SiC-Free Carbon-Silicon Alloys Prepared by Delithiation as Lithium

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SiC-Free Carbon-Silicon Alloys Prepared by Delithiation as Lithium-Ion Battery Negative Electrodes Leyi Zhao, J. C. Bennett, A. George, and M. N. Obrovac Chem. Mater., Just Accepted Manuscript • DOI: 10.1021/acs.chemmater.8b03898 • Publication Date (Web): 16 May 2019 Downloaded from http://pubs.acs.org on May 16, 2019

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Chemistry of Materials

SiC-Free Carbon-Silicon Alloys Prepared by Delithiation as Lithium-Ion Battery Negative Electrodes Leyi Zhao,a J.C. Bennettb, A. Georgec and M.N. Obrovac*,a,c aDepartment of Chemistry, Dalhousie University, Halifax, Nova Scotia B3H 4R2 Canada bDepartment of Physics, Acadia University, Wolfville, Nova Scotia B4P 2R6 Canada cDepartment of Physics and Atmospheric Science, Dalhousie University, Halifax, Nova Scotia B3H 4R2 Canada

ABSTRACT: Carbon‐silicon alloys in different stoichiometric ratios are synthesized by delithiation of carbon‐lithium‐silicon ternary alloys with ethanol followed by washing with HCl and distilled water. The as‐prepared carbon‐silicon materials are air and water stable. In contrast to mechanically milled or sputtered C‐Si alloys studied in the past, the method of synthesizing C‐Si alloys introduced in this work avoids the formation of inactive SiC even after two hours high energy ball milling. This results in C‐Si alloys with significantly greater volumetric and specific capacity. When cycled in Li half‐cells, C‐Si alloys exhibit good cycling performance and lower volume expansion compared to conventionally made Si alloys. This is attributed to the presence of void spaces in the structure which can accommodate some of the Si volume expansion.

INTRODUCTION Lithium‐ion batteries (LIB) occupy a primary position in the market of mobile electronics and electrical vehicles. Currently, the most widely used negative electrode material for LIB is graphite, which has low volume expansion (~10%), low average voltage (~150 mV), high capacity (719 Ah L‐1 and 372 mAh g‐1, corresponding to LiC6), good cycling stability, etc. However, the developing market de‐ mands for batteries with higher energy density. One path‐ way towards this goal is to utilize active alloys as negative electrodes. Si alloys are promising candidates because of the low average voltage (~0.4 V) and high capacity (2194 Ah L‐1 and 3579 mAh g‐1, corresponding to Li15Si4) of Si.1,2 As a result, the use of Si alloys in LIB can result in significant improvements in energy density.3,4 However, the commercialization of Si alloys as LIB nega‐ tive electrodes is hindered by the large volume expansion of Si after full lithiation (280% for pure crystalline Si).2 The large volume expansion can lead to electrode degradation and disconnection of active regions, resulting in capacity fade and poor cycling performance. One possible solution is the fabrication of carbon‐silicon (C‐Si) alloys. Ideally, in C‐Si alloys, a C matrix acts as a diluent to the volume expan‐ sion of Si, resulting in less volume expansion overall. It is hoped that C‐Si materials might combine the advantages of carbon (good electronic conductivity and good cycling sta‐ bility) and silicon (high capacity). C‐Si alloys containing micro‐porous and nano‐porous complex structures have been produced by chemical vapor deposition (CVD) and physical vapor deposition (PVD).5‐8 Recently, the introduction of void spaces in carbon matrices has been proposed to accommodate the large volume ex‐ pansion of Si. Such “core‐shell” (Si covered by C) materials

have been shown to have good cycling performance in Li half‐cells.9,10 High energy ball milling (HEBM) has also been widely applied to synthesize C‐Si alloys, including milling silicon powder with graphite11‐14 or other carbon precur‐ sors.15,16 Carbon and silicon are almost insoluble, according to their binary phase diagram.17 However, SiC forms readily when silicon and carbon are combined by HEBM or heated to high temperatures (e.g. 900 °C).11 The formation of inac‐ tive SiC greatly reduces the capacity of C‐Si alloys, since its formation consumes both active Si and active C.11 Another method to accommodate the large volume ex‐ pansion of Si is the demetallization of metal‐silicon interme‐ tallic compounds. Many recent studies have investigated the demagnesiation of Mg2Si to prepare nanostructured Si or porous Si materials.18‐21 Annou et al. prepared Si nano‐ materials by an approach based on the direct solution phase oxidation of a solid state Zintl phase NaSi with alcohols.22 Zeilinger et al. synthesized amorphous silicon (a‐Si) by chemical extraction of the alkali metal component from Li15Si4 and Li3NaSi6 with liquid ammonia and ethanol, re‐ spectively.23 In our previous study, we have shown a con‐ venient way to prepare bulk quantities of layered a‐Si by the delithiation of Li‐Si intermetallic compounds with alcohol.24 When utilized as electrodes in Li half‐cells, these materials exhibited lower volume expansion and superior cycling per‐ formance compared to that of crystalline silicon.20‐24 In the present study, SiC‐free C‐Si alloys are prepared by the delithiation of ball milled C‐Li‐Si alloys. This is the first use of this method to make C‐Si alloys. Using this method, the formation of inactive crystalline/nanocrystalline SiC is avoided, resulting in Si‐C alloys with significantly higher cy‐ cling capacity than Si‐C alloys made by HEBM or heating at high temperature.

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Table 1. Mass of Li12Si7 and MAG‐E graphite for the prep‐ aration of C‐Li‐Si alloys. C:Si

Li12Si7

MAG‐E Graphite

By mole ratio

(g)

(g)

90: 10

0.5433

1.4688

80: 20

0.8503

1.0217

70: 30

1.0476

0.7343

60: 40

1.1851

0.5340

50: 50

1.2865

0.3864

40: 60

1.1851

0.5340

EXPERIMENTAL SECTION Synthesis of C‐Si alloys. Li12Si7 was synthesized in an arc furnace by melting in total ~1.2 g of specific stoichiometric ratios of Li (99.90% Li ribbon, Sigma) and Si (98.40% Si lump, Alfa Aesar) with a 7% excess of Li (to offset Li loss from vaporization) under an Ar‐flow. After cooling, the re‐ sulting Li‐Si hemispherical slug was ground into powder with a mortar and pestle in an Ar‐filled glove box. The Li12Si7 powder was then ball milled with graphite (MAG‐E, Hitachi, average size of 20 μm) in a 65 mL hardened steel vial containing 115 g of 4.76 mm stainless steel balls under an argon atmosphere for two hours using a SPEX high en‐ ergy ball mill (Model 8000‐D, Spex Certiprip, Metuchen, M.J.). Table 1 shows the masses of Li12Si7 and MAG‐E graph‐ ite used for the preparation of different C‐Li‐Si alloys in this study. Bulk quantities of air‐ and water‐stable C‐Si powders were prepared by delithiation of C‐Li‐Si alloys with an ex‐ cess of ethanol, presumably according to the general reac‐ tion: LixSiyCz (s)

xEtOLi (s) + H2 (g) + SiyCz (s).

(1)

The delithiation was carried out under Ar gas by adding ~1.5 g of C‐Li‐Si powder into a round bottom flask affixed with a Vigreux condenser. About 120 mL of ethanol (99.89%, containing 0.10% H2O, Commercial Alcohols) was added to the flask and stirred continuously by a magnetic stirrer. The reaction proceeded overnight under room tem‐ perature. The whole process was conducted in an Ar‐filled glovebox. The lithium alkoxide byproducts of the delithia‐ tion reaction are insoluble in ethanol and were subse‐ quently washed away by transferring the reaction products to a Büchner funnel (with 44.25 mm filter paper, Whatman) and washing with distilled water three times, HCl (1 mol L‐ 1, prepared from 37 wt.% HCl, Aldrich) three times, and then distilled water again until the pH of the washing water was about seven. Any residual Li in the LizSixCy alloys after the delithiation reaction would also be washed away during this process. The air‐ and water‐stable products were then dried in a tube furnace at 120 °C for one hour and at 300 °C for three hours under Ar‐flow. The resulting C‐Si powders were produced with an average yield of 79 wt. %. For comparison, silicon (powder, ‐325 mesh, 99%, Al‐ drich) and carbon (MAG‐E, Hitachi, average size of 20 μm)

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were also ball milled using the same method to combine LI12Si7 and carbon. Material characterization. X‐ray diffraction (XRD) pat‐ terns were collected with a Rigaku Ultima IV diffractometer equipped with a copper target, a dual position graphite dif‐ fracted beam monochromator and a scintillation counter detector. XRD data were collected at a scattering angle (2θ) range of 20° to 80° with a single step of 0.05° and a three second dwell time. Particle size, morphology, and elemental distribution of C‐Si alloys were studied by scanning electron microscopy (SEM) and energy dispersive X‐ray spectros‐ copy (EDS) with a TESCAN MIRA 3 LMU Variable Pressure Schottky Field Emission Scanning Electron Microscope equipped with an INCA X‐max 80mm 2 EDS system. Trans‐ mission electron microscopy (TEM) images, selected area electron diffraction (SAED) patterns and elemental map‐ ping (512 400 pixel; 10 nm probe diameter) were taken using a Philips CM30 electron microscope. Samples for TEM imaging and selected area electron diffraction (SAED) pat‐ terns were prepared by dispersing powders in distilled wa‐ ter (ultrasonic for ~5 min) and placing a drop of the solu‐ tion onto a Lacey carbon film supported on a copper TEM grid. Samples for TEM elemental mappings were prepared by essentially the same method except placing a drop of sample solution onto a carbon coated TEM grid. The XPS instrument used for measurements was a Multilab 2000 (ThermoVG Scientific) using a non‐monochromatic Al source for the incident X‐rays (14kV, 25mA). Experiments were conducted at room temperature under ultra high vac‐ uum conditions (~2.7 x 10‐7 Pa). Sample densities were measured by a gas (helium) displacement pycnometry sys‐ tem (AccuPyc II 1340 Pycnometer, Micrometritics). Electrochemical characterization. Electrode slurries were prepared by mixing active materials (e.g. C‐Si alloys, ball milled C0.5Si0.5), carbon black (Super‐P, Erachem Eu‐ rope) and polyimide (PI, PI‐2555, HD Microsystems) in N‐ methyl pyrrolidinone (NMP, 99.50%, Sigma). The formula‐ tion of all electrode coatings was C‐Si alloy/carbon black/PI solids in a 62.5/18/19.5 volume ratio. For comparison, a blend of nano‐Si (nanopowder, 98%, Aldrich) and graphite (MAG‐E, Hitachi, average size of 20 μm) was also used as the active material (Si: C= 3: 7 by at., hereafter named nano‐Si0.3C0.7). The slurries were prepared with a planetary mill (Retsch PM 200) with four 11 mm tungsten carbide (WC) balls at 100 rpm for one hour. The resulting slurries were coated on Cu foil (Furukawa Electric, Japan) with a 0.1 mm coating bar and then dried at 120 °C in air for one hour. The coating thickness was 0.029‐0.032 mm after drying. The loading of active materials (Si and C) was 1.0‐ 1.2 mg cm‐2. Circular working electrodes (13.27 mm diam‐ eter) were punched from the coated foil and heated in a tube furnace at 300 °C under an Ar‐flow for three hours, in order to cure the PI binder. Ball milled graphite and a‐Si (by HEBM of graphite (MAG‐E, Hitachi) or silicon (325 mesh, 99%, Aldrich), respectively, using the same milling condi‐ tions as described above for the Li12Si7/graphite compo‐ sites) were also used as active materials and made into elec‐ trodes using the same procedure. Electrodes were transferred into an Ar‐filled glove box and assembled in 2325 coin‐type half‐cells, with a lithium foil reference/counter electrode. Two layers of Celgard


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2300 separators, one stainless steel spring and one spacer were used in each coin cell. 1M LiPF6 (BASF) in a solution of monofluoroethylene carbonate (FEC), ethylene car‐ bonate (EC) and diethyl carbonate (DEC) (in a volume ratio of 1:3:6, all from Novolyte Technologies) was used as elec‐ trolyte. Most coin cells were cycled at 30.0 ± 0.1 °C between 0.005 ‐ 0.9 V with a MACCOR Series 4000 Automated Test System at a C/10 rate and trickle discharged to C/20 rate for the first cycle, and at a C/5 rate and trickle discharged to a C/10 rate for the following cycles. Some coin cells were cycled at different C‐rates without trickle discharge. All coin cells had a rest period of 15 minutes at open circuit at the end of each half cycle. C‐rates were calculated by assuming Si has a capacity of 3579 mAh g‐1,1 and ball milled carbon has a capacity of 200 mAh g‐1, respectively. The value of 200 mAh g‐1 was determined by measuring the capacity of a ball milled graphite electrode. To measure the volume expansion of C‐Si electrodes, elec‐ trode thicknesses were measured in an Ar‐filled glove box to within ±1 μm with a Mitutoyo 293‐340 precision microm‐ eter before assembling cells and after one full lithiation (C/20 constant current, C/40 trickle, 5 mV voltage limit). Coating volume was determined by subtracting the thick‐ ness of the Cu current collector from the thickness of the whole electrode, and then multiplying by the coating area. The volume expansion (ξ) was then calculated by:

100%

Figure 1. The C‐Li‐Si ternary system according to Reference 25. Red dots (a‐f) represent the the target compositions of the pre‐ pared C‐Li‐Si precursors and blue dots (g‐l) represent the ex‐ pected final C‐Si sample compositions after delithiation.

(2)

Where υ and υl are the unlithiated and lithiated coating volumes, respectively. RESULTS AND DISCUSSION Structural and Morphological Characterization of C‐Si alloys. Figure 1 shows the target compositions of the pre‐ pared precursors and the expected final sample composi‐ tions after delithiation superimposed on the C‐Li‐Si ternary system according to Reference 25. According to the ternary phase diagram, the C‐Li‐Si precursor target compositions reside in a C‐SiC‐Li13Si4 3‐phase region, while the expected product compositions g‐j reside in a C‐SiC 2‐phase region, sample k has the same composition as SiC, and sample l re‐ sides in a SiC‐Si 2‐phase region. Figure 2a‐f show the XRD patterns of the prepared C‐Li‐Si alloys. There are two broad peaks in each of Figure 2a‐f. To identify these two peaks, Li12Si7 and MAG‐E graphite were ball milled separately using the same conditions as were used to synthesize the C‐Li‐Si alloys. XRD patterns of Li12Si7 and graphite before ball milling are shown in Figure 3a‐b; and after ball milling are shown in Figure 3c‐d, respectively. After milling, both Li12Si7 and graphite XRD patterns have broad peaks at similar positions (~25° and ~45°), indica‐ tive of amorphous phases. It is difficult to deconvolute these peaks in samples containing both phases. The two broad peaks in Figure 2a‐f are consistent with being either a combination of the ball milled Li12Si7 and ball milled graphite peaks (i.e. a mixture at the nm‐scale of two phases) or an amorphous C‐Li‐Si solid solution.

Figure 2. XRD patterns of (a‐f) C‐Li‐Si alloys, and (g‐l) the cor‐ responding C‐Si alloys prepared from their delithiation.

According to previous studies, SiC forms when silicon and carbon are combined by high energy ball milling (HEBM).11,26,27 However, no SiC peaks appear in Figure 2a‐ f. Crystalline Si and graphite powders were milled in a 1:1 mole ratio using the same conditions used to prepare C‐Li‐ Si alloys, in order to see if SiC could be detected in alloys



Figure 3. XRD patterns of (a) Li12Si7 prepared by arc melting, (b) MAG‐E graphite, (c) ball milled Li12Si7, (d) ball milled MAG‐ E graphite, (e) ball milled graphite and Si in a 1:1 mole ratio, and (f) amorphous Si prepared by delithiation of Li12Si7.

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prepared from Si and C without lithium. As shown in Figure 3e, pronounced SiC peaks can be easily observed from the product of milling Si and C, indicative of the formation of a nanocrystalline SiC phase. This implies that SiC forms when Si ball milled directly with C while SiC formation is avoided when Li‐Si compounds are ball milled with C, as shown in Figure 2a‐f. The absence of SiC is also confirmed by other techniques, as discussed below. Figure 2g‐l show XRD patterns of C‐Si alloys with differ‐ ent C:Si mole ratios made by the delithiation of C‐Li‐Si al‐ loys, corresponding to the blue dots in Figure 1g‐l. The XRD pattern of C0.9Si0.1, shown in Figure 2g, comprises a large broad peak near 26° and a less intense broad peak near 45°, which is similar to the XRD pattern of ball milled graphite shown in Figure 3d. As the Si content is increased, the XRD patterns are consistent with either a nm‐scale 2‐phase mix‐ ture of a‐Si and amorphous carbon (a‐C) or an amorphous C‐Si solid solution. The XRD pattern of the most Si‐rich sample, C0.4Si0.6, shown in Figure 2l is nearly the same as that of pure a‐Si, shown in Figure 3f which consists of two broad peaks, near 28° and 50°. In addition, XRD patterns of C‐Si alloys with 30 at. % Si, comprise a peak near 44°, which may be attributed to Fe contamination during the ball mill‐ ing process. No evidence of SiC formation was observed by XRD.

Figure 4. (a) An SEM image of a C0.8Si0.2 particle. (b) An SEM image of a cross‐sectioned C0.8Si0.2 particle and corresponding elemental mapping for (c) carbon and (d) silicon. (e) A scanning TEM image of a C0.8Si0.2 particle and elemental mapping for (f) carbon, (g) silicon, and (h) both carbon and silicon. (i‐j) TEM images of C0.8Si0.2 particles and (j, insert) an SAED pattern. (k‐l) Si2p and C1s XPS spectra of C0.8Si0.2.

All C‐Si alloys have similar particle morphology. The overall particle size of C‐Si alloys ranges from 5 µm to 50 µm. As an example, Figure 4a shows an SEM image of an

C0.8Si0.2 alloy particle. Large particles are composed of ag‐ gregates of smaller primary particles. Figure 4b shows an SEM image of a cross‐section formed at a broken edge of a


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C0.8Si0.2 particle. It confirms that large particles are com‐ posed of smaller primary particles with void spaces be‐ tween each other. Within the particle, C and Si elements are evenly distributed, as confirmed by the SEM‐EDS elemental mapping images shown in Figure 4c‐d. Figure 4e shows a scanning TEM image of a C0.8Si0.2 primary particle. The im‐ age is consistent with the primary particles consisting of re‐ gions of pure carbon and regions comprising pure Si or a C‐ Si solid solution. Both interpretations are consistent with the XRD patterns shown in Figure 2. No crystalline/nano‐ crystalline microstructures can be observed in high magni‐ fication TEM images or SAED patterns, as shown in Figure 4i‐j, indicating that the C‐Si alloys are completely amor‐ phous. According to the XPS Si2p and C1s spectra, the C0,8Si0.2 alloy contains C‐C, C‐O, Si‐Si and C‐Si bonds (Figure 4k‐l), but show no evidence of SiC formation. Here the C‐C and C‐ O bonds likely originate from typical XPS ambient surface contamination, whereas the Si‐Si and C‐Si bonds are con‐ sistent with the TEM results, indicating that C‐Si solid solu‐ tions exist, but without SiC formation. As a summary of the XRD, SEM, TEM and XPS results the C‐Si alloy samples are composed of porous secondary C‐Si particles comprising aggregates of smaller primary parti‐ cles. The primary C‐Si alloy particles are composed of coex‐ isting amorphous Si and amorphous C‐Si solid solution phases, but do not comprise any SiC. Figure 5 shows the measured densities of C‐Si alloys and calculated values based on a physical mixture of ball milled graphite and a‐Si. The density of a‐Si made by ethanol deli‐ thiation is 2.26 g mL‐1,24 and the density of ball milled graph‐ ite was measured to be 2.04 g mL‐1. The measured densities of the C‐Si alloys correspond well to a mixture of ball milled graphite and a‐Si. However, most samples had lower densi‐ ties than predicted from a simple physical mixture. This might arise from the aggregate nature of the alloy particles, as shown in Figure 4a‐b, since such aggregates comprise isolated voids, which would lower the apparent density as measured by pycnometry. Since most measured densities were less than that of a simple mixture of ball milled graph‐ ite and a‐Si, this further confirms that any Fe (7.87 g mL‐1) content in the C‐Si alloys is likely < 0.5 atomic %. Electrochemical characterization of C‐Si alloys. Fig‐ ure 6a‐b shows the voltage profiles of a ball milled graphite electrode, an a‐Si electrode, and C‐Si electrodes with differ‐ ent C:Si ratios for the first two cycles. All the C‐Si curves

have sloping plateaus which are characteristic of amor‐ phous Si‐based alloys, consistent with the XRD results. With the increasing Si content, the low voltage lithiation plateau becomes relatively longer, and the capacity increases ap‐ proximately linearly. Figure 6c‐d shows the differential ca‐ pacity curves of the same electrodes as shown in Figure 6a‐ b. These curves contain no features characteristic of graph‐ ite staging, suggesting the layered structure of graphite has been destroyed.28 Indeed, the differential capacity of ball milled graphite is relatively featureless. During discharge (lithiation) differential capacity of the a‐Si electrode and the C‐Si electrodes have two broad peaks at 0 – 0.1 V and ~0.22 V, corresponding to two sloping plateaus in voltage profiles. The high‐voltage peak (~0.22V) represents the Li‐Si neigh‐ bors filling during lithiation of Si and the low‐voltage peak (0 – 0.1V) represents both of the Li‐Li neighbors filling dur‐ ing lithiation of Si and the lithiation of disordered carbon.29‐ 31 During charge (delithiation), two analogous peaks in the differential capacity at 0.3V and 0.5V correspond to the del‐ ithiation of a‐LixSi.2 Sharp peaks in the differential capacity at about 0.45 V correspond to the delithiation of crystalline Li15Si4.1 The Li15Si4 peak generally decreases in area as the carbon content is increased, indicating that the formation of this phase is becoming suppressed during lithiation.

Figure 5. Density of C‐Si alloys. Red spots represent the meas‐ ured values and the black solid line represents the calculated values based on a physical mixture of ball milled graphite (measured to be 2.04 g mL‐1) and a‐Si (2.26 g mL‐1).24



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Figure 6. (a‐b) Voltage profiles and (c‐d) differential capacity curves of the ball milled graphite, amorphous Si, and C‐Si electrodes for the first two cycles. (e) Cycling performance and (f) coulombic efficiency (CE) of C‐Si electrodes. (g) Reversible, irreversible, 50th and 100th cycle capacities of C‐Si electrodes; (h) measured and calculated reversible capacities of C‐Si electrodes. The blue dash‐dot line and the black dashed line are calculated capacities based on the physical mixture of a‐Si and ball milled graphite. The red solid line is based on C‐Si phase diagram considering the maximum amount of SiC formation. (i) Rate capability of C‐Si electrodes. Ball milled C0.5Si0.5 is also shown for comparison.

The position of the two broad delithiation differential ca‐ pacity peaks is roughly constant for all C‐Si electrodes and the a‐Si electrode, suggesting the carbon matrix does not in‐ fluence the LixSi delithiation potential. However, as the car‐ bon proportion is increased, the two broad lithiation peaks generally shift to lower voltage. This peak‐shift phenome‐ non has been observed previously in sputtered C‐Si films11 and was thought to be due to increased impedance from SiC formation. However, no SiC is observed in the materials made here. More recently, alloy lithiation voltage depres‐ sion is thought to be caused by the compressive stress act‐ ing on Si as it is expanded during lithiation within a fixed matrix.11,32 This potential shift results in the suppression of crystalline Li15Si4 formation, as is also observed here.

Figure 6e‐g show the cycling performance of C‐Si elec‐ trodes in Li half‐cells. With increasing Si content, the re‐ versible capacity increases approximately linearly, while the irreversible capacity decreases, as shown in Figure 6g. For carbon‐rich (C0.9Si0.1, C0.8Si0.2, and C0.7Si0.3) electrodes, the cycling performance is stable: the retention after 100 cy‐ cles is high (~ 90%) and capacities of 50th and 100th cycles are similar (Figure 6g). For higher Si content alloys (C0.6Si0.4, C0.5Si0.5, and C0.4Si0.6), capacity fade becomes apparent and increases with increasing Si content. This may be due to the larger volume expansion of these compositions, resulting in greater disruption of the electrode coating, a greater disrup‐ tion of the SEI layer, and increased possibility of particle fracture during cycling. All these effects can lead to particles becoming disconnected during cycling, resulting in capacity fade and poor coulombic efficiency (as will be seen below). The appearance of capacity fade is also coincidental with


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the appearance of a significant Li15Si4 delithiation peak in differential capacity curves. The formation of Li15Si4 has been shown to be associated with Si that is not bound to a fixed substrate, suggesting particle fracture resulting in ac‐ tive Si becoming disconnected from the alloy during cy‐ cling.3 Ball milled C0.5Si0.5 has almost no capacity due to the formation of the inactive silicon carbide phase (Figure 3e). The nano‐Si0.3C0.7 electrode, made from a blend of nano‐Si and graphite, has a higher reversible capacity than the C0.7Si0.3 electrode made by the chemical delithiation method. However, the nano‐Si0.3C0.7 electrode fades quickly in com‐ parison. It was estimated its capacity may become lower than the C0.7Si0.3 electrode at ~110 cycles. As shown in Figure 6a, all of the samples have high irre‐ versible capacities. This is expected for ball milled carbons, which suffer from high irreversible capacities due to their high surface areas.33,34 The high irreversible capacity and low reversible capacity of the high carbon samples results in low initial coulombic efficiencies (CE) of about 37%. However, as the Si content is increased the reversible capac‐ ity increases, resulting initial CE values to also increase to about 80%. The coulombic efficiency of most carbon‐rich electrodes (C0.8Si0.2, C0.7Si0.3, and C0.6Si0.4) increases rapidly to 99.6% in the first 10 cycles and increases further. For the C0.5Si0.5 electrode, the CE jumps to >99.6% in the first 10 cy‐ cles, then keeps increasing to ~99.7% at 50th cycle, but de‐ creases afterwards. For the C0.4Si0.6 electrode, the CE jumps to ~99.6% in the first 15 cycles, but decreases afterwards. These electrodes (C0.5Si0.5 and C0.4Si0.6) also have the highest amount of Si and form the most Li15Si4. Therefore, this de‐ crease is likely due to the same volume change effects that cause capacity fade, as discussed above. In summary, a higher Si content in C‐Si alloys results in higher reversible capacity and initial CE. However, a high Si content, above ~30 at.%, also leads to less stable cycling performance. Figure 6h compares the calculated and measured reversi‐ ble capacities of the C‐Si alloys. The blue dash‐dot line is the calculated capacity based on a physical mixture of a‐Si and ball milled graphite. Here, 3579 mAh g‐1 was used for the Si capacity, corresponding to the theoretical capacity of Si con‐ sidering Li15Si4 formation, and 200 mAh g‐1 was used for the C capacity, as determined by measuring the capacity of ball milled graphite. Overall, the measured capacities are ~20% lower than the calculated capacities. In a previous study, we have shown that the a‐Si made by the delithiation of Li12Si7 has a capacity of only 3037 mAh g‐1.24 This may result from impurities, such as Si oxidation, during the delithiation pro‐ cess. The black dashed line in Figure 6h is the calculated capacity based on the physical mixture of a‐Si and ball milled carbon, based on a 3037 mAh g‐1 capacity for delithi‐ ated a‐Si and a 200 mAh g‐1 capacity for the ball milled car‐ bon. In this case, the measured capacities are close to the predicted values. The capacity was also calculated based on the Si reacting with C to form inactive SiC. This corresponds to the red solid line in Figure 6h. It is noteworthy that the measured capacities are much higher than the calculated values. This indicates that very little, if any, SiC is formed in these C‐Si materials, which is consistent with the XRD re‐ sults. This is significant, since when Si and C are directly ball milled, heated together or even co‐sputtered, inactive crys‐

talline/nanocrystalline SiC readily forms, resulting in a sig‐ nificant reduction in capacity.11 For example, the C0.4Si0.6 electrode made by conventional high energy mechanical milling only has a reversible capacity of ~800 mAh g‐1.11 In comparison, the C0.4Si0.6 electrode made in this study has much higher reversible capacity (> 2300 mAh g‐1). There‐ fore, the method introduced in this study is a good way to reduce or even avoid the formation of SiC when synthesiz‐ ing C‐Si alloys, resulting in significantly improved reversible capacities. Figure 6i compares the rate performance of C‐Si alloys and C0.5Si0.5 made by ball milling, which contains SiC (corre‐ sponding XRD shown in Figure 3e). The C‐Si alloy elec‐ trodes have improved relative rate capability as the carbon content is increased. This is likely due to the lower diffusion coefficient of nano‐Si (~10‐12 cm2 s‐1)35 compared to ball milled graphite (~10‐11 cm2 s‐1)36. The ball milled C0.5Si0.5 electrode has very low cycling capacity at all C‐rates, due to the formation of inactive SiC. To investigate the volume expansion of C‐Si electrodes, electrode thickness was measured before and after one full lithiation. As shown in Figure 7a, a higher Si content leads to larger volume expansion in general, with an exception that C0.5Si0.5 has a lower volume expansion than that of C0.6Si0.4. When the Si content is low (C0.9Si0.1, C0.8Si0.2, and C0.7Si0.3), the volume expansion is low (

SiC-Free Carbon-Silicon Alloys Prepared by Delithiation as Lithium

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