In Situ and Operando Investigations of Failure Mechanisms of the

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In Situ and Operando Investigations of Failure Mechanisms of the Solid Electrolyte Interphase on Silicon Electrodes Ravi Kumar, Anton Tokranov, Brian W. Sheldon, Xingcheng Xiao, Zhuangqun Huang, Chunzeng Li, and Thomas Müller ACS Energy Lett., Just Accepted Manuscript • DOI: 10.1021/acsenergylett.6b00284 • Publication Date (Web): 01 Sep 2016 Downloaded from http://pubs.acs.org on September 3, 2016

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ACS Energy Letters

In Situ and Operando Investigations of Failure Mechanisms of the Solid Electrolyte Interphase on Silicon Electrodes Ravi Kumar†, Anton Tokranov†, Brian W. Sheldon†,*, Xingcheng Xiao§, Zhuangqun Huang‡, Chunzeng Li‡, Thomas Mueller‡ †

Brown University - School of Engineering, 182 Hope Street, Box D, Providence, Rhode Island

02912, United States §

General Motors Global R&D Center, 30500 Mound Road, Warren, Michigan 48090, United

States ‡

Bruker Nano Surfaces, 112 Robin Hill Road, Goleta, California 93117, United States

AUTHOR INFORMATION Corresponding Author * [email protected]

ACS Paragon Plus Environment

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ABSTRACT The lifetime of rechargeable lithium-ion batteries is closely related to the formation and evolution of the solid electrolyte interphase (SEI). These passivation films undergo substantial deformations when the underlying electrode particles expand and contract during cycling. Directly probing these changes is extremely challenging. In this study we demonstrate a new approach for applying controlled strains to SEI films with patterned Si electrodes, in conjunction with direct observations of mechanical degradation using in operando atomic force microscopy. Monitoring both strain and crack formation in SEI provides new in-depth understanding of SEI fracture. These results verify that crack formation occurs during lithiation (this has been predicted previously, but not directly observed). Additional SEI formation at low potentials did not fill these cracks, which directly contradicts prior speculation. These experiments also made it possible to estimate the fracture toughness of the SEI (a key value which has not been previously measured).

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In Li-ion batteries, passivation films on the surfaces of electrode materials are critical. This layer, usually referred to as the solid electrolyte interphase (SEI), must retain its integrity when the underlying active material expands and contracts during electrochemical cycling. The coupled mechanical and chemical degradation of the SEI is a critical roadblock for developing Si and other high capacity negative electrodes in Li-ion batteries. A large variety of characterization techniques including Electrochemical impedance spectroscopy (EIS),1–3 Raman spectroscopy,4,5 X-ray scattering,6 X-ray photoelectron spectroscopy (XPS),7 Atomic force microscopy (AFM)8– 10

and Secondary ion mass spectrometry (SIMS)3,11 have been used to characterize the formation,

evolution and functionality of SEI layers. However, this existing knowledge is largely focused on the SEI chemistry. In contrast there is very little direct information about SEI deformation and mechanical failure mechanisms, linking these phenomena to SEI chemical composition and structure. There has however been considerable speculation about the mechanical degradation of SEI films, particularly with Si electrodes which undergo extremely large volume expansions during Li insertion (> 300%). Here, many researchers have proposed that electrochemical cycling in rechargeable batteries should lead to repeated tensile and compressive strains, and that this in turn could induce breaking and reforming of SEI, leading to poor passivation and thick films which ultimately lead to high impedance, irreversible Li consumption, and limited battery lifetime.12 While some kind of connection between strain and SEI degradation is logical, there are a number of steps between the basic behavior of the electrode material (i.e., expansion and contraction, SEI formation, etc.) and the cycle life of the battery. Mechanical degradation of SEI films is critical in this sequence (particularly with Si electrodes), but to date there is virtually no direct information about failure mechanisms in the SEI. This is largely because of the inherent


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difficulties that are associated with observing the SEI as Li is inserted and removed from the underlying electrode materials. Three primary challenges here are: (1) The film thicknesses are typically on the order of 100 nm, which means that high spatial resolution is needed to observe failure processes (i.e., nanoscale cracking, delamination, spallation, etc.). (2) In situ and operando measurements in liquid electrolytes are needed to observe how and when mechanical failures occur. (3) A well-defined specimen configuration is needed, to interpret the deformation that occurs in the SEI. We have successfully addressed all of these issues with in situ and operando atomic force microscopy (AFM) imaging of patterned Si films. Prior work on SEI does include some important information on mechanical behavior. For example, several recent efforts have explored mechanical mapping of the SEI layer with AFM-based indentation tests.13–15 These studies provide important information about SEI deformation, but specific failure mechanisms were not directly observed. Key questions here are whether the SEI layers fail during lithiation or delithiation, and whether it cracks, delaminates or spalls off from the anode surface. As noted above, most prior work on SEI failure is speculative, based largely on schematics which propose that intercalation induced expansion should induce cracking of the SEI layer.12,16,17 One notable exception is Zhang et al.18 who studied SEI morphology evolution on the surface of graphite particles using ex situ focused ion beam (FIB) milled specimens. These images reveal cracks in very thick SEI layers (450-1600 nm), which they attribute to internal pressure from gas production. Pioneering work by Aurbach et al.19,20 also employed in situ AFM to investigate composite graphite electrodes, where they proposed that stretching of the surface films caused by


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Li intercalation leads to continuous SEI formation and impedance rise. They also showed cracking of graphite particles (through the SEI layer), and subsequent crack filling during additional cycling (apparently due to further SEI formation). These prior investigations are important, but they do not provide direct information about failure mechanisms in the SEI. With ex situ methods, it is very difficult to identify the point where cracking in the SEI occurs – partly because the cracks are only observed after the fact, and also because the SEI layers are expected to change when they are removed from the liquid electrolyte and cut with the FIB. Either of these specimen preparation steps can introduce cracks in the SEI. Also, the graphite particles used in prior studies contain a substantial number of internal defects, which can initiate cracks that propagate into the SEI. While this type of failure is potentially important, it is distinctly different than direct fracture of the SEI. The difference between these failure modes also highlights the fact that a variety of SEI failure processes are potentially relevant. The use of in situ AFM to characterize the electrochemical-mechanical behavior of SEI layers improves on prior studies in several key ways. New instrumentation permits substantially higher spatial resolution in liquid electrolytes, compared to the initial in situ AFM work that was done more than a decade ago. Here, the unique PeakForce tapping mode was used to image the extremely fragile SEI layer, which is very challenging with conventional AFMs in organic solvents.21 More importantly, we employ a novel approach that is based on patterned Si island structures. We previously demonstrated that thin patterned Si islands can undergo lateral expansion and contraction due to a shear lag effect.22–24 This provides an ideal platform to investigate the SEI layer response to mechanical strains that are induced by the underlying Si, with a well-defined geometric configuration in an electrochemical environment. These studies


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show for the first time the in operando nanocracking of an SEI layer on Si during cycling, along with precise measurements of SEI formation. Figure 1 shows a schematic of the experimental configuration. This includes a depiction of the stress distribution in a patterned Si island during cycling. The aspect ratio in these islands is ~75 (i.e., much larger than in the schematics), and thus basic shear lag theory provides a reasonable approximation of the stress distribution in the Si islands. Near the free edge there is shear traction between the film and the substrate, while far from the edge the film is in a uniform biaxial stress state with no traction acting across the interface. In the center of the islands, expansion occurs normal to the substrate. In the shear lag zone near the edge, this out-of-plane Si expansion occurs along with lateral island expansion. Our prior work22 with similar Si islands shows that the in-plane extension in the shear-lag zone is accommodated by interfacial sliding (with some possible contributions from plastic deformation of the Cu current collector and plastic flow of Si25). These islands are a convenient model system for studying the effect of lithiation induced volume expansions on SEI formation, evolution and failure modes without the complications associated with more practical composite electrodes that contain a conductive matrix and binder. The height evolution, lateral size changes and deformation behavior of the patterned Si island during cycling (observed via AFM surface topographs) are plotted in Figure 2a, 2b and 2(c-g) respectively. The behavior in the center of the island is similar to what has been observed previously.8 Here, there is significant increase in the measured height after Li insertion, due to both SEI formation and irreversible volume expansion of the Si. The Li induced expansion produces in-plane compressive stress22 in the Si electrode, however significant in-plane stress in the SEI is not expected because the Si only expands in the out-of-plane direction. Here, a stable


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SEI thickness in the center region occurs relatively quickly. In contrast, very different height evolution occurs at the edges and corners. During the first cycle lithiation to 0.6V, the height evolution at the edge and center is similar. At this potential, the SEI formation is largely associated with the formation of carbonaceous electrolyte decomposition products (sometimes referred to as “organic” SEI).

At lower

potentials, particularly below 0.3V, the edges of the island move laterally by ~1 µm. This provides direct evidence of the shear lag effect depicted in Figure 1. In the near edge region, the Si is also expected to undergo substantially lower out-of-plane expansion compared to the center (i.e., because the expansion of the material in the shear lag zone is accommodated by both outof-plane and lateral expansion/interfacial sliding). This is also confirmed from the observation of a significant contrast in the AFM topograph between the center and edge regions of the island at fully lithiated state (Figure 2d). This observation of lateral expansion differs from Dahn’s26 in situ AFM observations of patterned Si island where lateral dimensional changes were not observed. One possible reason for this difference is the use of stainless steel as a substrate for Si, which has different interfacial properties compared to the LixSi/Cu interface.27 In another study of patterned a-Si electrodes, almost 30% change in lateral dimensions was observed with much thicker islands (500nm) and Ti as current collector instead of Cu.28 Another study by Becker et al.29 of patterned Si nanostructures on Ni current collectors also showed lateral dimensional changes, and they concluded that the size, geometry and the interface properties all affected the island shape evolution. We believe that the variations observed in these different studies indicate that subtle differences in interface properties can have a substantial impact on the interfacial sliding resistance, . The variability noted above is not a concern for the experiments reported


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here, where our intent is only to use the lateral expansion of Si along the current collector as a means of inducing in-plane deformation of the SEI layers. Based on a basic force balance, the size of the shear lag zone is approximately ℎ / , where is the flow stress of Si and is the interfacial strength between lithiated Si and the Cu current collector. The latter is bounded by the shear flow stress of Cu and the interfacial friction strength i.e. ( = min( , )).22 Inserting the measured size of the shear lag zone (~ 1 µm from Fig. 2b) in the scaling law along with = 1 and ℎ = 500, gives an estimate of ~ 500 MPa. This value is higher than the interfacial friction strength of the interface predicted with atomistic simulations27 (~100 MPa for XLi~0.6), and lower than the flow stress of e-beam deposited Cu films30 (~875 MPa). Based on these comparisons, the Si expansion in the shear lag zone is likely to be accommodated by interfacial sliding and/or plastic flow in the current collector, similar to prior work.22 Because most of the Si expansion occurs after SEI formation, the lateral in-plane expansion in the shear lag zone is expected to create in-plane tensile stress in the SEI. The underlying Si in the shear lag zone is still in compression (but lower in magnitude than that at the center23). In Figure 2d and f nanoscale cracking of the SEI is clearly seen in the edge and corner regions at full lithiation, during the 1st and 2nd cycles. This type of cracking normal to the surface is characteristic of tensile stress in the film, which is expected to occur as the lateral expansion of the Si stretches the SEI. The region where cracking occurs also matches the expected size of the shear lag zone based on the measured lateral expansion of ~1 µm. This confirms that the SEI is subjected to a different stress state in the shear lag zone. While there is considerable speculation about SEI cracking in the literature,12,16 we believe that this is the first direct observation of this phenomenon during active lithiation. Cracks were observed in Aurbach’s pioneering in situ AFM


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work with graphite electrodes19, but here the stress state of the SEI was ambiguous and it is not clear if these cracks were initiated in the SEI or the underlying graphite. In our experiments, it is important to note that cracking should not occur in the underlying Si which is in compression during lithiation. During the first 0.1V hold, cracking begins to occur in the edge and corner regions. Most of the SEI forms at higher potentials before significant Li insertion occurs, such that tensile stress then evolves in the SEI layer as the Si expands at lower potentials. This tensile state is further substantiated when the size of crack opening increases further as the Si expansion continues. Additional surface cracking appears in the edge and corner regions during further Li insertion at 0.05V (Figure 2d). During delithiation, the island contracts both laterally and perpendicular to the plane. Prior work indicates that this leads to tensile stress in the Si. The shrinkage of the underlying Si in the shear lag zone should relax the tensile stress that was induced in the SEI during lithiation, but its final stress state after delithiation is unclear. After delithiation, the AFM results show that edge and corners of the island are also taller compared to the center of island, in spite of the fact that less out-of-plane expansion occurs in the shear lag zone. This combination of results suggests that the SEI film is thicker at the edge and corners compared to the center of the island. With the large volume expansion that occurs in Si, an estimate of the state of charge can be obtained from the ratio of the electrode thickness at given state of charge, to the unlithiated thickness. Both of these can be obtained from the AFM measurements (Figure 2a) after the first cycle, by subtracting the measured SEI thickness from the full measured height. For example, after the 0.05V holds the measured volume expansion is ~256%. With partial molar volume of Li as 9 cm3/mole, this corresponds to an approximate capacity of ~2100 mAh/g of Si (Li2.20Si). This


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is consistent with standard electrochemical measurements at this potential, and thus verifies that the lithiated Si equilibrates with the applied potential. The SEI thickness can be obtained from the observation of irreversible height change after 1st cycle delithiation. This height difference comes from a combined effect of SEI formation on Si and irreversible Si expansion. Tokranov et al.8 reported that irreversible volume expansion of thin film a-Si structure is ~15%. By taking this into account, the resulting SEI thickness after 1st cycle is estimated to be ~270nm which is on the high end of reported values. Here, it is important to note that ex situ studies may underreport the thickness because of changes that occur when the material is removed from liquid electrolyte, and also that significant variations in the SEI thickness can occur with differences in the electrolyte chemistry and the forming conditions. For example, our initial AFM work8 demonstrates that cycling slower through higher potentials leads to thicker SEI. Yoon et al. also recently used in-situ AFM to measure the SEI thickness evolution on Si electrodes and reported values of 10-20nm with a different electrolyte composition. However, Cresce et al.10 used in-situ AFM to study SEI formation on graphitic surfaces and reported that the upper SEI thickness (mostly organic layer) lies in the range of 10 - 480 nm for electrolytes containing no additives. To obtain a more precise understanding of the surface SEI nanocracks in the edge and corner regions, the scan size was reduced to 7um X 7um during the 3rd cycle. This resulted in significantly better resolution (~14nm). Figure 3 shows evolution of SEI layer cracks during 3rd cycle at the edge (3a) and corner (3b) of patterned island. At the onset of lithiation the edge of the island starts to slide laterally, resulting in tensile stress in the SEI layer. This leads to opening of the existing cracks. From the succession of height profiles during lithiation at 0.2V and 0.05V in Figure 3a, it can be seen that these cracks become wider and deeper with increasing


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SOC. An important observation here is that these nanocracks do not fill up with new decomposition products (SEI) during lithiation, an effect which has been proposed in previous speculative discussions of SEI fracture. While one might expect new SEI to grow inside of cracks, it appears that any additional SEI growth is not sufficient to fill in the large cracks observed in our experiments. Two possible explanations for this, depicted in Figure 4, are based on a bilayer SEI structure that has been detected with a variety of methods. One possibility here is that the crack only runs through the top organic layer and doesn’t reach the interface of SEI and Si to cause further SEI formation. Another possible explanation is that the cracks run all the way to the SEI/Si interface, but that filling with much denser inorganic compounds only occurs at the bottom of the crack. The inherent limitations of the AFM imaging make it difficult to clarify behavior at the bottom of these cracks. However, the absence of crack closure that was observed on the edges and corners of islands in a large number of locations is reliable and reproducible. A quick comparison of the tip shape and crack profile shows that the tip angle is much smaller than the angle of the crack. This difference suggests that the tip profile is not influencing the contour in these topographic images (at least in the top portion of the cracks). It was also observed that the same surface cracks which appeared during the 1st cycle reopen and then close again during subsequent cycling. Figure 5 shows AFM topographs comparing crack opening and closure from the 1st to the 3rd cycles. This suggests that the surface cracks don’t completely close during delithiation although the underlying Si retracts back laterally. Also, Fig. 5 shows that the cracks become wider and deeper with cycling and debonding is also observed at the crack fronts during 3rd cycle. This indicates that once cracks appear in the SEI film, continued cycling leads to debonding which could further result in delamination and


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spallation of SEI film. The delamination and spallation of SEI film requires further experimentation and is planned for future publication. During delithiation, the crack openings became narrower at 0.4V. With further Li removal, a large buildup of material at the edge cracks was also observed. This buildup was not observed during the 1st and 2nd cycles, which suggest that these cracks in the SEI layer evolve dynamically with continued cycling. Also, the SEI cracks at the corner of the island (Figure 3b) are much more severe both in terms of crack opening and crack depth compared to the edge region. These differences in cracking behavior at edge and corners of islands can be explained by the difference in their deformation modes. At the corners, the Si is completely free to move parallel to the substrate resulting in true biaxial stress state along with out-of-plane expansion which resembles the behavior of Si nanoparticles during cycling (3D). However, the edge only goes through 2D mode of deformation. This difference in deformation modes at edge and corners leads to very different stress states and hence different crack formation and evolution. In addition to providing direct in situ information about height changes and fracture, the AFM results also provide a direct measurement of the SEI surface roughness and how it evolves during cycling. Figure 6 plots the root mean squared (RMS) roughness at the island edge and center. These values were obtained from 10 µm X 1 µm regions in both the center and edge of the island. At the island center the RMS roughness increases during initial SEI formation and then reaches a fixed value which is consistent with growth of the SEI thickness at higher voltages (0.6V hold). However, at the edge the RMS roughness values keep increasing during the lithiation process. This is consistent with the continuing SEI growth which is also observed in the shear lag region. These trends are also observed during the 2nd cycle, where the roughness in the center remains almost constant but increases in the edge region.


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The results obtained with the patterned islands demonstrate that this approach provides a welldefined configuration where SEI failure mechanisms can be investigated under controlled conditions. It is important to consider these results in the context of prior research on SEI films, which has largely focused on chemical characterization.

During initial battery cycling the

electrolyte undergoes reduction at the negatively polarized electrode surface. This generates a passivating layer consisting of inorganic and organic electrolyte decomposition products.31,32 For optimal performance this layer should be mechanically stable, allow fast Li+ ion transport, and be electronically insulating to prevent further electrolyte decomposition.33,34 These passivation films are expected to undergo substantial deformations when the underlying electrode particles expand and contract during electrochemical cycling. As noted before, this has led to considerable speculation about failure mechanisms in SEI films. However, without direct observations of SEI failure processes, it is difficult to apply meaningful analysis to the design of improved passivation films. Against this backdrop, the experiments described above provide the first direct observations of SEI fracture. Prior modeling of the mechanical behavior of SEI is limited, but this provides a starting point for evaluating our results. Most previous research on mechanical effects during cyclic aging in battery electrodes is based on diffusion induced stresses in the active material, leading to particle fracture. Woodford et al.35 developed failure criterion for individual particles of LixMn2O4 and showed that there are critical combinations of C-rate and particle size below which the material is fracture-safe. Cheng and Verbrugge36 developed analytic solutions for the evolution of concentration and stresses within insertion electrodes and used it to construct fracture initiation criteria. A recent paper by Takahashi and Srinivasan37 suggests that cracking of graphitic particles is unlikely to occur during typical vehicle operations. Particle cracking may


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be more prevalent in other materials. For example, in high capacity electrode materials like Si, there have been significant efforts to design fracture resistant nanostructures (e.g. nanoparticles, nanowires, yolk-shell structures, etc.16,38,39). Wang et al.40 carried out nanoindentation tests on lithiated a-Si structures and concluded that the fracture toughness increases with the lithium-tosilicon molar ratio. In contrast, prior investigation of SEI failure modes has been limited to hypothetical studies, because of the lack of direct experimental observations. Deshpande et al.41 developed a model to capture new Li loss due to SEI formation on freshly exposed surfaces due to fracture of graphite particles, but did not include the effects of SEI fracture. Laresgoiti et al.42 considered SEI damage as the main aging mechanism and implemented a single particle fatigue approach to calculate resulting stresses due to intercalation in both active material and SEI. They showed that qualitatively there was good agreement between the stress amplitudes in SEI and the experimentally observed capacity loss. However, in these studies, specific failure mechanisms in the SEI were not addressed. Tokranov et al.8,43 carried out a simple treatment of electrode and SEI as a core-shell system and calculated an upper bound for stress in the SEI layer. These authors also developed fracture criteria for different failure mechanisms of SEI layer such as through-thickness cracking and interfacial debonding. Verbrugge et al.44 treated the electrode and SEI layer as core-shell structure and showed how different SEI materials e.g. purely inorganic (Li2CO3) vs. purely organic (polypropylene) would affect the stability of the SEI interfaces. In general, these modeling efforts have emphasized that there is a severe lack of information about the key material properties that are needed for quantitative mathematical modeling of the mechanical behavior of SEI materials. Thus, although these initial models provide important insight they cannot provide detailed information about specific failure modes because of the lack of direct observations.


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The method that we report here for directly observing nanocracking and ultimately other failure mechanisms in SEI layers provides crucial information that is needed for the types of modeling efforts described above. For example, the strain at which the SEI layer starts to fail can be obtained from the AFM results. A full analysis of the strain state in the shear lag zone will require finite element simulations of the experiments. However, an estimate of this strain during cycling can be obtained directly from a set of successive AFM scans. This is shown in Figure 4, where specific surface features on top of the SEI were tracked in subsequent images (i.e., the position of these readily identifiable features in the AFM surface profiles were monitored as the island lithiates. These features move both laterally and in the out-of-plane direction). By tracking this motion, a strain map was generated for the SEI layer. Figure 7 plots these displacements during 3rd cycle lithiation at the beginning (1.5V), at 0.2V and at 0.05V (fully lithiated state). As expected, the lateral in-plane displacements are minimal in the center of island, resulting in very low in-plane strain. These displacements show a sharp increase at the inner edge of the shear lag region. The local displacement change (i.e., the strain) is maximum at the inner edge of the shear lag zone, which validates the observation that most of the cracks are concentrated here. The prevalent cracking at the inner edge could also be influenced by the complexity of the Si deformation in the shear lag region. Based on the displacement measurements, the strain at the top of the SEI can be obtained with = !"⁄!# , where " and # are the displacement and position of the feature, respectively. At 0.2V, this gives a strain of ~15% at the inner edge of the shear lag zone. At the fully lithiated state (0.05V), this increases to ~32%. Full analysis of strains and cracking in the SEI are expected to provide a wealth of information about failure mechanisms. As noted above, this will require more detailed mechanics analysis. A simplified version of this approach is presented here, to estimate the fracture energy of these SEI


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films. This is a critical property that has not been previously obtained for any type of SEI film. By obtaining strains from the series of AFM images (i.e., Figure 7), we can estimate the energy release rates associated with fracture. A through-thickness crack first forms in the SEI film at 0.2V during lithiation. For the simple case where both SEI and Si underneath are assumed to be fully elastic, the critical thickness ℎ for film cracking is45

(ℎ )%& =

. 2 Γ ()* + / +

where, Γ is the fracture energy of the SEI, / is stress in the film, -0 = - ⁄(1 − 2 + ) is the planestrain elastic modulus of the surface film and )* = 1.1215 is a numerical coefficient. Although, this equation applies only when the elastic moduli of the film and the substrate are the same, it is used here to provide a rough estimate. Measured elastic moduli for SEI films are on the order of ~1 GPa. Using this value along with a Poisson ratio of 0.3, the measured strain of 15% at 0.2V, and (ℎ )%& as the full SEI thickness (~270nm), the fracture energy of the SEI layer is estimated to be ~13 J/m2. Clearly there is significant uncertainly in this approximation, however, this estimate shows how direct observations of cracking can be used to determine fracture properties. Additional refinements will allow us to further improve the accuracy of these measurements. In the last two decades extensive research has been devoted to identifying the structure and chemical composition of SEI layers, however, the mechanical aspects of SEI degradation are not well understood. The current work demonstrates a novel approach for inducing controlled strains in SEI films. This allowed us to directly show how the volume expansion and contraction of Si electrodes leads to significant mechanical damage of the SEI layer. During lithiation Si


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expansion puts the SEI in tension, which leads to through thickness nanocracking in the surface layer. The key direct observations that were made during this initial study are: 1. SEI formation in the shear lag zone is significantly different than the island center, proving that lateral Si expansion significantly alters the SEI. 2.

Extensive cracking occurs during lithiation, where the SEI film in the shear lag region is subjected to tensile strains.

3. The surface nanocracks which form during 1st cycle lithiation do not fill in at low potential. These cracks then reopen and then close again with further cycling. This indicates that these cracks don’t fill up with new decomposition products, and that they remain open in the form of nano-channel cracks after delithiation (although the crack width decreases due to contraction of the SEI film during delithiation). 4. By measuring strains with a well-defined configuration, it was possible to estimate the fracture energy of the SEI.

In summary, this investigation enhances our mechanical understanding of these surface passivation films. Insights from these experiments are expected to contribute to ongoing work on improving the failure resistance of SEI. These results will also provide relevant information for chemical-mechanical degradation models that are used to predict capacity fade in Li-ion batteries.

EXPERIMENTAL METHODS The patterned Si island samples for AFM investigations were prepared on 500 µm thick quartz wafers (double-side polished, 40 mm X 40 mm in size). A bonding layer of 20 nm thick Ti and


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200 nm thick Cu current collector was deposited by electron beam evaporation, at a rate of 1 Å/s for both metals. Silicon patterned electrodes were fabricated by lift-off process through a standard lithographic process. The S1813 (Shipley Co.) photo resist was spin coated on the asprepared Ti/Cu multilayer electrode and pre-baked at 115oC for 2 minutes. The exposure process was carried out by an ultraviolet mask aligner system (Karl Suss MA6, Germany). The samples were then developed in 1:1 mixture of MF312 and DI water for 1 minute. Si thin film was deposited by e-beam evaporation of pure Si pieces (P-type, 99.999%) at a rate of 1.5 Å/s. After deposition the remaining photoresist was removed by dipping the samples in acetone using slow agitation ultrasonication. The in situ measurements were conducted with a Dimension ICON Electrochemical AFM setup inside an argon-filled glovebox (Nano Surfaces Division, Bruker), where both H2O and O2 were below 1 ppm. The unique PeakForce tapping mode was used with MLCT tips (Bruker AFM Probes), composed of a silicon nitride cantilever with a sharp silicon nitride tip (spring constant: 0.6 N/m, nominal tip radius: 20nm).

The electrolyte was a mixture of ethylene

carbonate (EC) and diethyl carbonate (DEC) (1:1 vol. ratio with 1 M LiPF6). The samples were cycled against Li metal foil, in an in-house electrochemical cell designed for lithium-ion battery materials, and sealed during AFM operation. Based on our experience with these in situ cells (previously reported in Tokranov et al.8,43), the current does not always provide an accurate measure of the state of charge. For this reason, cycling was conducted with a sequence of potentiostatic holds. The cell was held at each potential until the current reached an asymptotic value which was less than 10% of the value at the start of the hold. Continuous AFM scans during cycling show that the center of the islands lithiated/delithiated very uniformly, after the first cycle (i.e., after most of the SEI forms). Similar behavior occurred


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in our previous work using similar islands prepared with a slightly different procedure, where lateral motion of the edges was prevented.8 Because the variation of electrode height with time at the center and edge of these islands was very similar, we concluded that there were no edge effects due to nonuniform current distribution in these islands. Based on this we believe that the current distribution was also uniform across the island during the present investigation. However, here it should be noted that comparisons between the center and edge regions are more complex because of the lateral expansion/contraction of the edges and the resulting differences in the SEI thickness.


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Figure 1. Schematic representation of experimental configuration and shear lag model showing a lithiated Si island and free body diagrams depicting force balance and stress state in the island. During lithiation, edges of the island move laterally resulting in shear stress ( ) at LixSi/Cu interface, however center goes through only out-of-plane expansion resulting in true biaxial stress state.


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Figure 2. a) Height and b) width evolution of patterned a-Si electrode during cycling. Because of shear lag effect; the edges and corners of the island show smaller height change compared to the center of the island. 2D AFM topographs of patterned a-Si electrode during 1st cycle at: (c) 1.5V, (d) 0.05V, (e) 1.5V; and during 2nd cycle at (f) 0.05V and (g) 1.5V. The contrast between center and edge regions of island is due to height differences. The surface topographs at the fully lithiated states (d & f) clearly show cracking of the SEI layer in the shear lag region.


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Figure 3. Crack evolution at a) the edge and b) corner of patterned a-Si electrode during third cycle. Black arrows point out the position of crack in the SEI layer. Note that the SEI cracks remain open during the potentiostatic hold at 0.05 V (i.e., additional SEI formation here is not sufficient to fill the cracks). Solid contour lines represent the height profile along the dashed line section in each AFM topograph.


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Figure 4. Schematic showing deformation of patterned Si island during cycling, and the resulting impact of volume changes on SEI formation and failure.

Figure 5. AFM images showing opening and closing of SEI cracks which formed during: (a-c) 1st cycle; (c-e) 2nd cycle; and (e-g) 3rd cycle.


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Figure 6. Roughness evolution at the edge and center of a patterned Si island.

Figure 7. Schematic representation of surface displacements (c) calculated from a succession of AFM scans (a) during 3rd cycle lithiation. The arrows in (b) represent the position of interest for calculation of local strain. The dashed box in (c) shows the displacement in the shear lag region.


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AUTHOR INFORMATION Corresponding Author * [email protected] Notes The authors declare no competing financial interest. ACKNOWLEDGMENT This work was supported by GM-Brown Collaborative Research grant and by the Assistant Secretary for Energy Efficiency and Renewable Energy, Office of Vehicle Technologies of the US Department of Energy under Contract No. DE-AC02- 05CH11231, Subcontract No 7056410 under the Batteries for Advanced Transportation Technologies (BATT) Program. R.K. and B.W.S. also acknowledge Bruker Corp. for providing access to in situ AFM facilities.

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