Elucidating Relationships between Structural Properties of

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Elucidating Relationships between Structural Properties of Nanoporous Carbonaceous Shells and Electrochemical Performances of Si@Carbon Anodes for Lithium-Ion Batteries Jihoon Ahn, Kyung Jae Lee, Woojeong Bak, Jung-Joon Kim, Jin-Kyu Lee, Won Cheol Yoo, and Yung-Eun Sung J. Phys. Chem. C, Just Accepted Manuscript • DOI: 10.1021/acs.jpcc.5b02073 • Publication Date (Web): 16 Apr 2015 Downloaded from http://pubs.acs.org on April 23, 2015

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The Journal of Physical Chemistry

Elucidating

Relationships

between

Structural

Properties of Nanoporous Carbonaceous Shells and Electrochemical Performances of Si@carbon Anodes for Lithium-ion Batteries Jihoon Ahn,1‡ Kyung Jae Lee,2‡ Woojeong Bak,3 Jung-Joon Kim,2 Jin-Kyu Lee,1 Won Cheol Yoo,3* and Yung-Eun Sung2*

1

Department of Chemistry, Seoul National University, 1 Gwanak-ro, Gwanak-gu, Seoul 151-

747, Republic of Korea 2

Center for Nanoparticle Research Institute for Basic Science, School of Chemical &

Biological Engineering, Seoul National University, 1 Gwanak-ro, Gwanak-gu, Seoul 151-747, Republic of Korea 3

Department of Applied Chemistry, Hanyang University, Ansan, Gyeonggi-do 426-791,

Republic of Korea

1

ACS Paragon Plus Environment


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ABSTRACT

The encapsulation of silicon in hollow carbonaceous shells (Si@C) has been known to be a successful solution for silicon anodes in Li-ion batteries, resulting in many efforts to manipulate the structural properties of carbonaceous materials in order to improve their electrochemical performance. In this regard, we demonstrated how both shell thickness and pore sizes of nanoporous carbonaceous materials containing silicon anodes influence the electrochemical performance. Structurally well-defined Si@C with varying carbon-shell thicknesses and pore sizes were synthesized by a nanocasting method that manipulated the carbon-shell, and by a subsequent magnesiothermic reduction that converted amorphous silica cores into silicon nanocrystals. When employed as anode materials, it has been verified that two opposite effects exist with respect to the thickness of carbon-shell: the weight ratio of silicon and the electrical conductivity are simultaneously affected, so that the best electrochemical performance is not obtained either from the thickest or thinnest carbon-shell. Such a countervailing effect was carefully proved through a series of electrochemical performance tests and the use of electrochemical impedance spectroscopy. In addition, the pore size effect was elucidated by comparing Si@C samples with different pore sizes, revealing that larger pore can further improve the electrochemical performance assisted by enhanced Li-ion diffusion.

KEYWORDS. Li-ion battery, silicon, carbonaceous shell, magnesiothermic reduction, structure-performance relationship

2


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INTRODUCTION Among many candidates for anode materials in Li-ion batteries (LIBs), silicon is highly attractive because of its high theoretical specific capacity (i.e., 3580 mAh g-1, theoretical capacity of pure silicon at room temperature). However, silicon anodes have problems such as low electrical conductivity, pulverization, and the continuous formation of solid electrolyte interphases (SEIs) through charge/discharge cycles. Hence, much effort has been focused on the development of methods to deal with these problems.1,2 As an ideal structure that can overcome the aforementioned issues, it has been suggested that silicon nanocrystals be located inside a cavity wrapped up by an electrically conductive layer, while mass diffusion is still possible via nanopores. Such an ideal structure has been realized via the encapsulation of silicon nanocrystals in conductive and nanoporous carbonaceous materials.3-5 In these structures, it is expected that the void space inside carbon shells buffers the volume change of the silicon anode during the electrochemical cycles, while the conductive and nanoporous layer accommodates the delivery of electric charges and prompts the diffusion of Li-ions, respectively.6 Accompanied by the formation of a stable SEI layer on the carbonaceous shell, the enhanced electrochemical performance of these structures has been widely reported.7 The materialization of silicon encapsulated in hollow carbon shells and the demonstration of the advantages of these structures have been achieved by some pioneering works. Cui et al. synthesized the structure by adopting a sacrificial silica layer between the silicon nanoparticles and the carbon layer.8 Via an in situ TEM technique and SEM, they directly observed the buffering effect of the void space and the formation of a stable SEI on the carbon layer. Another approach, used by Lee et al., was to directly etch silicon nanoparticles inside the carbon shell via electroless etching.9 Via an in situ dilatometer technique, it was

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clarified that the void space mitigates the volume change of the entire cell, owing to the volume expansion of silicon during the electrochemical cycles. The degree of control over the structural properties of Si@C anode materials has been varied in order to enhance their electrochemical performances.10-13 However, explaining the structure-performance relationship still remains a challenge. Herein, we report how the pore size and the thickness of the carbonaceous shell that contains the silicon anode material affect the electrochemical performance of LIBs. Both the shell thickness and pore size of the carbonaceous materials were explicitly controlled via a nanocasting process, for which mesoporous silica (ca. 2 nm in pore size) layers with thicknesses ranging from 7.9 ± 1.3 nm to 60.3 ± 4.9 nm on a Stöber silica core were used as a template. The amorphous silica cores coated with mesoporous carbonaceous shells were then transformed from amorphous silica into silicon nanocrystals via a magnesiothermic reduction process. Because of the adjustability of the nanocasting method as well as the powerfulness of the magnesiothermic reduction process, tailor-made silicon nanocrystals encapsulated inside hollow mesoporous carbonaceous shell (Si@void@mC, where m denotes mesopore) materials were successfully produced. The as-produced Si@void@mC materials with different shell thicknesses were then employed as anode materials in LIBs in order to reveal the relationships between their structural properties and electrochemical performances. When the thickness of the carbonaceous shell was varied from around 7.6 ± 0.9 nm to 59.2 ± 3.8 nm, the best electrochemical performance was not achieved for the thinnest or thickest shell. In this regard, it was suggested that a trade-off exists between two contradictory trends: the expected specific capacity of a Si@mC composite determined from the weight ratio of silicon with respect to carbon increases with decreasing carbon contents, whereas the electrical conductivity for the composite decreases for thinner carbonaceous shells. Such a trend was 4


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verified by electrochemical performance tests for Si@void@mC samples with different carbon shell thicknesses (i.e., 7.6 ± 0.9, 15.5 ± 1.2, 59.2 ± 3.8 nm, and in the absence of a carbon shell) using a constant pore size of around 2 nm, and by electrochemical impedance spectroscopy (EIS), where the charge-transfer resistances (Rct) were determined for Si@void@mC samples with different carbon shell thicknesses. In addition, hollow microporous carbonaceous shell materials (Si@void@µC, where µ denotes micropore) were synthesized to investigate pore-size effect on electrochemical performance. When Si@void@mC and Si@void@µC materials with equal shell thicknesses (ca. 15 ~ 16 nm) were used as anode materials, a more pronounced rate performance decrease was identified for the Si@void@µC sample. Hence, it was elucidated that enhanced Li-ion diffusion is facilitated for larger pore sizes of the carbonaceous shells.

EXPERIMENTAL SECTION Materials & Reagents Aluminum

(III)

chloride

hexahydrate

(AlCl3·6H2O,

99%),

phenol

(≥

99%),

paraformaldehyde (95%), resorcinol (98%), and a formaldehyde aqueous solution (37 wt%) were purchased from Sigma Aldrich. Tetraethylorthosilicate (TEOS, ≥ 96%) and mesitylene (≥ 97%) were purchased from TCI Chemicals. Sodium hydroxide (NaOH), a 28~30% aqueous ammonia solution (NH4OH), and a 38~40% HCl concentrated aqueous solution were purchased from Samchun Chemical (Korea). Absolute ethanol was purchased from J. T. Baker. Magnesium powder with a 100~200 mesh was purchased from Alfa-Aesar. Cetyltrimethylammonium bromide (CTAB, 99%) was purchased from Daejung Chemicals. All the reagents were used without further purification. Deionized (DI) water purified with a Milli-Q purification system was used for all the experiments. 5



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Preparation of core silica particles Core SiO2 nanoparticles were synthesized by the Stöber method. For this, 12 ml of an aqueous ammonia solution (28~30 wt%), 12 ml of DI water, and 370 ml of absolute ethanol were mixed together under vigorous magnetic stirring. Then, after 10 min of mixing, 8 ml of TEOS was injected to the solution, and the mixture was allowed to react for 16 hours. Then, an additional 15 ml of TEOS was added dropwise to the solution for further growth of the particles to the targeted size. After another 16 h corresponding to the growth step, the resulting silica particles were collected by centrifugation and washed with ethanol. The centrifugation and washing cycles were repeated thrice to assure complete removal of the remaining reactants. Finally, the silica particle were dried at a reduced pressure at 80°C and stored as powders for later reactions.

Preparation of mesoporous silica shells coated on silica cores (SiO2@mSiO2-n) SiO2@mSiO2 was synthesized by introducing a mesoporous SiO2 shell around the core silica nanoparticles that were previously prepared. For this, a soft-template method based on cetyltrimethylammonium bromide (CTAB) was used.14 Typically, a ca. 60.3 ± 4.9 nm-sized mesoporous shell could be synthesized by the following reaction conditions. First, 1 g of core silica particles was fully added in a co-solvent composed of 420 g of DI water and 200 g of ethanol. Then, 6 g of an aqueous ammonia solution (28~30%) and 2.22 g of CTAB were added to the solution, while the latter was magnetically stirred. After 30 min of mixing, 4.08 g of TEOS was added to the solution, after which the mixture was allowed to react for 16 h. The resulting colloids were collected and purified by repeated centrifugation and re6


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dispersion in ethanol. The collected nanoparticles were dried at a reduced pressure at 80°C. In order to remove CTAB in the mesopores, SiO2@mSiO2 was calcined at 600°C for 10 h in a muffle furnace. In addition, the mesoporous shell thickness can be controlled to 7.9 ± 1.3 nm and 16.0 ± 1.2 nm by modifying the ratio of the employed reactants. Specifically, for 16.0 ± 1.2-nm-thick shells, 3 g of core silica, 108 g of ethanol, 225 g of DI water, 1.11 g of CTAB, 3 g of an aqueous ammonia solution, and 2 g of TEOS were used. For 7.9 ± 1.3-nm-thick shells, 3 g of core silica, 27 g of ethanol, 56.2 g of DI water, 0.28 g of CTAB, 0.75 g of an aqueous ammonium solution, and 0.5 g of TEOS were used.

Preparation of silica encapsulated in hollow mesoporous carbonaceous shells (SiO2@void@mC-n) Synthesis of SiO2@void@mC was achieved by using the mesoporous silica as a template for mesoporous carbon. The procedure begins with doping the inner wall of the mesopores in SiO2@mSiO2 with aluminum. The doped Al was then used as the catalyst for the polymerization of phenol-formaldehyde (PF) resin, which can be converted to amorphous carbon. In order to dope Al on the walls, 1 g of SiO2@mSiO2 was distributed in a co-solvent consisting of 20 ml of both ethanol and DI water, after which 1 g of AlCl3·6H2O was added to the solution. The mixture was placed in a sonicator while Al3+ ions adsorbed on the wall. After 3 h, the colloids were collected by centrifugation and dried at a reduced pressure. Then, the Al-incorporated SiO2@mSiO2 was annealed at 600°C for 3 h, resulting in Al-doped SiO2@mSiO2. PF resin was introduced into the mesopores by the following procedure. First, 1 g of Aldoped SiO2@mSiO2 was added to 100 ml of mesitylene. Then, 6 g of phenol and 1.5 g of paraformaldehyde were added to the mixture, after which the mixture was heated at 90°C for 7



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12 h. During the procedure, both phenol and paraformaldehyde polymerize in order to create PF resin in the mesopores where Al-catalytic sites exist. SiO2@mSiO2 filled with PF resin was washed three times with ethanol and then coated with another SiO2 shell to protect it from sintering during the following carbonization step. The resulting powder was carbonized at 800°C for 3 h under an Ar flow in a tube furnace. Finally, the void space between the silica core and the mesoporous carbonaceous shell was created by the partial etching of silica. For this, 700 mg of the carbonized product described above was added to a co-solvent consisting of 350 ml of both ethanol and DI water (1:1 v/v). The temperature of the vigorously stirred mixture solution was fixed at 40°C. Then, 70 ml of a 10 M aqueous NaOH solution was added to the mixture. After a certain time when the desired void space was created, the particles (SiO2@void@mC) were washed with fresh DI water until the pH of the dissolved particle solution reached 7. Typically, the etching times for SiO2@void@mC-1, SiO2@void@mC-2, and SiO2@void@mC-3 were 45 min, 45 min, and 90 min, respectively. Finally, SiO2@void@mC was dried at 80°C at a reduced pressure and stored as powder.

Preparation of silica encapsulated in hollow microporous carbonaceous shells (SiO2@void@µC) In order to produce microporous carbon, silica particles were directly coated with resorcinol formaldehyde (RF) resin. For this, 500 mg of silica particles were added to 380 ml of DI water. Then, 5 ml of a 10 mM aqueous CTAB solution and 0.2 ml of ammonia solution were added to the solution. Next, 10 ml of a solution containing 200 mg of resorcinol and 0.4 ml of formaldehyde (37 wt% in H2O) was injected to the mixture. The solution was allowed to react for 16 h, after which the protective anti-sintering silica shells were introduced to the 8


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outer surface. The resulting powder was carbonized at 800°C for 3 h under an Ar flow in a tube furnace. Finally, voids were created in the particles using the same etching method as for SiO2@void@mC. The etching time was 135 min.

Synthesis of Si@void@C via a two-step magnesiothermic reduction The remaining silica cores of SiO2@void@mC and SiO2@void@µC were reduced via a magnesiothermic reduction, resulting in Si@void@mC and Si@void@µC, respectively. First of all, SiO2@void@C was mixed with Mg powder (100~200 mesh) via grinding. The mole ratio between core SiO2 and Mg was 1:4.5 (the amount of core SiO2 in SiO2@void@C was determined by a thermogravimetric analysis). The mixture was placed in a covered boatshaped crucible and annealed in a tube furnace at 560°C for 40 h under an Ar flow. Next, the cover was removed and the mixture was further annealed at 580°C under a 5% H2/Ar flow. The annealing time of this second step ranged from 10 to 30 h, depending on each SiO2@void@C sample in order to assure complete conversion of Mg2Si into Si. After this reduction step, the side-product such as MgO was removed by immersing the product in a 1 M HCl solution for 10 h. The resulting particle radiuses corresponding to Si@void@mC-n and Si@void@µC were 178.9 ± 11.5, 227.4 ± 15.2, 318.2 ± 10.3, and 214.6 ± 9 nm.

Cell fabrication and electrochemical measurements The Si@void@C active material was mixed with Super P carbon black and a polyamideimide (PAI) binder with weight ratio of 60:20:20 in a N-Methylpyrrolidone (NMP) solvent to form a slurry. The prepared slurry was spread onto a copper foil using the doctor blade method, and the solvent was removed at a reduced pressure. To activate the PAI binder, the as-prepared electrode was heat-treated at 350°C in an Ar atmosphere for 90 min. The 9



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prepared electrodes were transferred to an Ar-filled glove box and after vacuum-drying at 120°C for 8 h, and the cells were assembled. The employed electrolyte was 1.5 M LiPF6 with EC/DEC/FEC (5/70/25 volume ratio, Panatech, Korea), and the separator was obtained from SK Innovation. Lithium metal was used as a counter electrode, and a cell of type 2016 was used. Galvanostatic charge/discharge tests were performed between 0.01 to 1.5 V vs. Li/Li+ with various current densities using a WBCS3000 cycler (Won-A Tech, Korea). Electrochemical impedance spectroscopy (Zahner, Germany) was conducted at 0.25 V with a 10 mV amplitude, and after three cycles, the frequency range was from 100 kHz to 50 mHz. The equilibrium was achieved by maintaining the voltage at 0.25 V for 12 h.

Materials Characterization Transmission Electron Microscope (TEM) images were acquired by a Hitachi-7600

system (Hitachi, Japan). The thermo-gravimetric analysis (TGA) was performed using an SDT Q600 device (TA Instruments Inc.). High-Resolution TEM (HR-TEM) images, Energy Dispersive Spectroscopy (EDS) mapping, and Selected Area Electron Diffraction (SAED) patterns were acquired using a Tecnai F20 system (FEI). Elemental Analysis (EA) data was acquired by a Flash EA 1112 (Thermo Electron Corp.) XRD (Rigaku, DMAX2500-PC, with Cu Kα radiation (λ = 1.5406 Å) at 40 kV and 100 mA) setup. The nitrogen sorption isotherms were measured at 77 K using liquid nitrogen on a Belsorp Mini-II device. All samples were degassed at 160°C in a static vacuum (p < 10-5 mbar) for 12 h. The Brunauer–Emmett–Teller (BET) method was applied to estimate the specific surface areas and the total pore volumes were obtained at P/P0 (= 0.99). The pore size distributions were 10


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obtained from the Barrett-Joyner-Halenda (BJH) method adopted from the adsorption branches of all samples.

RESULTS AND DISCUSSION Preparation of the SiO2@void@mC and SiO2@void@µC materials In an attempt to synthesize the Si@void@mC or Si@void@µC materials depicted in Scheme 1, first, carbonaceous shells that were coated on amorphous silica core composites (i.e., SiO2@void@mC and SiO2@void@µC) were produced. The Stöber method was employed to synthesize the amorphous silica particles.15 The average size of the silica particles was 205 ± 11 nm, as determined by TEM and SEM measurements (Figure S1), and these particles were used as core particles for the following reactions.

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Scheme 1. Schematic representation of the preparation of Si@void@mC-n and Si@void@µC samples. The encapsulation of amorphous silica cores was carried out via two different methods. The first one consisted of using the nanocasting method, in which a thickness-controlled mesoporous silica layer coated on the amorphous (and nonporous) silica core was used as a hard template for a carbonaceous replica, hence producing a series of SiO2@void@mC materials. The second process relied on the direct wrapping of a carbon shell around the silica core, resulting in SiO2@void@µC materials. In the case of SiO2@void@mC composites, a mesoporous silica layer was coated on the nonporous silica core. In this approach, the silica layer was created via the slightly modified Stöber method based on the use of a surfactant (herein, cetyltrimethylammonium bromide, CTAB), which, through micellization with the 12


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silica source, resulted in the generation of mesopores inside the silica layer.14 The thickness of the mesoporous silica shell was controlled by varying the experimental conditions (see Experimental Section for details), in which the three different mesoporous silica shell thicknesses identified by TEM were 7.9 ± 1.3, 16 ± 1.2, and 60.3 ± 4.9 nm (Figure S2). These

non-porous-silica-core@mesoporous-silica-shell

structures

were

denoted

as

SiO2@mSiO2-n (in which n = 1, 2, 3, in the order of increasing shell thickness). A nitrogen sorption analysis was then performed to identify the porosity as well as the pore size distribution (Figure S3). The typical type-IV isotherm shape obtained for all samples was attributed to the mesoporosity of the samples, while the specific surface areas that were determined by the BET method increased for thicker shells (Table S1), implying that the porosity originated from the mesoporous silica shell.16 In addition, the pore size distribution calculated by the BJH method, which was adopted from the adsorption branch of the isotherms, showed narrow pore size distributions of around 2.2-2.4 nm, thereby exhibiting almost equal sizes for all samples (Figure S3b). As-produced SiO2@mSiO2-n materials were then transformed to the SiO2@mC-n materials via the nanocasting method, in which phenolic resin was deposited inside the mesopores of the silica shell through an acid-site specific polymerization and subsequent carbonization. In order to bear acid catalysts on the surface of the mesopores for the polymerization process, incorporation of Al was achieved (see Experimental Section). Phenolic resin was polymerized in situ inside the mesopores, where the Al dopant acted as a polymerization catalyst,17,18 after which the resin-incorporated SiO2@mSiO2-n materials were wrapped by another silica layer in order to prevent particle aggregation during carbonization.19 The carbonaceous materials were synthesized at 800°C in an inert atmosphere, and the carbon containing SiO2@mSiO2-n samples were then immersed in a basic solution for partial silica 13



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etching in order to finally produce the SiO2@void@mC materials. It should be noted that the size of the void space and the remaining silica core particles could be controlled by variation of the etching time (Figure S4 and Experimental Section).20 In order to leave the thickness of the carbon shell as the only factor of influence during the electrochemical performance tests, the size of the remaining silica core of the SiO2@void@mC materials was fixed at 156.6 ± 8.6 nm. As-made composites are denoted as SiO2@void@mC-n and results from detailed TEM analyses are shown in Figure 1a-c.

Figure 1. TEM images of (a) SiO2@void@mC-1, (b) SiO2@void@mC-2, (c) SiO2@void@mC-3, and (d) SiO2@void@µC. Meanwhile, microporous carbonaceous shells can be introduced on the core silica particles by direct coating of phenolic resin on the silica core, without using the mesoporous silica template.21 The thickness of the phenolic resin that was directly coated on the silica cores was 17.6 ± 2.0 nm (Figure S5). After the coating of a protective silica layer, these structures were carbonized under the same conditions. The resulting carbonized product was also partially etched in a basic solution to produce a slightly etched silica core, whose size was 157.7 ±

14


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13.1 nm, and was denoted as SiO2@void@µC (Figure 1d). The thickness of carbon shell for SiO2@void@µC was 14.6 ± 1.0 nm. The textural properties of the SiO2@void@mC-n samples were characterized using nitrogen sorption measurements. It was observed that more of type-I features, which represent microporosity, were exhibited by the samples with thicker carbonaceous shells, while all samples showed type-IV features indicating mesoporosity (Figure 2a). The BJH pore size distributions extracted from the adsorption branches of the SiO2@void@mC-n samples revealed relatively narrow pore-size distributions around 2-3 nm, hence confirming that a successful replication from the parent templates was achieved. It should also be mentioned that the observed broadly distributed pores with sizes of tens of nm probably originated from the void space between the core silica and the carbon shell. Moreover, a gradual increase of the specific surface areas and more intense pore size distributions appeared for thicker carbon shell samples (Table S2). On the other hand, the SiO2@void@µC sample did not show any mesopores of less than 10 nm in size and exhibited a much smaller BET surface area, while this SiO2@void@µC sample also shows pores with sizes of tens of nanometer that seemed to be formed at the void space between the core and the shell (Figure 2b and Table S2). Therefore, it can be concluded that SiO2@void@mC-n and SiO2@void@µC samples resemble each other in appearance but have completely different carbonaceous shell porosities.

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Figure 2. (a) N2 sorption isotherm (the offset was set to 150, 300, 500 cm3/g for SiO2@void@mC-1, SiO2@void@mC-2, and SiO2@void@mC-3, respectively) (b) BJH pore diameter distributions extracted from the adsorption branches of the isotherms. Lastly, the amount of carbon contents in the SiO2@void@mC-n and SiO2@void@µC samples was determined by a thermogravimetric analysis (TGA) in air (Table S2, Figure S6). The weight loss during the TGA experiment can be attributed to carbon content, given the robust nature of silica under these conditions. According to these results, SiO2@void@mC-n samples with thicker carbonaceous shells clearly have higher carbon content.

Preparation of Si@void@C via the magnesiothermic reduction of SiO2@void@C 16


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Silicon nanocrystals were transformed by a two-step magnesiothermic reduction, starting from amorphous silica cores coated with either mesoporous or microporous carbon shell materials, and in such a way, that the morphologies remained unchanged (Figure S7). For the first step, it should be noted that an excess of magnesium was used for the complete and homogeneous reduction of silica to magnesium silicide (Mg2Si) with a molar ratio given by Mg:SiO2 = 4.5:1, resulting in the composite of magnesium silicide and magnesium oxide (MgO) obtained by Reaction (1) and (2).22,23

SiO2(s) + 4Mg(g) → Mg2Si(s) + 2MgO(s)

(1)

Si(s) + 2Mg(g) → Mg2Si(s)

(2)

According to XRD patterns of the SiO2@void@mC-3 sample after the first reduction step, the as-reduced sample was only composed of Mg2Si and MgO (Figure S8a), thereby strongly supporting the fact that instead of silicon, magnesium silicide is the major product between the silica and the excess of magnesium under the given conditions. Note that in the magnesiothermic reduction, magnesium silicide is thermodynamically favored when compared to the formation of silicon.24 In addition, volume expansion of inside contents of the as-reduced composite (i.e., Mg2Si+MgO@mC-3) with respect to the silica core was observed by TEM which is caused by the reaction between magnesium and silica (Figure S8b). In the second reduction step, the as-reduced sample was annealed again at a similar temperature, but in a different atmosphere, specified by 5% H2 in Ar. It has been reported that the hydrogen mixture was used to reduce silica to silicon via the magnesiothermic reduction 17



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process, thereby exploiting the inherent reducing nature of hydrogen gas during the reduction process.25-27 In the second reduction step of the Si@void@mC-3 sample, the conversion reaction was completely finished under the hydrogen mixture conditions and a reaction time of 10 h, whereas conversion under an Ar-only condition hardly occurred even for much longer reaction times (Figure S8a). The latter strongly indicates the beneficial effect of the hydrogen mixture for this second reduction step under the given conditions. Using the established two-step reduction method, all of the prepared SiO2@void@mC-n and SiO2@void@µC materials were reduced. To remove the unwanted magnesium oxide side-product, and the remaining magnesium silicide, the samples were treated with 1 M HCl. XRD patterns of the reduced products clearly showed only diffraction peaks attributed to silicon (Figure 3). Note that relatively smaller intensity observed for the SiO2@void@mC-3 sample can be attributed to its thicker shell thickness (approximately 60 nm), because the diffraction of the graphitic region of amorphous carbonaceous materials appeared around 2030°. The as-reduced products originating from the SiO2@void@mC-n and SiO2@void@µC samples were named Si@void@mC-n and Si@void@µC, respectively.

Figure 3. XRD patterns of a series of Si@void@mC-n and Si@void@µC samples.

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The morphologies of the Si@void@mC-n and Si@void@µC samples were confirmed by TEM measurements. After the two-step reduction process and the acid washing step, the thickness and the hollow spherical shape of the carbonaceous shells remained unchanged (Figure 4). However, when comparing the inner contents, entirely different shapes were found. According to results of the TEM-EDS mapping analysis displayed in the insets of Figure 4a-d, the indication of silicon intensity was obtained from the core part, while the carbon signal was mainly collected from the peripheral region, which corresponds to the carbon shell. According to the HR-TEM investigation, the dark objects inside the carbon shell were identified as silicon nanocrystals with a (111) silicon lattice plane and a d-spacing of 3.1 Å (Figure 4e-l). In addition, such small particles show a single crystalline silicon feature, which was identified by the electron diffraction analysis (insets in the Figure 4i-l).

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Figure 4. (a ~ d) Normal-TEM images (inset: EDS mapping) and (e ~ l) HR-TEM images (inset : SAED patterns) of Si@void@mC-n and Si@void@µC samples.

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For further characterization, the sizes of the as-produced silicon nanocrystals of the Si@void@mC-n and Si@void@µC samples were obtained using the Scherrer equation (Table 1). Because of the broad band in the range of 20-30° in the XRD patterns, which corresponds to the carbonaceous shell, not only the first peak at 28.4°, but also the second peak at 47.3° was used for the size-calculation of the silicon nanocrystals, whose sizes are around 16-20 nm for all samples.

Table 1. Crystalline size and carbon contents (weight percent) of Si@void@mC-n and Si@void@µC samples

Crystalline sizea)

Carbon weight percentb)

[nm]

[wt%]

Sample

Si weight percentc) [wt%]

1st peak

2nd peak

TGA

EA

SEM-EDS

TGA residue

Si@void@µC

20.0

16.8

40.0

37.2

39.6

60

Si@void@mC-1

20.7

18.8

17.1

14.8

23.0

82.9

Si@void@mC-2

16.3

16.3

43.8

35.2

43.0

56.2

Si@void@mC-3

17.7

17.3

81.9

60.9

74.4

18.1

a)

The crystalline size of Si@void@C was calculated by the Scherrer equation using the parameters of both the 1st peak (28.4°) and 2nd peak (47.3°) with a K factor of 0.94. b)

Carbon weight percent was extracted from results of the thermogravimetric analysis in air (TGA), the elemental analysis (EA), and the SEM Energy Dispersive X-ray Spectroscopy (SEM-EDS) analysis. c)

The weight percent of the residue after carbon combustion in the TGA experiment was considered as the Si weight percent of each sample.

Lastly, the carbonaceous shell contents were characterized for the samples by various techniques (Table 1, Figure S9). By combining the results from TGA measurements, elementary analysis (EA), and SEM-EDS analysis, it was confirmed that the carbon contents 21



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constantly increased for samples with thicker carbon shells (e.g., Si@void@mC-3 exhibits the highest carbon content in all measurements, as listed in Table 1). Furthermore, the carbon contents in Si@void@mC-2 and Si@void@µC samples were similar to each other, which is consistent with what was already observed for the parent samples SiO2@void@mC-2 and SiO2@void@µC (Table 1 and Table S2). When comparing the results of the TGA measurements for the parent and the produced samples, increased carbon contents for all the produced samples were observed, which is probably caused by the decrease of oxygen species in the silica core after the reduction process. Meanwhile, it was suggested that the silicon content in each sample was determined as the residual weight percentage from TGA experiments (Table 1). In this regard, it is worth mentioning that confirming silicon contents via TGA experiments has been broadly accepted, because silicon remains intact where carbon starts to burn out.28,29 Therefore, samples with different carbon shell thicknesses (a series of Si@void@mC-n samples) and samples with a constant carbon shell thickness but with different pore sizes (a series of Si@void@mC-2 and SiO2@void@µC samples), were successfully synthesized and further investigated as anode materials in order to figure out the relationships between their structural properties and electrochemical performances. The corresponding results are discussed in the next section.

Electrochemical properties of Si@void@mC-n and Si@void@µC As-made samples were utilized as anode materials for LIBs in order to reveal the effect of the carbon shell thickness and pore size on the electrochemical performances. First, the effect of the carbon shell thickness on the electrochemical performance was investigated by comparing a series of Si@void@mC-n samples and a sample without carbon shell, which was made by the same magnesiothermic reduction reaction. Upon measuring the cycling 22


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performance at a fixed rate of 400 mA g-1 (Figure S10), the capacity fading for the sample without a carbon shell was much faster than those corresponding to Si@void@mC-n samples, evidencing the fact that the carbon shell improves the stability of the active material during cycles, as previously reported by a number of research groups.3,30 In addition, the stability of cycling performance was not related with carbon shell thickness. The thinnest carbon shell can also ensure the stability of silicon anode material. Meanwhile, it was initially presumed that the specific capacity was only affected by the carbon shell thickness, that is, by the weight ratio of silicon over carbon. When this assumption was used to specify the theoretical capacity as a specific capacity averaged with both the carbon shell and the silicon, the first-cycle capacities for the samples Si@void@mC3 and Si@void@mC-2, given by with 514 mAh g-1 and 1184 mAh g-1, respectively, wellmatched the expected values (Figure 5). It should be noted that the specific capacities of the carbon shell and the silicon for the averaged specific capacity were obtained from experimental measurements and theoretical estimations considering an intrinsic silica layer located on the silicon nanocrystals (see Supporting information for details).

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Figure 5. Non-linear correlation between the Si-C content and electrochemical properties of Si@void@mC-n and Si without a carbon shell; the first delithiation capacities of the Si@void@mC-n samples and the maximum delithiation capacity of Si without a carbon shell are observed at 400 mA g-1. The theoretical capacities are calculated under the assumption that the silicon nanocrystals have a size of around 20 nm and an intrinsic silica layer (see Supporting Information for details).

However, the measured capacity of Si@void@mC-1 (1201 mAh g-1) deviated sharply from the theoretical expectation (Figure 5). In other words, capacity was not increased as compared to Si@void@mC-2, which was hardly accounted for when considering the difference of silicon content for Si@void@mC-1 and Si@void@mC-2. Similarly, for the sample without the carbon shell, despite the presumable activation process in the silicon active material during the first few cycles, which is due to electrolyte wetting,31,32 the maximum specific capacity after several cycles was 1172 mAh g-1, similar to those of the Si@void@mC-1 and Si@void@mC-2 samples (Figure 5). As a result, it can be concluded that the mass ratio of the active materials cannot be simply correlated with the specific 24


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capacity of the anode materials, strongly giving rise to speculations related to other parameters that could have an impact on the electrochemical performances. It is suggested that the non-linear correlation between the thickness of the carbon shell and the specific capacity can be explained by considering two contradictory influences of the carbonaceous materials. In the first place, it is obvious that an increase of the carbon shell thickness decreases the specific capacity because of the reduced mass ratio of silicon. Simultaneously, carbon coating of anode also enhances the capacity of electrodes due to an electrical conductivity enhancement; in this regard, an enhanced electrochemical performance was reported for thicker carbon shells coated on the electrode material.33,34 This effect is possibly explained by electrons from outside the conducting agent that move through the carbonaceous shell until they reach the silicon anode materials,35 where the resistance would be changed by the shell thickness since this parameter is inversely proportional to the crosssectional area of the carbon shell. In other words, more effective electron pathways would be formed for thicker carbonaceous shells. In this regard, it seems like the Si@void@mC-1 anode has more silicon but an insufficient carbonaceous shell thickness for effective electron transfer when compared to the Si@void@mC-2 anode, hence resulting in a limited use of active material. Such a countervailing relationship of the two effects is thought to be the reason for the nonlinear correlation between the Si-C mass ratio and the specific capacity. An effort to prove this hypothesis was done via EIS (Figure S11). The charge-transfer resistance (Rct), shown as the size of a semi-circle in the Nyquist plot, represents the amount of faradaic currents flowing under a potential perturbation. Thus, Rct normalized by the mass of active material (silicon) could represent the efficiency of the active-material use. In Figure S11, Rct increases as the carbonaceous shell thickness decreases, indicating that the reduced electrical conductivity limits the use of active materials. 25



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The influence of both the carbon shell and the electrical conductivity was also observed in power performance measurements of the Si@void@mC-n samples and the sample without carbon shell (Figure 6a). When the rate was increased in a stepwise manner from 400 mA g-1 to 3200 mA g-1, the sample without carbon shell showed the most severe decrease in normalized capacity ratio, while all Si@void@mC-n samples showed an enhanced power performance, thus confirming the beneficial effect of the carbon shell.36,37 On the other hand, it can unambiguously be observed that power performance enhancement is related to the thickness of carbon shell, since the sample with the thickest carbon shell (Si@void@mC-3) showed the smallest decrease in normalized capacity as a function of the current density (Figure 6b). The reason of this phenomenon could be explained by the enhanced electrical conductivity of thicker carbon shells. Note that at high rates, the conductivity becomes more crucial for the performance, suggesting that silicon covered with thicker carbon shells could exhibit a better performance. Consequently, when compromising the two contradictory effects for the carbon-shell thickness, the Si@void@mC-2 sample with a relatively higher capacity and a prominent rate performance turns out to be the most appropriate anode material from the series of Si@void@mC-n samples.

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Figure 6. (a) Rate-dependent power performance for Si@void@mC samples and a Si sample without carbon coating; (b) rate capability of the samples normalized by the capacity at 400 mA g-1; (c) peak potentials as a function of the rate, relating Si lithiation/delithiation in the dQ/dV plot. 27



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For a detailed analysis, differential capacity profiles (dQ/dV) were derived from the lithiation/delithiation profiles and the peak voltages were depicted in Figure S12, Figure S13, and Figure 6c. The peak voltages represent the voltages where the charge/discharge reactions occur. At a rate of 400 mA g-1, the peaks observed at around 0.2 V and 0.1 V, and those at around 0.3 V and 0.5 V were assigned to the lithiation and delithiation of silicon, respectively; these two reaction peaks are related to different Li insertion/deinsertion environments.38,39 All Si@void@mC-n samples showed a voltage shift for larger rates owing to an over-potential, but their extents were different for each sample. The Si@void@mC-2 and Si@void@mC-3 samples showed a relatively smaller shift of the potential peak than the Si@void@mC-1 sample, as depicted in Figure 6c. This is strongly related to the fact that the Si@void@mC-1 sample has the thinnest carbon shell, and therefore, the most limited electron conductivity, as already mentioned above. Therefore, it could be suggested that major contribution of over-potentials for the Si@void@mC-n series was electric conduction. It is noteworthy that ion diffusion could also be the source of the over-potential, however, it has been reported that the mesopores in carbon shells seem to significantly relieve problems arising from ion diffusion.40 Our observation and interpretation are consistent with the assumption that electron transfer to the silicon is important for a better capacity at high current densities, as has already been proposed in many other studies.41 Furthermore, the pore size effect was investigated using the Si@void@mC-2 and Si@void@µC samples. Both sample-types are characterized by similar shapes, carbon shell thicknesses, void spaces, Si/C contents ratio, and silicon crystallinity, but exhibit different pore sizes of around 2 nm (Si@void@mC-2) and micropores (Si@void@µC) (Table 1 and Table S2). Judging by the BJH pore-size distribution (Figure 2b), it is expected that the 28


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Si@void@mC-2 sample with larger pore sizes exhibits a faster Li-ion diffusion rate through the carbon shell, when compared to the Si@void@µC sample. First, the cycling performances of the Si@void@mC-2 and Si@void@µC samples were measured at a fixed rate of 400 mA g-1 (Figure S14). Both samples have almost the same initial capacity and retention, complying with the suspicion that an expected Li-ion diffusion rate difference could not be prominent at a rate of 400 mA g-1. Hence, the rate performances were monitored for rates ranging from 400 mA g-1 to 5400 mA g-1 (Figure 7a). From these measurements, it can be seen that the capacity retention of the Si@void@mC-2 sample is better than that of the Si@void@µC sample for rates above 1600 mA g-1.

29



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Figure 7. (a) Power performance at each rate for the Si@void@mC-2 and Si@void@µC samples; (b) rate-dependent peak potentials, relating Si lithiation/delithiation in the dQ/dV plot. At high rates, Li-ion diffusion as well as electrical conductivity likely becomes more important as mentioned above.42 First, the electrical conductivity for both samples is considered because the different precursors of resorcinol-formaldehyde (RF) and phenolformaldehyde (PF) were used for the Si@void@mC-2 and the Si@void@µC samples, respectively. It is reported that the different electrical conductivities were recorded from the different precursors used.

43

Hence, the electrical conductivity for both samples was

investigated through EIS measurements (Figure S11). In the Nyquist plot of Figure S11, the semi-circle radii in high frequency region (i.e. charge-transfer resistance) corresponding to Si@void@mC-2 and Si@void@µC were similar to each other. It could be interpreted that both samples have similar electrical conductivities, because the charge-transfer resistance in high frequency region is related with electron transfer efficiency.44 Therefore, deviation of power performances at the high rates could be attributed to the difference of Li-ion diffusion through carbon shells. The differential capacity profiles (dQ/dV) of the Si@void@mC-2 and Si@void@µC samples were also derived from the voltage profiles for each rate (Figure S15 and Figure S16). In the low rate region, the peak shifts for both electrode materials were not significant in the dQ/dV plot. However, as the rate increases, the peak shifts due to the over-potential prominently appear, as observed in the series of Si@void@mC-n samples. Moreover, it can be observed that the voltage shift is worse for the Si@void@µC sample as the rate increases (Figure 7b), which is indicative for the higher over-potential degree of the Si@void@µC sample when compared to the Si@void@mC-2 sample. Therefore, the larger over-potential at 30


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high rates could be ascribed to low Li-ion diffusion rate for the Si@void@µC sample, which is caused by the small pore sizes of the carbonaceous shell.

CONCLUSION In this article, we have shown how the structural properties of a carbonaceous shell coated on silicon nanocrystals influences the electrochemical performance when these materials are used as anode materials in LIBs. In an attempt to fine-control the structure-parameters such as the thickness and the pore size of the carbon shell materials, a nanocasting method was applied, in which thickness- and pore-size-controlled mesoporous silica layers coated on either non-porous or amorphous silica cores were employed as hard templates. After a twostep magnesiothermic reduction process, the silicon nanocrystals were converted from amorphous silica cores, while maintaining the structural features of the carbonaceous shells. Taylor-made series of Si@void@mC-n and Si@void@µC samples, as well as a sample without carbon shell, were employed as anode electrodes in LIBs. First, the role of the carbon shell itself was identified, which has already been reported by many research groups. In addition, we have proven the contradictory trend that exists when relating the thickness of carbon shell to the electrochemical performance; the specific capacity was strongly influenced by the mass ratio of the active material and carbon, whereas thicker carbon shells provided better electrical conductivity, the latter being caused by efficient electron transfer pathways. By combining both of these countervailing effects of carbonaceous shell materials, an optimal thickness of the carbon shell, which can provide sufficient electron-transfer enhancement as well as minimize the capacity decrease, is discussed in order to use the active material (herein, silicon) as a viable possibility. Furthermore, the pore-size effect was also addressed, which is based on the fact that larger pores inside the carbon shell further enhance 31



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the electrochemical properties when compared to smaller pores due to elevated Li-ion diffusion, especially at higher rates. Lastly, the designed carbon coating on the silicon anode could still have possibilities to enhance the electrochemical performance in order to improve the performance of LIBs.

ACKNOWLEDGMENT This work was supported by the Basic Science Research Program through the National Research Foundation of Korea (2014R1A1A2057204) and IBS-R006-G1.

AUTHOR INFORMATION Corresponding Authors *

Prof. Won Cheol Yoo, Tel: +82-31-400-5504, E-mail: [email protected]

*

Prof. Yung-Eun Sung, Tel: +82-2-880-1889 E-mail: [email protected]

Author Contribution ‡

These authors contributed equally.

ASSOCIATED CONTENT Supporting Information. Detailed material characterization (SEM, TEM, N2 sorption, and XRD) and electrochemical performance tests (cycling performance, voltage profile, dQ/dV

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plot, and Nyquist plot). This information is available free of charge via the Internet at http://pubs.acs.org.

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Elucidating Relationships between Structural Properties of

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