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Electrical percolation and infrared emissivity of pressureless sintered SiC-MoSi2 composites tailored by sintering temperature
Journal of the European Ceramic Society 39 (2019) 3981–3987
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Original Article
Electrical percolation and infrared emissivity of pressureless sintered SiCMoSi2 composites tailored by sintering temperature
T
⁎
Jia-Qi Zhenga,b, Jian Chena, , Bu-Hao Zhanga,b, Xue-Jian Liua, Zhong-Ming Chena, Hai-Bo Wua, ⁎ Zheng-Ren Huanga, a State Key Laboratory of High Performance Ceramics and Superfine Microstructure, Shanghai Institute of Ceramics, Chinese Academy of Sciences, Shanghai, 200050, China b University of Chinese Academy of Sciences, Beijing, 100049, China
A R T I C LE I N FO
A B S T R A C T
Keywords: Silicon carbide Molybdenum disilicide Electrical percolation Grain boundary Sintering temperature
SiC-MoSi2 composites with low electrical resistivity and high infrared emissivity were fabricated via pressureless sintering. The relationship between microstructure evolution and electrical behaviors along with infrared emission properties of the resulting composites is investigated at various sintering temperatures. The electrical resistivity undergoes two significant drops with increasing sintering temperature. Pore elimination bears responsible for the initial decrease in electrical resistivity. Transmission electron microscopy (TEM) observation manifests that the thinned amorphous layers at SiC/MoSi2 interface decrease grain boundary resistivity and allow for electrical percolation to occur when sintering temperature further rises. Additionally, increasing sintering temperature leads to a higher infrared emissivity owing to the formation of Mo4.8Si3C0.6 and the decreased boundaries. The lowest electrical resistivity of 7.2 Ω cm and the highest infrared emissivity of 0.721 are recorded for composite sintered at 2000 ℃. Overall, SiC-MoSi2 composites exhibit a promising prospect as infrared source elements that must endure harsh environments.
1. Introduction Among a wide range of structural and electrical ceramic materials, silicon carbide (SiC) stands out by virtue of its superb physical and chemical properties, such as superior mechanical strength, high thermal conductivity, excellent corrosion resistance and a large bandgap [1–5]. In addition, SiC exhibits high infrared emissivity (˜0.9 above 1200 K [6]), which shows great promise in high temperature electrothermal materials such as infrared source. In this type of applications, a linear current-voltage (I-V) characteristic is basically required. Unfortunately, SiC ceramics usually suffer from relatively low electrical conductivity and nonlinear I-V behavior, both of which are derived from a double Schottky barrier at grain boundaries [7]. With continued pursuit for developing SiC-based infrared source elements, it has become imperative to tailor the electrical properties of SiC ceramics at no expense of infrared emissivity. So far, incorporating an electrically conductive filler over a SiC matrix has been considered a more convenient and versatile method compared to elemental doping. Though ohmic contacts with low contact resistivity can be obtained with the addition of metals (Ni, Ti, etc. [8]), the tendency of SiC to react with metals to form carbide or ⁎
silicides may threat long term applications. There has been several recent success in producing highly conductive SiC ceramics by introducing TiN [9], ZrB2 [10], graphene nanoplatelets [11] and carbon nanotubes [12]. These conductive fillers could generally decrease the Schottky barrier height (SBH) once they are uniformly dispersed over the matrix, and even deliver linear electrical properties. Given the harsh working environment of infrared source, MoSi2 becomes a leading candidate for electrically conductive phase due to its high melting point (˜2030 ℃), very low electrical resistivity (2.15 × 10−5 Ω·cm), and distinctive oxidation resistance [13,14]. Much attention has been paid to the mechanical properties and oxidation behaviors of SiC-MoSi2 composites [15–17]. Recently, Mao et al. [18] successfully fabricated SiC-MoSi2 composite coatings with high infrared emissivity and high thermal shock resistance. Besides, some reports mentioned that MoSi2 phase effectively decreased the electrical resistivity of SiC ceramics [19,20]. However, the electrical conduction mechanism of SiC-MoSi2 composites is still left unanswered. Our previous work [21] reported a percolation phenomenon in SiCMoSi2 composites. There is a percolation threshold in the amount of MoSi2 particles over which the electrical behavior changes dramatically. And this phenomenon is strongly associated with the Schottky
Corresponding authors. E-mail addresses: [email protected] (J. Chen), [email protected] (Z.-R. Huang).
https://doi.org/10.1016/j.jeurceramsoc.2019.05.019 Received 4 April 2019; Accepted 14 May 2019 Available online 16 May 2019 0955-2219/ © 2019 Elsevier Ltd. All rights reserved.
Journal of the European Ceramic Society 39 (2019) 3981–3987
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Table 1 Relative density and electrical properties of SiC-MoSi2 composites as a function of sintering temperature. Sintering temperature (℃)
1700 1800 1900 2000
Relative density (%)
86.6 94.5 94.7 96.0
Electrical properties Electrical resistivity (Ω cm)
Nonlinearity exponent
Rg (Ω cm)
Rgb (Ω cm)
2519.2 266.0 35.2 7.2
1.271 1.208 1.104 1.034
173.1 ± 62.3 148.6 ± 8.8 28.1 ± 1.4 21.4 ± 1.1
(3.173 ± 0.011) ×104 (1.484 ± 0.004) ×104 – –
For electrical resistivity measurement, the silver paste was applied to both surfaces of the disk-shaped specimens of 15 mm in diameter and 2 mm in thickness followed by heated at 600 ℃ for 10 min. The impedance spectra in the frequency range from 20 Hz to 120 MHz were examined using an AC (alternating current) impedance spectrometer (Agilent HP4294A, Hewlett-Packard, Canada). The applied voltage was fixed at 0.5 V. Every component parameter in the equivalent circuit for the measured impedance plots was obtained via ZSimpWin software package (EChem Software, Michigan, USA). The DC (direct current) potential-current characteristics were measured in an Ar atmosphere from room temperature up to 700 ℃ by a high resistance meter (Keithley 6517A, Keithley, USA). The electrical resistivity of the specimens was determined from the I-V data at 1 mA·cm−2. The infrared emissivity properties of SiC-MoSi2 composites were obtained using a high temperature spectrum measurement system. The system mainly consisted of a Fourier-transform infrared (FTIR) spectrometer (NICOLET Is10, Thermo Fisher, USA) with a liquid nitrogencooled detector and a resistance furnace. The infrared spectrum of the specimens was detected within the range of 2.5–16 μm. Emission from the polished surface of each specimen in a vacuum at 25 ℃ and in the range of 100–700 ℃ was measured to minimize the effects of surface roughness and oxidation behaviors.
barrier at grain boundary of SiC ceramics. Based on the Schottky-Mott theory [22], the SBH of a SiC-MoSi2 system should be primarily dependent on the work function of MoSi2 and the electron affinity of SiC, but there is more to study. Recently, Kim et al. [23] indicated that the presence of thin intergranular films between SiC grains resulted in highly resistive SiC ceramics. Gu et al. [24] directly observed boron segregation to grain boundaries in hot-pressed SiC ceramics. These findings let us ponder how the quality of SiC/MoSi2 interface influences the electrical transport in SiC-MoSi2 composites, in terms of chemical interaction or microstructure evolution of grain boundaries with various sintering temperatures. In this study, pressureless sintering was performed for the synthesis of SiC-MoSi2 composites. Y2O3 along with AlN was used as sintering aids to secure densification. The objective of this work is to investigate the microstructure evolution, electrical behaviors and infrared emission properties of SiC-MoSi2 composites. The effect of sintering temperature on microstructure, electrical percolation, and infrared emissivity is systematically discussed. 2. Experimental procedure Commercially available submicron β-SiC (˜0.5 μm, ≥99% pure, EnoMaterial Co., Ltd., Hebei, China) and MoSi2 (˜1.0 μm, ≥99.9% pure, Haoxi Research Nanomaterials Inc., Shanghai, China) were used as the starting powders; 5 wt% AlN (˜2.0 μm, ≥99% pure, Tokuyama Co., Tokyo, Japan) and 2 wt% Y2O3 (≥99% pure, Aladdin Industrial Co., Ltd., Shanghai, China) powders were used as the sintering aids; phenolic resin (Shanghai Tiao Chemical Co., Ltd., Shanghai, China) was used as the binder for green compacts. Based on our previous research [21], the MoSi2 content in the composites was fixed at 15 wt% which was approximately close to the electrical percolation threshold sintered at 1900 ℃. The mixture was blended with 5 wt% phenolic resin in ethanol and ball-milled for 24 h using SiC media in a polypropylene jar. Then the slurry was dried at 80 ℃ for 12 h and sieved (100 mesh). The as-received powder mixture was pressed uniaxially into disks under 10 MPa followed by cold isostatically pressed under 200 MPa. The green bodies were kept at 1000 ℃ for 1 h in a vacuum to ensure pyrolysis of the phenolic resin. Finally, the pressureless sintering of the samples was carried out in a graphite resistance furnace at 1700 ℃, 1800 ℃, 1900 ℃, and 2000 ℃ for 1 h in an argon atmosphere, respectively. The heating and cooling rates were fixed at 10 ℃/min. The relative density of SiC-MoSi2 composites was measured by the Archimedes method. The corresponding theoretical density was determined according to the rule of mixtures. Phase identification for sintered samples was performed by an X-ray diffractometer (XRD, D8 Advance, Bruker, Germany) with Cu Kα radiation (λ =1.5406 Å). The diffractograms were scanned in 2θ from 10° to 80° at a rate of 0.04°/s. The microstructures of polished samples were characterized by scanning electron microscopy (SEM, Magellan 400, FEI, Netherland). Further detailed information about the microstructures was obtained via TEM (JEM-2100F, JEOL, Japan) equipped with energy-dispersive spectroscopy (EDS), which was carried out at 200 kV. Specimens for TEM were prepared conventionally by cutting 3 mm discs from the assintered bulk. The specimens were mechanically ground and further thinned by ion beam milling.
3. Results and discussion The relative densities of the as-prepared SiC-MoSi2 composites are listed in Table 1. Most of the samples can be densified to nearly 95% of the theoretical density except that sintered at 1700 ℃. During liquidphase sintering, Y2O3 and AlN react with the native SiO2 layer which is present on the surface of SiC and MoSi2 particles, to form Y-Al-Si-O-C-N melt with the dissolution of SiC [5]. The pyrolysis of phenolic resin yields ˜60% excess carbon, which is generally believed to facilitate densification process by reacting with the SiO2 layer and enhancing bulk self-diffusion of SiC as well [25]. Fig. 1 presents the XRD patterns of SiC-MoSi2 composites sintered at various temperatures. Apparently, all the samples exhibit characteristic peaks of β-SiC (3C), α-SiC (6H, 4H, and 2 H), AlN, and MoSi2 as major phases. The β → α phase transformation of SiC with increasing sintering temperature can be discerned. The introduction of Al or Al compounds could also promote this process by stimulating the formation of point defects in the SiC lattice to enhance mass transport [26]. There is no trace of Y2O3 or Y3Al5O12 presumably due to the detection limit. In addition, a small amount of Mo4.8Si3C0.6 compound (JCPDS 731380) is detected in all samples. The peaks of Mo4.8Si3C0.6 heighten slightly with those of MoSi2 weakening at the elevated sintering temperature. A synchronous shift of the peaks of SiC and MoSi2 in Fig. 1b shows evidence of the formation of Mo4.8Si3C0.6 phase. This compound possibly originates from the reactions of partial SiC and MoSi2 at high temperature [16]: 8MoSi2 → Mo5Si3 + Mo3Si + 12 Si
(1)
3/5 Mo5Si3 + 3/5 Mo3Si + 3/5 SiC → Mo4.8Si3C0.6
(2)
The typical SEM micrographs of the SiC-MoSi2 composites with elevated sintering temperature are illustrated in Fig. 2. The most 3982
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Fig. 1. XRD patterns of (a) SiC-MoSi2 composites with various sintering temperatures, and (b) magnified detail of 2θ range of [38°, 39°] in (a).
(Fig. 3f). Therefore, the increase in sintering temperature allows for the removal of the amorphous film present at SiC/MoSi2 interface. It is worth noting that the SAED pattern (inset of Fig. 3f) and the EDS result (Fig. 3g) taken from the dark phase in Fig. 3e provide direct evidence of the formation of Mo4.8Si3C0.6 phase, which is consistent with the XRD data shown in Fig. 1. The current density versus electrical field curves at room temperature for SiC-MoSi2 composites are shown in Fig. 4. Nonlinear I-V characteristics are observed in the specimens sintered below 2000 ℃. The following relation [29] can well fit these curves:
obvious differences in the microstructures are the grain size and shape of SiC (grey and dark phase). The increased sintering temperature not only leads to the grain growth of SiC but accelerates the occurrence of elongated grains, which indicates the presence of β → α phase transformation. A number of pores can be found in the composite sintered at 1700 ℃. Instead, the other samples display a more robust microstructure when sintering temperature becomes higher. Irregularly shaped MoSi2 particles are homogeneously dispersed over the SiC matrix. According to the EDS results in our recent report [21], insulating Y2O3 is mainly confined to the multigrain junctions. TEM was employed to further reveal the microstructure changes when sintering temperature rises from 1800 ℃ to 2000 ℃, as shown in Fig. 3. The results suggest that trace amount of sintering necks remain in the bulk sintered at 1800 ℃ (Fig. 3a). On the contrary, SiC and MoSi2 grains seem to be more strongly bonded at a higher sintering temperature (Fig. 3c and e). The corresponding HRTEM image (Fig. 3b) demonstrates a typical amorphous characteristic of SiC/MoSi2 interface sintered at 1800 ℃. The amorphous layer, about 2–3 nm in thickness, is estimated to be the Y-Al-Si-O-C-N melt It occurs during liquid-phase sintering and remains at grain boundaries after cooling [27,28]. It can be identified that the amorphous layer becomes thinner when sintering temperature rises (Fig. 3d), and finally disappears up to 2000 ℃
I ∝ Uα
(3)
where α is the nonlinearity exponent. The fitted α value accompanied by the electrical resistivity is presented in Table 1. It is suggested that increasing sintering temperature results in both the weakened electrical nonlinearity and the decreased electrical resistivity. The composite sintered at 2000 ℃ gives an ohmic nature, which implies electrical percolation is already obtained. And it also possesses the lowest electrical resistivity (7.2 Ω·cm) compared to the other samples. In Fig. 5a–b, the frequency dependence of the dielectric constant ε’ and the impedance complex plane plots for the measured samples are demonstrated. In Fig. 5a, the composites sintered below 1900 ℃ exhibit
Fig. 2. Typical microstructures of SiC-MoSi2 composites sintered at (a) 1700 ℃, (b) 1800 ℃, (c) 1900 ℃, and (d) 2000 ℃. 3983
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Fig. 3. TEM images of SiC-MoSi2 composites sintered at (a) 1800 ℃, (c) 1900 ℃, and (e) 2000 ℃; (b), (d), and (f) are HRTEM images enlarged from the yellow square in (a), (c), and (e); insets in (b), (d), and (f) are the corresponding SAED images; (g) EDS results taken from the dark phase in (e) (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article).
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Fig. 4. I-V characteristics of SiC-MoSi2 composites sintered at different temperatures.
Fig. 6. Electrical resistivity changes with testing temperature for SiC-MoSi2 composites sintered at different temperatures.
frequency dependent throughout the frequency range, which is a typical feature of a capacitive behavior [30]. Whereas the samples sintered at higher temperatures show an almost flat plateau in ε’, which is characteristic of conductive materials. The flat-to-dispersive transition in ε’ occurs at the critical frequency (fm). The Debye model [31] can be used to describe the dielectric relaxation:
temperature reaches 1900 ℃, which implies Rgb no longer dominates the overall resistivity and the onset of electrical percolation has been triggered. Further rising sintering temperature results in a gradual decrease in the electrical resistivity. The conductive pathways are deemed to curl or even spiral when concerning occurrence of inductance in the composites sintered above 1800 ℃. Accordingly, it is safe to believe there is an intimate relationship between the microstructure of the composite material and its electrical properties. The grain boundary effect is responsible for the enhanced electrical percolation behavior, whereas other aspects should be taken into consideration as well. The decreased electrical resistivity when sintering temperature rises from 1700 ℃ to 1800 ℃ might greatly benefit from pore elimination during densification. It is widely acknowledged that the removal of pores can enhance grain boundary structures and prevent gas phase from blocking current path [32]. However, given the similar relative density for the composites sintered above 1800 ℃ (Table 1), this factor could be nearly excluded. Specifically, the electrical nonlinearity weakens sharply and electrical percolation emerges when sintering temperature increases to 1900 ℃. From the microstructural perspective, the amorphous film with a high disorder and intrinsically insulating nature contributes to the formation of the depletion region at grain boundaries. It was reported that reduced depletion layer width would decrease potential barrier and consequently lead to notably electrical conduction [33]. Thus, this electrical behavior evolution is mainly attributed to the thinner amorphous layer at SiC/MoSi2 interface. At the final increment in sintering temperature up to 2000 ℃, the decrease in electrical resistivity still proceeds yet in a more gradual manner. Devitrification of the intergranular glass at SiC/ MoSi2 interface (Fig. 3f) is almost complete without any contribution of
ωm τ = 2πfm τ = 1
(4)
where ωm is the angular frequency of dielectric relaxation and τ is the relaxation time. Obviously, fm moves to a higher frequency with increasing sintering temperature, leading to decreased τ. This suggests the weakened interfacial polarization in the composite structures, i.e., capacitive coupling at grain boundaries. Fig. 5b further illustrates how grain boundary effect shapes the electrical behaviors. The plots for samples sintered below 1900 ℃ show a typical compressed semicircular arc at low frequency which is also attributed to the relaxation process at grain boundaries. To elaborate the contributions of grain and grain boundary to the electrical resistivity, accurate fitting to the equivalent circuit of the plots is performed, as shown in Fig. 5b. The equivalent series circuit is composed of one bulk component and one parallel grain boundary component where grain boundary resistance and corresponding constant phase element (CPE) are included. Hence, the bulk resistivity and grain boundary resistivity can be obtained, as listed in Table 1. Both the bulk resistivity Rg and the grain boundary resistivity Rgb decline as sintering temperature rises from 1700 ℃ to 1800 ℃. Remarkably, Rgb of the sample sintered at 1800 ℃ is only about half of that sintered at 1700 ℃. Moreover, the grain boundary response disappears when sintering
Fig. 5. (a) Frequency dependence of the dielectric constant ε’, and (b) impedance complex plane plots of SiC-MoSi2 composites with different sintering temperatures. 3985
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Fig. 7. (a) Spectral emissivity of the composite samples at room temperature, and (b) the corresponding total emissivity as a function of temperature.
infrared emissivity despite two aspects mentioned earlier. The decreased interfaces and the enhanced lattice distortion improve the emissivity, while the reduced porosity degenerates it. Therefore, the emissivity remains almost the same when sintering temperature rises from 1700 ℃ to 1800 ℃. Fig. 7b also shows that the emissivity of all SiC-MoSi2 composites rises with the increment in temperature. The composite sintered at 2000 ℃ maintains the highest emissivity over the testing temperature range, increasing continuously from 0.721 at room temperature to 0.851 at 700 ℃. This high temperature emissivity result is higher than that of C/SiC composites (˜0.8 at 1000 K) [39]. Due to the direct relationship between infrared emission and free electrons activities and lattice vibration in the near- and mid-infrared region according to the thermal radiation theory [40], a higher temperature guarantees enhanced transition among electron energy levels and lattice vibration, leading to an increase in the total emissivity of the composites.
grain boundary to the overall resistivity. On the other hand, a little higher doping level can be achieved due to a larger solubility limit of Al in α-SiC than in β-SiC [34]. Hence, it is assumed that the Al-doped SiC grains at elevated sintering temperature improve electrical conduction. Meanwhile, grain growth of SiC grains also plays a vital role in electrical transport since the number of grain boundaries decreases with increasing sintering temperature. Fig. 6 illustrates the electrical resistivity of the sintered composites measured at different temperatures. As is clearly observed, the increase in testing temperature causes the electrical resistivity to decrease drastically. Such a phenomenon is considered to follow the electrical transport mechanism in semiconductors: more electrons from the valence band are agitated into the conduction band at a higher temperature, leading to an exponential increase in the conductivity. The lowest electrical resistivity value of 5.4 × 10−4 Ω·cm is recorded at 700 ℃ for the composite sintered at 2000 ℃. Fig. 7a shows the spectral emissivity of SiC-MoSi2 composites at room temperature. A pronounced V-like spectrum during 10–14 μm exists in the spectral emissivity of all samples. This refers to an infrared emission feature of monolithic SiC [35]. In the lower wavelength range, these spectra fluctuate at the emissivity value of 0.8 and reach the peak value of ˜0.85 at ˜10 μm. There are merely slight differences except that the minimum value of spectral emissivity for composites sintered at 2000 ℃ is much higher compared to that of the other materials. Based on the definition of thermal radiation, the total emissivity at a given temperature εT from λ1 to λ2 can be deduced from the spectral emissivity via integration, as expressed below [36]: λ
εT =
∫λ1 2 ελ Ebλ dλ λ
∫λ1 2 Ebλ dλ
4. Conclusions In summary, the potential of pressureless sintered SiC-MoSi2 composites for bulk infrared source applications is studied. Sintering temperature was tuned to prepare materials with various microstructures and its effects on the electrical and infrared emission properties are systematically inspected. Increasing sintering temperature from 1700 ℃ up to 2000 ℃ causes continuously electrical resistivity reduction. The initial decrease in electrical resistivity from 1700 °C to 1800 °C is primarily ascribed to pore elimination during densification. As sintering temperature further ascends, the thinned amorphous layer located at SiC/MoSi2 interface is identified to minimize the contribution of grain boundaries to the overall resistivity and give rise to electrical percolation. Consequently, electrical linearity is obtained with a low resistivity of 7.2 Ω·cm when sintered at 2000 ℃. Moreover, the infrared emissivity becomes higher with increasing sintering temperature, presumably due to the reduced interface and the presence of Mo4.8Si3C0.6 compound. The emissivity of the composite sintered at 2000 ℃ is 0.721 at room temperature and reaches as high as 0.851 at 700 ℃.
λ
=
∑λ12 ελ Ebλ Δλ λ
∑λ12 Ebλ Δλ
(5)
where Ebλ represents the emission intensity of blackbody at wavelength λ and temperature T. As a result, the total emissivity of SiC-MoSi2 composites as a function of temperature can be obtained, as shown in Fig. 7b. At room temperature, increasing sintering temperature greatly enhances the total emissivity. It is supposed that the reduced number of interfaces induced by grain growth of SiC would decrease the reflection and scattering of electromagnetic waves and enhance the energy absorption, leading to the improved infrared emissivity. Moreover, the formation of Mo4.8Si3C0.6 might be another important factor. Though the infrared emission properties of this compound are rarely reported, Parthé et al. [37] stated that there is considerable irregular lattice in the crystal of Mo4.8Si3C0.6. Thus, lattice distortion in Mo4.8Si3C0.6 is assumed to enable asymmetry and lattice vibration of structure cell (phono absorption), which effectively enhances the emissivity. In the case of the composites sintered at 1700 ℃ and 1800 ℃, the total emissivity seems identical. Since a porous structure allows for the multiple reflection and absorption of electromagnetic waves [38], pore elimination may become a side effect of densification process on the
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sintered SiC ceramics, Ceram. Int. 41 (7) (2015) 9107–9114. [3] M. Herrmann, K. Sempf, M. Schneider, U. Sydow, K. Kremmer, A. Michaelis, Electrochemical corrosion of silicon carbide ceramics in H2SO4, J. Eur. Ceram. Soc. 34 (2) (2014) 229–235. [4] Y.-K. Dou, J.-B. Li, H.-B. Jin, X.-Y. Fang, M.-S. Cao, Enhanced photoconductivity of 3C-SiC by Al/N codoping, J. Appl. Phys. 114 (10) (2013). [5] K.J. Kim, K.-Y. Lim, Y.-W. Kim, X.D. Zhou, Control of electrical resistivity in silicon carbide ceramics sintered with aluminum nitride and yttria, J. Am. Ceram. Soc. 96 (11) (2013) 3463–3469. [6] G. Neuer, G. Jaroma-Weiland, Spectral and total emissivity of high-temperature materials, Int. J. Thermophys. 19 (3) (1998) 917–929. [7] E. Pike Gordon, Semiconducting polycrystalline ceramics, in: M.V. Swain (Ed.), Materials Science and Technology, VCH, Weinheim, 2006, pp. 731–754. [8] T.-Y. Zhou, X.-C. Liu, C.-C. Dai, W. Huang, S.-Y. Zhuo, E.-W. Shi, Effect of graphite related interfacial microstructure created by high temperature annealing on the contact properties of Ni/Ti/6H-SiC, Mater. Sci. Eng., B 188 (2014) 59–65. [9] T.-Y. Cho, R. Malik, Y.-W. Kim, K.J. Kim, Electrical and mechanical properties of pressureless sintered SiC-Ti2CN composites, J. Eur. Ceram. Soc. 38 (9) (2018) 3064–3072. [10] J. Chen, J. Chen, Z. Chen, X. Liu, Z. Huang, Y. Huang, Potential-current characteristics in SiC/ZrB2 composite ceramics, J. Eur. Ceram. Soc. 38 (6) (2018) 2477–2485. [11] C. Chen, Y. Tan, X. Han, H. Luo, S. Zeng, S. Peng, H. Zhang, Enhanced electromagnetic interference shielding properties of silicon carbide composites with aligned graphene nanoplatelets, J. Eur. Ceram. Soc. 38 (16) (2018) 5615–5619. [12] K. Yang, H. Mei, D. Han, L. Cheng, Enhanced electromagnetic shielding property of C/SiC composites via electrophoretically-deposited CNTs onto SiC coating, Ceram. Int. 44 (16) (2018) 20187–20191. [13] A.K. Vasudévan, J.J. Petrovic, A comparative overview of molybdenum disilicide composites, Mater. Sci. Eng., A 155 (1) (1992) 1–17. [14] A.A. Sharif, High-temperature oxidation of MoSi2, J. Mater. Sci. 45 (4) (2010) 865–870. [15] D. Sciti, L. Silvestroni, A. Balbo, S. Guicciardi, G. Pezzotti, High-strength and -toughness electroconductive SiC-based composites, Adv. Eng. Mater. 8 (10) (2006) 997–1001. [16] W. Guo, H. Xiao, P. Gao, W. Xie, Q. Li, J. Hu, Investigation of MoSi2 melt infiltrated RSiC and its oxidation behavior, Ceram. Int. 38 (1) (2012) 111–117. [17] R. Suresh Kumar, D. Sivakumar, K. Venkateswarlu, A.S. Gandhi, Mechanical behavior of molybdenum disilicide reinforced silicon carbide composites, Scripta Mater. 65 (9) (2011) 838–841. [18] J. Mao, S. Ding, Y. Li, S. Li, F. Liu, X. Zeng, X. Cheng, Preparation and investigation of MoSi2/SiC coating with high infrared emissivity at high temperature, Surf. Coat. Technol. 358 (2019) 873–878. [19] P.-Z. Gao, X.-L. Zhang, S.-T. Huang, J.-X. Liu, W.-X. Wang, H.-N. Xiao, Influence of MoSi2 content on the microstructure and properties of MoSi2–RSiC composites, Mater. Res. Bull. 54 (2014) 13–19. [20] X. Zhang, Z. Lu, Z. Jin, Electrical resistivity and microstructure of pressureless reactive sintered MoSi2–SiC composite, Mater. Chem. Phys. 86 (1) (2004) 16–20. [21] J.-Q. Zheng, J. Chen, X.-J. Liu, Z.-R. Huang, J.-J. Chen, New insights into linear electrical properties of pressureless sintered SiC-MoSi2-AlN composites, J. Eur.
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Original Article
Electrical percolation and infrared emissivity of pressureless sintered SiCMoSi2 composites tailored by sintering temperature
T
⁎
Jia-Qi Zhenga,b, Jian Chena, , Bu-Hao Zhanga,b, Xue-Jian Liua, Zhong-Ming Chena, Hai-Bo Wua, ⁎ Zheng-Ren Huanga, a State Key Laboratory of High Performance Ceramics and Superfine Microstructure, Shanghai Institute of Ceramics, Chinese Academy of Sciences, Shanghai, 200050, China b University of Chinese Academy of Sciences, Beijing, 100049, China
A R T I C LE I N FO
A B S T R A C T
Keywords: Silicon carbide Molybdenum disilicide Electrical percolation Grain boundary Sintering temperature
SiC-MoSi2 composites with low electrical resistivity and high infrared emissivity were fabricated via pressureless sintering. The relationship between microstructure evolution and electrical behaviors along with infrared emission properties of the resulting composites is investigated at various sintering temperatures. The electrical resistivity undergoes two significant drops with increasing sintering temperature. Pore elimination bears responsible for the initial decrease in electrical resistivity. Transmission electron microscopy (TEM) observation manifests that the thinned amorphous layers at SiC/MoSi2 interface decrease grain boundary resistivity and allow for electrical percolation to occur when sintering temperature further rises. Additionally, increasing sintering temperature leads to a higher infrared emissivity owing to the formation of Mo4.8Si3C0.6 and the decreased boundaries. The lowest electrical resistivity of 7.2 Ω cm and the highest infrared emissivity of 0.721 are recorded for composite sintered at 2000 ℃. Overall, SiC-MoSi2 composites exhibit a promising prospect as infrared source elements that must endure harsh environments.
1. Introduction Among a wide range of structural and electrical ceramic materials, silicon carbide (SiC) stands out by virtue of its superb physical and chemical properties, such as superior mechanical strength, high thermal conductivity, excellent corrosion resistance and a large bandgap [1–5]. In addition, SiC exhibits high infrared emissivity (˜0.9 above 1200 K [6]), which shows great promise in high temperature electrothermal materials such as infrared source. In this type of applications, a linear current-voltage (I-V) characteristic is basically required. Unfortunately, SiC ceramics usually suffer from relatively low electrical conductivity and nonlinear I-V behavior, both of which are derived from a double Schottky barrier at grain boundaries [7]. With continued pursuit for developing SiC-based infrared source elements, it has become imperative to tailor the electrical properties of SiC ceramics at no expense of infrared emissivity. So far, incorporating an electrically conductive filler over a SiC matrix has been considered a more convenient and versatile method compared to elemental doping. Though ohmic contacts with low contact resistivity can be obtained with the addition of metals (Ni, Ti, etc. [8]), the tendency of SiC to react with metals to form carbide or ⁎
silicides may threat long term applications. There has been several recent success in producing highly conductive SiC ceramics by introducing TiN [9], ZrB2 [10], graphene nanoplatelets [11] and carbon nanotubes [12]. These conductive fillers could generally decrease the Schottky barrier height (SBH) once they are uniformly dispersed over the matrix, and even deliver linear electrical properties. Given the harsh working environment of infrared source, MoSi2 becomes a leading candidate for electrically conductive phase due to its high melting point (˜2030 ℃), very low electrical resistivity (2.15 × 10−5 Ω·cm), and distinctive oxidation resistance [13,14]. Much attention has been paid to the mechanical properties and oxidation behaviors of SiC-MoSi2 composites [15–17]. Recently, Mao et al. [18] successfully fabricated SiC-MoSi2 composite coatings with high infrared emissivity and high thermal shock resistance. Besides, some reports mentioned that MoSi2 phase effectively decreased the electrical resistivity of SiC ceramics [19,20]. However, the electrical conduction mechanism of SiC-MoSi2 composites is still left unanswered. Our previous work [21] reported a percolation phenomenon in SiCMoSi2 composites. There is a percolation threshold in the amount of MoSi2 particles over which the electrical behavior changes dramatically. And this phenomenon is strongly associated with the Schottky
Corresponding authors. E-mail addresses: [email protected] (J. Chen), [email protected] (Z.-R. Huang).
https://doi.org/10.1016/j.jeurceramsoc.2019.05.019 Received 4 April 2019; Accepted 14 May 2019 Available online 16 May 2019 0955-2219/ © 2019 Elsevier Ltd. All rights reserved.
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Table 1 Relative density and electrical properties of SiC-MoSi2 composites as a function of sintering temperature. Sintering temperature (℃)
1700 1800 1900 2000
Relative density (%)
86.6 94.5 94.7 96.0
Electrical properties Electrical resistivity (Ω cm)
Nonlinearity exponent
Rg (Ω cm)
Rgb (Ω cm)
2519.2 266.0 35.2 7.2
1.271 1.208 1.104 1.034
173.1 ± 62.3 148.6 ± 8.8 28.1 ± 1.4 21.4 ± 1.1
(3.173 ± 0.011) ×104 (1.484 ± 0.004) ×104 – –
For electrical resistivity measurement, the silver paste was applied to both surfaces of the disk-shaped specimens of 15 mm in diameter and 2 mm in thickness followed by heated at 600 ℃ for 10 min. The impedance spectra in the frequency range from 20 Hz to 120 MHz were examined using an AC (alternating current) impedance spectrometer (Agilent HP4294A, Hewlett-Packard, Canada). The applied voltage was fixed at 0.5 V. Every component parameter in the equivalent circuit for the measured impedance plots was obtained via ZSimpWin software package (EChem Software, Michigan, USA). The DC (direct current) potential-current characteristics were measured in an Ar atmosphere from room temperature up to 700 ℃ by a high resistance meter (Keithley 6517A, Keithley, USA). The electrical resistivity of the specimens was determined from the I-V data at 1 mA·cm−2. The infrared emissivity properties of SiC-MoSi2 composites were obtained using a high temperature spectrum measurement system. The system mainly consisted of a Fourier-transform infrared (FTIR) spectrometer (NICOLET Is10, Thermo Fisher, USA) with a liquid nitrogencooled detector and a resistance furnace. The infrared spectrum of the specimens was detected within the range of 2.5–16 μm. Emission from the polished surface of each specimen in a vacuum at 25 ℃ and in the range of 100–700 ℃ was measured to minimize the effects of surface roughness and oxidation behaviors.
barrier at grain boundary of SiC ceramics. Based on the Schottky-Mott theory [22], the SBH of a SiC-MoSi2 system should be primarily dependent on the work function of MoSi2 and the electron affinity of SiC, but there is more to study. Recently, Kim et al. [23] indicated that the presence of thin intergranular films between SiC grains resulted in highly resistive SiC ceramics. Gu et al. [24] directly observed boron segregation to grain boundaries in hot-pressed SiC ceramics. These findings let us ponder how the quality of SiC/MoSi2 interface influences the electrical transport in SiC-MoSi2 composites, in terms of chemical interaction or microstructure evolution of grain boundaries with various sintering temperatures. In this study, pressureless sintering was performed for the synthesis of SiC-MoSi2 composites. Y2O3 along with AlN was used as sintering aids to secure densification. The objective of this work is to investigate the microstructure evolution, electrical behaviors and infrared emission properties of SiC-MoSi2 composites. The effect of sintering temperature on microstructure, electrical percolation, and infrared emissivity is systematically discussed. 2. Experimental procedure Commercially available submicron β-SiC (˜0.5 μm, ≥99% pure, EnoMaterial Co., Ltd., Hebei, China) and MoSi2 (˜1.0 μm, ≥99.9% pure, Haoxi Research Nanomaterials Inc., Shanghai, China) were used as the starting powders; 5 wt% AlN (˜2.0 μm, ≥99% pure, Tokuyama Co., Tokyo, Japan) and 2 wt% Y2O3 (≥99% pure, Aladdin Industrial Co., Ltd., Shanghai, China) powders were used as the sintering aids; phenolic resin (Shanghai Tiao Chemical Co., Ltd., Shanghai, China) was used as the binder for green compacts. Based on our previous research [21], the MoSi2 content in the composites was fixed at 15 wt% which was approximately close to the electrical percolation threshold sintered at 1900 ℃. The mixture was blended with 5 wt% phenolic resin in ethanol and ball-milled for 24 h using SiC media in a polypropylene jar. Then the slurry was dried at 80 ℃ for 12 h and sieved (100 mesh). The as-received powder mixture was pressed uniaxially into disks under 10 MPa followed by cold isostatically pressed under 200 MPa. The green bodies were kept at 1000 ℃ for 1 h in a vacuum to ensure pyrolysis of the phenolic resin. Finally, the pressureless sintering of the samples was carried out in a graphite resistance furnace at 1700 ℃, 1800 ℃, 1900 ℃, and 2000 ℃ for 1 h in an argon atmosphere, respectively. The heating and cooling rates were fixed at 10 ℃/min. The relative density of SiC-MoSi2 composites was measured by the Archimedes method. The corresponding theoretical density was determined according to the rule of mixtures. Phase identification for sintered samples was performed by an X-ray diffractometer (XRD, D8 Advance, Bruker, Germany) with Cu Kα radiation (λ =1.5406 Å). The diffractograms were scanned in 2θ from 10° to 80° at a rate of 0.04°/s. The microstructures of polished samples were characterized by scanning electron microscopy (SEM, Magellan 400, FEI, Netherland). Further detailed information about the microstructures was obtained via TEM (JEM-2100F, JEOL, Japan) equipped with energy-dispersive spectroscopy (EDS), which was carried out at 200 kV. Specimens for TEM were prepared conventionally by cutting 3 mm discs from the assintered bulk. The specimens were mechanically ground and further thinned by ion beam milling.
3. Results and discussion The relative densities of the as-prepared SiC-MoSi2 composites are listed in Table 1. Most of the samples can be densified to nearly 95% of the theoretical density except that sintered at 1700 ℃. During liquidphase sintering, Y2O3 and AlN react with the native SiO2 layer which is present on the surface of SiC and MoSi2 particles, to form Y-Al-Si-O-C-N melt with the dissolution of SiC [5]. The pyrolysis of phenolic resin yields ˜60% excess carbon, which is generally believed to facilitate densification process by reacting with the SiO2 layer and enhancing bulk self-diffusion of SiC as well [25]. Fig. 1 presents the XRD patterns of SiC-MoSi2 composites sintered at various temperatures. Apparently, all the samples exhibit characteristic peaks of β-SiC (3C), α-SiC (6H, 4H, and 2 H), AlN, and MoSi2 as major phases. The β → α phase transformation of SiC with increasing sintering temperature can be discerned. The introduction of Al or Al compounds could also promote this process by stimulating the formation of point defects in the SiC lattice to enhance mass transport [26]. There is no trace of Y2O3 or Y3Al5O12 presumably due to the detection limit. In addition, a small amount of Mo4.8Si3C0.6 compound (JCPDS 731380) is detected in all samples. The peaks of Mo4.8Si3C0.6 heighten slightly with those of MoSi2 weakening at the elevated sintering temperature. A synchronous shift of the peaks of SiC and MoSi2 in Fig. 1b shows evidence of the formation of Mo4.8Si3C0.6 phase. This compound possibly originates from the reactions of partial SiC and MoSi2 at high temperature [16]: 8MoSi2 → Mo5Si3 + Mo3Si + 12 Si
(1)
3/5 Mo5Si3 + 3/5 Mo3Si + 3/5 SiC → Mo4.8Si3C0.6
(2)
The typical SEM micrographs of the SiC-MoSi2 composites with elevated sintering temperature are illustrated in Fig. 2. The most 3982
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Fig. 1. XRD patterns of (a) SiC-MoSi2 composites with various sintering temperatures, and (b) magnified detail of 2θ range of [38°, 39°] in (a).
(Fig. 3f). Therefore, the increase in sintering temperature allows for the removal of the amorphous film present at SiC/MoSi2 interface. It is worth noting that the SAED pattern (inset of Fig. 3f) and the EDS result (Fig. 3g) taken from the dark phase in Fig. 3e provide direct evidence of the formation of Mo4.8Si3C0.6 phase, which is consistent with the XRD data shown in Fig. 1. The current density versus electrical field curves at room temperature for SiC-MoSi2 composites are shown in Fig. 4. Nonlinear I-V characteristics are observed in the specimens sintered below 2000 ℃. The following relation [29] can well fit these curves:
obvious differences in the microstructures are the grain size and shape of SiC (grey and dark phase). The increased sintering temperature not only leads to the grain growth of SiC but accelerates the occurrence of elongated grains, which indicates the presence of β → α phase transformation. A number of pores can be found in the composite sintered at 1700 ℃. Instead, the other samples display a more robust microstructure when sintering temperature becomes higher. Irregularly shaped MoSi2 particles are homogeneously dispersed over the SiC matrix. According to the EDS results in our recent report [21], insulating Y2O3 is mainly confined to the multigrain junctions. TEM was employed to further reveal the microstructure changes when sintering temperature rises from 1800 ℃ to 2000 ℃, as shown in Fig. 3. The results suggest that trace amount of sintering necks remain in the bulk sintered at 1800 ℃ (Fig. 3a). On the contrary, SiC and MoSi2 grains seem to be more strongly bonded at a higher sintering temperature (Fig. 3c and e). The corresponding HRTEM image (Fig. 3b) demonstrates a typical amorphous characteristic of SiC/MoSi2 interface sintered at 1800 ℃. The amorphous layer, about 2–3 nm in thickness, is estimated to be the Y-Al-Si-O-C-N melt It occurs during liquid-phase sintering and remains at grain boundaries after cooling [27,28]. It can be identified that the amorphous layer becomes thinner when sintering temperature rises (Fig. 3d), and finally disappears up to 2000 ℃
I ∝ Uα
(3)
where α is the nonlinearity exponent. The fitted α value accompanied by the electrical resistivity is presented in Table 1. It is suggested that increasing sintering temperature results in both the weakened electrical nonlinearity and the decreased electrical resistivity. The composite sintered at 2000 ℃ gives an ohmic nature, which implies electrical percolation is already obtained. And it also possesses the lowest electrical resistivity (7.2 Ω·cm) compared to the other samples. In Fig. 5a–b, the frequency dependence of the dielectric constant ε’ and the impedance complex plane plots for the measured samples are demonstrated. In Fig. 5a, the composites sintered below 1900 ℃ exhibit
Fig. 2. Typical microstructures of SiC-MoSi2 composites sintered at (a) 1700 ℃, (b) 1800 ℃, (c) 1900 ℃, and (d) 2000 ℃. 3983
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Fig. 3. TEM images of SiC-MoSi2 composites sintered at (a) 1800 ℃, (c) 1900 ℃, and (e) 2000 ℃; (b), (d), and (f) are HRTEM images enlarged from the yellow square in (a), (c), and (e); insets in (b), (d), and (f) are the corresponding SAED images; (g) EDS results taken from the dark phase in (e) (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article).
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Fig. 4. I-V characteristics of SiC-MoSi2 composites sintered at different temperatures.
Fig. 6. Electrical resistivity changes with testing temperature for SiC-MoSi2 composites sintered at different temperatures.
frequency dependent throughout the frequency range, which is a typical feature of a capacitive behavior [30]. Whereas the samples sintered at higher temperatures show an almost flat plateau in ε’, which is characteristic of conductive materials. The flat-to-dispersive transition in ε’ occurs at the critical frequency (fm). The Debye model [31] can be used to describe the dielectric relaxation:
temperature reaches 1900 ℃, which implies Rgb no longer dominates the overall resistivity and the onset of electrical percolation has been triggered. Further rising sintering temperature results in a gradual decrease in the electrical resistivity. The conductive pathways are deemed to curl or even spiral when concerning occurrence of inductance in the composites sintered above 1800 ℃. Accordingly, it is safe to believe there is an intimate relationship between the microstructure of the composite material and its electrical properties. The grain boundary effect is responsible for the enhanced electrical percolation behavior, whereas other aspects should be taken into consideration as well. The decreased electrical resistivity when sintering temperature rises from 1700 ℃ to 1800 ℃ might greatly benefit from pore elimination during densification. It is widely acknowledged that the removal of pores can enhance grain boundary structures and prevent gas phase from blocking current path [32]. However, given the similar relative density for the composites sintered above 1800 ℃ (Table 1), this factor could be nearly excluded. Specifically, the electrical nonlinearity weakens sharply and electrical percolation emerges when sintering temperature increases to 1900 ℃. From the microstructural perspective, the amorphous film with a high disorder and intrinsically insulating nature contributes to the formation of the depletion region at grain boundaries. It was reported that reduced depletion layer width would decrease potential barrier and consequently lead to notably electrical conduction [33]. Thus, this electrical behavior evolution is mainly attributed to the thinner amorphous layer at SiC/MoSi2 interface. At the final increment in sintering temperature up to 2000 ℃, the decrease in electrical resistivity still proceeds yet in a more gradual manner. Devitrification of the intergranular glass at SiC/ MoSi2 interface (Fig. 3f) is almost complete without any contribution of
ωm τ = 2πfm τ = 1
(4)
where ωm is the angular frequency of dielectric relaxation and τ is the relaxation time. Obviously, fm moves to a higher frequency with increasing sintering temperature, leading to decreased τ. This suggests the weakened interfacial polarization in the composite structures, i.e., capacitive coupling at grain boundaries. Fig. 5b further illustrates how grain boundary effect shapes the electrical behaviors. The plots for samples sintered below 1900 ℃ show a typical compressed semicircular arc at low frequency which is also attributed to the relaxation process at grain boundaries. To elaborate the contributions of grain and grain boundary to the electrical resistivity, accurate fitting to the equivalent circuit of the plots is performed, as shown in Fig. 5b. The equivalent series circuit is composed of one bulk component and one parallel grain boundary component where grain boundary resistance and corresponding constant phase element (CPE) are included. Hence, the bulk resistivity and grain boundary resistivity can be obtained, as listed in Table 1. Both the bulk resistivity Rg and the grain boundary resistivity Rgb decline as sintering temperature rises from 1700 ℃ to 1800 ℃. Remarkably, Rgb of the sample sintered at 1800 ℃ is only about half of that sintered at 1700 ℃. Moreover, the grain boundary response disappears when sintering
Fig. 5. (a) Frequency dependence of the dielectric constant ε’, and (b) impedance complex plane plots of SiC-MoSi2 composites with different sintering temperatures. 3985
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Fig. 7. (a) Spectral emissivity of the composite samples at room temperature, and (b) the corresponding total emissivity as a function of temperature.
infrared emissivity despite two aspects mentioned earlier. The decreased interfaces and the enhanced lattice distortion improve the emissivity, while the reduced porosity degenerates it. Therefore, the emissivity remains almost the same when sintering temperature rises from 1700 ℃ to 1800 ℃. Fig. 7b also shows that the emissivity of all SiC-MoSi2 composites rises with the increment in temperature. The composite sintered at 2000 ℃ maintains the highest emissivity over the testing temperature range, increasing continuously from 0.721 at room temperature to 0.851 at 700 ℃. This high temperature emissivity result is higher than that of C/SiC composites (˜0.8 at 1000 K) [39]. Due to the direct relationship between infrared emission and free electrons activities and lattice vibration in the near- and mid-infrared region according to the thermal radiation theory [40], a higher temperature guarantees enhanced transition among electron energy levels and lattice vibration, leading to an increase in the total emissivity of the composites.
grain boundary to the overall resistivity. On the other hand, a little higher doping level can be achieved due to a larger solubility limit of Al in α-SiC than in β-SiC [34]. Hence, it is assumed that the Al-doped SiC grains at elevated sintering temperature improve electrical conduction. Meanwhile, grain growth of SiC grains also plays a vital role in electrical transport since the number of grain boundaries decreases with increasing sintering temperature. Fig. 6 illustrates the electrical resistivity of the sintered composites measured at different temperatures. As is clearly observed, the increase in testing temperature causes the electrical resistivity to decrease drastically. Such a phenomenon is considered to follow the electrical transport mechanism in semiconductors: more electrons from the valence band are agitated into the conduction band at a higher temperature, leading to an exponential increase in the conductivity. The lowest electrical resistivity value of 5.4 × 10−4 Ω·cm is recorded at 700 ℃ for the composite sintered at 2000 ℃. Fig. 7a shows the spectral emissivity of SiC-MoSi2 composites at room temperature. A pronounced V-like spectrum during 10–14 μm exists in the spectral emissivity of all samples. This refers to an infrared emission feature of monolithic SiC [35]. In the lower wavelength range, these spectra fluctuate at the emissivity value of 0.8 and reach the peak value of ˜0.85 at ˜10 μm. There are merely slight differences except that the minimum value of spectral emissivity for composites sintered at 2000 ℃ is much higher compared to that of the other materials. Based on the definition of thermal radiation, the total emissivity at a given temperature εT from λ1 to λ2 can be deduced from the spectral emissivity via integration, as expressed below [36]: λ
εT =
∫λ1 2 ελ Ebλ dλ λ
∫λ1 2 Ebλ dλ
4. Conclusions In summary, the potential of pressureless sintered SiC-MoSi2 composites for bulk infrared source applications is studied. Sintering temperature was tuned to prepare materials with various microstructures and its effects on the electrical and infrared emission properties are systematically inspected. Increasing sintering temperature from 1700 ℃ up to 2000 ℃ causes continuously electrical resistivity reduction. The initial decrease in electrical resistivity from 1700 °C to 1800 °C is primarily ascribed to pore elimination during densification. As sintering temperature further ascends, the thinned amorphous layer located at SiC/MoSi2 interface is identified to minimize the contribution of grain boundaries to the overall resistivity and give rise to electrical percolation. Consequently, electrical linearity is obtained with a low resistivity of 7.2 Ω·cm when sintered at 2000 ℃. Moreover, the infrared emissivity becomes higher with increasing sintering temperature, presumably due to the reduced interface and the presence of Mo4.8Si3C0.6 compound. The emissivity of the composite sintered at 2000 ℃ is 0.721 at room temperature and reaches as high as 0.851 at 700 ℃.
λ
=
∑λ12 ελ Ebλ Δλ λ
∑λ12 Ebλ Δλ
(5)
where Ebλ represents the emission intensity of blackbody at wavelength λ and temperature T. As a result, the total emissivity of SiC-MoSi2 composites as a function of temperature can be obtained, as shown in Fig. 7b. At room temperature, increasing sintering temperature greatly enhances the total emissivity. It is supposed that the reduced number of interfaces induced by grain growth of SiC would decrease the reflection and scattering of electromagnetic waves and enhance the energy absorption, leading to the improved infrared emissivity. Moreover, the formation of Mo4.8Si3C0.6 might be another important factor. Though the infrared emission properties of this compound are rarely reported, Parthé et al. [37] stated that there is considerable irregular lattice in the crystal of Mo4.8Si3C0.6. Thus, lattice distortion in Mo4.8Si3C0.6 is assumed to enable asymmetry and lattice vibration of structure cell (phono absorption), which effectively enhances the emissivity. In the case of the composites sintered at 1700 ℃ and 1800 ℃, the total emissivity seems identical. Since a porous structure allows for the multiple reflection and absorption of electromagnetic waves [38], pore elimination may become a side effect of densification process on the
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