Alternating Silicon and Carbon Multilayer-Structured Anodes

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Alternating Silicon and Carbon Multilayer-structured Anodes Suppress Formation of the c-Li3.75Si Phase Sayed Youssef Sayed, W. Peter Kalisvaart, Brian C Olsen, Erik J. Luber, Hezhen Xie, and Jillian M. Buriak Chem. Mater., Just Accepted Manuscript • Publication Date (Web): 23 May 2019 Downloaded from http://pubs.acs.org on May 23, 2019

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Chemistry of Materials

Alternating Silicon and Carbon Multilayer-structured Anodes Suppress Formation of the c-Li3.75Si Phase

Sayed Youssef Sayed,†,‡* W. Peter Kalisvaart,†,‡ Brian C. Olsen,†,‡ Erik J. Luber,†,‡* Hezhen Xie,†,‡ Jillian M. Buriak,†,‡,* †Department of Chemistry, University of Alberta, 11227 Saskatchewan Drive, Edmonton, Alberta T6G 2G2, Canada ‡National Institute for Nanotechnology, National Research Council Canada, 11421 Saskatchewan Drive, Edmonton, Alberta T6G 2M9, Canada

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ABSTRACT Silicon-based anodes for Li-ion batteries have been gaining a great deal of attention due to their high theoretical gravimetric energy density. Approaches for overcoming the challenge of pulverization associated with Si-based electrodes are required for efficient, reversible, and stable operation of such high energy batteries. This study focuses on addressing the source of pulverization of amorphous silicon films upon cycling, which is typically attributed to the formation of the c-Li3.75Si phase. Cross-sectional samples prepared by focused-ion beam milling revealed a fractured sponge-like silicon structures after 150 cycles at a lithiation cutoff voltage of 5 mVLi, at which the c-Li3.75Si phase forms. Cycling at a higher lithiation cutoff voltage, 50 mVLi, however, resulted in a film with a higher degree of integrity, along with the absence of the c-Li3.75Si phase. These results clearly verify and underscore the deleterious effects of the c-Li3.75Si phase. Alternating carbon and silicon layers results in suppression of the formation of the cLi3.75Si phase, to a degree dependent upon the relative thicknesses of both the silicon and carbon layers. Best results were observed for multilayers of 8 nm Si/4 nm C, with which no evidence for the c-Li3.75Si phase up to 149 cycles was observed. Carbon interlayers were also found to beneficially lower the relative irreversible capacity loss due to SEI formation and associated electrical disconnection.

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INTRODUCTION The growing demand for efficient Li-ion batteries for widespread energy storage applications, including stationary and mobile electronic devices, calls for the development of new electrode materials with a high energy density and long-term stability.1–3 With respect to anodes, graphite, with a theoretical specific capacity of 372 mAh/g, has been used for decades.4 Silicon, however, has a much higher theoretical specific capacity of 3579 mAh/g, corresponding to Li3.75Si.5 Silicon undergoes a very large volume expansion of ~280% with the formation of c-Li3.75Si, a crystalline phase, at lithiation voltages < 50 mVLi compared to 160% for a less Li-rich phase such as Li2Si, which forms at higher voltages.6–8 The main challenge is that repeated (de-)lithiation and associated large changes of volume result in severe fracturing and pulverization of Sibased electrodes.9,10 Cycling of silicon anodes at higher voltages (such as 120 mVLi) has shown improved capacity retention, but results in drastically reduced capacities of 1030%, compared to those observed with the c-Li3.75Si phase.11,12 One approach to overcoming the detrimental pulverization behavior has been to use nanostructured silicon anodes. As observed by Liu et al.,7 silicon particles smaller than 150 nm did not crack during lithiation, which was attributed to these particles being smaller than the critical fracture size of silicon.13 The corresponding critical size for cracking of Si nanowires (NWs), amorphous nanoparticles (NPs), and thin amorphous films, have been reported to be 300 nm,14 870 nm,15 and 100-200 nm,16 respectively. The Li-Si phase transformation during lithiation of Si has been observed using in situ Li NMR analysis, where phase changes from amorphous a-Si to a-Li2Si, a-Li3.5Si, and c-Li3.75Si at ~300-250 mVLi, 120-50 mVLi, and < 50 mVLi, respectively.17 The


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formation of c-Li3.75Si as the fully lithiated phase at room temperature was also elucidated via in situ X-ray diffraction and transmission electron microscope (TEM) analyses.18–20 However, two fundamental questions remain unclear: (i) how severe is the fracturing that results from full lithiation of Si and formation of crystalline-Li3.75Si, formed at voltages < 50 mVLi, compared to amorphous-Li3.5Si and (ii) how can the formation of c-Li3.75Si be minimized or avoided? Some clues can be found in the literature. For instance, Iaboni et al. have shown that the induced compressive stress at the substrate-silicon interface can initially suppress the formation of c-Li3.75Si.21 Moreover, the addition of transition metals to Si was found to create sufficient internal stress, minimizing the formation of the c-Li3.75Si phase by lowering the lithiation voltage as the concentration of metal in the alloy increased.22–25 This stress-voltage coupling has been reported to be ~100-120 mV/GPA from in situ measurements of lithiated Si films.26 For carbon-based silicon electrodes, detailed differential capacity measurements for the first two cycles of Si1-xCx samples prepared by co-sputtering and high-energy mechanical milling showed the absence of the delithiation peak at 430 mVLi, characteristic for the two-phase dealloying of c-Li3.75Si, but, no data was provided for later cycles.27 Jiménez et al. studied magnetron sputtered thin films of Si(70 nm)/C(5, 10, 50 nm)/Si(70 nm), which displayed no delithiation peaks from the c-Li3.75Si phase in the differential capacity plot for the first cycle; however, no evidence was reported for the progression of the appearance of the c-Li3.75Si phase with cycling or any continued effect of the carbon interlayer on suppressing the c-Li3.75Si phase.28 Continuous fracturing of Si during cycling results in ongoing electrolyte decomposition, gradual consumption of the active Li ions that is highly detrimental for


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full-cells, continuous growth of the solid-electrolyte interphase (SEI), depletion of SEIforming additives, and capacity fade.29–31 To overcome these detrimental effects and improve the electronic conductivity within the electrode, surface coating of silicon with carbon via different methods has been studied, including chemical vapour deposition (CVD),32,33 spray pyrolysis,34,35 self-assembly of graphene nanosheets,36 liquidcoating/sintering,37,38 evaporation,39 and sputtering.40 Modification of silicon surfaces with inorganic materials has been also reported to improve the performance of Si-based anodes, including Al2O3,41,42 TiO2,43 LIPON,44 Ti,45 Co3O4,46 BaTiOx,47 and B4C.48 In this study, we present a comprehensive approach to investigating the effects of amorphous carbon on the performance of Si-based Li-ion battery anodes. Thin films of silicon and carbon were deposited as co-sputtered or thin sequential multilayers. In particular, our electrochemical measurements focused on investigating the durability and effectiveness of carbon in suppressing the formation of the c-Li3.75Si phase during cycling. Different Si thicknesses with respect to the carbon layers are investigated. This work also studies the effect of different combinations of Si and C on the cycling stability, Coulombic efficiency, irreversible capacity, formation of the c-Li3.75Si phase and polarizability during lithiation and delithiation steps. Moreover, we illustrate a clear correlation between the formation of c-Li3.75Si and pulverization. Therefore, our results could serve as a useful guide to understand and optimize Si and C composite films as anodes for Li-ion batteries. EXPERIMENTAL METHODS Materials and equipment. Fluoroethylene carbonate (FEC) (99%), metallic Li foil (99.9%), and 1.0 M Lithium hexafluorophosphate (LiPF6) solution in ethylene


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carbonate and diethyl carbonate (EC/DEC), 1:1 v/v% battery-grade, were purchased from Sigma-Aldrich. Copper foil was purchased from McMaster Carr. Stainless steel spacers, springs, and caps were purchased from MTI Corporation. Polypropylene–polyethylene– polypropylene separators with a porosity of 39% (Celgard™ 2325) were soaked in 90 μL of electrolyte solution. Sputtering targets were purchased from Plasmaterials, Inc. Sputtering was performed using an ATC Orion 8, AJA International Inc. instrument. Argon gas with a 5 N purity supplied by Praxair Canada Inc. was used during sputtering. Sample weights were determined after deposition using a Mettler Toledo XP6U balance which has a repeatability of 0.15 µg. Electrochemical measurements were performed using an Arbin BT2000 battery testing system. Electrode preparation and battery assembly. Copper foils of 15 mm diameter were cleaned by sequential sonication for 10 min in each of acetone and isopropanol before sputtering. Prior to sputtering of Si films, 10 nm of Ni was deposited on Cu foil (1.9 cm2) for better adhesion of the Si film during cycling.21,49 Sputtered films of 100 nm thickness were then deposited by either co-sputtering Si and C or sputtering consecutive layers of both elements. Gravimetric capacities were calculated based on the average weight of the 100 nm films. The average weight for each sample and the corresponding standard deviation are listed in Table S1. The average weight varies between 35 and 96 μg depending on the composition and thickness of each element. Silicon (n-type) was deposited with DC magnetron sputtering, while carbon deposition was performed by radio-frequency magnetron sputtering. Sputtering was carried out with continuous substrate rotation under argon gas at a pressure of 4 m Torr. For multilayer deposition, the sputtering rate is 1.5 nm/min and 0.36 nm/min for silicon and carbon, respectively.


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For the co-sputtered Si80C20 sample, Si and C were co-deposited at rates of 2.46 nm/min and 0.27 nm/min, respectively. Coin cells (2032) were assembled using a Li foil counter electrode, one layer of polypropylene-polyethylene-polypropylene separators, stainless steel spacers, springs, and caps. The electrolyte was 1.0 M LiPF6 in EC/DEC with a volume ratio of 1:1 mixed with FEC (10 % by weight). The battery assembly was done in an argon glovebox with oxygen and moisture levels below 0.2 ppm. In order to assess the uniformity of our film deposition and reproducibility of we prepared and tested three thin film batteries under identical conditions. The capacity plots and Coulombic efficiencies are shown in Figure S1, where it is seen that the capacity curves all agree within ~5% and the Coulombic efficiencies are essentially indistinguishable. We expect that one of the important sources of variance is due to uncertainty in measured film weight (Table S1). Due to the large number of different sample configurations tested in this work and the limited number of channels available, it is assumed that all the data presented in this work has similar amounts of uncertainty as shown in Figure S1. Electrochemical testing. Galvanostatic charge-discharge experiments were performed in the voltage range of 5 mVLi - 1.5 VLi. The first three cycles were run at a lower rate of 200 mA/gelectrode. The batteries were further cycled at a higher rate of 600 mA/gelectrode. The constant voltage (CV) step was performed at the cutoff voltage of each of the lithiation and delithiation constant current (CC) steps (i.e., at 5 mVLi and 1.5 VLi for the lithiation and delithiation steps, respectively). The current limit for the CV steps was set up at 20 mA/g. Capacity retention were measured with respect to the second cycle. The cell voltage will be designated as ‘VLi’ in all graphs to remind the reader that a


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Li metal counter electrode is being used. It should be kept in mind that since a current is flowing through the Li metal electrode, the value of VLi reflects overpotentials on both the working and counter electrode. Moreover, in half-cell measurements, the lithium foil counter electrode acts as an infinite reservoir of Li ions and that this is the reason the active material can have a CE of significantly less than 100%. This would not be the case in full cells. All measurements were performed in a temperature-controlled chamber maintained at 25.0 ± 0.1 °C. Cycled batteries were disassembled in an argon glovebox and soaked in diethyl carbonate to remove residual electrolyte prior to further characterization. Characterization. Scanning electron microscopy (SEM) analysis was carried out with a Zeiss Sigma Field Emission SEM at an accelerating voltage of 5 kV. A ZEISS Orion Helium ion microscope (HIM) equipped with a Ga-focused ion beam (FIB) column was also used to characterize the surface and cross-sectional morphology. FIB etching was performed at a tilt angle of 54°, followed by HIM imaging. X-ray diffraction (XRD) analysis was performed in glancing angle mode (incident angle of 3°) on a Rigaku IV diffractometer with Cu Kα radiation (λ = 1.5406 Å). The amorphous nature of the sputtered Si film was investigated using a Thermo DXR2 dispersive Raman microscope. A green laser with a wavelength 532 nm was used at an excitation power of 1 mW. The spectra were recorded by using a 100X magnification objective and an acquisition time of 10 x 40 s. RESULTS AND DISCUSSION The films used to study the effects of the lithiation cutoff voltage on silicon anodes were prepared by sputtering of Si films, with a nominal thickness of 100 nm. As seen in


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Figure S2, Raman spectroscopy reveals that the as-deposited films have an amorphous microstructure. Three peaks are observed at ~320 cm-1, 380 cm-1, and 477 cm-1, corresponding to the longitudinal acoustic (LA), longitudinal optical (LO), and transverse optical (TO) modes, respectively, for disordered Si.50 The two bands observed at 630 cm1

and 940 cm-1 are due to the second order 2LA and 2TO phonon Raman scattering,

respectively.51,52 Crystalline silicon is characterized by a Raman peak at 520 cm-1, which is not seen in the Raman spectra in Figure S2, revealing the amorphous nature of these sputtered silicon films. Furthermore, XRD patterns (Figure S4) of these films do not show evidence of any crystalline Si peaks. Upon full lithiation of amorphous silicon films, it has been widely found that the c-Li3.75Si phase is formed.5,17 In order to investigate the effects of the c-Li3.75Si phase on the performance and integrity of Si electrodes, we carried out electrochemical cycling at two different lithiation cutoff voltages, 50 mVLi and 5 mVLi, under a constant-current constant-voltage (CCCV) protocol. During lithiation with a 50 mVLi cutoff voltage, the voltage-capacity plot (Figure 1a) reveals two sloping plateaus – a characteristic behavior of a single-phase transformation of amorphous-Si to Li-rich amorphous-Si phases.17 The sloping plateaus in Figure 1a are more clearly identified as broad peaks in the corresponding differential capacity plot (dQ/dV vs. voltage) in Figure 1b. According to the NMR analyses provided by Grey et al.,17 the lithiation peak at 0.24 VLi in the differential capacity plot (Figure 1b) is characteristic for the formation of a-Li2Si, along with the corresponding delithiation peak at 0.49 VLi. The lithiation and delithiation peaks at 0.085 VLi and 0.28 VLi are related to the formation and dealloying of the a-Li3.5Si phase, respectively. On the other hand, Figure 1(c, d) show the voltage-capacity and


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differential capacity plots for a 100 nm Si film cycled using a CCCV protocol with a 5 mVLi lithiation cutoff voltage – a low cutoff voltage sufficient for the formation of cLi3.75Si.53 The formation of the c-Li3.75Si phase is clearly revealed from the characteristic lithiation and delithiation peaks at 20 mVLi and 0.42 VLi in the differential capacity plot (Figure 1d),17 as well as the corresponding plateaus in the voltage-capacity plots (Figure 1c). It is worth mentioning that an onset for the c-Li3.75Si lithiation peak can be seen at the end of the second cycle (insert Figure in Figure 1d, and Figure S3), however, after 20 cycles a significant shift to higher voltages with a complete lithiation peak appears at ~20 mVLi. The peak shift to higher voltages with cycling reveals abatement of the substrate-film induced stress-voltage coupling reported by Iaboni et al.21 and presumably a concomitant interfacial film detachment.


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Figure 1. Voltage vs. capacity (a, c) and differential capacity vs. voltage (b, d) for 100 nm Si films cycled using CCCV cycling protocols with lithiation cutoff voltages of 50 mVLi (a, b) and 5 mVLi (c, d), and a delithiation cutoff voltage of 1.5 VLi. The plateaus at the end of lithiation at 50 mVLi (a) and 5 mVLi (c), and the end of delithiation at 1.5 VLi (a, c) represent the contributed capacity from the constant voltage steps. Cycle number is indicated in each corresponding figure legend. The red arrow in Figure (d) refers to the delithiation peak for the c-Li3.75Si phase.

The capacity retention of Si electrodes cycled at 50 mVLi and 5 mVLi lithiation cutoff voltages, are provided in Figure 2a. The fade in capacity retention of both electrodes is clearly revealed from the decreasing capacity upon cycling Figure 1(a, c), and reduction of the differential capacity peak areas in Figure 1(b, d). However, the


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degree of capacity fade is significantly higher when the formation of c-Li3.75Si occurs. The electrode cycled at a lithiation cutoff voltage of 5 mVLi shows capacity retention of 57% after 150 cycles compared to 69% observed for a lithiation cutoff of 50 mVLi. Capacity fade is usually related to the irreversible capacity loss upon cycling.54– 56

For nanostructured Si-based electrodes used in a half-cell configuration, the origin of

the irreversible capacity has been mainly attributed to the formation of the solid electrolyte interface (SEI) and/or disconnection of the active material due to fracturing and pulverization.29,31,56–59 A rough comparison of irreversible capacities resulting from SEI formation (RICSEI) can be made by calculating the difference between lithiation ()*+,capacity at cycle 𝑛 + 1, 𝑄%&' , and delivered delithiation capacity of the previous cycle,

𝑄%()*+,- , relative to the delithiation capacity at the nth cycle according to Equation 1.31 9

∑RIC234 = 6 %:'

7+,𝑄%&' − 𝑄%()*+,-

𝑄%()*+,-

(1)

It is noted that Equation 1 is only an approximation of the irreversible capacity loss due to SEI formation. Specifically, it assumes that any silicon particle that becomes disconnected during the delithiation step has lost a negligible amount of lithium (charge) before it becomes disconnected. If it is indeed the case hat large fraction of the disconnected silicon particles are fully (or partially) delithiated prior to disconnection, then Equation 1 will provide a lower bound on the relative irreversible capacity due to SEI formation. Lastly, it is assumed that other parasitic reactions such as redox shuttling,60 metal plating/dissolution,61 and gas evolution62 can be ignored. This also includes any capacity loss due to sufficiently large increases in electrode polarization resistance.31


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The irreversible disconnection loss can also be approximated, where disconnection is defined as any active material that becomes electrically isolated from the current collector, from either physical disconnection from the substrate or accumulation of products that insulate the active materials either electronically or ionically (e.g trapping).31,56,59 The relative irreversible capacity due to disconnection (RICDisconn) can be estimated as the difference between delithiation capacity at the nth cycle and the lower delithiation capacity at the next cycle, relative to the delithiation capacity at the end of the nth cycle, according to Equation 2.31 9

∑RIC(+>?@AA = 6 %:'

()*+,𝑄%()*+,- − 𝑄%&'

𝑄%()*+,-

(2)

Equation 2 assumes that the differences in any potential parasitic reactions that occur during subsequent delithiation steps are negligible. The two irreversible capacities are represented versus cycle numbers in Figure 2(b, c). As seen in Figure 2b, limiting the lithiation cutoff voltage to 50 mVLi did not show any difference on the SEI irreversible capacity compared to 5 mVLi until the 75th cycle, at which point the difference in capacity retention remains constant. These results indicate that the formation of the c-Li3.75Si phase does not change electrochemical signature of SEI formation in these elemental Si anodes. However, the irreversible capacity due to disconnection of the active material clearly shows the contributed effect of the lithiation cutoff voltages. Figure 2c shows an increase of the RICDisconn by a factor of 2 for the electrode cycled at a 5 mVLi lithiation cutoff voltage compared to 50 mVLi after 50 cycles. As such, the irreversible capacity loss due to disconnection (Figure 2c)


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can be attributed to expansion/contraction associated with the c-Li3.75Si phase (opposed to electrical disconnection due to differences in SEI formation).

Figure 2. Cyclic performance, including the capacity retention (a) and cumulative relative irreversible capacity for the formation of SEI (b) and disconnection (c) vs. the number of cycles for Si electrodes cycled under CCCV protocol with 5 mVLi and 50 mVLi lithiation cutoff voltages denoted by CCCV(5 mV) and CCCV(50 mV), respectively.


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In order to depict the extent of disconnection as a result of the 5 mVLi lithiation cutoff voltage and formation of the c-Li3.75Si phase, we investigated the morphology and integrity of the Si electrodes after 150 cycles by both plan-view SEM and cross-sectional HIM/FIB micrographs. For the electrode cycled at a 50 mVLi lithiation cutoff voltage, the SEM micrograph in Figure 3a shows a cracked film with spacings in the range of tens of microns. The HIM micrograph in Figure 3c reveals a cross-section with some internal voids in the film. On the other hand, for the electrode cycled at a 5 mVLi lithiation cutoff voltage, the plan-view SEM micrograph in Figure 3b reveals a significant disturbance of the film integrity and existence of fractured islands of different thicknesses and widths. The HIM micrograph in Figure 3d reveals that the islands have a porous sponge-like structure. These significant morphological differences between lithiation cutoffs of 5 and 50 mVLi correlate with the measured electrochemical differences. Specifically, with a 50 mVLi lithiation cutoff, the formation of the c-Li3.75Si phase is not observed in the electrochemical data and the morphology of the cycled films is found to be relatively compact in comparison with films cycled at a 5 mVLi cutoff, which did show the formation of the c-Li3.75Si phase. It is worth considering the implications of our results on the influence of the terminal voltage on anode capacity retention for application of Si in full cells. Commercial cells are also charged using a CCCV protocol. During charging of a full cell, only the voltage difference between the electrodes is known whereas the potential vs. Li/Li+ of either electrode is not. During the constant-current stage, the anode potential has been shown to reach less than 50 mV vs. Li/Li+, even falling below 0 V for short periods of time, using a micro-reference electrode technique in commercial graphite-


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LiCoO2 cells63. The detrimental effects of crossing the 50 mV potential threshold, below which c-Li15Si4 can be formed and ways to avoid them are therefore of high practical importance. Keeping the anode above 50 mV vs. Li/Li+ may be achieved by properly adjusting the mass of the anode with respect to the cathode64, but that necessarily involves sacrificing (i.e. not using) part of the anode capacity. Alternatively, ways to alleviate the effects of c-Li15Si4 formation on capacity retention should be sought, as will be shown in the remainder of the paper.

Figure 3. Plan-view SEM micrographs (a, b) and cross-sectional HIM/FIB micrographs (c, d) of 100 nm Si films after 150 cycles with a CCCV cycling protocol with lithiation cutoff voltages of 50 mVLi (a, c) and 5 mVLi (b, d).

Silicon and carbon nanocomposite films In order to suppress the formation of the c-Li3.75Si phase, we sputtered both silicon and carbon in three different combinations: co-sputtering, insertion of a silicon layer in


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between two carbon layers in a sandwich-like arrangement, and consecutive multilayers of Si and carbon with the first silicon layer in contact with the current collector and carbon as the final top layer, as illustrated in Figure 4.

Figure 4. Schematic illustrating the different battery architectures that were utilized to study pulverization of silicon-based electrodes. All films were deposited using magnetron sputtering. Various thicknesses (d) of Si and C interlayers were sputtered: for multilayers samples, dSi= 16, 8, or 2 nm, and dC= 4, or 2 nm; and for sandwich-like samples, dSi= 240, 160, or 80 nm, and dC= 10 nm.

Figure 5 shows the voltage-capacity and differential capacity plots for 100 nm films of Si80C20, 10 nm C/80 nm Si/10 nm C, and 5 layers of 16 nm Si/4 nm C. All electrodes were run under a CCCV protocol with a 5 mVLi lithiation cutoff voltage. Each electrode showed suppression of the c-Li3.75Si phase in the initial cycles with the appearance of the two lithiation/delithiation peaks, characteristic for the amorphous Lirich Si phases, in the differential capacity plots, Figure 5(d, e, f). The characteristic delithiation peak for the c-Li3.75Si phase appears, however, after 25, 80, and 85 cycles for Si80C20, 10 nm C/80 nm Si/10 nm C, and 16 nm Si/4 nm C, respectively. For both the co-


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sputtered (Figure 5d) and sandwich-like (Figure 5e) samples, doublet delithiation peaks at 0.38 and 0.42 VLi were clearly seen. Increasing the Si thickness to 160 nm and 240 nm in the sandwich-like films (Figure S4), with 10 nm C for each of the bottom and top layers, resulted in the appearance of the delithiation peak for the c-Li3.75Si phase after 38, and 15 cycles, respectively, compared to 80 cycles for 80 nm Si (Figure 5e). Therefore, suppression of the c-Li3.75Si phase in the sandwich-like films is dependent on the thickness of the Si layer.

Figure 5. Voltage vs. capacity (a, b, c) and differential capacity vs. voltage (d, e, f) for 100 nm films of Si80C20 (a, d), 10 nm C/80 nm Si/10 nm C (b, e), and 5 layers of 16 nm Si/4 nm C(top-layer) cycled using a CCCV protocol with 5 mVLi and 1.5 VLi cutoff voltages for lithiation and delithiation, respectively. The plateaus at the end of lithiation at 5 mVLi and delithiation at 1.5 VLi, in (a, b, c), represent the contributed capacity from the constant voltage steps. Cycle number is indicated in each corresponding figure legend. The red arrows in (d, e, f) represent the onset for the delithiation peak from the c-Li3.75Si phase.


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Figure 6 shows plan-view SEM micrographs for 100 nm films of Si80C20, 10 nm C/80 nm Si/10 nm C, and 5 layers of 16 nm Si/4 nm C before and after cycling. All the as-deposited samples are seemingly of the same morphology, as directed by the roughness of the Cu foil. After 150 cycles, Figure 6e, the sandwich-like sample showed a ruptured top layer with a layer of pulverized worm-like features underneath. The cosputtered (Figure 6d) and multilayers (Figure 6f) samples, however, showed films of disconnected islands with a relatively higher degree of pulverization and internal porosity for the co-sputtered sample, as compared to the multilayer sample, as revealed from the inset SEM micrographs in Figure 6(d, f) and Figure S6(a, c), which might be due to the fact that the multilayer electrode showed the formation of the c-Li3.75Si phase after 85 cycles compared to 25 cycles for the co-sputtered one. These results are consistent with the data presented for the pure silicon films (Figure 3) where the formation of the cLi3.75Si phase resulted in a greater degree of porosity.

Figure 6. SEM micrographs for 100 nm films of Si80C20 (a, d), 10 nm C/80 nm Si/10 nm C (b, e), and 5 layers of 16 nm Si/4 nm C(top-layer) (c, f). SEM micrographs (a, b, c) are for as-deposited electrodes and after 150 cycles (d, e, f).


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Recent in situ TEM studies for CVD-grown 2D carbon coated silicon sheets by Park and co-workers revealed three effects for the carbon coating.33 Firstly, the carbon coating acts as a lateral constraint, converting the tensile stress in the Si layer during lithiation into a compressive stress. Such compressive stresses might be responsible for the suppression of the c-Li3.75Si observed here, based on the stress-voltage coupling effect. Secondly, cycling resulted in Si sheets with a ripple morphology upon releasing the internal stress during delithiation. Thirdly, finite element analyses revealed that thin carbon coatings can be fractured by the Si if the internal stresses are sufficiently large.33 These reported results reveal a requirement for a sufficient thickness of the carbon coating, relative to that of Si, for better suppression of cracks. In order to investigate the thickness effect of both Si and C in multilayer films, we reduced the studied thicknesses to half of those in Figure 5(c, f), i.e., 8 nm Si/2 nm C. Figure 7(a, d) shows the voltage-capacity and differential capacity plots for a 100 nm film of 10 layers of 8 nm Si/2 nm C. Interestingly, the delithiation peak for the c-Li3.75Si phase appeared at the 35th cycle (Figure 7d), however, this peak does not appear until after 85 cycles in multilayers of 16 nm Si/4 nm C (Figure 5f). This result might indicate that with a reduced Si thickness (8 nm), the 2 nm C layer was not thick enough to overcome the tensile lithiation stress within the silicon layer and prevent formation of the crystalline Li3.75Si phase. On the other hand, with the 8 nm Si layers, increasing the thickness of the carbon to 4 nm resulted in suppressing the formation of c-Li3.75Si up to 149 cycles, as shown in Figure 7(b, e). The 4 nm C layer, in contact with the 8 nm Si, appears to maintain compressive stress sufficient to retard the formation of c-Li3.75Si at the applied bias, due to stress-voltage coupling. Moreover, and surprisingly, reducing the


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Si thickness to 2 nm with 2 nm C was also sufficient to suppress the c-Li3.75Si phase upon cycling, even up to 400 cycles as revealed from the absence of a delithiation peak for the crystalline phase at 0.43 VLi in Figure 7(d).

Figure 7. Voltage vs. capacity (a, b, c) and differential capacity vs. voltage (d, e, f) for films of 10 layers of 8 nm Si/2 nm C(top-layer) (a, c), 8 layers of 8 nm Si/4 nm C(top-layer) (b, e), and 25 layers of 2 nm Si/2 nm C(top-layer) (c, f) cycled using a CCCV protocol with 5 mVLi and 1.5 VLi cutoff voltages for lithiation and delithiation, respectively. The plateaus at the end of lithiation at 5 mVLi and delithiation at 1.5 VLi, in (a, b, c), represent the contributed capacity from the constant voltage steps. The red arrows in (d, e) represent the onset for the delithiation peak from the c-Li3.75Si. Cycle number is indicated in each corresponding figure legend. From all of the data presented thus far, there appears to be significant variability with respect to the appearance of the c-Li3.75Si phase, ranging from cycle 15 for the 10 nm C/240 nm Si/10 nm C sandwich architecture to cycle 149 for the 8 nm Si/4 nm C multilayers. Shown in Figure 8 is a scatter plot of the cycle number showing when the c-


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Li3.75Si phase first appears versus the total C thickness to total Si thickness ratio - there is a clear positive correlation (Pearson correlation coefficient of 0.95) between these two values. During lithiation, the silicon film must undergo significant volume expansion, but, the film is initially constrained at the interface, which gives rise to compressive stresses in the Si film.26 As shown by Sandu et al.,65 coating of Si nanowires with a conformal layer of nickel results in the buffering of volume expansion of the Si during lithiation via deformation of the coating layer. Moreover, it was shown that the buffering effect increases with the ratio of the shell thickness to the nanowire radius.65 As such, we observe a similar effect in Figure 8 where the larger relative thickness of the C layer(s) may be buffering the volume expansion of the Si layer and suppressing the formation of the c-Li3.75Si phase. It is also worth noting that the observed linear trend in Figure 8 is not universal at all thickness ratios, i.e. the 2 nm Si/2 nm C is not included in this plot, as the c-Li3.75Si phase is never observed, even after 400 cycles. This result suggests that the relationship becomes highly non-linear at higher thickness ratios. Moreover, it should also be cautioned that this relationship breaks down at sufficiently thin C layer thicknesses. As seen in Figure 8, there is one obvious outlier to the linear trend at a thickness ratio of 0.25, which corresponds to the 8 nm Si/2 nm C multilayer stack. Presumably this is due to the change in mechanical properties of C layer at 2 nm, where the carbon is unlikely to even form a continuous film. Lastly, it should be pointed out that the initial film thicknesses used in this study are near or below the so-called critical fracture thickness of silicon,13 and as such, the energetics of c-Li3.75Si nucleation could change at larger Si film thicknesses.


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Figure 8. Scatter plot of the cycle number when the c-Li3.75Si phase first appears versus the C to Si layer thickness ratio for all different Si/C architectures studied in this work. The solid line is a linear fit, with a Pearson correlation coefficient of 0.95.

The morphologies of the cycled electrodes are shown in the SEM micrographs in Figure S7. After 150 cycles, the multilayer electrode composed of 8 nm Si/2 nm C (Figure S7d) showed the existence of interconnected islands. The SEM micrographs for multilayers of 8 nm Si/4 nm C and 2 nm Si/2 nm C electrodes after 150 and 400 cycles, respectively, are shown in Figure S7(e, f). Although in spite of the existence of cracks on the surface of the latter electrodes, both surfaces seemingly maintain the surface morphology of the corresponding as-deposited films. Capacity retention of the 100 nm thick silicon-carbon electrodes is shown in Figure 9a. A capacity retention of 76% is observed after 150 cycles, which is very close to the reported capacity retention 0f 74% by Jiménez et al. for sputtered silicon films.28 Upon cycling, it is clearly seen that the presence of carbon enhances the capacity retention of all the electrodes, compared to elemental Si. Moreover, the multilayer electrode composed of 8 nm Si/4 nm C shows the best capacity retention, 85%, after 150 cycles (Figure 9a). The effect of carbon was also noticeable through minimization of the


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irreversible capacity resulting from both the SEI formation and disconnection (Figure 9(b, c)). In Figure 9b, the co-sputtered Si80C20 and multilayer (8 nm Si/2 nm C) electrodes, revealing the c-Li3.75Si phase after fewer than 40 cycles (vide supra), have irreversible capacities that are quite close to that of the elemental Si electrode, indicating a correlation between the formation of the c-Li3.75Si and the accompanying fracturing and SEI growth.The observed irreversible capacities for SEI for the Si80C20 and multilayer (8 nm Si/2 nm C) electrodes are then followed by those for the two electrodes (10 nm C/80 nm Si/10 nm C, and 16 nm Si/4 nm C), which show formation of the c-Li3.75Si phase between 80 and 85 cycles. Finally, the multilayer electrodes, (8 nm Si/4 nm C) and (2 nm Si/2 nm C), showing no evidence for the c-Li3.75Si phase up to 149 and 400 cycles, respectively, exhibit the lowest relative irreversible capacity for the formation of SEI. This observation may be due to the dual function of the carbon interlayer, which both protects the Si layer and creates an internal stress-voltage coupling effect that is enough to prevent the formation of the c-Li3.75Si. The cumulative relative irreversible capacity for disconnection (Figure 9c) for the (2 nm Si/ 2 nm C) electrode, however, is the second highest after the elemental Si film, which is most likely due to the thin nature of the Si interlayers. Lastly, the 8 nm Si/4 nm C multilayer electrode has the lowest irreversible capacity due to disconnection, which is probably a result of having the highest C to Si thickness ratio (excluding 2 nm Si/2 nm C for the previously discussed reasons of incomplete C layers), making it the most effective at buffering stresses associated with lithiation/delithiation.


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Figure 9. Capacity retention (a), and cumulative relative irreversible capacity for the formation of SEI (b) and disconnection (c) vs. number of cycles.

Commercial Li-ion batteries are usually charged at a constant current until reaching a limiting voltage.66 A constant voltage (CV) step is usually applied after the constant current (CC) step to prevent overcharging of fully lithiated anodes and


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conversion of the charging electrical energy to thermal energy, in the absence of any ionic intercalation or alloying.66,67 Therefore, the CCCV protocol needs to be implemented for safe practical applications of Li-ion batteries.67 The CCCV protocol might also be beneficial to overcome any sources of polarization, particularly with samples with high RICSEI and RICDisconn values, vide infra. For both graphite- and siliconbased electrodes for Li-ion batteries, the CCCV protocol typically results in decreased capacity fade and long-term cycling.11,68 In order to understand the effects of the CV step, we ran our samples under a CCCV protocol for both the lithiation and delithiation steps. Figure 10a illustrates the percentage of the contributed capacity from the constant voltage steps (%QCV) during lithiation, at 5 mVLi, versus the number of cycles for the different aforementioned silicon/carbon electrodes. The %QCV can be calculated according to Equation 3 %𝑄DE =

𝑄DE × 100% (3) 𝑄DD + 𝑄DE

where QCC and QCV are the contributed capacities from the constant current and constant voltage steps, respectively. As seen in Figure 10a, after 15 cycles, %QCV of the pure Si electrode increases continuously up to the 70th cycle, and then becomes constant at ~6%. The rise in %QCV indicates an increasing electrode polarization and resistance, which might be due to electrolyte decomposition and formation of SEI with inherent insulating components.30 The constant value of %QCV after 70 cycles is in agreement with the deviation from linearity in the plot of the RICSEI vs. the number of cycles for the Si electrode in Figure 2b, indicating a lower rate for SEI formation that might be due to cracking and repairing of the existing SEI. Wetjen et al. have observed a similar


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inflection in the slope of the irreversible capacity plot for anodes composed of silicon nanoparticles and graphite.30

Figure 10. (a) Percentage capacity contributed from constant voltage steps for lithiation at 5 mVLi vs. number of cycles. (b) Percentage capacity, at the 36th cycle, vs. number of multilayers (b). All cells were cycled at 600 mA/g. The first three cycles were omitted to exclude the effect of the lower applied rate at these initial cycles.

In the case of silicon/carbon electrodes, all the electrodes have an initial decrease in the %QCV, similar to elemental Si (Figure 10a). The reason for this initial decrease in %QCV is not clear. However, after 15-20 cycles, an earlier stage of cycling compared to


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elemental Si, nearly constant values are observed with no noticeable variation except for the 2 nm Si/2 nm C electrode that shows a decrease from 12% to 10% of contributed capacity after 150 cycles. The employment of the CV step is advantageous for the latter electrode as revealed from its highest %QCV compared to the other electrodes; more than one tenth of the lithiation capacity is achieved from the CV step. The 2 nm Si/2 nm C electrode includes 50 consecutive layers of silicon and carbon, with a total thickness of 100 nm. Each layer has two interfacial boundaries. Although glancing-angle X-ray diffraction does not reveal the presence of silicon-carbides, Figure S3, each interfacial boundary might represent a source of polarization. To interrogate the effect of interfacial boundaries, we plot the total number of multilayers versus the %QCV at the 36th cycles Figure 10b, and at different cycles Figure S8, for each of the multilayer Si/C electrodes. The linear relation between the %QCV and the number of multilayers is clearly revealed from Figure 10b, indicating the importance of applying the CCCV protocol, especially in composite electrodes with a high number of interfacial boundaries. The CCCV protocol is not usually applied to the discharge process, the delithiation of the anodes in full Li-ion batteries, as this process is actually dependent on the user’s needs. However, fundamental studies on the effect of the CCCV protocol on delithiation have shown significant %QCV values (~10% after 100 cycles) for electrodes that encountered degradation of the active materials.30 For all of the carbon-based silicon electrodes in this study, Figure S9, the %QCV values for delithiation are lower than those for lithiation. The overpotential at the end of the CC step during delithiation, with respect to the delithiation voltage for c-Li3.75Si ( ~0.43 V), was > 1 V, which is much higher as compared to that during lithiation and explains the lower %QCV during delithiation steps.


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However, upon cycling, the elemental Si electrode has similar %QCV values for both delithiation and lithiation steps. Moreover, upon cycling, %QCV for delithiation of the elemental Si electrode is higher than the silicon-carbon multilayer anodes, which suggests that carbon interlayers improve the delithiation kinetics of silicon anodes. CONCLUSIONS The pulverization of Si-based electrodes upon cycling is one of the major challenges for implementing these high energy electrodes in Li-ion batteries. Our microscopic and electrochemical results clearly showed that the formation of the c-Li3.75Si phase is highly detrimental for the integrity, porosity of Si-based films, and associated irreversible capacity loss. In this work, we show that a multilayer architecture, alternating thin carbon and silicon interlayers suppresses the formation of the c-Li3.75Si phase to a degree that is linearly related to the thickness ratio of the carbon to silicon layers. The presence of 4 nm carbon in a sequential multilayer arrangement with 8 nm Si resulted in suppressing the formation of the c-Li3.75Si up to 149 cycles and has shown the lowest relative irreversible capacity. In addition, there is a linear relationship between the number of multilayers and contributed capacity from the constant voltage steps. These findings provide a reference for the rational design of silicon-based nanostructured electrodes for high performance Li-ion batteries. Author Information Supporting Information Raman spectrum of as-deposited Si films, voltage profiles and differential capacity plots, SEM micrographs of cycled Si-C films, capacity retention and Coulombic efficiency


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plots, table summary of capacity and Coulombic efficiency, XRD of as deposited films, percentage capacity during constant voltage steps. Corresponding Authors *(J.M.B.) E-mail: [email protected] *(E.J.L) E-mail: [email protected] *(S.Y.S.) E-mail: [email protected] ORCIDs Sayed Youssef Sayed: 0000-0003-1575-676X W. Peter Kalisvaart: 0000-0003-1228-906X Brian C. Olsen: 0000-0001-9758-3641 Erik J. Luber: 0000-0003-1623-0102 Jillian M. Buriak: 0000-0002-9567-4328 Notes The authors declare no competing financial interest. ACKNOWLEDGMENTS This work was supported by grants from the Western Economic Diversification Canada (grant number 000014328), Future Energy Systems (https://futureenergysystems.ca; Grant Nos. T12-P04), and Alberta Innovates Technology Futures (grant number AITF iCORE IC50-T1 G2013000198. HX thanks Alberta Innovates Technology Futures for a Graduate Scholarship. Electron microscopy and X-ray diffraction were carried out at the University of Alberta nanoFAB. We would also like to thank Dr. Shihong Xu from University of Alberta nanoFAB for assistance with the He-ion microscopy. We would


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like to thank Tate Hauger for valuable discussions. The authors would like to thank Dr. Michael D. Fleischauer for the use of his lab facilities.

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Alternating Silicon and Carbon Multilayer-Structured Anodes

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