Carbon nanofibers-based nanocomposites with silicon oxy-carbide matrix

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Carbon nanofibers-based nanocomposites with silicon oxy-carbide matrix Krystian Sokolowski, Piotr Palka, Stanislaw Blazewicz, Aneta Fraczek- Szczypta∗ Faculty of Materials Science and Ceramics, AGH University of Science and Technology, Mickiewicza Av. 30, 30-059, Cracow, Poland

A R T I C LE I N FO

A B S T R A C T

Keywords: Nanocomposites Electrical properties Silicon oxy-carbides Functional applications

The subject of the study were nanocomposites made by infiltration of carbon nanofibers in the form of mat (CNF) with a poly(methylphenylsiloxane) resin (PMPS) solution followed by heat-treatment up to 1000 °C. The nanocomposites were studied using scanning electron microscopy (SEM-EDS), X-ray diffraction (XRD), Raman spectroscopy and by means of surface measurement tests. The morphology, microstructure, structure and selected properties of the nanocomposites were analysed. The influence of the CNF oxidative surface treatment on interface properties in CNF/resin and CNF/resin-derived ceramics nanocomposites was investigated. The high surface tension occurring at the nanofiber/resin interface required preliminary chemical treatment of carbon nanofiber surface. The surface chemical state of CNF had a significant influence on their interaction with the resin and formation of SiOC protective layer on nanofibers surface. The SiOC phase formed on the CNF surface at 1000 °C improved thermal stability of the nanocomposites in air. Due to the formation of the free carbon phase during resin heat-treatment the resulting ceramic nanocomposite matrix showed a lower electrical resistance compared to pure carbon mats.

1. Introduction Carbon nanomaterials, such as carbon nanotubes, graphene or carbon nanofibers, are the subject of research in terms of their applicability as components of nanocomposites [1]. Their interesting physicochemical properties make them the subject of an extensive research aimed at potential applications in modern electronics as multifunctional nanocomposites, electrochemical power sources for next generation of Li-ion batteries, flexible electrodes, fuel cells and supercapacitors [2–5]. Although carbon nanomaterials have already been applied in these areas with documented reports, improvement of their resistance to oxidation and corrosion remain an open issue [6,7]. One of the solutions for improving their oxidation resistance is to introduce a carbon component to the silicon compounds-based matrix, thus limiting the oxygen access [8–10]. The dispersion of the nanometric carbon phase in the ceramic matrix is one of the problems in obtaining homogeneous nanocomposites [11]. Searching for new ways of manufacturing nanocomposite containing silicon-based compounds and carbon, aims to develop modern energy storage systems [12], new biomaterials and scaffolds for tissue engineering [13], high-temperature sensors and micro-electromechanical systems (MEMS) [14,15]. Fibrous ceramic

nanocomposites, in the form of membranes, are particularly interesting for application in environment protection, and for anode materials for lithium-ion batteries (LiB) [16,17]. Silicon-containing polymers are often used as precursors for the manufacture of ceramic materials in the form of silicon oxy-carbide (SiOC) formed up to 1000 °C, whereas at higher temperatures (above 1500 °C) pure silicon carbide (SiC) is formed [18]. In the case of ceramic matrix composites, it is important to prepare the fiber surface to create a proper fiber-matrix interphase bonding. Serious problems associated with the use of polysiloxanes as ceramic matrix precursors are their mass loses and pyrolysis-induced shrinkage that can lead to fiber detachment and the formation of voids in the matrix [19–21]. So far, this problem has not been described in the literature for carbon nanofibers used as a nanocomponent with a polysiloxane-derived matrix precursor. The as-received CNF are rather chemically inert material and like in case of carbon microfibers, their use in nanocomposite technology requires appropriate surface preparation [22,23]. The aim of the research was to develop a method for combining CNF in the form of an isotropic mat with a polysiloxane resin, which was a precursor of the silicon oxy-carbide matrix. The influence of the surface preparation of CNF on the properties and structural parameters of the resulting CNF/silicon oxy-carbide nanocomposite was investigated. The

∗

Corresponding author. AGH - University of Science and Technology, Faculty of Materials Science and Ceramics, Department of Biomaterials and Composites, Mickiewicza Av. 30, 30-059, Cracow, Poland. E-mail addresses: [email protected] (K. Sokolowski), [email protected] (P. Palka), [email protected] (S. Blazewicz), [email protected] (A. Fraczek- Szczypta). https://doi.org/10.1016/j.ceramint.2019.09.069 Received 28 May 2019; Received in revised form 9 August 2019; Accepted 6 September 2019 0272-8842/ © 2019 Elsevier Ltd and Techna Group S.r.l. All rights reserved.

Please cite this article as: Krystian Sokolowski, et al., Ceramics International, https://doi.org/10.1016/j.ceramint.2019.09.069

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Table 1 Chemical characteristics of PMPS resin. Trade name

Lukosil 4102

Properties

Molar fraction of elements 3

Solvent

Viscosity (mPa·s)

Density(kg/m )

Drying time(min/°C)

C/Si

Si

O

C

H

Acetone

180–230

1150 ± 20

300/20

3,11

0,105

0123

0,326

0446

study focused on determining the mechanism for the conversion of polysiloxane into silicon oxy-carbide in the presence of CNF. Selected physical properties of the obtained nanocomposites were determined.

nanofibers (before carbonization). An average thickness of the obtained single-layer carbon mat (from Fig. 1b) was 18 ± 2 μm with an average porosity of 82 ± 6%.

2. Experimental

2.2. Preparation of PMPS resin-derived SiOC/CNF nanocomposites

2.1. Materials and carbon nanofibers manufacture

The CNF-containing nanocomposites were manufactured by PIP method, i.e. infiltration of the carbon mat with a liquid resin solution, followed by controlled heat-treatment to convert the PMPS resin to SiOC matrix. The mat coated/impregnated with the PMPS resin precursor was processed by complete heat-treatment at 1000 °C. The stages of the fabrication procedure of PIP-ceramic nanocomposites (Fig. 2), which were applied for obtain different types of SiOC/CNF nanocomposite, are presented in Table 2. According to the first procedure (nanocomposite denoted SiOC1/ CNF), the CNF mat was impregnated with a liquid PMPS resin solution (aerosolized resin solution). This method allowed uniform impregnation of the carbon mat without compromising its integrity. To obtain the highest possible volume fraction of CNF in the SiOC nanocomposite matrix, and facile the process of polymer impregnation, the resin was diluted with acetone (boiling point 56 °C), to obtain a 10% resin solution. The carbon nanofibers covered with the resin were left at room temperature for 24 h, to obtain the dry polymer nanocomposites. Another type of nanocomposites was obtained by the second procedure (sample denoted SiOC2/CNF), wherein the PMPS resin-coated carbon nanofibers was thermally stabilized, first at 150 °C for 18 h, and then at 250 °C for 24 h. The cross-linking conditions such as temperature and the length of cross-linking time, were selected based on previous tests [26]. Cross-linking of resin in PIP-ceramic nanocomposites was aimed at limiting the problem of polymer instability at the initial processing stage. According to Table 2, in the procedure 3 (sample denoted SiOC3/ CNF), the PAN-based carbon nanofibers were surface oxidized with liquid H2SO4/HNO3 mixture. The process was aimed at introducing additional oxygen–containing functional groups to improve adhesion between the polymer matrix and the surface of the carbon mat. The modification of CNF in acidic medium consisted of the treatment of CNF with H2SO4/HNO3 (3:1) mixture, for 60 min at 60 °C. The surface modified carbon nanofibers (MCNF) were subsequently rinsed with deionized water until they reached pH~7, and then dried at 100 °C for about 12 h. The processing conditions for the oxidation treatment of the CNF were determined in the preliminary tests. Oxidized carbon mats under these conditions were completely saturated with the resin. The effectiveness of saturation of the carbon mats with the resin solution

A poly[methyl(phenyl)siloxane] resin (PMPS) under the trade name, Lukosil L4102®, produced by Lučební Závody a.s. Kolín (Czech Republic), was used as a precursor of the ceramic matrix (SiOC). The polymer solution (78%), with acetone as the solvent was used. Selected physicochemical properties of the PMPS resin and its elemental composition are presented in Table 1. The PMPS resin chain contains aliphatic (-CH3) and unsaturated aromatic (-C6H5) groups. The properties of the final product are influenced by the C/Si ratio(for the selected resin 3.11), as it promotes the formation of free carbon in the ceramized SiOC phase. Polyacrylonitrile (PAN) copolymer under the trade name of Mavilon (Zoltek, Hungary) consisting of acrylonitrile (93–94 wt%), methyl methacrylate (5–6 wt%) and sodium allylsulfonate (1 wt%) was used as the CNF precursor. The PAN-based CNF precursors, were prepared via electrospinning method of a 11% polymer solution. The solution was obtained by dissolving dry PAN in N’N-dimethylformamide (DMF > 99.5%, POCH, Poland) with a molecular weight M = 73.10 (g/ mol), under ambient conditions (RT). In order to achieve the desired level of homogenization, 24 h mixing operation was carried out on the PAN solution. The PAN nanofibers were electrospun under the processing conditions established in previous work [24]. The PAN nanofibers (with an average diameter of single nanofibers of 277 ± 70 nm), were collected on the surface of an aluminum foil and subsequently subjected to a two-stage heat-treatment according to the previously established parameters of CNF manufacture [24,25]. The first stage consisted of the oxidation (stabilization) of PAN mats in the temperature range from 250 to 290 °C, at a constant heating rate of 4 °C/min. The oxidized PAN mats were subsequent carbonized at 1000 °C (heating rate: 15 °C/min) under a nitrogen atmosphere. The measurement of shrinkage in oxidized PAN samples (discs with a diameter of 30 mm) caused by the carbonization process to 1000 °C was about 6%. Fig. 1 shows SEM microphotographs of PAN-derived CNF and a diagram presenting the distribution of nanofiber diameters in the mat. After the carbonization process, CNF with an average diameter of 176 ± 51 nm were obtained. The diameter of the nanofibers after carbonization decreased by about 35%, compared to the stabilized PAN

Fig. 1. a, b) SEM microphotographs of carbon nanofibers and c) distribution of nanofiber diameters. 2

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Fig. 2. Schematic view of the manufacture of SiOC/CNF nanocomposites.

Ltd., England). Depth of analysis was about 5 nm. Mg Ka X-ray radiation with 200 W energy was used as the excitation source. Structural parameters of the nanocomposite samples were evaluated using X-ray diffraction (XRD- X'Pert Pro Philips diffractometer). To evaluate the content of carbon phases in the nanocomposites Raman spectroscopy (HORIBA LabRAM HR Raman Spectrometer) was used. The microstructures of PMPS/CNF and SiOC/CNF nanocomposites were studied using a SEM- JEOL JMS-5400 microscope. The diameters distribution of the nanofibers forming the mat was determined using the ImageJ 1.52i program. The electrical properties of the materials have been determined on the basis of the two-point resistance measurements. In experiments, the resistance dependence on temperature, R(T) in carbon nanofibers was studied within a temperature interval from 80 K to 323 K. A measurement system for the tests included a thermal chamber placed in a container with liquid nitrogen, adjustable heating system and measuring system (multimeter - type M 3660D, Metex). Samples in the form of strips (35 × 5 mm), cut from the mats, were fixed with a conductive silver paste on an alumina plate between electrodes (distanced by 30 mm). The chamber with the sample was heated from the temperature of liquid nitrogen to the maximum temperature (100 °C) at the rate of 5 °C/min. The dependence of R(T) was described by the Mott type of the variable-range hopping (VRH) charge transfer mechanism. This permitted identification of the differences in electrical resistances of pure CNT in the form of mats, modified with PMPS resin and heated to 1000 °C.

Table 2 Denotation of nanocomposites. Denotation of samples

Procedure

Manufacture steps (from Fig. 2)

SiOC1/CNF SiOC2/CNF SiOC3/CNF

1 2 3

1 → 3→5 1 → 3→4 → 5 1 → 2→3 → 5

was evaluated by analyzing the microscopic images of the samples. The SiOC/CNF nanocomposites were obtained by heating the CNFcontaining methylphynylsiloxane nanocomposites in a nitrogen atmosphere to 1000 °C, with a constant heating rate of 5 °C/min. Established process conditions related to the phase of impregnation of the CNF mat as well as the subsequent processing phase, i.e. stabilization and heattreatment, allowed to obtain PIP-ceramic nanocomposites with preserved integrity and shape. 2.3. Characterization The first step in the preparation of SiOC/C nanocomposites was to investigate the changes occurring after the cross-linking and heattreatment of resin and resin in the presence of CNF. Thermal analysis (TG, DTG and DSC) were made using the NETZSCH STA 449F3 thermoanalyser. The measurements were carried out under an argon atmosphere in the temperature range from 20 °C to 1000 °C at a heating rate of 10 °C/min. Structural changes after the cross-linking and heattreatment of the samples were examined by Fourier transform infrared spectroscopy (FTIR) using a Fourier FTS-60 Bio-Rad Excalibur spectrometer. The measurements were carried out in the mid-infrared range (4000-400 cm−1) using the samples preparation in powder form and the KBr/sample mixture. To improve carbon nanofibers wettability by the liquid polymer and interaction at the nanofiber-resin interface, the carbon nanofibers were subjected to surface oxidation treatment. To identify the average chemical composition and functional groups on the surface of the carbon nanofibers, XPS analysis was carried out (Vacuum Systems Workshop

3. Results and discussion 3.1. Thermal properties of the as-received and cross-linked resin An initial stage of these studies, we compared the thermal performance of two states of resin, i.e. as-received and after thermal crosslinking in the conditions discussed earlier (in section 2.2). The main goal was to answer the question whether in the process of manufacturing a ceramic nanocomposite it is necessary to initial cross-linking 3

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temperature at which the 1% weight loss occurs for as-received resin (210.1 °C for N2 and 204.3 °C for air) refers to the cross-linking process. This mechanism does not occur for cross-linked resin as evidenced by increased thermal stability in both atmospheres (426.5 °C for N2 and 394.7 °C for air). The crosslinked resin has a lower maximum rate of decomposition temperature (MRDT) at which a ceramic phase is formed. The total weight loss of the sample in N2 was 25.42% (for asreceived resin) and 24.07% (for cross-linked resin). In contrast to the inert atmosphere, the mass losses during heattreatment in the air are more intense (as-received resin = 47.92%; cross-linked resin = 44.78%). The larger ceramic residue after heattreatment of the cross-linked resin (in N2) indicates the formation of volatile decomposition products in a smaller amount that may be responsible for the formation of the porous matrix of the nanocomposite. The cross-linked resin also showed a larger residual mass after heattreatment in the air, which can affect the oxidation resistance of the nanocomposite.

Fig. 3. Calorimetric curves of PMPS resin. Heating rate 10 °C/min in a nitrogen atmosphere.

of the resin, whether this process can be omitted, and whether the cross-linked polymer matrix nanocomposite allows to increase the ceramic residue. The DSC test was aimed to examination of the extent of cross-linking of PMPS resin which directly influences on thermal effects recorded during low-temperature heat-treatment. Two-stage heating and cooling were used for the analysis during which expect crosslinking of the sample. The DSC curves, recorded from −75 °C during the first and the second heating cycles (after cooling at 10 °C/min), are shown in Fig. 3. The shapes of the DSC curves for the as-received and cross-linked PMPS resin differ. The curve for the as-received resin sample shows five exothermic effects with maxima at the temperatures of 53.7 °C, 71.7 °C, 91.4 °C, 150.6 °C and 200.3 °C. The crosslinking of the polymer causes that the exothermic effects become less visible, and the peaks shift to higher temperatures, i.e. 58.1 °C, 132.7 °C, 144.7 °C, 147.4 °C and 208.7 °C. The observed changes in the DSC curves can be explained by the decrease in the mobility of the elastic Si-O-Si chain segments and the increase in their stiffness as a result of the cross-linking of the PMPS resin structure [27,28]. The thermal effects visible on the thermograms may also be attributed to depolymerization followed by relaxation of the chain segments. Such a mechanism involves internal or intermolecular arrangements in the polymer structure activated by rotation and polarization of the Si-O bond, and results in the formation of an intermediate product without loss of mass [28–30]. The width and shape of the heat energy changes, recorded for the cross-linked resin, suggest that they may originate from the breakdown of non-crosslinked resin and subsequent relaxation of its structure. Differences in the shape of the calorimetric curves are also visible during the second heating cycle after the sample has been cooled (recrystallized) to about −75 °C. Due to ordering crystalline domains, the thermal effects initially visible on the thermogram disappear and the remaining peaks are flattened and shifted to higher temperatures (e.g. up to 84.0 °C and 204.2 °C) for the as-received polymer structure. The influence of nitrogen and air atmosphere on thermal stability and kinetics parameters of decomposition of the both states of resin (asreceived and cross-linked) was investigate via TGA analysis up to 1000 °C. As shown in Table 3, thermal decomposition parameters and ceramic residues are different for the both states of resin. The

3.2. Surface treatment of carbon nanofibers Preliminary experiments of impregnation of the as-received carbon mats with a resin solution have shown that single nanofibers in the mat are covered with the resin only partially, close to the surface of the mat. To improve the wettability of carbon mats with a liquid polymer and to achieve full supersaturation of the carbon mat, the samples were subjected to oxidation treatment. The XPS method was used to measure the chemical composition of CNF and MCNF. Table 4 compares the chemical compositions of the CNF before and after the oxidative treatment. Elemental analysis of untreated CNF showed, that besides carbon (88.9 at.%), they contain also some amount of oxygen (7.8 at.%) and nitrogen (3.3 at.%). Their presence in the CNF results from the chemical composition of the PAN nanofiber precursor, as well as from its stabilization process (PAN in the form of mat), which was carried out in the air. After the oxidative treatment changes in the chemical composition of CNF (MCNF) can be observed. The carbon content decrease to 72.5 at.% whereas the oxygen content increases to 21.6 at.% and the nitrogen content increases from 3.3 at.% to 5.8 at.% (Table 4). The structural changes accompanying the process were assessed by FTIR. The set of FTIR spectra for the as-received CNF and oxidized MCNF are shown in Fig. 4. The bands of CNF spectra at 1245 and 1188 cm−1 are derived from the C-C stretching vibrations, and the band at 862 cm−1 is derived from the vibrations of the C-H groups in the aromatic ring. At higher wave numbers, i.e., at 2921, 2852, the C-H bands brought about by the CH2/CH3 stretching vibrations are visible. Within the wavelength range of 3660–2970 cm−1, the O-H oscillation vibration groups also appear [31]. After treating the CNF in H2SO4/ HNO3 mixture, the IR spectrum is similar to the spectrum of unmodified CNF. There are a number of slightly shifted bands in the spectrum, i.e., 1596, 1261, 1188 and 869 cm−1, all of which are derived from the vibrations of the C=C, O-H, C-C and C-H groups, respectively [32–34]. It is worth noting the appearance of a 1770 cm−1 band, that is typical for the band attributed to the C=O stretching vibrations of carboxylic/lactone groups. Moreover, the band at 1396 cm−1 was due to sulphate groups ν(OSO3H) and δ(OH) bending vibration of COOH. The appearance of C=O and OSO3H bonds suggests that oxidation of the

Table 3 Thermal and thermo-oxidative stability of resins obtained in nitrogen and air atmospheres. Sample designation

As-received resin Cross-linked resin

Nitrogen atmosphere parameters(°C)

Air atmosphere parameters (°C)

Residue yield at 1000 °C (wt.%)

T1%

T5%

MRDT

T1%

T5%

MRDT

Inert atmosphere

Air atmosphere

210.1 426.5

279.4 463.8

501.4 457.5

204.3 394.7

287.8 469.7

607.7 549.7

74.6 75.9

52.1 55.2

4

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Optimal conditions for the CNF resin impregnation process, ensuring full saturation of carbon nanofibers was determined on the basis of SEM images. Fig. 6 shows carbon mats covered with different amounts of PMPS resin. SEM micrographs (Fig. 6) show the polysiloxane-based nanocomposites containing from 10 to 30 wt% of CNFs. With the increase in the resin content, a decrease in porosity and an increase in the density of the nanocomposite (PMPS/CNF) was observed. When the polymer's weight exceeded 80 wt%, the fibrous carbon shape was less visible (Fig. 6(c,d)), as a result of the non-wetting PMPS resin layer forming on CNF surface. The ability to infiltration the CNF mats with PMPS resin was the same in procedures (1) and (2). In the case of the procedure (3), samples with a resin content greater than 80 wt% showed an increased susceptibility to cracking and brittleness. As a result of these experiments it was established that the 80 wt%PMPS/ 20 wt%CNF nanocomposite was the optimal composition to obtain SiOC nanocomposite for each procedure. The SiOC/CNF nanocomposites have been prepared by cross-linking the resin followed by heating to 1000 °C under a nitrogen atmosphere. To determine the possible impact of the carbon nanofibers on PMPS resin conversion to the ceramic phase, the CNF-containing PMPS nanocomposites were subjected to thermal analysis up to 1000 °C (heating rate: 10 °C/min) (Fig. 7). The TGA curves of the PMPS resin registered during heat-treatment to 1000 °C in nitrogen atmosphere show two characteristic inflections. The first one appears at about 250 °C and is associated with a condensation reaction, leading to cross-linking of the polymer network, which is accompanied by a weak DTG effect (designated as “a" in Fig. 7). Absence of this inflection for the cross-linked resin indicates a well stabilized polymer structure. The second inflection occurs at the temperature range of 450–700 °C (marked as “b" in Fig. 7) and is assigned to the degradation of the organic side groups. In this stage of heat-treatment the highest weight loss of the sample was observed. These results suggest that the polymer's mass loss during heat-treatment to 1000 °C takes place in the three stages of thermal decomposition. At the final step of heat-treatment (above 800 °C), both TG curves for pure resins exhibit similar plato without further mass changes indicating completed decomposition. As can be seen from Fig. 7, the PMPS resin losses significantly more weight (25.42%), as compared to the CNF-containing resin samples (21.38%), and to the cross-linked resin (24.07%) during heat-treatment up to 1000 °C. The presence of CNF in nanocomposite samples affects the resin's ceramization process, forming the SiOC phase with a greater residue. This is probably due to the formation of interfacial boundaries between the surface of nanofibers and the resin. Due to the presence of oxygen-containing groups on the surface of CNF, the bonds at the interface are probably chemical in nature. The heat-treatment of resincoated CNFs is accompanied by an increase in the weight percentage of nanofibers resulting from the transformation of the polymer matrix into a SiOC-based matrix. Thus, it can be assumed that the mass losses of the resin/nanofiber samples are mainly related to the resin decomposition, as the carbon nanofibers themselves do not change their mass during heating up to 1000 °C. Quantitative analysis showed that the weight fraction of nanofibers in SiOC/CNF nanocomposites was 40%, while in the cross-linked resin-based nanocomposite around 39%. Contrary to the TG curves for pure resins, the Tg curves for nanocomposite samples indicate that the pyrolysis process to 1000 °C has not been completed.

Table 4 Elemental composition date of CNF and MCNF obtained from XPS analysis. Samples

As-received CNF Modified CNF (MCNF)

Atomic concentration (at.%) C1s

O1s

N1s

O/C

N/C

88.9 72.5

7.8 21.6

3.3 5.8

0.09 0.30

0.04 0.08

Fig. 4. FT-IR spectra of as-received carbon nanofibers and modified with oxidizing agent.

CNFs successfully introduced COOH, OH and OSO3H groups onto the surface of the carbon nanofibers [33,34]. The intensity of the wide band in the range of 3660–2970 cm−1 from O-H bonds does not show any noticeable changes [32–34]. The analysis of MCNF showed that the amount of acidic groups after the oxidative treatment is 4.85·10−4 mol g−1, and is distinctly higher in comparison to untreated carbon nanofibers (1.24·10−5 mol g−1). The increase in the number of functional groups on the surface of modified carbon nanofibers is also confirmed by the results of the XPS analysis. Fig. 5 shows the signals (areas) at binding energy 284.6 eV characteristic for C1s and at 532 eV for O1s. The main band for asreceived CNF at 284.6 eV is derived from sp2 carbon atoms occurring in the form of C=C or C-C bonds (Fig. 5a) [35]. The deconvolution of the C1s band shows that carbon atoms are also present in oxygen-containing groups, such as C-OH/C-O-C, C-O, C=O and O=C-O, which have higher binding energies, i.e. 285.8 eV, 286.8 eV, 288.2 eV, and 290.3 eV, respectively. In the as-received CNF, three oxygen (O1s) bonds (Fig. 5b) can be distinguished, i.e. at 530.1 eV corresponding to group C=O, at 532.2 eV corresponding to functionalized oxygen-containing groups, such as C-OH or C-O-C, at 533.7 eV assigning moisture [35,36]. Considering the XPS spectra for modified carbon nanofibers (MCNF) (Fig. 5c and d) almost identical energy configurations of these bonds can be observed in the C1s (280–294 eV) and O1s (526–538 eV) regions. Deconvolution of the C1s band showed an increase (in percentage) of C-O groups from 21.4% (for as-received CNF) to 32.7% (for MCNF), while C=O groups from 4.7% (as-received CNF) to 11.5% (for MCNF). The number of functional groups on the surface of the nanofiber after oxidation (MCNF) increases, which may indicate a typical process of oxidation of CNF as observed for other types of carbons, e.g. carbon fibers and carbon nanotubes [35–37].

3.4. Mechanism of SiOC formation on carbon nanofibers surface Fig. 8 presents SEM microphotographs and EDS spectrum of the SiOC2/CNF nanocomposites obtained at 1000 °C. The pure CNF mat (Fig. 8a) contains pores into which a liquid polymer solution can penetrate. After impregnation of CNF with PMPS resin according to procedure 2, followed by heat-treatment to 1000 °C, the nanofibers are covered with SiOC. The cross-section of the coated carbon nanofibers in Fig. 8b shows that the layer depth is about 6.5 μm. Partial penetration

3.3. Heat-treatment of the CNF-containing poly(methylphenylsiloxane) nanocomposites The PMPS-based nanocomposites containing carbon mats were the starting material for the formation of a SiOC/CNF nanocomposite. 5

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Fig. 5. High-resolution XPS spectra of carbon nanofibers before and after surface treatment. Deconvoluted C1s (a, c) and O1s (b, d).

Fig. 6. SEM images of PMPS/CNF nanocomposites containing: a) 30 wt%, b) 20 wt%, c) 15 wt% and d) 10 wt% of CNF.

pores between the nanofibers. SEM study of samples showed an increase in the average diameter of the carbon nanofiber (about 53%) after the heat-treatment process, which is associated with the formation of SiOC layers, with an average thickness of about 47 nm on carbon nanofibers. The microanalysis of the EDS surface showed that the ceramic coating contains C, Si and O (Fig. 8d). After impregnation of carbon nanofibers with the polysiloxane solution according to procedure 1 (Table 2) and further heat-treatment to 1000 °C, the SiOC ceramic layers covered the surface of the mat to a depth of about 11 μm (Fig. 9(c,e)). The SEM image (Fig. 9a) clearly shows the fibrous morphology of a SiOC1/CNF nanocomposite with an average diameter of a single nanofiber of about 206 nm. At higher magnification (Fig. 9e), it was observed that part of the carbon nanofibers was detached from the SiOC matrix. The debonding of the nanofibers was probably caused by the lack of polysiloxane-derived SiOC binder due to poor resin infiltration into the carbon mat. In order to improve the interaction of the resin with the carbon surface, the CNF surface was prepared according to procedure 3. After impregnation of the modified nanofiber mat with a liquid resin followed by heat-treatment at 1000 °C, nanofibers were completely covered with SiOC (Fig. 9(b,d,f)). In this case, the MCNF surfaces contained increased concentration of oxygen functional groups that ensured wetting and full depth infiltration of carbon mat by resin solution, and

Fig. 7. TG/DTG curves recorded during heat-treatment PMPS-impregnated carbon nanofibers to 1000 °C.

of the resin into the carbon mat results from the high surface tension of the solution at the CNF/resin interface. The SEM morphology (Fig. 8c) of the surface of the SiOC2/CNF nanocomposites reveal their fibrous form. The transformation of polymer matrix nanocomposites into SiOC-matrix nanocomposites occurs during heat-treatment, which is accompanied by the formation of 6

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Fig. 8. SEM images of a) pure CNF, b, c) SiOC-coated CNF (procedure 2), d) EDS spectrum of the average chemical composition of the nanocomposite.

vibrations of the -Si-O-Si- groups in the main siloxane chain. Bands at 1450 cm−1, attributed to the deformation vibrations of silicon and the side aliphatic group (Si-CH3), can also be observed. The Si-C6H5 stretching vibrations of aromatic phenyl substituents occurs at 1413 cm−1 and is additionally confirmed by the wide band at 3070 cm−1, typical for the C-H bonds in the aromatic ring [38,39]. The spectra of pyrolysates derived from pure resin and from CNF-based nanocomposites are typical for amorphous materials. The SiOC/CNF samples for FTIR study were prepared by the KBr pellet method. The remaining samples were obtained using the ATR-FTIR technique with ATR correction preceding the spectral measurement. The band at 3400 cm−1 is derived from the carbon dioxide background as a result of sample's contact with atmosphere. The spectra of precursors heat-treated at 1000 °C contain a series of bands related mainly to the stretching vibrations of Si-O and Si-C molecules. They indicate that SiOC bonds are formed in the sample. The band registered at the 440 cm−1 derives from the Si-O stretching vibrations, whereas the band at 796 cm−1 derives from the Si-C stretching vibrations. Due to the range of heat-treatment of the samples studied and the specific nature of polymer-ceramic conversion process, its presence is attributed to Si-CH3 groups. The intense, wide band within the wavelength range of 1100–1000 cm−1, is caused the stretching bonds in Si-O-Si groups. The elementary analysis of a polysiloxane resin shown in Table 1 indicates that band, corresponding to 1607 cm−1, can be assigned to C=C stretching vibrations from carbon phase. Its occurrence is associated with the turbostratic structure of free carbon phase formed during the heat-treatment of the resin-based nanocomposites [39,40]. Bands within the 3700-3650 cm−1 wavenumber range, assigned to stretching of aromatic phenyl groups -C6H5 are also seen in the spectra [39]. The residues obtained from pure resin and the SiOC1/CNF samples were examined by X-ray diffraction. The diffractograms are shown in Fig. 12. As presented in Fig. 12, the XRD of the heat-treated polysiloxane-based samples contain broad band resulting from the

subsequently the ceramic binder was formed during the heat-treatment. The different ability to infiltration of PMPS resin in a CNF mat is caused by, among others surface tension arising at the fiber/polymer boundary. The use of resin cross-linking (in procedure 2) as compared to the treatment of the oxidizing nanofiber (procedure 3) does not change the mechanism of interaction of the resin with the carbon nanofibers. The increase in the amount of oxygen functional groups on the surface of the MCNF beneficially affects the chemical state of the surface, which improves the hydrophilic properties and infiltration of the CNF mat by the resin solution. SiOC-coated carbon nanofibers topography was examined by atomic force microscopy (AFM) (Fig. 10). The surfaces of nanocomposites obtained in procedures 1 and 2 show slight differences in topography (Fig. 10(a, b)). SiOC derived from polysiloxane covers the CNF mats at the low depth. The greater accumulation of SiOC on the surface of these samples results from the limited infiltration of the PMPS resin in the CNF mat at the stage of obtaining polymer nanocomposites. The topography of the nanocomposite obtained in the procedure 3 differs from the others (Fig. 10c). In this case, nanofibers having oxygen functional groups (MCNF) have better infiltration properties of the PMPS resin. Thus, the SiOC homogeneously covers the surface of carbon nanofibers over the entire depth of the MCNF mat, thanks to which phase separation is not observed as in the case of samples obtained in procedures 1 and 2.

3.5. Structure evolution during processing of PMPS-based nanocomposite In order to determine the composition and the structure of the ceramic matrix, the silicon oxycarbide-based nanocomposites obtained by procedure 1 (Table 2) were studied by FTIR spectroscopy. The FTIR spectra (Fig. 11) also reveal the changes taking place in the precursor's structure after the thermal decomposition process at 1000 °C. The cured polysiloxane spectrum contains bands in the range of 1100–1000 cm−1 and 600-520 cm−1. The bands are brought about by the stretching 7

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Fig. 9. SEM micrographs of the SiOC-based nanocomposites obtained according to procedure 1 (a, c, e) and procedure 3 (b, d, f).

curves recorded during oxidation of the SiOC-coated carbon nanofibers at a constant heating rate of 10 °C/min, are shown in Fig. 14. The as-received CNF were used as a reference sample. The TG curves indicate that the onset of as-received CNF oxidation starts at about 370 °C, and their complete oxidation, with the formation of a solid residue (13 wt%), takes place at 1000 °C. The SiOC/CNF nanocomposites, show a significant improvement in oxidation resistance. The mass losses and initial oxidation temperatures changed, depending on the CNF surface preparation before impregnated with PMPS resin. The nanocomposite samples impregnated with cross-linked PMPS according to procedure 2, lost 39% of the initial mass. The relatively lower resistance to oxidation observed for SiOC nanocomposites obtained using the primary PMPS resin (procedure 1) is caused by its decomposition mechanism (Table 3), resulting in a significant higher porosity (52.7%) compared to nanocomposites obtained from the crosslinked resin matrix. These losses are related to the oxidation of free carbon formed in SiOC matrix as well as to the oxidation of the CNF mat, by gaseous oxygen diffusing through the porous SiOC matrix. The study shows that the nanocomposites containing pre-oxidized CNF (procedure 3) have lower thermal stability in air as compared to the nanocomposite obtained according to procedure 2. This is probably caused by the too strong initial adhesion between the polymer layer and the surface of functionalized CNF. When heating the nanocomposite, the polymer matrix shrinks as a result of pyrolysis combined with the release of gaseous products. Thermal stresses are created at the

formation of a residue of amorphous character. By deconvolution of the band, which appears within the 2θ 20–30° angle range, by means of the Gaussian function, the presence of two turbostratic carbon phases [bC] at 23.0° and 25.4° were identified. These phases were assigned to the carbon nanofibers and the Cfree residue deriving from the heat-treated resin [41,42]. In order to determine the content of free carbon in the obtained SiOC1/CNF nanocomposite, a Raman spectroscopic analysis was performed. The Raman spectrum (Fig. 13) of the SiOC ceramic-coated carbon nanofibers (procedure 1), shows two peaks associated with the presence of the carbon phase: the so-called D-band at 1350 cm−1, corresponding to disordered structure of carbon, and the G-band (graphite structure) at 1580 cm−1, typical for higher-order carbon phase. The intensity ratio of D-peak and G-peak (ID/IG), was 0.99 for CNF and 1.04 for SiOC1/CNF. A higher ID/IG ratio for the SiOC/CNF nanocomposite results from the formation of the turbostratic carbon [bC] during the processes of ceramization of polysiloxane and phase separation, as indicated by a slight increase in the intensity of D peak at 1350 cm−1.

3.6. Oxidation resistance of nanocomposites The lifetime of SiOC/CNF nanocomposites under thermo-oxidative conditions was determined on the basis of thermogravimetric studies (TGA) conducted in air at the temperature range of 20–1000 °C. The TG 8

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Fig. 10. AFM topography (2D and 3D) images of the SiOC/CNF nanocomposites obtained according to procedure 1 (a - CNF, and resin without cross-linking), procedure 2 (b – CNF, and cross-linked resin) and procedure 3 (c – oxidized CNF and resin without cross-linking).

3.7. Electrical conductivity of nanocomposites Low-temperature measurements of electrical properties are a valuable tool in material research. Lowering the temperature of a sample allows observation of phenomena that are normally masked by thermal interactions. Electrical conductivity measurements are sensitive to the amount and distribution of free carbon phase crystallizing during polymer heat-treatment resin in the SiOC matrix. The SiOC matrix precursor was a polysiloxane polymer with a high C/Si ratio (3.11). Thus, the heat-treatment residue contains a certain amount of free carbon can affect the electrical conductivity of CNF in the form of mats. The relationship between the resistance and the temperature of samples is shown in Fig. 15. The electrical properties of carbon mats depend on a number of factors, including the type of raw material used for their manufacture, the method of their treatment and the final processing temperature [43,44]. Carbon nanofibers made from the PAN precursor have a similar structure to the carbon fiber structure and are composed of crystallites of various degree of structural ordering [24]. Their turbostratic structure is characterized by a high intrinsic lattice disorder and to describe their electrical properties, models for strongly disordered structures can be applied, including the Mott variable-range hopping (VRH) conduction model [45,46]. Such materials usually show a marked increase in electrical resistance at low temperatures, which is caused by the electrical charge localization on the conductive elements of the structure. Thus, the factors favoring the VRH charge transfer are connected with decreasing temperature and increasing degree of the disorder [47]. The function R(T) of such a material can then be described by the following expression:

Fig. 11. FT-IR spectra of PMPS resin-derived residues and PMPS resin-based nanocomposites after heat-treatment to 1000 °C.

interphase of both composite components, which may lead to partial separation of the forming ceramic phase and the creation of damages that facilitate diffusion of oxygen to the carbon component. A similar mechanism occurs when the bond between carbon fiber and the carbon matrix is too strong during C/C composite manufacturing [21,23]. The results of stability tests in air indicate the need for further optimization of the CNF surface functionalization process to improve the thermal resistance of SiOC matrix nanocomposites.

RT = RS ⋅exp 4

T0M T

(1)

where, RT – material resistance at temperature T, [Ω]; RS – constant; T0M – excitation energy of charge carriers on hopping in the Mott model [K]; T – temperature [K]. After standardization of variable RT, equation (1) may be expressed, as follows: 9

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Fig. 12. a) XRD patterns of carbon nanofibers, polysiloxane, and SiOC/CNF nanocomposite (procedure 1), b) Gaussian deconvolution of the (002) band into two separate bands.

Fig. 13. Raman spectra for carbon nanofibers and SiOC1/CNF nanocomposite.

Fig. 15. Temperature dependence of electrical resistance for pure CNF, and CNF-based nanocomposites (from procedure 1).

Fig. 14. wt residue of the SiOC-matrix nanocomposites during oxidizing up to 1000 °C in air.

Fig. 16. Temperature dependence of electrical resistances for the Mott VRH conduction for pure CNF and modified with resin; the lines are linear fits.

R R ln ⎛ T ⎞ = ln ⎛ S ⎞ + R 0 ⎝ ⎠ ⎝ R0 ⎠ ⎜

⎟

⎜

⎟

4

T0M ⋅ 4

1 T

Table 5 Parameters of samples obtained from electrical measurements.

(2)

where, R0 is the resistance of sample at 80 K. The plots R(T) and their linear fitting are shown in Fig. 16. As can be seen from this figure the function R(T) of all samples is similar, exhibiting a distinctly activated behavior within the whole temperature range, available for the resistance measurements. The plots display broad linear intervals, related to the Mott VRH conduction according to equation (1). The values of T0M can be determined from the slopes of these plots. The temperature intervals, ΔT, of the Mott VRH charge transfer, were obtained with the linear intervals of the

Sample

T0M [eV]

R2

RRT [Ω]

RRT/R0

CNF PMPS1/CNF SiOC1/CNF

0.019 0.011 0.007

0.986 0.991 0.993

937.0 547.0 108.9

0.560 0.598 0.674

corresponding plots. The values of T0 determined for the temperature interval from 80 to 323 K, as well as correlation coefficients, R2, for functions fitting are gathered in Table 5. Table 5 shows that the highest 10

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advantageous properties, due to the uniform surface distribution of the preceramic polymer. Further research on the developed nanocomposite will refer to the evaluation of its functional properties as a new form of material for electrical energy storage application.

increase in resistance is observed for CNF, then for polymer-impregnated nanofibers, and for SiOC1/CNF samples. Additionally, it can be seen that the excitation energy, T0M decreases in the following order: CNF, PMPS1/CNF, and SiOC1/CNF. This suggests that the addition of a polysiloxane resin facilitates the charge transfer process. The mat composed of pure carbon nanofibers changes its resistance in a wider range than the composite mat (Table 5, RRT/R0). The addition of a resin to the nanofibers and subsequent heat-treatment cause a reduction in the variability of the resistance with a change in temperature. The excitation energy values (T0M) required for the electron hopping between the crystallites in the carbon nanofibers are higher, compared to the composite samples i.e., PMPS1/CNF and SiOC1/CNF. By combining the CNF with the polymer matrix the electrical resistance at room temperature (RRT) of these nanocomposites is lower and T0M is reduced as compared to pure CNT. This behavior can be explained by an increase in electrical contacts between individual carbon nanofibers in the mat. Although the resin itself is an electrical insulator, ρRT > 1014 (Ω∙cm) [48], due to the increase in electrical contacts between individual carbon nanofibers and creating new conductive paths in the carbon mat, the temperature dependence of the resistance changes. Even more pronounced changes in resistance with temperature increase are observed for nanocomposites obtained at 1000 °C. In this case, the polymer precursor is converted to SiOC structure accompanied by a precipitation of free carbon phase. Excessive carbon nucleates in the form of clusters called basic structural units (BSU). The formation of continuous, turbostratic carbon lattice promotes percolation and improves the electrical conductivity of SiOC1/CNF. Its presence is reflected in the further increase in electrical contacts in the carbon mats as compared PMPS1/CNF. The initial resistance (RRT) of such a material is significantly lower compared to the resistance of the samples containing pure carbon nanofibers and is reduced with increasing temperature (Table 5). CNF are today the subject of extensive research of their potential applications in technique and medicine. A particularly interesting field of research is the use of this fibrous form in membrane and filtration techniques in which high thermal and chemical stability is required. The study carried out in this work has shown that the use of polysiloxane resins as precursors of carbide compounds for CNF modification may be one of the possible approaches to solve the problem of improving their resistance to oxidation. The use of polysiloxane resins allows, on the one hand, to form continuous carbide layers on the CNF substrates, retaining the high-porous microstructure of the initial carbon form, i.e. mats, and on the other hand, the proposed approach allows obtaining flexible SiOC/C nanocomposites. Although the experiments in this work were not conducted to obtain non-porous nanocomposites, their manufacture is possible by repeatedly impregnating the carbon mat. The research has shown that an important factor in the impregnation phase of the porous carbon structure is the appropriate selection of polymer concentration in the solution and the chemical state of the carbon surface. The research also indicates that the effective impregnation of the porous carbon mat requires pre-treatment of the surface of the carbon mat. It is to be expected that subsequent impregnating phases may be carried out using a solution with a higher polymer concentration. Another interesting result of the research is the demonstration of the effect of free carbon formed in the silicon oxy-carbide matrix on the electrical properties of the modified carbon mat. The study has shown that the free carbon phase forming during the resin heat-treatment increases the contact between single carbon nanofibers in the mat and consequently improves the electrical conductivity of porous SiOC/C nanocomposites. Compared to the method of electrospinning of silicon oxy-carbide precursors, the method of CNF impregnation with the Sicontaining polymer allows the manufacture of fibrous ceramic nanocomposites with an increased carbon fraction and potentially more

4. Conclusions We applied PIP-method for the manufacture of polysiloxanes-derived SiOC/CNF nanocomposites by impregnating the carbon nanofibers mat with a liquid solution of poly[methyl(phenyl)siloxane] resin (PMPS) followed by heat-treatment up to 1000 °C. The surface chemistry of CNF had a significant influence on the interaction with PMPS resin and thermal stability of the SiOC-coated carbon nanofibers. The study has shown, that the impregnation of CNF mat with resin solution, without prior functionalization of the surface of carbon nanofibers, is inhibited due to too high surface tension occurring between the carbon surface and the liquid polymer phase. Carbon nanofibers cannot contain too high concentration of oxygencontaining functional groups that form a strong bond with the polymer matrix, because the heat-treatment process can lead to the detachment of the layer and the creation of additional routes for easy diffusion of oxygen to the carbon surface. Study of electrical properties has shown the semiconductor nature of protective SiOC coatings. Resin-impregnated carbon mat reduces resistivity due to the creation of new electrical contacts between the carbon nanofibers. Ceramic nanocomposites consisted from the CNF and SiOC matrix have a lower electrical resistance compared to pure nanofiber mats due to the formation of the free carbon phase during heat-treatment the PMPS resin. Declarations of interest The authors declare that they have no conflict of interest. Acknowledgments This work was financed from the statute funds of AGH University of Science and Technology, Faculty of Materials Science and Ceramics, Project no. 11.11.160.182. References [1] I. Ahmad, B. Yazdani, Y. Zhu, Recent advances on carbon nanotubes and graphene, Nanomaterials 5 (2015) 90–114, https://doi.org/10.3390/nano5010090. [2] B. Zhang, F. Kang, J. Tarascon, J. Kim, Progress in Materials Science Recent advances in electrospun carbon nanofibers and their application in electrochemical energy storage, Prog. Mater. Sci. 76 (2016) 319–380, https://doi.org/10.1016/j. pmatsci.2015.08.002. [3] B. Kim, K. Seung, H. Woo, K. Oshida, Supercapacitor performance of porous carbon nanofiber composites prepared by electrospinning polymethylhydrosiloxane (PMHS)/polyacrylonitrile (PAN) blend solutions, Synth. Met. 161 (2011) 1211–1216, https://doi.org/10.1016/j.synthmet.2011.04.005. [4] Q. Song, H. Yan, K. Liu, K. Xie, W. Li, W. Gai, G. Chen, H. Li, C. Shen, Q. Fu, S. Zhang, L. Zhang, B. Wei, Vertically grown edge-rich graphene nanosheets for spatial control of Li nucleation, Adv. Energy Mater. 8 (2018) 1800564, https://doi. org/10.1002/aenm.201800564. [5] Q. Song, F. Ye, X. Yin, W. Li, H. Li, Y. Liu, K. Li, K. Xie, X. Li, Q. Fu, L. Cheng, L. Zhang, B. Wei, Carbon nanotube-multilayered graphene edge plane core-shell hybrid foams for ultrahigh-performance electromagnetic-interference shielding, Adv. Mater. 29 (2017) 1800564, https://doi.org/10.1002/adma.201701583. [6] H. Oh, K. Kim, Y. Ko, H. Kim, Effect of chemical oxidation of CNFs on the electrochemical carbon corrosion in polymer electrolyte membrane fuel cells, Int. J. Hydrogen Energy 35 (2010) 701–708, https://doi.org/10.1016/j.ijhydene.2009.10. 105. [7] T. Yuanjian, W. Xiaoqian, S. Hua, X. Lianghua, Oxidation kinetics of polyacrylonitrile-based carbon fibers in air and the effect on their tensile properties, Corros. Sci. 53 (2011) 2484–2488, https://doi.org/10.1016/j.corsci.2011.04.004. [8] B. Kim, C. Hyo, K. Seung, K. Kim, Y. Lee, SiC/SiO2 coating for improving the oxidation resistive property of carbon nanofiber, Appl. Surf. Sci. 257 (2010) 1607–1611, https://doi.org/10.1016/j.apsusc.2010.08.104. [9] K. Sokolowski, A. Fraczek-Szczypta, J. Tomala, S. Blazewicz, Organosilicon resin-

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