In Situ Observation of Epitaxial LiâSi ... - ACS Publications

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In Situ Observation of Epitaxial Li−Si-Nanostructure Formation on Si(111) F. Grosse,* A. Proessdorf, M. Hanke, and O. Bierwagen Paul-Drude-Institut für Festkörperelektronik, Hausuogteiplatz 5−7, 10117 Berlin, Germany ABSTRACT: The formation kinetics of lithium−silicon phases was studied in a molecular beam epitaxy environment under ultrahigh vacuum (UHV) conditions. A heated Si(111) substrate was exposed to an incoming lithium flux at a gradually decreasing substrate temperature. The onset of the nucleation and the formation of Li−Si phases were simultaneously observed by in situ synchrotron X-ray diffraction (XRD), azimuthal reflection high-energy electron diffraction (ARHEED), and line-of-sight quadrupole mass spectrometry (QMS) of the desorbing Li. First, Li-induced reconstructions formed, followed by the appearance of nanostructures which grow epitaxially with respect to the Si substrate. Both crystal structures, namely Li12Si7 and Li21Si5, which are thermodynamical stable at elevated temperatures were identified. They form with an epitaxial relationship such that [111]Si ∥ [010]Li12Si7 and [112̅]Si ∥ [100]Li12Si7 as well as [111]Si ∥ [111]Li21Si5 and [011]̅ Si ∥ [011]̅ Li21Si5. Finally, three-dimensional Li structures appeared on the surface. The in situ measurements were supplemented by ex situ atomic force microscopy and scanning electron microscopy showing the real-space information on three-dimensional Li−Si structures. Similarities and differences compared to a Li-ion battery structure containing a Si anode are discussed. A significant mobility of Si and Li to form nm-sized crystal structures is concluded.

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INTRODUCTION The combination of only two chemical elements in a solid seems at a first glance scientifically not too challenging. One or more crystal structures, e.g., varying in stoichiometry, could be formed in a chemical reaction of the two elements. A further possible scenario is the intermixing of both elements to form an alloy. Thermodynamic data are available already for a large variety of binary element systems from which phase diagrams can be constructed, i.e., the thermodynamically stable structures at a given concentration and temperature are well-known. But even for a simple two element system, growth scenarios can be very complex;1 various factors like kinetic properties of the constituents, strain, surface properties, including reconstructions, and many more might play a decisive role. Epitaxial layers are grown on crystalline substrates which, to achieve flat interfaces, should remain unaffected in the process, with the possible exception of the topmost atomic layer. In molecular beam epitaxy (MBE), the material for the growing layer is conventionally provided by incoming atomic or molecular fluxes which then form a crystalline layer on the substrate. Silicon plays an important and even growing role in modern technology. It is the central chemical element in complementary metal-oxide−semiconductor (CMOS) technology. Recently, Si has also attracted attention due to its potential for application as anode in lithium-ion based batteries.2 A major challenge here is the up to 3-fold increase of the volume upon lithiation.3,4 Therefore, the basic mechanisms of the formation of a Li−Si anode in a battery are of great importance for understanding failure mechanisms and performance issues. An © XXXX American Chemical Society

unresolved issue is the crystallographic relationship of the crystalline Li−Si phases during the lithiation process to the Si anode. Submonolayer deposition of Li on Si leads to surface reconstructions. On the (111) surface a honeycomb chain channel reconstruction Li/Si(111)-(3 × 1) has been reported.5,6 This surface exhibits a true Si double bond.7 Desorption of lithium from the Si(111) surface has been investigated by thermal programmed desorption (TPD) spectroscopy:8 Three Li desorption channels from the Si surface have been identified. A submonolayer zero order desorption with an activation energy of 2.6 eV and a first order desorption with 1.66 eV. For larger coverages above 1.4 monolayer the determined activation energy 1.70 eV of the zeroth order desorption is close to the value of bulk Li (1.64 eV). Various Li−Si crystal phases have been reported in literature, which vary in crystal symmetry and stoichiometry.9−12 LiSi, Li12Si7, Li7Si3, Li13Si4, and Li21Si5 phases have been determined by density functional theory (DFT) to be thermodynamically stable.12 Recently, investigations of Li12Si7,13 Li15Si4, Li13Si4, and Li7Si314 phases by solid state nuclear magnetic resonance have been conducted. Possible Li incorporation mechanisms on the (001) and (111) Si surface were studied by DFT calculations.15,16 The main conclusion is that a Li atom has Received: September 25, 2013 Revised: June 27, 2014

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dx.doi.org/10.1021/jp5025735 | J. Phys. Chem. C XXXX, XXX, XXX−XXX

The Journal of Physical Chemistry C

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desorbing from the substrate surface was measured by a Hiden HAL3 QMS, allowing a mass range of 0−300 μ. The electron energy of the ionizer was taken to be 40 eV at an emission filament current of 500 μA. The substrate was heated by a resistive wire heater at a few mm distance. The substrate temperature was measured by a noncontact thermocouple between heater and substrate.

to overcome a relatively high energetic barrier to be incorporated even into the topmost atomic Si layers. This contradicts in part the intuitive picture that mainly the Li kinetics determines the Li−Si phase formation together only with a local rearrangement of the Si atoms. The kinetics of crystalline Li−Si phase formation is not understood although essential for the usage of Si as battery anode material. Specifically, it is unclear if a significant Si mass transport is necessary to form the Li−Si crystal phases. Either a Li mass transport into and inside the Si crystal is sufficient, or both atoms, Li as well as Si, are mobile and contribute to the crystal formation. To shed light on the basic mechanisms involved in the formation of Li−Si crystal phases, we performed a growth experiment by exposing a Si(111) substrate to an Li flux and simultaneously monitoring the structure by in situ X-ray diffraction (XRD), azimuthal reflection high energy electron diffraction (ARHEED), and measuring the Li desorption and incorporation by line-of-sight quadrupole mass spectrometry (QMS). After giving the experimental details, we discuss the various stages of the Li−Si phase formation starting with the in situ observations. The observed surface symmetries are analyzed by ARHEED. In-plane and out-of-plane crystal structure determination was done by XRD during the growth as well as in more detail post growth. The evolution of the Li− Si layer and its characterization by scanning electron microscopy (SEM) and atomic force microscopy (AFM) is described afterward. Special focus is given to the formation of the Li12Si7 and Li21Si5 phases and their epitaxial relationship to the underlying Si(111) substrate. Although the experiments are carried out at elevated temperatures under ultrahigh vacuum (UHV) conditions, the elementary processes observed are similar to the ones during lithiation of Si anodes in advanced Li-ion batteries.

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RESULTS An initial pristine Si(111) surface was prepared.23 It exhibited a (7 × 7) reconstruction, which indicates that it was free of Li atoms. The substrate was then kept at 750 °C for 10 min. The following opening of the Li shutter turned the initial (7 × 7) immediately into a (3 × 1) pattern. This pattern corresponds to the well-known honeycomb chain-channel Li stabilized Si(111) reconstruction7 with a Li coverage of 1/3 ML determined by TPD.8 The substrate temperature was then lowered by 0.1°s−1 down to 150 °C with the Li shutter open. At 150 °C, the Li shutter was closed again. During the whole experiment, the Li flux desorbing from the substrate was monitored by line-ofsight QMS. Only atomic Li was observed, whereas both isotopes (6 and 7u) were monitored, as shown in Figure 1.

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EXPERIMENTAL DETAILS The experiments were carried out in an MBE chamber which is connected to the beamline U125/2-KMC at the BESSYII synchrotron (Helmholtz Center for Materials and Energy, Berlin).17 Grazing incidence X-ray diffraction (GID) with an incidence angle of 0.2° is enabled by a Huber (4 + 2) six-circle diffractometer.18−20 A photon energy of 12 keV with a resolution of 10−4 is obtained by a double crystal Si(111) monochromator. Reflection high-energy electron diffraction (RHEED) was used with an energy of 20 keV and an emission current of 40 μA. A charge-coupled device (CCD) camera recorded the signal from the phosphorus screen. Further processing was carried out by a digital RHEED image acquisition system. By rotating the sample and simultaneously recording the RHEED images, an azimuthal pattern of the surface was obtained, allowing an unambiguous determination of surface symmetry.21,22 SEM images were acquired by a Hitachi S4800 at an electron energy of 10 keV. The sample distance was about 6 mm. The atomic force microscopy was carried out in the tapping mode. Lithium was evaporated by a hot-lip effusion cell with a tantalum crucible. A dedicated inset was used to focus the Li beam to the substrate and distribute it homogeneously across the sample surface. The cell was operated with a hot-lip temperature of 675 °C and a base temperature of 475 °C. Under this condition the flux corresponded to a beamequivalent pressure of 10−8 mbar, measured by a retractable nude ion gauge in front of the substrate. The lithium flux

Figure 1. Desorbing Li flux measured by QMS during (a) cooling of the Si(111) substrate from 750 °C down to 150 °C under constant impinging Li flux and (b) subsequent heating of the substrate without Li flux back to 750 °C at a rate of 0.1 °Cs−1. Values are given for 6 and 7 representing the stable isotopes of Li.

Except for the small amount of Li responsible for the reconstructions at substrate temperatures above Tsub = 450 °C all other Li atoms were desorbed: The QMS signal was constant within the error margins of the experiment and corresponded to the impinging Li effusion cell flux. A second change of the surface symmetry to a (1 × 1) pattern was B



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observed by RHEED at Tsub = 630 °C, which we relate to a surface reconstruction with 1 ML of Li as observed by TPD.8 Below a substrate temperature of 450 °C, the QMS signal decreased exponentially indicating an increasing amount of Li being incorporated. At this stage, the RHEED pattern changed to an approximate (3 × 3) diffraction pattern with respect to the Si lattice. Simultaneously in-plane XRD profiles along the Si[2̅11] direction were recorded. Three diffraction peaks, besides the Si ones, appeared in the substrate temperature range 400 °C > Tsub > 300 °C as shown in Figure 2. These

Figure 2. In-plane XRD profiles along the Silicon [2̅11] direction during cooling of a Si(111) substrate under constant Li flux. The onsets of diffraction peaks are marked by arrows and the onset temperature is labeled. The observed (ideal24,25) interplanar distance values are Li21Si5(4̅22) d = 3.81(3.82) Å, Li21Si5(6̅60) d = 2.23(2.21) Å, and Li12Si7(400) d = 2.17(2.14) Å.

Figure 3. ARHEED pattern of the surface symmetry obtained by interrupting the cooling and the Li flux at (a) 350 °C and (b) 300 °C. The reciprocal (1 × 1)-Si(111) unit cell is marked by red lines, the (1 × 1)-Li21Si5(111) by blue lines. The corresponding diffraction maxima marked by colored circles. The surface unit cell is chosen with the main axis along the cubic [112̅] for Si(10) and [12̅ 1̅] for Si(01) as indicated.

peaks, and therefore the associated crystalline phases, were present down to room temperature. A full identification of the crystal structure was carried out at the final sample (see below). The identified diffraction peaks are labeled in the figure and further details are mentioned in the figure caption. At 250 °C and below, only a background Li QMS signal was measured, indicating that all impinging Li atoms stick to the Si surface. Finally, when the substrate temperature reached 150 °C the Li shutter was closed. Subsequently, the Li was desorbed from the Si(111) substrate again by ramping, the substrate temperature at the same absolute rate back to 750 °C. Similarities can be seen by comparing to the TPD experiments8 with the desorption observed by reevaporation of the grown layer, as shown in Figure 1(b). The low temperature QMS peak arises from the desorption of the Li overlayer. The broadened double peak in the QMS signal around 600 °C is likely due to the 1 and 0.3 ML coverage of the surface reconstruction in agreement with previous TPD experiments.8 We attribute the broadening of the desorption peaks (compared to low coverage desorption from Si) to the growth induced disorder on the surface. To study the intermediate crystal and surface structure in more detail further samples were prepared by repeating the same recipe as above but interrupting the Li flux at a higher temperature and rapidly cooling down the sample. At a substrate temperature below Tsub = 400 °C, RHEED indicated the onset of a new phase with a changed lattice parameter. Complete surface in-plane information were obtained by ARHEED scans whose maps are shown in Figure 3. The ARHEED map Figure 3(a) taken for the sample for which the growth was stopped at the substrate temperature Tsub = 350 °C

shows a diffraction pattern with an approximately (3 × 3) enlarged surface unit cell which is identified as the (111) surface of the Li21Si5 crystal (see Discussion). Its in-plane surface orientations [11̅0] and [112̅] are aligned with those of the Si substrate. At Tsub = 300 °C (Figure 3(b)), the ARHEED pattern clearly indicates three-dimensional growth with a 6-fold rotational symmetry. The diffraction of the Si(1 × 1) surface unit cell is still visible indicating uncovered area of the substrate. An in-plane XRD map, see Figure 4, and a symmetric crystal truncation rod XRD scan presented in Figure 5 were taken for the sample stopped at Tsub = 150 °C. Both, in-plane and out-ofplane measurements, contain diffraction peaks originating from an additional crystal structure in the sample. The inter planar distance values d are given in the map which can be compared with powder diffraction data. Besides the relatively strong reflections, the map shows diffraction patterns arranged on a concentric ring centered at the origin. The (00L) scan contains distinctive peaks which together with the in-plane information is used to determine the crystal structure of the grown layer. For further analysis, the samples were studied by AFM and SEM at the various intermediate stages. The AFM image in Figure 6 shows 10 nm high islands with in-plane diameters of up to 100 nm for the sample whose growth was interrupted at 350 °C. The in situ ARHEED map of the identical sample is given in Figure 3(a). The platelets are homogeneously C



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(post growth) XRD in-plane map (Figure 4) and out-of-plane (00L) scan (Figure 4) could be utilized to confirm the presence of both phases: The measured d spacings extracted from both agree well with the tabulated values for the Li12Si7 and Li21Si5 crystal structures as indicated by the numerical values in the corresponding figure captions. The alignment of the Li12Si7 crystallites with the Si(111) orientation is determined from Figure 4 to be the [010] direction. The Li12Si7(020) as and (0 10 0) diffraction peaks are observed. Together with the in-plane data the complete epitaxial relationship for the Li12Si7 crystals is determined to be [111]Si ∥ [010]Li12Si7 and [112̅]Si ∥ [100]Li12Si7. The Li21Si5 crystallites have a (111) out-of-plane orientation as indicated by the Li21Si5(555) peak. The intensities of all other diffraction orders are significantly lower than the strongest (555) peak consistent with the structure factors. All observed diffraction peaks can be associated with one of the two phases. Therefore, it is safe to assume that the nanometer sized islands visible in the AFM images in Figure 6 contain either the Li12Si7 or Li21Si5 crystal structure, or both. The in situ XRD scan shown in Figure 2 indicates that first the Li12Si7 crystals form. At lower temperatures additionally the Li21Si5 phase forms. The appearance of the Li21Si5(111) surface symmetry (ARHEED Figure 3(a)) shows that its phase formation sets in at higher temperature and starts from the surface. Its lattice constant (aLi21Si5 = 18.71 Å)25 is approximately 3.5 times as large as that of Si, which explains the enlarged surface unit cell area: The RHEED pattern shows a (1 × 1) surface unit cell of the Li21Si5 crystal. Further information about the epitaxial relationship can be deduced from the XRD inplane map (Figure 4) where the {224}, {448}, and {6 6 12} peaks are aligned with the equivalent {224} peaks of the Si substrate. However, the in situ measurement shows also that for lower temperatures, the crystal grows polycrystalline by inplane rotations as indicated by the appearance of additional weaker Li21Si5(066) peaks aligned with Si(2̅11) direction at 300. The majority of the Li21Si5 crystallites have an epitaxial relationship to the Si(111) substrate which is [111]Si ∥ [111]Li21Si5 and [011̅]Si ∥ [011̅]Li21Si5. A close lattice match between two crystals enables the interface formation with low strain energy for pseudomorphic growth and consequently lower interface defect density. The lattices of materials with dissimilar crystal structures can be combined by matching multiples of their unit cells as closely as possible resulting in an enlarged in-plane interface unit cell. The observed epitaxial relationship is visualized in Figure 9. Single layers of Si, Li12Si7, and Li21Si5 are shown in top view to its inplane orientation scaled by their bulk lattice parameters. Due to the different surface symmetry of the Li12Si7 and Li21Si5 compounds, it is worthwhile to compare the unit cell length along the Si ⟨112⟩ and ⟨110⟩ directions. One finds along the Si [112̅] direction (a) aSi = 6.65 Å, (b) aLi12Si7 = 8.60 Å, (c) aLi21Si5 = 22.91 Å and along the [1̅10] direction (a) aSi = 3.84 Å, (b) aLi12Si7 = 14.32 Å, (c) aLi21Si5 = 13.23 Å. A possible way to minimize the epitaxial strain between Si and the Li21Si5 is the formation of a coincidence lattice along both directions by matching 7 Si to 2 Li21Si5 in-plane unit cells. The resulting mismatch between both cubic crystal structures is only 1.5% which is sufficiently small to favor the observed epitaxial relationship. However, the matching of Li12Si7 to the Si substrate is less optimal. Constructing the lattice by aligning 3

Figure 4. In-plane XRD map of the Li covered Si(111) substrate. The measured (ideal24,25) interplanar distances d correspond to the reflections: (1) Li21Si5(224) d = 3.82(3.82) Å, (2) Li21Si5(066) d = 2.23(2.21) Å/Li12Si7(400) d = 2.17(2.14) Å, (3) Li21Si5(448) d = 1.90(1.91) Å, (4) Li21Si5 (12 0 0) d = 1.55(1.56) Å, (5) Li21Si5(6 6 12) d = 1.26(1.27) Å, and (6) Li21Si5(0 12 12) d = 1.09(1.10) Å.

Figure 5. Out-of-plane XRD of the final Li covered Si(111) substrate obtained at room temperature. The measured (ideal24,25) interplanar distances d correspond to the reflections: Li12Si7(0 2 0) d = 10.08(9.85) Å, Li21Si5(555) d = 2.19(2.16) Å, and Li12Si7(0 10 0) d = 2.02(1.97) Å.

distributed as it can be seen in the lower magnification. The SEM image of the sample whose growth was stopped at 300 °C is shown in Figure 7. Here, additional rod-like structures with a length of up to 5 μm appear with a 120° symmetry in their orientation compared to the sample whose growth was interrupted at 350 °C. The main difference in the corresponding ARHEED maps in Figure 3(a),(b) is the presence of additional lines with hexagonal symmetry indicating a relation to the observed rod-like structures. The SEM image of the fully Li-covered sample is shown in Figure 8.

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DISCUSSION The appearance of an additional XRD diffraction peak (Figure 2) at Tsub = 390 °C and two more between 320 and 300 °C indicates the formation of additional crystal phases. By comparing to all known Li−Si phases, only the data for the Li12Si724 (JCPDS # 04-004-4549) and Li21Si525 (JCPDS #04007-0820) phase agrees sufficiently well. Since the diffraction peaks persisted and no additional ones were found, the final D



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Figure 6. Atomic force microscopy images of Li21Si5/Si(111) nanostructures with a height of about 10 nm. The sample growth was interrupted at T = 350 °C. The ARHEED measurement of the same sample is given in Figure 3(a).

energies (see Chevrier et al.12 Figure 14) is only slightly above of the values of Li7Si3 and Li13Si4 at that temperature, and this difference decreases with increasing temperature. This indicates a destabilization of Li7Si3 and Li13Si4 at higher temperature. Since we find no indication of Li7Si3 and Li13Si4 crystal phases in our experiment, it could indicate that their destabilization temperature is lower than the one predicted by theory. However, we cannot rule out kinetic limitations or epitaxial strain as sources for the suppression of these crystal phases. Both, the Li12Si7 and Li21Si5 phase, are thermodynamically stable structures at elevated temperature, and their appearance is consistent with the DFT results. In addition to crystal structure and epitaxial relation, the experimental results also allow conclusions about the underlying growth kinetics at substrate temperatures between 300 and 390 °C: The initial Si substrates exposed to the Li flux converted initially to Li12Si7 and later to Li21Si5 crystallites which formed on the Si(111) substrate. The necessary Si had to originate from the surface where it reacted with the deposited Li. The nanostructures observed in AFM show that there is no significant Li diffusion into the Si substrate. Details of the phase formaton are determined by the relationship between the Si processes, i.e., removal of Si atoms from, e.g., step edges and their diffusion, and the Li processes, diffusion, and desorption. The chosen conditions lead to epitaxial islands with a height of

Li12Si7 and 4 Si cells along the [112̅] direction leads to a +3% expansion, whereas a 2 to 7 ratio along [110] gives −6.6% compression of the Li12Si7 lattice. Due to the opposite signs along both directions the strain would be partially compensated, which in turn reduces the necessary strain energy in the interface formation. Only the Li12Si7 and Li21Si5 crystal structures are observed. Surprisingly, we did not find any hint of further reported Li−Si stoichiometries. This can be explained by considering that (I) the stability of the Li−Si phases is temperature dependent, (II) the substrate temperature also influences the available amount of Li (the chemical potential) and by that the Li−Si concentration ratio, and (III) the formed phases are not aligned along the XRD orientation. If a sufficiently fast kinetics allows it to reach thermodynamic equilibrium by varying the substrate temperature and by that the chemical potential, then all stable stoichiometries LixSi1−x should be observed. In a recent article by Chevrier et al.,12 the thermodynamic stability for Li−Si phases was predicted by DFT. The Li12Si7 and the Li21Si5 Helmholtz formation energies have only a weak dependence on temperature compared to others. This leads to a destabilization, e.g., of the LiSi phase above a temperature T = 273 °C. At 327 °C (600 K), only Li12Si7, Li7Si3, Li13Si4, and Li21Si5 are stable in the order of increasing Li content. The tieline between the Li12Si7 and Li21Si5 Helmholtz formation E



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Figure 8. SEM images of the Li covered Si(111) surface. Parts (a) and (b) show a bird’s-eye view (45° inclination). Part (c) shows a side cross-sectional view of the layers formed on the Si(111) substrate.

Figure 7. Bird’s-eye view SEM images of the sample obtained by interrupting cooling and Li flux at 300 °C. The corresponding ARHEED pattern is given in Figure 3(b). Two magnifications are shown.

Sb on GaSb(111)A substrates27 exhibiting the same in-plane symmetry. The final room temperature structure, shown in the SEM micrograph Figure 8, still shows a strong anisotropy originating from the crystalline substrate and the formed nanostructures: The elongation of the hillocks is along the ⟨110⟩ direction. The presented investigation of the Li−Si reaction under idealized conditions, i.e., UHV conditions, is discussed in the following in comparison to Si anodes used in Li-ion batteries. Whereas in the presented case there is a chemically clean situation involving only Li and Si, the complexity of a real battery structure is, due to the existence of other elements much higher and less well-defined. Obviously, the comparison can only focus on the situation when Li is in close contact with Si. Besides the chemically more complex situation in a realistic battery structure further differences exist. The resulting Li chemical potential is given in the UHV situation by the applied flux and substrate temperature which determines the amount of desorbing Li. The amount of Li available at the Si anode for the formation of a Li−Si phase in Li-ion batteries is determined by the applied electric field, geometry factors, and diffusion, e.g., through the solid−electrolyte interface (SEI) layer. The UHV experiment results in the formation of thermodynamically stable phases according to the DFT calculations.12 The formation of the metastable Li15Si4 has been observed in nanowire (NW) based Si anodes28 and battery structures by in situ XRD.29 The stoichiometry of the Li15Si4 phase is in between phases observed in the UHV experiment. Therefore, our UHV experiment covers the Li chemical potential range where the Li15Si4 phases should be present. The stabilization of

10 nm and an extension of 100 nm in diameter, much smaller than the μm sized particles in Si anodes used in battery structures.26 A detailed analysis of the atomic processes is out of the scope of the present paper, but the results show that the conversion of the Si substrate is driven by surface processes and not by bulk processes. This includes a significant necessary Si mobility. This conclusion is consistent with DFT results showing that Li has to overcome high energetic barriers especially for the (111) surface of Si to get incorporated.15,16 By further reduction of the substrate temperature, the thermally driven Si and Li processes slowed down. The constant Li flux to the surface lead to an increase of the Li coverage which was not completely used for the Li21Si5 formation and remained at the end as a metallic Li layer on the surface burying the epitaxial structures. The in situ investigation by XRD was essential to show that these epitaxial structures form at higher temperatures and persisted during cooling. Three-dimensional μm sized structures were formed at 300 °C (Figure 7). A single object is visible in Figure 7 (a), a larger scale is shown in (b). The protrusions are elongated preferentially along the equivalent ⟨110⟩ directions, having a 120° rotational symmetry appearing as approximately 5 μm long rods. One-dimensional rods in real space possess a planar two-dimensional diffraction pattern in reciprocal space. The hexagonal lines in the ARHEED diffraction pattern, see Figure 3(b), are cuts of the Ewald sphere through these planes. A similar diffraction pattern is found also for epitaxial growth of F



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at an elevated temperature to Li vapor. Epitaxial nanostructured Li12Si7 and Li21Si5 crystals formed within a temperature range between 425 °C < Tsub < 325 °C. An epitaxial relation of [111]Si ∥ [010]Li12Si7 and [112̅]Si ∥ [100]Li12Si7, as well as [111]Si ∥ [111]Li21Si5 and [011̅]Si ∥ [011̅]Li21Si5 was determined. The observed orientation can be explained by formation of coincidence lattices between Si, Li12Si7, and Li21Si5. The (111) Li21Si5 surface exhibits a approximately nine times larger (1 × 1) surface unit cell than the unreconstructed Si(111) surface. The kinetics of the nanostructure formation is influenced by both elements: The Si release mechanism from the crystalline substrate and its diffusion as well as the Li desorption and diffusion kinetics contribute. The Li−Si crystal phases form by surface diffusion and not Li bulk intercalation. The finally formed Li structures on the surface still preserve aspects of the two-dimensional hexagonal symmetry of the underlying Si(111) substrate. It is expected that similar elementary processes influence the lithiation in Si-based Libatteries.

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS

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REFERENCES

We would like to thank Bernd Jenichen for critically reading the manuscript, and S. Behnke and C. Stemmler for technical support. BESSYII (Helmholtz-Zentrum Berlin) provided the synchrotron beamtime. This work was supported by the Leibniz-Gemeinschaft under P roject No. SAW-2011-PDI-230.

Figure 9. Comparison of (a) Si(111), (b) Li12Si7(010), and (c) Li21Si5(111) layers in top view representing the epitaxial relationship (identical scale and oriented as observed). For each structure, only a single layer is depicted. Si atoms in green, Li atoms in gray. Rectangular surface unit cells are marked in order to explain the epitaxial relationship. For details, see the text.

(1) Croke, E. T.; Grosse, F.; Vajo, J. J.; Gyure, M. F.; Floyd, M.; Smith, D. J. Substitutional C fraction and the influence of C on Si dimer diffusion in Si1-yCy alloys grown on (001) and (118) Si. Appl. Phys. Lett. 2000, 77, 1310−1312. (2) Zhang, W.-J. A review of the electrochemical performance of alloy anodes for lithium-ion batteries. J. Power Sources 2011, 196, 13−24. (3) McDowell, M. T.; Woo Lee, S.; Wang, C.; Cui, Y. The effect of metallic coatings and crystallinity on the volume expansion of silicon during electrochemical lithiation/delithiation. Nano Energy 2012, 1, 401−410. (4) Wu, H.; Cui, Y. Designing nanostructured Si anodes for high energy lithium ion batteries. Nano Today 2012, 7, 414−429. (5) Weitering, H. H.; Shi, X.; Erwin, S. C. Band dispersions of the πbonded-chain reconstruction of Si(111)−Li: A critical evaluation of theory and experiment. Phys. Rev. B 1996, 54, 10585−10592. (6) Olthoff, S. High-temperature scanning tunneling microscopy study of the Li/Si(111) surface. J. Vac. Sci. Technol. B Microelectron. Nanom. Struct. 1996, 14, 1019. (7) Erwin, S. C.; Weitering, H. H. Theory of the Honeycomb ChainChannel Reconstruction of M/Si(111)-(3−1). Phys. Rev. Lett. 1998, 81, 2296−2299. (8) Weindel, C.; Jänsch, H.; Paggel, J.; Veith, R.; Fick, D. Thermal desorption of Li from Si(111). Surf. Sci. 2003, 543, 29−35. (9) Nesper, R. Structure and chemical bonding in zintl-phases containing lithium. Prog. Solid State Chem. 1990, 20, 1−45. (10) Okamoto, H. Li−Si (Lithium−Silicon). J. Phase Equilibria Diffus. 2009, 30, 118−119.

the Li15Si4 phase in a battery structure might be due to the chemical environment. A further important difference is that Si(111) substrates with a high crystalline quality are used in the UHV experiment, whereas for battery anodes nanostructured Si is advantageous. Nanostructured Si exposes different facets of Si for Li incorporation. Despite the significant differences in the environment and the observed crystal phases, there is a common important statement to make. In both situations, under UHV conditions and in real batteries, a significant mobility of both chemical elements, Si and Li, is necessary to form nanoscale crystalline phases. Due to the significantly different Si density in Si and the various Li−Si crystal phases, Si has to be able to diffuse at least along the dimensions of the observed nanocrystals. Therefore, even if there is an intercalation of Li into the Si, which we do not observe in the case of the Si(111) surface, it has to be accompanied by diffusion of Si. Independent of the crystalline phase, it can be concluded that Li−Si crystal structures form by kinetic surface processes involving Li and Si diffusion.

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SUMMARY In our experiment, the formation of a Li−Si compound by lithiation of Si was investigated by exposing a Si(111) substrate G



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