Electromagnetic interference shielding effectiveness of carbon foam containing in situ grown silicon carbide nanowires

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Electromagnetic interference shielding effectiveness of carbon foam containing in situ grown silicon carbide nanowires Shameel Farhan a,n, Rumin Wang a,n, Kezhi Li b a b

Department of Applied Chemistry, School of Science, Northwestern Polytechnical University, Xi'an, Shaanxi 710072, China School of Materials Science and Engineering, Northwestern Polytechnical University, Xi'an 710072, China

art ic l e i nf o

a b s t r a c t

Article history: Received 26 March 2016 Received in revised form 10 April 2016 Accepted 11 April 2016

Electromagnetic interference (EMI) shielding effectiveness of the carbon foam containing different amounts of silicon carbide nanowires (SiC NWs) was determined over X-band (8.2–12.4 GHz). Carbon foam was prepared from powdered precursors containing polyurethane, novolac and pitch. In situ SiC NWs were grown inside the porous body by using different weight ratios of silicon powder ranging from zero to 20 wt% and pyrolyzing at 1500 °C under argon atmosphere. Decorated nanowires led to a different cellular structure compared to pristine carbon foam and improved the absorptivity of EM radiations due to a very high level of dielectric loss (loss tangent of 3.1 at 12.4 GHz). A maximum EMI specific shielding effectiveness of
79.50 dB cm3/g (at 8.2 GHz) was achieved in the carbon foam containing the highest amount of SiC NWs, which was more than two times higher than that of the pristine carbon foam. Presence of NWs with stacking faults, heterostructure interfaces and long tortuous paths in the carbon foam caused a strong dielectric loss and induced higher dielectric permittivity. The study shows that by controlling the amount of SiC NWs, it is possible to improve multiple properties while achieving lightweight material for stealth technology. & 2016 Elsevier Ltd and Techna Group S.r.l. All rights reserved.

Keywords: Carbon foam In situ SiC NWs Dielectric properties Microstructure Electromagnetic shielding (absorber)

1. Introduction Carbon foam is a promising candidate for different applications owing to its several fascinating properties such as low density, large surface area, controlled porosity, good thermal/electrical properties and mechanical stability [1–4]. There is a tunable range of thermal and electrical properties of carbon foams derived from thermosetting and thermoplastic polymers by various techniques [5–7]. Based on their highly porous interconnected framework, such foams were found to have many promising applications as electromagnetic interference (EMI) shields or microwave absorber. However, it is not easy to find extensive literature about EMI shielding effectiveness (SE) of carbon foams. Yang et al. studied the microwave (2–18 GHz) absorption characteristics of mesopitch derived carbon foam and found better microwave absorption properties with less reflection losses [8]. Wang et al. prepared a dense mesoporous carbon and fused silica composite and found EMI absorption was much higher than reflection in the X band [9]. Zhou et al. also designed carbon composites C–SiO2–Fe as novel EMI absorbers that showed considerable reflection loss in the n

Corresponding authors. E-mail addresses: [email protected] (S. Farhan), [email protected] (R. Wang).

frequency range of 2–18 GHz [10]. Moglie et al. studied the EMI SE of carbon foam in the frequency band of 1–4 GHz and found that the total SE increased with a rise in thickness [11]. Blacker et al. developed electrically graded carbon foam whose electrical resistivity increased with increasing the thickness and can be used as radar absorbers [12]. Kuzhir et al. reported the EMI SE of bio based carbon foam exceeding 23 dB in the 26–40 GHz band for a very low density of 0.150 g/cm3 [13]. Zhang et al. developed novel syntactic foam using carbon microspheres in resole resin and used carbon nanofibers as fillers. A maximum EMI shielding of 25 dB was achieved for the composite foam containing 2.0% filler [14]. For the effective EMI shielding systems, being lightweight is an important technological requirement. Very recently, we prepared a new class of carbon foams from powdered precursors containing polyurethane, novolac and pitch [15,4], thus opening up new avenues for a multifunctional porous carbon produced on a fullyenabled large scale and large sizes. Obviously, this carbon foam possesses unique feature of in situ grown nanowires and nanomaterials. Consequently, it is envisioned as playing an important role in various kinds of applications, and thus there is a dire need to investigate further in order to explore and completely exploit its advantageous properties especially EMI SE. EMI shielding refers to the reflection and/or absorption of EM radiations by a shielding material. Reflection is by far the simplest and historical shielding mechanism and in order to perform

http://dx.doi.org/10.1016/j.ceramint.2016.04.054 0272-8842/& 2016 Elsevier Ltd and Techna Group S.r.l. All rights reserved.

Please cite this article as: S. Farhan, et al., Electromagnetic interference shielding effectiveness of carbon foam containing in situ grown silicon carbide nanowires, Ceramics International (2016), http://dx.doi.org/10.1016/j.ceramint.2016.04.054i

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shielding, the material with high electrical conductivity must interact with EM field in radiation [16]. Different metals and their nanoparticles have been studied as EM shielding materials or coatings; however, they have drawbacks of high density, corrosion and costly processing [17]. These reflected EM radiations may interfere with the other surrounding devices and electronic equipment. Therefore, materials with a larger absorption (4 25 dB) and a lower reflection (o 10 dB) are needed. Absorption is due to the different interactive energy dissipation processes of Ohmic loss, polarization loss and magnetic loss [18]. In recent years, with the expanded applications of EM technology, there is an explosive increase in gigahertz (GHz) band range by telecommunication and it is becoming vital to design and fabricate EM absorbers for protecting environment, human and sensitive circuits [19]. It is generally accepted that carbon foam is composed of two phases: solid and void. If the size of voids is much smaller than an incident wavelength, scattering will not occur and the material will behave as an “effective medium”, whose effective permittivity can be determined by Maxwell–Garnet theory [20,21]. Based on this theory, it can be concluded that the presence of high surface area pores can decrease the effective permittivity and optimize characteristic impedance, resulting in the improved microwave absorption [22– 24]. In addition, it has been revealed that multilayered structure of materials is beneficial for improving the EMI SE [25–27]. Carbon and ceramic hybrids are also a promising choice. Silicon carbide (SiC) is an attractive high temperature structural material because of superior properties such as excellent anti oxidation, stable chemical properties and high strength [28–31]. It can meet the technical requirements of large electric breakdown field at a wide frequency range, but pure SiC cannot be used as absorption material due to its weak dielectric permittivity [32,33]. Theoretically, the electrical conductivity of the pure SiC is considerably low (
10  13 S/m) owing to the band gap energies ranging from ∼2.4 to 3.4 eV [34,35]. Due to the impedance match between the SiC and free air, the high reflection of incident waves on SiC surface lead to its low absorption [36,37]. SiC nanowires (NWs) appear higher electrical conductivity than SiC grains because of the onedimensional structure on a nanometer scale, which can greatly increase the carrier concentration and high-saturated carrier drift velocity and thus contribute to the higher electrical conductivity [38,39]. In addition, owing to its chemical inertness, higher surface area and ability to resist radiations, this semiconductor is an especially suitable electronic material to be used in high-temperature and high frequency EMI shielding applications [40]. The average reflectivity of SiC fabricated by chemical vapor infiltration is around 5.7 dB and that of SiC NWs fabricated by the pyrolysis of preceramic polymer is as low as 3 dB [41,42]. Chiu et al. found that due to the formation of dielectric network structure, β-SiC nanowires exhibited a better EM wave absorption than those of SiC microparticles [43]. Wu et al. ascribed the high dielectric permittivity of the synthesized SiC NWs to the combined action of stacking faults (SFs), twin faults, and the interfacial polarization between the SiC and binder [44]. We suppose that in situ formed SiC NWs in the carbon foam may be an alternative method to increase the dielectric loss and absorption property of carbon foam. Herein, different amounts of SiC NWs were grown and the resultant carbon foams were characterized by scanning electron microscopy (SEM), transmission electron microscopy (TEM), X-ray diffraction (XRD) and vector network analyzer (VNA). The developed carbon foam containing in situ grown SiC NWs is unique of its kinds and by accurately understanding the EMI SE, it will be conducive to discuss the underpinning mechanism of EM absorption and design a new class of microwave absorbers.

2. Experimental procedure 2.1. Raw materials Semi-rigid polyurethane (PU) foam (0.08 g/cm3 density, 0.40 mm average pore size, 8–12 wt% carbon contents as determined by TGA) was purchase commercially, washed with distilled water and dried in air. Commercial novolac, silicon powder (99.0%, crystalline, particle size less than 300 mesh), activated carbon (AC, purity Z99.5%, according to manufacturer specifications has a Brunauer–Emmett–Teller (BET) surface area of 385 m2/g) was purchased from Sinopharm Chemical Reagent Beijing Co., Ltd. (Beijing, China). AC was first ground into a fine powder passed through 300 mesh screen, and then ultrasonically washed with distilled water for 1 h, filtered and finally dried in a box furnace at 110 °C. Powdered coal-tar pitch (Taiyuan China) with a softening point 120–130 °C, a density of 1.25 g/cm3 and carbon content of 88–92% was used to increase the carbon yield of carbon foam during the carbonization process. 2.2. Carbon foam preparation The carbon foam was synthesized according to our previous technique based on a novel powder molding method [15]. Herein, we have described the preparation procedure briefly. Initially, PU foam was soaked in a dissolved novolac and dried for more than 24 h. It was then crushed and ground along with coal tar pitch into a very fine powder. This powder became an instant precursor (IP) of carbon foam, which contained the sacrificial soft template (PU) and carbon precursors (novolac and pitch). The ratio of PU: novolac: pitch was maintained at 50:33:17. In order to in situ grow SiC nanowires, five batches with different quantities of Si were prepared as given in Table 1. After making the required powdered recipes, the mixture was ball milled in a tungsten carbide jar for about 4 h using a QM-1SP2 planetary-type ball mill with a planetary rotation speed of 450 rev min  1. The objective was to further reduce the particle sizes and make a closer contact of all powdered constituents. The final mixture with controlled particle sizes and metered weight was molded in a perforated steel mold and a slight pressurization was exerted using hands (Fig. 1b). Afterwards, the molded mixtures were cured thermally in an aircirculating oven using a curing cycle of 120 °C/2 hþ140 °C/ 2 hþ 160 °C/2 hþ180 °C/2 h. After curing, the foam, generally called green foams shrank conformably (preserving the initial shape) and remained monolithic (Fig. 1c). Carbonization process was carried out in inert gas environment and under the cover of coke breeze. In the coking furnace, the samples were placed in a steel container and coke breeze was poured over them to completely cover them. The coke breeze (size r10 mm) is a by-product of coke production from a low ash bituminous coal after heat treatment at 550 °C. During heating, volatiles and unstable components are driven off while fused, stable and fixed-carbon is left behind which is termed as metallurgical coke breeze. At the first stage of the carbonization process, called pre-carbonization, a slow heating Table1 Composition of carbon foam containing various amounts of Si powder for in situ grown SiC NWs. Sample

IP (g)

AC (g)

Ferrocene (g)

Si (g)

IPSi0 IPSi5 IPSi10 IPSi15 IPSi20

100 100 100 100 100

0 3 6 9 12

0 2 4 6 8

0 5 10 15 20

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⎛ρ −ρ ⎞ a %P =⎜ t ⎟× 100 ⎝ ρt ⎠

Fig. 1. Digital camera images of carbon foam processing; (a) steel mold, (b) powder mixture molded and pressed, (c) green foam after curing and (d) carbon foam after heat treatment at 1500 °C.

rate of 10°/h was selected and after 450 °C, the rate was increased to 40°/h up to 1000 °C and held for 4 h. The slow rate minimized the risk of cracks produced due to evolution of volatiles from the sample, reaction of various additives with carbon precursors and the pyrolysis gases to form in situ nanowires and other compounds. The coke breeze, being reusable, served multiple functions other than producing a reduced environment. It controlled the flow of volatiles, homogenized temperature distribution, exerted a slight pressurization on the sample and preserved heat during the cooling cycle. To convert residual Si into SiC, the pyrolyzed samples were placed in a high temperature furnace (∅500 mm, 1000 mm in height) and heated to 1500 °C at 10 °C/ min in a flowing argon environment, maintained under these conditions for 4 h, and then cooled to room temperature naturally. SiC already present in the foam was expected to mitigate the exothermic reaction between silicon and carbon as shown in Eq. (1). Fig. 1 shows various stages of processing carbon foam containing in situ grown SiC NWs.

(2)

where %P is the bulk porosity, ρt and ρa are the true and apparent densities of the samples respectively. Thermal behavior was measured with a thermo gravimetric analyzer (Shimadzu, TG/DTA50). The samples were heated from 25 to 1500 °C at a ramping rate of 10 °C min  1 under oxygen atmosphere, and constantly weighed. Crystal phase assessment of the carbon foams was performed with X-ray diffraction (XRD, Philips analytical PW 1710) using Cu Kα radiations of 1.5418 Å at 40 kV and 30 mA. The XRD patterns were recorded at 2θ values in the range of 10–90° with a step size of 0.07° using monochromatic X-rays. The microstructure of the specimens was studied by scanning electron microscopy (SEM, JEOL JSM-6490LV) with an operating voltage of 20 kV and by transmission electron microscopy (TEM and HRTEM, Tecnai G2– F30, America). Bulk electrical conductivity of specimens was measured by four probe method using computer controlled current source (Keithley 6221 DC, Ohio, USA) and nano-voltmeter (Model 2182 A) with dimensions of 22.86l  10.16w  4.00t mm3. Samples were attached to a circuit board and four contacts were made on each end using silver paste for eliminating the contact resistance. The EMI shielding properties were measured using the rectangular waveguide method with a two-port vector network analyzer (VNA; MS4644A, Anritsu, Kana-gawa, Japan) in collaboration with the coaxial transmission line. The rectangular carbon composite foam specimens of dimensions 22.86l mm  10.16w mm  4.00t mm were fixed in the cavity of the wave-guide. The total shielding effectiveness (SET), absorption shielding effectiveness (SEA) and reflection shielding effectiveness (SER) were determined by calculations of the magnitude of complex scattering parameters (S-parameters) that correspond to reflection (S11 or S22) and transmission (S21 or S12) coefficients [45]. An electromagnetic wave was directly injected onto the specimen and frequency was scanned from 8.2–12.4 GHz (X band). Transmission and reflection technique introduced by Nicolson–Ross and Weir–NRW algorithm was used to calculate permittivity and permeability from measured S-parameters (ASTM D5568-08). Various researchers have reported reflection loss using back metal plate method but here, back metal plate was not used and, instead, transmission between the port 1 and port 2 has been performed and analyzed through S21 parameter. At least three samples were prepared and tested to minimize the effect of air gaps existing within sample – holder region and to eradicate artificial errors caused in measurement process.

3. Results and discussion 3.1. Physical, morphology and crystalline phase Table 2 shows the density, porosity and residual mass of carbon

Si (s) þC (s)-SiC (s)

ΔH ¼  73 kJ/mol

(1)

2.3. Characterization The densities were calculated from the weight and dimensions of rectangular carbon foam bodies. True density and apparent porosity were examined by helium gas displacement pycnometer (Pentapyc 5200e Quantachrome Instruments, Florida, USA). The true density is the mass per unit volume of the material excluding all the voids or pores. The porosity was calculated using the following expression:

Table 2 Properties of SiC NWs containing carbon foam; bulk density, porosity and residual masses at 800 and 1500 °C. ID

Bulk density (g/cm3) Open porosity (%)

IPSi0 IPSi5 IPSi10 IPSi15 IPSi20

0.56 0.56 0.57 0.58 0.59

72.14 70.5 71.81 70.7 69.53 70.8 66.18 71.0 64.33 71.0

Residual mass at 800 °C (%)

Residual mass at 1500 °C (%)

6.90 22.75 40.90 51.75 60.73

5.50 28.80 52.50 67.50 80.00

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Fig. 2. Weight change curves of the SiC NWs containing carbon foam in air with the increasing temperature from room temperature to 1500 °C at the rate of 10 °C/ min.

foam containing various amounts of SiC NWs. Fig. 2 shows the air reactivity of carbon foam specimens containing different amount of SiC NWs. The IPSi0 carbon foam showed a thermo-oxidative stability up to 494 °C with only 1% weight loss. Later the mass loss increased quickly to a ca. 7.5% residual mass of the initial value at 795 °C. Beyond this temperature, the mass loss was very slow with a residual mass of 5.5% at 1415 °C and remained constant until 1500 °C. The reaction of carbon strongly depends on temperature and porosity. The mechanism of oxidation in carbon foam involved several steps including physisorption, chemisorption and desorption on the active sites, like edges, defects, cracks and pores [46]. These entire phenomena contribute to the weight loss of material and are mostly influenced by the diffusion coefficient of oxygen in the pores and cracks. The results indicate that above 500 °C, pores became wider presenting more reactive sites for oxygen attack, facilitating the diffusion of air in the foam and increasing the oxidation reaction. At a lower temperature (o500 °C), the oxidation loss was due to the surface reaction (machined surface being more active) between carbon and oxygen and was controlled by a chemical reaction: C(s) þO2(g)-CO2(g) þ activation energy (195–205 kJ/mol)

(3)

At an intermediate temperature (500–800 °C), it was controlled both by a chemical reaction and gaseous diffusion [47]. The oxidation of carbon resulted in the formation of gaseous carbon oxides and thus, this gasification reaction caused weight loss and geometrical change. At this stage, the chemical reaction rate was the highest and the pores of the foam were filled with CO and CO2 gases [48]. For all other specimens containing SiC NWs, as the amount of SiC NWs increased, the residual weight also increased in the temperature range of 725–825 °C. The IPSi20 carbon foam showed a delayed oxidation starting at 512 °C and continued till 694 °C at a faster rate. The residual mass also increased from 5.5% to 80% as the amount of Si increased from zero to 20 wt% in the carbon foam recipe. It can be seen that after 825 °C, oxidation of all the specimens containing SiC NWs was a continuous weight gain. Due to passive oxidation, the rate became sharp in the temperature range of 1250–1300 °C as shown by the following reaction [49]: 2SiC(s) þ 3O2(g)-2SiO2(s) þ 2CO(g)

(4)

The SiO2 formed during passive oxidation get deposited over the surface of pores leading to a net increase in the residual mass. After formation of SiO2, further oxidation proceeded through the diffusion of oxygen molecules through amorphous oxide layer and reaction with underlying SiC and carbon at the oxide/SiC interface, thus slowing down the oxidation reaction [50,51]. The IPSi0 carbon foam collapsed completely forming porous skeleton being debris, which suggests its poor oxidation resistance; while the SiC NWs containing specimens maintained their integrity in the same oxidation environment, IPSi20 being the strongest among all. Fig. 3 shows SEM images of the pristine and as-grown SiC NWs in the pores of carbon foam. The surface morphology of the IPSi0 were non-uniform spherical and distorted spherical with open and interconnected pores (Fig. 3a). Due to the complexity and nonuniformity of the cell walls, it is difficult to obtain accurate dimensions and distribution of pore sizes. It falls in broad ranges up to a maximum of 200 μm diameter and a wall thickness up to 50 μm. This was due to the carbonization shrinkages in the carbon foam containing PU particles as pore former which transformed the foam into a coarse, rough and uneven morphology. Some pores in the image are seen broken which were actually damaged during cutting, grinding and sample preparation. The pore morphology changed by the addition of Si powder in the carbon foam recipe. After heat treatment at 1500 °C, the pores were covered with sponge like webs containing thousands of SiC NWs. The amount of SiC NWs increased in the samples from IPSi5 to IPSi20 as the amount of Si increased in the recipe. The clump that looked tangled up is actually many NWs that are difficult to distinguish individually and some of them even fused together into thicker structures especially in IPSi20 (Fig. 3e). Among the jungle of kinked and twisted curly wires, several very large fairly straight and smooth tubes are also seen. The NWs with diameter between 20 to 40 nm have been observed but there are also some large wires with diameter 400 to 600 nm. Most of them are smoothly curved, while some of them possess bends and kinks. The growth mechanism being vapor–liquid–solid (VLS) has already been described in details in [4]. The phase composition of carbon foams containing various amount of SiC NWs is shown in Fig. 4. The small peaks located at about 33.6° (2θ), marked as SF, reflect the spontaneously formed stacking faults (SFs) during the NWs growth and all the other peaks are indexed to those of 3C–SiC (JCPDF 29-1127) [52]. In addition, no amorphous scattering curves and diffraction peaks of trace Si were detected. Hence, the conversion of Si into the β-SiC was rather completed. In our case, the SiC NWs were grown in situ in the carbon foam via VLS mechanism. In such a case, formation of SFs can remarkably promote the growth of NWs along 〈1 1 1〉 direction due to less energy consumption [53]. Thus, high density of SFs was formed spontaneously during the growth of SiC NWs in the carbon foam. To investigate more details of the structure of β-SiC nanowires, TEM combined with energy dispersive X-ray (EDX) analyses was performed as shown in Fig. 5. The diameter of SiC nanowires was in the range of 30–100 nm, consistent with the result of SEM images (Fig. 3). The inset image in Fig. 5d shows the high-resolution TEM (HRTEM) image where numerous SFs (stripes) were visible in a nanowire perpendicular to the growth direction. The electron diffraction pattern clearly exhibited featureless streaks typical of a disordered layer structure. As these streaks are always perpendicular to the stacking fault planes 〈1 1 1〉, the main growth direction can be concluded as parallel to the 〈1 1 1〉 direction. The EDX analysis reveals that the main chemical elements in the samples are Si and C (Fig. 5c). The quantitative analysis indicates that the mass percentage of elements is 22.95 wt% Si and 70.49 wt% C.

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Fig. 3. Surface morphology of carbon foam; (a) IPSi0, (b) IPSi5, (c) IPSi10, (d) IPSi15 (e) IPSi20 and (f) low magnification view of IPSi20.

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Fig. 4. X-ray diffraction patterns of the carbon foam; (a) IPSi0 and (b) IPSi5, IPSi10, IPSi15 and IPSi20.

Fig. 5. TEM image of β-SiC nanowires grown in situ in the carbon foam; (a) curvy and branched NWs, (b) different sizes, (c) stacking faults along with EDX inset and (d) HRTEM with SAED patterns.

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Fig. 6. Electrical conductivity of IPSi0 to IPSi20 carbon foam samples as a function of Si contents.

3.2. Shielding effectiveness Although a shield tends to be electrically conducting but a high conductivity is not a scientific criterion for shielding. Conduction requires connectivity in the conduction path where shielding does not but it is enhanced by connectivity, so it can be seen that better shielding effectiveness can be achieved from moderate conducting materials [54]. In case of homogenous or single–phase porous bodies, electrical conductivity increases linearly with the increase of relative density due to decreasing tortuosity of the current path in the foam [55]. This is reasonable when the material is homogenous or composed of single phase. Herein the IPSi0 showed a conductivity of 2.23  102 S m  1, which was ascribed to the locally interconnected graphitic structures formed during high temperature treatment. Other samples containing SiC NWs showed somewhat lower electrical conductivities as shown in Fig. 6. It is not contradictory when taking the porosity and density of samples into consideration. In this work, since the as-prepared carbon foams containing SiC NWs are heterogeneous in nature, the relative content of each phase may also have significant influence on the electrical conductivity [56]. A semi-conductor phase (SiC and NWs) mixed with the conductive phase and a decrease of relative content of conductive phase inside the foam decreased the electrical conductivity as evident in specimens IPSi5 to IPSi20. The relative permittivity (ɛr ¼ɛ/ε0 ¼ɛ′  jɛ″) is the key parameter for the characterization of dielectric properties of materials. For most dielectric materials, conductivity loss is the most important factor that affects the permittivity. The real (ɛ′) and the imaginary (ε″) parts are related to the storage and loss of electric energy respectively [57,58]. As shown in Fig. 7, the real part decreased while the imaginary part increased non-linearly with increasing frequency in the X band. The decrease was due to the reduction of the electric field induced in the porous body in response to the reversing external electric field by the delay in the molecular polarization at a higher frequency. This phenomenon is known as dielectric relaxation, and it reflects the irreversible degradation of free energy [59]. Presence of SiC NWs affected the real part of permittivity, which indicates energy storage when EM wave impinges upon the material. The values of ɛ′ and ɛ″ increased nonlinearly with the increase in SiC NWs in the carbon foam. Additionally, ɛ′ decreased while ɛ″ increased slightly with the frequency in X band (Fig. 7a and b). This is conducive to the increase of dielectric loss. For IPSi0, the permittivity was small (14  i11.12)

7

at 8.2 GHz, which changed slightly to (11.12  i16.52) at 12.4 GHz. Concerning the IPSi20, the permittivity increased remarkably (90.7  i70) at 8.2 GHz, and somewhat a frequency dispersion effect was found. In situ grown SiC NWs offered the channels of mobile charge carriers which interacted with the EM field over the short range. The average ɛ′ value of the specimens IPSi0 to IPSi20 increased from 12.5 to 75 with the average ɛ″ value increasing from 28 to 95. Therefore, the ɛ′ and ɛ″ could be adjusted by changing the Si contents in the carbon foam recipe. From IPSi0 to IPSi20 carbon foams, both ɛ′ and ɛ″ were simultaneously improved, especially for IPSi20, whose ɛ′ and ɛ″ even exceeded 95 dB. This should be attributed to the formation of more SiC NWs and the consequent increase of dipolar polarization resulting from the difference in conductivity of carbon foam, air and SiC. Stacking faults also contributed to the increase in permittivity [60]. The accumulated charges in the SF/3C–SiC interfaces induced dipole moments, which mean that the increased amount of SFs will result in the enhanced dipole moments. Meanwhile, the electrons confined in the SF/3C–SiC interfaces can move freely along the SF/3C–SiC plane [61–63]. As the SiC NWs grew, there existed a plenty of interfaces between the carbon matrix and NWs. The incident EM wave would reflect back and forth between these interfaces and dissipate gradually [64]. It is possible to consider that the generated SiC NWs offered the mobile charge carriers channels to move and interact with the EM field over a short range [38]. From the SEM images (Fig. 3), it can be seen that the length of the SiC NWs is very long. So when the EM field was applied to the specimen, the mobile charge carriers moved along with NWs, resulting in higher electrical loss and giving higher dielectric constant. In addition, it is discovered that the real (μ′) and imaginary parts (μ″) of complex permeability of these foams are also quite different. The μ′ and μ″ of IPSi0 are approximately constant and μ″ is close to zero, indicating a negligible magnetic loss for incident EM wave. The μ′ of all the specimens is below one as carbon is diamagnetic. Very interestingly, the μ″ gradually decreased and became negative as the frequency increased in the X band. For IPSi10, IPSi15 and IPSi20, it remained negative, which in theory, could be possible for porous diamagnetic materials [65]. Additionally, some published papers have also proposed that negative μ″ can produce geometrical effects to improve microwave absorption [66]. Recently, some results related to negative permeability are reported in composites containing fillers as multiwalled carbon nanotube, Fe3O4/ZnO core/shell nanorods or mesoporous C/SiO2 [67–69]. However, the absolute values of these negative μ″ are very small. This phenomenon suggests that the interaction between NWs and incident EM wave can induce variation of magnetic field and little magnetic energies can be radiated out. Moreover, it should be considered that the permittivity and permeability are obtained from S-parameters through going some modeling and sometimes these modeling may not provide a correct response. The variations in SET, SEA and SER of the fabricated composites in the X band are shown in Fig. 6c–e respectively. EMI SE is a combined result of SER, SEA and SEM. According to the Schelkunoff formula based on the Transmission line model, EMI SE for a shielding material, is described by Eq. (5):

SET = SER + SEA + SEM

(5)

The correction term SEM can be ignored in all practical applications when SET 415 dB. In another view, SEM can be considered as SE induced by multiple reflections inside the shield, which in the case of foams, and especially in this case of tertiary phase, can be significant. However, from the simple S parameters, it is not possible to decouple SEM and SEA. Moreover,  10 log(T/(1  R)) and

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Fig. 7. Relative complex permittivity, permeability and SE of carbon foam; (a) real permittivity, (b) imaginary permittivity, (c) SET, (d) SEA and (e) SER in X band.

10 log(1  R) are different from SEA and SER due to any nonnegligible SEM. But, it is convenient and practical to use these expressions for assessing the degree of contribution of SER and SEA to the total EMI SE as follows [70]:

SER ≈ − 10 × log10 (1 − R)

(6)

SEA ≈ − 10 × log10 (T /(1 − R))

(7)

SET ≈ SEA + SER = 10 × log10 ( Pin/Pout )

(8)

where R and T are calculated according to following equation:

T = S12 2 = S21 2 = Pout /Pin

(9)

A = 1−R−T

(10)

R = S11 2 = S22 2 = Pref /Pin

(11)

Pref, Pin and Pout represent the reflected power, absorbed power, Please cite this article as: S. Farhan, et al., Electromagnetic interference shielding effectiveness of carbon foam containing in situ grown silicon carbide nanowires, Ceramics International (2016), http://dx.doi.org/10.1016/j.ceramint.2016.04.054i

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Fig. 8. The dielectric loss tangent (a) and magnetic loss tangent (b) of carbon foam containing various amounts of SiC NWs.

and transmitted power; R and T represent the reflection coefficient and transmission coefficient. The above equations, being derived from dense single-phase conductive materials, can be used as criterions to anticipate the general effects of different parameters on the EMI SE [71]. The SET value of IPSi10, IPSi15 and IPSi20 were always above 30 dB over the entire measured band, while the values for IPSi0 and IPSi05 were below 25 dB. The value was not less than 20 dB in all the specimens containing SiC NWs, which is the target value, required for commercial application. The carbon foams with more quantities of SiC NWs showed a high SEA values, implying that the absorption is dominant EMI shielding mechanism [72]. In fact, SER is also increasing with the rise in SiC NWs in the foam. In the present work, electrical conductivity of samples was low due to in situ reactions forming SiC. The SEA values are lower in case of low electrical conductivity. Nevertheless, due to higher contents of SiC NWs, the polarization loss increased generating a positive contribution to the SEA. In the porous carbon foam, partial absorption, reflection, and multi-reflection occurs repeatedly and in this regard, we can conclude that the EMI SER is contributed by reflection from the surface and multi-reflection within the pores. Multiple reflections within the pores inside the shield can also contribute to the EM wave attenuation inside the shield by prolonging the EM wave path and eventually reflected as a part of SEA. For EMI shielding materials, the dielectric loss is a significant determinant of the SEA. The dielectric loss (imaginary permittivity or Ohmic loss) comes from the dissipation of electrical energy by nomadic charges in phase with the applied electric field. Generally, the conductivity loss is mainly attributed to the intrinsic conductivity while the polarization loss can be regarded as an integrated factor resulting from the defects and interfaces within the samples [73]. Polarization loss comes from the reorientation of dipoles in each half cycle of alternating field and is out of phase with the applied field. Both of which contribute to EMI shielding through Ohmic loss and polarization loss. In this work, the intrapores and inter-pores as shown in Fig. 3, are less than 200 μm, while the wavelength of the EM waves in X band remains in the order of centimeter (2.5–3.75 cm). The losses of incident EM waves into the pores are simply caused by the dissipation of electric current in the pore walls. Because of aerospace applications related to EMI SE, the specific SE (SSET, EMI shielding effectiveness divided by the density) is more appropriate for use in comparing the shielding performance between typical metals and foam composites. In this work, SSET of the IPSi20 was calculated to be around 79 dB cm3/g at 8.2 GHz and 66 dB cm3/g at 12.4 GHz, which is much higher than that of typical metals (compared to

10 dB cm3/g for solid copper). The pore structure decreased the effective permittivity, and made a positive contribution to impedance match. This resulted in a lower reflection of the EM field achieved by the porous structures than by solid structures. Moreover, the pore structure increased the multiple reflections causing more energy to be dissipated inside because they make the microwave travel a longer distance in the porous carbon materials. The porous structures and SiC NWs have a combined effect on the increment of the absorption loss and SET. The SER is also an outcome of combined action exerted by multiple factors. With the increase of SiC, the resistance of carbon foam increased correspondingly, which is favorable to the improvement of impedance matching between the specimen and the air. The improved impedance matching, meanwhile, weakens the reflection ability and will make more EM waves enter the foam. On the other hand, as shown in Fig. 3, the pores were gradually filled by SiC NWs leading to the reduction of porosity. Due to very high surface area of NWs the EM wave impedance mismatch may enlarge resulting in the increase of SER. As a result, the SER of the sample as a function of SiC NWs also increased. To further understand the incident EM wave attenuation in an EM absorbent, the dielectric loss tangent (tan δE ¼ ε″/ε′, electric dissipation of EM energy) and magnetic loss tangent (tan δM ¼ μ ″/μ′, magnetic dissipation of EM energy) were calculated based on the measured complex permittivity and permeability. Fig. 8 illustrates tan δE and tan δM of the carbon foam containing various amounts of SiC NWs. The tan δE was in the range of 0.67–3.17 between 8.2 and 12.4 GHz, which is very high. The tan δE also increased with the increase of SiC NWs contents and frequency in X band. Generally, higher tan δE means better microwave attenuation capability [74]. Conductive micro-network structure of NWs made the micro current exhausted in the foam under alternating EM field, which generated a strong conductive loss and led to the increase of EM attenuation capability. We can say carbon foam containing SiC NWs have the highest dielectric loss tangent, which indicated that it could be a potential absorbing material. The excellent tan δE properties could be concluded from the Debye dipolar polarization relaxation and interfacial polarization relaxation etc. [75,76]. Based on the surface effect of SiC NWs, with the decrease in size, the number of surface atoms with unsaturated bonds increased greatly, causing an increase of dipoles. Consequently, the dipole polarizations and electrons could not match up with the change of EM field in the high frequency, which led to the Debye relaxation contributing to very high dielectric loss. Moreover, excessive stacking faults acted as polarized centers for the dipole

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polarization relaxation, enhancing the dielectric loss and electromagnetic energy dissipation [77]. The results imply that SiC NWs induced a higher dielectric permittivity and loss tangent than that of the carbon foam without SiC NWs. A very low magnetic loss tangent (Fig. 8b) indicated that the EM loss of these materials was dielectric type and the reasons for the negative value of magnetic loss tangent, is also referred.

4. Conclusions

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[15]

[16]

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In summary, SiC NWs were in situ formed within a carbon foam by using a powder molding method for the processing of carbon foam. The formation of SiC NWs plays a great role in determining EM parameters and SE along with the relative complex permittivity that improved greatly. IPSi20 displayed a specific EMI SE up to 79 dB cm3/g – almost more than two times higher than that of the pristine carbon foam, indicating its advantages as a kind of structural and functional materials for radar absorption. A very small size of NWs, stacking faults, heterostructure interfaces and long tortuous paths in the carbon foam caused a strong dielectric loss and induced higher dielectric permittivity. The developed NWs network enabled the foam lossy and diamagnetic, and led to an enhanced dielectric loss of 3.17 at 12.4 GHz. This suggests that the carbon foam filled with different quantities of NWs could be used as shielding materials; however, more studies are needed to study the role of pore size, pore filling with NWs and their optimum quantities.

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Acknowledgments The authors acknowledge the financial support (CX201627) of Innovation Foundation for Doctor Dissertation of Northwestern Polytechnical University and National Natural Science Foundation of China Grant no. 51472202.

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