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Insight into silicon-carbon multilayer films as anode materials for lithium-ion batteries: A combined experimental and first principles study
Acta Materialia 178 (2019) 173e178
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Insight into silicon-carbon multilayer films as anode materials for lithium-ion batteries: A combined experimental and first principles study Zhen Zhang, Ningbo Liao*, Hongming Zhou, Wei Xue College of Mechanical & Electrical Engineering, Wenzhou University, Wenzhou, 325035, PR China
a r t i c l e i n f o
a b s t r a c t
Article history: Received 13 February 2019 Received in revised form 26 July 2019 Accepted 6 August 2019 Available online 8 August 2019
The combination of silicon and carbon layers exhibits superior lithium capacity and rate performance; however, the corresponding lithiation mechanism on the atomic-scale is not clear. In this work, the impact of the carbon layer on the electrochemical performance of silicon-carbon film systems as the lithium anode is investigated by a combination of experiments and first principles calculations. Experimental results show that the sample with the thickest carbon layer presents the smallest first cycle discharge capacities (2814 mAhg1); however, this sample also results in the largest capacity retentions after 100 cycles (69%) and the rate capability test (48.4%). Based on first principles calculations, the average length of the LieSi bond near the silicon-carbon interface is significantly shorter than that in silicon, indicating an irreversible capacity loss. The structure with the largest carbon layer thickness corresponds to enhanced reversible capacity, electronic conductivity and lithium diffusion coefficient, which is consistent with experimental results. Our calculations provide a deeper understanding of irreversible capacity loss and how the primary nanostructure contributes to superior rate performance for silicon-carbon film anode materials. © 2019 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
Keywords: Lithium-ion battery Silicon-carbon Film anode materials First principles calculation Lithiation mechanism
1. Introduction Lithium-ion batteries show large potential for the storage and utilization of energy in many applications such as electric vehicles, microelectronic devices and microsensors [1]. In particular, the development of electric vehicles requires higher specific capacity than current commercial graphitic anodes with a capacity of 372 mAhg1 [2]. Another key issue for graphitic anodes is the limited rate performance attributed to slower lithium diffusion and increased resistance caused by solid electrolyte interface. Silicon is the most promising alternative to graphite as a next-generation anode material, with large theoretical capacity of 3579 mAhg1 and low cost [3]. In addition, silicon has a lower operating potential of 0.4 V, which is important to enhance the battery voltage and obtain high energy efficiency with low voltage hysteresis between lithiation and delithiation. However, there is a huge obstacle in the application of silicon due to the large volume change (>300%) during lithiation, which results in the destruction of the electrode
* Corresponding author. E-mail address: [email protected] (N. Liao). https://doi.org/10.1016/j.actamat.2019.08.009 1359-6454/© 2019 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
system and a drastic attenuation of capacity [4e6]. Therefore, developing high capacity electrode materials with excellent cycling performance is still an unsolved key issue for the further applications of lithium-ion batteries. Several approaches have been employed to overcome the volume expansion problem of silicon [7e11]. Silicon-based nanomaterials, including nanowires and nanoparticles, are applicable solutions. A combination of silicon- and carbon-based materials could be a perfect way to solve the problem. Yolk-shell structures with hollow carbon and silicon cores were prepared, and the results show different degrees of improvement in performance [12e14]. Preparing silicon-based film anode materials is another direction, where the compositions and feature size can be precisely controlled by film preparation technologies. The general idea is to reduce volume expansion by decreasing the feature size of the materials and adding carbon as buffer layers to stabilize the lithiated structures during charging/discharging cycles. Takamura demonstrated that a very thin amorphous silicon film with a thinness of 100 nm can maintain stable capacity over 100 cycles [15]. A silicon-carbon composite prepared with the optimized nanoscale building block size and carbon coating temperature presents a lithiation capacity
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of 1200 mAhg1 after stable 600 cycles, with a rate of 1.2 Ag-1 [16]. Although the available experimental studies show the excellent electrochemical performance of silicon-based film anodes, the lithiation mechanism on the atomic-scale is not clear. Further study on the origin of the desirable electrochemical properties is crucial for the future design of silicon-based anode materials. A powerful approach to investigate the properties of electrode materials is using first principles calculations [17e21], which were successfully used to evaluate the electrochemical performance of silicon oxycarbide anodes materials in our previous studies [22,23]. A combination of experimental technologies and numerical approaches will be beneficial to understand the nanostructure-performance relationship for electrode materials. In this work, the impact of the carbon layer on the electrochemical performance of siliconcarbon film systems as the lithium anode is investigated by a combination of experiments and first principles calculations. Amorphous silicon and carbon film anode samples are prepared by radio frequency magnetron sputtering and their electrochemical performance are evaluated. First principles calculations are then conducted to simulate the lithiation process of silicon-carbon interfaces with different carbon layer thicknesses. The predicted properties are consistent with the experimental results and provide a deeper understanding of irreversible capacity loss and how the primary nanostructure contributes to the rate performance for silicon-carbon film anode materials. 2. Method 2.1. Experimental methods Amorphous silicon and carbon films were prepared by radio frequency (RF) magnetron sputtering. Double targets of N-type mono-crystalline silicon (purity 99.99%) and amorphous carbon (purity 99.99%) were used, and the distances between targets and substrate were 8 cm. The silicon and carbon films were deposited at constant RF power of 100 W and 60 W. Both depositions were performed at 25 C and pre-sputtered with argon to remove surface impurities of the target. In the sputtering process, the substrate was rotated at 10 rpm and water-cooled to make the thickness of the deposited film uniform under a constant temperature. After the base pressure reached 7 104 Pa, the argon flow was adjusted to 30 sccm to ensure a working pressure of 1.5 Pa in the vacuum chamber. The deposition rate was determined by measuring the thickness of the film for a certain period of time; thus, the film thickness of the sample can be controlled by adjusting the deposition time. The amorphous silicon film was deposited on copper foil with a thickness of approximately 250 nm, and the amorphous carbon film was deposited on silicon film with three designed thicknesses of 10 nm, 50 nm and 100 nm for sample1, sample2 and sample3, respectively. The specific capacity is calculated based on the thickness and area of the films, assuming an estimated theoretical density of 2.62 gcm3 and 2.33 gcm3 for carbon and silicon, respectively. The electrochemical performance of the film anodes were tested with charging/discharging experiments by a coin cell setup (Cr2032). The cells were assembled in an argon-filled glove box with O2 and H2O purities of less than 1 ppm. The prepared films and lithium metal were used as working electrode and counter electrode, respectively, and 1 M (1 molL-1) LiPF6/EC: DEC (1:1 by volume) with 2% VC additive was used as the electrolyte. The charge and discharge tests were implemented in an electrochemical workstation (CHI660B, Shanghai Chenhua Instrument Co.) with constant and different current rates of 0.05C, 1C, 2C, 3C and 4C, with cut off voltage between 0.04 and 1.5 V. A small current test at a rate of 0.05C was carried out to accurately measure the
capacity of the cell. 2.2. Computational method The atomistic simulations were based on the density functional theory (DFT) ultrasoft pseudopotential method and conducted with a CASTEP module [24] in Materials Studio. The generalized gradient approximation (GGA) [25] of Perdew-Becke-Ernzerhof (PBE) [26] was used for the exchange correlation functional. The unit cell dimensions were 10.7 Å 10.7 Å 38 Å for all three models and the initial distance between silicon and amorphous carbon was 1.5 Å. A geometry optimization was conducted to minimize the energy at the interface and an optimal distance between the two materials was obtained. The Brillouin zone was sampled using a 3 3 1 Monkhorst-Pack k-points grid. Periodic boundary conditions were applied in all directions. The optimized structures were obtained by relaxing all of the atomic configurations until the interatomic forces were less than 0.01 eVÅ1 and the convergence of energy change per atom was less than 105 eV, with a cutoff of 400 eV for kinetic energy. The formation energy was calculated by the following [27]:
Ef ¼ ELix Si=C xELi ESi=C
(1)
where ELix Si=C and ESi=C are the total energies per Si atom of the Lix Si =C and Si =C systems, respectively, and ELi is the per atom energy of the body-centered-cubic Li (bcc-Li) system. By a fitted EfeLi concentration profile, the voltage-Li concentration curve was obtained by the following [27]:
V¼
dEf ðxÞ dx
(2)
The diffusion properties were investigated by Ab intio molecular dynamics (AIMD) simulations with canonical ensemble (NVT) and a -Hoover thermostat was used to control the time step of 1 fs A Nose temperature. The structures were annealed at 500 K to obtain sufficient atomic rearrangement, and the mean squared displacements of lithium in silicon-carbon structures were then calculated by NVT simulations at a temperature of 300 K. The CASTEP module was used to conduct all the AIMD simulations. 3. Result and discussion Fig. 1 shows the results of scanning electron microscopy (SEM), Raman spectrum and XRD pattern of the silicon-carbon films deposited on Cu foil. The SEM cross-section morphology of siliconcarbon multilayered films is shown in Fig. 1 (a), which is used to obtain structural information and determine the deposition rate in sample preparation. It can be observed that silicon and carbon are uniformly deposited on the copper foil, with obvious layering. The Raman spectra shows broad peaks at approximately 150 cm1, 300 cm1, 400 cm1 and 480 cm1, corresponding to transverse acoustic, longitudinal acoustic, longitudinal optic and transverse optic respectively, indicating that amorphous silicon is sputtered on the Cu substrate [28]. The typical D (at 1360 cm1) and G (at 1580 cm1) bands correspond to amorphous carbon [29]. The deposition rate under the preparation conditions is then determined and used to evaluate the thickness of the films. The X-ray diffraction (XRD) results of the silicon-carbon films are shown in Fig. 1 (c). All of the diffraction peaks are attributed to the Cu foil, and the peaks for silicon and carbon do not appear, indicating an amorphous state for the films system. The cycling performances for the silicon-carbon films anodes with different thickness of carbon layer are obtained at a rate of 0.05C and potential window of 0.04e1.5 V, as shown in Fig. 2. When
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Fig. 1. (a) Cross-section, (b) Raman spectrum and (c) XRD pattern of the silicon-carbon films deposited on Cu foil. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)
Fig. 2. Cycling performance of a-Si/C with different carbon thicknesses.
the thickness of the carbon layer increases, the silicon-carbon film anode shows first cycle discharge capacities of 3014 mAhg1, 2875 mAhg1 and 2814 mAhg1. As the carbon contents of the silicon-carbon films samples increases, the capacity retentions are 76%, 83% and 88% for the second cycle and are 54%, 61% and 69% after 100 cycles. The electrochemical performance of silicon-carbon film anodes with a potential window of 0.04e1.5 V and different C rates (from 0.05C to 4C) are shown in Fig. 3. As the thickness of carbon increases, the silicon-carbon system shows improved rate capability. The structure with a carbon layer of 100 nm delivers a specific capacity of 1160 m Ahg1 at the 4C rate; this is 48.4% of the capacity at the 0.05C rate and significantly larger than the structure with a carbon layer of 10 nm, which is 32.7% of the capacity at the 0.05C rate. The improved performance on rate capability is attributed to the enhanced electrical conductivity resulting from the increased carbon thickness [30]. The results indicate that the presence of the carbon layer plays a key role in the cycling performance of the film anode; however, the corresponding mechanism is not clear. To obtain deeper insights into the lithium capacity and cycling performance at the silicon-carbon interface, the lithiation processes are simulated by the step-by-step insertion of the lithium atom. An initial model of crystalline silicon is established, it shows similar results with those of amorphous silicon based on our test simulations and some previous works [31,32], as crystalline silicon will transfer to amorphous state even with very small amount of lithium. The Si (001) surface is chosen as the adsorption surface of
Fig. 3. Specific capacity for silicon-carbon films anodes with different C rates. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)
lithium, and an amorphous carbon layer obtained from the Materials Studio structure library is added on the surface. Three siliconcarbon interfacial models with different thickness of carbon are established, and the dimensions of the simulation box and thickness of the carbon layer (in Å) are shown in Fig. 4. The C/Si mass ratios for the models are 0.25, 0.36, 0.52, corresponding to similar C/Si experimental mass ratios. A set of configurations of the siliconcarbon interfaces upon different lithium concentrations are shown in the lower part of Fig. 4. It can be observed that silicon-silicon bonds gradually break when the lithium concentration increases, and the general trend of the amorphization of silicon is similar for the structures with different carbon thicknesses. Moreover, as the thickness of the carbon layer increases, a more significant limitation on the volume expansion of silicon can be observed. To obtain more detailed information on the structural properties of the lithiated silicon-carbon interface, the lithium concentration dependent pair distribution functions (PDFs) of SieLi are calculated and shown in Fig. 5. The lithium content is represented by x, the number of Li atoms per silicon atom. The first PDF peak for SieLi decreases with as the lithium content increases, relating to the formatting of the SieC bond at the silicon and carbon interface. In addition, a thicker carbon layer generally corresponds to a smaller nearest neighbor distance of SieLi, resulting in a limitation on the volume of silicon and is consistent with the above observations. Thus, the carbon layer shows a major contribution to the improved cycling performance of the silicon-carbon system, and various
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Fig. 4. Dimensions of the simulation box (in Å), the thickness of carbon layer (in Å), and the lithiated configurations of the silicon-carbon interfacial models at different stages of lithiations. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)
Fig. 5. The PDFs of LieSi at different Li concentrations.
thicknesses of carbon would result in different capabilities of volume limitation. Fig. 6 shows some typical zoomed-in atomic configurations in fully lithiated silicon-carbon interfacial structures (x ¼ 4). It can be observed that the SieC bonds form in the highly lithiated siliconcarbon interface, consistent with the above PDF results. The average length for the LieSi bonds near the silicon-carbon interface
is approximately 2.5 Å, which is significantly shorter than the 2.7 Å for those in the silicon portion (Fig. 6 (b)) and lithiated pure silicon [18]. The shorter length of the LieSi bond at the silicon-carbon interface corresponds to an irreversible capacity loss due to the strong LieSi connections [33e36]. Moreover, LieC bonds with lengths of approximately 1.8 Å form near the silicon-carbon interface, which are also shorter than the 2.2 Å found in LieC in lithiated carbon [37], indicating an irreversible capacity loss. As the thickness of carbon increases, the formation of shorter LieSi bonds related to irreversible capacity loss is reduced, which provides an explanation for the improved cycling performance of the sample with a thicker carbon layer. The formation energies (Ef) at different lithiation stages are calculated and used to evaluate the theoretical capacities of the silicon-carbon systems. Moreover, the reversible capacities are also predicted by excluding the amount of lithium atoms with shorter LieSi bond lengths at the silicon-carbon interface and are compared to experimental results with similar C/Si mass ratios, as shown in Fig. 7. The formation energy generally decreases monotonically as the lithium concentration increases. The structure with a thicker carbon layer shows a steeper descending trend for the formation energy, indicating it is more energetically favorable for lithium insertion. As the structures reach saturation for lithium (x ¼ 4), the theoretical capacities correspond to 3067 mAhg1 (C/ Si ¼ 0.25), 2810 mAhg1 (C/Si ¼ 0.36) and 2513 mAhg1 (C/ Si ¼ 0.52), which are close to the experimental values of 2875 mAhg1 (C/Si ¼ 0.22) and 2814 mAhg1 (C/Si ¼ 0.45). The reversible lithium capacity is then calculated by excluding the lithium atoms with shorter LieSi bonds. The obtained reversible capacities are 2315 mAhg1 (C/Si ¼ 0.25), 2125 mAhg1 (C/ Si ¼ 0.36) and 2185 mAhg1 mAhg1 (C/Si ¼ 0.52), which are comparable to the experimental values of 2385 (C/Si ¼ 0.22) and 2509 mAhg1 (C/Si ¼ 0.45). The theoretical capacity generally
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Fig. 6. Some typical zoomed-in atomic configurations in the fully lithiated silicon-carbon interfacial structures. (a) Formation of SieC the bond and lengths of SieLi bonds (in Å) at the silicon-carbon interface, (b) bonds lengths for the SieLi bond in the silicon portion, (c) the lithiated interfacial structures for different models. The black, yellow and violet spheres represent the C, Si and Li atoms respectively. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)
Fig. 7. Calculated formation energies, theoretical capacities, and reversible capacities of silicon-carbon systems, compared with our experimental results. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)
decreases as the carbon thickness increases, as silicon is the primary source for theoretical capacity. However, the structure with an increasing C/Si mass ratio shows increased reversible capacities, consistent with the above conclusion on irreversible capacity loss. The electronic conductivity and diffusion property are believed to be main factors that influence the rate performance of electrode materials [36,38,39]. The rate capabilities of the silicon-carbon films are represented by capacity retention after different C rate tests, as shown in Fig. 3. The electronic conductivity of the lithiated silicon-carbon interfacial structures are analyzed by calculating their density of states (DOS) and compared with the experimental results, as shown in Fig. 8 (a). The calculated band gaps of the silicon-carbon interfaces are 0.301, 0.107 and 0.021 eV as carbon thickness increases. As a smaller band gap means an easier transfer of electrons, the calculation results demonstrate that the thicker carbon layer will lead to an improved electronic conductivity and rate capability, consistent with our experimental conclusions. The mean squared displacements (MSDs) of lithium in the three siliconcarbon structures are obtained from molecular dynamics simulations and are used to calculate the diffusion coefficient, as show in
Fig. 8. Calculated density of states, band gaps, mean squared displacements, and diffusion coefficients of silicon-carbon systems with different carbon thickness are compared to the experimental results of capacity retentions. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)
Fig. 8 (b). The diffusion coefficient for model 3 is 10.59 104 cm2s1, which is greater than the diffusion coefficients of 5.24 104 cm2s1 and 7.50 104 cm2s1 for model 1 and model
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2, respectively. A greater diffusion coefficient means an easier diffusion of lithium in the structure with a thicker carbon layer, which is also in accordance with the results of rate capability in our experiments. 4. Conclusion In this work, experiments and first principles calculations are combined to investigate the impact of the carbon layer on the electrochemical performance of silicon-carbon film anodes. The experimental results show that the sample with the thickest carbon layer presents the smallest first cycle discharge capacities (2814 mAhg1); however, this sample also corresponds to the largest capacity retentions after 100 cycles (69%) and rate capability test (48.4%). First principles calculations reveal that the average length of the LieSi bonds near the silicon-carbon interface is significantly shorter than that in silicon, relating to irreversible capacity loss. The calculated theoretical capacities are 3067 mAhg1 (C/Si ¼ 0.25), 2810 mAhg1 (C/Si ¼ 0.36) and 2513 mAhg1 (C/Si ¼ 0.52), which are close to the capacities of 2875 mAhg1 (C/Si ¼ 0.22) and 2814 mAhg1 (C/Si ¼ 0.45) found in our experiment. The structure with the largest carbon layer thickness corresponds to enhanced reversible capacity, electronic conductivity and lithium diffusion coefficient, consistent with the experimental results. The proposed approach can be applied to evaluate and predict reversible capacity and rate capability of silicon-based multilayer interfaces and other film anode materials for lithium-ion batteries. Acknowledgements The authors would like to acknowledge the support of the National Natural Science Foundation of China (51675384). Appendix A. Supplementary data Supplementary data to this article can be found online at https://doi.org/10.1016/j.actamat.2019.08.009. References [1] N. Nitta, F. Wu, J.T. Lee, G. Yushin, Li-ion battery materials: present and future, Mater. Today 18 (5) (2015) 252e264. [2] R.B. Cervera, N. Suzuki, T. Ohnishi, M. Osada, K. Mitsuishi, T. Kambara, K. Takada, High performance silicon-based anodes in solid-sate lithium batteries, Energy Environ. Sci. 7 (2) (2014) 662e666. [3] M.N. Obrovac, V.L. Chevrier, Alloy negative electrodes for Li-ion batteries, Chem. Rev. 114 (23) (2014) 11444e11502. [4] H. Wu, G. Chan, J.W. Choi, I. Ryu, Y. Yao, M.T. Mcdowell, S.W. Lee, A. Jackson, Y. Yang, L. Hu, Stable cycling of double-walled silicon nanotube battery anodes through solid-electrolyte interphase control, Nat. Nanotechnol. 7 (5) (2012) 310e315. [5] W.J. Zhang, A review of the electrochemical performance of alloy anodes for lithium-ion batteries, J. Power Sources 196 (1) (2011) 13e24. [6] S. Huang, F. Fan, J. Li, S. Zhang, T. Zhu, Stress generation during lithiation of high-capacity electrode particles in lithium ion batteries, Acta Mater. 61 (12) (2013) 4354e4364. [7] M. Fang, Z. Wang, X. Chen, S. Guan, Sponge-like reduced graphene oxide/silicon/carbon nanotube composites for lithium ion batteries, Appl. Surf. Sci. 436 (2018). [8] Z. Liu, P. Guo, B. Liu, W. Xie, D. Liu, D. He, Carbon-coated Si nanoparticles/ reduced graphene oxide multilayer anchored to nanostructured current collector as lithium-ion battery anode, Appl. Surf. Sci. 396 (2017) 41e47. [9] W. He, H. Tian, S. Zhang, H. Ying, Z. Meng, W. Han, Scalable synthesis of Si/C anode enhanced by FeSix nanoparticles from low-cost ferrosilicon for lithiumion batteries, J. Power Sources 353 (2017) 270e276. [10] D. Kowalski, J. Mallet, S. Thomas, A.W. Nemaga, J. Michel, C. Guery, M. Molinari, M. Morcrette, Electrochemical synthesis of 1D core-shell Si/TiO2 nanotubes for lithium ion batteries, J. Power Sources 361 (2017) 243e248. [11] A.F. Bower, E. Chason, P.R. Guduru, B.W. Sheldon, A continuum model of deformation, transport and irreversible changes in atomic structure in amorphous lithiumesilicon electrodes, Acta Mater. 98 (2015) 229e241.
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Insight into silicon-carbon multilayer films as anode materials for lithium-ion batteries: A combined experimental and first principles study Zhen Zhang, Ningbo Liao*, Hongming Zhou, Wei Xue College of Mechanical & Electrical Engineering, Wenzhou University, Wenzhou, 325035, PR China
a r t i c l e i n f o
a b s t r a c t
Article history: Received 13 February 2019 Received in revised form 26 July 2019 Accepted 6 August 2019 Available online 8 August 2019
The combination of silicon and carbon layers exhibits superior lithium capacity and rate performance; however, the corresponding lithiation mechanism on the atomic-scale is not clear. In this work, the impact of the carbon layer on the electrochemical performance of silicon-carbon film systems as the lithium anode is investigated by a combination of experiments and first principles calculations. Experimental results show that the sample with the thickest carbon layer presents the smallest first cycle discharge capacities (2814 mAhg1); however, this sample also results in the largest capacity retentions after 100 cycles (69%) and the rate capability test (48.4%). Based on first principles calculations, the average length of the LieSi bond near the silicon-carbon interface is significantly shorter than that in silicon, indicating an irreversible capacity loss. The structure with the largest carbon layer thickness corresponds to enhanced reversible capacity, electronic conductivity and lithium diffusion coefficient, which is consistent with experimental results. Our calculations provide a deeper understanding of irreversible capacity loss and how the primary nanostructure contributes to superior rate performance for silicon-carbon film anode materials. © 2019 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
Keywords: Lithium-ion battery Silicon-carbon Film anode materials First principles calculation Lithiation mechanism
1. Introduction Lithium-ion batteries show large potential for the storage and utilization of energy in many applications such as electric vehicles, microelectronic devices and microsensors [1]. In particular, the development of electric vehicles requires higher specific capacity than current commercial graphitic anodes with a capacity of 372 mAhg1 [2]. Another key issue for graphitic anodes is the limited rate performance attributed to slower lithium diffusion and increased resistance caused by solid electrolyte interface. Silicon is the most promising alternative to graphite as a next-generation anode material, with large theoretical capacity of 3579 mAhg1 and low cost [3]. In addition, silicon has a lower operating potential of 0.4 V, which is important to enhance the battery voltage and obtain high energy efficiency with low voltage hysteresis between lithiation and delithiation. However, there is a huge obstacle in the application of silicon due to the large volume change (>300%) during lithiation, which results in the destruction of the electrode
* Corresponding author. E-mail address: [email protected] (N. Liao). https://doi.org/10.1016/j.actamat.2019.08.009 1359-6454/© 2019 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
system and a drastic attenuation of capacity [4e6]. Therefore, developing high capacity electrode materials with excellent cycling performance is still an unsolved key issue for the further applications of lithium-ion batteries. Several approaches have been employed to overcome the volume expansion problem of silicon [7e11]. Silicon-based nanomaterials, including nanowires and nanoparticles, are applicable solutions. A combination of silicon- and carbon-based materials could be a perfect way to solve the problem. Yolk-shell structures with hollow carbon and silicon cores were prepared, and the results show different degrees of improvement in performance [12e14]. Preparing silicon-based film anode materials is another direction, where the compositions and feature size can be precisely controlled by film preparation technologies. The general idea is to reduce volume expansion by decreasing the feature size of the materials and adding carbon as buffer layers to stabilize the lithiated structures during charging/discharging cycles. Takamura demonstrated that a very thin amorphous silicon film with a thinness of 100 nm can maintain stable capacity over 100 cycles [15]. A silicon-carbon composite prepared with the optimized nanoscale building block size and carbon coating temperature presents a lithiation capacity
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of 1200 mAhg1 after stable 600 cycles, with a rate of 1.2 Ag-1 [16]. Although the available experimental studies show the excellent electrochemical performance of silicon-based film anodes, the lithiation mechanism on the atomic-scale is not clear. Further study on the origin of the desirable electrochemical properties is crucial for the future design of silicon-based anode materials. A powerful approach to investigate the properties of electrode materials is using first principles calculations [17e21], which were successfully used to evaluate the electrochemical performance of silicon oxycarbide anodes materials in our previous studies [22,23]. A combination of experimental technologies and numerical approaches will be beneficial to understand the nanostructure-performance relationship for electrode materials. In this work, the impact of the carbon layer on the electrochemical performance of siliconcarbon film systems as the lithium anode is investigated by a combination of experiments and first principles calculations. Amorphous silicon and carbon film anode samples are prepared by radio frequency magnetron sputtering and their electrochemical performance are evaluated. First principles calculations are then conducted to simulate the lithiation process of silicon-carbon interfaces with different carbon layer thicknesses. The predicted properties are consistent with the experimental results and provide a deeper understanding of irreversible capacity loss and how the primary nanostructure contributes to the rate performance for silicon-carbon film anode materials. 2. Method 2.1. Experimental methods Amorphous silicon and carbon films were prepared by radio frequency (RF) magnetron sputtering. Double targets of N-type mono-crystalline silicon (purity 99.99%) and amorphous carbon (purity 99.99%) were used, and the distances between targets and substrate were 8 cm. The silicon and carbon films were deposited at constant RF power of 100 W and 60 W. Both depositions were performed at 25 C and pre-sputtered with argon to remove surface impurities of the target. In the sputtering process, the substrate was rotated at 10 rpm and water-cooled to make the thickness of the deposited film uniform under a constant temperature. After the base pressure reached 7 104 Pa, the argon flow was adjusted to 30 sccm to ensure a working pressure of 1.5 Pa in the vacuum chamber. The deposition rate was determined by measuring the thickness of the film for a certain period of time; thus, the film thickness of the sample can be controlled by adjusting the deposition time. The amorphous silicon film was deposited on copper foil with a thickness of approximately 250 nm, and the amorphous carbon film was deposited on silicon film with three designed thicknesses of 10 nm, 50 nm and 100 nm for sample1, sample2 and sample3, respectively. The specific capacity is calculated based on the thickness and area of the films, assuming an estimated theoretical density of 2.62 gcm3 and 2.33 gcm3 for carbon and silicon, respectively. The electrochemical performance of the film anodes were tested with charging/discharging experiments by a coin cell setup (Cr2032). The cells were assembled in an argon-filled glove box with O2 and H2O purities of less than 1 ppm. The prepared films and lithium metal were used as working electrode and counter electrode, respectively, and 1 M (1 molL-1) LiPF6/EC: DEC (1:1 by volume) with 2% VC additive was used as the electrolyte. The charge and discharge tests were implemented in an electrochemical workstation (CHI660B, Shanghai Chenhua Instrument Co.) with constant and different current rates of 0.05C, 1C, 2C, 3C and 4C, with cut off voltage between 0.04 and 1.5 V. A small current test at a rate of 0.05C was carried out to accurately measure the
capacity of the cell. 2.2. Computational method The atomistic simulations were based on the density functional theory (DFT) ultrasoft pseudopotential method and conducted with a CASTEP module [24] in Materials Studio. The generalized gradient approximation (GGA) [25] of Perdew-Becke-Ernzerhof (PBE) [26] was used for the exchange correlation functional. The unit cell dimensions were 10.7 Å 10.7 Å 38 Å for all three models and the initial distance between silicon and amorphous carbon was 1.5 Å. A geometry optimization was conducted to minimize the energy at the interface and an optimal distance between the two materials was obtained. The Brillouin zone was sampled using a 3 3 1 Monkhorst-Pack k-points grid. Periodic boundary conditions were applied in all directions. The optimized structures were obtained by relaxing all of the atomic configurations until the interatomic forces were less than 0.01 eVÅ1 and the convergence of energy change per atom was less than 105 eV, with a cutoff of 400 eV for kinetic energy. The formation energy was calculated by the following [27]:
Ef ¼ ELix Si=C xELi ESi=C
(1)
where ELix Si=C and ESi=C are the total energies per Si atom of the Lix Si =C and Si =C systems, respectively, and ELi is the per atom energy of the body-centered-cubic Li (bcc-Li) system. By a fitted EfeLi concentration profile, the voltage-Li concentration curve was obtained by the following [27]:
V¼
dEf ðxÞ dx
(2)
The diffusion properties were investigated by Ab intio molecular dynamics (AIMD) simulations with canonical ensemble (NVT) and a -Hoover thermostat was used to control the time step of 1 fs A Nose temperature. The structures were annealed at 500 K to obtain sufficient atomic rearrangement, and the mean squared displacements of lithium in silicon-carbon structures were then calculated by NVT simulations at a temperature of 300 K. The CASTEP module was used to conduct all the AIMD simulations. 3. Result and discussion Fig. 1 shows the results of scanning electron microscopy (SEM), Raman spectrum and XRD pattern of the silicon-carbon films deposited on Cu foil. The SEM cross-section morphology of siliconcarbon multilayered films is shown in Fig. 1 (a), which is used to obtain structural information and determine the deposition rate in sample preparation. It can be observed that silicon and carbon are uniformly deposited on the copper foil, with obvious layering. The Raman spectra shows broad peaks at approximately 150 cm1, 300 cm1, 400 cm1 and 480 cm1, corresponding to transverse acoustic, longitudinal acoustic, longitudinal optic and transverse optic respectively, indicating that amorphous silicon is sputtered on the Cu substrate [28]. The typical D (at 1360 cm1) and G (at 1580 cm1) bands correspond to amorphous carbon [29]. The deposition rate under the preparation conditions is then determined and used to evaluate the thickness of the films. The X-ray diffraction (XRD) results of the silicon-carbon films are shown in Fig. 1 (c). All of the diffraction peaks are attributed to the Cu foil, and the peaks for silicon and carbon do not appear, indicating an amorphous state for the films system. The cycling performances for the silicon-carbon films anodes with different thickness of carbon layer are obtained at a rate of 0.05C and potential window of 0.04e1.5 V, as shown in Fig. 2. When
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Fig. 1. (a) Cross-section, (b) Raman spectrum and (c) XRD pattern of the silicon-carbon films deposited on Cu foil. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)
Fig. 2. Cycling performance of a-Si/C with different carbon thicknesses.
the thickness of the carbon layer increases, the silicon-carbon film anode shows first cycle discharge capacities of 3014 mAhg1, 2875 mAhg1 and 2814 mAhg1. As the carbon contents of the silicon-carbon films samples increases, the capacity retentions are 76%, 83% and 88% for the second cycle and are 54%, 61% and 69% after 100 cycles. The electrochemical performance of silicon-carbon film anodes with a potential window of 0.04e1.5 V and different C rates (from 0.05C to 4C) are shown in Fig. 3. As the thickness of carbon increases, the silicon-carbon system shows improved rate capability. The structure with a carbon layer of 100 nm delivers a specific capacity of 1160 m Ahg1 at the 4C rate; this is 48.4% of the capacity at the 0.05C rate and significantly larger than the structure with a carbon layer of 10 nm, which is 32.7% of the capacity at the 0.05C rate. The improved performance on rate capability is attributed to the enhanced electrical conductivity resulting from the increased carbon thickness [30]. The results indicate that the presence of the carbon layer plays a key role in the cycling performance of the film anode; however, the corresponding mechanism is not clear. To obtain deeper insights into the lithium capacity and cycling performance at the silicon-carbon interface, the lithiation processes are simulated by the step-by-step insertion of the lithium atom. An initial model of crystalline silicon is established, it shows similar results with those of amorphous silicon based on our test simulations and some previous works [31,32], as crystalline silicon will transfer to amorphous state even with very small amount of lithium. The Si (001) surface is chosen as the adsorption surface of
Fig. 3. Specific capacity for silicon-carbon films anodes with different C rates. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)
lithium, and an amorphous carbon layer obtained from the Materials Studio structure library is added on the surface. Three siliconcarbon interfacial models with different thickness of carbon are established, and the dimensions of the simulation box and thickness of the carbon layer (in Å) are shown in Fig. 4. The C/Si mass ratios for the models are 0.25, 0.36, 0.52, corresponding to similar C/Si experimental mass ratios. A set of configurations of the siliconcarbon interfaces upon different lithium concentrations are shown in the lower part of Fig. 4. It can be observed that silicon-silicon bonds gradually break when the lithium concentration increases, and the general trend of the amorphization of silicon is similar for the structures with different carbon thicknesses. Moreover, as the thickness of the carbon layer increases, a more significant limitation on the volume expansion of silicon can be observed. To obtain more detailed information on the structural properties of the lithiated silicon-carbon interface, the lithium concentration dependent pair distribution functions (PDFs) of SieLi are calculated and shown in Fig. 5. The lithium content is represented by x, the number of Li atoms per silicon atom. The first PDF peak for SieLi decreases with as the lithium content increases, relating to the formatting of the SieC bond at the silicon and carbon interface. In addition, a thicker carbon layer generally corresponds to a smaller nearest neighbor distance of SieLi, resulting in a limitation on the volume of silicon and is consistent with the above observations. Thus, the carbon layer shows a major contribution to the improved cycling performance of the silicon-carbon system, and various
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Fig. 4. Dimensions of the simulation box (in Å), the thickness of carbon layer (in Å), and the lithiated configurations of the silicon-carbon interfacial models at different stages of lithiations. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)
Fig. 5. The PDFs of LieSi at different Li concentrations.
thicknesses of carbon would result in different capabilities of volume limitation. Fig. 6 shows some typical zoomed-in atomic configurations in fully lithiated silicon-carbon interfacial structures (x ¼ 4). It can be observed that the SieC bonds form in the highly lithiated siliconcarbon interface, consistent with the above PDF results. The average length for the LieSi bonds near the silicon-carbon interface
is approximately 2.5 Å, which is significantly shorter than the 2.7 Å for those in the silicon portion (Fig. 6 (b)) and lithiated pure silicon [18]. The shorter length of the LieSi bond at the silicon-carbon interface corresponds to an irreversible capacity loss due to the strong LieSi connections [33e36]. Moreover, LieC bonds with lengths of approximately 1.8 Å form near the silicon-carbon interface, which are also shorter than the 2.2 Å found in LieC in lithiated carbon [37], indicating an irreversible capacity loss. As the thickness of carbon increases, the formation of shorter LieSi bonds related to irreversible capacity loss is reduced, which provides an explanation for the improved cycling performance of the sample with a thicker carbon layer. The formation energies (Ef) at different lithiation stages are calculated and used to evaluate the theoretical capacities of the silicon-carbon systems. Moreover, the reversible capacities are also predicted by excluding the amount of lithium atoms with shorter LieSi bond lengths at the silicon-carbon interface and are compared to experimental results with similar C/Si mass ratios, as shown in Fig. 7. The formation energy generally decreases monotonically as the lithium concentration increases. The structure with a thicker carbon layer shows a steeper descending trend for the formation energy, indicating it is more energetically favorable for lithium insertion. As the structures reach saturation for lithium (x ¼ 4), the theoretical capacities correspond to 3067 mAhg1 (C/ Si ¼ 0.25), 2810 mAhg1 (C/Si ¼ 0.36) and 2513 mAhg1 (C/ Si ¼ 0.52), which are close to the experimental values of 2875 mAhg1 (C/Si ¼ 0.22) and 2814 mAhg1 (C/Si ¼ 0.45). The reversible lithium capacity is then calculated by excluding the lithium atoms with shorter LieSi bonds. The obtained reversible capacities are 2315 mAhg1 (C/Si ¼ 0.25), 2125 mAhg1 (C/ Si ¼ 0.36) and 2185 mAhg1 mAhg1 (C/Si ¼ 0.52), which are comparable to the experimental values of 2385 (C/Si ¼ 0.22) and 2509 mAhg1 (C/Si ¼ 0.45). The theoretical capacity generally
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Fig. 6. Some typical zoomed-in atomic configurations in the fully lithiated silicon-carbon interfacial structures. (a) Formation of SieC the bond and lengths of SieLi bonds (in Å) at the silicon-carbon interface, (b) bonds lengths for the SieLi bond in the silicon portion, (c) the lithiated interfacial structures for different models. The black, yellow and violet spheres represent the C, Si and Li atoms respectively. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)
Fig. 7. Calculated formation energies, theoretical capacities, and reversible capacities of silicon-carbon systems, compared with our experimental results. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)
decreases as the carbon thickness increases, as silicon is the primary source for theoretical capacity. However, the structure with an increasing C/Si mass ratio shows increased reversible capacities, consistent with the above conclusion on irreversible capacity loss. The electronic conductivity and diffusion property are believed to be main factors that influence the rate performance of electrode materials [36,38,39]. The rate capabilities of the silicon-carbon films are represented by capacity retention after different C rate tests, as shown in Fig. 3. The electronic conductivity of the lithiated silicon-carbon interfacial structures are analyzed by calculating their density of states (DOS) and compared with the experimental results, as shown in Fig. 8 (a). The calculated band gaps of the silicon-carbon interfaces are 0.301, 0.107 and 0.021 eV as carbon thickness increases. As a smaller band gap means an easier transfer of electrons, the calculation results demonstrate that the thicker carbon layer will lead to an improved electronic conductivity and rate capability, consistent with our experimental conclusions. The mean squared displacements (MSDs) of lithium in the three siliconcarbon structures are obtained from molecular dynamics simulations and are used to calculate the diffusion coefficient, as show in
Fig. 8. Calculated density of states, band gaps, mean squared displacements, and diffusion coefficients of silicon-carbon systems with different carbon thickness are compared to the experimental results of capacity retentions. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)
Fig. 8 (b). The diffusion coefficient for model 3 is 10.59 104 cm2s1, which is greater than the diffusion coefficients of 5.24 104 cm2s1 and 7.50 104 cm2s1 for model 1 and model
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2, respectively. A greater diffusion coefficient means an easier diffusion of lithium in the structure with a thicker carbon layer, which is also in accordance with the results of rate capability in our experiments. 4. Conclusion In this work, experiments and first principles calculations are combined to investigate the impact of the carbon layer on the electrochemical performance of silicon-carbon film anodes. The experimental results show that the sample with the thickest carbon layer presents the smallest first cycle discharge capacities (2814 mAhg1); however, this sample also corresponds to the largest capacity retentions after 100 cycles (69%) and rate capability test (48.4%). First principles calculations reveal that the average length of the LieSi bonds near the silicon-carbon interface is significantly shorter than that in silicon, relating to irreversible capacity loss. The calculated theoretical capacities are 3067 mAhg1 (C/Si ¼ 0.25), 2810 mAhg1 (C/Si ¼ 0.36) and 2513 mAhg1 (C/Si ¼ 0.52), which are close to the capacities of 2875 mAhg1 (C/Si ¼ 0.22) and 2814 mAhg1 (C/Si ¼ 0.45) found in our experiment. The structure with the largest carbon layer thickness corresponds to enhanced reversible capacity, electronic conductivity and lithium diffusion coefficient, consistent with the experimental results. The proposed approach can be applied to evaluate and predict reversible capacity and rate capability of silicon-based multilayer interfaces and other film anode materials for lithium-ion batteries. Acknowledgements The authors would like to acknowledge the support of the National Natural Science Foundation of China (51675384). Appendix A. Supplementary data Supplementary data to this article can be found online at https://doi.org/10.1016/j.actamat.2019.08.009. References [1] N. Nitta, F. Wu, J.T. Lee, G. Yushin, Li-ion battery materials: present and future, Mater. Today 18 (5) (2015) 252e264. [2] R.B. Cervera, N. Suzuki, T. Ohnishi, M. Osada, K. Mitsuishi, T. Kambara, K. Takada, High performance silicon-based anodes in solid-sate lithium batteries, Energy Environ. Sci. 7 (2) (2014) 662e666. [3] M.N. Obrovac, V.L. Chevrier, Alloy negative electrodes for Li-ion batteries, Chem. Rev. 114 (23) (2014) 11444e11502. [4] H. Wu, G. Chan, J.W. Choi, I. Ryu, Y. Yao, M.T. Mcdowell, S.W. Lee, A. Jackson, Y. Yang, L. Hu, Stable cycling of double-walled silicon nanotube battery anodes through solid-electrolyte interphase control, Nat. Nanotechnol. 7 (5) (2012) 310e315. [5] W.J. Zhang, A review of the electrochemical performance of alloy anodes for lithium-ion batteries, J. Power Sources 196 (1) (2011) 13e24. [6] S. Huang, F. Fan, J. Li, S. Zhang, T. Zhu, Stress generation during lithiation of high-capacity electrode particles in lithium ion batteries, Acta Mater. 61 (12) (2013) 4354e4364. [7] M. Fang, Z. Wang, X. Chen, S. Guan, Sponge-like reduced graphene oxide/silicon/carbon nanotube composites for lithium ion batteries, Appl. Surf. Sci. 436 (2018). [8] Z. Liu, P. Guo, B. Liu, W. Xie, D. Liu, D. He, Carbon-coated Si nanoparticles/ reduced graphene oxide multilayer anchored to nanostructured current collector as lithium-ion battery anode, Appl. Surf. Sci. 396 (2017) 41e47. [9] W. He, H. Tian, S. Zhang, H. Ying, Z. Meng, W. Han, Scalable synthesis of Si/C anode enhanced by FeSix nanoparticles from low-cost ferrosilicon for lithiumion batteries, J. Power Sources 353 (2017) 270e276. [10] D. Kowalski, J. Mallet, S. Thomas, A.W. Nemaga, J. Michel, C. Guery, M. Molinari, M. Morcrette, Electrochemical synthesis of 1D core-shell Si/TiO2 nanotubes for lithium ion batteries, J. Power Sources 361 (2017) 243e248. [11] A.F. Bower, E. Chason, P.R. Guduru, B.W. Sheldon, A continuum model of deformation, transport and irreversible changes in atomic structure in amorphous lithiumesilicon electrodes, Acta Mater. 98 (2015) 229e241.
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