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Article Cite This: ACS Omega 2018, 3, 4006−4016

Bulk Polymer-Derived Ceramic Composites of Graphene Oxide Chen Shen,†,‡ Elizabeth Barrios,†,‡ and Lei Zhai*,†,‡,§ †

NanoScience Technology Center, University of Central Florida, Orlando, Florida 32826, United States Department of Material Science and Engineering and §Department of Chemistry, University of Central Florida, Orlando, Florida 32816, United States

‡

S Supporting Information *

ABSTRACT: Bulk polymer-derived ceramic (PDC) composites of SiCO with an embedded graphene network were produced using graphene-coated poly(vinyl alcohol) (PVA) foams as templates. The pyrolysis of green bodies containing cross-linked polysiloxane, PVA foams, and graphene oxide (GO) resulted in the decomposition of PVA foams, compression of GO layers, and formation of graphitic domains adjacent to GO within the SiCO composite, leading to SiCO composites with an embedded graphene network. The SiCO/GO composite, with about 1.5% GO in the ceramic matrix, offered an increase in the electrical conductivity by more than 4 orders of magnitude compared to that of pure SiCO ceramics. Additionally, the unique graphene network in the SiCO demonstrated a drop in the observed thermal conductivity of the composite (∼0.8 W m−1 K−1). Young’s modulus of the as-fabricated SiCO/GO composites was found to be around 210 MPa, which is notably higher than the reported values for similar composites fabricated from only ceramic precursors and PVA foams. The present approach demonstrates a facile and cost-effective method of producing bulk PDC composites with high electrical conductivity, good thermal stability, and low thermal conductivity.

1. INTRODUCTION Silicon-based polymer-derived ceramics (PDCs) based on the pyrolysis of polymeric precursors were first demonstrated in the 1960s and have since gained momentum in the materials research community due to their ease of fabrication, highly tailorable nano- and microstructures, and unique material properties.1−7 The polymer-to-ceramic conversion process has demonstrated easy integration of functional groups into the ceramic material. Additionally, the polymeric precursors enable materials to be formed using common polymer-processing strategies such as extrusion, blow molding, injection molding, and spin-coating.8−12 Utilizing such techniques, ceramics in the forms of fibers,13,14 porous materials,15−17 and membrane coatings18,19 have been fabricated, which are otherwise known to be difficult or even impossible to obtain through conventional routes. Common polymeric precursors used for PDC fabrication include oligosilazanes and polysiloxanes, both of which demonstrate a unique nanostructure, leading to the observed material properties. Typically, nanodomains of Si3N4 or SiO2 surrounded by carbon nanodomains are found within the structure of an oligosilazane- or polysiloxane-based PDC, respectively.1,7,20−22 This type of nanostructure has offered PDCs various attractive material properties such as low density, high-temperature thermomechanical properties, high piezo resistivity, and chemical/thermal stability.20,21,23,24 Applications of such PDCs have been found in micro-electro-mechanical systems,25−27 energy storage devices,12,28−30 and high-temperature sensors.31−34 © 2018 American Chemical Society

Whereas many applications require materials with a high electrical conductivity, PDCs possess a low electrical conductivity of 10−12−10−4 S cm−1 due to their semiconductive characteristics.35,36 Adding conductive components such as carbon nanotubes,20,37−39 carbon fibers,40 carbon black,41,42 graphite flakes,43 reduced graphene oxide (GO),44,45 and metal salts46,47 into an insulating PDC matrix can effectively increase the electrical conductivity of the PDC-based composites. In addition, introducing percolated networks of conductive materials into PDC composites is found to be a highly effective method of increasing the electrical conductivity.48−50 For example, additions of embedded graphene aerogels were reported to be able to increase the electrical conductivity of composites to 1.57 S cm−1.49 Here, we report the introduction of conductive percolated graphene oxide (GO) networks into bulk PDC composites through a facile templating method, where polysiloxane is infiltrated into GO-coated poly(vinyl alcohol) (PVA) foams to produce the green bodies used for pyrolysis. PVA foams provided a continuous network to support GO and served as a sacrificial component that decomposed during pyrolysis, leading to a GO network embedded in SiCO PDCs. The impact of the percolated GO networks on the preparation of bulk PDC composites, graphitic domain formation, electrical conductivity, and thermal conductivity is investigated. Received: March 15, 2018 Accepted: April 5, 2018 Published: April 10, 2018 4006

DOI: 10.1021/acsomega.8b00492 ACS Omega 2018, 3, 4006−4016

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ACS Omega Scheme 1. Schematic Illustration of the Preparation of GO Network Embedded PDC Composites

Figure 1. (a, b) Transmission electron microscopy (TEM) images of GO flakes. (c) SAED pattern of GO flakes. Four sets of 6-fold SAED pattern regarding the (1100) lattice fringe of graphitic structures are labeled in red, blue, green, and purple circles. (d) Raman spectrum of GO flakes. The positions of the G peak, D peak, and two-dimensional (2D) peak are at 1563.8, 1343.9, and 2693.6 cm−1, respectively.

GO flakes were deposited onto PVA foam walls. The liquid SiCO precursor (SILRES 62C) was filled into the voids of the PVA foam and cross-linked to produce SILRES 62C/GO/PVA green bodies. The PVA foam was decomposed during the

2. RESULTS AND DISCUSSION Scheme 1 illustrates the fabrication process of SiCO/GO composites. The graphene network was obtained by immersing the preshaped PVA foam into a GO−ethanol dispersion where 4007


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Figure 2. SEM images of (a, b) GO-attached PVA foams after the immersion of PVA foams in the GO dispersion, (c, d) SILRES 62C/GO/PVA composite green bodies, and (e, f) SiCO/GO composites obtained from SILRES 62C/GO/PVA composites pyrolyzed at 1000 °C.

smaller than that for GO synthesized from Hummer’s method (0.9−1.9).58−61 The intensity ratio of the D peak and the G peak is related to the ratio between sp3 carbon and sp2 carbon. A smaller value of the intensity ratio reflects a higher percentage of sp2 carbon in the samples (i.e., more ordered graphitic structures with fewer defects). Therefore, GO is observed to have a more ordered crystalline structure than GO synthesized by Hummer’s method. The graphitic structure along the basal plane of the GO flakes yield GO with electrical conductivities of around 10 S cm−1, which are much higher than that of GO synthesized from Hummer’s method (∼4 × 1010 Ω sq−1).62 Meanwhile, the functional groups on the edges facilitate the dispersity of GO in solvents such as ethanol, acetone, and water, making GO an excellent candidate for solution-based processing. It should be noted that ethanol is the best medium to disperse this GO, which can be attributed to the abundance of hydroxyl groups on the edges of the GO flakes. Therefore, ethanol was used as the dispersant for GO in the preparation of the PDC composites. GO flakes were attached to the surface of PVA networks through van der Waals interaction when the PVA foams were immersed into GO−ethanol dispersions. Figure 2a,b shows the scanning electron microscopy (SEM) images of PVA foams covered by GO. High-magnification SEM images (Figure 2b) show the excellent coverage of GO on the PVA foam walls compared to the SEM image of bare PVA foam (Figure S3).

pyrolysis, which led to different shapes of SiCO/GO composites with a continuous GO network. The GO used in this study is a commercially available multilayered graphene oxide produced through ball-milling with hydrogen peroxide. The edges of the GO flakes are functionalized with hydroxyl, carbonyl, and carboxylate groups, and the oxygen content is reported to be about 5 wt %. The GO flakes have sizes below 1 μm (Figure 1a), with the average size of flakes used in this work being around 350 nm, as measured from dynamic light scattering (DLS) (Figure S1). Figure 1b clearly shows the multilayer structure of GO. Additionally, four sets of 6-fold selected area electron diffraction (SAED) patterns circled in Figure 1c indicate the existence of the stacked multilayer graphitic structure in GO.52−54 The arcs in the SAED image corresponding to spacings of 2.10 and 1.23 nm are attributed, respectively, to the (1100) and (1120) lattice fringes of graphitic structures randomly oriented in GO flakes.55−57 The typical thickness of GO flakes is found to be around 8 nm from the atomic force microscopy (AFM) height profile (Figure S2), suggesting that GO flakes have about eight layers of stacked graphene. The Raman spectrum verifies the existence of graphitic domains in the GO flakes due to the presence of the characteristic G, D, and 2D peaks at 1563.8, 1343.9, and 2693.6 cm−1, respectively (Figure 1d). The intensity ratio between the D peak and the G peak is 0.39, which is much 4008


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dramatic mass loss from 300 to 350 °C, and only 4% of the original mass remains at 1000 °C. On the contrary, the mass of GO remains nearly constant through the heating process. The minimal mass loss at 100−150 °C is due to the evaporation of residual moisture, with 96% of its original mass remaining at 1000 °C. The major mass loss in the SILRES 62C decomposition happens from 300 to 800 °C, and around 75% of its original mass remains at 1000 °C. This mass loss is primarily attributed to the decomposition of organic components in SILRES 62C during the heat treatment. The breakage of Si−H bonds and Si−C bonds in the polymeric precursors leads to a release of gases such as methane and hydrogen.64−68 The major mass loss of the green body composite is in the temperature ranging from 200 to 800 °C, and about 63% of its original mass remains at 1000 °C. These TGA studies suggest that most of the PVA decomposed with a minimal loss of the GO during pyrolysis. The loading percentage of GO in SiCO/GO composites is then calculated from the density of each component and is found to be around 1.4−1.9 wt % (density of composites ∼1.76 g cm−3, Table S3; GO 1.8 g cm−3 and pure SiCO 1.768 g cm−3, Table S4). The elemental compositions of SiCO/GO composites obtained at different pyrolysis temperatures were investigated by energy-dispersive X-ray spectroscopy (EDS) (Table 1). The atomic percentage of carbon decreases as the pyrolysis temperature increased, corresponding to the decomposition of the organic components into carbon-containing gases such as methane and ethane.69 At the same time, the atomic ratio between silicon and oxygen in ceramic composites is around 0.5 during the entire pyrolysis process, suggesting the existence of SiO2 clusters in SiCO ceramics, which is in accordance with the observation of the structural evolution of PDCs in previous reports.35,70,71 The dissociation energy of the Si−O bond (452 kJ mol−1) is higher than that of the C−C bond (346 kJ mol−1), Si−C bond (318 kJ mol−1), and Si−H bond (319 kJ mol−1) in the composites.72 Therefore, SiO2 clusters are the dominant component in SiCO ceramics. However, the hydrogen contents in composites cannot be monitored by EDS, but it is believed that the hydrogen contents decrease by applying higher pyrolysis temperatures due to the decomposition of Si−H bonds. In this study, Raman spectroscopy was used to monitor the generation of the graphitic domains within pure SILRES 62C, SILRES 62C/GO composites, and SILRES 62C/GO/PVA composites at different temperatures. It has been reported that graphene facilitates the formation of graphitic domains in PDCs during pyrolysis.49 Figure 4a shows that SILRES 62C pyrolyzed at or above 800 °C shows the G peak and the D peak, which are the signature Raman peaks of graphitic domains in Raman spectra. Although GO itself has Raman signals originating from the graphitic domains (Figure 1d), it is interesting to note that SILRES 62C/GO composites pyrolyzed below 700 °C show no discernible G or D peaks (Figure 4b). It is believed that low GO concentration in the composites makes them undiscernible in the Raman spectra of SILRES 62C/GO composites pyrolyzed below 700 °C. Therefore, the signature G and D peaks in the Raman spectra of SILRES 62C/GO composites pyrolyzed at or above 700 °C are not attributed to the GO component, but to the graphitic domains generated from the pyrolysis of SILRES 62C precursor. Furthermore, the G peaks shift from 1586.4 to 1614.7 cm−1 when the pyrolysis temperature increased from 700 to 1000 °C. The blue shift (28.3 cm−1) of the G peaks indicates the effective transition

The GO-coated PVA foams were filled with the polymeric precursor in a vacuum environment and then cured at 160 °C to obtain the green bodies for pyrolysis. The vacuum-assisted SiCO precursor infiltration produced the fully dense green bodies (Figure 2c) by removing the trapped air between the PVA foam and the SiCO precursor. The thickness of GOcoated PVA foam networks within the green bodies is around 10−25 μm, as seen from Figure 2d. The green bodies were pyrolyzed at 1000 °C to produce SiCO/GO composites. The linear shrinkage of the composites after pyrolysis at 1000 °C is 22−24% compared to that of green bodies (Figure S6). The cross-sectional area of the PDC composites shows the interface between the GO component and the SiCO matrix (Figure 2e,f). The breakage of the SiCO/ GO composites happens at the interface due to the poor interactions between the matrix and graphitic domains. It is interesting to notice that the thickness of the GO component is less than 2 μm, suggesting that the PVA was decomposed during pyrolysis. The SEM images of the cross section (Figure 2e,f) and the surface (Figure S7) of SiCO/GO composites indicate the formation of fully dense SiCO/GO composites due to the successful infiltration of SILRES 62C and complete PVA decomposition. In contrast, SiCO samples pyrolyzed from pure cross-linked SILRES 62C have obvious cracks due the stresses generated in the pyrolysis process. We believe that the embedded GO network plays two roles in producing crackfree PDC composites: (1) providing channels to release the gas produced by the pyrolysis of SILRES 62C and decomposition of PVA and (2) functioning as a scaffold to hold together the SiCO matrix. The measured Young’s moduli of the bulk fully dense SiCO/GO composites pyrolyzed at 1000 °C have an average value of 213.4 MPa (Table S2). In comparison, bulk SiCN ceramics produced from oligosilazane using poly(vinyl alcohol) sponges as templates63 have a low strength (80 MPa) due to 10% pores in the ceramic. The higher strength of the SiCO/GO composites observed in the current study is likely attributed to lower porosity. The mass change of SILRES 62C/GO/PVA composites and each component in the PDC composites (i.e., PVA, SILRES 62C precursor, and GO) during pyrolysis was investigated by thermal gravimetric analysis (TGA). The TGA curves in Figure 3 show the mass loss of materials from room temperature to 1000 °C. The thermal decomposition of PVA foam causes a

Figure 3. TGA curves of the SILRES 62C/GO/PVA composite and each component in the composite. 4009


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ACS Omega Table 1. Elemental Compositions of SiCO/GO Composites Obtained at Different Pyrolysis Temperatures carbon (%) oxygen (%) silicon (%)

300 °C

400 °C

500 °C

600 °C

700 °C

800 °C

900 °C

1000 °C

68 21 11

67 21 12

63 23 14

57 29 15

56 29 15

51 33 16

49 33 17

47 36 17

blue shift of the G band (3.9 cm−1) of the SILRES 62C samples is much smaller than that of the pyrolyzed SILRES 62C/GO composites (23.1 cm−1) and SILRES 62C/GO/PVA composites (28.3 cm−1) (Table S5). Therefore, GO-coated PVA foams are shown to facilitate the formation of graphitic domains in the composites. During the pyrolysis of preceramic polymers, carbon atoms are produced between SiO2 domains, which can crystallize into graphitic domains. If graphene is presented in the matrix, the transition of carbon atoms generated in the area close to graphene may have different values of the thermodynamic variables of state, which could lead to a reduction of crystallization activation energy and an overall lowering crystallization temperature. X-ray photoelectron spectroscopy (XPS) studies were conducted to investigate the composition evolution during the pyrolysis process. Because the transition to graphitic carbon in the composites happened at around 700 °C, as indicated by both SEM imaging and Raman spectroscopy, XPS study was focused on the composition change at 700 °C. The XPS peaks at the carbon and silicon binding energy regions are presented in Figure 5. The XPS peaks of the composites pyrolyzed at 700 °C have a small red shift at the carbon region compared to those of the green body sample, indicating the formation of graphitic “free carbon” domains.75−77 Deconvoluted XPS spectrum at the carbon region for the composites pyrolyzed at 1000 °C shows a peak with the binding energy of 284.6 eV, suggesting the formation of graphitic domains. On the contrary, the XPS peak in the silicon region shifts to a lower energy region when the composites were pyrolyzed at 700 °C and further shifts to a lower energy region when the composites were pyrolyzed at 1000 °C. The breakage of Si−C bond and Si−H bond decreases the valance of the silicon element in the composites and leads to the observed shift of the silicon XPS peak to the lower energy region. Therefore, the silicon shift can be correlated to the decomposition of the organic component. To understand the effect of graphitic domains on the electrical conductivity of SiCO/GO composites produced from SILRES 62C/GO/PVA, the electrical conductivities of the samples obtained at different pyrolysis temperatures were measured at room temperature (25 °C) (Figure 6). The composites obtained below 300 °C show electrical resistance beyond the measurement range even with embedded conductive GO networks. The charge accumulation on the bare samples in SEM studies also illustrates their low electrical conductivity (Figure S8). However, the electrical conductivity increased significantly when the pyrolysis temperature was above 300 °C due to the removal of functional groups from GO.78 The electrical conductivity of the composites pyrolyzed at the temperature ranging from 300 to 600 °C increased with higher pyrolysis temperatures. Such an increase was not caused by the formation of graphitic domains in composites because the Raman study showed that the formation of conductive graphitic domains in the SILRES 62C precursor only happened above 700 °C. It is also interesting to note that the electrical conductivity increment slope in Figure 6 is similar to the increase slope of linear shrinkage of the composites. Therefore,

Figure 4. Raman spectra of (a) pure SILRES 62C, (b) SILRES 62C/ GO composites with 2 wt % GO, and (c) SILRES 62C/GO/PVA composites pyrolyzed at different temperatures.

from amorphous carbon to graphitic domains in composites during pyrolysis.73,74 Similarly, the G peak and the D peak appear in the Raman spectra of SILRES 62C/GO/PVA composites pyrolyzed at and above 700 °C (Figure 4c). The 4010


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Figure 6. Room-temperature electrical conductivity of SiCO/GO composites obtained at different pyrolyzed temperatures and the corresponding linear shrinkage after pyrolysis.

Scheme 2. SiCO/GO Composite Pyrolyzed at 1000 °C Remains Highly Electrical Conductive and Can Light Up Four White LED Diodes with an Output of 9 V in the Flame of a Propane Guna

a

The dimension of the composite is 5 mm × 5 mm × 40 mm.

electrical conductivity (∼0.19 S cm−1) using various models by considering the electrical conductivity of GO (∼10 S cm−1), SiCO matrix (1.8 × 10−5 S cm−1), and the concentration of GO (∼1.4−1.9%) (see the Supporting Information).79−82 We believe that the higher observed electrical conductivity is attributed to the formation of percolated graphitic domains guided by the GO network (Figure S9). Scheme 3 illustrates the graphitic domains formed in composites during pyrolysis. The formation of such graphitic domains in the composites follows two mechanisms, namely, self-crystallization (Scheme 3a) and GO-induced crystallization (Scheme 3b,c). The selfcrystallization process generates graphitic domains when pure SILRES 62C was pyrolyzed (Scheme 3a). The abundance of graphitic domains increases with increased pyrolysis temperatures, as demonstrated by Raman spectroscopic analysis (Figure 4c). The electrical conductivity of SiCO obtained at 1000 °C was around 1.8 × 10−5 S cm−1. Both self-crystallization and GO-induced crystallization happened when SILRES 62C/ GO was pyrolyzed (Scheme 2b). The abundance of graphitic domains within the matrix increases with increasing pyrolysis temperatures while GO induces the formation of graphitic domains from the surface of GO flakes. However, the contribution to the overall electrical conductivity (8.7 × 10−4

Figure 5. XPS spectra of SiCO/GO composites at (a) carbon region and (b) silicon region before pyrolysis (room temperature), at the transition state of pyrolysis (700 °C), and after pyrolysis (1000 °C).

we believe that the increase of electrical conductivity of composites pyrolyzed between 300 and 600 °C was caused by the shrinkage of the composites, which pushed the graphene sheets closer to each other, facilitating easier charge transport. Furthermore, the electrical conductivity of composites pyrolyzed above 700 °C continued to increase, whereas the shrinkage of the samples diminished. The increased electrical conductivity in this temperature range was attributed to the formation of graphitic carbons in ceramic matrices. It is important to note that the SiCO/GO composites obtained at 1000 °C are very stable at high temperature in air, as shown in Scheme 2 and the video. The electrical conductivity of GO ceramic composites produced from pyrolyzing SILRES 62C/GO/PVA green bodies at 1000 °C (∼1 S cm−1) is much higher than that of the recently reported SiCO/graphene composite paper (0.05 S cm−1).28 Moreover, it is much higher than the calculated 4011


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of SiCO/GO composites, which is likely due to an increase in phonon scattering at the interfaces. Understanding the mechanism of high electrical conductivity and low thermal conductivity in bulk SiCO/GO composites is important in their applications as thermal insulator, energy sources, and hightemperature sensors. The reported method provides a facile and effective approach to produce bulk PDC composites with multiple functions.

Scheme 3. Schematic Illustration of Graphitic Domains Evolution in (a) SILRES 62C, (b) SILRES 62C/GO, and (c) SILRES 62C/GO/PVA Composites during the Pyrolysis Processes

4. EXPERIMENTAL PROCEDURE 4.1. Materials. The polysiloxane precursor (SILRES 62C, Wecker), poly(vinyl alcohol) (PVA) foam (Sponge King), Pyro-Duct 597-A conductive adhesive (Aremco), ethanol (Fisher), and copper wire (Fisher) were used as received. Argon was purchased from Air Liquide and used without further purification. Edge functionalized graphene oxide (GO) is a few layers of graphene oxide from Garmor Inc. with functional groups such as hydroxyls, carbonyls, and carboxylic acids found only on the edges of the graphene flakes. These functional groups enable good dispersity of GO in solvents such as ethanol, acetone, and water. Meanwhile, the electrical conductivity of this kind of GO (∼10 S cm−1) is much higher than those synthesized from Hummer’s method, which are considered as insulators. This difference is attributed to retaining the graphitic structure along the basal plane of the GO flakes, which facilitates charge transport by interplane hopping among GO flakes. 4.2. Materials Synthesis. 4.2.1. Preparation of SiCO/GO Composites from SILRES 62C/GO/PVA Composites. A mixture containing 700 mg of GO and 100 mL of ethanol was sonicated (Branson digital sonifer equipped with stainless steel tips) for 30 min to obtain a homogenous GO dispersion. Preshaped PVA foams were immersed into the GO dispersion for 2 h. The GO-coated PVA foams were then heated in a vacuum oven at 50 °C for 12 h to completely evaporate the residual ethanol within PVA foams. The dried PVA foams were immersed in the SILRES 62C precursor under vacuum for 12 h to ensure that all voids were filled with the precursor. The filled samples were heated at 160 °C for 16 h to cross-link the polysiloxane, followed by pyrolysis of this green body (i.e., cross-linked composite). In a typical pyrolysis process, the green bodies were pyrolyzed in a Lindberg Blue M Furnace equipped with a one inch quartz tube following a heating schedule of 2 h at each 300, 400, 600, and 1000 °C. The ramping rate for all heating steps was 1 °C min−1, and the composites were allowed to cool naturally after pyrolysis. The whole pyrolysis process was conducted in argon atmosphere. To study the composition evolution of ceramic composites during pyrolysis processes, composites with different final pyrolysis temperatures (300, 400, 500, 600, 700, 800, and 900 °C) were obtained. The details of the pyrolysis conditions are summarized in Table S1. 4.2.2. Preparation of SiCO/GO Composites from SILRES 62C/GO Composites. A mixture containing SILRES 62C and GO of a mass ratio of 50:1 was sonicated by a horn sonifer for 30 min. After sonication, the mixture was poured into cylindershaped molds made from aluminum foil. The filled molds were heated in an oven at 160 °C for 16 h to cross-link the precursor. After peeling the aluminum foil off of the cured samples, the samples were taken through the same pyrolysis procedure as described above for the SILRES 62C/GO/PVA ceramic composites. Similarly, SILRES 62C/GO ceramic composites with different final pyrolysis temperatures (300, 400, 500, 600, 700, 800, 900, and 1000 °C) were prepared to study the

S cm−1) is limited because graphene is randomly dispersed in the SiCO matrix (Figure S10). Both self-crystallization and GO-induced crystallization happened when SILRES 62C/GO/ PVA composites were pyrolyzed (Scheme 2c). Because the GO has formed a network in the matrix, the graphitic domains produced on GO networks greatly enhanced the electrical conductivity (1 S cm−1). Although SiCO/GO composites produced from SILRES 62C/GO and SILRES 62C/GO/PVA have similar amounts of GO, the electrical conductivity of SiCO/GO produced from SILRES 62C/GO/PVA is much higher due to the GO network formed on PVA foams. Investigation of the thermal conductivity of SiCO/GO composites showed that embedded GO had a negative impact on the thermal conductivity. The SiCO/GO composite pyrolyzed at 1000 °C has a thermal conductivity of 0.7654 ± 0.1091 and 0.8676 ± 0.1289 W m−1 K−1 at room temperature (25 °C) and at 300 °C, respectively. The thermal conductivity of SiCO/GO composites is much lower than the reported thermal conductivity of pure SiCO (1.5−3 W m−1 K−1).83,84 The observed decrease in thermal conductivity is probably due to the increase in phonon scattering sites at ceramic/GO interfaces when GO is embedded into the composites.85 Although the in-plane thermal conductivity of graphene is among the highest of any known material at around 2000− 4000 W m−1 K−1,86,87 the obtained thermal conductivity is still heavily dependent on substrates and defects. The mechanism of reduction of thermal conductivity by GO is currently under investigation.

3. CONCLUSIONS In summary, bulk SiCO/GO composites are produced using GO-coated PVA foams as templates. The embedded GO network enables the production of crack-free PDC composites by providing channels to release the gas produced in the pyrolysis and scaffolds to hold together the SiCO matrix. The GO networks in the composites greatly increase the electrical conductivity by providing effective pathways for electron transport, facilitated by the formation of conductive graphitic domains in the ceramic matrix. It is interesting to find that the GO network has a negative impact on the thermal conductivity 4012


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(Figure S5). The morphology and elemental composition of the ceramic composites were examined with a scanning electron microscope (SEM, ZEISS Ultra 55) coupled with energy-dispersive X-ray spectroscopy (EDS). The decomposition and mass loss of the ceramic composites and starting materials were acquired by thermal gravimetric analysis (TGA, TA Instruments) in argon atmosphere. Raman spectra of the ceramic composites were obtained from a WiTec Raman system equipped with a 532 nm laser operated at a power of 0.175 mW for all characterizations. The chemical states of the elements in the ceramic composites were tested by X-ray photoelectron spectroscopy (XPS, PHI 5400).

evolution of the ceramic structure. These composites have GO randomly dispersed in SiCO matrices. 4.2.3. Preparation of SiCO Ceramics from Pure SILRES 62C. The liquid SILRES 62C precursor was poured into molds made from aluminum foil and placed in an oven at 160 °C for 16 h to cross-link the precursor. After removing the aluminum foil, the cross-linked SILRES 62C samples were pyrolyzed at different final pyrolysis temperatures (300, 400, 500, 600, 700, 800, 900, and 1000 °C) using the same pyrolysis conditions stated in the previous two sections. 4.3. Characterization. 4.3.1. Measurement of Electrical Conductivity. Electrical resistance was measured across the two ends of cylinder-shaped samples using an electrochemistry workstation (CH Instruments). The resistance was then used to calculate the electrical conductivity. The resistance of each sample was measured three times for error analysis. 4.3.2. Measurement of Thermal Conductivity. The thermal conductivities of the samples were calculated from eq 1, where κ is the thermal conductivity (W m−1 K−1), α is the thermal diffusivity (m2 s−1), Cp is the specific heat capacity (J kg−1 K−1), and ρ is the density (kg m−3).51

κ = α ·Cp·ρ

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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/acsomega.8b00492. Dynamic light scattering and AFM characterization of GO, SEM image of a PVA foam, 1H NMR spectrum of SILRES 62C preceramic polymer, FTIR spectrum of SILRES 62C precursor before and after curing process, SEM images and EDS of SiCO/GO composites obtained at different pyrolyzed temperatures, Raman spectra of pure SILRES 62C and SILRES 62C/GO composites with 2 wt % GO pyrolyzed at different temperatures, SEM image of SILRES 62C/GO/PVA composites treated in air at 500 °C, SEM images of GO dispersed in SiCO matrix, specific heat capacities of green bodies and 1000 °C pyrolyzed SiCO/GO composites at different temperatures, tables of pyrolysis details, Young’s moduli of SiCO/GO composites, density of SiCO/GO composites obtained at different pyrolyzed temperatures, density of SiCO, Raman G peak positions of pyrolyzed composites, calculation of the room-temperature electrical conductivity of SiCO/GO composites produced from SILRES 62C/GO/PVA composites by applying different models for the electrical conductivity of composites (PDF) SiCO/GO composites obtained at 1000 °C are very stable at high temperature in air (AVI)

(1)

The thermal diffusivity was measured using the laser flash method (FlashLine 5000 by Anter Corporation). Disks with a 12.5 mm diameter and 1−2 mm thickness were made from the PDC composites before and after pyrolysis at 1000 °C. To make the disks opaque to the laser radiation, they were sputter coated with ∼15 nm platinum (Pt) followed by coating with a graphite aerosol spray. Specific heat capacity was measured separately by using differential scanning calorimetry (DSC, TA Instruments). About 10 mg of the PDC composites before and after pyrolysis at 1000 °C was used in the DSC tests. The ramping rate was 5 °C min−1 with scanning temperatures ranging from 10 to 170 °C for green bodies and 10 to 350 °C for pyrolyzed PDC composites. The acquired endothermal curves were used to calculate heat capacities at specific temperatures (Figure S11). 4.3.3. Measurement of Linear Shrinkage. Linear shrinkage was obtained by measuring the change of dimension from multiple cylinder-shaped samples after pyrolysis processes at different temperatures. 4.3.4. Measurement of Young’s Modulus. Young’s moduli of the composites were acquired by an Instron 50 kN Electromechanical Load Frame Universal Testing Machine at room temperature. For the test, 20 mm × 20 mm samples of the SILRES 62C/GO/PVA composites pyrolyzed at 1000 °C were used. The compression experiments were performed at room temperature until sample breakage was detected with a cross-head speed ranging from 40 to 86 μm min−1. Young’s modulus was calculated from the ratio of the force applied per unit of area to the strain of the sample. 4.3.5. Characterizations. The size distribution of the GO flakes was examined using dynamic light scattering (DLS, Malvern ns290). TEM images and selected area electron diffraction (SAED) patterns were obtained from a JEOL 1011 transmission electron microscope. The thickness of the GO flakes was measured by a Veeco Dimension 3100 atomic force microscope system. The proton nuclear magnetic resonance (1H NMR) spectrum of the SiCO precursor was obtained from a Bruker Avance III 400 (400 MHz) (Figure S4). The Fouriertransform infrared (FTIR) spectrum of polysiloxane precursor was investigated by a Perkin Elmer Spectrum 100 spectrometer

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. Tel: +1 (407)882-2847. Fax: +1 (407) 882-2819. ORCID

Lei Zhai: 0000-0002-3886-2154 Author Contributions

The manuscript was written through the contributions of all authors. All authors have given approval to the final version of the manuscript. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS

The authors are grateful to the support from Garmor, Inc. and NASA (NNX16AM88H) through NASA Space Technology Research Fellowships (NSTRF) and Materials Characterization Facility (MCF) at University of Central Florida. 4013


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(22) Ramakrishnan, P. A.; Wang, Y. T.; Balzar, D.; An, L.; Haluschka, C.; Riedel, R.; Hermann, A. M. Silicoboron-carbonitride ceramics: a class of high-temperature, dopable electronic materials. Appl. Phys. Lett. 2001, 78, 3076−3078. (23) Zhang, L.; Wang, Y.; Wei, Y.; Xu, W.; Fang, D.; Zhai, L.; Lin, K.C.; An, L. A silicon carbonitride ceramic with anomalously high piezoresistivity. J. Am. Ceram. Soc. 2008, 91, 1346−1349. (24) Sarkar, S.; Chunder, A.; Fei, W.; An, L.; Zhai, L. Superhydrophobic mats of polymer-derived ceramic fibers. J. Am. Ceram. Soc. 2008, 91, 2751−2755. (25) Liew, L.-A.; Saravanan, R. A.; Bright, V. M.; Dunn, M. L.; Daily, J. W.; Raj, R. Processing and characterization of silicon carbon-nitride ceramics: application of electrical properties towards mems thermal actuators. Sens. Actuators, A 2003, 103, 171−181. (26) Liew, L.-A.; Liu, Y.; Luo, R.; Cross, T.; An, L.; Bright, V. M.; Dunn, M. L.; Daily, J. W.; Raj, R. Fabrication of sicn mems by photopolymerization of pre-ceramic polymer. Sens. Actuators, A 2002, 95, 120−134. (27) Asthana, A.; Kim, K.-O.; Perumal, J.; Kim, D.-M.; Kim, D.-P. Facile single step fabrication of microchannels with varying size. Lab Chip 2009, 9, 1138−1142. (28) David, L.; Bhandavat, R.; Barrera, U.; Singh, G. Silicon oxycarbide glass-graphene composite paper electrode for long-cycle lithium-ion batteries. Nat. Commun. 2016, 7, No. 10998. (29) David, L.; Bhandavat, R.; Barrera, U.; Singh, G. Polymer-derived ceramic functionalized mos2 composite paper as a stable lithium-ion battery electrode. Sci. Rep. 2015, 5, No. 9792. (30) Oschatz, M.; Kockrick, E.; Rose, M.; Borchardt, L.; Klein, N.; Senkovska, I.; Freudenberg, T.; Korenblit, Y.; Yushin, G.; Kaskel, S. A cubic ordered, mesoporous carbide-derived carbon for gas and energy storage applications. Carbon 2010, 48, 3987−3992. (31) Zhao, R.; Shao, G.; Cao, Y.; An, L.; Xu, C. Temperature sensor made of polymer-derived ceramics for high-temperature applications. Sens. Actuators, A 2014, 219, 58−64. (32) Nagaiah, N. R.; Kapat, J. S.; An, L.; Chow, L. Novel polymer derived ceramic-high temperature heat flux sensor for gas turbine environment. J. Phys.: Conf. Ser. 2006, 34, 458−463. (33) Li, Y.; Yu, Y.; San, H.; Wang, Y.; An, L. Wireless passive polymer-derived sicn ceramic sensor with integrated resonator/ antenna. Appl. Phys. Lett. 2013, 103, No. 163505. (34) Zhao, R.; Shao, G.; Cao, Y.; An, L.; Xu, C. Temperature sensor made of polymer-derived ceramics for high-temperature applications. Sens. Actuators, A 2014, 219, 58−64. (35) Cordelair, J.; Greil, P. Electrical conductivity measurements as a microprobe for structure transitions in polysiloxane derived si−o−c ceramics. J. Eur. Ceram. Soc. 2000, 20, 1947−1957. (36) Shao, G.; Peng, W.; Ma, C.; Zhao, W.; Guo, J.; Feng, Y.; Wang, H.; Zhang, R.; et al. An, L. Enhanced electric conductivity of polymerderived sicn ceramics by microwave post-treatment. J. Am. Ceram. Soc. 2017, 100, 842−847. (37) Mamunya, Y. Carbon Nanotubes as Conductive Filler in Segregated Polymer Composites-Electrical Properties. In Carbon Nanotube-Polymer Nanocomposites; Yellampalli, S., Ed.; InTech: Rijeka, 2011; pp 10−15. (38) Shopova-Gospodinova, D.; Burghard, Z.; Dufaux, T.; Burghard, M.; Bill, J. Mechanical and electrical properties of polymer-derived si− c−n ceramics reinforced by octadecylamine-modified single-wall carbon nanotubes. Compos. Sci. Technol. 2011, 71, 931−937. (39) Ionescu, E.; Francis, A.; Riedel, R. Dispersion assessment and studies on ac percolative conductivity in polymer-derived si−c−n/cnt ceramic nanocomposites. J. Mater. Sci. 2009, 44, 2055−2062. (40) Shibuya, M.; Sakurai, M.; Takahashi, T. Preparation and characteristics of a vapor-grown carbon fiber/ceramic composite using a methylsilicone precursor. Compos. Sci. Technol. 2007, 67, 3338− 3344. (41) Xu, H.-P.; Dang, Z.-M.; Shi, D.-H.; Bai, J.-B. Remarkable selective localization of modified nanoscaled carbon black and positive temperature coefficient effect in binary-polymer matrix composites. J. Mater. Chem. 2008, 18, 2685−2690.

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