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Polymer Derived Silicon Oxycarbide Ceramics as Promising Next Generation Sustainable Thermoelectrics Adhimoolam Bakthavachalam Kousaalya, Xiaoyu Zeng, Mehmet Karakaya, Terry M. Tritt, Srikanth Pilla, and Apparao M. Rao ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.7b17394 • Publication Date (Web): 08 Jan 2018 Downloaded from http://pubs.acs.org on January 8, 2018

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ACS Applied Materials & Interfaces

Polymer Derived Silicon Oxycarbide Ceramics as Promising Next Generation Sustainable Thermoelectrics Adhimoolam Bakthavachalam Kousaalya$,§, Xiaoyu Zeng#, Mehmet Karakaya#, Terry Tritt #,†, Srikanth Pilla$,†,§,*, Apparao M Rao# $

Department of Automotive Engineering, Clemson University, Greenville, SC, USA § #

†

Clemson Composites Center, Greenville, SC, USA

Department of Physics & Astronomy, Clemson University, Clemson, SC, USA

Department of Materials Science and Engineering, Clemson University, Clemson, SC, USA ∗Corresponding author: [email protected]; Tel: 1-864-283-7216

Abstract: We demonstrate the potential of polymer-derived ceramics (PDC) as next-generation sustainable thermoelectrics. Thermoelectric behavior of polymer-derived silicon oxycarbide (SiOC) ceramics (containing hexagonal boron nitride (h-BN) as filler) was studied as a function of measurement temperature. SiOC, sintered at 1300 °C exhibited invariant low thermal conductivity (~ 1.5 W/m-K) over 30-600 °C, coupled with a small increase in both Seebeck coefficient and electrical conductivity, with increase in measurement temperature (30-150 °C). SiOC ceramics containing 1 wt. % h-BN showed the highest Seebeck coefficient (-33 µV/K) for any PDC thus far. Keywords: Polymer-derived ceramic, thermoelectric, h-BN, SiOC, free carbon, Seebeck coefficient Thermoelectrics (TEs) have been the subject of extensive research in physics and material science communities over the past two decades for their potential use as alternative green and clean energy technologies1–4. However, current TEs are woefully deficient due to poor conversion efficiencies, exhibition of their highest ZT values at temperatures close to melting point, presence of toxic and/or expensive elements (e.g. PbTe, Yb in Half-Heusler alloys), and 1 ACS Paragon Plus Environment

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high brittleness2,4. Together, these issues render existing TEs less suitable and environmentally harmful for commercial usage. To address these concerns, polymer-derived ceramics (PDCs) processed via thermolysis of silicon-based polymers have been evaluated in this study. PDCs possess the ability to become the next-generation sustainable thermoelectrics due to their inherent advantages of: (a) Bottomup synthesis that enables fine-tuning of ceramic microstructure and properties; (b) Presence of non-toxic, inexpensive elements; (c) Excellent high-temperature mechanical properties and high thermal stability (> 2000 °C); and (d) Ease of low-temperature processing (< 800 °C)5,6. Further, PDCs can exhibit thermoelectric properties due to the simultaneous presence of an amorphous matrix that lowers thermal conductivity (κ)7–9, and free carbon (i.e. graphitic-like carbon) phase in a percolating network that enhances electrical conductivity (σ)5,10,11. Although many studies have been undertaken to understand both these properties for PDCs7,10–13, there has been little if any published data that fully elucidates the entirety of their TE behavior. The novel contribution of this study lies in determining thermoelectric properties of polymer-derived SiOC ceramics, namely, their thermal conductivity (κ), electrical conductivity (σ), and Seebeck coefficient (S), and evaluate the impact of adding an external filler on these properties. Hexagonal boron nitride (h-BN) was chosen as filler for its ability to enhance S, and the authors analyzed the role of the aforementioned phases – free carbon, amorphous matrix and h-BN – to elucidate thermoelectric behavior of PDCs. Thermo-gravimetric analysis of cross-linked polymethyl-hydrosiloxane (PMHS), filled with varying wt. % of h-BN (see supporting information for experimental details), showed thermal transition from organic (polymer) to inorganic (ceramic) phase over 250 to 550 °C (Figure 1a), and subsequent completion of ceramization at 850 °C. However, the extent of loss of

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organic moieties during polymeric decomposition in the transition region showed significant reduction with increase in h-BN content (Figure 1a). This is due to increase in crosslinking density of PMHS upon the addition of h-BN. Similar increase in ceramic yield was observed (from 72 to 79 %) with the addition of 1 wt. % h-BN, followed by further increase in ceramic yield (to 87 %) upon adding of 5 wt. % h-BN. Analogous behavior of increase in crosslinking density of polysilazane upon addition of h-BN has been reported by David et. al14. They also hypothesized the occurrence of chemical functionalization of h-BN due to the wetting of h-BN sheets by the polymer. Hence, the observed increase in both crosslinking density and ceramic yield in this study could be due to the possibility of chemical functionalization of h-BN sheets, despite their chemical inertness. To analyze and verify if there is any linkage between crosslinking density and functionalization of h-BN sheets, Fourier Transform Infra-red (FTIR) spectra (Figure 1b) were recorded for the crosslinked polymer.

a

b

Figure 1: a) Thermogravimetric analysis and b) Fourier Transform Infra-red spectra of crosslinked Polymethyl-hydrosiloxane with varying wt. % of h-BN Generally, cross-linking of PMHS occurs via dehydrocoupling reaction that encompasses the cleavage of Si-H and C-H bonds, along with the loss of H215. Hence, enhanced crosslinking 3 ACS Paragon Plus Environment


of PMHS upon the addition of h-BN can be corroborated with decrease in FTIR peak intensity at ~ 2169 cm-1 (Si-H vibration) and 2966 cm-1 (C-H vibration) (Figure 1b). In contrast, increase in peak intensity at 1360 cm-1 indicates increase in h-BN concentration. Thus, the aforementioned chemical functionalization of h-BN can be envisaged as the reaction between the catalyst (triethylenediamine – Lewis base) and defective boron atoms (Lewis acid) in h-BN layers, forming a Lewis acid-base adduct16. A broad peak at 1050 cm-1 corresponds to the stretching mode of Si-O-Si bond, while peaks at 758 cm-1 and 1250 cm-1 correspond to Si-C and Si-CH3 bonds respectively. In order to determine the thermoelectric behavior of SiOC ceramic, cross-linked PMHS (with varying wt. % of h-BN) was thermolyzed at 1000 °C under argon atmosphere and subjected to pulsed electric current sintering (PECS) (see supporting information for details on experimental conditions). XRD and Raman spectra of these as-thermolyzed samples indicates their completely amorphous nature (Figures S1 and S2). However, after PECS, SiOC ceramics (C1 and C2) were observed to be X-ray amorphous (Figure 2a), while also indicating the presence of free carbon phase, as exemplified by the broad peak at 26°. In contrast, sintering of all SiOC samples (with and without h-BN) at 1300 °C resulted in the initiation of SiC crystallization (C3, C4 and C5), as shown by the peak at 36°, in addition to the broad peak at 26° for free carbon (Figure 2a). Subsequently, the nature and size of free carbon phase in PECS samples (C1-C5) was ascertained via Raman Spectra. All PECS samples exhibited two distinct signals (Figure 2b): one at ~ 1600 cm-1 corresponding to the G-band (ordered peak), and the other at ~ 1338 cm-1 corresponding to the D-band (disordered peak, corresponding to the A1g breathing mode)17. Intensities of D-band ( ) and G-band ( ) were used to calculate the cluster


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size of free carbon () using the TK (Tunistra and Koenl) equation (Equation 1), where = 2.41 17. =

………….(1)

From Figure 2b, it can be clearly seen that initially, cluster size () increases with increase in sintering temperature (from 1100 to 1200 °C), indicating the dominance of agglomeration of carbon atoms (i.e. nucleation and growth of free carbon clusters). However, subsequent increase in sintering temperature (C3) resulted in reduction of , indicating the opening up of free carbon and its reaction with Si via carbothermal reduction reaction to form SiC crystals18, as confirmed by XRD (Figure 2a). However, on adding h-BN (C4 and C5), free carbon cluster size showed an increase compared to C3. This can be attributed to the effect of retention of carbon by h-BN sheets14,19 due to increased crosslinking density, as observed from TGA and FTIR data (Figure 1a and 1b). Further, increase in h-BN content from 1 wt. % (C4) to 5 wt. % (C5) was accompanied by decrease in free carbon cluster size. This may have been due to the clustering of free carbon atoms over a significantly higher number of retention (nucleating) sites (due to higher h-BN content) in C5. In contrast, clustering of free carbon atoms over much fewer retention sites may have led to the higher free carbon cluster size in C4 compared to C5. Thermal conductivity (κ) of PECS samples (C1-C5) (sample nomenclature: Table S1) exhibited temperature-independent behavior over 25-600 °C (Figure 2c). Samples C1 and C2 showed low κ (< 1 W/m-K), most likely due to their higher extent of open porosity (Figure 2d) and their amorphous nature (Figure 2b). Increase in sintering temperature (to 1300 °C) was observed to increase κ to 1.5 W/m-K due to two reasons. While the first reason is significant reduction in open porosity (Figure 2d), the second – initiation of crystallization of SiC due to carbothermal reduction (Figure 2a) – is highly critical, as crystalline SiC possesses very high κ 5 ACS Paragon Plus Environment


(~ 262 W/m-K)20. While such invariant κ is typical of amorphous materials, some exhibit an exception to this trend21. For instance, while temperature-invariant κ behavior was reported for amorphous polymer-derived SiOC9, increase in κ has been observed for amorphous non-oxide PDCs (such as SiCN and SiBCN) with increase in measurement temperature22. Conversely, κ of SiBCN ceramic, containing crystalline SiO2 (due to oxygen contamination during processing), exhibited temperature-invariant κ till 1000 °C7. This indicates that SiO2 (amorphous/crystalline) significantly influences thermal conduction behavior of PDCs. In this regard, the exact role of SiO2 in influencing the mechanism of thermal conductivity of PDCs is yet to be fully understood. Further, such temperature-invariant κ has been attributed in literature to the scattering of phonons due to tunneling states – a common feature of amorphous materials possessing an open structure23. This explains the temperature-invariant behavior of κ observed in this study, as PDCs are well-known for their open structure predominantly consisting of the amorphous matrix5. However, addition of varying amounts of h-BN (in wt. %) was observed to result in increase in κ of samples (C3-C5), which can be ascribed to the combined effect of reduction in open porosity (due to the addition of filler) and high thermal conductivity of h-BN.


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Figure 2: a) X-ray diffractogram; b) Raman spectra; c) Thermal conductivity; and d) Open porosity of PECS SiOC ceramic (C1 – C5) at different sintering temperature and varying h-BN content (%). Electrical conductivity (σ) and Seebeck coefficient (S) of SiOC (C3) and SiOC/1 wt. % h-BN (C4) are shown in Table 1. Measurements for σ and S were not carried out for samples C1 and C2 due to their higher open porosity (> 15 %), and for Sample C5 due to its higher κ (> 2 W/m-K). Polymer-derived SiOC is well-known to exhibit a wide range of σ from the insulating (10-14 Ω-1cm-1) to the conducting regime (100 Ω-1 cm-1), depending upon material composition, free carbon content, processing method and measurement temperature5,11. However, both C3 and 7 ACS Paragon Plus Environment


C4 samples showed σ values in the semiconducting regime (10-3-10-2 Ω-1cm-1) that increased with measurement temperature. Of the three conduction mechanisms (direct conduction, tunneling percolation and hopping of charge carriers/semi-conducting mechanism)24 used to explain electrical conduction in PDCs, hopping of charge carriers was found to be the predominant mechanism for both systems (C3 and C4), which is in sync with the behavior observed for amorphous semiconductors24. This behavior is confirmed through Figures 3a and 3b that shows linear variation in σ with T-1/4– a behavior observed in previous studies on PDCs13. While some studies have ascribed this behavior to be the result of Mott’s three-dimensional variable range hopping (VRH)12,13 (based on Equation 2), others attribute it to the occurrence of band tail hopping (BTH)25. However, addition of 1 wt. % h-BN was observed to cause reduction in σ, despite the simultaneous increase in free carbon cluster size. This reduction can be attributed to the insulating behavior of h-BN and its role in suppressing conductivity of free carbon through its entrapment between the h-BN layers14. In addition, low measurement temperatures (< 150 °C) may also have resulted in low σ. ∝ −

! ⁄# " … … … … … . (2) !

Where, T0 is characteristic temperature and T is measurement temperature (both are in K). Table 1: Electrical conductivity and Seebeck coefficient of PECS (@ 1300 °C) SiOC (C3) and 1 % h-BN filled (C4) ceramic Temperature (°C) 30 50 100 150

Seebeck Coefficient (µV/K) Electrical Conductivity (Ω-1m-1) 1 % h-BN SiOC 1 % h-BN SiOC -33 -10 2.0 E-03 9.5 E-02 -25 -11 2.3 E-03 1.0 E-01 -22 -11 3.0 E-03 1.2 E-01 -28 -12 4.0 E-03 1.4 E-01


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Seebeck coefficient (S) showed a small increase with increase in measurement temperature for SiOC (C3), while the addition of 1 wt. % h-BN (C4) was observed to cause contradictory behavior (Table 1). The small increase in S of C3 is commonly known as “metallic diffusion thermopower”, and is said to be the result of electron-phonon interactions26. This indicates that the “metallic diffusion thermopower” behavior of SiOC (C3) is possibly due to the interaction of electrons – hopping across the graphitic-like free carbon phase – with phonons (in the amorphous phase). In contrast, addition of h-BN (C4) led to SiOC exhibiting thermopower behavior typical of a semiconductor with a small band gap, initially showing reduction in S (magnitude) with increase in measurement temperature (till 100 °C), followed by increase in S with measurement temperature (Table 1). This increase is due to the high band gap of h-BN (5.95 eV)27, which would have enhanced the band gap of SiOC ceramic upon its addition (albeit in small quantity of 1 wt. %). However, further research must be undertaken to thoroughly elucidate the contribution of amorphous matrix to Seebeck coefficient of PDCs to comprehensively understand their thermoelectric behavior. Interestingly, pure SiOC (C3) showed simultaneous increase in both S and σ with increase in temperature. Previous studies have shown such anomalous behavior for materials that exhibit electrical conduction mainly due to hopping of charge carriers28. Our literature review28 showed that for amorphous materials, the ratio of S to temperature (T) is observed to exhibit a linear relationship with T-1/4, a trend that was also observed in our study, albeit with some error (Figures 3c and 3d). These findings further confirm that electrical conduction in our study was mainly via hopping conduction. Interestingly, room-temperature S obtained for SiOC/1 wt. % hBN (C4) is the highest value (-33 µV/K) reported thus far for any PDC (in terms of magnitude), when compared to the values reported elsewhere for SiCN (-10 µV/K) and SiBCN (2 µV/K)12.



a

b

c

d

Figure 3: a) and b) Electrical conductivity c) and d) Seebeck coefficient of PECS SiOC (C3) and SiOC/1 % h-BN (C4) respectively The negative sign of S observed in this study indicates the presence of n-type charge carriers in this system, which is in stark contrast to the presence of p-type charge carriers as reported by Kim et. al29 in polymer-derived SiOC, based on Hall effect measurements. Such an anomaly between the nature of majority charge carriers (obtained from Hall effect measurements) and S has been observed in earlier studies on amorphous materials, such as boron carbide and amorphous silicon30. This anomaly can be attributed to the localization of charge carriers that contribute to σ, a consequence of the hopping-conduction mechanism. 10 ACS Paragon Plus Environment

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SiOC ceramics are known to contain mixed bonds between silicon, oxygen and carbon that eventually undergo phase separation to form amorphous SiO2, amorphous SiC and graphiticlike free carbon phases5. This study shows that addition of h-BN results in increase in S of SiOC, even as its σ remains in the semiconducting regime. This increase is due to the synergistic effect of presence of h-BN (insulator), graphitic-like free carbon (conducting phase), amorphous SiO2 (insulator), and amorphous SiC (conductor). In contrast, both systems showed poor κ due to the amorphous nature of the aforementioned phases. In summary, addition of h-BN was found to enhance thermal stability of SiOC ceramic due to the possible chemical functionalization of h-BN. Additionally, thermal conductivity of SiOC was much lower < 1 W/m-K when the sample was highly amorphous and porous (> 15 %). However, highly dense SiOC/1 % h BN (C4) samples exhibited conductivity of 1.5 W/m-K along with the highest thermopower ( -33 µV/K) observed thus far for PDCs. By highlighting the combination of low thermal conductivity and simultaneous increase in both electrical conductivity and Seebeck coefficient, the authors have demonstrated the potential of PDCs as thermoelectric materials. Improvement in thermopower of these systems is expected to yield a series of non-toxic, easily processable thermoelectrics in the future. While this study remains the first in determining all three thermoelectric properties of any PDC processed under the same conditions, systematic studies are needed to understand the process-property relationship of these ceramics, particularly variation in parameters such as polymeric precursor, chemical composition of ceramic, thermolysis temperature, and free carbon content. Acknowledgement: The authors wish to acknowledge the financial support provided by the Southern Automotive Women’s Scholarship Fund. The authors are also grateful to Ms. Kimberly Ivey for her assistance in conducting TGA and FTIR measurements, and to Dr. Colin McMillen for XRD results.



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Abstract Graphic


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