Nanoscale Mechanics of the Solid Electrolyte Interphase on Lithiated

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Nanoscale Mechanics of the Solid Electrolyte Interphase on Lithiated-Silicon Electrodes Haoran Wang, and Huck Beng Chew ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.7b07626 • Publication Date (Web): 21 Jul 2017 Downloaded from http://pubs.acs.org on July 21, 2017

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ACS Applied Materials & Interfaces

Nanoscale Mechanics of the Solid Electrolyte Interphase on Lithiated-Silicon Electrodes Haoran Wang, Huck Beng Chew* Department of Aerospace Engineering, University of Illinois Urbana-Champaign, Urbana, IL 61801

*Corresponding author: [email protected] Abstract The ~300% volume changes of lithiated silicon electrodes (LixSi) during electrochemical cycling lead to cracking of the solid electrolyte interface (SEI). Here, we report how strain is transferred from LixSi to two primary inorganic SEI components: LiF and Li2O. Our first principle calculations show that LiF, effectively bonded on LixSi at x > 1, enables the entire interface structure to deform plastically by forming delocalized stable voids. In contrast, Li2O tightly bonded to LixSi is stiffer, and deforms rigidly across all x. Our results explain the significantly improved ductility of SEI with higher LiF versus Li2O content observed experimentally. KEYWORDS: Solid electrolyte interphase; silicon electrodes; mechanical deformation; lithium ion batteries; density functional theory

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Silicon electrodes hold the promise for next generation lithium-ion batteries because of the order-of-magnitude higher energy density compared to conventional graphite electrodes.1 A fully-lithiated silicon electrode, however, expands volumetrically by ~300%.2 This leads to massive cracking of the silicon electrode and its subsequent delamination from the current collector.2-5 Studies have proposed the use of nanostructured silicon to mitigate cracking,6 but such electrodes have much larger relative surface area exposed to the electrolyte. The capacity retention of the battery then depends on the integrity of the ~10 to 50 nm thick passivation film separating the silicon electrode from the electrolyte, termed the solid electrolyte interphase (SEI).7,8 The SEI formed on conventional graphite electrodes is very stable and can protect the electrodes for thousands of cycles. In contrast, the ~300% volume changes of silicon electrodes during charge cycling can crack the SEI film, which exposes the underlying electrode surface to the electrolyte causing new SEI to form.9 Repeated cracking and re-forming of SEI leads to the irreversible consumption of Li ions, and subsequent capacity fade.10 The SEI comprises of a soft porous organic layer on the side of the electrolyte and a stiffer, denser inorganic layer on the side of the silicon electrode.10,11 The inorganic layer provides the primary functionality of the SEI, and is made up of LiF, Li2O, Li2CO3, and lithium alkyl carbonate phases which are believed to be assembled in a mosaic-like pattern.8,10,12 Recent studies have also proposed the engineering of artificial SEIs on LixSi based primarily on LiF or Li2O.13,14 Here, we also focus only on LiF and Li2O as the primary SEI components of interest, since Li2CO3 and lithium alkyl carbonate are less stable phases which decompose to Li2O + CO2.15-17 There is now increasing evidence relating the structural integrity of the SEI to the compositional proportions of LiF and Li2O (or Li2CO3) within the inorganic layer.7,15,18 For example, addition of fluoroethylene carbonate (FEC) in the electrolyte has been shown to 2 ACS Paragon Plus Environment

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improve the chemical stability and structural integrity of the SEI due to the more abundant presence of LiF.19,20 The composition of the SEI was also demonstrated to depend on the initial cycling conditions, with stiffer SEI films attributable to Li2O forming at lower potentials, while softer and more strain-tolerant SEI films associated with LiF forming at higher potentials.7,12 To elucidate the mechanisms underlying the above experimental observations, we perform density functional theory (DFT) calculations to characterize the interfacial properties of LiF and Li2O on LixSi. Our LiF-LixSi and Li2O-LixSi supercells comprise of (100)-oriented crystalline slabs of LiF or Li2O directly bonded to slabs of amorphous LixSi, as shown in Fig. 1, with x = 0, 0.5, 1, and 3.75 representing four different lithiation stages. These interface structures were created via a rapid heating and quenching scheme based on DFT formalism,21 as outlined in Methods. Note that two distinct interfaces (dashed lines) exist for each periodic model structure. For LiF-LixSi, both these interfaces exhibit similar bonding characteristics since the amorphous LixSi is exposed to the same (100)-oriented LiF surface structure. For Li2O-LixSi, however, Liand O-terminated interfaces are possible on the side of the (100)-oriented Li2O crystal and we consider both these interface types in our periodic model structures. To provide a general overview of the atomic interactions, we also include contours of the electron localized function (ELF) for the various LiF-LixSi and Li2O-LixSi interface structures in Fig. 1. Qualitatively, a large ELF value close to 1 (red) corresponds to a region with high probability of finding electron localization as in covalent bonding, whereas an ELF close to 0.5 (green) corresponds to a region of electron gas-like behavior as in metallic bonding. An intermediate ELF value of 0.75 (yellow) can be interpreted to correspond to ionic bonding. For LiF-Si model structure (Fig. 1a-i), Si-Si covalent bonding and Li-F ionic bonding are dominant within the bulk Si and LiF phases respectively, while there is lack of significant 3 ACS Paragon Plus Environment


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interfacial bonding across these two phases with exception of isolated Si-F bonds. Increasing extents of lithiation (x), however, allows Li atoms to segregate to the LiF-LixSi interface to form an increasingly seamless interface (Fig. 1a-ii to 1a-iv). These segregated Li atoms act as “glue” between LiF and LixSi,18 by forming strong ionic bonds with crystalline LiF on one side of the interface and with Si atoms on the other side of the interface. In the case of Li2O-Si (Fig. 1b-i), strong Si-O and Si-Li ionic bonding is observed across both the O- and Li-terminated interfaces. Increasing x also increases the segregation of Li atoms to the O-terminated interface (Fig. 1b-ii to 1b-iv); the Si-O bonds across the interface are gradually replaced by O-Li-Si bonds. Li atoms completely saturate the O-terminated interface for Li2O-Li15Si4 (Fig. 1b-iv), which effectively transforms this interface type to a Li-terminated one. We quantify the work of separation Γ along the interfaces of LiF-LixSi and Li2O-LixSi, as a function of lithiation extent x, in Fig. 2. Here, Γ is defined as the change of energy per unit area caused by rigid separation along the selected interface (dashed lines in Fig. 1) followed by quantum-mechanical relaxation of the separated structure in DFT. We average Γ of the two interfaces within each LiF-LixSi model structure, while we separately calculate Γ of the Li- and O-terminated interfaces for Li2O-LixSi. For comparison purposes, we also include the average Γ taken along bulk LixSi in Fig. 2.21 Observe that Γ for LiF-Si is close to 0 due to negligible bonding between LiF and Si. However, Γ increases with x because of increasing Li segregation to the interface (blue line). Once Li atoms fully saturate the LiF-LixSi interface, as in x = 3.75, Γ associated with Li-Li and Li-Si bonding across the interface then approaches that of bulk Li15Si4. In the case of Li-terminated interface for Li2O-LixSi (green line), Γ decreases with increasing x, since stronger ionic Li-Si bonds are increasingly replaced by weaker metallic Li-Li bonds. While the same general trend is exhibited by bulk LixSi (black line), Γ for the Li-terminated interface of 4 ACS Paragon Plus Environment

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Li2O-LixSi remains consistently higher because of the highly-ordered arrangement of Li atoms on the side of the Li2O crystal. The interfacial energy associated with the O-terminated interface for Li2O-Si is almost 3-fold higher than that of bulk Si, because of the highly-ordered Si-O bonding across the interface. At x = 0.5, the segregation of Li atoms to the interstitial sites along the Li2O-LiSi2 interface further improves the interfacial bonding, as shown by the slight increase in Γ. However, increased segregation of Li atoms to the interface results in a dramatic drop in Γ at x = 1, due to preferential Li-O bonding (bond length ~2.1 Å) over Si-O bonding (bond length ~1.6 Å) associated with the high coordination number of Li atoms at the interface (see ELF contours in Fig. 1b-iii). Once sufficient Li atoms have segregated to the interface at x = 3.75, Γ for both interface types are now similar, confirming that the original O-terminated interface is now effectively Li-terminated. Herein, we focus only on the O-terminated interface of Li2OLixSi because of the similarity in bonding structure between O-terminated Li2O-Li15Si4 and Literminated Li2O-LixSi. We subject the LiF-LixSi and Li2O-LixSi interface structures to uniaxial tensile deformation by uniformly stretching the out-of-plane x2 dimension of the supercells in 1% strain increments, while constraining the in-plane x1-x3 dimensions. After each strain increment, atoms bounding the top/bottom interface of both LiF-LixSi and Li2O-LixSi are rigidly fixed, while the remaining atoms are subjected to quantum-mechanical relaxation in DFT. Fig. 3 shows the von Mises effective strain-strain relationship for LiF-LixSi (blue line) and Li2O-LixSi (red line), as compared to bulk LixSi (black line),21 at various lithiation instances. The applied tensile strain of the interface structures can be accommodated by both the bulk LixSi and Li2O/LiF phases, as well as by interface separation. For LiF-Si (Fig. 3a), the weak interface separates easily resulting in low effective strength. The interfacial bonding between the LiF and LixSi phases becomes 5 ACS Paragon Plus Environment


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stronger at higher x, but the effective stress remains consistently lower than that of bulk LixSi at the same applied strain. Bulk LixSi is known to exhibit significant plasticity by nanopore growth at x ≥ 1; these nanopores initiate by the breaking of Li-Si and Li-Li bonds.21,22 For LiF-LiSi and LiF-Li15Si4 (Fig. 3c and 3d), the ~2-fold lower tensile strength suggests that considerable strain is accommodated by the LiF phase and/or by interface separation, in addition to the plasticallydeforming LixSi. In contrast, the effective stress-strain curves for Li2O-LixSi structures up to x = ~0.5 are comparable to bulk LixSi (Fig. 3a and 3b). Beyond x ≥ 1, the consistently higher effective stress for Li2O-LixSi versus the plastically-deforming bulk LixSi suggests that the companion Li2O phase deforms rigidly with limited interface separation. We examine the deformed atomic configurations for LiF-LixSi and Li2O-LixSi at a fixed applied strain of ߝଶଶ = 0.2 in Fig. 4. For clarity, we highlight the porous low-bond-density regions in grey.26 Observe that the weakly-bonded LiF-Si interface separates cleanly (Fig. 4a-i). In the case of LiF-LiSi2 (Fig. 4a-ii), distinct interfacial voids separated by chains of F-Li-Si bonds now develop. For LiF-LiSi, deformation is no longer confined to the interface, allowing significant strain to be accommodated by both the LiF and LiSi phases (Fig. 4a-iii). Pores now develop along the LiF-LiSi interface, as well as within the LiF phase by the breaking of Li-F bonds, but limited porosity is observed within the LiSi phase which suggests the earlier onset of plasticity for LiF versus LiSi. For LiF-Li15Si4, stable pores now develop homogeneously within LiF and Li15Si4, as well as along the interface (Fig. 4a-iv. This enhanced plasticity of LiF-Li15Si4 results in significant softening of the stress-strain response in Fig. 3d, since deformation can now be accommodated in-part by the presence of these voids. In contrast to the weakly-bonded LiF-Si or LiF-LiSi2 interface structures, the much stronger interfacial bonding for Li2O-LixSi across all x allows strain to be accommodated by both Li2O and LixSi phases. Unlike LiF, however, the 6 ACS Paragon Plus Environment

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much stiffer Li2O phase only deforms elastically even at applied strains of ߝଶଶ = 0.2. For Li2OSi, the high ~8 GPa yield strength of Si allows both phases to elastically stretch (Fig. 4b-i).25 Increasing x, however, significantly lowers the yield strength to ~3.4 GPa for LiSi2, ~2.5 GPa for LiSi, and ~1.4 GPa for Li15Si4.25 Coupled with the two- to three-fold larger stiffness of Li2O versus LixSi, strain becomes increasingly localized within the LixSi phase with increasing x, causing fracture to consistently develop within the LixSi phase (Fig. 4b-ii to 4b-iv). Recent experiments report rapid SEI growth between 10% (Li0.4Si) and 30% (Li1.1Si) lithiation state of the silicon electrode, after which the SEI film stabilizes and Li insertion in the electrode dominates Li uptake.23 Our simulations clearly show the segregation of Li atoms along both LiF-LixSi and Li2O-LixSi interfaces; this segregation of Li atoms along the respective interfaces should precede significant lithiation of the silicon electrodes, in agreement with experiments. In the case of LiF-LixSi, the weak bonding between LiF and LixSi for ‫ < ݔ‬0.5 suggests that LiF can only form on LixSi with ‫ ≥ ݔ‬0.5. In contrast, Li2O strongly bonds with LixSi across all x. At ‫ < ݔ‬0.5, ionic Si-O bonds form between the O-terminated interface of Li2O and LixSi; at ‫ ≥ ݔ‬0.5, Li segregates at the interface and form Li-Si bonds by breaking Si-O bonds. Therefore, the formation of Li2O over LiF is expected to be more favorable during the early stages of lithiation (small x) or at slow lithiation rates. This is in agreement with experimental studies which report the formation of more uniform SEI containing higher compositional proportion of LiF with fast lithiation of the silicon electrode.7 Studies have also subjected fully-lithiated Si nanoparticles to fluorinated compounds to create LiF-based artificial SEI, which has been proven to be highly effective in maintaining the cyclic capacity;13 our results suggest that pre-lithiation is in fact a necessary condition to achieve proper adhesion between LiF and LixSi. We remark that our results here pertain to the lithiation of pure 7 ACS Paragon Plus Environment


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amorphous Si without its native oxide layer. Previous studies have shown that the SiO2 layer on Si can be very stable, and cause high first-cycle capacity loss.20,24 On the other hand, reducing this SiO2 layer by chemical etching significantly improves the electrochemical performance of the battery.25 The contributing effect of this native oxide layer is a subject of future work. Experiments have consistently shown that increasing the composition of LiF in SEI, via the addition of FEC in the electrolyte or by tuning the cycling conditions, significantly improves the fracture resistance of SEI and the cycle life of the battery.19,20,26 For a strongly-bonded SEI film on LixSi, the stress built-up in the SEI caused by volumetric changes of LixSi scales with the stiffness of the SEI, considering the deformation to be elastic. The [100]-oriented LiF has an elastic modulus of ~81.9 GPa, which is comparable to that of LixSi. As shown in Fig. 4, LiF readily deforms plastically through the formation of delocalized stable voids, which significantly limits the build-up of stresses and results in improved damage tolerance. Compared to LiF, [100]-oriented Li2O has a two-fold higher elastic modulus of ~197.6 GPa, and deforms only elastically. The rigid Li2O phase in SEI therefore places significant constraints on the expanding/contracting of LixSi electrode, which exacerbates the build-up of interfacial stresses to cause cracking. In summary, our first principle calculations reveal contrasting mechanical responses of LiFLixSi and Li2O-LixSi interface structures, which are relevant to the deformation of SEI on silicon electrodes. We show that LiF only bonds with LixSi after sufficient Li segregation to the LiFLixSi interface with x > ~1. Once formed, LiF imparts significant ductility to the SEI film by inducing plasticity through the formation of delocalized stable voids within the entire interfacial structure. In contrast, Li2O imparts rigidity to the SEI film, and thus is less strain-tolerant and promotes brittle-like interfacial failure. Our results explain the experimentally-observed 8 ACS Paragon Plus Environment

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improvements in fracture toughness of the SEI films with increasing compositional proportion of LiF versus Li2O. Understanding the relationship between the composition of the SEI and its deformation response provides important insights into the engineering of artificial SEI layers on silicon electrodes with improved damage tolerance. Methods:

Our DFT calculations are performed using the Vienna Ab initio Simulation

Package (VASP), with the Projector-Augmented-Wave (PAW) method and Perdew-BurkeErnzerhof (PBE) form of the generalized gradient approximation (GGA) for exchange and correlation. To create the LiF-LixSi and Li2O-LixSi interfacial structures, we start with initial structures comprised of evenly mixed LixSi slabs and the crystalline LiF or Li2O slabs placed above. The predefined initial dimensions of the LiF-LixSi and Li2O-LixSi supercells (Table 1) are determined based on the stress-free volume of crystalline LiF and Li2O, as well as the expanded volume of the Si electrodes when lithiated.27,28 Using ab-initio molecular dynamics (MD) in VASP with a plane wave energy cutoff of 300 eV, we freeze the top LiF or Li2O slabs and create the amorphous LixSi structure by subjecting the bottom LixSi slabs to a temperature of 3000 K while keeping the dimensions of the supercell fixed; the temperature is maintained with a Nose thermostat for 3000 MD time steps, with each time step of 3 fs. This temperature far exceeds the melting point of Si, Li, and their respective compounds and allows for sufficient intermixing between the Si and Li atoms, while confined within the fixed, periodic LiF or Li2O slabs. Thereafter, we quench the temperature of the LixSi slabs to the target temperature of 300 K, at a rate of 1 K per MD time step. The entire structure, including the LiF and Li2O slabs, is then equilibrated at 300 K for an additional 1000 MD time steps, which allows significant rearrangement of atoms near the LiF-LixSi or Li2O-LixSi interfaces. Thereafter, we relax the shape and volume of the supercell to achieve a stress-free configuration, with the final 9 ACS Paragon Plus Environment


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dimensions of the supercell as shown in Table 1. Finally, the supercell is quantum-mechanically relaxed to its local minimum energy state with DFT, using an atomic force tolerance of 0.01 eV/nm and a plane wave expansion kinetic energy cutoff of 500 eV. This higher plane wave energy cutoff allows for more precise calculation of the stress tensors and interfacial energies. In all our calculations, we employ a Gamma-centered 3×3×1 uniform Monkhorst-Pack k-point sampling taken over the entire Brillouin zone. Table 1. Modeling details of the LiF-LixSi and Li2O-LixSi interface structures

Structure

Li/Si ratio

No. of atoms Si Li F/O

Initial model dimensions (x×y×z Å3)

Final stress-free model dimensions (x×y×z Å3)

LiF-Si

0.00

64

75

75

9.9×9.9×23.0

10.3×10.3×25.3

LiF-LiSi2

0.50

52

101

75

9.9×9.9×26.0

10.0×10.4×26.2

LiF-LiSi

1.00

40

115

75

9.9×9.9×26.0

10.3×10.1×26.1

LiF-Li15Si4

3.75

20

150

75

9.9×9.9×29.0

10.4×10.2×28.1

Li2O-Si

0.00

64

96

48

9.2×9.2×25.0

9.5×9.5×25.4

Li2O-LiSi2

0.50

52

122

48

9.2×9.2×28.0

9.3×9.4×28.6

Li2O-LiSi

1.00

40

136

48

9.2×9.2×28.0

9.5×9.4×29.0

Li2O-Li15Si4

3.75

20

171

48

9.2×9.2×31.0

9.3×9.3×31.7

Acknowledgements The authors acknowledge the support of National Science Foundation Grant No. NSF-CMMI1300805, as well as computational time provided by TACC Grant No. TG-MSS130007, and Blue Waters sustained-petascale computing project which is supported by the National Science Foundation (Award Nos. OCI-0725070 and ACI-1238993) and the state of Illinois. Blue Waters is a joint effort of the University of Illinois at Urbana–Champaign and its National Center for Supercomputing Applications. 10 ACS Paragon Plus Environment

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(19) Lindgren, F.; Xu, C.; Niedzicki, L.; Marcinek, M.; Gustafsson, T.; Bjorefors, F.; Edstrom, K.; Younesi, R. SEI Formation and Interfacial Stability of a Si Electrode in a LiTDI-Salt Based Electrolyte with FEC and VC Additives for Li-Ion Batteries. ACS Appl. Mater. Interfaces 2016, 8 (24), 15758-15766. (20) Young, B. T.; Heskett, D. R.; Nguyen, C. C.; Nie, M.; Woicik, J. C.; Lucht, B. L. Hard X-ray Photoelectron Spectroscopy (HAXPES) Investigation of the Silicon Solid Electrolyte Interphase (SEI) in Lithium-Ion Batteries. ACS Appl. Mater. Interfaces 2015, 7 (36), 20004-20011. (21) Wang, H.; Wang, X.; Xia, S.; Chew, H. B. Brittle-to-ductile transition of lithiated silicon electrodes: Crazing to stable nanopore growth. J. Chem. Phy. 2015, 143 (10), 104703. (22) Wang, H.; Chew, H. B. Molecular dynamics simulations of plasticity and cracking in lithiated silicon electrodes. Extreme Mech. Lett. 2016, 9 (3), 503-513. (23) Bordes, A.; De Vito, E.; Haon, C.; Secouard, C.; Montani, A.; Marcus, P. Investigation of Lithium Insertion Mechanisms of a Thin-Film Si Electrode by Coupling Time-of-Flight Secondary-Ion Mass Spectrometry, X-ray Photoelectron Spectroscopy, and Focused-Ion-Beam/SEM. ACS Applied Materials & Interfaces 2015, 7 (50), 27853-27862. (24) Ferraresi, G.; Czornomaz, L.; Villevieille, C.; Novak, P.; El Kazzi, M. Elucidating the Surface Reactions of an Amorphous Si Thin Film as a Model Electrode for Li-Ion Batteries. ACS Appl. Mater. Interfaces 2016, 8 (43), 29791-29798. (25) Xun, S.; Song, X.; Wang, L.; Grass, M. E.; Liu, Z.; Battaglia, V. S.; Liu, G. The Effects of Native Oxide Surface Layer on the Electrochemical Performance of Si Nanoparticle-Based Electrodes. Journal of the Electrochemical Society 2011, 158(12), A1260-A1266. (26) Xu, C.; Lindgren, F.; Philippe, B.; Gorgoi, M.; Bjorefors, F.; Edstrom, K.; Gustafsson, T. Improved Performance of the Silicon Anode for Li-Ion Batteries: Understanding the Surface Modification Mechanism of Fluoroethylene Carbonate as an Effective Electrolyte Additive. Chem. Mater. 2015, 27 (7), 2591-2599. (27) Jain, A.; Ong, S. P.; Hautier, G.; Chen, W.; Richards, W. D.; Dacek, S.; Cholia, S.; Gunter, D.; Skinner, D.; Ceder, G.; Persson, K. A. Commentary: The Materials Project: A materials genome approach to accelerating materials innovation. APL Mater. 2013, 1 (1), 011002. (28) Johari, P.; Qi, Y.; Shenoy, V. B. The Mixing Mechanism during Lithiation of Si Negative Electrode in Li-Ion Batteries: An Ab lnitio Molecular Dynamics Study. Nano Lett. 2011, 11 (12), 5494-5500.


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Figure Captions Figure 1. Atomic configurations and contours of the electron localized function for (a) LiF-LixSi and (b) Li2O-LixSi interface structures, with (i)-(iv) representing x = 0, 0.5, 1, and 3.75. Blue dashed lines denote planar cuts representing LiF-LixSi interfaces; magenta and red dashed lines denote planar cuts representing Li- and O-terminated Li2O-LixSi interfaces. Figure 2. Interface energy Γ versus Li fraction within the amorphous LixSi phase for bulk LixSi (black lines), LiF-LixSi (blue), and O- and Li-terminated Li2O-LixSi (red and magenta lines). Calculation of Γ taken along planar cuts denoted by corresponding colored dashed lines in Fig. 1.

Figure 3. Von Mises effective stress versus applied tensile strain for LiF-LixSi (blue curves), Li2O-LixSi (red curves), and bulk LixSi (black curves) with (a)-(d) representing x = 0, 0.5, 1, and 3.75. Figure 4. Deformed atomic configurations and corresponding porosity distributions (highlighted in gray) within (a) LiF-LixSi and (b) Li2O-LixSi interface structures at a fixed applied strain of ɛ22 = 0.2, with (i)-(iv) representing x = 0, 0.5, 1, and 3.75.



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Figure 1. Atomic configurations and contours of the electron localized function for (a) LiF-LixSi and (b) Li2O-LixSi interface structures, with (i)-(iv) representing x = 0, 0.5, 1, and 3.75. Blue dashed lines denote planar cuts representing LiF-LixSi interfaces; magenta and red dashed lines denote planar cuts representing Li- and O-terminated Li2O-LixSi interfaces.

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Figure 2. Interface energy Γ versus Li fraction within the amorphous LixSi phase for bulk LixSi (black lines), LiF-LixSi (blue), and O- and Li-terminated Li2O-LixSi (red and magenta lines). Calculation of Γ taken along planar cuts denoted by corresponding colored dashed lines in Fig. 1.



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Figure 3. Von Mises effective stress versus applied tensile strain for LiF-LixSi (blue curves), Li2O-LixSi (red curves), and bulk LixSi (black curves) with (a)-(d) representing x = 0, 0.5, 1, and 3.75.


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Figure 4. Deformed atomic configurations and corresponding porosity distributions (highlighted in gray) within (a) LiF-LixSi and (b) Li2O-LixSi interface structures at a fixed applied strain of ɛ22 = 0.2, with (i)-(iv) representing x = 0, 0.5, 1, and 3.75.



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Abstract Figure 188x113mm (150 x 150 DPI)


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Nanoscale Mechanics of the Solid Electrolyte Interphase on Lithiated

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