Variation in Crystalline Phases: Controlling the Selectivity between

Feb 7, 2016 - Jihoon Ahn†, Hee Soo Kim∥, Jung Pyo†, Jin-Kyu Lee⊥, and Won Cheol Yoo ... Xuan-Manh Pham , Choong-Nyeon Park , and Chan-Jin Park...
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Variation in Crystalline Phases: Controlling the Selectivity between Silicon and Silicon Carbide via Magnesiothermic Reduction using Silica/Carbon Composites Jihoon Ahn,† Hee Soo Kim,∥ Jung Pyo,† Jin-Kyu Lee,⊥ and Won Cheol Yoo*,∥ †

Department of Chemistry, Seoul National University, 1 Gwanak-ro, Gwanak-gu, Seoul 151-747, Republic of Korea Department of Applied Chemistry, Hanyang University, Ansan 426-791, Republic of Korea ⊥ LG Chem., 188 Munji-ro, Yuseong-gu, Daejeon 305-738, Republic of Korea ∥

S Supporting Information *

ABSTRACT: Magnesiothermic reduction of various types of silica/carbon (SiO2/C) composites has been frequently used to synthesize silicon/carbon (Si/C) composites and silicon carbide (SiC) materials, which are of great interest in the research areas of lithium-ion batteries (LIBs) and nonmetal oxide ceramics, respectively. Up to now, however, it has not been comprehensively understood how totally different crystal phases of Si or SiC can result from the compositionally identical parent materials (SiO2/C) via magnesiothermic reduction. In this article, we propose a formation mechanism of Si and SiC by magnesiothermic reduction of SiO2/C; SiC is formed at the interface between SiO2 and carbon when silicon intermediates, mainly in situ-formed Mg2Si, encounter carbon through diffusion. Otherwise, Si is formed, which is supported by an ex situ reaction between Mg2Si and carbon nanosphere that results in SiC. In addition, the resultant crystalline phase ratio between Si and SiC can be controlled by manipulating the synthesis parameters such as the contact areas between silica and carbon of parent materials, reaction temperatures, heating rates, and amount of the reactant mixtures used. The reasons for the dependence on these synthesis parameters could be attributed to the modulated chance of an encounter between silicon intermediates and carbon, which determines the destination of silicon intermediates, namely, either thermodynamically preferred SiC or kinetic product of Si as a final product. Such a finding was applied to design and synthesize the hollow mesoporous shell (ca. 3−4 nm pore) SiC, which is particularly of interest as a catalyst support under harsh environments.

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with carbon.17 For example, Si@skeletal structures provided by carbon hollow shells with void spaces can afford the volume expansion of Si anode materials, enhance the electron transferring efficiency due to high electrical conductivity of carbon, and reduce the instability by which SEI layer forms on the carbon surface instead of Si surface.17 Therefore, designed structures of SiO2/C composites that are pseudomorphically transformed to Si/C materials via magnesiothermic reduction have received considerable attention in the field of LIBs.16,18−24 At the same time, magnesiothermic reduction can also be used to produce metal carbide ceramic materials with controlled structural features, such as SiC from SiO2/C composites.25−29 This carbide material has a very strong covalent bond between Si and carbon, thus exhibiting outstanding mechanical strengths and chemical stabilities. For example, SiC has a hardness of 9 on Mohs scale, which is nearly the same as a diamond,30 and is hardly attacked by concentrated acids and bases, as well as high temperature.31

INTRODUCTION Magnesiothermic reduction is defined as a reaction where vaporized or liquefied magnesium at elevated temperatures reduces various metal oxides (e.g., SiO2, ZrO2, TiO2, and GeO2).1−4 Especially, the reduction of SiO2 to form Si has obtained popularity since Bao et al. reported that the threedimensionally structured Si replicas could be produced from parent SiO2 diatom by magnesiothermic reduction.5 Magnesiothermic reduction has been widely employed in many applications due to the powerfulness of the reduction process that produces pseudomorphically transformed Si structures.6−11 Among them, researchers working on Li-ion batteries (LIBs) have shown tremendous enthusiasm in employing magnesiothermic reduction because Si is known to have the largest theoretical capacity (3580 mAh g−1 at room temperature), making it promising as an anode material of LIBs.12−16 When looking inside the characteristics of Si as an anode material, it has critical problems such as volume expansion during cycling, low intrinsic electrical conductivity, and continuous formation of solid electrolyte interphase (SEI) layer. Hence, Si is frequently used as a form of hybrid structure © XXXX American Chemical Society

Received: December 31, 2015 Revised: February 1, 2016

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suggested. Silicon intermediates, mainly in situ-formed Mg2Si, have a crucial role for determining the resultant crystalline phase: when the silicon intermediates encounter carbon, SiC is formed; otherwise, Si is formed. The proof-of-concept was achieved by a designed reaction between Mg2Si and carbon nanopshere. In addition, we have tried to determine the synthesis parameters that affect the formation of either Si or SiC. First, the effect of different contact areas of SiO2 with carbon on the formation of different crystalline phases under the same experimental conditions is elucidated. For this purpose, three different types of SiO2/C composites were designed. These include SiO2 within mesoporous carbon with around 2 nm pore size, SiO2 inside hollow carbon shell, and SiO2 particles on the chunk of carbon. The results from the three different models revealed that the contact area influences the crystalline phases of the products. Furthermore, the resultant crystalline phases also depend on the reaction conditions such as the reaction temperatures, heating rates, and amount of the reactants used. The underlying reasons of variation in the crystalline phase could be accounted for by a solid-state diffusion limited process that produces thermodynamically more stable SiC and its dependence on the reaction temperature. Moreover, material preparation based on the understandings of the reaction mechanism is demonstrated. Hollow mesoporous SiC nanoparticles could be synthesized since the formation of SiC occurs only in the restricted regions where sufficient contact between SiO2 and carbon is attained.

In addition, magnesiothermic reduction is a much less energy consuming process in the production of SiC from SiO2/C composites as it is usually performed around 500 °C or above, compared to the conventional carbothermal reaction that requires at least 1200 °C.32 In addition, SiC is a semiconductor with a wide bandgap.33 Therefore, formation of SiC with controlled morphology and porosity from designed SiO2/C composites via magnesiothermic reduction is attractive in areas of electronic devices and catalyst supports under harsh environments.34 The reason for the wide use of magnesiothermic reduction in the synthesis of Si/C and SiC is that the field of the SiO2/C composites has already been matured as they are used as intermediate materials in the preparation of SiO2 templated carbon replicas. For example, carbon molecular sieves of CMK series, reported by Ryoo and co-workers, were synthesized via nanocasting of carbonaceous material in the parent mesoporous SiO2 with pore sizes less than 10 nm.35,36 Zhao’s group had employed triconstituent coassembly of ordered mesostructured SiO2/C composites to form mesoporous carbon.36 Also, three dimensionally ordered mesoporous carbon (3DOmC) was successfully synthesized through nanocasting of carbon materials into a close-packed SiO2 nanoparticle array with sizes ranging from 10 to 40 nm.37,38 In addition, hollow carbon sphere was also prepared through an introduction of carbon shell onto a SiO2 nanoparticle template and seubsequent etching of SiO2 template.39 Thus, various types of SiO2/C composites with diverse sizes can readily be designed and fabricated for synthesizing carbon replica using SiO2 preform. On reviewing the previous studies on the synthesis of Si/C or SiC materials through magnesiothermic reduction of SiO2/C composites, the fact that compositionally equivalent parent SiO2/C materials could be used to synthesize both Si/C materials and SiC materials cannot be easily rationalized. A representative case of this problem is given as follows. When initial utilization of magnesiothermic reduction method using SiO2/C composites was conducted by Stucky and co-workers, the resultant crystalline phase was of SiC instead of Si/C composite.27 A triconsituent mixture of tetraethylorthosilicate (TEOS), phenol-formaldehyde (PF) resin and F127 nonionic surfactant, and close-packed polystyrene sphere array were used to prepare three dimensionally ordered macroporous SiO2/C framework with mesopores on the wall, which was pseudomorphically transformed to macro-/mesoporous SiC. Meanwhile, when a similar approach to prepare SiO2/C composites using triconsituent coassembly of TEOS, PF resin and F127 was employed using the magnesiothermic reduction by the Zhao’s group, small Si nanoparticles of sizes around 3 nm embedded inside carbon framework were successfully produced, which were used as anode materials for LIBs.18 Unfortunately, no proper explanation is available, which comprehensively addresses the differences in the systems that yield Si or SiC from the compositionally identical parent system. Only limited explanations on the respective cases of Si or SiC formation were provided.25,40−42 With respect to the previous studies in the literature, it is noteworthy that the size of SiO2 within carbon framework, in other words, the contact area of SiO2 with the carbon framework, and a few other experimental conditions seem crucial for the determination of the final product (either Si or SiC). In this study, a possible formation mechanism of Si and SiC via magnesiothermic reduction of SiO2/C composites is

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EXPERIMENTAL SECTION

Materials and Reagents. Aluminum(III) chloride hexahydrate (AlCl3·6H2O, 99%), phenol (≥99%), paraformaldehyde (95%), resorcinol (98%), oxalic acid (98%), pluronic P123, magnesium silicide (≥99%), and aqueous formaldehyde solutions (37 wt %) were purchased from Sigma-Aldrich. TEOS (≥96%) was purchased from TCI Chemicals. Sodium hydroxide (NaOH), 28−30% aqueous solutions of ammonia (NH4OH), and 38−40% concentrated HCl solutions were purchased from Samchun Chemical (Korea). Absolute ethanol was purchased from J. T. Baker. Magnesium powder (100− 200 mesh) was purchased from Alfa-Aesar. L-lysine (99%) was purchased from Acros Organics. Cetyltrimethylammonium bromide (CTAB, 99%) was purchased from Daejung Chemicals (Korea). Furfuryl alcohol was purchased from JUNSEI (Japan). Deionized (DI) water purified with a Milli-Q purification system was used for all the experiments Reaction between Mg2Si and Carbon Nanosphere. Purchased pristine Mg2Si and synthetic carbon nanospheres were used. In order to synthesize carbon nanosphere, resorcinol (0.76 g), formaldehyde (0.88 g), and aqueous ammonia solution (0.7 g) were dissolved in 200 mL of DI water.43 The solution was allowed to react with stirring at room temperature for 24 h, then aged in 90 °C oven for 24 h without stirring. Colloidal products were collected and washed by repeated centrifugation. Finally, collected products were carbonized at 800 °C in a N2 atmosphere for 3 h and carbon nanospheres of the size of 328 ± 27 nm were finally synthesized. To perform the reaction between Mg2Si and carbon nanospheres, 76 mg of Mg2Si and 24 mg of carbon nanospheres were mixed together using mortar and pestle. Pelletized mixture was then annealed at 750 °C for 5 h with heating rate of 10 °C/min in Ar atmosphere. To remove remaining Mg2Si and formed Si, the result product was immersed in 1 M HCl for 5 h. After purification, it was immersed in 2 M NaOH at 75 °C for 16 h. Synthesis of Mesoporous SiO2 Containing Carbon (mSiO2− C). In order to prepare mSiO2−C, mesoporous SiO2 nanoparticles were initially synthesized.44 Subsequently, 420 g of DI water, 200 g of absolute ethanol, 2.22 g of CTAB, 6 g of aqueous ammonia solution, and 4.08 g of TEOS were allowed to react for 16 h to produce mesoporous SiO2 nanoparticles. Colloidal product of the mixture was washed by repeated centrifugation. The product was dried at 80 °C B

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Chemistry of Materials under reduced pressure and then calcined at 600 °C for 3 h in a muffle furnace to remove the remaining CTAB in the mesopores. Mesopores of the SiO2 template were filled with phenol-paraformaldehyde resin by a gas-phase polymerization. In a cosolvent of 20 mL of ethanol and 20 mL of DI water, 2 g of SiO2 template was mixed with 2 g of aluminum(III) chloride hexahydrate. Under sonication of the mixture, Al ions were adsorbed on the pore wall. After drying at 80 °C under reduced pressure, the product was calcined at 700 °C for 3 h ina muffle furnace. Then, 200 mg of calcined product, 95 mg of phenol, and 100 mg of paraformaldehyde were sealed in glass container. The container was then heated at 100 °C for 12 h in an oven to allow gas-phase polymerization. Finally, it was carbonized at 800 °C for 3 h in Ar atmosphere. Synthesis of SiO2@mSiO2−C and SiO2 Encapsulated in Hollow Carbon Shell (SiO2@void@C). Synthesis of SiO2@void@ C was performed by a previously reported procedure.20 Briefly, 200 nm SiO2 nanoparticles were synthesized by the Stöber process. Then, SiO2 nanoparticles were encapsulated by mesoporous shells. Three grams of core SiO2 nanoparticle, 108 g of ethanol, 225 g of DI water, 1.11 g of CTAB, 3 g of aqueous ammonia solution, and 2 g of TEOS were used for the encapsulation reaction. The products were purified by centrifugation and calcined at 600 °C for 3 h. Then, the pores of the mesoporous SiO2 were filled with phenol-formaldehyde resin by a gas phase polymerization using a procedure similar to that of mSiO2− C (SiO2@mSiO2−C indicates the material at this point). In order to alleviate sintering of materials during carbonization step, a thin SiO2 layer (3−5 nm thick) was coated on SiO2@mSiO2 containing carbon source.6 Then, it was carbonized at 800 °C for 3 h in Ar atmosphere. Finally, SiO2 was partially etched in NaOH solution. The products were reannealed at 800 °C for 1 h in Ar atmosphere to remove dangling bonds formed on the carbon shell and SiO2 surface. Preparation of SiO2−C Mix. SiO2 nanoparticles and carbon chunks were separately prepared and mixed to obtain SiO2−C composites. Then, 200 nm SiO2 nanoparticles were synthesized by Stöber process. The colloidal products were purified and calcined at 600 °C for 3 h. The carbon chunks were prepared by the following procedure. Two grams of resorcinol and 2 mL of 37% formaldehyde were mixed., and 0.3 mL of aqueous ammonia solution was injected to the mixture. The resulting solid products were purified by repeated filtration and then carbonized at 800 °C for 3 h. Preparation of SBA15-C. Initially, SBA15 was prepared using a block copolymer templating method. Four grams of pluronic P123, 30 g of DI water, and 120 g of 2 M aqueous HCl were mixed. After complete dissolution of P123 polymer, 8.4 g of TEOS was injected to the reaction mixture. After stirring at 35 °C for 24 h, the mixture was heated at 100 °C for another 24 h without stirring. As-prepared solid product was washed with ethanol three times by centrifugation and then calcined at 700 °C for 3 h in a muffle furnace. One gram of prepared SBA15 was mixed with AlCl3·6H2O in a cosolvent of 20 mL of ethanol and 20 mL of DI water by sonication for 1 h. Solid products were collected by centrifugation and dried at 80 °C under reduced pressure, and the product was calcined at 700 °C for 3 h in a muffle furnace. Then, 200 mg of calcined product, 95 mg of phenol, and 100 mg of paraformaldehyde were sealed in glass container, and the container was heated at 100 °C for 12 h in an oven to allow gas-phase polymerization. Finally, it was carbonized at 800 °C for 3 h in Ar atmosphere. Preparation of SiO2@3DOmC Composites. SiO2 nanoparticles of size 33 nm were prepared using a modified Stöber method using a basic amino acid (L-lysine) as catalyst, following previous reports.37,38 Briefly, around 15 nm sized SiO2 nanoparticle seeds were first synthesized. Eighty milligrams of L-lysine was dissolved in 70 g of DI water. Then, 5.3 g of TEOS was added to the solution with vigorous stirring under 90 °C for 2 days. In order to obtain 33 nm sized SiO2 nanoparticles, seeded growth using the 15 nm SiO2 nanoparticles was employed. Then, 21.2 g of TEOS was added to the 15 nm seed solution, and the synthesis was allowed to continue at 90 °C for 2 days. After filtration, the solution mixture was dried at 80 °C overnight in a vibration free environment to obtain closely packed SiO2 nanoparticles, which was then calcined at 500 °C for 6 h to remove

L-lysine

molecules on the SiO2 surface. Mixture of furfuryl alcohol (12 g) and oxalic acid (0.06 g) as a carbon source was infiltrated into the close-packed SiO2 particles at least two times in order to complete the infiltration. Furfuryl alcohol infiltrated close-packed SiO2 particles were transferred to an oven kept at 80 °C for 24 h for polymerization. The polymer/close-packed SiO2 composites were then carbonized at 800 °C for 3 h, resulting in the composites of close-packed SiO2@three dimensionally ordered mesoporous carbon (SiO2@3DOmC). Typical Reduction Procedure. Fifty milligrams of SiO2/C composite and 31 mg of magnesium were ground together. The molar ratio of magnesium was around 2.1:1. The mixture was placed in an alumina crucible, and the crucible was covered by another alumina crucible. Meanwhile, for the experiments with respect to the amount of reactants mixture, 10, 25, or 50 mg of mixture composed of SiO2@ void@C and Mg where the molar ratio between SiO2 and Mg was fixed as 2.1:1 was used. Also, for the experiments regarding the limit specific contact area for formation of Si crystal, 13 piles of mixture composed of SiO2/C composite (SBA15-C or SiO2@3DOmC) and Mg where the molar ratio between SiO2 and Mg was fixed as 2.1:1 was placed in the crucible. The each pile of mixture was 3 mg and separated from each other with distance at least 2 cm. Then, it was annealed in a horizontal tube furnace at a programmed temperature (600, 650, or 750 °C) for 5 h under Ar flow. The heating rate was 10 °C/min unless otherwise specified. Rietveld Quantification. The Rietveld refinement calculation was performed by the well-known Rietveld program named as “Fullprof” with an interface provided by the software “Match!”.45 The XRD patterns obtained from the HCl treated samples of reduced SiO2@ void@C at various reduction conditions were used for the refinement. Referred crystalline phase information on Si (Crystallography Open Database (COD) number: 2102763) and SiC (COD number: 1010995) were provided as initial models. Refined scale factors of each crystalline phase were used for obtaining weight ratio between Si and SiC. Preparation of Hollow Mesoporous Shell SiC Nanoparticle (HMS-SiC). Two hundred milligrams of SiO2@mSiO-C (synthetic method of SiO2@mSiO2−C is described above) was mixed with 134 mg of magnesium. The mixture was annealed at 650 °C for 5 h in Ar atmosphere with heating rate of 10 °C/min. The result mixture was treated with 1 M HCl for 5 h, and then it was washed by repeated centrifugation. Additionally, the washed product was immersed in 2 M NaOH at 75 °C for 16 h, followed by washing through repeated centrifugation. Finally, residual carbon was removed by calcination at 600 °C for 3 h in muffle furnace. Characterizations. Transmission Electron Microscope (TEM) images were acquired by a Hitachi-7600 system (Hitachi, Japan). Scanning Electron Microscope (SEM) images were recorded using a Hitachi S-4300 instrument (Hitachi, Japan). High-Resolution TEM (HR-TEM) images were obtained with a JEM-2100F (JEOL). The thermogravimetric analysis (TGA) was performed using an SDT Q600 device (TA Instruments Inc.), XRD measurement was performed by D8 Advance (Bruker), with Cu radiation (λ = 1.5406 Å) at 40 kV and 40 mA. 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 static vacuum (p < 10−5 bar) 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 obtained by the Barrett−Joyner− Halenda (BJH) method applied to the adsorption branches of all samples.

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RESULTS AND DISCUSSION Understanding of Reaction System from Literature. Prior to discussing about the synthesis parameters that affect the final crystalline phases, we present the thermodynamics of the magnesiothermic reduction of SiO2/C composite in order to understand the fundamentals of the reaction.46 When the formation reactions of Si and SiC are assumed to occur as in reaction (1) and (2), the Gibbs free energy changes as a C

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Chemistry of Materials function of the temperature could be calculated (Figure S1, see Supporting Information for details). SiO2 (s) + 2Mg(g) + C(s) → Si(s) + C(s) + 2MgO(s) (1)

SiO2 (s) + 2Mg(g) + C(s) → SiC(s) + 2MgO(s)

(2)

As depicted in Figure S1, the free energy change for the formation of SiC is lower than that of Si in the entire temperature range. In other words, SiC formation is thermodynamically preferred, implying that kinetic considerations should account for the formation of Si in this system. In order to find the kinetic aspect influential to the formation of Si or SiC, the proposed formation mechanisms of SiC during the magnesiothermic reduction of SiO2/C composite were surveyed in the literature. It was noted that direct reaction between produced Si and carbon is hardly considered as the formation reaction of SiC, because it is known that the direct reaction requires much higher reaction temperature (at least 1100 °C) to be initiated than ordinary temperatures used by magnesiothermic reduction.47 Therefore, it is widely accepted that intermediate chemical species should participate in the formation of SiC. One of suggested intermediates is magnesium silicide (Mg2Si). Nguyen et al. reported that crystalline phases of Mg 2 Si, tetragonal MgC 2 /SiC, and cubic SiC were sequentially observed during magnesiothermic reduction of chiral nematic SiO2/C composite.42 It implies that SiO2 is initially reduced by magnesium to form Mg2Si then reacts with carbon. In addition, the direction of mass transfer is from SiO2 to carbon because the morphology of SiC product resembles the parent carbon structure.26,29 As a consequence, based on the SiC formation mechanism, it can be supposed that the silicon intermediates (mainly in situ-formed Mg2Si) should encounter carbon by a solid state diffusion process to form SiC. The diffusion process that takes place in the solid state is commonly thought to be a kinetically limiting process, so SiC formation could strongly depend on the diffusion of silicon intermediates. At the same time, according to numerous reports, Mg2Si has been considered as intermediates of Si during magnesiothermic reduction of SiO2.40,41 On the basis of these facts, one can imagine that the kinetic product of Si could be formed if conversion from SiO2 to Si is completed before the silicon intermediates reach carbon. In other words, Si could be synthesized if insufficient diffusion of Si intermediates to carbon is prompted. Schematic representation of the suggested mechanism of Si and SiC formation by magnesiothermic reduction of SiO2/C composite is depicted in Figure 1. In order to prove the aforementioned discussion, reaction between silicon intermediate (Mg2Si) and carbon nanosphere were carried out to find the evidence of SiC formation. Pelletized mixture of prestine Mg2Si and carbon nanospheres (328 ± 27 nm, Figure 2a, see more information at experimental section) was annealed at 750 °C for 5 h in Ar flowing tube furnace. Then, the product was treated by HCl and NaOH to remove remaining Mg2Si and the side product of Si. Monitored by TEM analysis, it was clearly observed that crystalline fringes of (111) planes of SiC was identified on the surface of carbon nanospheres after the reaction (Figure 2b and 2c). Also, the formation of SiC crystal was observed in the XRD pattern (Figure 2d) and the size of the crystal with around 7.3 nm determined by the Scherrer equation is well matched with TEM investigation, strongly evidencing the role of Mg2Si that can

Figure 1. Schematic representation of Si and SiC formation mechanism by magnesiothermic reduction of SiO2/C composite.

Figure 2. (a) TEM image of carbon nanosphere; (b,c) TEM image (0.25 nm refers to the interlayer spacing of SiC (111) plane); and (d) XRD pattern of HCl and NaOH treated sample from the reaction between Mg2Si and carbon nanoparticle at 750 °C.

generate SiC when encounting carbon. On the other hand, it is noted that the peak located around 43° could be assigned to MgC2 which was also reported by Nguyen et al.42 Variation in Crystalline Phases: Effect of Contact Area between SiO2 and Carbon. In the previous section, we proposed that encounter between silicon intermediates and carbon results in the formation of SiC, otherwise Si can be formed. It implies that formation of Si and SiC in magnesiothermic reduction of SiO2/C composite could be controlled if the encounter between silicon intermediates and carbon is manipulated. In order to confirm this presumption experimentally, we first manipulated the actual contact area of SiO2 and carbon. If the contact area is large enough for all carbon surfaces to face the intermediates effectively, SiC formation occurs vigorously. However, if the actual contact is quite small, encounter of the intermediates and carbon could be significantly restricted because most of the silicon intermediates could be located much farther from the carbon surface than the length of diffusion at a given condition. Therefore, low SiO2/C contact materials could increase the chances of Si formation. As model parent SiO2/C composites, three types of SiO2/C composites with different SiO2/C contact areas (mSiO2−C, D

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Chemistry of Materials SiO2@void@C, and SiO2−C mix) were prepared (Figure 3). For mSiO2−C, carbon was confined inside mesoporous SiO2

The SiO2/C models were then used as reactants for magnesiothermic reduction. In a typical reaction, SiO2/C was mixed with magnesium, and annealed at 650 °C for 5 h in Ar flow with a heating rate of 10 °C/min. The magnesium to SiO2 molar ratio was kept as 2.1:1, nearly stoichiometric. XRD was employed to analyze the products (Figure 4). The XRD

Figure 4. XRD patterns of samples reduced at 650 °C from SiO2/C composites with different contact areas (heating rate: 10 °C/min, reactant weight: 81 mg).

patterns clearly show that the resultant crystalline structures are different for the products obtained from different contact areas of the parent SiO2/C composites. Solely Si and SiC products were obtained from SiO2−C mix and mSiO2−C, respectively, but a mixture of Si and SiC was prepared from SiO2@void@C. These observations strongly support the assumption that the crystalline phase of the products in magnesiothermic reduction of SiO2/C composite is strongly related to the solid state diffusion of the silicon intermediates to encounter carbon. In the case of mSiO2−C with the highest contact surface area, silicon intermediates could meet most of carbon atoms at a very short diffusion length, resulting in thermodynamically preferred SiC formation. On the other hand, for SiO2@void@C and SiO2−C mix, there was a part of the silicon intermediates could not encounter carbon due to the diffusion limits at the given conditions. Thereby, the part of silicon intermediates that are away from the carbon surface could be converted to Si. To avoid confusion, it should be emphasized that the low contact material, i.e., SiO2−C mix, does not totally prevent the formation of SiC, thereby forming both Si and SiC. Indeed, SiC was observed in the reduced product of the SiO2−C mix (Figure 5a). The reason why the SiC feature was not observed in the XRD pattern should be attributed to the relative amounts of the two crystalline phases. In addition, the evidence supporting the silicon intermediates are the main subjects of diffusion were observed. According to the TEM images of HCl-treated products obtained from reduced SiO2@void@C (Figure S6), the shape of carbon shell remains intact, while the SiO2 particle inside the shell is totally deformed. Furthermore, SiC in the carbon shell part and Si in the core part are observed under TEM measurements (Figure 5b). Therefore, it is likely that the silicon intermediates are diffused into the carbon shell to form SiC. Gao et al. has already reported a similar observation.26,29

Figure 3. Graphic representations and TEM images of (a) mSiO2−C, (b) SiO2@void@C, and (c) SiO2−C mix (inset: magnified image of the marginal region).

nanoparticles. The pore size of the mesoporous SiO 2 nanoparticle (mSiO2) was mostly 2 nm as determined by the BJH method using nitrogen sorption isotherm (Figure S2). Energy-dispersive X-ray (EDX) mapping of mSiO2−C revealed that the carbon is well-dispersed in the entire mSiO2−C particle (Figure S3). SiO2@void@C has a yolk−shell structure where hollow carbon shell encapsulates the SiO2 nanoparticle. SiO2− C mix is a simple mixture of SiO2 nanoparticles and carbon chunks with sizes over tens of micrometers (Figure S4). mSiO2−C should have the largest contact between SiO2 and carbon, and SiO2−C mix should have the smallest. SiO2@ void@C was thought to have the midrange contact compared to the other models. We have adopted a criterion that numerically expresses the contact area between SiO2 and carbon, which is in the dimension of contact surface area per SiO2 mass. The contact areas of our models were estimated by performing simple calculations with assumptions or by nitrogen sorption experiment (see Supporting Information for details). The contact areas of mSiO2−C, SiO2@void@C, and SiO2−C mix were estimated to be 1.6 × 103, 1.5 × 101, 8.9 × 10−2 m2/g, respectively. It is noteworthy that the surface areas of the composites differed by at least 2 orders of magnitude; therefore, the model SiO2/C composites were appropriately designed and prepared. Meanwhile, atomic ratios between SiO2 and carbon in the three models were similarly controlled in order to exclude the stoichiometric issues. The atomic ratios of SiO2:C were determined to be around 1:1.8 using TGA as depicted in Figure S5. E

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Thermodynamic preference of the crystalline phase of the product in this system cannot be changed by temperature, as seen in the calculation in Figure S1. However, the diffusion could be affected by the reaction temperature. The dependence of solid-state diffusion on temperature could be expressed with the following equations. The diffusion coefficient (D) of solid state follows the Arrhenius equation.48 D = D0 ·e−Ea / kT

where D0, Ea, k, T are the maximum diffusion coefficient, activation energy, Boltzmann constant, absolute temperature, respectively. Mean diffusion length (L) of the diffusion species in the solid state is given as follows.48 L = (D·t )0.5

In the equation above, t is the lifetime of diffusing chemical species. These equations imply that the temperature change affects the diffusion coefficient, thereby changing the diffusion length of the silicon intermediates. Therefore, control of the temperature would help in controlling the resultant crystalline phases. The SiO2/C models were annealed at two different temperatures of 600 and 750 °C, while the other parameters were not manipulated compared to the experiments performed at 650 °C. The result product was analyzed by XRD (Figure 6). In the case of mSiO2−C and SiO2@void@C reduction at 600 °C, the peak identification of the product was not clear, so the XRD patterns of the corresponding 1 M HCl-treated samples are presented in Figure S7. Some important points could be noted from the XRD patterns in Figure 6. SiO2@void@C exhibits a temperature-dependent phase generation behavior. At 750 °C, the relative peak intensity of SiC to Si is increased compared to the case at 650 °C. On the other hand, at 600 °C, only a peak corresponding to Si was observed. In order to quantify the amounts of Si and SiC obtained at different reaction temperatures from SiO2@void@C, the Rietveld refinement and TGA were employed (Table 1 and Figure S8). The Rietveld refinement was performed on the XRD patterns of HCl-treated samples (the acid treatment was used for MgO removal). In addition, TGA data were collected for the samples treated with HCl and NaOH in order to

Figure 5. HR-TEM images of HCl treated samples of reduced (a) SiO2−C mix, (b) SiO2@void@C, and (c) mSiO2−C mix at 650 °C (heating rate: 10 °C/min, reactant weight: 81 mg, 0.25 and 0.31 nm refer to the interlayer spacing of SiC (111) plane and Si (111) plane, respectively).

Variation in Crystalline Phases: Effect of Reaction Temperature. In the previous section, we observed that final crystalline phases could be varied if the chance of encounter between silicon intermediates and carbon is manipulated by controlling the SiO2/C contact. In this section, we found that reaction temperature also can be influential to the final crystalline phases.

Figure 6. XRD patterns of as-reduced products from (a) mSiO2−C, (b) SiO2@void@C, and (c) SiO2−C mix reduced at various reaction temperatures with a heating rate of 10 °C/min and reactant weight of 81 mg. F

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However, there was an unsolved question that the dominant formation of Si from the SiO2/C composites with two digits of contact areas from the literature was performed with the temperature over 650 °C, at which considerable amount of SiC was also obtained with major product of Si in this study as seen in Figure 6b. Therefore, other parameters could be involved in determining the crystalline phases during magnesiothermic reduction. Variation in Crystalline Phases: Effect of Heating Rate. We investigated the influences of the heating rate on our reaction system. SiO2@void@C was chosen as a model system due to its variation in crystalline phases of the products on changing the reduction temperatures. The heating rates of 1 °C/min and 20 °C/min were employed to compare with the result of 10 °C/min heating rate, which was already verified. According to XRD (Figure 7a), the formation of Si and SiC is

Table 1. Quantitative Analysis Results for the Reduced SiO2@void@C at Various Reaction Conditions quantification results (atomic %)

reduction parameters temp (°C)

ramp rate (°C/min)

weight of reactant (mg)

750 650 600 650

10

81

20 10 1 20

81

650

50 25 10

Rietveld

TGA

Si

SiC

Si

SiC

43.5 60.9 94.6 17.8 60.9 94.3 3.5 26.6 83.4

56.5 39.1 5.4 82.2 39.1 5.7 96.5 73.4 16.6

39.0 67.4 91.3 9.9 67.4 90.4 8.6 28.0 81.8

61.0 32.6 8.7 90.1 32.6 9.6 91.4 72.0 18.2

remove MgO and Si, respectively. TGA provided the ratio between carbon and SiC, which could be eventually converted to the ratio of Si and SiC by a simple calculation (see Supporting Information for more details). As seen in the Table 1, SiC was mostly produced at higher reaction temperatures, whereas the amount of Si was increased at lower reaction temperatures. Both the analyses by the Rietveld method and TGA are found to be in good agreement. These observations are as expected from the equations; higher temperature enhances the diffusion length of the silicon intermediates, resulting in the formation of SiC as the major product. In other words, part of SiO2, which could not participate in the formation of SiC at lower temperatures, could take part in the SiC formation at higher temperatures. On the other hand, diffusion at lower temperature was restricted, and this suppressed the formation of SiC, while reduction of SiO2 to Si by magnesium was still possible. Therefore, only Si could be observed when the experiment was performed at 600 °C. Additionally, it can be noticed that the observable results were not affected by temperature if the contact area between SiO2 and carbon was either very small (1.6 × 103 m2/g, in the case of mSiO2−C) at all reaction temperatures (Figure 6a,c). mSiO2−C was a limiting case for large contact areas because the thickness of SiO2 layer between carbon mesopores was supposed to be around several nanometers, so that the diffusion length at lower temperature was still enough to make all of intermediates to react with carbon. Meanwhile, SiO2−C mix seemed to be an oppositely limiting case of small contact; therefore, the amount of the produced ratio of SiC is always much smaller than Si in spite of an increased SiC formation at higher temperatures. Consequently, SiC was not observed under XRD investigations. The validity of our assumption was examined with the results reported in the literature.18,19,21−24,26−29,42 Estimations of the contact areas between SiO2 and carbon and the resulting crystalline phases of the products obtained from various studies are listed in Table S1. According to the data in Table S1, the contact area and the resultant crystalline phase are related. For example, parent SiO2/C composites which were dominantly reduced to SiC usually have contact areas over 102 m2/g, whereas Si was mainly produced when the contact areas were around 10° ∼ 101 m2/g except in one case (entry 8 in Table S1). This clearly shows that the crystalline phase of the product is definitely dependent on the contact areas.

Figure 7. XRD patterns of reduced SiO2@void@C at 650 °C with a heating rate of 1 °C/min, 10 °C/min, and 20 °C/min (81 mg of mixture).

affected by the heating rate. Higher heating rate induces an increased SiC formation, whereas, under the lower heating rate conditions, less SiC is formed. Interestingly, the peak of SiC obtained at a heating rate of 20 °C/min exhibits higher intensities than those of the samples obtained at 750 °C with heating rate of 10 °C/min (Figure 6b). This implies that the effects of heating rate could be more crucial than the reaction temperature. Quantification of Si and SiC using the Rietveld refinement and TGA was carried out (Table 1 and Figure S9). Atomic percentages of SiC were 82.2% and 90.1% from the Rietveld refinement and TGA, respectively, for a sample prepared at a heating rate of 20 °C/min, whereas the sample prepared with 1 °C/min heating rate showed atomic percentages of SiC as 5.7% (the Rietveld refinement) and 9.6% (TGA), which strongly supported the findings from XRD measurements. Influence of heating rate on the results of magnesiothermic reduction has been already presented by the Cui group.49 They reported that the nanostructure of silicon reduced from rice husk can be modified by heating rate. It was suggested that heating rate can change the local accumulation of heat released by magnesiothermic reduction (ΔH = −593 kJ/molSi for reaction (1)), thereby higher heating rate resulted in more coarsoned silicon products. Likewise, the effect of heating rate G

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explore the limit in terms of the contact area between SiO2 and carbon, in which formation of crystalline Si is allowed. Hence, a model system that has different contact areas of SiO2/C composites including mSiO2−C, SBA15-C, and SiO2@ 3DOmC (Figure 9) is proposed to find the limit of contact

on our system can be also rationalized by adopting the heat accumulation effect modified by heating rate. The heat arising from the reaction cannot be efficiently dissipated under a high heating rate because sufficient time is not allowed, resulting in a higher local temperature and the formation of thermodynamically favored product (SiC). On the other hand, low heating rate allows enough heat-dissipation time and the local temperature is not largely elevated. Therefore, the actual reaction temperature can vary according to the heating rate, resulting in different crystalline products. Variation in Crystalline Phases: Effect of the Weight of Reactants. We also found that the crystalline phase of the products of magnesiothermic reduction could be affected by the amount of the reactants used. When analyzing the products obtained with different amounts of reactant mixtures (i.e., 10, 25, and 50 mg) under the same heating rate condition of 20 °C/min by XRD, different crystalline phases were identified (Figure 8). Large amounts of the reactant mixture resulted in

Figure 9. Graphic representations and TEM images of (a) SBA15-C (inset: HR-TEM image) and (b) SiO2@3DOmC.

area, which can provide Si from magnesiothermic reduction under specific conditions. First, SBA15-C was synthesized with carbon confined in the 2D hexagonally arrayed SiO2 mesopores of around 6 nm (Figure 9a). The contact area was measured as 5.0 × 102 m2/g from nitrogen sorption experiment (Figure S11). The molar ratio between SiO2 and carbon was 1:1.8 from TGA plot (Figure S11e). In addition, SiO2@3DOmC is a composite of closely packed SiO2 nanoparticles with the interparticle spaces filled with carbon (Figure 9b). The contact area of these materials, measured by nitrogen sorption, are 1.7 × 102 m2/g for SiO2@3DOmC. Carbon content of SiO2@ 3DOmC was slightly lower than other models (Figure S11e), but carbon is still excess to the SiO2 (SiO2:C = 1:1.2). On the basis of the findings that restrict the formation of SiC, magnesiothermic reductions of the model composites were performed at 600 °C using a heating rate of 1 °C/min and 3 mg of reactant mixture. According to the XRD patterns of the asreduced product and HCl treated sample of mSiO2−C (Figure 10), SiC is hardly observed. However, Raman signal corresponding to Si−C bonding at around 800 cm−1 was observed (Figure S12), suggesting that SiO2 and carbon reacted with each other at the interface to form SiC, even though no detectable peak was observed in XRD. The reason why XRD peaks of SiC was not detected seemed that the reaction conditions do not allow SiC crystallite to have enough crystallinity to be detected in XRD. Meanwhile, XRD peaks corresponding to crystalline Si were also not observed, implying that the encounter between the silicon intermediates and carbon readily occurs due to the large contact area (1.6 × 103 m2/g) between SiO2 and carbon for the mSiO2−C sample. In addition, crystalline Si peak was also not observed in the XRD patterns of the product of SBA15-C and Raman signal of SiC was detected as similar to mSiO2−C (Figure 10 andFigure S12). Whereas, crystalline Si peaks were observed for the product from SiO2@3DOmC sample (Figure 11). Therefore, in conclusion SiO2/C composites with contact areas lower than

Figure 8. XRD patterns of reduced SiO2@void@C at 650 °C for 5 h with a reactant weight of 10, 25, and 50 mg experiments and (heating rate: 20 °C/min).

the dominant formation of SiC, but the lower amounts of the mixture resulted in formation of less SiC in the product. In addition, quantification by the Rietveld refinement and TGA was performed (Figure S10, Table 1). As observed in the XRD, mostly SiC (96.5% from the Rietveld refinement and 91.4% by the TGA) was identified when 50 mg of the reactants were used; on the other hand, mostly Si (83.4% from the Rietveld refinement and 81.8% by the TGA) was formed 10 mg of weight of the reactants. It could be suggested that the larger heat release from the larger amount of reactant resulted in higher elevation of local temperature, thereby more formation of SiC could occur. Limit of Contact Area of SiO2/C Composite for Formation of Si Crystal. Up to now, we have examined the parameters that affect the crystalline phases of the products from magnesiothermic reduction using various kinds of SiO2/C composites. It has been revealed that the parameters such as the low contact area, low temperature, low heating rate, and low reactant weight, all of which were associated with the chance of encounter between silicon intermediates and carbon, restricted the formation of SiC, resulting in an increase of the relative amounts of Si. On the basis of these findings, we tried to H

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Figure 11. (a) Graphic representation of SiO2@mSiO2−C, (b) XRD patterns of reduced SiO2@mSiO2−C at each preparation step, (c) TEM image (inset: magnified image), (d) HR-TEM image (0.25 nm refers to the interlayer spacing of SiC (111) plane, inset: SAED pattern), (e) TGA plot, and (f) BJH pore distribution of HMS-SiC.

Figure 10. XRD patterns of (a) as-prepared products reduced at 600 °C for 5 h with heating rate of 1 °C/min and 3 mg reactant weight for mSiO2−C, SBA15-C, and SiO2@3DOmC, (b) HCl-treated materials.

1.7 × 102 m2/g could provide Si crystal at the given reaction conditions. In addition, it should be mentioned that 29.7% of SiC compared to Si (70.3%) was detected for reduced SiO2@ 3DOmC by TGA quantification method (Figure S13). Preparation of Hollow Mesoporous Shell SiC Nanoparticle (HMS-SiC). In this section, we demonstrate how these understandings on magnesiothermic reduction of SiO2/C composites could be utilized to prepare functional materials. As an example, hollow mesoporous shell SiC nanoparticle (HMS-SiC) was prepared by using SiO2@mSiO2−C as a parent SiO2/C composite (Figure 11a). SiO2@mSiO2−C is formed in the synthetic process of SiO2@void@C before the partial etching of SiO2. The core of SiO2@mSiO2−C constitutes nonporous SiO2, but the part of shell is composed of mesoporous SiO2 with carbonaceous materials present inside the mesopores (Figure 11a and Figure S14). Because it has been determined that SiC is formed only in the confined regions of SiO 2 /C mesopores, we expect SiC to be pseudomorphically formed in the mesoporous SiO 2/C composite shell while the inside part could be converted to Si, which can be selectively removed. First, SiO2@mSiO2−C was reduced at 650 °C with a heating rate of 10 °C/min, and subsequent HCl and NaOH treatments were performed in order to remove MgO, Si, and the remaining SiO2. The final product formed was SiC, as identified from the XRD patterns (Figure 11b). Calcination at 600 °C for 3 h was performed to remove carbon residues (ca. 3.8 wt % of free carbon was detected by TGA, Figure S15). No significant weight loss was observed in TGA after calcination (Figure 11e), indicating the thermal stability of the HMS-SiC sample.

The morphology of HMS-SiC after calcination was confirmed by TEM analysis (Figure 11c,d). HMS-SiC showed a hollow shell structure with a void core as expected. Crystallites with d-spacing corresponding to SiC (111) were observed in the shell part, confirming the successful pseudomorphic transformation of SiO2/C composite shell into the SiC shell (Figure 11d). The selected area diffraction pattern (SAED) clearly shows the ring patterns of SiC, implying the polycrystalline feature of HMS-SiC. In addition, nitrogen sorption experiment revealed that HMS-SiC showed a BET surface area of 223 m2/g, and most of pore sizes were in the range of ∼3−4 nm in the BJH pore size distribution, strongly indicating a successful transformation of mesoporosity from the parent SiO2/C composite (Figure 11f). In addition, it is worthwhile to mention that the thickness of the carbon shell and diameter of the hollow core can be manipulated by adjusting the corresponding dimensions of the parent material with already established methods.20 Therefore, this procedure could be utilized to synthesize HMS-SiC with controlled thickness of shell and hollow core diameter.

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CONCLUSIONS It is suggested that silicon intermediates (mainly in situ-formed Mg2Si) show a pivotal role for determining resultantant crystalline phases during magnesiothermic reduction of SiO2/ C composites. When in situ-formed Mg2Si readily faces carbon, SiC is preferentially formed at the interface between silica and the carbon surface. This hypothesis is supported by a designed reaction of Mg2Si and carbon nanosphere that resulted in SiC I

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ACKNOWLEDGMENTS This work was supported by the Basic Science Research Program through the National Research Foundation of Korea (2014R1A1A2057204).

formation on the carbon surface. In contrast, Si is considered as a kinetic product which can be usually formed away from the carbon surface because the encounter between silicon intermediates and carbon surface is more restricted when the distance between silica and carbon surface becomes longer. The formation mechnism suggested above implies that the relative formation ratio between Si and SiC can be controlled by manipulating the chance of encounter between silicon intermediates and carbon. Indeed, some synthesis parameters affect the resultant crystalline phases of products in magnesiothermic reduction of SiO2/C composites. It is elucidated that the contact area between SiO2 and carbon is the prime factor in deciding the resultant crystalline phase. It is revealed that the thermodynamically unfavored product of Si instead of SiC can be formed during magnesiothermic reduction, if the contact area is below a certain value. Also, it was found that lower contact area results in more formation of Si. In addition, the crystalline phase of the product could be controlled by altering reduction temperatures when the SiO2/C composite (herein, SiO2@void@C sample) with a medium contact area was used. Furthermore, faster heating rate caused a much higher actual temperature compared to the programmed temperature, due to the acceleration of self-heating of the reaction system; hence, the ratio between Si and SiC in the product could be adjusted. The amount of the reactant mixtures was also found to be an influential factor because the amount of heat released was varied with the amount of the reactants used, which affects the local temperature. All of synthesis parameters strongly affect the chance of encounter between silicon intermediates and carbon, which regulates the formation of thermodynamically favored product of SiC. When controlling the chance of encounter using these synthetic parameters, the kinetic product of Si could be synthesized. Finally, based on the understanding of the system, as an example, pseudomorphically transformed hollow mesoporous shell SiC (HMS-SiC) was successfully fabricated using SiO2@ mSiO2−C as a parent composite, which could be of great interest as a catalyst support under harsh environments. Our findings supported by the reaction parameters could provide a reliable guidance for designing magnesiothermic reaction system.

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REFERENCES

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.chemmater.5b05037. Details of the calculations, additional SEM images, TEM images, XRD patterns, the Rietveld refinement data, Raman spectra, and nitrogen sorption experiment data (PDF)

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AUTHOR INFORMATION

Corresponding Author

*(W.C.Y.) E-mail: [email protected]. Tel.: +82-31-4005504. Notes

The authors declare no competing financial interest. J

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