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A low-cost and advanced SiOx–C composite with hierarchical structure as an anode material for lithium-ion batteries† Wenjun Wu, Jing Shi, Yunhui Liang, Fang Liu, Yi Peng and Huabin Yang* A cost-efficient and scalable method is designed to prepare a SiOx–C composite with superior cyclability and excellent rate performance. The glucose addition in a two-step way induces a hierarchical structure, where individual SiOx nanoparticles are wrapped by a conductive carbon layer and these agglomerated particles are further wrapped by a carbon shell functioning as an electrolyte blocking layer. Instrumental analysis indicates that the SiOx domains are comprised of SiO2 and SiO. The SiOx–C anode exhibits a high reversible specific capacity of 674.8 mA h g
1
after 100 cycles at 100 mA g
1
with a capacity retention
Received 28th February 2015, Accepted 24th April 2015
of about 83.5%. The excellent electrochemical performance is due to the hierarchical structure, the
DOI: 10.1039/c5cp01212k
process, all of which can immensely relieve the volume expansion induced by the lithiation of silicon.
well-dispersed conductive carbon network, and the Li2O and Li4SiO4 generated in the initial discharge This hierarchical SiOx–C composite has a promising prospect of practical application given its adequate
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storage capacity, good cycling stability, commercially available materials and simple equipment.
Introduction With the rapid development of portable electronic devices and electric vehicles, Li-ion batteries have become the first priority because of its high energy density and long cycle life.1,2 The development of high energy Li-ion batteries requires new anode materials with higher energy density than that of graphitebased materials, which play a central role of anode materials in commercial lithium-ion batteries. There is no doubt that Si-based materials have the greatest potential to substitute for graphite-based materials, because of its abundance and high theoretical capacity of B4200 mA h g 1.3,4 Nevertheless, two major problems delay the practical use of Si-based materials as anode electrodes: the low intrinsic electric conductivity and massive volume changes during lithium alloying and dealloying, giving rise to poor cycling performance.5–8 To solve these problems, SiOx-based materials have been proposed for alternative anode materials because of its improved cycling stability compared with Si.9–11 Lithium insertion into SiOx in the initial
Institute of New Energy Material Chemistry, Tianjin Key Laboratory of Metal and Molecule Based Material Chemistry, Collaborative Innovation Center of Chemical Science and Engineering, Nankai University, Tianjin 300071, China. E-mail: [email protected]; Fax: +86-22-23502604; Tel: +86-22-23508405 † Electronic supplementary information (ESI) available: The 29Si-NMR spectra for the SiOx–C composite, the element analysis of EDS for the SiOx–C composite, HRTEM image of the SiOx–C composite after discharging to 0.01 V, and magnified HRTEM image of (b) region 1 and (c) region 2. See DOI: 10.1039/c5cp01212k
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lithiation process generates Li2O and Li4SiO4 matrix, which could alleviate the volume expansion of Si during lithiation– delithiation. SiOx–C composites were reported to exhibit a much better cycle performance than the bare SiOx owing to the existence of a carbon-coating layer with good electronic conductivity.12 SiO1.39–C composites synthesized by using an evaporation, condensation, and calcination process exhibited a reversible capacity of about 426 mA h g 1 up to 20 cycles.13 The nano-sized SiOx–C ¨ber method composite synthesized through a modified Sto showed high specific capacity (ca. 800 mA h g 1), excellent cycling stability, good rate-capability but a low initial coulombic efficiency of 47.3%.14 To realize the commercialization of SiOx-based materials, a facile route to large-scale production with low cost is urgently needed. In this work, cheap tetraethyl orthosilicate (TEOS) and glucose were used as the silicon source and carbon source, respectively. SiOx–C composites were synthesized via a modi¨ber method, which combined the acid and alkaline fied Sto sol–gel method, high-energy ball-milling, and the high temperature pyrolysis process. A two-step way of glucose addition was applied to produce a hierarchical structure. The SiOx–C composite with unique structure demonstrates excellent cycling stability and good rate-capability when used as an anode material for LIBs. Our method provides a simple way for the synthesis of Si-based composites with excellent performance on a large scale, which is promising for practical applications.
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Experimental
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Materials synthesis All the chemicals and solvents were of analytical grade and used without further purification. The SiOx–C composite was ¨ber method with tetraethyl synthesized using a modified Sto orthosilicate (TEOS), absolute ethanol, hydrochloric acid, ammonia and glucose as starting materials. An aqueous solution (containing 2.5 g glucose, 3 mL of deionized (DI) water and 22 mL of ethanol) was added dropwise to a homogeneous mixture (8.36 mL of TEOS and 44 mL of ethanol) under magnetic stirring. Subsequently, 1.35 mL of hydrochloric acid solution (0.1 mol L 1) was mixed with the solution followed by hydrolysis of TEOS for 10 h under magnetic stirring. Afterward, 0.14 mL of ammonia solution (0.1 mol L 1) was slowly added to the solution under magnetic stirring to form a gel precursor with 24 h of aging. The gel was dried at 80 1C in air. After 2.5 g glucose had been added into the obtained dried gel, the mixture was ball milled for 5 h using a planetary mill (QM-3SP2, Nanjing University Instrument Plant, China) with a rotation speed of 400 rpm and the mass ratio of ball to powder was 10 : 1. Then the mixture was pressured into a wafer at 10 MPa before heat treatment at 950 1C for 30 min in air. The product was ground and ball milled to get a powder. Finally the SiOx–C composite was obtained. Materials characterization X-ray diffraction (XRD) measurements were carried out using a Rigaku D/max-2500 X-ray diffractometer using Cu-Ka radiation to investigate the crystal structures of the resultant materials. The particle morphology and dispersion of components in particles of the composite were characterized using a JEM-7500F scanning electron microscope (SEM) and a Tecnai 20 high-resolution transmission electron microscope (HR-TEM). Solid-state 29Si NMR spectra were obtained using a Varian unity INOVA-400 NMR spectrometer. The valence state of the Si was determined by measuring using an X-ray photoelectron spectroscopy (XPS, Axis Ultra DLD, England). The electrode materials were scraped off the Cu current collector of the cells first discharged to 0.01 V and rinsed with dimethyl carbonate (DMC) in an Ar-filled glove box before being characterized by HRTEM.
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constant current density of 100 mA g 1 between 0.01 and 3.0 V (vs. Li+/Li) using the Land CT2100 battery test system. Cyclic voltammograms (CV) were measured at a scan rate of 0.1 mV s 1 in the potential range of 0–3.0 V using a Solartron 1287 electrochemical workstation. Electrochemical impedance spectroscopy (EIS) measurements were carried out on a Zahner-Elektrik IM6e electrochemical workstation with a signal amplitude of 5 mV in the frequency from 10 2 Hz to 100 kHz. All the above measurements were conducted at room temperature.
Results and discussion The X-ray diffraction (XRD) patterns of the SiOx–C composite and SiO2 prepared under similar conditions without the addition of glucose are shown in Fig. 1. The broad bands indicate that the structure of the SiOx–C composite and pure SiO2 nanoparticles is amorphous. With regard to the SiOx–C composite, the broad peaks at about 221 and 431 are assignable to the amorphous carbon coating layer.15 Obviously, no peak corresponding to silicon is detected, indicating that silicon was not obtained by carbothermic reduction and a homogeneous amorphous SiOx phase was formed in this experiment. The 29Si-NMR technique has been used to obtain detailed information about the composition of the SiOx–C composite. The 29Si-NMR spectrum of the SiOx–C composite is provided in Fig. S1.† The spectrum gives an inconspicuous resonance signal at 59.2 ppm and an intensive resonance signal at 108.5 ppm, attributed to the SiO and the synergistic effect of SiO2 and SiO, respectively.16 To confirm the chemical characteristics of the Si in the SiOx–C composite, XPS analysis was conducted with different argon ions sputtering time. Their respective Si 2p core level spectra are illustrated in Fig. 2. Fig. 2a shows Si 2p spectra for the SiOx–C composite before ion bombardment. It can be seen that there is only one broad peak at around 103 eV corresponding to Si4+. The result demonstrates the existence of the native SiO2 layer at its surface,17 which is unavoidable especially after heat treatment in air. To avoid surface impurities of samples for the XPS analysis, spectra with 2 min and 5 min ion bombardment were measured. As seen in Fig. 2b and c, deconvolution of the Si 2p peak demonstrates the two states of Si: Si2+ (SiO, B102 eV), and Si4+
Electrochemical measurement The composite electrode was prepared by coating the slurry with 75 wt% of active material powder, 10 wt% of carbon black (Super P) and 15 wt% of polyvinylidene fluoride (PVdF) onto copper foil. Afterward the electrode was dried for 15 h at 120 1C in vacuum. The copper foil with the electrode materials was then punched into circular discs with a diameter of 13 mm for use as a working electrode. Coin-type cells (CR2032) were assembled in an Ar-filled glove box by using Celgard 2400 as a separator, metallic lithium as a counter electrode, and 1 M LiPF6 dissolved in a mixed solvent of ethylene carbonate (EC) and dimethyl carbonate (DMC) (1 : 1 in volume ratio) as an electrolyte. The cell performance was evaluated by the galvanostatic charge–discharge cycling test, which was carried out at a
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Fig. 1
XRD patterns of the SiOx–C composite and pure SiO2.
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Fig. 3 (a, b) SEM images of the SiOx–C composite. (c) Schematic of the preparation process for the SiOx–C composite.
Fig. 2 XPS spectra of the Si 2p of SiOx–C composite with different argon ions sputtering time: (a) 0 min, (b) 2 min, (c) 5 min.
Table 1 The relative content (wt%) of Si2+ and Si4+ obtained after different sputtering times and the corresponding forms of SiOx
Content (wt%)
Sputtering time (min)
Si2+
Si4+
Form of SiOx
0 2 5
0 14.9 33.3
100 85.1 66.7
SiO2 SiO1.85 SiO1.67
(SiO2, B103 eV), indicating that a reduction reaction occurred during the carbonization process. The relative content (wt%) of Si2+ and Si4+, combined with the corresponding form of SiOx, have been listed as a function of the argon ions sputtering time and the results are shown in Table 1. Obviously, 5 min ion bombardment can result in the largest amount of Si2+, because the deeper layer is less prone to be oxidized by air. Moreover, no peaks centered at around 99 eV corresponding to Si0 (metallic Si), are detected, which is in good agreement with the XRD results.18 The size, morphology, and structure of the as-prepared samples were examined by SEM and TEM. Random-shaped particles with size ranging from hundreds of nanometers to several micrometers are observed in Fig. 3a, consistent with the characteristic of ball-milling technology. The inhomogeneity of size is owing to the short ball-milling time. A particle of micron-size (secondary particle) is composed of small particles with B100 nm size, as shown in the circular damaged region of Fig. 3b, and an outermost carbon shell. The hierarchical structure can be assigned to the addition of glucose by ball-milling, resulting in secondary carbon coating of the agglomerated primary particles. Fig. 3c explicitly depicts the schematic of the preparation process for the SiOx–C composite. In regard to the TEM image, Fig. 4a shows the agglomerated SiOx–C particles with B100 nm size, which is consistent with the SEM images. In the high resolution TEM image (Fig. 4b), it is observed that the amorphous SiOx domain is covered with a uniform carbon layer. The energy dispersive X-ray spectroscopy (EDS) line scan analysis (Fig. 4c) reveals that the carbon layer about 10 nm in
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Fig. 4 (a) TEM image of the SiOx–C composite. (b) HRTEM image of the primary SiOx–C particle and the corresponding SAED pattern (inset). (c) EDS line scan result. (d) EDS spectrum of the SiOx–C composite.
thickness is present at the outer shell and SiOx in the core. HRTEM result indicates that a single SiOx–C particle has a characteristic core–shell structure. However, due to the hierarchical structure, the lower resolution TEM image (Fig. 4a) does not reveal the core–shell structure of the SiOx–C particle as a typical feature. The selected area electron diffraction (SAED) pattern exhibits a dispersed feature, which further proves that the SiOx domains are amorphous. EDS has been performed to further elucidate the chemical composition of the SiOx–C composite. The EDS analysis demonstrates that the composite is composed of C, O and Si (Fig. 4d), and it should be noted that the element Cu is originated from the Cu substrate.19 In addition, the element analysis shows that the atom ratio of O/Si in the SiOx–C composite is approximately 1.9 (Table S1†). Typical galvanostatic discharge–charge voltage profiles of the SiOx–C electrode for the 1st, 2nd, 5th cycles at 100 mA g 1 between 0.01 and 3.0 V (vs. Li+/Li) are shown in Fig. 5a. An obvious difference exists between the first discharge voltage profile and the subsequent ones. It is a well-established fact that the high irreversible capacity loss during the initial charge– discharge is inevitable, which can be identified in accordance with the dQ/dV differential curve of the first-discharge profile in Fig. 5b. The broad peak at 0.75 V can be ascribed to the
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Fig. 5 (a) The galvanostatic discharge–charge voltage profiles of the SiOx–C electrode and (b) the dQ/dV differential curve of the firstdischarge profile.
electrolyte decomposition and accompanying solid electrolyte interface (SEI) formation.20 It should be noted that a dramatic change takes place at around 0.25 V, attributed to the electrochemical reactions between the lithium ions and the nano-SiOx (SiO2 and SiO) in the composite.9,21 An intensive peak at 0.17 V can be discerned, which should be associated with the alloying process of the generated silicon.22 Verifying the structural change of the SiOx regions after the initial discharge process is necessary to identify the reaction mechanism. By means of the XRD data, the crystallized Li4SiO4 is difficult to identify and no distinguishable peaks corresponding to the crystallized Li4SiO4 appear.10 According to the literature, the crystallized Li4SiO4 formed during the first discharge process is generally confirmed by XPS,23 HRTEM,10 and NMR.16 In our work, the HRTEM imaging was carried out to verify the crystallized Li4SiO4, and the image is presented in Fig. S2.† The HRTEM imaging was carried out after discharging to 0.01 V. The result shows that the d-spacings of 0.204 nm and 0.243 nm can be assigned to the planes of (2% 20) and (2% 10) of Li4SiO4, respectively. Although Li2O is undetected in the HRTEM image as its amorphous state,21,24 the reaction mechanism still can be summarized as follows:15,25–27 SiO2 + 4Li+ + 4e - 2Li2O + Si
(1a)
2SiO2 + 4Li+ + 4e - Li4SiO4 + Si
(1b)
SiO + 2Li+ + 2e - Li2O + 2Si
(1c)
4SiO + 4Li+ + 4e - Li4SiO4 + 3Si
(1d)
Si + xLi+ + xe 2 LixSi
(2)
Reactions (1a)–(1d) only happen in the initial discharge process, which is an important reason for the high irreversible capacity loss of the first cycle. The charge–discharge curves of the second and fifth cycle show a high overlapping degree, indicating little capacity loss occur during the first few cycles except the initial cycle. The underlying reason is that nanosized SiOx particles significantly alleviate the volume changes during Li+ insertion and extraction, and meanwhile Li4SiO4 and Li2O generated in the first discharge serve as an inert matrix which can support and disperse active Si formed by reduction, and relieve volume changes during the alloying–dealloying process.28 The potential profile in the delithiation process is
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slant without a distinct plateau, most probably due to the poor crystallinity of the active material in the electrode.29 A plateau was detected at around 0.25 V during the first discharging step, it can be attributed to the large overpotential related to the metal oxide electrode during the lithiation process.10,30 In order to testify the cycling stability of the SiOx–C composite electrode, the SiOy@C composite is produced under the same conditions as those of SiOx–C except that the whole glucose is added before ball-milling and the schematic of the preparation process for SiOy@C composite is illustrated in Fig. S3.† The cyclic performance of the SiOx–C and SiOy@C composite at a current density of 100 mA g 1 and in the voltage window of 0.01–3 V is plotted in Fig. 6a. The capacity of the SiOy@C composite electrode rapidly decreases to 558.6 mA h g 1 after 40 cycles, due to the pulverization caused by volumetric expansion of the generated Si particles. By contrast, the SiOx–C composite still retains a capacity of 674.8 mA h g 1 after 100 cycles, with 83.5% charge capacity retention. Moreover, its coulombic efficiency maintains above 99% after the 18th cycle, indicating a good reversibility. In the cycling test at a high current density of 500 mA g 1 (Fig. S4†), the SiOx–C electrode still maintains a capacity of about 485 mA h g 1 after 100 cycles, still higher than the theoretical capacity of graphite (372 mA h g 1). Hence, it demonstrates that the addition of half of the total glucose during the sol–gel process ensures monodispersion of glucose or uniform carbon coating, giving rise to a significant electrochemical improvement. The remarkable improvement can be attributed to three aspects. Firstly, a uniform carbon coating layer guarantees the nano-SiOx particles homodisperse instead of agglomeration and renders the individual SiOx particle to expand and contract inside the carbon layer. With regard to the SiOy@C composite, the SiOy particles are agglomerated instead of homogeneously dispersed. The agglomerate SiOy particles are just coated by the outmost carbon layer. Secondly, the outermost carbon shell of the secondary particle
Fig. 6 (a) Cyclic performance of the SiOx–C and SiOy@C composite electrode at a current density of 100 mA g 1. (b) Rate capability of the SiOx–C composite at various rates.
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prevents the electrolyte from being in direct contact with the primary SiOx–C particle, reducing the formation of the SEI film. Indeed, the SEI film is a downside for revesible capacity or cycling stability.31,32 Thirdly, the generated Li2O and Li4SiO4 acting as a shielding layer can accommodate the volume changes upon cycling, which have been mentioned above. Of course, for the SiOy@C composite, the Si particles are involved in the same behaviors as the SiOx–C one, as shown in the chemical eqn (1a)–(1d). Additionally, the excellent electronic conductivity of amorphous carbon and good Li+ ionic conductivity of Li4SiO4 as well as Li2O play an important role in the outstanding cycling stability.33,34 Actually, the capacity loss at each cycle is unrecoverable in an actual lithium battery, and the coulombic efficiency of each cycle along with capacity retention is so crucial that excellent cycling stability can reasonably be expected in a lithium battery using a SiOx–C composite anode.35 The rate capability is a significant assessment index for practical LIBs. Fig. 6b shows the charge–discharge curves of the SiOx–C composite electrodes as the current density increases from 100 mA g 1 to 500 mA g 1. As observed, the electrode displays an excellent cycling stability with a capacity of B500 mA h g 1 even at a high current density of 500 mA g 1. When the current density returns to 100 mA g 1, the capacity returns to B680 mA h g 1. The superior rate capability may result from the high electronic conductivity of the carbon layer and Li4SiO4 combined with Li2O as a good Li+-ion conductor, allowing fast electron and Li+-ion transport through the entire electrode. Fig. 7a shows the cyclic voltammetry curves of the SiOx–C composite at a scanning rate of 0.1 mV s 1 between 0 and 3 V for the first 5 cycles. A distinct broad reduction peak at around
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0.75 V in the first cathodic scanning process is due to the formation of SEI films, which is consistent with the result of the dQ/dV differential curve of the first-discharge profile. Additionally, the cathodic peak located at around 0.41 V corresponds to the formation of a Li–Si alloy. During the charge process, the broad anodic peak at 0.48 V can be attributed to the phase transition from the Li–Si alloy to Si. Moreover, the CV curves are almost overlapped after the 3rd cycle, demonstrating an excellent reversible behavior to a certain degree. Fig. 7b depicts Nyquist plots of the SiOx–C electrode after different cycles. Each Nyquist plot presents a compressed semicircle in the high frequency region and an inclined line at low frequency, corresponding to charge transfer and lithium ion diffusion, respectively.36,37 The SiOx–C electrode exhibits the charge-transfer resistance in the range of 40 to 45 O with tiny variations during 30 cycles, manifesting the stable interface impedance.38,39 The EIS result implies that the SiOx–C electrode undergoes stable faradic reactions during cycling, which supports the enhanced cycling stability and high-rate performance of the hierarchical SiOx–C anode.
Conclusions A unique SiOx–C composite with hierarchical structure is synthesized using a low-cost method by the sol–gel process, ball-milling and heat treatment, which have been widely used in the industrial community. Instrumental analysis manifests that the valence state of the Si in the as-prepared composite material is made up of Si2+ (SiO) and Si4+ (SiO2), while the atom ratio of O/Si is approximately 1.9. The SiOx–C composite exhibits high reversible specific capacity, good cycling stability and outstanding rate capability. The excellent electrochemical properties are ensured by the improved SiOx particle dispersion ability, the hierarchical structure and the generated Li2O and Li4SiO4 in the initial discharge process, all of which make significant contribution to accommodating the volume change of generated Si upon cycling and preventing the pulverization of the electrode. Although the initial coulombic efficiency should be improved, the SiOx–C composite can be considered a promising candidate as an anode material in commercial LIBs.
Acknowledgements This work was supported by the Program of China (2011AA11A255), Natural Science Foundation of Tianjin, China (13JCZDJC32000) and the MOE Innovation Team (IRT13022).
Notes and references
Fig. 7 (a) Cyclic voltammetry curves of the SiOx–C composite from 0 to 3.0 V for the first 5 cycles. (b) Electrochemical impedance spectra of the SiOx–C composite electrode.
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A low-cost and advanced SiOx–C composite with hierarchical structure as an anode material for lithium-ion batteries† Wenjun Wu, Jing Shi, Yunhui Liang, Fang Liu, Yi Peng and Huabin Yang* A cost-efficient and scalable method is designed to prepare a SiOx–C composite with superior cyclability and excellent rate performance. The glucose addition in a two-step way induces a hierarchical structure, where individual SiOx nanoparticles are wrapped by a conductive carbon layer and these agglomerated particles are further wrapped by a carbon shell functioning as an electrolyte blocking layer. Instrumental analysis indicates that the SiOx domains are comprised of SiO2 and SiO. The SiOx–C anode exhibits a high reversible specific capacity of 674.8 mA h g
1
after 100 cycles at 100 mA g
1
with a capacity retention
Received 28th February 2015, Accepted 24th April 2015
of about 83.5%. The excellent electrochemical performance is due to the hierarchical structure, the
DOI: 10.1039/c5cp01212k
process, all of which can immensely relieve the volume expansion induced by the lithiation of silicon.
well-dispersed conductive carbon network, and the Li2O and Li4SiO4 generated in the initial discharge This hierarchical SiOx–C composite has a promising prospect of practical application given its adequate
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storage capacity, good cycling stability, commercially available materials and simple equipment.
Introduction With the rapid development of portable electronic devices and electric vehicles, Li-ion batteries have become the first priority because of its high energy density and long cycle life.1,2 The development of high energy Li-ion batteries requires new anode materials with higher energy density than that of graphitebased materials, which play a central role of anode materials in commercial lithium-ion batteries. There is no doubt that Si-based materials have the greatest potential to substitute for graphite-based materials, because of its abundance and high theoretical capacity of B4200 mA h g 1.3,4 Nevertheless, two major problems delay the practical use of Si-based materials as anode electrodes: the low intrinsic electric conductivity and massive volume changes during lithium alloying and dealloying, giving rise to poor cycling performance.5–8 To solve these problems, SiOx-based materials have been proposed for alternative anode materials because of its improved cycling stability compared with Si.9–11 Lithium insertion into SiOx in the initial
Institute of New Energy Material Chemistry, Tianjin Key Laboratory of Metal and Molecule Based Material Chemistry, Collaborative Innovation Center of Chemical Science and Engineering, Nankai University, Tianjin 300071, China. E-mail: [email protected]; Fax: +86-22-23502604; Tel: +86-22-23508405 † Electronic supplementary information (ESI) available: The 29Si-NMR spectra for the SiOx–C composite, the element analysis of EDS for the SiOx–C composite, HRTEM image of the SiOx–C composite after discharging to 0.01 V, and magnified HRTEM image of (b) region 1 and (c) region 2. See DOI: 10.1039/c5cp01212k
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lithiation process generates Li2O and Li4SiO4 matrix, which could alleviate the volume expansion of Si during lithiation– delithiation. SiOx–C composites were reported to exhibit a much better cycle performance than the bare SiOx owing to the existence of a carbon-coating layer with good electronic conductivity.12 SiO1.39–C composites synthesized by using an evaporation, condensation, and calcination process exhibited a reversible capacity of about 426 mA h g 1 up to 20 cycles.13 The nano-sized SiOx–C ¨ber method composite synthesized through a modified Sto showed high specific capacity (ca. 800 mA h g 1), excellent cycling stability, good rate-capability but a low initial coulombic efficiency of 47.3%.14 To realize the commercialization of SiOx-based materials, a facile route to large-scale production with low cost is urgently needed. In this work, cheap tetraethyl orthosilicate (TEOS) and glucose were used as the silicon source and carbon source, respectively. SiOx–C composites were synthesized via a modi¨ber method, which combined the acid and alkaline fied Sto sol–gel method, high-energy ball-milling, and the high temperature pyrolysis process. A two-step way of glucose addition was applied to produce a hierarchical structure. The SiOx–C composite with unique structure demonstrates excellent cycling stability and good rate-capability when used as an anode material for LIBs. Our method provides a simple way for the synthesis of Si-based composites with excellent performance on a large scale, which is promising for practical applications.
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Experimental
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Materials synthesis All the chemicals and solvents were of analytical grade and used without further purification. The SiOx–C composite was ¨ber method with tetraethyl synthesized using a modified Sto orthosilicate (TEOS), absolute ethanol, hydrochloric acid, ammonia and glucose as starting materials. An aqueous solution (containing 2.5 g glucose, 3 mL of deionized (DI) water and 22 mL of ethanol) was added dropwise to a homogeneous mixture (8.36 mL of TEOS and 44 mL of ethanol) under magnetic stirring. Subsequently, 1.35 mL of hydrochloric acid solution (0.1 mol L 1) was mixed with the solution followed by hydrolysis of TEOS for 10 h under magnetic stirring. Afterward, 0.14 mL of ammonia solution (0.1 mol L 1) was slowly added to the solution under magnetic stirring to form a gel precursor with 24 h of aging. The gel was dried at 80 1C in air. After 2.5 g glucose had been added into the obtained dried gel, the mixture was ball milled for 5 h using a planetary mill (QM-3SP2, Nanjing University Instrument Plant, China) with a rotation speed of 400 rpm and the mass ratio of ball to powder was 10 : 1. Then the mixture was pressured into a wafer at 10 MPa before heat treatment at 950 1C for 30 min in air. The product was ground and ball milled to get a powder. Finally the SiOx–C composite was obtained. Materials characterization X-ray diffraction (XRD) measurements were carried out using a Rigaku D/max-2500 X-ray diffractometer using Cu-Ka radiation to investigate the crystal structures of the resultant materials. The particle morphology and dispersion of components in particles of the composite were characterized using a JEM-7500F scanning electron microscope (SEM) and a Tecnai 20 high-resolution transmission electron microscope (HR-TEM). Solid-state 29Si NMR spectra were obtained using a Varian unity INOVA-400 NMR spectrometer. The valence state of the Si was determined by measuring using an X-ray photoelectron spectroscopy (XPS, Axis Ultra DLD, England). The electrode materials were scraped off the Cu current collector of the cells first discharged to 0.01 V and rinsed with dimethyl carbonate (DMC) in an Ar-filled glove box before being characterized by HRTEM.
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constant current density of 100 mA g 1 between 0.01 and 3.0 V (vs. Li+/Li) using the Land CT2100 battery test system. Cyclic voltammograms (CV) were measured at a scan rate of 0.1 mV s 1 in the potential range of 0–3.0 V using a Solartron 1287 electrochemical workstation. Electrochemical impedance spectroscopy (EIS) measurements were carried out on a Zahner-Elektrik IM6e electrochemical workstation with a signal amplitude of 5 mV in the frequency from 10 2 Hz to 100 kHz. All the above measurements were conducted at room temperature.
Results and discussion The X-ray diffraction (XRD) patterns of the SiOx–C composite and SiO2 prepared under similar conditions without the addition of glucose are shown in Fig. 1. The broad bands indicate that the structure of the SiOx–C composite and pure SiO2 nanoparticles is amorphous. With regard to the SiOx–C composite, the broad peaks at about 221 and 431 are assignable to the amorphous carbon coating layer.15 Obviously, no peak corresponding to silicon is detected, indicating that silicon was not obtained by carbothermic reduction and a homogeneous amorphous SiOx phase was formed in this experiment. The 29Si-NMR technique has been used to obtain detailed information about the composition of the SiOx–C composite. The 29Si-NMR spectrum of the SiOx–C composite is provided in Fig. S1.† The spectrum gives an inconspicuous resonance signal at 59.2 ppm and an intensive resonance signal at 108.5 ppm, attributed to the SiO and the synergistic effect of SiO2 and SiO, respectively.16 To confirm the chemical characteristics of the Si in the SiOx–C composite, XPS analysis was conducted with different argon ions sputtering time. Their respective Si 2p core level spectra are illustrated in Fig. 2. Fig. 2a shows Si 2p spectra for the SiOx–C composite before ion bombardment. It can be seen that there is only one broad peak at around 103 eV corresponding to Si4+. The result demonstrates the existence of the native SiO2 layer at its surface,17 which is unavoidable especially after heat treatment in air. To avoid surface impurities of samples for the XPS analysis, spectra with 2 min and 5 min ion bombardment were measured. As seen in Fig. 2b and c, deconvolution of the Si 2p peak demonstrates the two states of Si: Si2+ (SiO, B102 eV), and Si4+
Electrochemical measurement The composite electrode was prepared by coating the slurry with 75 wt% of active material powder, 10 wt% of carbon black (Super P) and 15 wt% of polyvinylidene fluoride (PVdF) onto copper foil. Afterward the electrode was dried for 15 h at 120 1C in vacuum. The copper foil with the electrode materials was then punched into circular discs with a diameter of 13 mm for use as a working electrode. Coin-type cells (CR2032) were assembled in an Ar-filled glove box by using Celgard 2400 as a separator, metallic lithium as a counter electrode, and 1 M LiPF6 dissolved in a mixed solvent of ethylene carbonate (EC) and dimethyl carbonate (DMC) (1 : 1 in volume ratio) as an electrolyte. The cell performance was evaluated by the galvanostatic charge–discharge cycling test, which was carried out at a
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Fig. 1
XRD patterns of the SiOx–C composite and pure SiO2.
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Fig. 3 (a, b) SEM images of the SiOx–C composite. (c) Schematic of the preparation process for the SiOx–C composite.
Fig. 2 XPS spectra of the Si 2p of SiOx–C composite with different argon ions sputtering time: (a) 0 min, (b) 2 min, (c) 5 min.
Table 1 The relative content (wt%) of Si2+ and Si4+ obtained after different sputtering times and the corresponding forms of SiOx
Content (wt%)
Sputtering time (min)
Si2+
Si4+
Form of SiOx
0 2 5
0 14.9 33.3
100 85.1 66.7
SiO2 SiO1.85 SiO1.67
(SiO2, B103 eV), indicating that a reduction reaction occurred during the carbonization process. The relative content (wt%) of Si2+ and Si4+, combined with the corresponding form of SiOx, have been listed as a function of the argon ions sputtering time and the results are shown in Table 1. Obviously, 5 min ion bombardment can result in the largest amount of Si2+, because the deeper layer is less prone to be oxidized by air. Moreover, no peaks centered at around 99 eV corresponding to Si0 (metallic Si), are detected, which is in good agreement with the XRD results.18 The size, morphology, and structure of the as-prepared samples were examined by SEM and TEM. Random-shaped particles with size ranging from hundreds of nanometers to several micrometers are observed in Fig. 3a, consistent with the characteristic of ball-milling technology. The inhomogeneity of size is owing to the short ball-milling time. A particle of micron-size (secondary particle) is composed of small particles with B100 nm size, as shown in the circular damaged region of Fig. 3b, and an outermost carbon shell. The hierarchical structure can be assigned to the addition of glucose by ball-milling, resulting in secondary carbon coating of the agglomerated primary particles. Fig. 3c explicitly depicts the schematic of the preparation process for the SiOx–C composite. In regard to the TEM image, Fig. 4a shows the agglomerated SiOx–C particles with B100 nm size, which is consistent with the SEM images. In the high resolution TEM image (Fig. 4b), it is observed that the amorphous SiOx domain is covered with a uniform carbon layer. The energy dispersive X-ray spectroscopy (EDS) line scan analysis (Fig. 4c) reveals that the carbon layer about 10 nm in
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Fig. 4 (a) TEM image of the SiOx–C composite. (b) HRTEM image of the primary SiOx–C particle and the corresponding SAED pattern (inset). (c) EDS line scan result. (d) EDS spectrum of the SiOx–C composite.
thickness is present at the outer shell and SiOx in the core. HRTEM result indicates that a single SiOx–C particle has a characteristic core–shell structure. However, due to the hierarchical structure, the lower resolution TEM image (Fig. 4a) does not reveal the core–shell structure of the SiOx–C particle as a typical feature. The selected area electron diffraction (SAED) pattern exhibits a dispersed feature, which further proves that the SiOx domains are amorphous. EDS has been performed to further elucidate the chemical composition of the SiOx–C composite. The EDS analysis demonstrates that the composite is composed of C, O and Si (Fig. 4d), and it should be noted that the element Cu is originated from the Cu substrate.19 In addition, the element analysis shows that the atom ratio of O/Si in the SiOx–C composite is approximately 1.9 (Table S1†). Typical galvanostatic discharge–charge voltage profiles of the SiOx–C electrode for the 1st, 2nd, 5th cycles at 100 mA g 1 between 0.01 and 3.0 V (vs. Li+/Li) are shown in Fig. 5a. An obvious difference exists between the first discharge voltage profile and the subsequent ones. It is a well-established fact that the high irreversible capacity loss during the initial charge– discharge is inevitable, which can be identified in accordance with the dQ/dV differential curve of the first-discharge profile in Fig. 5b. The broad peak at 0.75 V can be ascribed to the
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Fig. 5 (a) The galvanostatic discharge–charge voltage profiles of the SiOx–C electrode and (b) the dQ/dV differential curve of the firstdischarge profile.
electrolyte decomposition and accompanying solid electrolyte interface (SEI) formation.20 It should be noted that a dramatic change takes place at around 0.25 V, attributed to the electrochemical reactions between the lithium ions and the nano-SiOx (SiO2 and SiO) in the composite.9,21 An intensive peak at 0.17 V can be discerned, which should be associated with the alloying process of the generated silicon.22 Verifying the structural change of the SiOx regions after the initial discharge process is necessary to identify the reaction mechanism. By means of the XRD data, the crystallized Li4SiO4 is difficult to identify and no distinguishable peaks corresponding to the crystallized Li4SiO4 appear.10 According to the literature, the crystallized Li4SiO4 formed during the first discharge process is generally confirmed by XPS,23 HRTEM,10 and NMR.16 In our work, the HRTEM imaging was carried out to verify the crystallized Li4SiO4, and the image is presented in Fig. S2.† The HRTEM imaging was carried out after discharging to 0.01 V. The result shows that the d-spacings of 0.204 nm and 0.243 nm can be assigned to the planes of (2% 20) and (2% 10) of Li4SiO4, respectively. Although Li2O is undetected in the HRTEM image as its amorphous state,21,24 the reaction mechanism still can be summarized as follows:15,25–27 SiO2 + 4Li+ + 4e - 2Li2O + Si
(1a)
2SiO2 + 4Li+ + 4e - Li4SiO4 + Si
(1b)
SiO + 2Li+ + 2e - Li2O + 2Si
(1c)
4SiO + 4Li+ + 4e - Li4SiO4 + 3Si
(1d)
Si + xLi+ + xe 2 LixSi
(2)
Reactions (1a)–(1d) only happen in the initial discharge process, which is an important reason for the high irreversible capacity loss of the first cycle. The charge–discharge curves of the second and fifth cycle show a high overlapping degree, indicating little capacity loss occur during the first few cycles except the initial cycle. The underlying reason is that nanosized SiOx particles significantly alleviate the volume changes during Li+ insertion and extraction, and meanwhile Li4SiO4 and Li2O generated in the first discharge serve as an inert matrix which can support and disperse active Si formed by reduction, and relieve volume changes during the alloying–dealloying process.28 The potential profile in the delithiation process is
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slant without a distinct plateau, most probably due to the poor crystallinity of the active material in the electrode.29 A plateau was detected at around 0.25 V during the first discharging step, it can be attributed to the large overpotential related to the metal oxide electrode during the lithiation process.10,30 In order to testify the cycling stability of the SiOx–C composite electrode, the SiOy@C composite is produced under the same conditions as those of SiOx–C except that the whole glucose is added before ball-milling and the schematic of the preparation process for SiOy@C composite is illustrated in Fig. S3.† The cyclic performance of the SiOx–C and SiOy@C composite at a current density of 100 mA g 1 and in the voltage window of 0.01–3 V is plotted in Fig. 6a. The capacity of the SiOy@C composite electrode rapidly decreases to 558.6 mA h g 1 after 40 cycles, due to the pulverization caused by volumetric expansion of the generated Si particles. By contrast, the SiOx–C composite still retains a capacity of 674.8 mA h g 1 after 100 cycles, with 83.5% charge capacity retention. Moreover, its coulombic efficiency maintains above 99% after the 18th cycle, indicating a good reversibility. In the cycling test at a high current density of 500 mA g 1 (Fig. S4†), the SiOx–C electrode still maintains a capacity of about 485 mA h g 1 after 100 cycles, still higher than the theoretical capacity of graphite (372 mA h g 1). Hence, it demonstrates that the addition of half of the total glucose during the sol–gel process ensures monodispersion of glucose or uniform carbon coating, giving rise to a significant electrochemical improvement. The remarkable improvement can be attributed to three aspects. Firstly, a uniform carbon coating layer guarantees the nano-SiOx particles homodisperse instead of agglomeration and renders the individual SiOx particle to expand and contract inside the carbon layer. With regard to the SiOy@C composite, the SiOy particles are agglomerated instead of homogeneously dispersed. The agglomerate SiOy particles are just coated by the outmost carbon layer. Secondly, the outermost carbon shell of the secondary particle
Fig. 6 (a) Cyclic performance of the SiOx–C and SiOy@C composite electrode at a current density of 100 mA g 1. (b) Rate capability of the SiOx–C composite at various rates.
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prevents the electrolyte from being in direct contact with the primary SiOx–C particle, reducing the formation of the SEI film. Indeed, the SEI film is a downside for revesible capacity or cycling stability.31,32 Thirdly, the generated Li2O and Li4SiO4 acting as a shielding layer can accommodate the volume changes upon cycling, which have been mentioned above. Of course, for the SiOy@C composite, the Si particles are involved in the same behaviors as the SiOx–C one, as shown in the chemical eqn (1a)–(1d). Additionally, the excellent electronic conductivity of amorphous carbon and good Li+ ionic conductivity of Li4SiO4 as well as Li2O play an important role in the outstanding cycling stability.33,34 Actually, the capacity loss at each cycle is unrecoverable in an actual lithium battery, and the coulombic efficiency of each cycle along with capacity retention is so crucial that excellent cycling stability can reasonably be expected in a lithium battery using a SiOx–C composite anode.35 The rate capability is a significant assessment index for practical LIBs. Fig. 6b shows the charge–discharge curves of the SiOx–C composite electrodes as the current density increases from 100 mA g 1 to 500 mA g 1. As observed, the electrode displays an excellent cycling stability with a capacity of B500 mA h g 1 even at a high current density of 500 mA g 1. When the current density returns to 100 mA g 1, the capacity returns to B680 mA h g 1. The superior rate capability may result from the high electronic conductivity of the carbon layer and Li4SiO4 combined with Li2O as a good Li+-ion conductor, allowing fast electron and Li+-ion transport through the entire electrode. Fig. 7a shows the cyclic voltammetry curves of the SiOx–C composite at a scanning rate of 0.1 mV s 1 between 0 and 3 V for the first 5 cycles. A distinct broad reduction peak at around
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0.75 V in the first cathodic scanning process is due to the formation of SEI films, which is consistent with the result of the dQ/dV differential curve of the first-discharge profile. Additionally, the cathodic peak located at around 0.41 V corresponds to the formation of a Li–Si alloy. During the charge process, the broad anodic peak at 0.48 V can be attributed to the phase transition from the Li–Si alloy to Si. Moreover, the CV curves are almost overlapped after the 3rd cycle, demonstrating an excellent reversible behavior to a certain degree. Fig. 7b depicts Nyquist plots of the SiOx–C electrode after different cycles. Each Nyquist plot presents a compressed semicircle in the high frequency region and an inclined line at low frequency, corresponding to charge transfer and lithium ion diffusion, respectively.36,37 The SiOx–C electrode exhibits the charge-transfer resistance in the range of 40 to 45 O with tiny variations during 30 cycles, manifesting the stable interface impedance.38,39 The EIS result implies that the SiOx–C electrode undergoes stable faradic reactions during cycling, which supports the enhanced cycling stability and high-rate performance of the hierarchical SiOx–C anode.
Conclusions A unique SiOx–C composite with hierarchical structure is synthesized using a low-cost method by the sol–gel process, ball-milling and heat treatment, which have been widely used in the industrial community. Instrumental analysis manifests that the valence state of the Si in the as-prepared composite material is made up of Si2+ (SiO) and Si4+ (SiO2), while the atom ratio of O/Si is approximately 1.9. The SiOx–C composite exhibits high reversible specific capacity, good cycling stability and outstanding rate capability. The excellent electrochemical properties are ensured by the improved SiOx particle dispersion ability, the hierarchical structure and the generated Li2O and Li4SiO4 in the initial discharge process, all of which make significant contribution to accommodating the volume change of generated Si upon cycling and preventing the pulverization of the electrode. Although the initial coulombic efficiency should be improved, the SiOx–C composite can be considered a promising candidate as an anode material in commercial LIBs.
Acknowledgements This work was supported by the Program of China (2011AA11A255), Natural Science Foundation of Tianjin, China (13JCZDJC32000) and the MOE Innovation Team (IRT13022).
Notes and references
Fig. 7 (a) Cyclic voltammetry curves of the SiOx–C composite from 0 to 3.0 V for the first 5 cycles. (b) Electrochemical impedance spectra of the SiOx–C composite electrode.
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