delivered an initial charge/ discharge capacity of 1972/2560 mA h g−1 (the capacity is based on the total mass of nanocomposite), corresponding to a Coulombic efficiency of 77%. The initial capacity loss can be attributed to the reductive decomposition of electrolyte for the formation of surface SEI film and the partial irreversible insertion of Li+ in the Si lattice. As the current density was increased from 100 mA g−1 to 500 mA g−1, the reversible capacity decreased slightly to 1825 mA h g−1 at the second cycle, and kept stably at 1800 mA h g−1 after 50 cycles. Meanwhile, the Coulombic efficiency of the electrode rose up
troscopy (XPS). Figure 2a compares the FTIR spectra of pure PMMA polymer and Si@PMMA composite, in which a clear difference can be observed between these two examples. In the case of Si@PMMA composite, a new absorption band appears at 1679 cm−1, which should result from the shift of stretching vibration absorption peak of carbonyl group toward lower wavenumber, indicating that the PMMA molecules interact with surface Si atoms to form silicon carboxylates.25 These carboxylates produce a binding force to prevent the escaping of thermally degraded species from the Si surface, thus forming carbon residue and SixO in the pyrolysis process. XPS analyses further confirmed the formation of SixO. As revealed by the Si 2p spectra in Figure 2b, the pristine nano-Si exhibits a strong peak at ∼99 eV and a much weaker peak at ∼103 eV, which can be ascribed to the elemental Si and slightly oxidized surface Si, respectively. However, these two peaks disappear in the XPS spectra of Si@SiOx@C example. Instead, a strong peak, corresponding to a higher oxidation degree of Si (SiOx, x < 226), emerges at a higher binding energy of ∼104 eV. This comparison suggests that the nano-Si is partially oxidized to SiOx in the pyrolysis process, which forms a dense surface layer to construct a core−shell structure. Based on the above results, the formation mechanism of dual core−shelled Si@SiOx@C nanocomposite can be schematically illustrated as in Figure 2c. Owing to the interaction between the ester groups on PMMA polymer and the hydroxyl groups on the silicon surface, a part of PMMA molecules are immobilized on the nanosilicon surface to form chemically bonded carboxylates. In the following pyrolysis process, these carboxylates undergo a thermal decomposition with their alkyl chains converted to carbon and the residue oxygen recombining with Si to form SiOx, thus forming dual-shells on the nano-Si surface. D
DOI: 10.1021/acsami.6b09775 ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX
Research Article
ACS Applied Materials & Interfaces quickly to ≥96% at the second cycle and continuously promoted to ≥99.4% after a few cycles. The high-rate performance of the Si@SiOx@C composite electrode is shown in Figure 3b. As the current was increased stepwise from 100 to 500 mA g−1 and further to 1000 mA g−1, the capacity of the electrode decreases slightly from 1880 to 1620 and 1400 mA h g−1. Even at very high currents of 2000 and 4000 mA g−1, the composite electrode can still achieve a reversible capacity of 1030 mA h g−1 and 670 mA h g−1, respectively, demonstrating a good rate capability. Apparently, such a high rate capability of the Si@SiOx@C composite should be benefited from the outer carbon layer that provides not only significantly enhanced electrical conductivity, but also fast Li+ transport channels for the inner alloying reactions of Si cores. Besides, the Si@SiOx@C composite also exhibited excellent capacity retention and high Coulombic efficiency upon extended cycling. As shown in Figure 3c, the composite electrode can still deliver a high reversible capacity of 1030 mA h g−1 over 500 cycles at a current density of 1000 mA h g−1, corresponding to an average capacity fade rate of only 0.07% per cycle. At the same time, the Coulombic efficiency maintained steadily at ≥99.5% after 50 cycles. In contrast, the pristine nano-Si anode shows a very poor cycling stability. As can be seen from Figure S2, the nano-Si anode can also exhibit an extra-high capacity of 2850 mA h g−1 at the first cycle, but the cyclable capacity decreases rapidly to 540 mA h g−1 in the first 100 cycles and then continuously decays to
