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Characterization of silicon carbide ceramics obtained from porous carbon structure achieved by plant carbonization
Materials Chemistry and Physics 245 (2020) 122768
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Characterization of silicon carbide ceramics obtained from porous carbon structure achieved by plant carbonization Vladimir Dodevski a, *, Maja C. Pagnacco b, Ivana Radovi�c a, Milena Rosi�c a, Bojan Jankovi�c c, Marija Stojmenovi�c a, Vojislav V. Miti�c d, e a
University of Belgrade, Institute of Nuclear Sciences “Vin�ca”, Laboratory for Materials Sciences, Mike Petrovi�ca Alasa 12-14, P.O. Box 522, 11001, Belgrade, Serbia University of Belgrade, Institute of Chemistry, Technology and Metallurgy, Center for Catalysis and Chemical Engineering, Belgrade, Serbia University of Belgrade, Institute of Nuclear Sciences “Vin�ca”, Department of Physical Chemistry, Mike Petrovi�ca Alasa 12-14, P.O. Box 522, 11001, Belgrade, Serbia d Faculty of Electronic Engineering, University of Ni�s, Aleksandra Medvedeva 14, Ni�s, Serbia e Institute of Technical Sciences of SASA, Belgrade, Serbia b c
H I G H L I G H T S
� Formation of bio-SiC from biomass feedstock precursor by high temperature processing. � β-SiC with cubic structure and the small amount of hexagonal forms were identified. � Detection of polytype inclusions during nano-cubic β-SiC growth by TEOS treatments. � The formations of SiC polycrystals and nanowires become competitive. � SiC nanowires growth dependent from gaseous phase compositions and reaction condition. A R T I C L E I N F O
A B S T R A C T
Keywords: Wood-fruit precursor Silica impregnation Graphitic material Stacking faults β-SiC nano-wires
The aim of this research was to obtain a carbon solid residue by the carbonization process of biomass in an inert atmosphere which, through physical activation and chemical treatment (using TEOS - tetraethyl orthosilicate) would allow creation of highly porous and spatially distinct ordered bio-SiC ceramics. The results of carbon ization experiments at several operating temperatures and activation of carbons with multiple-cycle treatments TEOS clearly showed the possibility of obtaining SiC nano-structures, after performing the carbothermal reduction at 1400 � C. The increase in the activation temperature and the duration time starts the development of the SiC particles inside the porous structure. The XRPD analysis showed that the major SiC polytype has cubic SiC (β-SiC) structure and remainder is hexagonal SiC polytypic (α-SiC) structure. It was established that the carbons obtained from carbonization of the Platanus orientalis L. plane tree fruit (PTF) precursor and activated at 850 � C with longer holding times (1 and 2 h) exhibit β-SiC (cubic) nano-wires. A possible nano-wires increment mechanism was suggested. The obtained results represent significant contribution in understanding the process as well as the main characteristics of SiC nano-materials and their possible applications.
1. Introduction
accordingly, SiC is also contemplated for in vivo membrane applications [7]. Porous ceramics/composites derived from wood (biomorphic ce ramics), sawdust and paper have received particular attention in the past decade [8–11]. Bio-SiC and SiC micro-fiber rods are most promising for medical applications, especially for orthopedic implants, composite SiC–C electrodes, etc. Synthesized microporous SiC ceramics are also
Physical properties, along with improved hemo-compatibility [1–3] make SiC ceramics appropriate for titanium-based hip and dentistry implants [4,5]. Experiments in vitro and in vivo [6] demonstrate excel lent biocompatibility of the SiC ceramic fabricated from wood owing to the interrelated hierarchical porous composition of material;
* Corresponding author. E-mail addresses: [email protected] (V. Dodevski), [email protected] (M.C. Pagnacco), [email protected] (I. Radovi�c), [email protected] (M. Rosi�c), [email protected] (B. Jankovi�c), [email protected] (M. Stojmenovi�c), [email protected] (V.V. Miti�c). https://doi.org/10.1016/j.matchemphys.2020.122768 Received 23 July 2019; Received in revised form 11 December 2019; Accepted 4 February 2020 Available online 6 February 2020 0254-0584/© 2020 Published by Elsevier B.V.
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considered for filtering aggressive liquids and gases. The distinction in the characteristics of SiC and bio-SiC are minimal because the biomorphic properties of bio-SiC mimic the performance of SiC. Although bio-SiC behaves as a ceramic material, its microstructure and properties depend on the raw material used for the synthesis [12]. The similar structure of bio-SiC to SiC allows bio-SiC to retain important properties as high conductivity and other wide band gap (WBG) classi fications. Bio-SiC is more porous and the density of the material is also affected by this porosity as soon as the total hardness and density of bio-SiC are reduced [13]. However, insulation and thermal resistance are increased due to porosity, which is desirable for the electronic ap plications of bio-SiC. Materials of variable porosity derived from pyrolyzed wood or wood parts are employed as carbon matrices for the production of SiC ce ramics through silicon impregnation, hardness, and porosity control. Systematic investigations of impregnation process, a minutely appre hension of SiC ceramics synthesis, and the influence of synthesis con ditions on the physical characteristics of the end product are required to produce materials with predefined properties. The high pyrolyzing temperature (above 600 � C) of wood results in decomposition of the poly-aromatic constituents forming the carbon pattern with an original cellular structure. This structure is suitable for the synthesis of SiC with highly anisotropic porosity, using tetraethyl orthosilicate (TEOS) and the subsequent high-temperature carbothermal reduction. The synthesis procedure of SiC with high surface area modified in order to get higher purities of SiC was developed earlier using sol-gel synthesis from an organosilicon precursor [14]. Some earlier methods led to the formation of excess carbon. Namely, phenyl-trimethoxysilane (PTMS) provides carbon by decomposing the phenyl group, which then bonds with a Si atom to form SiC. However, the phenyl group leads to excess carbon and, thus, an extra source of Si is required to match the Si/C ratio. Hence, tetraethyl orthosilicate (TEOS) was added in order to obtain the exact Si/C ratio. TEOS does not furnish any carbon atoms to be bonded with silicon, but it only provides Si atoms to the gel, and has a significant role in a one-step sol-gel synthesis of porous flow-through carbon-silica monoliths [15]. The aim of this study was to obtain highly porous bio-SiC ceramics, originating from a wood precursor (plane tree seeds (PTS), namely fruit of Platanus orientalis L.), and based on the activated carbon (AC) tem plates (carbon pre-forms), using TEOS as a pore generator. Fouriertransform-infrared (FTIR) spectroscopy, X-ray powder diffraction (XRPD) analysis, Raman spectroscopy (RS) analysis, as well as scanning electron microscopy (SEM) were performed for the characterization of bio-SiC ceramics. The characteristics of derived bio-SiC ceramics were discussed.
42781 Haan, Germany). Particle size fractions within the range of 0.5–1.5 mm were used for the carbonization process. The carbonization was set up in a quartz horizontal tubular reactor with open plate pellets (Protherm Furnaces, model PTF 16/38/250, Turkey). During the process, purified nitrogen at a flow rate of 500 cm3/ min 1 was used as purge gas. The temperature in the reactor was increased from room temperature to operating temperature of 850 � C. Heating rate was constant and maintained at 5 � C min 1. When oper ating temperature reached the desired value, about 20 g of powder was placed into the reactor and held there for 60 min. At the end of the process, the gaseous flow of N2 in the reactor was maintained until it cooled down to room temperature.
2. Experimental
2.4. Treatment of AC templates with TEOS (tetraethyl orthosilicate) and synthesis of SiC
2.3. Activation process In the same horizontal furnace, the activation of the carbonized PTS was performed at fixed temperatures, i.e. T ¼ 650, 750 and 850 � C for 0.3, 1 and 2 h, under 0.5 L min 1 flow of CO2. Active carbons obtained by this procedure were denoted as AC. Since carbons obtained at 650 � C and 750 � C for 0.3 and 1 h did not show the representative results (samples have not been activated), their results were not taken into consideration. The following activated carbons were included in the analysis: 750 � C/2 h (AC-1), 850 � C/0.3 h (AC-2), 850 � C/1 h (AC-3), and 850 � C/2 h (AC-4). During the course of the activation process, first, nitrogen was allowed to flow through the reactor and then the CO2 was released. To maintain the equal flow rate and the pressure of both gases, the gas mixture was allowed to transit through the mixing chamber. The flow of nitrogen was then blocked and only the pure CO2 was allowed to flow through the reactor. Sorption processes are carried out through the adsorption/desorp tion isotherms measurements defined by the adsorption/desorption of nitrogen at 77 K using a Sorptomatic 1990 Thermo Finnigan device. Before measuring, the samples were degassed under the vacuum pres sure [16]. The results indicate that the isotherms correspond with type I (in the low pressure area) and type II (in the high pressure area), with rather high specific surface areas, ranging from 573 to 965 m2 g 1 [16]. The results show that the specific surface area follows an increasing sequence: 750 � C (2 h) < 850 � C (0.3 h) < 850 � C (1 h). The activation process for 1 h in CO2 at 850 � C (850 � C/1 h) demonstrated the devel opment of a highly specific surface area of 965 m2 g 1 with a relatively high specific surface as a result of micro-pore formation [16]. The development of both micro- and meso porosity was the consequence of an increase in the activation temperature from 750 to 850 � C due to the faster C–CO2 reaction at a higher temperature [16]. About 1 g of each obtained AC was taken for further analysis.
2.1. Material
The AC pre forms impregnated with silica were implemented using organosilane such as tetraethyl orthosilicate (TEOS) as a silicon (Si) source with the following characteristics: (Si(OC2H5)4, 98%, M ¼ 208.33 g mol 1, liquid state, ρ ¼ 0.933 g cm 3 at 20 � C, Tb (boiling point) ¼ 168 � C, Tm (melting point) ¼ 85.5 � C) (Merck, Darmstadt, Germany). Each individually activated sample (about 1 g) was post on the filter paper which was then placed on the Büchner funnel connected to a vacuum pump. Filtration was carried out using 50 ml of TEOS (poured onto the funnel). For each sample, pro-filtration was performed three times using the same volume of TEOS. After this procedure, the sample was placed in the furnace with the programmed heating mode of 850 � C for 2 h and at the heating rate 5 � C min 1, under nitrogen atmosphere. The filtration was performed to “enrich” the sample with TEOS, since the selected reagent represents, not only a favorable generator of pores with the larger dimensions, but also serves to increase the activated carbon material due to a higher porosity.
The material used in this study was collected from individual plane tree specimens (Platanus orientalis L.) growing in parks around Belgrade, Serbia. The plane tree seed size of fruit was about 2–3 cm in a diameter. Achenes with its thin fiber bristles were used for the experiment. The dimensions of bristles were approximately 1 cm in length and 1 mm in diameter. 2.2. Carbonization process Raw plane tree seed (PTS) was first cleaned and washed with water to remove dirt and impurities and then it was dried at T ¼ 80 � C for 24 h in furnace (Carbolite Gero GmbH & Co. KG, Hesselbachstraβe 15, 75242 Neuhausen, Germany) to decrease the moisture content. Dried PTS is then milled (using Planetary Micro Mill Pulverisette 7, Fritsch, Indus triestrasse 8 55743 Idar-Oberstein, Germany) and sieved (Vibratory Sieve Shaker AS 200 basic (0.5–1.5 mesh), Retsch, Retsch-Allee 1–5, 2
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Also, the additional thermal treatment is very important in ensuring that the final material is purely SiC rather than a mixture of SiC/C. All steps in the procedure described above were cyclically performed three times for all of the AC samples, 750 � C/2 h (AC-1), 850 � C/0.3 h (AC-2), 850 � C/1 h (AC-3), and 850 � C/2 h (AC-4). After the impregnation process of activated carbons (samples: AC-1, AC-2, AC-3 and AC-4) using TEOS, the synthesis of SiC was performed thermally at high temperature. Approximately half the mass (~300 mg) of the AC sample that was treated with TEOS (each sample separately) was set into the furnace in an inert argon atmosphere at operating temperature of 1400 � C, at the heating rate of 3 � C min 1, where the sample was held for 2 h. The in fluence of preparation conditions, including the reduction in tempera ture, holding time, and heating rate, is of significant importance in processing bio-SiC samples. Thus, for example, the electrical resistance of SiC varies considerably when sintered in a nitrogen atmosphere, and the rate of SiC production can, thus, be obtained based on the processing retention time and temperature. Furthermore, due to the large difference between the coefficients of SiC thermal expansion and the pure Si (that of SiC being higher), the effect of fast heating rates on the homogeneity of mechanical properties of the carbonized layer is more pronounced than the slower heating rates. Our hypothesis was that a faster rate, especially at higher tem peratures, may induce cracks in SiC layers, which may result in a partially open or not completely coalesced SiC layer that would leave the room for silicon to out diffuse from the substrate into the gas phase. Due to these facts, for the purpose of this study, the lowest heating rates were used. Fig. 1 summarizes the whole processing scheme.
2.5.2. XRD analysis The composition of phase and crystal structure of bio-SiC samples was recorded at room temperature by X-ray diffractometer -Rigaku (Rigaku International Corporation, Tokyo, Japan) Ultima IV with nickel filtered CuKα radiation and the step-scan mode (2θ - range: 10� - 80� in a continuous step with a scan mode width of 0.02� and 0.5� min 1). Structural analysis was performed using the program Powder Cell [17]. The TCH pseudo-Voigt profile function gave the best fit to the experi mental data.
2.5. Characterization techniques
3. Results and discussion
2.5.1. FTIR spectroscopy analysis The FTIR spectra of high temperature processing AC pre forms were collected for the purpose of obtaining the bio-SiC ceramics using a PerkinElmer Spectrum Two FTIR spectrometer (PerkinElmer, Inc., Waltham, USA) in a transmission mode. Using a KBr tablets (1:100) compression technique, the measurement pattern was prepared. In the range of 4000 to 400 cm 1, the spectrum was recorded, with the reso lution mode of 4 cm 1, in order to obtain the spectrum lines.
3.1. FTIR results
2.5.3. Raman spectroscopy (RS) analysis The Raman spectra of the raw (PTS) material and impregnated/ thermally treated AC pre forms samples were recorded on a DXR Raman microscope (Thermo Scientific, USA) equipped with an Olympus optical microscope and a CCD detector, with a diode-pumped solid-state highbrightness laser (532 nm) and a 10� objective. The powdered sample that was measured was put on a X–Y motorized sample stage. The power of the laser was 1 mW. Using the spectrograph with a grille of 900 lines mm 1, a bulk light analysis was performed. The observed Raman spectra were within the range of 200 cm 1 to 3500 cm 1. 2.5.4. SEM analysis The micrographs of derived bio-SiC ceramics were analyzed using the SEM JEOL, JSM 5800LV operated at 5 kV. The SEM analysis was implemented in order to obtain information about changes in porosity and anisotropic shrinkage of the porous macro- and micro structures of thermally treated AC performs which undergo impregnation with silica using TEOS.
FTIR analysis was performed in order to investigate thermal stability of previously estimated ACs under various carbonization temperatures through different holding times and subsequent formation of bio-SiC after high temperature treatment. After the break down of the –C–C– chains in the bio-polymer structures, an aromatic poly-nuclear carbon structure starts to form at temperatures above 600 � C [18]. The major mechanism [19] can be presented by the following steps: a) desorption of adsorbed water up to
Fig. 1. Processing scheme of manufacturing bio-SiC ceramics from plane tree seeds (PTS) (Platanus orientalis). 3
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150 � C; b) splitting of wood structural water between 150 � C and 260 � C; c) chain scissions, or depolymerization, and breaking of C–O and C–C bonds within ring units evolving water, CO and CO2 between 260 � C and 400 � C; d) aromatization forming graphitic layers above 400 � C and e) above 800 � C, thermally induced decomposition and rearrangement reactions. Three major pseudo-components of wood, hemicelluloses, cellulose, and lignin, break down in a stepwise manner at 200–280 � C, 260–350 � C and 280–500 � C, respectively [20,21]. Between 260 � C and 400 � C almost 80% of total mass loss occurs, due to the evolution of H2O, CO2 and volatile hydrocarbon species from the fragmentation reactions of poly-aromatic constituents through wood with an open-pore channel system [22]. Fig. 2 shows the curves of FTIR AC pre-forms 750 � C/2 h (AC-1), 850 � C/20 min (AC-2), 850 � C/1 h (AC-3) and 850 � C/2 h (AC-4) at 1400 � C, for derivation of bio-SiC ceramics. Stretching of –C–O–C appeared at about 1117 cm 1 [23,24]. The surface groups from carbon, that was previously activated using the 750 � C/2 h (AC-1) sample, were almost completely decomposed after the high-temperature treatment at 1400 � C. The 850 � C/2 h (AC-4) sample showed the highest thermal stability, which can be identified by the presence of high-intensity peaks. The most prominent peak for this series of the spectra has appeared at 800 cm 1 and it was assigned to the Si–C bond [25,26], thus, indicating that the procedure is suitable for the production of bio-SiC, heated at 1400 � C in an argon atmosphere. The existence of Si–C fundamental streching vibration at ~800 cm 1 clearly suggests that high-temperature bio-SiC processing has gone to completion. However, an intense broader band located at about 800/825 cm 1 was observed and characterized by Si–C fundamental stretching vibra tion. The broad band which was located between 825 and 898 cm 1 belongs to SiC prepared from a carbon/silica hydrogels [27]. However, according to other research reports [28,29], the main band was located between 789 and 794 cm 1. The discrepancy between these values may be attributed to difference in morphology of analyzed SiC and impreg nation procedures [30]. With respect to the above-presented FTIR analysis results, proof of the benefits of activated carbon obtained through a pyrolytic process at temperatures above 600 � C for the production of high-performance bioceramics can be viewed in Table 1 which clearly shows the relevance of obtained FTIR results (Fig. 2). Table 1 illustrates positions of infrared bands characteristic for some organosilicon groups.
Table 1 The positions of infrared bands which are characteristic of some organosilicon groups [31]. Group
Wave number (cm 1)
Additional comments
Si–O–Si
1130–1000
Si–H in amorphous silicon CERAMICS Silicon carbide
2150–2000
In the region of 1130–1000 cm 1 siloxanes determine very strong infrared bands. Single Si–O–Si band was shown by disiloxanes and small-ring cyclo-siloxanes. Amorphous silicon (a-Si) has been found in the range of 2150–2000 cm 1 Si–H bands at three or more frequencies. Silicon carbide occurs in a bewildering number of crystal modifications. It also occurs as amorphous material (a-SiC). Detailed SPECTRA-STUCTURE CORRELATIONS for these forms are not available. It can be expected that different forms will have small differences in the infrared spectra, but all these forms of SiC show strong absorption at or near 800 cm 1.
ca. 800
3.2. XRD analysis of studied samples The XRD diffraction patterns of the bio-SiC ceramics obtained for 750 � C/2 h (AC-1)-a, 850 � C/20 min (AC-2)-b, 850 � C/1 h (AC-3)-c, and 850 � C/2 h (AC-4)-d after carbothermal reduction at 1400 � C are pre sented in Fig. 3. The peak at ~33.55� corresponds to (1010) planes of hexagonal SiC, while peaks at ~ 26 and 43� correspond to (002) and (111) planes of the graphite, where all other peaks originate from cubic β-SiC. The peaks marked by Miller indices (111), (200), (220), and (311) on diffracto grams indicate the crystalline structure of cubic silicon carbide. The position of observed peaks indicated the presence of β-SiC (cubic space group F-43 m, No. 216, 2θ ¼ 35.650� , d111 ¼ 2.516). The d (interplanar spacing) values of four peaks are 2.516, 2.179, 1.541, and 1.314 Å, respectively. The peak detected at 2θ ¼ 33.554� , d ¼ 2.668 Å, belongs to hexagonal α-SiC (space group P 63mc, No. 186) and was characteristic of the collision error on the (111) planes in the β-SiC phase [32,33]. The crystal defects, such as stacking faults, are observed which altered the steps of sequence (111) planes to yield an additional peak at 2θ ~ 33.55� . The weak, broadly rounded peaks of amorphous carbon at ~26 and 43� are visible in all XRD patterns except in the last one (d) which is sometimes associated with the sample-processing temperature, whereby the lower temperature means a smaller angle. The obtained diffracto grams also indicated partial graphitization of the material with residual
Fig. 2. Thermally treated AC pre-forms 750 � C/2 h (AC-1), 850 � C/20 min h (AC-2), 850 � C/1 h (AC-3) and 850 � C/2 h (AC-4) after the impregnation process using TEOS at 1400 � C, in order to obtain the bio-SiC ceramics.
Fig. 3. The XRD pattern of the bio-SiC samples produced at 1400 � C. 4
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amorphous carbon and match to the (002) and (111) planes of graphite [34,35]. All peaks are broad and with low intensities which differ slightly with changes in time interval, indicating nano-crystalline nature of powders. With the increase in holding time, the intensities of XRD peaks which belong to β-SiC were moderately increased. At the same time, the XRD peaks β-SiC phase became more intense. The diffracto grams of the samples (a) and (b) are different from the other two based on the peaks with lower intensity, indicating a gradual transformation from a micro-to nano-particle sizes that can be observed in SEM images (see the later results). Based on the d-values and 2θ angles from the ICSD database, card number #28389 (the cubic β-SiC) and card number #164970 (the hexagonal α-SiC), a very good agreement was obtained. The XRD pat terns indicate that the carbon template is completely converted into SiC, at a sufficiently high processing temperature. The bio-SiC synthesis occurring through a reduction reaction is practically accomplished and no essential amount of the crystalline SiO2 phase can be detected [34]. Thus, for the products obtained at 1400 � C (Fig. 3), the peaks of the major cubic type β-SiC phase appear. The diffraction peaks of crystalline SiO2 (cristobalite) were not observed which indicate that nearly the whole amount of silica has been consumed during carbothermal reduction at 1400 � C. This means that the whole amount of silica was mineralized. During quantitative treatment, it reacted quantitatively. It is evident that, after carbothermal reduction at 1400 � C, the β-SiC phase, as a major phase with cubic structure and a small amount of the hexagonal structure, can be identified. If we consider the synthetic approach presented here, then, it follows that the bio-SiC predominantly consists of the cubic structure but a small amount of the volume fraction of hexagonal structure is always present under applied processing con ditions. Results from Fig. 3 show that partial transformation, hexagonal [1 h 850 � C] → cubic [2 h 850 � C], took place during the synthesis. Ortiz et al. [36] found that defects such as twins and mismatch errors always exist in the cubic SiC phase. The β to α transformation is a kinetically controlled process and, often, several days are needed to achieve the complete conversion [37], where α to β is preferably governed by diffusion, with lower barrier energy and higher conversion rates. In transformation from α to β structures, it can be assumed that the diffusion-controlled process (growth) can happen in this ceramic because of the presence of long range rearrangement process of the atoms in a silicon matrix in order to form precipitates. In the SiC powder product, which was synthesized at 1400 � C during 2 h from the carbon perform activated at 850 � C, the nucleation of hexagonal SiC structure has just started but, obviously, high activated temperature and high reduction processing temperature are probably quite enough to induce fast conversion where trans formation from the cubic to the hexagonal form is very small. Namely, β-structure is more stable at lower temperatures and transforms irre versibly to the α-form at very high temperature (~2000 � C). However, the products obtained at nearly 2000 � C consist of α-SiC [38].
Fig. 4. a) - d) Raman spectra of 750 � C/2 h (AC-1), 850 � C/20 min (AC-2), 850 � C/1 h (AC-3), and 850 � C/2 h (AC-4) samples, respectively, which are treated with TEOS and further annealed at 1400 � C for 2 h.
consisting of a sequence of Si–C double layers stacked on top of each other. The order of agreement of double layers is called “hexagonal” (α-SiC) if bonds in adjacent double layers are in an eclipsed orientation; while, for staggered configuration, the order of agreement is called “cubic” (β-SiC (3C–SiC)) [40]. β-SiC has characteristic peaks at around 796 cm 1 and 850 cm 1, corresponding to the zone center transversal optical-phonon mode TO, as well as at 950 cm 1, corresponding to the longitudinal optical phonon LO [39,41,42]. For α-SiC two additional folded phonon modes are typical TO at around 767 cm 1 and the most intense TO mode at 789 cm 1, being identified as α-SiC inclusions in β-SiC [41]. The Raman peaks originating from Si–C vibration appear only in the samples 850 � C/1 h (AC-3) and 850 � C/2 h (AC-4) treated by TEOS at 1400 � C, as shown in Fig. 4 and enlarged in Fig. 5. The obtained bio-SiC Raman peaks for samples 850 � C/1 h (AC-3) and 850 � C/2 h (AC-4), demonstrate detection and quantification of the polytype inclusions during the growth of the nano-cubic β-SiC. Namely, the peak obtained for sample 850 � C/1 h (AC-3) at around 788 cm 1 could be assigned to the intense peak of the α-SiC TO mode, which is used for identification of the α-SiC polytype inclusion in β-SiC or, in other words, the transformation of cubic SiC to a hexagonal structure; while peak at 917 cm 1 could be assigned to the LO phonon peaks of cubic SiC (Fig. 5a). The obvious shift obtained (33 cm 1 for the LO β-SiC bulk) is an indication of the nano-wired morphology of SiC [42,43]. Furthermore, the Raman spectrum related to the 850 � C/2 h (AC-4) sample shows three peaks at 787 cm 1, 853 cm 1, and 948 cm 1 indi cating the extra “combined” SiC morphology: the β-SiC with the pres ence of nano-wires, accompaning β-SiC with the small amount transformed to the hexagonal SiC (less then in the 850 � C/1 h (AC-3) sample) [41,43]. Most intense peak exists around 853 cm 1, originating from β-SiC [44], and Fig. 5b might be considered as a representation of better crystallization. Raman spectra prove the transformation of β-SiC to α-SiC for the 850 � C/1 h (AC-3) sample as well as the existence of
3.3. Raman spectroscopy The Raman spectra of activated carbon (AC) samples 750 � C/2 h (AC1), 850 � C/20 min (AC-2), 850 � C/1 h (AC-3), and 850 � C/2 h (AC-4) treated with TEOS and further annealed at 1400 � C for 2 h are presented in Fig. 4 a-d, respectively. The four Raman spectra look very similar, except for the region from 600 cm 1 to 1100 cm 1 obtained from carbon samples activated at 850 � C for 1 h and 2 h. The 750 � C/2 h (AC-1) and 850 � C/20 min (AC-2) samples show two main peaks at around 1350 cm 1 and 1580 cm 1, plus two small peaks at around 2690 cm 1 and 2930 cm 1. The samples 850 � C/1 h (AC-3) and 850 � C/2 h (AC-4) except all previously mentioned peaks, exhibited additional small peaks around 788 cm 1, 853 cm 1, 917 cm 1, and 940 cm 1. The vibration range from 760 cm 1 to 950 cm 1 is characteristic for the Si–C bond [39]. SiC crystallizes in polytype crystal structures, all 5
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could be an indication that a number of graphene layers (or presented carbon clusters) decreases in investigated samples (Fig. 4). It is well known that the Raman spectral shape of a carbon material drastically changes not only due to the kind of abundant allotropic forms of carbon but to the fine structural transformation of the individual allotrope. Thus, Fig. 6 represents Raman spectra of the activated 1 h 850 � C carbon without TEOS treatment and further annealing, and the obtained Raman spectrum for the 850 � C/1 h (AC-3) sample treated with TEOS and annealed at 1400 � C for 2 h. The red-shift and narrowing of the G-peak are obvious. This is an undoubted indication of an increase in ordering of the carbon phase (presented in the form of few layers or an organized carbon cluster) in bio-SiC samples. Due to a better Raman scattering efficiency of C–C bonds compared to Si–C bond, the Raman lines, characteristic for crystalline SiC, could be almost unobservable, especially if SiC is in the form of very small crys tallites [44]. This is the reason why the Raman spectroscopy is unable to detect the peaks of the SiC in the bio-SiC 750 � C/2 h (AC-1) and the 850 � C/20 min (AC-2) samples at 1400 � C. However, it is evident that an increase in the activation temperature and holding time causes growth of the SiC crystallite phase. Such results are in accordance with those obtained through the XRD analysis. Furthermore, the residual carbon which is not transformed to SiC could have many beneficial properties in catalysis, especially in photocatalytic processes, where it can serve as an electron acceptor, facili tating the interface charge separation, etc [44]. The mechanism of SiC formation from the PTS precursor after carbonization and activation processes, and the subsequent carbo thermal reduction, at 1400 � C, in an argon atmosphere, may be explained by the following chemical reactions, such as: TEOS (at elevated temperature on C-surface) → SiO2 (and SiO) þ C (and CO) þ H 2↑ (1) SiO2 þ 2C → SiC þ CO2↑
(2)
SiO þ 2C → SiC þ CO↑ (possible formation of polycrystals)
(3)
SiO þ 3CO→ SiC þ 2CO2↑ (possible formation of nano-wires)
(4)
3C–SiC (β-SiC) from PTS, identified through the characterization analysis in this study, can obviously be used as high-temperature bioceramics. It may be supposed that the two above-indicated reactions can occur (reactions [3,4]) in the conversion process following their competition with one to another.
Fig. 5. a) – b) Raman peaks originated from bio-SiC samples at 850 � C/1 h (AC3) and 850 � C/2 h (AC-4), respectively, treated with TEOS, annealed at 1400 � C for 2 h.
nano-wires, whereas the Raman spectra for the 850 � C/2 h (AC-4) sample consists mainly of a large amount of β-SiC nano-wires, which are also documented in the XRD (see the previously shown results) and SEM measurements (see the later results). Raman analysis also suggests that some amount of free carbon is present in all obtained samples, because of the existence of two intensive bands, i.e. G-band around 1580 cm 1 and D-band around 1350 cm 1, which are generally ascribed to E2g and A1g in-plane vibration modes in polycrystalline graphite, respectively [42,45]. The D-band of carbon or the so-called “disorder band” is ascribed to structural imperfections. The G-band (or graphitic) arises from stretching the C–C bond in graphitic forms and is attributed to all sp2 carbons. The peak at 2690 cm 1 is of a second order overtone of the D-mode (also named 2D) whereas the peak at 2930 cm 1 is a combination of the peak obtained scattering from G and D (G þ D) [45]. The activation time and elevated temperature have more impact on the position of G-band than the D-band position as it may be seen in Fig. 4. The red-shift (from 1591 cm 1 to 1578 cm 1) of the G-band might indicate an increase in the ordering of the carbon phase [45]. The visual inspection of the 2D peak shows no clear asym metry, nor a shoulder structure. The small intense and broad 2D peak (around 2690 cm 1), as observed here, could be an indication of the existence of carbon in the form of few graphene layers or some carbon domains – clusters [45]. Furthermore, the obtained 2D peak red-shift
Fig. 6. The comparison of the Raman spectra obtained from non-TEOS treated and the 850 � C/1 h (AC-3) TEOS treated sample, annealed at 1400 � C. 6
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3.4. SEM image analysis of TEOS infiltration process in AC samples after carbothermal reduction The SEM images of the AC samples which undergo carbothermal reduction at 1400 � C, designated as 750 � C/2 h (AC-1), 850 � C/20 min (AC-2), 850 � C/1 h (AC-3), and 850 � C/2 h (AC-4) are shown in Fig. 7 ad, respectively. The SEM image, presented in Fig. 7a, shows the morphology of 750 � C/2 h (AC-1) sample including examination of the cross-section along the axial direction. At the lowest activation temperature (750 � C for 2 h) and, after three cycles of TEOS soaking, the presence of non-sealed surface pores (mainly, of a round or an elliptic shape) can be identi fied in the carbonized precursor, exactly where a small extent of con version of SiO2 into SiC occurred. Since the quantity of silica in material is smaller than the quantity of carbon, the formation of SiC will be associated with the quantity of SiO2 mineralized on PTS. Furthermore, the whole cellular structures as cells, pits, and lumens mainly remain after thermal treatment. At a higher activation temperature (850 � C for 20 min) (Fig. 7b), the 850 � C/20 min (AC-2) sample that undergoes carbothermal reduction, shows structure approximately resembling that of the honeycomb where large cavities were gradually repeatedly dipped in TEOS. Therefore, SiO2 starts to cover the surface and fills the cavities and pores of tra cheids and pits, suggesting that SiO2 solution penetrated the cell wall structures and condensed around cell-shaped tissues. Upon conversion to SiC, reaction between molten silicon and amorphous carbon to form SiC was accompanied by a change in volume. It can be supposed that the overall sample dimensions did not undergo drastic changes, but the volume change within the pores was accommodated. This can resulted in decreased pore diameters but a certain number of pore walls was damaged, which led to some effectively larger pores. For the 850 � C/1 h (AC-3) sample (Fig. 7c), the porosity decreased and nearly completely disappeared once the material was converted into SiC. However, in the case of the 850 � C/20 min (AC-2) a 20 min (0.3 h) sample, the fact that it was held at an activation temperature of 850 � C was insufficient to completely convert the entire carbon scaffold into SiC. As carbon in smaller pores reacted, the pores were closed due to the volume expansion, impeding further carbon conversion into SiC. Most of the residual carbon was converted into SiC after this reaction step was repeated. The presence of small-diameter pores (Fig. 7a) was identified and they were pinched close (Fig. 7b) due to the volume expansion which occurs during C (at Fig. 7c we may see expanded graphite) → SiC (generation of SiC product - Fig. 7c) reaction. The 850 � C/1 h (AC-3) sample (Fig. 7c) exhibits formation of nano-wires with obviously com bined structural morphology, including the reservation of tubular interpenetrating pores from PTS. In this reaction stage, the SiO2 nanoparticles form droplets in touch with the activated carbon, thus, form ing SiO vapor. The SiO vapor further reacts with carbon to form β-SiC particles, which act as the nuclei for the growth of the nano-wires. Once the droplets are saturated with β-SiC, crystallization of β-SiC occurs, followed by the growth in one direction to produce nano-wires. In the case of 850 � C/2 h (AC-4) (Fig. 7d) amount of SiC nano-wires increases. Both, straight and curved nano-wires grow into a massive network (Fig. 7d), where nano-wires are distinguished by their great length. This behavior was confirmed by FTIR, XRD as well as Raman measurements. Fig. 8 shows the SEM image at higher magnification of the 850 � C/2 h (AC-4) sample. It may be seen in Fig. 8 that a certain number of nano-wires contain spherical caps at the tip of nano-wires, which are more uniform than SiC nano-wires previously synthesized with the vapor-liquid-solids (VLS) mechanism in other researches [46–48]. It is worth noting that no extra metal catalysts were introduced in our experiments. Therefore, we can assume that some metal catalyst would probably accelerate the forma tion of SiC nanowires. It can be assumed that observed spherical caps (Fig. 8) probably consists of silicon, oxygen, and carbon and some of alkali or alkaline earth metals, which may play the catalytic role in SiC
Fig. 7. SEM images of 750 � C/2 h (AC-1) (a), 850 � C/20 min (AC-2) (b), 850 � C/1 h (AC-3) (c), and 850 � C/2 h (AC-4) (d) samples which undergo carbo thermal reduction at 1400 � C for derivation of bio-SiC from activated carbons of the carbonized PTS precursor.
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development of SiC particles inside the porous structure. The X-ray powder diffraction (XRPD) analysis showed that major SiC polytype is the cubic SiC (β-SiC) and remainder is the hexagonal SiC polytypic (α-SiC). It was found that the carbon obtained through carbonizing the PTS precursor and activated at 850 � C during longer holding time (1 and 2 h) in a reducing atmosphere at 1400 � C possesses β-SiC (cubic) nanowires. Besides the XRPD analysis, the formation of β-SiC nano-wires was also confirmed by the Fourier-transform infrared (FTIR) and Raman spectroscopy analyses. The process described in this paper is rather complex for the industrial application, because it involves three pro cessing steps at high temperatures: pyrolysis, activation, and, finally, reduction. However, the described procedure allows the production of ultra-long SiC nano-wires, which have pre-disposition for unique optical properties, favorable thermal stability properties, as well as flexible nano-mechanical properties. Therefore, in regard to these characteris tics, further research will be carried out. Fig. 8. The SEM image of the 850 � C/2 h (AC-4) sample at 1400 � C at a higher magnification.
Acknowledgement This research work was partially supported by the Serbian Ministry of Education, Science and Technological Development under the pro jects numbered 172015, III45005, TR34011 and III45001.
nano-wire formation. From the presented results, it is evident that the vapor-solid (VS) mechanism probably occurs which can be attributed to increased tem perature and partial pressure creating reaction between CO(gas) and Si(solid) or CO(gas) and SiO(gas) [49]. Considering the fact that the (111) plane possesses the lowest surface energy among SiC surfaces, a SiC nano-wire, thus, tends to grow along a [111] direction. The diffraction peaks in XRD patterns at (111), (220), and (311) planes (Fig. 3) confirm that a nano-wire consists of β-SiC phase; the (111) lattice plane of a nano-wire is exposed to the outside for the generation of a branched nano-wire. Growth of SiC nano-wires is highly dependent on the phase of gaseous compositions and reaction conditions. Namely, the phases CO (gas), SiO(gas), SiO2(solid), and SiC(solid) can be formed through the Si–C–O–N system and through a complex scheme of reactions, which requires additional thermodynamic calculations. It should be noted that, within VS mechanism, the height of SiC nano-wires is restricted by re action conditions. Further thermodynamic calculations may help in estimating whether SiC is a stable phase when the partial pressure of CO is lower than 0.25 atm [50]. Presence of hydrogen reaction [1] in the gas phase may favor the formation of SiO from SiO2. SiO reacts with carbon to form SiC. The basic reactions involved in the formation of SiC nano-wires are as follows:
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(5)
SiO2 þ C → SiO þ CO↑
(6)
SiO þ 2C → SiC þ CO↑ 1
Formed β-SiC nano-wires show a strong FTIR band at 800 cm due to a Si–C stretching band and this is confirmed by FTIR (Fig. 2). Raman spectroscopy also shows existence of nano-wires. Inter-connected porous structure from the PTS precursor, after carbonization and during a carbothermal reduction, plays an important role in the SiC synthesis because pores accelerate the passage of gaseous SiO, CO, and CO2. 4. Conclusions Silicon carbide (SiC) ceramics have been synthesized through a carbothermal reduction of activated carbons (AC) obtained by carbon izing Platanus orientalis L (plane tree fruits). The activated samples were “enriched” with TEOS (tetraethyl orthosilicate) in a procedure for the impregnation of silicon particles into carbon templates. The various obtained porous products significantly depend on the activation tem perature and holding time, before the application of carbothermal reduction. The 750 � C/2 h (AC-1) sample synthesis of SiC begins at 1400 � C, but it is not completed after three cycles of TEOS soaking. The in crease in the activation temperature and the holding time initiates the 8
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Materials Chemistry and Physics journal homepage: www.elsevier.com/locate/matchemphys
Characterization of silicon carbide ceramics obtained from porous carbon structure achieved by plant carbonization Vladimir Dodevski a, *, Maja C. Pagnacco b, Ivana Radovi�c a, Milena Rosi�c a, Bojan Jankovi�c c, Marija Stojmenovi�c a, Vojislav V. Miti�c d, e a
University of Belgrade, Institute of Nuclear Sciences “Vin�ca”, Laboratory for Materials Sciences, Mike Petrovi�ca Alasa 12-14, P.O. Box 522, 11001, Belgrade, Serbia University of Belgrade, Institute of Chemistry, Technology and Metallurgy, Center for Catalysis and Chemical Engineering, Belgrade, Serbia University of Belgrade, Institute of Nuclear Sciences “Vin�ca”, Department of Physical Chemistry, Mike Petrovi�ca Alasa 12-14, P.O. Box 522, 11001, Belgrade, Serbia d Faculty of Electronic Engineering, University of Ni�s, Aleksandra Medvedeva 14, Ni�s, Serbia e Institute of Technical Sciences of SASA, Belgrade, Serbia b c
H I G H L I G H T S
� Formation of bio-SiC from biomass feedstock precursor by high temperature processing. � β-SiC with cubic structure and the small amount of hexagonal forms were identified. � Detection of polytype inclusions during nano-cubic β-SiC growth by TEOS treatments. � The formations of SiC polycrystals and nanowires become competitive. � SiC nanowires growth dependent from gaseous phase compositions and reaction condition. A R T I C L E I N F O
A B S T R A C T
Keywords: Wood-fruit precursor Silica impregnation Graphitic material Stacking faults β-SiC nano-wires
The aim of this research was to obtain a carbon solid residue by the carbonization process of biomass in an inert atmosphere which, through physical activation and chemical treatment (using TEOS - tetraethyl orthosilicate) would allow creation of highly porous and spatially distinct ordered bio-SiC ceramics. The results of carbon ization experiments at several operating temperatures and activation of carbons with multiple-cycle treatments TEOS clearly showed the possibility of obtaining SiC nano-structures, after performing the carbothermal reduction at 1400 � C. The increase in the activation temperature and the duration time starts the development of the SiC particles inside the porous structure. The XRPD analysis showed that the major SiC polytype has cubic SiC (β-SiC) structure and remainder is hexagonal SiC polytypic (α-SiC) structure. It was established that the carbons obtained from carbonization of the Platanus orientalis L. plane tree fruit (PTF) precursor and activated at 850 � C with longer holding times (1 and 2 h) exhibit β-SiC (cubic) nano-wires. A possible nano-wires increment mechanism was suggested. The obtained results represent significant contribution in understanding the process as well as the main characteristics of SiC nano-materials and their possible applications.
1. Introduction
accordingly, SiC is also contemplated for in vivo membrane applications [7]. Porous ceramics/composites derived from wood (biomorphic ce ramics), sawdust and paper have received particular attention in the past decade [8–11]. Bio-SiC and SiC micro-fiber rods are most promising for medical applications, especially for orthopedic implants, composite SiC–C electrodes, etc. Synthesized microporous SiC ceramics are also
Physical properties, along with improved hemo-compatibility [1–3] make SiC ceramics appropriate for titanium-based hip and dentistry implants [4,5]. Experiments in vitro and in vivo [6] demonstrate excel lent biocompatibility of the SiC ceramic fabricated from wood owing to the interrelated hierarchical porous composition of material;
* Corresponding author. E-mail addresses: [email protected] (V. Dodevski), [email protected] (M.C. Pagnacco), [email protected] (I. Radovi�c), [email protected] (M. Rosi�c), [email protected] (B. Jankovi�c), [email protected] (M. Stojmenovi�c), [email protected] (V.V. Miti�c). https://doi.org/10.1016/j.matchemphys.2020.122768 Received 23 July 2019; Received in revised form 11 December 2019; Accepted 4 February 2020 Available online 6 February 2020 0254-0584/© 2020 Published by Elsevier B.V.
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considered for filtering aggressive liquids and gases. The distinction in the characteristics of SiC and bio-SiC are minimal because the biomorphic properties of bio-SiC mimic the performance of SiC. Although bio-SiC behaves as a ceramic material, its microstructure and properties depend on the raw material used for the synthesis [12]. The similar structure of bio-SiC to SiC allows bio-SiC to retain important properties as high conductivity and other wide band gap (WBG) classi fications. Bio-SiC is more porous and the density of the material is also affected by this porosity as soon as the total hardness and density of bio-SiC are reduced [13]. However, insulation and thermal resistance are increased due to porosity, which is desirable for the electronic ap plications of bio-SiC. Materials of variable porosity derived from pyrolyzed wood or wood parts are employed as carbon matrices for the production of SiC ce ramics through silicon impregnation, hardness, and porosity control. Systematic investigations of impregnation process, a minutely appre hension of SiC ceramics synthesis, and the influence of synthesis con ditions on the physical characteristics of the end product are required to produce materials with predefined properties. The high pyrolyzing temperature (above 600 � C) of wood results in decomposition of the poly-aromatic constituents forming the carbon pattern with an original cellular structure. This structure is suitable for the synthesis of SiC with highly anisotropic porosity, using tetraethyl orthosilicate (TEOS) and the subsequent high-temperature carbothermal reduction. The synthesis procedure of SiC with high surface area modified in order to get higher purities of SiC was developed earlier using sol-gel synthesis from an organosilicon precursor [14]. Some earlier methods led to the formation of excess carbon. Namely, phenyl-trimethoxysilane (PTMS) provides carbon by decomposing the phenyl group, which then bonds with a Si atom to form SiC. However, the phenyl group leads to excess carbon and, thus, an extra source of Si is required to match the Si/C ratio. Hence, tetraethyl orthosilicate (TEOS) was added in order to obtain the exact Si/C ratio. TEOS does not furnish any carbon atoms to be bonded with silicon, but it only provides Si atoms to the gel, and has a significant role in a one-step sol-gel synthesis of porous flow-through carbon-silica monoliths [15]. The aim of this study was to obtain highly porous bio-SiC ceramics, originating from a wood precursor (plane tree seeds (PTS), namely fruit of Platanus orientalis L.), and based on the activated carbon (AC) tem plates (carbon pre-forms), using TEOS as a pore generator. Fouriertransform-infrared (FTIR) spectroscopy, X-ray powder diffraction (XRPD) analysis, Raman spectroscopy (RS) analysis, as well as scanning electron microscopy (SEM) were performed for the characterization of bio-SiC ceramics. The characteristics of derived bio-SiC ceramics were discussed.
42781 Haan, Germany). Particle size fractions within the range of 0.5–1.5 mm were used for the carbonization process. The carbonization was set up in a quartz horizontal tubular reactor with open plate pellets (Protherm Furnaces, model PTF 16/38/250, Turkey). During the process, purified nitrogen at a flow rate of 500 cm3/ min 1 was used as purge gas. The temperature in the reactor was increased from room temperature to operating temperature of 850 � C. Heating rate was constant and maintained at 5 � C min 1. When oper ating temperature reached the desired value, about 20 g of powder was placed into the reactor and held there for 60 min. At the end of the process, the gaseous flow of N2 in the reactor was maintained until it cooled down to room temperature.
2. Experimental
2.4. Treatment of AC templates with TEOS (tetraethyl orthosilicate) and synthesis of SiC
2.3. Activation process In the same horizontal furnace, the activation of the carbonized PTS was performed at fixed temperatures, i.e. T ¼ 650, 750 and 850 � C for 0.3, 1 and 2 h, under 0.5 L min 1 flow of CO2. Active carbons obtained by this procedure were denoted as AC. Since carbons obtained at 650 � C and 750 � C for 0.3 and 1 h did not show the representative results (samples have not been activated), their results were not taken into consideration. The following activated carbons were included in the analysis: 750 � C/2 h (AC-1), 850 � C/0.3 h (AC-2), 850 � C/1 h (AC-3), and 850 � C/2 h (AC-4). During the course of the activation process, first, nitrogen was allowed to flow through the reactor and then the CO2 was released. To maintain the equal flow rate and the pressure of both gases, the gas mixture was allowed to transit through the mixing chamber. The flow of nitrogen was then blocked and only the pure CO2 was allowed to flow through the reactor. Sorption processes are carried out through the adsorption/desorp tion isotherms measurements defined by the adsorption/desorption of nitrogen at 77 K using a Sorptomatic 1990 Thermo Finnigan device. Before measuring, the samples were degassed under the vacuum pres sure [16]. The results indicate that the isotherms correspond with type I (in the low pressure area) and type II (in the high pressure area), with rather high specific surface areas, ranging from 573 to 965 m2 g 1 [16]. The results show that the specific surface area follows an increasing sequence: 750 � C (2 h) < 850 � C (0.3 h) < 850 � C (1 h). The activation process for 1 h in CO2 at 850 � C (850 � C/1 h) demonstrated the devel opment of a highly specific surface area of 965 m2 g 1 with a relatively high specific surface as a result of micro-pore formation [16]. The development of both micro- and meso porosity was the consequence of an increase in the activation temperature from 750 to 850 � C due to the faster C–CO2 reaction at a higher temperature [16]. About 1 g of each obtained AC was taken for further analysis.
2.1. Material
The AC pre forms impregnated with silica were implemented using organosilane such as tetraethyl orthosilicate (TEOS) as a silicon (Si) source with the following characteristics: (Si(OC2H5)4, 98%, M ¼ 208.33 g mol 1, liquid state, ρ ¼ 0.933 g cm 3 at 20 � C, Tb (boiling point) ¼ 168 � C, Tm (melting point) ¼ 85.5 � C) (Merck, Darmstadt, Germany). Each individually activated sample (about 1 g) was post on the filter paper which was then placed on the Büchner funnel connected to a vacuum pump. Filtration was carried out using 50 ml of TEOS (poured onto the funnel). For each sample, pro-filtration was performed three times using the same volume of TEOS. After this procedure, the sample was placed in the furnace with the programmed heating mode of 850 � C for 2 h and at the heating rate 5 � C min 1, under nitrogen atmosphere. The filtration was performed to “enrich” the sample with TEOS, since the selected reagent represents, not only a favorable generator of pores with the larger dimensions, but also serves to increase the activated carbon material due to a higher porosity.
The material used in this study was collected from individual plane tree specimens (Platanus orientalis L.) growing in parks around Belgrade, Serbia. The plane tree seed size of fruit was about 2–3 cm in a diameter. Achenes with its thin fiber bristles were used for the experiment. The dimensions of bristles were approximately 1 cm in length and 1 mm in diameter. 2.2. Carbonization process Raw plane tree seed (PTS) was first cleaned and washed with water to remove dirt and impurities and then it was dried at T ¼ 80 � C for 24 h in furnace (Carbolite Gero GmbH & Co. KG, Hesselbachstraβe 15, 75242 Neuhausen, Germany) to decrease the moisture content. Dried PTS is then milled (using Planetary Micro Mill Pulverisette 7, Fritsch, Indus triestrasse 8 55743 Idar-Oberstein, Germany) and sieved (Vibratory Sieve Shaker AS 200 basic (0.5–1.5 mesh), Retsch, Retsch-Allee 1–5, 2
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Also, the additional thermal treatment is very important in ensuring that the final material is purely SiC rather than a mixture of SiC/C. All steps in the procedure described above were cyclically performed three times for all of the AC samples, 750 � C/2 h (AC-1), 850 � C/0.3 h (AC-2), 850 � C/1 h (AC-3), and 850 � C/2 h (AC-4). After the impregnation process of activated carbons (samples: AC-1, AC-2, AC-3 and AC-4) using TEOS, the synthesis of SiC was performed thermally at high temperature. Approximately half the mass (~300 mg) of the AC sample that was treated with TEOS (each sample separately) was set into the furnace in an inert argon atmosphere at operating temperature of 1400 � C, at the heating rate of 3 � C min 1, where the sample was held for 2 h. The in fluence of preparation conditions, including the reduction in tempera ture, holding time, and heating rate, is of significant importance in processing bio-SiC samples. Thus, for example, the electrical resistance of SiC varies considerably when sintered in a nitrogen atmosphere, and the rate of SiC production can, thus, be obtained based on the processing retention time and temperature. Furthermore, due to the large difference between the coefficients of SiC thermal expansion and the pure Si (that of SiC being higher), the effect of fast heating rates on the homogeneity of mechanical properties of the carbonized layer is more pronounced than the slower heating rates. Our hypothesis was that a faster rate, especially at higher tem peratures, may induce cracks in SiC layers, which may result in a partially open or not completely coalesced SiC layer that would leave the room for silicon to out diffuse from the substrate into the gas phase. Due to these facts, for the purpose of this study, the lowest heating rates were used. Fig. 1 summarizes the whole processing scheme.
2.5.2. XRD analysis The composition of phase and crystal structure of bio-SiC samples was recorded at room temperature by X-ray diffractometer -Rigaku (Rigaku International Corporation, Tokyo, Japan) Ultima IV with nickel filtered CuKα radiation and the step-scan mode (2θ - range: 10� - 80� in a continuous step with a scan mode width of 0.02� and 0.5� min 1). Structural analysis was performed using the program Powder Cell [17]. The TCH pseudo-Voigt profile function gave the best fit to the experi mental data.
2.5. Characterization techniques
3. Results and discussion
2.5.1. FTIR spectroscopy analysis The FTIR spectra of high temperature processing AC pre forms were collected for the purpose of obtaining the bio-SiC ceramics using a PerkinElmer Spectrum Two FTIR spectrometer (PerkinElmer, Inc., Waltham, USA) in a transmission mode. Using a KBr tablets (1:100) compression technique, the measurement pattern was prepared. In the range of 4000 to 400 cm 1, the spectrum was recorded, with the reso lution mode of 4 cm 1, in order to obtain the spectrum lines.
3.1. FTIR results
2.5.3. Raman spectroscopy (RS) analysis The Raman spectra of the raw (PTS) material and impregnated/ thermally treated AC pre forms samples were recorded on a DXR Raman microscope (Thermo Scientific, USA) equipped with an Olympus optical microscope and a CCD detector, with a diode-pumped solid-state highbrightness laser (532 nm) and a 10� objective. The powdered sample that was measured was put on a X–Y motorized sample stage. The power of the laser was 1 mW. Using the spectrograph with a grille of 900 lines mm 1, a bulk light analysis was performed. The observed Raman spectra were within the range of 200 cm 1 to 3500 cm 1. 2.5.4. SEM analysis The micrographs of derived bio-SiC ceramics were analyzed using the SEM JEOL, JSM 5800LV operated at 5 kV. The SEM analysis was implemented in order to obtain information about changes in porosity and anisotropic shrinkage of the porous macro- and micro structures of thermally treated AC performs which undergo impregnation with silica using TEOS.
FTIR analysis was performed in order to investigate thermal stability of previously estimated ACs under various carbonization temperatures through different holding times and subsequent formation of bio-SiC after high temperature treatment. After the break down of the –C–C– chains in the bio-polymer structures, an aromatic poly-nuclear carbon structure starts to form at temperatures above 600 � C [18]. The major mechanism [19] can be presented by the following steps: a) desorption of adsorbed water up to
Fig. 1. Processing scheme of manufacturing bio-SiC ceramics from plane tree seeds (PTS) (Platanus orientalis). 3
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150 � C; b) splitting of wood structural water between 150 � C and 260 � C; c) chain scissions, or depolymerization, and breaking of C–O and C–C bonds within ring units evolving water, CO and CO2 between 260 � C and 400 � C; d) aromatization forming graphitic layers above 400 � C and e) above 800 � C, thermally induced decomposition and rearrangement reactions. Three major pseudo-components of wood, hemicelluloses, cellulose, and lignin, break down in a stepwise manner at 200–280 � C, 260–350 � C and 280–500 � C, respectively [20,21]. Between 260 � C and 400 � C almost 80% of total mass loss occurs, due to the evolution of H2O, CO2 and volatile hydrocarbon species from the fragmentation reactions of poly-aromatic constituents through wood with an open-pore channel system [22]. Fig. 2 shows the curves of FTIR AC pre-forms 750 � C/2 h (AC-1), 850 � C/20 min (AC-2), 850 � C/1 h (AC-3) and 850 � C/2 h (AC-4) at 1400 � C, for derivation of bio-SiC ceramics. Stretching of –C–O–C appeared at about 1117 cm 1 [23,24]. The surface groups from carbon, that was previously activated using the 750 � C/2 h (AC-1) sample, were almost completely decomposed after the high-temperature treatment at 1400 � C. The 850 � C/2 h (AC-4) sample showed the highest thermal stability, which can be identified by the presence of high-intensity peaks. The most prominent peak for this series of the spectra has appeared at 800 cm 1 and it was assigned to the Si–C bond [25,26], thus, indicating that the procedure is suitable for the production of bio-SiC, heated at 1400 � C in an argon atmosphere. The existence of Si–C fundamental streching vibration at ~800 cm 1 clearly suggests that high-temperature bio-SiC processing has gone to completion. However, an intense broader band located at about 800/825 cm 1 was observed and characterized by Si–C fundamental stretching vibra tion. The broad band which was located between 825 and 898 cm 1 belongs to SiC prepared from a carbon/silica hydrogels [27]. However, according to other research reports [28,29], the main band was located between 789 and 794 cm 1. The discrepancy between these values may be attributed to difference in morphology of analyzed SiC and impreg nation procedures [30]. With respect to the above-presented FTIR analysis results, proof of the benefits of activated carbon obtained through a pyrolytic process at temperatures above 600 � C for the production of high-performance bioceramics can be viewed in Table 1 which clearly shows the relevance of obtained FTIR results (Fig. 2). Table 1 illustrates positions of infrared bands characteristic for some organosilicon groups.
Table 1 The positions of infrared bands which are characteristic of some organosilicon groups [31]. Group
Wave number (cm 1)
Additional comments
Si–O–Si
1130–1000
Si–H in amorphous silicon CERAMICS Silicon carbide
2150–2000
In the region of 1130–1000 cm 1 siloxanes determine very strong infrared bands. Single Si–O–Si band was shown by disiloxanes and small-ring cyclo-siloxanes. Amorphous silicon (a-Si) has been found in the range of 2150–2000 cm 1 Si–H bands at three or more frequencies. Silicon carbide occurs in a bewildering number of crystal modifications. It also occurs as amorphous material (a-SiC). Detailed SPECTRA-STUCTURE CORRELATIONS for these forms are not available. It can be expected that different forms will have small differences in the infrared spectra, but all these forms of SiC show strong absorption at or near 800 cm 1.
ca. 800
3.2. XRD analysis of studied samples The XRD diffraction patterns of the bio-SiC ceramics obtained for 750 � C/2 h (AC-1)-a, 850 � C/20 min (AC-2)-b, 850 � C/1 h (AC-3)-c, and 850 � C/2 h (AC-4)-d after carbothermal reduction at 1400 � C are pre sented in Fig. 3. The peak at ~33.55� corresponds to (1010) planes of hexagonal SiC, while peaks at ~ 26 and 43� correspond to (002) and (111) planes of the graphite, where all other peaks originate from cubic β-SiC. The peaks marked by Miller indices (111), (200), (220), and (311) on diffracto grams indicate the crystalline structure of cubic silicon carbide. The position of observed peaks indicated the presence of β-SiC (cubic space group F-43 m, No. 216, 2θ ¼ 35.650� , d111 ¼ 2.516). The d (interplanar spacing) values of four peaks are 2.516, 2.179, 1.541, and 1.314 Å, respectively. The peak detected at 2θ ¼ 33.554� , d ¼ 2.668 Å, belongs to hexagonal α-SiC (space group P 63mc, No. 186) and was characteristic of the collision error on the (111) planes in the β-SiC phase [32,33]. The crystal defects, such as stacking faults, are observed which altered the steps of sequence (111) planes to yield an additional peak at 2θ ~ 33.55� . The weak, broadly rounded peaks of amorphous carbon at ~26 and 43� are visible in all XRD patterns except in the last one (d) which is sometimes associated with the sample-processing temperature, whereby the lower temperature means a smaller angle. The obtained diffracto grams also indicated partial graphitization of the material with residual
Fig. 2. Thermally treated AC pre-forms 750 � C/2 h (AC-1), 850 � C/20 min h (AC-2), 850 � C/1 h (AC-3) and 850 � C/2 h (AC-4) after the impregnation process using TEOS at 1400 � C, in order to obtain the bio-SiC ceramics.
Fig. 3. The XRD pattern of the bio-SiC samples produced at 1400 � C. 4
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amorphous carbon and match to the (002) and (111) planes of graphite [34,35]. All peaks are broad and with low intensities which differ slightly with changes in time interval, indicating nano-crystalline nature of powders. With the increase in holding time, the intensities of XRD peaks which belong to β-SiC were moderately increased. At the same time, the XRD peaks β-SiC phase became more intense. The diffracto grams of the samples (a) and (b) are different from the other two based on the peaks with lower intensity, indicating a gradual transformation from a micro-to nano-particle sizes that can be observed in SEM images (see the later results). Based on the d-values and 2θ angles from the ICSD database, card number #28389 (the cubic β-SiC) and card number #164970 (the hexagonal α-SiC), a very good agreement was obtained. The XRD pat terns indicate that the carbon template is completely converted into SiC, at a sufficiently high processing temperature. The bio-SiC synthesis occurring through a reduction reaction is practically accomplished and no essential amount of the crystalline SiO2 phase can be detected [34]. Thus, for the products obtained at 1400 � C (Fig. 3), the peaks of the major cubic type β-SiC phase appear. The diffraction peaks of crystalline SiO2 (cristobalite) were not observed which indicate that nearly the whole amount of silica has been consumed during carbothermal reduction at 1400 � C. This means that the whole amount of silica was mineralized. During quantitative treatment, it reacted quantitatively. It is evident that, after carbothermal reduction at 1400 � C, the β-SiC phase, as a major phase with cubic structure and a small amount of the hexagonal structure, can be identified. If we consider the synthetic approach presented here, then, it follows that the bio-SiC predominantly consists of the cubic structure but a small amount of the volume fraction of hexagonal structure is always present under applied processing con ditions. Results from Fig. 3 show that partial transformation, hexagonal [1 h 850 � C] → cubic [2 h 850 � C], took place during the synthesis. Ortiz et al. [36] found that defects such as twins and mismatch errors always exist in the cubic SiC phase. The β to α transformation is a kinetically controlled process and, often, several days are needed to achieve the complete conversion [37], where α to β is preferably governed by diffusion, with lower barrier energy and higher conversion rates. In transformation from α to β structures, it can be assumed that the diffusion-controlled process (growth) can happen in this ceramic because of the presence of long range rearrangement process of the atoms in a silicon matrix in order to form precipitates. In the SiC powder product, which was synthesized at 1400 � C during 2 h from the carbon perform activated at 850 � C, the nucleation of hexagonal SiC structure has just started but, obviously, high activated temperature and high reduction processing temperature are probably quite enough to induce fast conversion where trans formation from the cubic to the hexagonal form is very small. Namely, β-structure is more stable at lower temperatures and transforms irre versibly to the α-form at very high temperature (~2000 � C). However, the products obtained at nearly 2000 � C consist of α-SiC [38].
Fig. 4. a) - d) Raman spectra of 750 � C/2 h (AC-1), 850 � C/20 min (AC-2), 850 � C/1 h (AC-3), and 850 � C/2 h (AC-4) samples, respectively, which are treated with TEOS and further annealed at 1400 � C for 2 h.
consisting of a sequence of Si–C double layers stacked on top of each other. The order of agreement of double layers is called “hexagonal” (α-SiC) if bonds in adjacent double layers are in an eclipsed orientation; while, for staggered configuration, the order of agreement is called “cubic” (β-SiC (3C–SiC)) [40]. β-SiC has characteristic peaks at around 796 cm 1 and 850 cm 1, corresponding to the zone center transversal optical-phonon mode TO, as well as at 950 cm 1, corresponding to the longitudinal optical phonon LO [39,41,42]. For α-SiC two additional folded phonon modes are typical TO at around 767 cm 1 and the most intense TO mode at 789 cm 1, being identified as α-SiC inclusions in β-SiC [41]. The Raman peaks originating from Si–C vibration appear only in the samples 850 � C/1 h (AC-3) and 850 � C/2 h (AC-4) treated by TEOS at 1400 � C, as shown in Fig. 4 and enlarged in Fig. 5. The obtained bio-SiC Raman peaks for samples 850 � C/1 h (AC-3) and 850 � C/2 h (AC-4), demonstrate detection and quantification of the polytype inclusions during the growth of the nano-cubic β-SiC. Namely, the peak obtained for sample 850 � C/1 h (AC-3) at around 788 cm 1 could be assigned to the intense peak of the α-SiC TO mode, which is used for identification of the α-SiC polytype inclusion in β-SiC or, in other words, the transformation of cubic SiC to a hexagonal structure; while peak at 917 cm 1 could be assigned to the LO phonon peaks of cubic SiC (Fig. 5a). The obvious shift obtained (33 cm 1 for the LO β-SiC bulk) is an indication of the nano-wired morphology of SiC [42,43]. Furthermore, the Raman spectrum related to the 850 � C/2 h (AC-4) sample shows three peaks at 787 cm 1, 853 cm 1, and 948 cm 1 indi cating the extra “combined” SiC morphology: the β-SiC with the pres ence of nano-wires, accompaning β-SiC with the small amount transformed to the hexagonal SiC (less then in the 850 � C/1 h (AC-3) sample) [41,43]. Most intense peak exists around 853 cm 1, originating from β-SiC [44], and Fig. 5b might be considered as a representation of better crystallization. Raman spectra prove the transformation of β-SiC to α-SiC for the 850 � C/1 h (AC-3) sample as well as the existence of
3.3. Raman spectroscopy The Raman spectra of activated carbon (AC) samples 750 � C/2 h (AC1), 850 � C/20 min (AC-2), 850 � C/1 h (AC-3), and 850 � C/2 h (AC-4) treated with TEOS and further annealed at 1400 � C for 2 h are presented in Fig. 4 a-d, respectively. The four Raman spectra look very similar, except for the region from 600 cm 1 to 1100 cm 1 obtained from carbon samples activated at 850 � C for 1 h and 2 h. The 750 � C/2 h (AC-1) and 850 � C/20 min (AC-2) samples show two main peaks at around 1350 cm 1 and 1580 cm 1, plus two small peaks at around 2690 cm 1 and 2930 cm 1. The samples 850 � C/1 h (AC-3) and 850 � C/2 h (AC-4) except all previously mentioned peaks, exhibited additional small peaks around 788 cm 1, 853 cm 1, 917 cm 1, and 940 cm 1. The vibration range from 760 cm 1 to 950 cm 1 is characteristic for the Si–C bond [39]. SiC crystallizes in polytype crystal structures, all 5
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could be an indication that a number of graphene layers (or presented carbon clusters) decreases in investigated samples (Fig. 4). It is well known that the Raman spectral shape of a carbon material drastically changes not only due to the kind of abundant allotropic forms of carbon but to the fine structural transformation of the individual allotrope. Thus, Fig. 6 represents Raman spectra of the activated 1 h 850 � C carbon without TEOS treatment and further annealing, and the obtained Raman spectrum for the 850 � C/1 h (AC-3) sample treated with TEOS and annealed at 1400 � C for 2 h. The red-shift and narrowing of the G-peak are obvious. This is an undoubted indication of an increase in ordering of the carbon phase (presented in the form of few layers or an organized carbon cluster) in bio-SiC samples. Due to a better Raman scattering efficiency of C–C bonds compared to Si–C bond, the Raman lines, characteristic for crystalline SiC, could be almost unobservable, especially if SiC is in the form of very small crys tallites [44]. This is the reason why the Raman spectroscopy is unable to detect the peaks of the SiC in the bio-SiC 750 � C/2 h (AC-1) and the 850 � C/20 min (AC-2) samples at 1400 � C. However, it is evident that an increase in the activation temperature and holding time causes growth of the SiC crystallite phase. Such results are in accordance with those obtained through the XRD analysis. Furthermore, the residual carbon which is not transformed to SiC could have many beneficial properties in catalysis, especially in photocatalytic processes, where it can serve as an electron acceptor, facili tating the interface charge separation, etc [44]. The mechanism of SiC formation from the PTS precursor after carbonization and activation processes, and the subsequent carbo thermal reduction, at 1400 � C, in an argon atmosphere, may be explained by the following chemical reactions, such as: TEOS (at elevated temperature on C-surface) → SiO2 (and SiO) þ C (and CO) þ H 2↑ (1) SiO2 þ 2C → SiC þ CO2↑
(2)
SiO þ 2C → SiC þ CO↑ (possible formation of polycrystals)
(3)
SiO þ 3CO→ SiC þ 2CO2↑ (possible formation of nano-wires)
(4)
3C–SiC (β-SiC) from PTS, identified through the characterization analysis in this study, can obviously be used as high-temperature bioceramics. It may be supposed that the two above-indicated reactions can occur (reactions [3,4]) in the conversion process following their competition with one to another.
Fig. 5. a) – b) Raman peaks originated from bio-SiC samples at 850 � C/1 h (AC3) and 850 � C/2 h (AC-4), respectively, treated with TEOS, annealed at 1400 � C for 2 h.
nano-wires, whereas the Raman spectra for the 850 � C/2 h (AC-4) sample consists mainly of a large amount of β-SiC nano-wires, which are also documented in the XRD (see the previously shown results) and SEM measurements (see the later results). Raman analysis also suggests that some amount of free carbon is present in all obtained samples, because of the existence of two intensive bands, i.e. G-band around 1580 cm 1 and D-band around 1350 cm 1, which are generally ascribed to E2g and A1g in-plane vibration modes in polycrystalline graphite, respectively [42,45]. The D-band of carbon or the so-called “disorder band” is ascribed to structural imperfections. The G-band (or graphitic) arises from stretching the C–C bond in graphitic forms and is attributed to all sp2 carbons. The peak at 2690 cm 1 is of a second order overtone of the D-mode (also named 2D) whereas the peak at 2930 cm 1 is a combination of the peak obtained scattering from G and D (G þ D) [45]. The activation time and elevated temperature have more impact on the position of G-band than the D-band position as it may be seen in Fig. 4. The red-shift (from 1591 cm 1 to 1578 cm 1) of the G-band might indicate an increase in the ordering of the carbon phase [45]. The visual inspection of the 2D peak shows no clear asym metry, nor a shoulder structure. The small intense and broad 2D peak (around 2690 cm 1), as observed here, could be an indication of the existence of carbon in the form of few graphene layers or some carbon domains – clusters [45]. Furthermore, the obtained 2D peak red-shift
Fig. 6. The comparison of the Raman spectra obtained from non-TEOS treated and the 850 � C/1 h (AC-3) TEOS treated sample, annealed at 1400 � C. 6
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3.4. SEM image analysis of TEOS infiltration process in AC samples after carbothermal reduction The SEM images of the AC samples which undergo carbothermal reduction at 1400 � C, designated as 750 � C/2 h (AC-1), 850 � C/20 min (AC-2), 850 � C/1 h (AC-3), and 850 � C/2 h (AC-4) are shown in Fig. 7 ad, respectively. The SEM image, presented in Fig. 7a, shows the morphology of 750 � C/2 h (AC-1) sample including examination of the cross-section along the axial direction. At the lowest activation temperature (750 � C for 2 h) and, after three cycles of TEOS soaking, the presence of non-sealed surface pores (mainly, of a round or an elliptic shape) can be identi fied in the carbonized precursor, exactly where a small extent of con version of SiO2 into SiC occurred. Since the quantity of silica in material is smaller than the quantity of carbon, the formation of SiC will be associated with the quantity of SiO2 mineralized on PTS. Furthermore, the whole cellular structures as cells, pits, and lumens mainly remain after thermal treatment. At a higher activation temperature (850 � C for 20 min) (Fig. 7b), the 850 � C/20 min (AC-2) sample that undergoes carbothermal reduction, shows structure approximately resembling that of the honeycomb where large cavities were gradually repeatedly dipped in TEOS. Therefore, SiO2 starts to cover the surface and fills the cavities and pores of tra cheids and pits, suggesting that SiO2 solution penetrated the cell wall structures and condensed around cell-shaped tissues. Upon conversion to SiC, reaction between molten silicon and amorphous carbon to form SiC was accompanied by a change in volume. It can be supposed that the overall sample dimensions did not undergo drastic changes, but the volume change within the pores was accommodated. This can resulted in decreased pore diameters but a certain number of pore walls was damaged, which led to some effectively larger pores. For the 850 � C/1 h (AC-3) sample (Fig. 7c), the porosity decreased and nearly completely disappeared once the material was converted into SiC. However, in the case of the 850 � C/20 min (AC-2) a 20 min (0.3 h) sample, the fact that it was held at an activation temperature of 850 � C was insufficient to completely convert the entire carbon scaffold into SiC. As carbon in smaller pores reacted, the pores were closed due to the volume expansion, impeding further carbon conversion into SiC. Most of the residual carbon was converted into SiC after this reaction step was repeated. The presence of small-diameter pores (Fig. 7a) was identified and they were pinched close (Fig. 7b) due to the volume expansion which occurs during C (at Fig. 7c we may see expanded graphite) → SiC (generation of SiC product - Fig. 7c) reaction. The 850 � C/1 h (AC-3) sample (Fig. 7c) exhibits formation of nano-wires with obviously com bined structural morphology, including the reservation of tubular interpenetrating pores from PTS. In this reaction stage, the SiO2 nanoparticles form droplets in touch with the activated carbon, thus, form ing SiO vapor. The SiO vapor further reacts with carbon to form β-SiC particles, which act as the nuclei for the growth of the nano-wires. Once the droplets are saturated with β-SiC, crystallization of β-SiC occurs, followed by the growth in one direction to produce nano-wires. In the case of 850 � C/2 h (AC-4) (Fig. 7d) amount of SiC nano-wires increases. Both, straight and curved nano-wires grow into a massive network (Fig. 7d), where nano-wires are distinguished by their great length. This behavior was confirmed by FTIR, XRD as well as Raman measurements. Fig. 8 shows the SEM image at higher magnification of the 850 � C/2 h (AC-4) sample. It may be seen in Fig. 8 that a certain number of nano-wires contain spherical caps at the tip of nano-wires, which are more uniform than SiC nano-wires previously synthesized with the vapor-liquid-solids (VLS) mechanism in other researches [46–48]. It is worth noting that no extra metal catalysts were introduced in our experiments. Therefore, we can assume that some metal catalyst would probably accelerate the forma tion of SiC nanowires. It can be assumed that observed spherical caps (Fig. 8) probably consists of silicon, oxygen, and carbon and some of alkali or alkaline earth metals, which may play the catalytic role in SiC
Fig. 7. SEM images of 750 � C/2 h (AC-1) (a), 850 � C/20 min (AC-2) (b), 850 � C/1 h (AC-3) (c), and 850 � C/2 h (AC-4) (d) samples which undergo carbo thermal reduction at 1400 � C for derivation of bio-SiC from activated carbons of the carbonized PTS precursor.
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development of SiC particles inside the porous structure. The X-ray powder diffraction (XRPD) analysis showed that major SiC polytype is the cubic SiC (β-SiC) and remainder is the hexagonal SiC polytypic (α-SiC). It was found that the carbon obtained through carbonizing the PTS precursor and activated at 850 � C during longer holding time (1 and 2 h) in a reducing atmosphere at 1400 � C possesses β-SiC (cubic) nanowires. Besides the XRPD analysis, the formation of β-SiC nano-wires was also confirmed by the Fourier-transform infrared (FTIR) and Raman spectroscopy analyses. The process described in this paper is rather complex for the industrial application, because it involves three pro cessing steps at high temperatures: pyrolysis, activation, and, finally, reduction. However, the described procedure allows the production of ultra-long SiC nano-wires, which have pre-disposition for unique optical properties, favorable thermal stability properties, as well as flexible nano-mechanical properties. Therefore, in regard to these characteris tics, further research will be carried out. Fig. 8. The SEM image of the 850 � C/2 h (AC-4) sample at 1400 � C at a higher magnification.
Acknowledgement This research work was partially supported by the Serbian Ministry of Education, Science and Technological Development under the pro jects numbered 172015, III45005, TR34011 and III45001.
nano-wire formation. From the presented results, it is evident that the vapor-solid (VS) mechanism probably occurs which can be attributed to increased tem perature and partial pressure creating reaction between CO(gas) and Si(solid) or CO(gas) and SiO(gas) [49]. Considering the fact that the (111) plane possesses the lowest surface energy among SiC surfaces, a SiC nano-wire, thus, tends to grow along a [111] direction. The diffraction peaks in XRD patterns at (111), (220), and (311) planes (Fig. 3) confirm that a nano-wire consists of β-SiC phase; the (111) lattice plane of a nano-wire is exposed to the outside for the generation of a branched nano-wire. Growth of SiC nano-wires is highly dependent on the phase of gaseous compositions and reaction conditions. Namely, the phases CO (gas), SiO(gas), SiO2(solid), and SiC(solid) can be formed through the Si–C–O–N system and through a complex scheme of reactions, which requires additional thermodynamic calculations. It should be noted that, within VS mechanism, the height of SiC nano-wires is restricted by re action conditions. Further thermodynamic calculations may help in estimating whether SiC is a stable phase when the partial pressure of CO is lower than 0.25 atm [50]. Presence of hydrogen reaction [1] in the gas phase may favor the formation of SiO from SiO2. SiO reacts with carbon to form SiC. The basic reactions involved in the formation of SiC nano-wires are as follows:
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(5)
SiO2 þ C → SiO þ CO↑
(6)
SiO þ 2C → SiC þ CO↑ 1
Formed β-SiC nano-wires show a strong FTIR band at 800 cm due to a Si–C stretching band and this is confirmed by FTIR (Fig. 2). Raman spectroscopy also shows existence of nano-wires. Inter-connected porous structure from the PTS precursor, after carbonization and during a carbothermal reduction, plays an important role in the SiC synthesis because pores accelerate the passage of gaseous SiO, CO, and CO2. 4. Conclusions Silicon carbide (SiC) ceramics have been synthesized through a carbothermal reduction of activated carbons (AC) obtained by carbon izing Platanus orientalis L (plane tree fruits). The activated samples were “enriched” with TEOS (tetraethyl orthosilicate) in a procedure for the impregnation of silicon particles into carbon templates. The various obtained porous products significantly depend on the activation tem perature and holding time, before the application of carbothermal reduction. The 750 � C/2 h (AC-1) sample synthesis of SiC begins at 1400 � C, but it is not completed after three cycles of TEOS soaking. The in crease in the activation temperature and the holding time initiates the 8
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