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Ab initio molecular dynamics study of SiO2 lithiation
Journal Pre-proofs Research paper Ab initio Molecular Dynamics Study of SiO2 Lithiation Iwnetim Iwnetu Abate, Chunjing J. Jia, Brian Moritz, Thomas P. Devereaux PII: DOI: Reference:
S0009-2614(19)30914-5 https://doi.org/10.1016/j.cplett.2019.136933 CPLETT 136933
To appear in:
Chemical Physics Letters
Received Date: Revised Date: Accepted Date:
2 October 2019 31 October 2019 2 November 2019
Please cite this article as: I. Iwnetu Abate, C.J. Jia, B. Moritz, T.P. Devereaux, Ab initio Molecular Dynamics Study of SiO2 Lithiation, Chemical Physics Letters (2019), doi: https://doi.org/10.1016/j.cplett.2019.136933
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© 2019 Published by Elsevier B.V.
Ab initio Molecular Dynamics Study of SiO2 Lithiation Iwnetim Iwnetu Abate,1,2 Chunjing J. Jia,2 Brian Moritz2, Thomas P. Devereaux1,2* 1 Department
of Materials Science and Engineering, Stanford University, Stanford, CA 94305, United States
2Stanford
Institute for Materials and Energy Sciences, Stanford University, Stanford, CA 94305,
United States and SLAC National Accelerator Laboratory, Menlo Park, CA 94025, United States * Corresponding author Abstract Li-ion batteries (LIBs) are sought to meet the demand for high energy storage applications. Due to its high specific charge capacity and low discharge potential, SiO2 is a promising candidate material for LIBs anodes. To design high performance anodes and coating materials using SiO2, understanding the structural transformation and Li-ion kinetics in different structural forms of SiO2 is essential. Here, we performed ab initio molecular dynamics to study the lithiation mechanism for crystalline and amorphous SiO2 and the effect of surface termination to elucidate the lithiation process of conversion oxides and contribute to the development of future LIBs.
Introduction The effort to make better batteries has been growing steadily due to their use in electric vehicles and grid storage. One of the most promising alternative anode materials, silicon, offers ten times more theoretical capacity than graphite; however, the 400% volume change during charge/discharge and poor cyclability have hindered widespread manufacturing and commercialization of this anode material1-7. In parallel, research efforts have been devoted to the development and commercialization of anode materials based on oxides of tin and silicon which have less change in volume.8–13 Silicon dioxide, SiO2 or silica, has played an important role in materials history – from glasses to everyday electronic devices, silica is an attractive candidate material due to its natural abundance. Silica also naturally forms as a native oxide on silicon that is used as anodes, affecting the lithiation process at the anode (solid)-electrolyte interface (SEI).14-16 We have recently demonstrated how native oxide on silicon anode contributes to the SEI composition and property by combining first-principle calculations and experimental techniques such as x-ray reflectivity and x-ray photoelectron spectroscopy.14 In addition, researchers have used SiO2 as a coating layer, via atomic layer deposition (ALD), on Si anodes to accommodate the large volume expansion and enhance the cyclability of Si electrodes.17-20 Artificial ALD coatings provide a resistance against fracture of Si-based electrodes during discharge/charge processes compared to natively formed oxides. Silicon oxides also may serve as a solid electrolyte in batteries. Transitioning from liquid to solid electrolytes represents an essential safety measure
to prevent explosions in case of shorting. To this end, Nava et al.21 first demonstrated the use of silica as a solid electrolyte. Quantum chemical calculations for SiO2 to study diffusion and lithiation mechanisms usually involve the lithiation of a single Li atom17,22 or adding one atom per calculation.23 The dynamics of lithiation in SiO2 may be complicated further by a crystal-to-amorphous phase transformation, which itself depends on SiO2 structure types. Thus, the detailed lithiation process must be investigated to gain quantitative insight into the fundamental interactions between Li, Si, and O and the diffusion kinetics. As demonstrated by Johari et al.24, this level of detail can be achieved by studying details of the mixing process during lithiation of SiO2 and the crystalline-toamorphous phase transition. Understanding the ion transport and lithiation mechanisms in different SiO2 forms, e.g. crystalline and amorphous, is useful for tailoring an SiO2 form for specific application and wider adoption. In this work, we performed ab initio molecular dynamics to study the lithiation mechanism in crystalline and amorphous, c-SiO2 and a-SiO2. In addition, since SiO2 is terminated by either H or OH groups in the presence of electrolyte, we also investigated lithiation process in hydrogen termination c-SiO2 and a-SiO2; hereafter referred to as c-SiO2-H, and a-SiO2-H, respectively. We examined inter-mixing kinetics and track the motion of Li in crystalline and amorphous silica. In particular, we have analyzed the evolution of the structure to elucidate how Li breaks the Si-O network during the Li-SiO2 mixing process; and we have studied the effect of the termination, if any, on the lithiation mechanism. To this end, we studied the radial distribution functions (RDFs) at various stages of lithiation and analyze the evolution of Si-O, Si-
Li, and Li-O bonds to identify the formation of different phases; and we also computed from first principles and compare the diffusivity of Li at room temperature in c-SiO2, c-SiO2-H, a-SiO2 and aSiO2-H. Our results enhance our understanding of lithiation mechanism and ionic conduction in silica, in particular, and motivate further study in conversion oxides, in general. Theoretical method To simulate the Li mixing process in SiO2, we performed ab initio molecular dynamics (AIMD) calculations at finite temperature within the framework of density functional theory (DFT), as implemented in the Vienna Ab Initio Simulation Package (VASP).25,26 Projectoraugmented-wave (PAW) potentials27 are used to mimic the ionic cores, while the generalized gradient approximation (GGA) in the Perdew Burke Ernzerhof28 (PBE) flavor provides the exchange and correlation functional. All calculations were performed with a plane wave cutoff of 500 eV and a k-point mesh of 2x2x1. The optimized coordinates and lattice parameters of c-SiO2 were obtained from the Materials Project and similar information for a-SiO2 was obtained by quenching a cubic crystalline silica phase using nonequilibrium molecular dynamics simulations.29 We considered a 2x2x1 supercell of c-SiO2 and a-SiO2, both containing 72 atoms, to keep the same atomic density for all cases with periodic boundary conditions. The hydrogen terminated structures (c-SiO2-H and a-SiO2-H) were obtained by terminating all dangling bonds on the SiO2 surface with hydrogen. In addition, the SiO2-Li interface lies perpendicular to the z-direction, along which we use only a single k-point. Similar to experimental studies of lithiation in materials, AIMD calculations were performed at high temperatures, in particular 900 K, 1050 K, 1200 K, 1350 K and 1500 K. Substantial mixing of
Li with the Si and O atoms in SiO2 can be observed at these high temperatures after ~ 6000 AIMD time steps, where each step is 1.5 fs, Figure 1 and SI 1. We determined diffusivities for different temperatures by monitoring mean square displacements (MSD) for the Li and Si atoms as a function of time, and extrapolated the results to predict room temperature diffusivity. To sustain both the temperature and pressure of the system via Langevin dynamics, friction coefficients of 10 ps-1 for atoms and 5 ps-1 (fictitious mass of 500 amu) for lattice degrees of freedom were used. The temperature oscillations were controlled every 40 time steps. Virtual NanoLab software was used to calculate radial distribution function and coordination number. Virtual NanoLab software was used to calculate radial distribution function and coordination number. Virtual NanoLab software was used to calculate coordination number (CN) and radial distribution function (RDF). CN and RDF were obtained by integrating 30 steps to represent the structure at given time of the evolution (t1, t2 or t3).
Results and Discussion Figure 1 shows the structural evolution of the c-SiO2 and a-SiO2 during lithiation as function of time at 1200 K. The application of periodic boundary conditions creates two Li/SiO2 interfaces in the unit cell. The negative heat of formation leads to mixing of Li and SiO2 and the insertion of Li transforms the c-SiO2 to an amorphous phase. The driving force to penetrate into SiO2 is larger for Li atoms near the surface than Li atoms away from the interface. Si-O bonds break as Li atoms migrate into and disrupt the SiO2 structure creating different local structures14: LixSiO2 + LixSi during lithiation via the formation of Li-Si and Li-O bonds, as shown in Figure 1.
0 ps
0.075 ps
1.5 ps
6 ps
9 ps
Figure 1 Structures at various stages of lithiation of crystalline (top) and amorphous (bottom) SiO2 at 1200 K. Structures at t = 0 ps and t ∼ 9.0 ps correspond to the starting and final configurations of our calculations, respectively, while structures at t = 0.075, 1.5 and 6ps depict the various intermediate stages of lithiation. In the figure, spheres in red, yellow and magenta color represent O, Si and Li atoms, respectively. Li and SiO2 mix due to the negative heat of formation disrupting the SiO2 structure creating different local structures.
First, we consider the radial distribution function (RDF or g(r)) as a function of time to understand how the structure evolves during the alloying process. Figure 2 shows g(r) for Si-O, Si-Li, and Li-O pairs at 1200 K. From left to right, the panels depict g(r) for c-SiO2, c-SiO2-H, a-SiO2 and a-SiO2H, respectively. The structures do not evolve beyond t3 (the RDFs after t3 overlap). During lithiation and for all initial SiO2 structures, the number of Si-O bonds decreases while the number of Si-Li and Li-O neighbors increases over time, indicating evolution of Li-SiO2 mixing. Figure 2(d) shows a broadening and decrease of the second nearest neighbor peak in c-SiO2, indicating amorphization during lithiation and the conversion of c-SiO2 to a-SiO2. In addition, the broad
peak of lithiated c-SiO2 after t1 resembles that of lithiated a-SiO2, further validating formation of the amorphous structure after t1. Similar to gSi-O (r), we find that gSi-Li (r) and gLi-O (r) for cSiO2 after mixing match well with those for a-SiO2. We observe a similar structural evolution in the case of c-SiO2-H and a-SiO2-H, as shown in Figures 2(e)-(g) and (k)-(m), respectively, confirming that surface termination has little impact on the structural transformation, as one might expect. Following lithiation, some Si-O bonds remain (see Figure 1); however, the Si-O bond length slightly increases after lithiation, from ~ 1.61 Å in c-SiO2 to ~ 1.65 Å in c-SiO2-H (see Figures 2(a) and (e)), due to charge transfer from Li to Si and O, which introduces a negative charge and a repulsive force. Using site-projected density of states, Ben et al. have previously demonstrated the bonds of Si atoms atoms are saturated by lone-electron pairs from the Li atoms.22 The significant overlap between the states associated with Li, Si and O atoms, both below and above the band gap, confirms the charge transfer from Li to Si and O.
c-SiO2
c-SiO2-H
a-SiO2
a-SiO2-H
a
e b
h \ g
k
b
f
i
l
c
g
j
m
d
Figure 2 Radial distribution function as function of lithiation time at 1200 K. Lithiation of four different initial structures were studied at t1=0 ps, t2 3 ps and t3= 9 ps. (a)-(d) for c-SiO2, (e)-(g) for c-SiO2-H, (h)-(j) a-SiO2-H for a-SiO2 and (k)-(m) for a-Si-O-H. In all four structures, Si-O bond breaks while Si-Li and Li-O bond forms as the lithiation time increases.
We also monitored the evolution of atomic coordination for different Li concentrations. Figure 3 shows the average values of Si-O coordination ⟨Si−O⟩ and Li-O coordination ⟨Li−O⟩ as a function of Li composition for c-SiO2. The coordination number includes Si−O atomic bonds whose distance lies within 15% of the covalent length in the bulk crystal SiO2, which corresponds to bond distances less than 1.87 Å. A similar definition for Li-O bonds, taken from the Li-O bond length in the lithium silicates Li2Si2O5 and Li4SiO4,30 sets the upper cutoff in that case at 2.15 Å. The Si-O coordination steadily decreases and changes almost linearly, which indicates a constant rate of Si−O bond breaking. In addition, the Li-O coordination remains approximately 2 for a wide range of Li concentrations. This could indicate the presence of local characteristics of Li2O upon lithium insertion into SiO2. The decrease in Si-O coordination number and formation of Li2O clusters due to changes of Si-O bond and local reduction of an O atom by two Li atoms, respectively, has been reported by Ban et al.12 In addition, Zhang et al.13 reported a change in coordination number for a-SiO2 similar to our finding for c-SiO2. In summary, structural evolution during lithiation appears to be similar for each SiO2 structural type, i.e. the Si-O bond breaks during Li-Si and Li-O formation. Our result compliments the conclusions reached by Johari et al.24 on the structural transformation of crystalline and amorphous silicon during lithiation.
Figure 3. The average values of Si−O coordination and Li−O coordination as function of Li content in crystalline SiO2.
Finally, we also studied the diffusivity for Li ions at room temperature in c-SiO2, c-SiO2-H, a-SiO2 and a-SiO2-H. To determine the diffusivity, we calculated the average MSD of Li atoms as a function of AIMD time step for different temperatures as shown in Figure S1 of the Supplementary Information. The simulation temperatures were chosen to lie below the melting point of SiO2, 1983 K. Figure S2 demonstrates the steps we followed to extrapolate room temperature (300 K) diffusion coefficients using the Einstein relation MSD = 6Dt and the standard expression for diffusivity, D = Doexp(-Ea/kBT), where Ea is the activation energy, kB is the Boltzmann constant, and T is the temperature. Figure 4 shows the diffusion coefficient for Li ions at room temperature as well as the extrapolation for c-SiO2, c-SiO2-H, a-SiO2 and a-SiO2-H. We found the diffusivity of Li, DLi, in c-SiO2 and a-SiO2 at room temperature to be 2.368 x 10-9 and
0.086 x 10-9 m2/s with an energy barrier of 0.046 eV and 0.146 eV, respectively. These are higher diffusion constant values for lithium transport compared to other oxides, (anatase TiO2: 4.7x1016 m2/s, amorphous Al2O3: (2.7 × 10–14) m2/s, ZrO2: 1.7x10-22 m2/s , MgO: 1.7x10-30 m2/s).31-33 The diffusion in c-SiO2 is anisotropic and that the value reported here is along the fast diffusion path <001>.17 Our result of a lower diffusion barrier for c-SiO2 than a-SiO2 agrees with a previous report by van Duin et al.17 using ReaxFF Reactive Force Field Modeling. A similar trend in the diffusion barrier and coefficient were observed in hydrogen terminated SiO2: 2.748 x 10-9 m2/s and 0.040 eV for c-SiO2-H and 0.421 x 10-9 m2/s and 0.099 eV for a-SiO2-H.
a
b
c
d
Figure 4. Derived diffusivity of Li at T=300 K in Li2SiO2. (a) c-SiO2, (b) c-SiO2-H and their respective insets (right). (c)a-SiO2 and (d) a-SiO2-H. The diffusivity was obtained by fitting the mean square displacement using D = Do exp(-Ea /kbT) . c-SiO2 has lower diffusion than a-SiO2 with or without termination.
Conclusion In summary, using ab initio molecular dynamics simulations at finite temperature, we investigated the structural evolution and diffusion kinetics of crystalline and amorphous SiO2 with and without termination during lithiation. Our results reveal several key findings. First, the mechanism and structural evolution are similar in each structural type. The Si-O bond breaks, while the Li-O bond forms with a consistent coordination number of 2, similar to Li2O. Second, both c-SiO2 and c-SiO2-H are amorphized after lithiation. The third finding indicates that while amorphous oxides generally offer faster ion transport, Li diffusion along the <001> direction in cSiO2 remains faster than that in a-SiO2 by an order of magnitude irrespective of surface termination. These suggests that termination has negligible impact on both on the rate of lithium transport through SiO2 and the mechanism for structural evolution during lithiation. Last but not least, the amorphization of c-SiO2 and c-SiO2-H suggests that the kinetics for Li transport are similar to amorphous structures after the first cycle of lithiation. In general, this work enhances our understanding of lithiation mechanism and ionic conduction in silica. In addition, the results lend insights and motivation to further study conversion oxides in quest to find high performance anodes, solid electrolytes, and coatings in Li-ion batteries.
Acknowledgment I.I.A., C.J., B.M., and T.P.D. acknowledge support from the U.S. Department of Energy (DOE), Office of Basic Energy Sciences, Division of Materials Sciences and Engineering, under Contract No. DE-AC02-76SF00515, Division of Materials Sciences and Engineering at Stanford and SLAC for theoretical analysis. The computational work used resources at the National Energy Research Scientific Computing Center (NERSC), a U.S. Department of Energy Office of Science User Facility operated under Contract No. DE-AC02-05CH11231. We thank C. D. Pemmaraju, M.F. Toney, C. Cao and H.-S. Steinrück for helpful discussions. I.I.A. also would like to thank Stanford EDGE fellowship program. Author contributions I.I.A., C. J. J., B. M. and T. P. D. conceived the idea and designed the project. I.I.A., performed first principles calculations and discussed the results with C.J.J, B.M., T.P.D. All authors discussed the results, co-wrote and contributed to the manuscript. Declaration of interests The authors declare no competing interests. References [1] W.J Zhang, Journal of Power Sources 196 (2011) 13-24. [2] M. Ko, S. Chae, J. Cho, ChemElectroChem 11 (2015) 1645-1651. [3] X. Zuo, J. Zhu, P. Müller-Buschbaum, Y.J. Cheng, Nano Energy 31 (2017) 113-143.
[4] A. Casimir, H. Zhang, O. Ogoke, J.C. Amine, J. Lu, G. Wu, Nano Energy 27 (2016) 359-376. [5] S.C. Jung, J.W. Choi, Y.K. Han, Nano letters 12 (2012) 5342-5347. [6] J. Zhu, T. Wang, F. Fan, L. Mei, B. Lu, ACS nano 10 (2016) 8243-8251. [7] J. Yang, Y.-X. Wang, S.-L. Chou, R. Zhang, Y. Xu, J. Fan, W.-X. Zhang, H. K. Liu, D. Zhao, S. X. Dou, Nano Energy 18 (2015) 133-142. [8] Y. Yu, J. Zhang, L. Xue, T. Huang, A. Yu, Journal of Power Sources 196 (2011) 10240-10243. [9] J. Yang, Y. Takeda, N. Imanishi, C. Capiglia, J. Y. Xie, and O. Yamamoto, Solid State Ionics 152 (2002) 125-129. [10] D. Su, S. Dou, G. Wang, Chemistry of Materials 27 (2015) 6022-6029. [11] F. Belliard, P. A. Connor, J. T. S. Irvine, Solid State Ionics 135 (2000) 163-167. [12] Y. Idota, T. Kubota, A. Matsufuji, Y. Maekawa, T. Miyasaka, Science 276 (1997) 1395-1397. [13] Y.-N. Nuli, S.-L. Zhao, Q. -Z. Qin, Journal of Power Sources 114 (2003) 113-120. [14] C. Cao, I.I. Abate, E. Sivonxay, B. Shyam, C. Jia, B. Mortiz, T.P. Devereaux, K.A. Persson, H.G. Steinrück, M.F. Toney, Joule 3 (2019) 762-781. [15] M. T. Morita, O. E. Hasegawa, M. Kawakami, M. Ohwada, Journal of Applied Physics 68 (1990) 12721281. [16] Y. He, D.M. Piper, M. Gu, J.J. Travis, S.M. George, S.H. Lee, A. Genc, L. Pullan., J. Liu, S.X. Mao, J.G. Zhang, C. Ban, C Wang, Acs Nano 8 (2014) 11816-11823. [17] O. Alireza, S.Y. Kim, E.D. Cubuk, Y. Qi, A.C.T. van Duin, The Journal of Physical Chemistry A 120 (2016) 2114-2127. [18] J.S. Chul, and Y.K. Han, The Journal of Physical Chemistry Letters 4 (2013) 2681-2685. [19] H. Yu, X. Yu, Y. Wang, H. Li, X. Huang, Advanced Materials 23 (2011) 4938-4941. [20] L. Juchuan, X. Xiao, Y.T. Cheng, M. W. Verbrugge, The Journal of Physical Chemistry Letters 4 (2013) 3387-3391.
[21] A. Nava, G.Ceder, D.R. Sadoway, E.A. Fitzgerald, Journal of applied physics 98 (2005) 023516. [22] B. Chunmei, B.B. Kappes, Q. Xu, C. Engtrakul, C.V. Ciobanu, A.C. Dillon, Y. Zhao, Applied Physics Letters 100 (2012) 243905. [23] Z. Yuefei, Y. Li, Z. Wang, K. Zhao, Nano letters 14 (2014) 7161-7170. [24] J. Priya, Y. Qi, V.B. Shenoy, Nano letters 11 (2011) 5494-5500. [25] G. Kresse, J. Furthmüller, Physical review B 54 (1996) 11169. [26] G. Kresse, J. Furthmüller, Computational materials science 6 (1996) 15-50. [27] G. Kresse, D. Joubert, Physical Review B 59 (1999) 1758. [28] J.P. Perdew, K.Burke,M. Ernzerhof, Physical review letters 77 (1996) 3865. [29] S. Plimpton, Journal of computational physics 117 (1995) 1-19. [30] C.Y. Chou, G.S. Hwang, Chemistry of Materials 25 (2013) 3435-3440. [31] M. Wagemaker, R.v.d. Krol, A.P.M. Kentgens, A. A.V. Well, F.M. Mulder, Journal of the American Chemical Society 123 (2001) 11454-11461. [32] H. Shiqiang, C.Wolverton, The Journal of Physical Chemistry C 117 (2013) 8009-8013. [33] S. Xu, R. M. Jacobs, H. M. Nguyen, S. Hao, M. Mahanthappa, C. Wolverton, D. Morgan, Journal of Materials Chemistry A 3 (2015) 17248-17272.
Graphical abstract
Highlights
SiO2 is a promising candidate material for the anode of Li-ion batteries Ab initio molecular dynamics was used to study the lithiation mechanism in crystalline and amorphous SiO2 Effect of surface termination on lithiation was investigated Li diffusion along the <001> direction in c-SiO2 is fastest compared to other oxides Our results lend insights and motivation to further study conversion oxides in quest to find high performance anodes, solid electrolytes, and coatings in Li-ion batteries.
Declaration of interests The authors declare no competing interests.
S0009-2614(19)30914-5 https://doi.org/10.1016/j.cplett.2019.136933 CPLETT 136933
To appear in:
Chemical Physics Letters
Received Date: Revised Date: Accepted Date:
2 October 2019 31 October 2019 2 November 2019
Please cite this article as: I. Iwnetu Abate, C.J. Jia, B. Moritz, T.P. Devereaux, Ab initio Molecular Dynamics Study of SiO2 Lithiation, Chemical Physics Letters (2019), doi: https://doi.org/10.1016/j.cplett.2019.136933
This is a PDF file of an article that has undergone enhancements after acceptance, such as the addition of a cover page and metadata, and formatting for readability, but it is not yet the definitive version of record. This version will undergo additional copyediting, typesetting and review before it is published in its final form, but we are providing this version to give early visibility of the article. Please note that, during the production process, errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
© 2019 Published by Elsevier B.V.
Ab initio Molecular Dynamics Study of SiO2 Lithiation Iwnetim Iwnetu Abate,1,2 Chunjing J. Jia,2 Brian Moritz2, Thomas P. Devereaux1,2* 1 Department
of Materials Science and Engineering, Stanford University, Stanford, CA 94305, United States
2Stanford
Institute for Materials and Energy Sciences, Stanford University, Stanford, CA 94305,
United States and SLAC National Accelerator Laboratory, Menlo Park, CA 94025, United States * Corresponding author Abstract Li-ion batteries (LIBs) are sought to meet the demand for high energy storage applications. Due to its high specific charge capacity and low discharge potential, SiO2 is a promising candidate material for LIBs anodes. To design high performance anodes and coating materials using SiO2, understanding the structural transformation and Li-ion kinetics in different structural forms of SiO2 is essential. Here, we performed ab initio molecular dynamics to study the lithiation mechanism for crystalline and amorphous SiO2 and the effect of surface termination to elucidate the lithiation process of conversion oxides and contribute to the development of future LIBs.
Introduction The effort to make better batteries has been growing steadily due to their use in electric vehicles and grid storage. One of the most promising alternative anode materials, silicon, offers ten times more theoretical capacity than graphite; however, the 400% volume change during charge/discharge and poor cyclability have hindered widespread manufacturing and commercialization of this anode material1-7. In parallel, research efforts have been devoted to the development and commercialization of anode materials based on oxides of tin and silicon which have less change in volume.8–13 Silicon dioxide, SiO2 or silica, has played an important role in materials history – from glasses to everyday electronic devices, silica is an attractive candidate material due to its natural abundance. Silica also naturally forms as a native oxide on silicon that is used as anodes, affecting the lithiation process at the anode (solid)-electrolyte interface (SEI).14-16 We have recently demonstrated how native oxide on silicon anode contributes to the SEI composition and property by combining first-principle calculations and experimental techniques such as x-ray reflectivity and x-ray photoelectron spectroscopy.14 In addition, researchers have used SiO2 as a coating layer, via atomic layer deposition (ALD), on Si anodes to accommodate the large volume expansion and enhance the cyclability of Si electrodes.17-20 Artificial ALD coatings provide a resistance against fracture of Si-based electrodes during discharge/charge processes compared to natively formed oxides. Silicon oxides also may serve as a solid electrolyte in batteries. Transitioning from liquid to solid electrolytes represents an essential safety measure
to prevent explosions in case of shorting. To this end, Nava et al.21 first demonstrated the use of silica as a solid electrolyte. Quantum chemical calculations for SiO2 to study diffusion and lithiation mechanisms usually involve the lithiation of a single Li atom17,22 or adding one atom per calculation.23 The dynamics of lithiation in SiO2 may be complicated further by a crystal-to-amorphous phase transformation, which itself depends on SiO2 structure types. Thus, the detailed lithiation process must be investigated to gain quantitative insight into the fundamental interactions between Li, Si, and O and the diffusion kinetics. As demonstrated by Johari et al.24, this level of detail can be achieved by studying details of the mixing process during lithiation of SiO2 and the crystalline-toamorphous phase transition. Understanding the ion transport and lithiation mechanisms in different SiO2 forms, e.g. crystalline and amorphous, is useful for tailoring an SiO2 form for specific application and wider adoption. In this work, we performed ab initio molecular dynamics to study the lithiation mechanism in crystalline and amorphous, c-SiO2 and a-SiO2. In addition, since SiO2 is terminated by either H or OH groups in the presence of electrolyte, we also investigated lithiation process in hydrogen termination c-SiO2 and a-SiO2; hereafter referred to as c-SiO2-H, and a-SiO2-H, respectively. We examined inter-mixing kinetics and track the motion of Li in crystalline and amorphous silica. In particular, we have analyzed the evolution of the structure to elucidate how Li breaks the Si-O network during the Li-SiO2 mixing process; and we have studied the effect of the termination, if any, on the lithiation mechanism. To this end, we studied the radial distribution functions (RDFs) at various stages of lithiation and analyze the evolution of Si-O, Si-
Li, and Li-O bonds to identify the formation of different phases; and we also computed from first principles and compare the diffusivity of Li at room temperature in c-SiO2, c-SiO2-H, a-SiO2 and aSiO2-H. Our results enhance our understanding of lithiation mechanism and ionic conduction in silica, in particular, and motivate further study in conversion oxides, in general. Theoretical method To simulate the Li mixing process in SiO2, we performed ab initio molecular dynamics (AIMD) calculations at finite temperature within the framework of density functional theory (DFT), as implemented in the Vienna Ab Initio Simulation Package (VASP).25,26 Projectoraugmented-wave (PAW) potentials27 are used to mimic the ionic cores, while the generalized gradient approximation (GGA) in the Perdew Burke Ernzerhof28 (PBE) flavor provides the exchange and correlation functional. All calculations were performed with a plane wave cutoff of 500 eV and a k-point mesh of 2x2x1. The optimized coordinates and lattice parameters of c-SiO2 were obtained from the Materials Project and similar information for a-SiO2 was obtained by quenching a cubic crystalline silica phase using nonequilibrium molecular dynamics simulations.29 We considered a 2x2x1 supercell of c-SiO2 and a-SiO2, both containing 72 atoms, to keep the same atomic density for all cases with periodic boundary conditions. The hydrogen terminated structures (c-SiO2-H and a-SiO2-H) were obtained by terminating all dangling bonds on the SiO2 surface with hydrogen. In addition, the SiO2-Li interface lies perpendicular to the z-direction, along which we use only a single k-point. Similar to experimental studies of lithiation in materials, AIMD calculations were performed at high temperatures, in particular 900 K, 1050 K, 1200 K, 1350 K and 1500 K. Substantial mixing of
Li with the Si and O atoms in SiO2 can be observed at these high temperatures after ~ 6000 AIMD time steps, where each step is 1.5 fs, Figure 1 and SI 1. We determined diffusivities for different temperatures by monitoring mean square displacements (MSD) for the Li and Si atoms as a function of time, and extrapolated the results to predict room temperature diffusivity. To sustain both the temperature and pressure of the system via Langevin dynamics, friction coefficients of 10 ps-1 for atoms and 5 ps-1 (fictitious mass of 500 amu) for lattice degrees of freedom were used. The temperature oscillations were controlled every 40 time steps. Virtual NanoLab software was used to calculate radial distribution function and coordination number. Virtual NanoLab software was used to calculate radial distribution function and coordination number. Virtual NanoLab software was used to calculate coordination number (CN) and radial distribution function (RDF). CN and RDF were obtained by integrating 30 steps to represent the structure at given time of the evolution (t1, t2 or t3).
Results and Discussion Figure 1 shows the structural evolution of the c-SiO2 and a-SiO2 during lithiation as function of time at 1200 K. The application of periodic boundary conditions creates two Li/SiO2 interfaces in the unit cell. The negative heat of formation leads to mixing of Li and SiO2 and the insertion of Li transforms the c-SiO2 to an amorphous phase. The driving force to penetrate into SiO2 is larger for Li atoms near the surface than Li atoms away from the interface. Si-O bonds break as Li atoms migrate into and disrupt the SiO2 structure creating different local structures14: LixSiO2 + LixSi during lithiation via the formation of Li-Si and Li-O bonds, as shown in Figure 1.
0 ps
0.075 ps
1.5 ps
6 ps
9 ps
Figure 1 Structures at various stages of lithiation of crystalline (top) and amorphous (bottom) SiO2 at 1200 K. Structures at t = 0 ps and t ∼ 9.0 ps correspond to the starting and final configurations of our calculations, respectively, while structures at t = 0.075, 1.5 and 6ps depict the various intermediate stages of lithiation. In the figure, spheres in red, yellow and magenta color represent O, Si and Li atoms, respectively. Li and SiO2 mix due to the negative heat of formation disrupting the SiO2 structure creating different local structures.
First, we consider the radial distribution function (RDF or g(r)) as a function of time to understand how the structure evolves during the alloying process. Figure 2 shows g(r) for Si-O, Si-Li, and Li-O pairs at 1200 K. From left to right, the panels depict g(r) for c-SiO2, c-SiO2-H, a-SiO2 and a-SiO2H, respectively. The structures do not evolve beyond t3 (the RDFs after t3 overlap). During lithiation and for all initial SiO2 structures, the number of Si-O bonds decreases while the number of Si-Li and Li-O neighbors increases over time, indicating evolution of Li-SiO2 mixing. Figure 2(d) shows a broadening and decrease of the second nearest neighbor peak in c-SiO2, indicating amorphization during lithiation and the conversion of c-SiO2 to a-SiO2. In addition, the broad
peak of lithiated c-SiO2 after t1 resembles that of lithiated a-SiO2, further validating formation of the amorphous structure after t1. Similar to gSi-O (r), we find that gSi-Li (r) and gLi-O (r) for cSiO2 after mixing match well with those for a-SiO2. We observe a similar structural evolution in the case of c-SiO2-H and a-SiO2-H, as shown in Figures 2(e)-(g) and (k)-(m), respectively, confirming that surface termination has little impact on the structural transformation, as one might expect. Following lithiation, some Si-O bonds remain (see Figure 1); however, the Si-O bond length slightly increases after lithiation, from ~ 1.61 Å in c-SiO2 to ~ 1.65 Å in c-SiO2-H (see Figures 2(a) and (e)), due to charge transfer from Li to Si and O, which introduces a negative charge and a repulsive force. Using site-projected density of states, Ben et al. have previously demonstrated the bonds of Si atoms atoms are saturated by lone-electron pairs from the Li atoms.22 The significant overlap between the states associated with Li, Si and O atoms, both below and above the band gap, confirms the charge transfer from Li to Si and O.
c-SiO2
c-SiO2-H
a-SiO2
a-SiO2-H
a
e b
h \ g
k
b
f
i
l
c
g
j
m
d
Figure 2 Radial distribution function as function of lithiation time at 1200 K. Lithiation of four different initial structures were studied at t1=0 ps, t2 3 ps and t3= 9 ps. (a)-(d) for c-SiO2, (e)-(g) for c-SiO2-H, (h)-(j) a-SiO2-H for a-SiO2 and (k)-(m) for a-Si-O-H. In all four structures, Si-O bond breaks while Si-Li and Li-O bond forms as the lithiation time increases.
We also monitored the evolution of atomic coordination for different Li concentrations. Figure 3 shows the average values of Si-O coordination ⟨Si−O⟩ and Li-O coordination ⟨Li−O⟩ as a function of Li composition for c-SiO2. The coordination number includes Si−O atomic bonds whose distance lies within 15% of the covalent length in the bulk crystal SiO2, which corresponds to bond distances less than 1.87 Å. A similar definition for Li-O bonds, taken from the Li-O bond length in the lithium silicates Li2Si2O5 and Li4SiO4,30 sets the upper cutoff in that case at 2.15 Å. The Si-O coordination steadily decreases and changes almost linearly, which indicates a constant rate of Si−O bond breaking. In addition, the Li-O coordination remains approximately 2 for a wide range of Li concentrations. This could indicate the presence of local characteristics of Li2O upon lithium insertion into SiO2. The decrease in Si-O coordination number and formation of Li2O clusters due to changes of Si-O bond and local reduction of an O atom by two Li atoms, respectively, has been reported by Ban et al.12 In addition, Zhang et al.13 reported a change in coordination number for a-SiO2 similar to our finding for c-SiO2. In summary, structural evolution during lithiation appears to be similar for each SiO2 structural type, i.e. the Si-O bond breaks during Li-Si and Li-O formation. Our result compliments the conclusions reached by Johari et al.24 on the structural transformation of crystalline and amorphous silicon during lithiation.
Figure 3. The average values of Si−O coordination and Li−O coordination as function of Li content in crystalline SiO2.
Finally, we also studied the diffusivity for Li ions at room temperature in c-SiO2, c-SiO2-H, a-SiO2 and a-SiO2-H. To determine the diffusivity, we calculated the average MSD of Li atoms as a function of AIMD time step for different temperatures as shown in Figure S1 of the Supplementary Information. The simulation temperatures were chosen to lie below the melting point of SiO2, 1983 K. Figure S2 demonstrates the steps we followed to extrapolate room temperature (300 K) diffusion coefficients using the Einstein relation MSD = 6Dt and the standard expression for diffusivity, D = Doexp(-Ea/kBT), where Ea is the activation energy, kB is the Boltzmann constant, and T is the temperature. Figure 4 shows the diffusion coefficient for Li ions at room temperature as well as the extrapolation for c-SiO2, c-SiO2-H, a-SiO2 and a-SiO2-H. We found the diffusivity of Li, DLi, in c-SiO2 and a-SiO2 at room temperature to be 2.368 x 10-9 and
0.086 x 10-9 m2/s with an energy barrier of 0.046 eV and 0.146 eV, respectively. These are higher diffusion constant values for lithium transport compared to other oxides, (anatase TiO2: 4.7x1016 m2/s, amorphous Al2O3: (2.7 × 10–14) m2/s, ZrO2: 1.7x10-22 m2/s , MgO: 1.7x10-30 m2/s).31-33 The diffusion in c-SiO2 is anisotropic and that the value reported here is along the fast diffusion path <001>.17 Our result of a lower diffusion barrier for c-SiO2 than a-SiO2 agrees with a previous report by van Duin et al.17 using ReaxFF Reactive Force Field Modeling. A similar trend in the diffusion barrier and coefficient were observed in hydrogen terminated SiO2: 2.748 x 10-9 m2/s and 0.040 eV for c-SiO2-H and 0.421 x 10-9 m2/s and 0.099 eV for a-SiO2-H.
a
b
c
d
Figure 4. Derived diffusivity of Li at T=300 K in Li2SiO2. (a) c-SiO2, (b) c-SiO2-H and their respective insets (right). (c)a-SiO2 and (d) a-SiO2-H. The diffusivity was obtained by fitting the mean square displacement using D = Do exp(-Ea /kbT) . c-SiO2 has lower diffusion than a-SiO2 with or without termination.
Conclusion In summary, using ab initio molecular dynamics simulations at finite temperature, we investigated the structural evolution and diffusion kinetics of crystalline and amorphous SiO2 with and without termination during lithiation. Our results reveal several key findings. First, the mechanism and structural evolution are similar in each structural type. The Si-O bond breaks, while the Li-O bond forms with a consistent coordination number of 2, similar to Li2O. Second, both c-SiO2 and c-SiO2-H are amorphized after lithiation. The third finding indicates that while amorphous oxides generally offer faster ion transport, Li diffusion along the <001> direction in cSiO2 remains faster than that in a-SiO2 by an order of magnitude irrespective of surface termination. These suggests that termination has negligible impact on both on the rate of lithium transport through SiO2 and the mechanism for structural evolution during lithiation. Last but not least, the amorphization of c-SiO2 and c-SiO2-H suggests that the kinetics for Li transport are similar to amorphous structures after the first cycle of lithiation. In general, this work enhances our understanding of lithiation mechanism and ionic conduction in silica. In addition, the results lend insights and motivation to further study conversion oxides in quest to find high performance anodes, solid electrolytes, and coatings in Li-ion batteries.
Acknowledgment I.I.A., C.J., B.M., and T.P.D. acknowledge support from the U.S. Department of Energy (DOE), Office of Basic Energy Sciences, Division of Materials Sciences and Engineering, under Contract No. DE-AC02-76SF00515, Division of Materials Sciences and Engineering at Stanford and SLAC for theoretical analysis. The computational work used resources at the National Energy Research Scientific Computing Center (NERSC), a U.S. Department of Energy Office of Science User Facility operated under Contract No. DE-AC02-05CH11231. We thank C. D. Pemmaraju, M.F. Toney, C. Cao and H.-S. Steinrück for helpful discussions. I.I.A. also would like to thank Stanford EDGE fellowship program. Author contributions I.I.A., C. J. J., B. M. and T. P. D. conceived the idea and designed the project. I.I.A., performed first principles calculations and discussed the results with C.J.J, B.M., T.P.D. All authors discussed the results, co-wrote and contributed to the manuscript. Declaration of interests The authors declare no competing interests. References [1] W.J Zhang, Journal of Power Sources 196 (2011) 13-24. [2] M. Ko, S. Chae, J. Cho, ChemElectroChem 11 (2015) 1645-1651. [3] X. Zuo, J. Zhu, P. Müller-Buschbaum, Y.J. Cheng, Nano Energy 31 (2017) 113-143.
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Graphical abstract
Highlights
SiO2 is a promising candidate material for the anode of Li-ion batteries Ab initio molecular dynamics was used to study the lithiation mechanism in crystalline and amorphous SiO2 Effect of surface termination on lithiation was investigated Li diffusion along the <001> direction in c-SiO2 is fastest compared to other oxides Our results lend insights and motivation to further study conversion oxides in quest to find high performance anodes, solid electrolytes, and coatings in Li-ion batteries.
Declaration of interests The authors declare no competing interests.