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Heat treatment of commercial polydimethylsiloxane PDMS precursors: Part II. Thermal properties of carbon-based ceramic nanocomposites
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Heat treatment of commercial polydimethylsiloxane PDMS precursors: Part II. Thermal properties of carbon-based ceramic nanocomposites Srisaran Venkatachalama,1, Stéphane Lenfanta, Michael Depriesterb, Abdelhak Hadj Sahraouib, Djamila Hourliera,* a b
IEMN (CNRS, UMR 8520), Université Lille 1, Avenue Poincaré, CS60069, F-59652, Villeneuve d'Ascq Cedex, France Laboratoire de Thermophysique de la Matière Condensée, Université du Littoral Côte d’Opale, 145 Avenue Maurice Schumann, F-59140, Dunkerque Cedex, France
A R T I C LE I N FO
A B S T R A C T
Keywords: Polymer-derived ceramics Polydimethylsiloxane Thermal conductivity Electrical conductivity Thermoelectric response
Organosilicon polymer-derived ceramic (PDC) materials have been extensively studied for their various properties such as mechanical, chemical and electrical properties. One of the least explored properties of PDC is their thermal properties, required for designing a heat energy converter. In this work, a commercially available polydimethylsiloxane elastomer is used as starting polymeric precursor. Pyrolysis of cross-linked gels conducted in inert atmosphere above 1000 °C yielded silicon oxycarbide SiOxCy, as revealed by X-ray diffraction analysis. The complex shapes of the gels are retained even after pyrolysis up to 1500 °C, however a linear shrinkage ratio of 25–30% is observed. Both thermal and electrical conductivity of free carbon based SiOxCy residues are found to rise simultaneously as the heat treatment temperature increased. The sample heat-treated at 1500 °C exhibits the room temperature thermal conductivity of 2.24 Wm-1K-1 and electrical conductivity of 0.021 S cm-1. These properties of PDCs suggest their appealing applications as a functional material in thermoelectric devices.
1. Introduction The silicon oxycarbide issued from pyrolysis of organo-polysiloxanes has seen much promise for its uses in many applications due to their intrinsic properties: chemical stability, high temperature stability in harsh atmospheres, low density, hardness, etc., [1–5]. Very recently, we have demonstrated that polydimethylsiloxane (PDMS) yields materials which exhibit variable absorption properties in terahertz band, making them attractive for bolometer applications [6]. The electromagnetic (EM) wave’s absorption is tightly controlled at desired values by means of control of the polymer-composition, the carbon content, the state of organization of the free carbon phase resulting from hydrocarbon groups present in polymer [7], and finally, shaping-structuring polymeric precursor. Actually, the main advantage of using polymers lies in their ability to form complex small parts in a near net-shape by the casting process. A broadband terahertz microbolometer with responsivity of 0.76 V/W, time constant of 180 ms, and noise equivalent power of 2 nW/√Hz, has been demonstrated as a proofof-concept device operating in atmospheric air [7]. However, the temperature sensing element used in the proof-of concept bolometer was based on the principle of change of an electrical resistance in close
contact with the rear face of the material and requires an external electrical power supply. In this paper we show that silicon oxycarbide issued from pyrolysis of organo-polysiloxanes are able to generate their own electrical voltage when submitted to a temperature gradient. Thus, active bolometers and energy harvesting converters based on silicon oxycarbide thermoelectric materials can be expected. These systems can be molded into any desired shape and allow a versatile use. To assess the capabilities of these materials for these applications, it would be reasonable to investigate other physical properties as their transport coefficients, such as electrical, thermal conductivity, to be of paramount importance. In theory, a thermoelectric material must possess high electrical conductivity (σ), high Seebeck coefficient (α), and low thermal conductivity (κ) as given by the thermoelectric figure-of-merit (zT), α2σ
zT = κ A few studies on the electrical conductivity of PDMS and thermal conductivity of silicon oxycarbide have been found in the literature. Johnson et al. [8] measured the electrical conductivity (d.c.) of the PDMS gel during the course of pyrolysis up to 700 °C, by supplying a very high voltage of 100–1000 V. The electrical conductivity of the polymer (measured in situ) increased linearly up to ~ 1 × 10-11 S cm-1
*
Corresponding author. E-mail address: [email protected] (D. Hourlier). 1 Present Address - Department of Mechatronics, Bannari Amman Institute of Technology, Sathyamangalam, India https://doi.org/10.1016/j.ceramint.2019.07.143 Received 11 January 2019; Received in revised form 23 June 2019; Accepted 12 July 2019 0272-8842/ © 2019 Elsevier Ltd and Techna Group S.r.l. All rights reserved.
Please cite this article as: Srisaran Venkatachalam, et al., Ceramics International, https://doi.org/10.1016/j.ceramint.2019.07.143
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and dense, the samples were fully covered with silver paint on the rear side, whereas a circular electrode was formed on the front side. The samples were dried at 120 °C for 30 min to eliminate the solvents present in the paint. After drying, the diameters of silver electrodes were 2–4 mm, mostly elliptical rather than circular. Semiconductor parameter analyzer (Agilent 4155C) was used to bias the sample in the voltage range 0–50 V. The measured current showed ohmic (linear) behavior. The electrical conductivity was calculated from the slope of the current versus voltage curve, when multiplied to the ratio of sample thickness to area of the front electrode. The mean sample thickness values were obtained using ten test measurements with a Mitutoyo thickness gauge, and the surface area of the electrodes was calculated by measuring the diameter under optical microscope. The thermal conductivity of samples was measured using photothermal radiometry. Briefly, the principle of measurement is explained as follows: a sample absorbs the incident laser beam, and generates heat, thereby acting as blackbody source that emits infrared radiations. Because of the inherent opacity of the sample, the emitted infrared radiation is proportional to the surface temperature of the sample. In the constructed experimental set-up, a laser source (Laser Quantum, Ventus HP532, wavelength = 532 nm) delivers a modulated beam directed towards the sample that is mounted horizontally on a holder. Two parabolic mirrors were used to collect the infrared radiations emitted by the sample, and focused towards a liquid N2-cooled HgCdTe infrared detector (Judson Technologies). The output of the detector signal was then amplified using a Lock-in amplifier (Ametek, Signal Recovery7260). The resulting signal is a frequency dependent complex signal S(f) which is proportional to the temperature variation (Tm) of the sample surface as a function of laser modulation frequency (f), according to equation (1).
at 200 °C in N2 atmosphere. Once the polymer started to degrade, the electrical conductivity varied non-linearly and reached a value of ~ 1.5 × 10-10 S cm-1 at 700 °C [8]. They found a correlation between the thermal decomposition and the electrical conductivity characteristics of the materials. Qiu et al. [9] measured the thermal conductivity of macro-porous SiOxCy issued from a polysiloxane precursor (without specifying the molecular structure) using the free-standing 3ω technique. The thermal conductivity varied between 0.041 Wm-1K-1 and 0.062 Wm-1K-1 for the samples pyrolysed at 650 °C and 1000 °C, respectively. The authors compared their experiemental values with the intrinsic thermal conductivity (2.44 Wm-1K-1) estimated from the law of mixtures using the thermal conductivity values of bulk phases (SiO2 = 1.34 Wm-1K-1, SiC = 1.44 Wm-1K-1 and C = 4.18 Wm-1K-1) and weight percent of each phase (SiO2 = 27.5 wt%, SiC = 35 wt% and C = 37.5 wt%). The large difference in thermal conductivity values was attributed to the presence of porosity in SiOxCy samples. Gurlo et al. [10] investigated thermal conductivity at different temperatures (up to 1200 K) of sintered powders of SiOxCy (pyrolyzed polymethylsilsesquioxane powders) using laser flash analysis. The sample sintered at 1100 °C (warm-pressed at 120 MPa, 180 °C) containing SiO2 = 72.0 vol%, SiC = 9.11 vol %, free carbon = 6.2 vol%, showed a thermal conductivity of 0.5 Wm-1K-1 that remains unchanged whatever the temperature of measurement, whereas the sample with composition SiO2 = 77.5 vol%, SiC = 9.11 vol%, free carbon = 13.35 vol%, obtained from hot-presssing (30 MPa) at 1600 °C exhibited a non-linear response. The values increased from ~ 0.8 Wm-1K-1 to ~ 1.7 Wm-1K-1 when the temperature increased from 150 K to 1200 K. Nevertheless, at room temperature, thermal conductivity increased from 0.5 Wm-1K-1 to 1.2 Wm-1K-1 when the volume % of free-carbon increased from 6.2 vol% to 13.35 vol% [10]. The main difference between the two ceramic samples lies in the content of free carbon. For PDMS Sylgard, we have demonstrated that the free carbon phase embedded into SiOxCy matrix is the most efficient phase for absorption of THz radiation [6]. An understanding of free carbon formation-organisation, and its dependance on thermal and electrical conductivity is therefore valuable for the development of efficient devices. Only scant data on physical properties of PDMS-derived SiCxOy has been reported in the literature [9,10]. The objective of this study is thus to determine electrical and thermal conductivity of carbon-embedded SiOxCy ceramic materials derived from commercial Sylgard®184 polydimethylsiloxane PDMS. These data will allow us to optimize pyrolysis conditions and also to adapt the chemical architecture of precursors to obtain the desired properties for specific applications. An attempt has been made to evaluate the thermoelectric properties of thermally converted PDMS.
S (f ) = K (f ) Tm =
K (f ) ηIo 1 + Rsbexp(−2σs Ls ) × 2κs σs 1 − Rsb exp(−2σs Ls )
(1)
Equation (1) was derived under the assumptions that the sample is opaque to both the excitation laser and the emitted infrared radiation, and using a one-dimensional heat propagation model across the different planes of the sample. The factor K(f) is an instrumental constant that depends on geometrical factors, spectral average emissivity of the sample surface, and Stephan-Boltzmann constant. This constant K(f) can be circumvent by a normalization procedure, which involves recording in two different configurations: one where the sample is suspended in the air, and the other where the sample stands on a liquid substrate. f is the laser modulation frequency, η is the absorbed fraction of the incident intensity Io, κs is the thermal conductivity of the sample Ls is the thickness of the sample σs is a factor related to the thermal diffusivity (αs) of the sample
2. Experimental The starting polymers were prepared using commercially available Sylgard®184 elastomer. The curing agent and base were mixed in a ratio of 1:10, part by volume, and stirred thoroughly in a clean beaker. The air bubbles and solvents were degassed in a vacuum chamber. The sol was transferred to a polystyrene dish, and allowed to cross-link at room temperature for minimum of 60 h. The gels in square sizes 1.5 × 1.5 cm2 were pyrolyzed in an alumina horizontal tube furnace at a heating rate of 2 °C/min in a flow of 50 ml/min of Argon (99.999% purity) at different temperatures i.e., 800 °C, 1000 °C, 1300 °C, and 1500 °C, as separate pieces. X-ray diffraction (XRD) patterns were recorded on a Bruker’s D8 Advance spectrometer between 10° and 90° with an increment of 0.05° and acquisition time of 2 s. For measurements of electrical conductivity, thermal conductivity and Seebeck coefficients, samples heated at various temperatures, for a short dwell time of 5 min, were cooled down to room temperature. A new specimen of the sample is used for each test measurement: thermal conductivity and electrical conductivity. The electrical conductivity (d.c.) of pyrolyzed samples was measured along their thickness in a two probe configuration. As the material is compact
σs = (1 + j )
πf αs
, j = √-1
Rsb is the reflection coefficient of the thermal waves between the sample s and the substrate b (air or water). Moreover, Rsb is related ⎛ es − eb ⎞ ⎝ eb eb ⎠ ⎜
to the thermal effusivity (e) as, Rsb =
⎟
⎛ es + eb ⎞ ⎝ eb eb ⎠ ⎜
.
⎟
For samples of thickness > 220 μm, the modulation frequency varies between 1 and 100 Hz. The thermal diffusivity αs and effusivity es of the samples can be extracted by fitting the normalized complex voltage to equation (1) using the Levenberg-Marquardt fit routine [11]. Thermal conductivity κs is related to thermal diffusivity and the thermal effusivity as κs = es√αs. The Seebeck coefficient was estimated via the photothermoelectric technique (PTE). Although it was initially designed for the thermal characterization of materials [12], this approach was successfully 2
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applied in the past to the determination of thermoelectric power with the assistance of IR thermography [13]. Both sides of the samples were sputtered with gold electrodes of diameter 0.8 cm, and then the electrodes were attached to low resistance cables. The cables were connected to a lock-in amplifier (EG&G 7260, 10 MΩ input impedance) allowing thermoelectric voltage |ΔV| measurement. The front side of the sample was illuminated using a frequency-modulated laser beam. As the laser radiation is being absorbed by the material, heat induced in the sample generates a temperature gradient along the thickness of the material. Here, instead of being measured as in Ref. [13], the temperature difference was estimated. When the sample is suspended in air, temperature difference ΔT between both sides of an opaque sample was calculated from the following relationship [12], assuming a fully absorbing medium:
ΔT =
1 + exp(−2σs Ls ) − 2exp(−σs Ls ) I0 × 2ks σs 1 − exp(−2σs Ls )
Fig. 1. Scanning electron microscope images of SiOxCy issued from pyrolysis of Sylgard ®184 PDMS at 1500 °C in flowing argon atmospheres (a) pyramidal structure (b) inverted pyramids. Influence of dwell time on the surface morphology (c) no dwell time (d) dwell time 15 min at 1500 °C.
(2)
Thermal diffusivity αs of the sample was determined from the photothermal radiometry experiment and the laser incident intensity I0 was given by the supplier data sheet. Seebeck coefficient is then obtained from the ratio of the amplitude of the photothermoelectric voltage |ΔV| to the estimated temperature difference modulus |ΔT|. A corrective factor was introduced due to the relative important value of the sample electrical resistance (Rs) compared to the internal input impedance of the lock-in amplifier (Rlock-in). Gold contribution to the Seebeck coefficient was neglected in the calculation. Absolute Seebeck coefficient value is then given by the following relation (3).
S =
|ΔV | Rlock − in × ΔT Rs + Rlock − in
(Fig. 1b). The formation of such structure is a consequence of carbothermal reduction reactions between the silicon oxycarbide and free carbon phases, leading to crystalline β-SiC solid phase and CO gas as byproduct that has been detected by mass spectrometry at m/z 28. The distribution of SiC and free carbon phases were identified by Hourlier et al. [15] using Raman mapping for the ceramic residue issued from pyrolysis of polyhydridomethylsiloxane (CH3-Si(H)-O-)n at 1450 °C for a dwell time of 150 min. During such a long dwell time, low crystalline SiC were homogeneously distributed in the matrix, whereas high crystalline SiC were observed on selected sites where the free carbon phase is totally consumed [15]. Depending on the dwell time, the content of high crystalline SiC and carbon can be controlled. In PDMS-derived gels, the presence of crystalline β-SiC phases were also identified using X-ray diffractograms (Fig. 2), which corroborates with the results obtained by Raman spectroscopy. Fig. 2 shows X-ray diffraction patterns of the ceramic residues as a function of heat-treatment temperature. The heat-treated sample at 1000 °C is totally amorphous. Further pyrolysis temperature increase triggers the formation of crystalline nuclei. For samples heated between 1300 °C and 1500 °C, the pattern consists of two broad signals centered at 2θ ≈ 60° and 72°. At 1500 °C and holding time of 15 min, the signals sharpen and new peaks appear due to the crystallization of β-SiC. This result well correlates with the results obtained by Raman and TG/MS analysis. The SiC formation is related to the well-known carbothermal reaction, in which either free carbon phase reacts with SiO issued from the decomposition of oxycarbide network SiOxCy, or free carbon phase and SiO2 phase present in
(3)
One of the advantages of the PTE technique is the small temperature difference (lower than 1 °C) between faces, allowing a good resolution. Nevertheless, the indirect estimation of the temperature difference from the received light flux is still a drawback in comparison with what was done in Ref. [13] This approach is a preliminary trial to estimate the potentialities of SiOxCy materials as thermoelectric elements. Other experiments will be held in the future with an especially dedicated setup to improve the accuracy on the measured Seebeck coefficient values. 3. Results and discussions 3.1. Shaping of ceramics and microstructure The main advantage of Sylgard PDMS is its ability to form desired shape needed for specific applications. As an example, scanning electron microscope (SEM) images of pyramidal net-shaped ceramics obtained up on pyrolysis of Sylgard PDMS at 1500 °C, is shown in Fig. 1 a and b. For terahertz/microwave absorbers, pyramidal structures either in normal and inverse configurations are necessary for the maximum trapping of light. Design and fabrication of such structures facilitate the incident beam to encounter several time on the pyramidal surfaces, so that the reflected beam from one pyramid is absorbed on the surface of neighboring pyramids [14]. Pyramidal structures were obtained by casting the Sylgard mixed precursors in appropriate mold, followed by demolding the cross-linked gel and subsequent pyrolysis. The molds having pyramidal heights ranging from 236 μm to 1 mm were used. Due to the mass loss of the gels during pyrolysis, the height of the pyramids were reduced from 850 μm to 642 μm and from 1 mm to 750 μm. Linear shrinkage ratios were found to vary between 25 and 30%. The surface of ceramic obtained at 1500 °C is initially smooth (Fig. 1c), but holding the sample only 15 min (dwell time) at the same temperature produced considerable changes on the surface. The sample surface exhibited different pore sizes ranging from 30 nm to 100 nm
Fig. 2. Influence of pyrolyzed conditions (temperature and dwell time) on XRD patterns of PDMS-derived ceramic residues. 3
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Fig. 3. Normalized amplitude and phase of thermal radiation collected from the sample surfaces.
the material. In either case the CO is eliminated at high temperature as shown by mass spectrometry in Part I [16]. It is rather surprising to see that the PDMS-derived ceramic does not exhibit any broad signal at 2θ between 22 and 26° characteristic of amorphous SiO2. Almost organosiloxane family [17,18] including MTES (CH3-SiO1.5)n, VTES (C2H3-SiO1.5)n, PTES (C6H5-SiO1.5)n, copolymers of DH (CH3-Si(H)-O-)n and TH (H-SiO1.5)n, methyl-substituted (SR350) [18] or methyl/phenyl-substituted (SR355) [18] polysiloxane resins, heated below 1500 °C indicate the presence of amorphous SiO2. The XRD pattern of the PDMS Sylgard heated at 1500 °C for 15 min resembles to that MTES, VTES and PTES heated at the same temperature but for a longer time (24 h) [19]. 3.2. Photothermal radiometry Thermal diffusivity and effusivity of the samples were extracted by fitting both amplitude and phase components received from the measured data. Fig. 3 depicts the normalized amplitude and phase signals obtained from PDMS samples pyrolyzed at various temperatures. For heat-treated sample at 800 °C, because the normalized amplitude is nearly flat, one can evaluate thermal parameters only from the normalized phase that contains the required information. A data fitting with argument of equation (1) was used to determine thermal parameters (thermal diffusivity, effusivity, and conductivity) values. At other temperatures, a combined data fitting of the normalized amplitude and phase according to equation (1) was applied. In normalized phase signal, the minimum crossing the zero line could be identified. As the treatment temperature increases, the minimum shifts from lower frequency to higher frequency.
Fig. 4. The room temperature electrical conductivity and thermal conductivity of SiOxCy ceramics issued from pyrolysis of Sylgard®184 polydimethylsiloxane.
resulting from thermal conversion of hybrid organic-inorganic precursor PDMS have been described as a composite made up of an intimate mixture of free carbon phase embedded into silicon oxycarbide SiOxCy network [20]. The starting Sylgard PDMS gel is an electrical insulator, having an ambient temperature electrical conductivity of 3.44 × 10-15 S cm-1 [8]. In contrast, for pyrolyzed materials, there is a progressive transition from insulating state to conductor. The electrical conductivity increases up to 0.021 S cm-1 for heat treated material at 1500 °C. Kim et al. [21] showed an improvement in electrical conductivities of SiOxCy system by sintering process of powders synthesized via thermal conversion of polysiloxane (YR3370, Momentive Performance Materials Japan Inc.). The electrical conductivity of sintered SiOxCy – based materials varies from 0.43 S cm-1 to 7.1 S cm-1 when the powders were sintered between 1450 °C and 1650 °C, respectively. The electrical conductivity values showed in Fig. 4 present statistical average of at least 5 measurements
3.3. Electrical and thermal conductivity Fig. 4 depicts the electrical conductivity and thermal conductivity, measured at room temperature 298 K, of residue materials produced at different pyrolysis temperatures from PDMS. These pyrolyzed materials 4
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Fig. 5. Thermoelectric response of PDMS heat treated at 1250 °C as a function of modulation frequency (a) the amplitude of the photothermoelectric voltage (b) apparent Seebeck coefficient.
whereas in amorphous carbon, it is proportional to the number of ordered polyaromatic rings. For the samples pyrolyzed between 1000 °C and 1250 °C, a non-linear increase in the ID/IG ratio with increasing pyrolysis temperature can be explained by the so-called carbonization process in which rearrangement of carbon units, and also dehydrogenation of polyaromatic rings take place simultaneously. Above 1250 °C, the Cfree phase embedded within amorphous SiOxCy is responsible for the conduction. The electrical conductivity increases with increasing pyrolysis temperature suggests that the electrically conductive of free carbon (Cfree) phase is connected through a percolation network. Indeed, the sharp monotonous increase in several orders of magnitude of electrical conductivity following the power law is indicative of percolated structure [22]. In contrast, the thermal conductivity follows a simple linear relationship with the pyrolysis temperature and no signature is observed for percolation in the thermal conduction path. The electrical percolation has no effect on thermal transport, which is usually dominated by phonons. Such relationships between thermal and electrical conductions are well documented in literature for carbon nanotubes/polymer-matrix composite [25–27]. In composite materials, due to the addition of carbon fillers, the enhancement in thermal conductivity is lower than the enhancement in electrical conductivity. This difference is attributed to the participation in thermal conduction of both phases matrix and fillers, whereas only fillers participate in the electrical conduction [26]. The interfacial and contact thermal resistances between free carbon domains and the SiOxCy matrix completely eliminate the effect of the percolation on the thermal conductivity [26,27].
taken on the same sample. Electrical conductivity increases with rising heat treatment temperature, suggesting that a network of percolated free carbon phase is being formed and undergoes subsequent clustering, as pointed out in earlier work [22]. The electrical conduction mechanism involves charge transport via variable range hopping mechanism at cryogenic temperatures, and tunneling of charge carriers at temperatures ≥300 K, as detailed elsewhere [22]. Fig. 4 also shows the thermal conductivity of SiOxCy measured using photothermal radiometry. The starting polymer has a low thermal conductivity of 0.15 Wm-1K-1 [23], whereas for heat-treated materials, their thermal conductivity linearly increases and reaches 2.24 Wm-1K-1 when PDMS was heated at 1500 °C. This behavior is related to the presence of thermally conductive Cfree phase in the form of turbostratic structure, which is a stack of two or three poly-aromatic layers that are randomly orientated to each other [6], as shown earlier in part I [16]. Similar to glasses and oxides, the thermal conduction mechanism in SiOxCy is due to lattice vibrations [10], which in turn depends on several extrinsic factors including the distribution of the individual phase, thermal resistance at the interferential sites (grain boundaries), and also on the porosity. However, the thermal conductivity of the sample heat-treated at 1500 °C (2.24 Wm-1K-1) is higher than the SiOxCy powders hot pressed at 1600 °C (1.5 Wm-1K-1) [10] and also than those with intentionally added HfO2 or ZrO2 filler-based SiOxCy ceramics [10]. This suggests that the synthesized monolithic SiOxCy contains significant fraction of thermally conductive phase and with relatively high apparent density of 2.24 g cm-3 (buoyancy floatation method) for the sample heat-treated at 1500 °C. Li et al. [24] showed that the thermal conductivity of binary phase SiC and Cfree material, resulting from pyrolysis (1300 °C, 5 h) of polymeric precursors namely polycarbosilane and divinyl benzene (ratio of 80:20) is 30 Wm-1K-1. Such high thermal conductivity has been attributed not only to the fractional ratio of SiC and Cfree but also to the grain size of the individual phases [24]. The authors claimed that by increasing the grain/domain size of the thermal conductive phase, the thermal conduction in SiOxCy based materials can be significantly improved [24]. For the samples heat-treated between 800 °C and 1000 °C, the increase of electrical, thermal conductivity and also the departure of H atoms as detected in mass spectrometry imply the increase of coherence diameter of the polyaromatic rings with the heat treatment temperature. As discussed earlier, indeed, the onset temperature at which the sp2 carbon is formed is situated around 1000 °C [6]. The characteristic peaks of disordered carbon are graphitic (G) band at 1600 cm−1 and disorder induced (D) band at 1365 cm−1 [6]. The G-band corresponds to the presence of sp2 carbon either in chains or rings, whereas the Dband originates from the breathing vibration mode of the polyaromatic rings. The intensity ratio of D to G-band (ID/IG) is a structural parameter that evaluates the in-plane crystallite size of graphitic materials
3.4. Thermoelectric voltage Thermoelectric voltage |ΔV| was measured with the procedure described in the experimental part. Fig. 5 depicts the thermoelectric voltage |ΔV| of SiOxCy when a laser beam hit the sample surface at three different powers (300, 550 and 800 mW) for modulation frequencies ranging from 1.0 Hz to 2.2 Hz. Sample was blackened using a black ink (thick < 1 μm) to avoid as much as possible light reflection on its own surface. The influence of the ink layer on the signal shape is negligible as it is thermally thin when compared to the sample thickness. In other words, the top absorbing layer is non significantly affected by thermal gradient, thus leading to null Seebeck voltage from the ink layer. From the knowledge of thermal parameters (Table 1) and assuming a beam diameter of 1 cm, amplitude of the temperature difference between faces was calculated using equation (2), allowing the estimation of the Seebeck coefficient S. Fig. 5b shows the Seebeck coefficient S calculated from equation (3) after the measurement of the electrical resistance (Rs = 17 MΩ) of the sample. For PDMS heat-treated at 1250 °C, the absolute Seebeck 5
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Table 1 Thermoelectric properties of SiOxCy derived from polydimethylsiloxane polymer obtained at ambient temperature. Heat treatment temperature (°C)
Thickness Ls (μm)
1000 1250 1500
619 981 1069
Thermal diffusivity σs (m2 s-1)
Electrical conductivity σd.c. (S cm-1)
7.389 × 10-7 9.147 × 10-7 8.706 × 10-7
2.56 × 10-5 9.36 × 10-4 0.0218
coefficient |S| was found to be 18 μV K-1, which remains constant with the modulation frequency in the range 1–2.2 Hz. The thermoelectric voltage of SiOxCy as a function of heat treatment temperature is shown in Table 1. The Seebeck coefficient decreases from 390 μV K-1 to 7 μV K-1 with the increase of heat-treatment temperature. Such behavior is expected as the electrical conductivity increases with heat treatment temperature. Higher values of Seebeck coefficient can be achieved for lower values of electrical conductivity. Although the thermal diffusivity (α) remain constant for all the three samples, the thermal conductivities (κ) vary more than a factor of two, meaning that significant changes in the heat capacity (cp). When the density (ρ) of the material is known, the heat capacity cp can be calculated using the relation below,
0.99 1.81 2.27
390 18 7
The commercial polydimethylsiloxane (Sylgard®184) has the ability to form complex shapes not only in its gel form but also after their thermal conversion. A linear shrinkage ratio of 25–30% must be considered while designing the near net-shape ceramic components. X-ray diffraction studies reveal that the amorphous SiO2 phase, typical to all organopolysiloxane precursors, is absent in Sylgard PDMS-derived SiOxCy. Moreover, the onset time of SiC nucleation starts earlier than in materials derived from R (methyl, vinyl, and phenyl) substituted triethoxysilane (R-(Si≡(OEt)3). The increase in electrical and thermal conductivities corroborates with the structural evolution of the free-carbon phase as observed by Raman spectroscopy already shown in previous study. In addition to variable electrical and thermal conductivities, the variety of perspectives on the use of this kind of oxycarbide material is highly promising. We have recently reported on their absorbing properties in the terahertz domain, and in this work the preliminary trials carried out to estimate the magnitude of the Seebeck coefficient seems to indicate the polysiloxane-derived ceramic’s potential as a thermoelectric material. Further development of the instrumental setup and the measurement procedure for determination of the thermoelectric figure-of-merit in a vacuum and at high temperatures (> 800 °C) may open up new opportunities in the design of materials for thermoelectric energy converters.
The densities of the heat treated PDMS samples were evaluated by helium pycnometry, (Fig. 6). It can be observed that the density (g/ cm3) increases with increasing heat treatment temperature. This behaviour is in agreement with results reported by Gallis et al. [30] on the density of silicon oxycarbide thin films via X-ray photospectrometry studies. The increasing tendency of density is attributed to the formation of SiC4 units at higher temperatures, in comparison with lower densities of SiOxCy and SiO2. Likewise, the heat capacity Cp (J.Kg/K) of the pyrolysed PDMS also increases with temperature. The thermoelectric figure-of-merit (zT) is calculated
S 2σd . c .Ta κs
Absolute Seebeck coefficient |S| (μV K-1)
4. Conclusion
κ ρ cp
zT =
Thermal conductivity κs (WmK )
temperature, purge gas atmosphere, and porosity. Organic polymerbased composites, for example, carbon nanotubes in polyaniline matrix show zT value of 2.0 × 10-4 at room temperature [12]. Although the zT of thermoelectric materials would be improved at high temperatures (> 800 °C) [28] where the Carnot efficiency will be maximum, organic matrix composites and the most other materials (Bi-Te) are not stable at such temperature as it undergoes severe thermal degradation. For this purpose, the polymer-derived ceramics seems to a persistence choice due to its high temperature stability. Furthermore, the thermal conductivity of polymer-derived SiOxCy is independent of temperature [10] unlike other bulk crystalline materials which has T3 dependence [29]. It is worthy to mention that PDMS-derived carbon-based nanocomposites have good thermal stability and chemical inertness together with thermoelectric response as demonstrated in Fig. 5. Improving ZT, however, will not be a trivial task given that the maximization of one parameter is typically achieved at the cost of the others two; the electrical conductivity and the thermal conductivity (k) of the material [30]. The nanostructure of the solid materials makes possible the decoupling of electrical and thermal conductivities [31]. The thermoelectric performance zT could be improved, by adjusting the chemical composition of starting polymers, the heat treatment process parameters, and by adding suitable conductive fillers [32].
Fig. 6. Density and heat capacity of PDMS derived ceramics heat treated at various temperatures.
α=
1 -1
(4)
Table 1 summarizes the thermoelectric voltage, electrical and thermal conductivities of SiOxCy as a function of heat treatment temperature. From these values, the figure of merit (zT) can be estimated at room temperature (298 K) using equation (4); this last is rather modest zT ~10-7 – 10-8 and must be improved to several orders of magnitude to reach the prerequisites needed to be used in thermoelectric energy convertors. Especially the electrical conductivity must be largely enhanced without degrading the Seebeck coefficient value. This operation could be done in the future by changing the precursors, by adding metallic nanoparticles and by selecting the proper heat treatment
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PDMS precursors: Part I. Towards conversion of patternable soft gels into hard ceramics, Ceram. Int. 45 (2019) 6255–6262, https://doi.org/10.1016/j.ceramint. 2018.12.106. J. Latournerie, Ceramiques nanocomposites SiCO : Synthese, caracterisation et stabiblite thermique, Unversite de Limoges, 2002. G.D. Sorarù, S. Modena, E. Guadagnino, P. Colombo, J. Egan, C. Pantano, Chemical durability of silicon oxycarbide glasses, J. Am. Ceram. Soc. 85 (2002) 1529–1536, https://doi.org/10.1111/j.1151-2916.2002.tb00308.x. J. Latournerie, P. Dempsey, D. Hourlier-Bahloul, J.-P. Bonnet, Silicon oxycarbide glasses: Part 1—thermochemical stability, J. Am. Ceram. Soc. 89 (2006) 1485–1491, https://doi.org/10.1111/j.1551-2916.2005.00869.x. V. Gualandris, D. Hourlier-Bahloul, F. Babonneau, Structural investigation of the first stages of pyrolysis of Si-C-O preceramic polymers containing Si-H bonds, J. Sol. Gel Sci. Technol. 14 (1999) 39–48, https://doi.org/10.1023/a:1008771729921. K.J. Kim, J.-H. Eom, Y.-W. Kim, W.-S. Seo, Electrical conductivity of dense, bulk silicon-oxycarbide ceramics, J. Eur. Ceram. Soc. 35 (2015) 1355–1360, https://doi. org/10.1016/j.jeurceramsoc.2014.12.007. J. Cordelair, P. Greil, Electrical conductivity measurements as a microprobe for structure transitions in polysiloxane derived Si-O-C ceramics, J. Eur. Ceram. Soc. 20 (2000) 1947–1957, https://doi.org/10.1016/S0955-2219(00)00068-6. J.E. Mark, Polymer Data Handbook, Oxford University Press, Seiten, 1999. Z. Li, Y. Wang, L. An, Control of the thermal conductivity of SiC by modifying the polymer precursor, J. Eur. Ceram. Soc. 37 (2017) 61–67, https://doi.org/10.1016/j. jeurceramsoc.2016.08.023. M.B. Bryning, D.E. Milkie, M.F. Islam, J.M. Kikkawa, A.G. Yodh, Thermal conductivity and interfacial resistance in single-wall carbon nanotube epoxy composites, Appl. Phys. Lett. 87 (2005) 161909, https://doi.org/10.1063/1.2103398. N. Shenogina, S. Shenogin, L. Xue, P. Keblinski, On the lack of thermal percolation in carbon nanotube composites, Appl. Phys. Lett. 87 (2005) 133106, https://doi. org/10.1063/1.2056591. A. Yu, M.E. Itkis, E. Bekyarova, R.C. Haddon, Effect of single-walled carbon nanotube purity on the thermal conductivity of carbon nanotube-based composites, Appl. Phys. Lett. 89 (2006) 133102, https://doi.org/10.1063/1.2357580. J.H. Goldsmid, Introduction to Thermoelectricity, Springer, 2010. A.A. Balandin, Thermal properties of graphene and nanostructured carbon materials, Nat. Mater. 10 (2011) 569–581. A.J. Minnich, M.S. Dresselhaus, Z.F. Ren, G. Chen, Bulk nanostructured thermoelectric materials: current research and future prospects, Energy Environ. Sci. 2 (2009) 466–479, https://doi.org/10.1039/B822664B. R. Kim, S. Datta, M.S. Lundstrom, Influence of dimensionality on thermoelectric device performance, J. Appl. Phys. 105 (2009) 034506, , https://doi.org/10.1063/ 1.3074347. K. Touati, Photothermoélectricité: Modélisation en régime harmonique et caractérisation de matériaux thermoélectriques solides et liquides, Universite Littoral Cote d'Opale, 2016.
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Heat treatment of commercial polydimethylsiloxane PDMS precursors: Part II. Thermal properties of carbon-based ceramic nanocomposites Srisaran Venkatachalama,1, Stéphane Lenfanta, Michael Depriesterb, Abdelhak Hadj Sahraouib, Djamila Hourliera,* a b
IEMN (CNRS, UMR 8520), Université Lille 1, Avenue Poincaré, CS60069, F-59652, Villeneuve d'Ascq Cedex, France Laboratoire de Thermophysique de la Matière Condensée, Université du Littoral Côte d’Opale, 145 Avenue Maurice Schumann, F-59140, Dunkerque Cedex, France
A R T I C LE I N FO
A B S T R A C T
Keywords: Polymer-derived ceramics Polydimethylsiloxane Thermal conductivity Electrical conductivity Thermoelectric response
Organosilicon polymer-derived ceramic (PDC) materials have been extensively studied for their various properties such as mechanical, chemical and electrical properties. One of the least explored properties of PDC is their thermal properties, required for designing a heat energy converter. In this work, a commercially available polydimethylsiloxane elastomer is used as starting polymeric precursor. Pyrolysis of cross-linked gels conducted in inert atmosphere above 1000 °C yielded silicon oxycarbide SiOxCy, as revealed by X-ray diffraction analysis. The complex shapes of the gels are retained even after pyrolysis up to 1500 °C, however a linear shrinkage ratio of 25–30% is observed. Both thermal and electrical conductivity of free carbon based SiOxCy residues are found to rise simultaneously as the heat treatment temperature increased. The sample heat-treated at 1500 °C exhibits the room temperature thermal conductivity of 2.24 Wm-1K-1 and electrical conductivity of 0.021 S cm-1. These properties of PDCs suggest their appealing applications as a functional material in thermoelectric devices.
1. Introduction The silicon oxycarbide issued from pyrolysis of organo-polysiloxanes has seen much promise for its uses in many applications due to their intrinsic properties: chemical stability, high temperature stability in harsh atmospheres, low density, hardness, etc., [1–5]. Very recently, we have demonstrated that polydimethylsiloxane (PDMS) yields materials which exhibit variable absorption properties in terahertz band, making them attractive for bolometer applications [6]. The electromagnetic (EM) wave’s absorption is tightly controlled at desired values by means of control of the polymer-composition, the carbon content, the state of organization of the free carbon phase resulting from hydrocarbon groups present in polymer [7], and finally, shaping-structuring polymeric precursor. Actually, the main advantage of using polymers lies in their ability to form complex small parts in a near net-shape by the casting process. A broadband terahertz microbolometer with responsivity of 0.76 V/W, time constant of 180 ms, and noise equivalent power of 2 nW/√Hz, has been demonstrated as a proofof-concept device operating in atmospheric air [7]. However, the temperature sensing element used in the proof-of concept bolometer was based on the principle of change of an electrical resistance in close
contact with the rear face of the material and requires an external electrical power supply. In this paper we show that silicon oxycarbide issued from pyrolysis of organo-polysiloxanes are able to generate their own electrical voltage when submitted to a temperature gradient. Thus, active bolometers and energy harvesting converters based on silicon oxycarbide thermoelectric materials can be expected. These systems can be molded into any desired shape and allow a versatile use. To assess the capabilities of these materials for these applications, it would be reasonable to investigate other physical properties as their transport coefficients, such as electrical, thermal conductivity, to be of paramount importance. In theory, a thermoelectric material must possess high electrical conductivity (σ), high Seebeck coefficient (α), and low thermal conductivity (κ) as given by the thermoelectric figure-of-merit (zT), α2σ
zT = κ A few studies on the electrical conductivity of PDMS and thermal conductivity of silicon oxycarbide have been found in the literature. Johnson et al. [8] measured the electrical conductivity (d.c.) of the PDMS gel during the course of pyrolysis up to 700 °C, by supplying a very high voltage of 100–1000 V. The electrical conductivity of the polymer (measured in situ) increased linearly up to ~ 1 × 10-11 S cm-1
*
Corresponding author. E-mail address: [email protected] (D. Hourlier). 1 Present Address - Department of Mechatronics, Bannari Amman Institute of Technology, Sathyamangalam, India https://doi.org/10.1016/j.ceramint.2019.07.143 Received 11 January 2019; Received in revised form 23 June 2019; Accepted 12 July 2019 0272-8842/ © 2019 Elsevier Ltd and Techna Group S.r.l. All rights reserved.
Please cite this article as: Srisaran Venkatachalam, et al., Ceramics International, https://doi.org/10.1016/j.ceramint.2019.07.143
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and dense, the samples were fully covered with silver paint on the rear side, whereas a circular electrode was formed on the front side. The samples were dried at 120 °C for 30 min to eliminate the solvents present in the paint. After drying, the diameters of silver electrodes were 2–4 mm, mostly elliptical rather than circular. Semiconductor parameter analyzer (Agilent 4155C) was used to bias the sample in the voltage range 0–50 V. The measured current showed ohmic (linear) behavior. The electrical conductivity was calculated from the slope of the current versus voltage curve, when multiplied to the ratio of sample thickness to area of the front electrode. The mean sample thickness values were obtained using ten test measurements with a Mitutoyo thickness gauge, and the surface area of the electrodes was calculated by measuring the diameter under optical microscope. The thermal conductivity of samples was measured using photothermal radiometry. Briefly, the principle of measurement is explained as follows: a sample absorbs the incident laser beam, and generates heat, thereby acting as blackbody source that emits infrared radiations. Because of the inherent opacity of the sample, the emitted infrared radiation is proportional to the surface temperature of the sample. In the constructed experimental set-up, a laser source (Laser Quantum, Ventus HP532, wavelength = 532 nm) delivers a modulated beam directed towards the sample that is mounted horizontally on a holder. Two parabolic mirrors were used to collect the infrared radiations emitted by the sample, and focused towards a liquid N2-cooled HgCdTe infrared detector (Judson Technologies). The output of the detector signal was then amplified using a Lock-in amplifier (Ametek, Signal Recovery7260). The resulting signal is a frequency dependent complex signal S(f) which is proportional to the temperature variation (Tm) of the sample surface as a function of laser modulation frequency (f), according to equation (1).
at 200 °C in N2 atmosphere. Once the polymer started to degrade, the electrical conductivity varied non-linearly and reached a value of ~ 1.5 × 10-10 S cm-1 at 700 °C [8]. They found a correlation between the thermal decomposition and the electrical conductivity characteristics of the materials. Qiu et al. [9] measured the thermal conductivity of macro-porous SiOxCy issued from a polysiloxane precursor (without specifying the molecular structure) using the free-standing 3ω technique. The thermal conductivity varied between 0.041 Wm-1K-1 and 0.062 Wm-1K-1 for the samples pyrolysed at 650 °C and 1000 °C, respectively. The authors compared their experiemental values with the intrinsic thermal conductivity (2.44 Wm-1K-1) estimated from the law of mixtures using the thermal conductivity values of bulk phases (SiO2 = 1.34 Wm-1K-1, SiC = 1.44 Wm-1K-1 and C = 4.18 Wm-1K-1) and weight percent of each phase (SiO2 = 27.5 wt%, SiC = 35 wt% and C = 37.5 wt%). The large difference in thermal conductivity values was attributed to the presence of porosity in SiOxCy samples. Gurlo et al. [10] investigated thermal conductivity at different temperatures (up to 1200 K) of sintered powders of SiOxCy (pyrolyzed polymethylsilsesquioxane powders) using laser flash analysis. The sample sintered at 1100 °C (warm-pressed at 120 MPa, 180 °C) containing SiO2 = 72.0 vol%, SiC = 9.11 vol %, free carbon = 6.2 vol%, showed a thermal conductivity of 0.5 Wm-1K-1 that remains unchanged whatever the temperature of measurement, whereas the sample with composition SiO2 = 77.5 vol%, SiC = 9.11 vol%, free carbon = 13.35 vol%, obtained from hot-presssing (30 MPa) at 1600 °C exhibited a non-linear response. The values increased from ~ 0.8 Wm-1K-1 to ~ 1.7 Wm-1K-1 when the temperature increased from 150 K to 1200 K. Nevertheless, at room temperature, thermal conductivity increased from 0.5 Wm-1K-1 to 1.2 Wm-1K-1 when the volume % of free-carbon increased from 6.2 vol% to 13.35 vol% [10]. The main difference between the two ceramic samples lies in the content of free carbon. For PDMS Sylgard, we have demonstrated that the free carbon phase embedded into SiOxCy matrix is the most efficient phase for absorption of THz radiation [6]. An understanding of free carbon formation-organisation, and its dependance on thermal and electrical conductivity is therefore valuable for the development of efficient devices. Only scant data on physical properties of PDMS-derived SiCxOy has been reported in the literature [9,10]. The objective of this study is thus to determine electrical and thermal conductivity of carbon-embedded SiOxCy ceramic materials derived from commercial Sylgard®184 polydimethylsiloxane PDMS. These data will allow us to optimize pyrolysis conditions and also to adapt the chemical architecture of precursors to obtain the desired properties for specific applications. An attempt has been made to evaluate the thermoelectric properties of thermally converted PDMS.
S (f ) = K (f ) Tm =
K (f ) ηIo 1 + Rsbexp(−2σs Ls ) × 2κs σs 1 − Rsb exp(−2σs Ls )
(1)
Equation (1) was derived under the assumptions that the sample is opaque to both the excitation laser and the emitted infrared radiation, and using a one-dimensional heat propagation model across the different planes of the sample. The factor K(f) is an instrumental constant that depends on geometrical factors, spectral average emissivity of the sample surface, and Stephan-Boltzmann constant. This constant K(f) can be circumvent by a normalization procedure, which involves recording in two different configurations: one where the sample is suspended in the air, and the other where the sample stands on a liquid substrate. f is the laser modulation frequency, η is the absorbed fraction of the incident intensity Io, κs is the thermal conductivity of the sample Ls is the thickness of the sample σs is a factor related to the thermal diffusivity (αs) of the sample
2. Experimental The starting polymers were prepared using commercially available Sylgard®184 elastomer. The curing agent and base were mixed in a ratio of 1:10, part by volume, and stirred thoroughly in a clean beaker. The air bubbles and solvents were degassed in a vacuum chamber. The sol was transferred to a polystyrene dish, and allowed to cross-link at room temperature for minimum of 60 h. The gels in square sizes 1.5 × 1.5 cm2 were pyrolyzed in an alumina horizontal tube furnace at a heating rate of 2 °C/min in a flow of 50 ml/min of Argon (99.999% purity) at different temperatures i.e., 800 °C, 1000 °C, 1300 °C, and 1500 °C, as separate pieces. X-ray diffraction (XRD) patterns were recorded on a Bruker’s D8 Advance spectrometer between 10° and 90° with an increment of 0.05° and acquisition time of 2 s. For measurements of electrical conductivity, thermal conductivity and Seebeck coefficients, samples heated at various temperatures, for a short dwell time of 5 min, were cooled down to room temperature. A new specimen of the sample is used for each test measurement: thermal conductivity and electrical conductivity. The electrical conductivity (d.c.) of pyrolyzed samples was measured along their thickness in a two probe configuration. As the material is compact
σs = (1 + j )
πf αs
, j = √-1
Rsb is the reflection coefficient of the thermal waves between the sample s and the substrate b (air or water). Moreover, Rsb is related ⎛ es − eb ⎞ ⎝ eb eb ⎠ ⎜
to the thermal effusivity (e) as, Rsb =
⎟
⎛ es + eb ⎞ ⎝ eb eb ⎠ ⎜
.
⎟
For samples of thickness > 220 μm, the modulation frequency varies between 1 and 100 Hz. The thermal diffusivity αs and effusivity es of the samples can be extracted by fitting the normalized complex voltage to equation (1) using the Levenberg-Marquardt fit routine [11]. Thermal conductivity κs is related to thermal diffusivity and the thermal effusivity as κs = es√αs. The Seebeck coefficient was estimated via the photothermoelectric technique (PTE). Although it was initially designed for the thermal characterization of materials [12], this approach was successfully 2
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applied in the past to the determination of thermoelectric power with the assistance of IR thermography [13]. Both sides of the samples were sputtered with gold electrodes of diameter 0.8 cm, and then the electrodes were attached to low resistance cables. The cables were connected to a lock-in amplifier (EG&G 7260, 10 MΩ input impedance) allowing thermoelectric voltage |ΔV| measurement. The front side of the sample was illuminated using a frequency-modulated laser beam. As the laser radiation is being absorbed by the material, heat induced in the sample generates a temperature gradient along the thickness of the material. Here, instead of being measured as in Ref. [13], the temperature difference was estimated. When the sample is suspended in air, temperature difference ΔT between both sides of an opaque sample was calculated from the following relationship [12], assuming a fully absorbing medium:
ΔT =
1 + exp(−2σs Ls ) − 2exp(−σs Ls ) I0 × 2ks σs 1 − exp(−2σs Ls )
Fig. 1. Scanning electron microscope images of SiOxCy issued from pyrolysis of Sylgard ®184 PDMS at 1500 °C in flowing argon atmospheres (a) pyramidal structure (b) inverted pyramids. Influence of dwell time on the surface morphology (c) no dwell time (d) dwell time 15 min at 1500 °C.
(2)
Thermal diffusivity αs of the sample was determined from the photothermal radiometry experiment and the laser incident intensity I0 was given by the supplier data sheet. Seebeck coefficient is then obtained from the ratio of the amplitude of the photothermoelectric voltage |ΔV| to the estimated temperature difference modulus |ΔT|. A corrective factor was introduced due to the relative important value of the sample electrical resistance (Rs) compared to the internal input impedance of the lock-in amplifier (Rlock-in). Gold contribution to the Seebeck coefficient was neglected in the calculation. Absolute Seebeck coefficient value is then given by the following relation (3).
S =
|ΔV | Rlock − in × ΔT Rs + Rlock − in
(Fig. 1b). The formation of such structure is a consequence of carbothermal reduction reactions between the silicon oxycarbide and free carbon phases, leading to crystalline β-SiC solid phase and CO gas as byproduct that has been detected by mass spectrometry at m/z 28. The distribution of SiC and free carbon phases were identified by Hourlier et al. [15] using Raman mapping for the ceramic residue issued from pyrolysis of polyhydridomethylsiloxane (CH3-Si(H)-O-)n at 1450 °C for a dwell time of 150 min. During such a long dwell time, low crystalline SiC were homogeneously distributed in the matrix, whereas high crystalline SiC were observed on selected sites where the free carbon phase is totally consumed [15]. Depending on the dwell time, the content of high crystalline SiC and carbon can be controlled. In PDMS-derived gels, the presence of crystalline β-SiC phases were also identified using X-ray diffractograms (Fig. 2), which corroborates with the results obtained by Raman spectroscopy. Fig. 2 shows X-ray diffraction patterns of the ceramic residues as a function of heat-treatment temperature. The heat-treated sample at 1000 °C is totally amorphous. Further pyrolysis temperature increase triggers the formation of crystalline nuclei. For samples heated between 1300 °C and 1500 °C, the pattern consists of two broad signals centered at 2θ ≈ 60° and 72°. At 1500 °C and holding time of 15 min, the signals sharpen and new peaks appear due to the crystallization of β-SiC. This result well correlates with the results obtained by Raman and TG/MS analysis. The SiC formation is related to the well-known carbothermal reaction, in which either free carbon phase reacts with SiO issued from the decomposition of oxycarbide network SiOxCy, or free carbon phase and SiO2 phase present in
(3)
One of the advantages of the PTE technique is the small temperature difference (lower than 1 °C) between faces, allowing a good resolution. Nevertheless, the indirect estimation of the temperature difference from the received light flux is still a drawback in comparison with what was done in Ref. [13] This approach is a preliminary trial to estimate the potentialities of SiOxCy materials as thermoelectric elements. Other experiments will be held in the future with an especially dedicated setup to improve the accuracy on the measured Seebeck coefficient values. 3. Results and discussions 3.1. Shaping of ceramics and microstructure The main advantage of Sylgard PDMS is its ability to form desired shape needed for specific applications. As an example, scanning electron microscope (SEM) images of pyramidal net-shaped ceramics obtained up on pyrolysis of Sylgard PDMS at 1500 °C, is shown in Fig. 1 a and b. For terahertz/microwave absorbers, pyramidal structures either in normal and inverse configurations are necessary for the maximum trapping of light. Design and fabrication of such structures facilitate the incident beam to encounter several time on the pyramidal surfaces, so that the reflected beam from one pyramid is absorbed on the surface of neighboring pyramids [14]. Pyramidal structures were obtained by casting the Sylgard mixed precursors in appropriate mold, followed by demolding the cross-linked gel and subsequent pyrolysis. The molds having pyramidal heights ranging from 236 μm to 1 mm were used. Due to the mass loss of the gels during pyrolysis, the height of the pyramids were reduced from 850 μm to 642 μm and from 1 mm to 750 μm. Linear shrinkage ratios were found to vary between 25 and 30%. The surface of ceramic obtained at 1500 °C is initially smooth (Fig. 1c), but holding the sample only 15 min (dwell time) at the same temperature produced considerable changes on the surface. The sample surface exhibited different pore sizes ranging from 30 nm to 100 nm
Fig. 2. Influence of pyrolyzed conditions (temperature and dwell time) on XRD patterns of PDMS-derived ceramic residues. 3
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Fig. 3. Normalized amplitude and phase of thermal radiation collected from the sample surfaces.
the material. In either case the CO is eliminated at high temperature as shown by mass spectrometry in Part I [16]. It is rather surprising to see that the PDMS-derived ceramic does not exhibit any broad signal at 2θ between 22 and 26° characteristic of amorphous SiO2. Almost organosiloxane family [17,18] including MTES (CH3-SiO1.5)n, VTES (C2H3-SiO1.5)n, PTES (C6H5-SiO1.5)n, copolymers of DH (CH3-Si(H)-O-)n and TH (H-SiO1.5)n, methyl-substituted (SR350) [18] or methyl/phenyl-substituted (SR355) [18] polysiloxane resins, heated below 1500 °C indicate the presence of amorphous SiO2. The XRD pattern of the PDMS Sylgard heated at 1500 °C for 15 min resembles to that MTES, VTES and PTES heated at the same temperature but for a longer time (24 h) [19]. 3.2. Photothermal radiometry Thermal diffusivity and effusivity of the samples were extracted by fitting both amplitude and phase components received from the measured data. Fig. 3 depicts the normalized amplitude and phase signals obtained from PDMS samples pyrolyzed at various temperatures. For heat-treated sample at 800 °C, because the normalized amplitude is nearly flat, one can evaluate thermal parameters only from the normalized phase that contains the required information. A data fitting with argument of equation (1) was used to determine thermal parameters (thermal diffusivity, effusivity, and conductivity) values. At other temperatures, a combined data fitting of the normalized amplitude and phase according to equation (1) was applied. In normalized phase signal, the minimum crossing the zero line could be identified. As the treatment temperature increases, the minimum shifts from lower frequency to higher frequency.
Fig. 4. The room temperature electrical conductivity and thermal conductivity of SiOxCy ceramics issued from pyrolysis of Sylgard®184 polydimethylsiloxane.
resulting from thermal conversion of hybrid organic-inorganic precursor PDMS have been described as a composite made up of an intimate mixture of free carbon phase embedded into silicon oxycarbide SiOxCy network [20]. The starting Sylgard PDMS gel is an electrical insulator, having an ambient temperature electrical conductivity of 3.44 × 10-15 S cm-1 [8]. In contrast, for pyrolyzed materials, there is a progressive transition from insulating state to conductor. The electrical conductivity increases up to 0.021 S cm-1 for heat treated material at 1500 °C. Kim et al. [21] showed an improvement in electrical conductivities of SiOxCy system by sintering process of powders synthesized via thermal conversion of polysiloxane (YR3370, Momentive Performance Materials Japan Inc.). The electrical conductivity of sintered SiOxCy – based materials varies from 0.43 S cm-1 to 7.1 S cm-1 when the powders were sintered between 1450 °C and 1650 °C, respectively. The electrical conductivity values showed in Fig. 4 present statistical average of at least 5 measurements
3.3. Electrical and thermal conductivity Fig. 4 depicts the electrical conductivity and thermal conductivity, measured at room temperature 298 K, of residue materials produced at different pyrolysis temperatures from PDMS. These pyrolyzed materials 4
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Fig. 5. Thermoelectric response of PDMS heat treated at 1250 °C as a function of modulation frequency (a) the amplitude of the photothermoelectric voltage (b) apparent Seebeck coefficient.
whereas in amorphous carbon, it is proportional to the number of ordered polyaromatic rings. For the samples pyrolyzed between 1000 °C and 1250 °C, a non-linear increase in the ID/IG ratio with increasing pyrolysis temperature can be explained by the so-called carbonization process in which rearrangement of carbon units, and also dehydrogenation of polyaromatic rings take place simultaneously. Above 1250 °C, the Cfree phase embedded within amorphous SiOxCy is responsible for the conduction. The electrical conductivity increases with increasing pyrolysis temperature suggests that the electrically conductive of free carbon (Cfree) phase is connected through a percolation network. Indeed, the sharp monotonous increase in several orders of magnitude of electrical conductivity following the power law is indicative of percolated structure [22]. In contrast, the thermal conductivity follows a simple linear relationship with the pyrolysis temperature and no signature is observed for percolation in the thermal conduction path. The electrical percolation has no effect on thermal transport, which is usually dominated by phonons. Such relationships between thermal and electrical conductions are well documented in literature for carbon nanotubes/polymer-matrix composite [25–27]. In composite materials, due to the addition of carbon fillers, the enhancement in thermal conductivity is lower than the enhancement in electrical conductivity. This difference is attributed to the participation in thermal conduction of both phases matrix and fillers, whereas only fillers participate in the electrical conduction [26]. The interfacial and contact thermal resistances between free carbon domains and the SiOxCy matrix completely eliminate the effect of the percolation on the thermal conductivity [26,27].
taken on the same sample. Electrical conductivity increases with rising heat treatment temperature, suggesting that a network of percolated free carbon phase is being formed and undergoes subsequent clustering, as pointed out in earlier work [22]. The electrical conduction mechanism involves charge transport via variable range hopping mechanism at cryogenic temperatures, and tunneling of charge carriers at temperatures ≥300 K, as detailed elsewhere [22]. Fig. 4 also shows the thermal conductivity of SiOxCy measured using photothermal radiometry. The starting polymer has a low thermal conductivity of 0.15 Wm-1K-1 [23], whereas for heat-treated materials, their thermal conductivity linearly increases and reaches 2.24 Wm-1K-1 when PDMS was heated at 1500 °C. This behavior is related to the presence of thermally conductive Cfree phase in the form of turbostratic structure, which is a stack of two or three poly-aromatic layers that are randomly orientated to each other [6], as shown earlier in part I [16]. Similar to glasses and oxides, the thermal conduction mechanism in SiOxCy is due to lattice vibrations [10], which in turn depends on several extrinsic factors including the distribution of the individual phase, thermal resistance at the interferential sites (grain boundaries), and also on the porosity. However, the thermal conductivity of the sample heat-treated at 1500 °C (2.24 Wm-1K-1) is higher than the SiOxCy powders hot pressed at 1600 °C (1.5 Wm-1K-1) [10] and also than those with intentionally added HfO2 or ZrO2 filler-based SiOxCy ceramics [10]. This suggests that the synthesized monolithic SiOxCy contains significant fraction of thermally conductive phase and with relatively high apparent density of 2.24 g cm-3 (buoyancy floatation method) for the sample heat-treated at 1500 °C. Li et al. [24] showed that the thermal conductivity of binary phase SiC and Cfree material, resulting from pyrolysis (1300 °C, 5 h) of polymeric precursors namely polycarbosilane and divinyl benzene (ratio of 80:20) is 30 Wm-1K-1. Such high thermal conductivity has been attributed not only to the fractional ratio of SiC and Cfree but also to the grain size of the individual phases [24]. The authors claimed that by increasing the grain/domain size of the thermal conductive phase, the thermal conduction in SiOxCy based materials can be significantly improved [24]. For the samples heat-treated between 800 °C and 1000 °C, the increase of electrical, thermal conductivity and also the departure of H atoms as detected in mass spectrometry imply the increase of coherence diameter of the polyaromatic rings with the heat treatment temperature. As discussed earlier, indeed, the onset temperature at which the sp2 carbon is formed is situated around 1000 °C [6]. The characteristic peaks of disordered carbon are graphitic (G) band at 1600 cm−1 and disorder induced (D) band at 1365 cm−1 [6]. The G-band corresponds to the presence of sp2 carbon either in chains or rings, whereas the Dband originates from the breathing vibration mode of the polyaromatic rings. The intensity ratio of D to G-band (ID/IG) is a structural parameter that evaluates the in-plane crystallite size of graphitic materials
3.4. Thermoelectric voltage Thermoelectric voltage |ΔV| was measured with the procedure described in the experimental part. Fig. 5 depicts the thermoelectric voltage |ΔV| of SiOxCy when a laser beam hit the sample surface at three different powers (300, 550 and 800 mW) for modulation frequencies ranging from 1.0 Hz to 2.2 Hz. Sample was blackened using a black ink (thick < 1 μm) to avoid as much as possible light reflection on its own surface. The influence of the ink layer on the signal shape is negligible as it is thermally thin when compared to the sample thickness. In other words, the top absorbing layer is non significantly affected by thermal gradient, thus leading to null Seebeck voltage from the ink layer. From the knowledge of thermal parameters (Table 1) and assuming a beam diameter of 1 cm, amplitude of the temperature difference between faces was calculated using equation (2), allowing the estimation of the Seebeck coefficient S. Fig. 5b shows the Seebeck coefficient S calculated from equation (3) after the measurement of the electrical resistance (Rs = 17 MΩ) of the sample. For PDMS heat-treated at 1250 °C, the absolute Seebeck 5
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Table 1 Thermoelectric properties of SiOxCy derived from polydimethylsiloxane polymer obtained at ambient temperature. Heat treatment temperature (°C)
Thickness Ls (μm)
1000 1250 1500
619 981 1069
Thermal diffusivity σs (m2 s-1)
Electrical conductivity σd.c. (S cm-1)
7.389 × 10-7 9.147 × 10-7 8.706 × 10-7
2.56 × 10-5 9.36 × 10-4 0.0218
coefficient |S| was found to be 18 μV K-1, which remains constant with the modulation frequency in the range 1–2.2 Hz. The thermoelectric voltage of SiOxCy as a function of heat treatment temperature is shown in Table 1. The Seebeck coefficient decreases from 390 μV K-1 to 7 μV K-1 with the increase of heat-treatment temperature. Such behavior is expected as the electrical conductivity increases with heat treatment temperature. Higher values of Seebeck coefficient can be achieved for lower values of electrical conductivity. Although the thermal diffusivity (α) remain constant for all the three samples, the thermal conductivities (κ) vary more than a factor of two, meaning that significant changes in the heat capacity (cp). When the density (ρ) of the material is known, the heat capacity cp can be calculated using the relation below,
0.99 1.81 2.27
390 18 7
The commercial polydimethylsiloxane (Sylgard®184) has the ability to form complex shapes not only in its gel form but also after their thermal conversion. A linear shrinkage ratio of 25–30% must be considered while designing the near net-shape ceramic components. X-ray diffraction studies reveal that the amorphous SiO2 phase, typical to all organopolysiloxane precursors, is absent in Sylgard PDMS-derived SiOxCy. Moreover, the onset time of SiC nucleation starts earlier than in materials derived from R (methyl, vinyl, and phenyl) substituted triethoxysilane (R-(Si≡(OEt)3). The increase in electrical and thermal conductivities corroborates with the structural evolution of the free-carbon phase as observed by Raman spectroscopy already shown in previous study. In addition to variable electrical and thermal conductivities, the variety of perspectives on the use of this kind of oxycarbide material is highly promising. We have recently reported on their absorbing properties in the terahertz domain, and in this work the preliminary trials carried out to estimate the magnitude of the Seebeck coefficient seems to indicate the polysiloxane-derived ceramic’s potential as a thermoelectric material. Further development of the instrumental setup and the measurement procedure for determination of the thermoelectric figure-of-merit in a vacuum and at high temperatures (> 800 °C) may open up new opportunities in the design of materials for thermoelectric energy converters.
The densities of the heat treated PDMS samples were evaluated by helium pycnometry, (Fig. 6). It can be observed that the density (g/ cm3) increases with increasing heat treatment temperature. This behaviour is in agreement with results reported by Gallis et al. [30] on the density of silicon oxycarbide thin films via X-ray photospectrometry studies. The increasing tendency of density is attributed to the formation of SiC4 units at higher temperatures, in comparison with lower densities of SiOxCy and SiO2. Likewise, the heat capacity Cp (J.Kg/K) of the pyrolysed PDMS also increases with temperature. The thermoelectric figure-of-merit (zT) is calculated
S 2σd . c .Ta κs
Absolute Seebeck coefficient |S| (μV K-1)
4. Conclusion
κ ρ cp
zT =
Thermal conductivity κs (WmK )
temperature, purge gas atmosphere, and porosity. Organic polymerbased composites, for example, carbon nanotubes in polyaniline matrix show zT value of 2.0 × 10-4 at room temperature [12]. Although the zT of thermoelectric materials would be improved at high temperatures (> 800 °C) [28] where the Carnot efficiency will be maximum, organic matrix composites and the most other materials (Bi-Te) are not stable at such temperature as it undergoes severe thermal degradation. For this purpose, the polymer-derived ceramics seems to a persistence choice due to its high temperature stability. Furthermore, the thermal conductivity of polymer-derived SiOxCy is independent of temperature [10] unlike other bulk crystalline materials which has T3 dependence [29]. It is worthy to mention that PDMS-derived carbon-based nanocomposites have good thermal stability and chemical inertness together with thermoelectric response as demonstrated in Fig. 5. Improving ZT, however, will not be a trivial task given that the maximization of one parameter is typically achieved at the cost of the others two; the electrical conductivity and the thermal conductivity (k) of the material [30]. The nanostructure of the solid materials makes possible the decoupling of electrical and thermal conductivities [31]. The thermoelectric performance zT could be improved, by adjusting the chemical composition of starting polymers, the heat treatment process parameters, and by adding suitable conductive fillers [32].
Fig. 6. Density and heat capacity of PDMS derived ceramics heat treated at various temperatures.
α=
1 -1
(4)
Table 1 summarizes the thermoelectric voltage, electrical and thermal conductivities of SiOxCy as a function of heat treatment temperature. From these values, the figure of merit (zT) can be estimated at room temperature (298 K) using equation (4); this last is rather modest zT ~10-7 – 10-8 and must be improved to several orders of magnitude to reach the prerequisites needed to be used in thermoelectric energy convertors. Especially the electrical conductivity must be largely enhanced without degrading the Seebeck coefficient value. This operation could be done in the future by changing the precursors, by adding metallic nanoparticles and by selecting the proper heat treatment
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