Maxime
Balestrat
*a,
Maxime
Cheype
a,
Christel
Gervais
b,
Xavier
Deschanels
c and
Samuel
Bernard
*a
aIRCER, CNRS, Univ. Limoges, Limoges, France. E-mail: maxime.balestrat@unilim.fr; samuel.bernard@unilim.fr
bSorbonne Université, CNRS, Collège de France, Laboratoire Chimie de la Matière Condensée de Paris, 4 Place de Jussieu, 75005 Paris, France
cUniv. Montpellier, ICSM, CEA, CNRS, ENSCM, Marcoule, France
First published on 13th January 2023
TiC/SiC nanocomposites have been prepared and fully characterized at various length scales before exploring their spectral selectivity for their exploitation as solar selective absorbers. The full design process consists of (i) performing the reaction between a polysilazane, namely Durazane® 1800, and tetrakis(dimethylamino)titanium, (ii) shaping the as-synthesized polymer into pellets, (iii) pyrolyzing in flowing argon the raw pellets in order to achieve the conversion of the polymer into a single-phase ceramic and (iv) performing a heat-treatment at higher temperature to generate nanocomposite structures while developing the porosity of the pellets. The evolutive material is characterized at each step of the process via a combination of complementary techniques including FTIR and solid-state NMR spectroscopies, elemental analyses, thermogravimetric analysis (TGA) coupled with mass spectroscopy (MS), X-ray diffraction (XRD), Raman spectroscopy, scanning electron microscopy (SEM), energy-dispersive X-ray spectroscopy (EDX) and transmission electron microscopy (TEM). The optical properties (0.25–25 μm wavelength range) of the materials have been explored as a function of the temperature at which the final material has been isolated (1000–1800 °C). The material prepared at 1800 °C made of TiC nanocrystals in situ immobilized in a highly crystallized SiC matrix with an open porosity of 56% exhibited a boosted room temperature spectral selectivity of 1.91 which has been so far the highest referenced among SiC-based materials.
A first requirement to design SSAs relies on a high absorbance (α > 90%, low reflectance R) in the wavelength range of 0.3–2.5 μm (spectral region of the solar radiation) and a low thermal emittance (ε < 10%, high reflectance R) in the infrared region (2.5–25 μm). This will contribute to an appropriate spectral selectivity (α/ε) which is the parameter to qualify the material composing SSAs.4 A second requirement is its working temperature. Indeed, the current maximum operating temperature of SSAs is limited to 600 °C due to the degradation of its components above this temperature.5 This significantly limits the full exploitation of the photothermal conversion of solar radiation for CSP. Higher working temperatures (typically 900 °C is a common target) will have the consequence of improving the efficiency of the thermal-to-electricity conversion of the power generation system which is capped by the Carnot efficiency; ηc = 1 − Tc/Th in which Th is the working temperature and Tc is the ambient temperature,6 thereby boosting the efficiency of SSAs.
Silicon carbide (SiC) offers unique features for high temperature applications compatible with the operating conditions of further CSP systems.7 It is light weight, and has a high thermal conductivity, an excellent thermal shock resistance performance, a high strength and an oxidation resistance because of its capability to form a passivating oxide layer upon oxidation which provides a thermal stability up to ∼1400 °C.7–9 However, despite its good sunlight absorption properties, SiC displays a high thermal emittance in the infrared region leading to poor optical selectivity.10
Ultra-High Temperature Ceramics (UHTC, e.g. borides, carbides and some nitrides of early transition metals (TM)) can be considered as a class of promising materials in the field of solar thermal power. They possess favorable properties as solar absorbers such as a very high melting point, a good thermo-mechanical properties at high temperatures and, most of them are intrinsically selective.10–17 Among those materials, TiC and TiN exhibit an inherent spectral selectivity which can be used for delivering high-performance SSAs.17–19 However, they tend to oxidize in the targeted operating temperature range.20–22 As most single-phase materials do not naturally have the required behavior (in terms of both thermal stability and spectral selectivity) as SSAs, how can we further enhance the working temperature of SSAs beyond 600 °C?
Engineered composites combining SiC with its unique features as a matrix surrounding the UHTC phase with its inherent spectral selectivity and a lower thermal emittance would display a combination of properties which can be advantageously exploited to increase the operating temperature of SSAs in the next generation of CSP systems. If such materials can be developed as porous components, they could be able to produce the so-called volumetric effect, which means that the irradiated side of the absorber is at a lower temperature than the medium leaving the absorber. Thus, the porous structure would act as a convective heat exchanger where the heat transfer fluid (i.e., air) is imposed to maximize the absorption of the solar radiation by convection heat transfer mode. However, designing UHTC/SiC composites while developing the porosity is highly challenging.
The Polymer-Derived Ceramics (PDCs) route – via in situ growth and materials nano-structuring23,24 represents an effective precursor approach for designing and synthesizing porous SiC – acting as a matrix – containing in situ generated nanoscale UHTC such as nanoTiC25 or nanoTiN.26
Herein, the synthesis and processing of porous nanocomposites consisting of TiC nanocrystals which are in situ generated in a SiC matrix upon modification of preceramic polymers with Ti-containing precursors (tetrakis(dimethylamino)titanium, Ti[N(CH3)2]4) is addressed. As a preceramic polymer, we have selected a commercially available polysilazane, Durazane® 1800, which is well known to produce silico carbonitride (SiCN) ceramics27,28 at intermediate temperatures and form SiC at higher temperatures because of the occurrence of carbothermal reduction reactions releasing nitrogen while developing the porosity of the material. The schematic representation of our approach is described in Fig. 1. The material has been characterized at each step of the process and we propose extensive chemical, structural and textural characterizations of the final nanocomposites. As a proof of concept, their spectral selectivity has been investigated at room temperature.
Fig. 1 Schematic diagram of the general process of designing SiC-based nanocomposite foams from Ti-modified polysilazanes and the provided chemical, structural and optical fingerprints. |
In order to design foams, PSZTi5 is introduced in a sealed stainless-steel bowl in the glove-box to be then ball milled at 200 rpm for 40 min. Then, 1 g of each of the ball-milled PSZTi5 powder is introduced in a heated steel die (AINT-25-H with a 25 mm internal diameter) in the glove-box then uniaxially pressed using a Specac Atlas Autotouch Press-40T at 80 MPa and heated to the set temperature (80 °C) at a heating rate of 5 °C min−1. The selected temperature and pressure are held for 0.5 h. The pressure applied on the samples is then slowly deloaded during cooling down to RT within 1 h. Typical dimensions of the cylindrical green bodies are of a diameter of 25 mm and a thickness of 1.5 mm. The pyrolysis procedure differs from those applied for powders: the pelleted samples are subjected to a cycle of ramping of 1 °C min−1 to 1000 °C, dwelling there for 2 h, and then cooling down to RT at 2 °C min−1. The high temperature treatment procedure matches that applied for powders.
(1) |
(2) |
Fig. 2 (a) FTIR spectrum of PSZ, TDMAT and PSZTi5, (b) 13C CP MAS and (c) 29Si MAS NMR spectra of PSZTi5 and (d) chemical environments identified in PSZTi5. |
It also highlights the typical bands of preceramic polymers containing (i) vinyl silyl groups (CH2 = CH–Si–) via the bands attributed to C–H vibrations at 3058 cm−1, CC stretching at 1594 cm−1, and scissoring of terminal methylene at 1404 cm−1, (ii) Si-methyl groups with two bands located at 2954 and 2902 cm−1 and assigned to the C–H stretching and one more intense band at 1251 cm−1 (δ(Si–CH3)) and (iii) Si–H groups through the strong absorption band at 2123 cm−1. The largest bands in the 1000–500 cm−1 range can be attributed to Si–N stretching in Si–N–Si units at 897 cm−1 and Si–C bond stretching at 787 cm−1 overlapping – in part – the deformation of vinyl and Si–H units.
A marked decrease in the intensity of absorption bands corresponding to N–H groups in –Si–NH–Si– units is observed in the ATR-FTIR spectrum of PSZTi5 while new bands related to TDMAT can be detected in the wavenumber ranges of 2880–2725 cm−1 and 1500–1425 cm−1. They can be attributed to –N(CH3)x (1 ≤ x ≤ 3) units and more particularly to the C–H vibrations and deformation in such units, respectively. These results suggest the successful incorporation of Ti in the PSZ backbone via the reaction between TDMAT and NH groups in PSZ. Additionally, a decrease in the intensity of the band related to Si–H bonds is also observed, while the absorption bands at 3058 and 1594 cm−1 attributed to the vinyl groups strongly decreased in intensity. If considering both units, this could be explained by hydrosilylation reactions between Si–H and –Si–CHCH2. Although the hydrosilylation reaction usually takes place above 120 °C in pure polysilazanes,29 previous works showed that the presence of inorganic catalysts such as transition metals or metal complexes could remarkably lower this reaction temperature.25,30–32 However, taken separately, this can also reflect the reaction of Si–H units with TDMAT while vinyl groups polymerize.
The chemical structure of PSZTi5 was further established by 13C and 29Si solid-state NMR spectroscopy (Fig. 2b and c). The cross-polarization (CP) technique has been used for 13C-NMR experiments to obtain spectra with reasonable acquisition times and signal-to-noise ratios. The solid-state 13C CP MAS NMR spectrum of PSZTi5 (Fig. 2b) exhibits signals generally observed with titanium-modified polysilazane.33 The very broad signal emerging around 3 ppm with shoulders present around 6 and 14 ppm is typical of carbon atoms of aliphatic groups bonded to a silicon atom, i.e., –SiCHx (1 ≤ x ≤ 3).34,35 The resonances at 39 and 45 ppm are assigned to –TiNCH3 groups in N2Si(CH3)–N(CH3)–TiN3 units – thus indicating the occurrence of the reaction between the silicon center of PSZ and –N(CH3)2 groups in TDMAT with the concomitant evolution of methane – and in Si2N–TiN3 units, respectively.25,33 Signals at 124 and 138 ppm are attributed to the carbon of the vinyl groups present in PSZ which tend to demonstrate that hydrosilylation – suggested by FTIR – is uncomplete during the synthesis of the PSZTi5 sample. This means that the decrease of the intensity of the Si–H band in the spectrum of PSZTi5 is most probably ascribed to both the reaction between the silicon center of PSZ and –N(CH3)2 groups in TDMAT, hydrosilylation reaction and vinyl polymerization. The 29Si MAS spectrum of the PSZTi5 sample (Fig. 2c) is composed of a main broad resonance at around −24 ppm that can be assigned to HSiN2C environments, i.e., CH3–Si(H)N2 units as present in PSZ.34,35 In addition, there are two shoulders at around −9 ppm and −35 ppm that can be attributed to SiN2C2 and SiN3C environments, respectively. These results confirm the chemical reaction between TDMAT and PSZ via two main mechanisms as proposed in Fig. 2d.
Our findings are in line with recently proposed mechanisms for the reaction of polycarbosilanes (Si–H) and polysilazanes (Si–H and N–H) with –N(R)2 (R = CH3 and/or CH2CH3) ligands of the TM precursors (with TM = Ti, Zr and Hf).36,37 The proposed reaction mechanisms also agree with our previous report on the mechanism involved in the reaction of polysilazanes with TDMAT25,33 although, in the present paper, the hydrosilylation reaction seems to be more limited most probably because of a too low Ti content in PSZTi5 ([Si1.0Ti0.2C2.2N1.4H6.7]n; oxygen content in the sample is 1.6 wt% and can be therefore neglected). Compared to the chemical formula of PSZ ([Si1.0C1.5N1.1H5.5]n (oxygen content in the sample is 0.4 wt% and can be therefore neglected), the elemental analysis data (Table 1) of PSZTi5 confirms the successful incorporation of Ti in the backbone of PSZ in agreement with the fixed Si:Ti ratio (5) and especially the presence of –Ti(NCH3)x (x = 1, 2) groups according to the increase of the carbon, nitrogen and hydrogen contents. The successful incorporation of –Ti(NCH3)x (x = 1, 2) groups modifies the chemical environment around Si and N as schematically shown in Fig. 2d. It is expected to increase the crosslinking degree of PSZ (x = 1) while forming side groups which are more likely to be released in the early stage of the thermo-chemical conversion of the precursors into nanocomposites (x = 2).
The main difference between the two polymers occurs in the low temperature regime of the TG experiments in which volatilization of low molecular weight species takes place with PSZ. As expected, the formation of –Si–N(R)–Ti– (R = CH3 or Si≡) units in PSZTi5 upon chemical modification of PSZ with TDMAT tends to reduce the segment mobility of oligomeric chains; thus, limiting volatilization of small oligomeric fragments in the low temperature regime of the polymer-to-ceramic conversion. As a consequence, PSZTi5 exhibits a reduced weight loss after decomposition at 1000 °C compared to PSZ. However, this reduced weight loss is somewhat limited most probably due to the presence of pendant groups in the form of –Ti[(NCH3)2]x (x = 1, 2 or 3) units that tend to be easily released as dimethylamine upon pyrolysis.
Indeed, MS (Fig. 4) demonstrated that the decomposition of PSZTi5 is accompanied with the evolution of gaseous H2 (m/z = 2), CH4 and/or NH3 (m/z = 15 and 16), ethane (m/z = 27), methylamine (m/z = 30) and dimethylamine (m/z = 44). The evolution of amine derivatives, i.e., methylamine and dimethylamine, occurred earlier in the process whereas NH3 and/or CH4 evolutions preferentially occurred at intermediate temperatures. Amine release indicates condensation reaction between as-formed –NCH3– bridges in N2Si(CH3)–N(CH3)–TiN3 units and –N(CH3)2 groups in pendant –Ti[(NCH3)2]x (x = 1, 2 or 3) groups with residual N–H units. They represent a continuation of the reaction which occurred during the PSZTi5 synthesis. NH3 evolution corresponds to the transamination reaction while the CH4 evolution originates from the decomposition of Si–CH3. The evolution of H2 indicates the occurrence of dehydrocoupling reactions (low temperature) as well as recombination of radicals (higher temperature) and allows forming the ceramic network labelled PSZTi5_10 after heat-treatment at 1000 °C under argon.
The 29Si MAS NMR spectrum (Fig. 5a) of the PSZTi5_10 sample shows a large signal ranging from 0 to −50 ppm indicating a distribution of SiCxN4−x (0 ≤ x ≤ 4) units. It can nonetheless be noticed that the main peak is centered at −48 ppm corresponding to a predominant SiN4 environment. In parallel, the X-ray diffractogram reveals that PSZTi5_10 is poorly crystallized as shown by the diffuse peaks identified in the XRD pattern (Fig. 5b). They indicate the slow nucleation of face-centered cubic (fcc) TiCxN1−x (0 ≤ x ≤ 1) (ICDD PDF number: 00-042-1488) – which represents the complete solid solution between TiC (ICDD PDF: 00-032-1383) and TiN (ICDD PDF: 00-038-1420) – according to the presence of reflections at 2θ around 36.7° (111), 42.1° (200) and 61.5° (220). Raman spectroscopy (Fig. 5c) allows distinguishing the bands of the defect-induced D (originates from the breathing mode A1g of the sp2 rings) and graphitic G bands (arises from the stretching mode E2g of the sp2 C–C bonds) at ∼1355 cm−1 and ∼1590 cm−1, respectively.38
Fig. 5 (a) 29Si MAS NMR spectra, (b) XRD patterns and (c) Raman spectra of PSZTi5_X samples with X = 10, 14 and 18. |
The presence of free carbon is confirmed through the determination of the chemical formula (Si1.0Ti0.2C1.3N1.4H0.2) of PSZTi5_10 (Table 2) allowing calculating the phase content in the sample: TiN (17 wt%), Si3N4 (57.6 wt%), SiC (5.4 wt%) and sp2 C (20 wt%) phases.
Samples | Ti (wt%) | Si (wt%) | C (wt%) | N (wt%) | O (wt%) | H (wt%) | Chemical formulab |
---|---|---|---|---|---|---|---|
a Measured by EDX. b Because of its low content (<2 wt%), oxygen has been omitted in the chemical formula. | |||||||
PSZ_10 | — | 43.6 | 19.1 | 23.2 | 0.2 | 1.1 | Si1.0C1.0N1.1H0.7 |
PSZTi5_10 | 14.3 | 37.9 | 20.2 | 26.0 | 1.4 | 0.2 | Si1.0Ti0.2C1.3N1.4H0.2 |
PSZTi5_14 | 18.7a | 40.8a | 14.9 | 24.2 | 1.4 | — | Si1.0Ti0.3C0.9N1.2 |
PSZTi5_18 | 15.8a | 55.3a | 27.4 | 1.3 | 0.2 | — | Si1.0Ti0.2C1.1N0.04 |
Analyses showed that the Si:Ti ratio fixed at the polymer level is retained in the ceramic sample prepared at 1000 °C and that the carbon and nitrogen contents increased compared to the reference PSZ_10 sample produced from PSZ under the same pyrolysis conditions. This is clearly due to the introduction of TDMAT in the PSZ network in the early stage of the process which enriches the Si environment with nitrogen and carbon.
Solid State NMR (Fig. 5a) shows a 29Si MAS NMR spectrum displaying SiN4 units present in the PSZTi5_14 sample (Fig. 5a). The X-ray diffractogram of the PSZTi5_14 sample (Fig. 5b) confirmed the results obtained with the PSZTi5_10 samples: the broad peaks at 2θ around 36.4°, 42.3°, 61.4° and 73.4° are attributed to a TiCxN1−x phase (0 ≤ x ≤ 1). Peaks related to this carbonitride phase are at the mid-position between the two references (i.e., TiC and TiN) that could be approximated by Vegard's law39 to a TiC0.5N0.5 phase. The Raman spectrum (Fig. 5c) confirms the presence of free carbon while elemental analyses show that the nitrogen content of the material is relatively stable compared to the PSZTi5_10 sample. This confirmed that no decomposition occurred in the temperature range 1000–1400 °C as shown by TG experiments (Fig. 3).
By assessing the phase micro-/nanostructure of the PSZTi5_14 sample by transmission electron microscopy (TEM, Fig. 6), a sample made of homogeneously dispersed small precipitates (appearing as a dark cluster in the TEM image) embedded in a featureless medium characteristic of an amorphous network is delivered. The Selected Area Electron Diffraction (SAED) pattern (insets in Fig. 6) of these small precipitates identifies poorly defined continuous rings which can be nevertheless indexed as the (111), (200) and (220) planes of the TiCxN1−x phase (0 ≤ x ≤ 1) at 0.251, 0.216 and 0.152 nm, respectively. Therefore, the material prepared at 1400 °C can be described as a nanocomposite made of TiCxN1−x nanocrystals surrounded by a poorly defined matrix; most probably made of SiCxN4−x (0 ≤ x ≤ 4) and sp2 carbon phases as identified by solid-state NMR and Raman spectroscopy, respectively.
Fig. 6 High-resolution TEM image of the surface of the PSZTi5_14 sample with the corresponding selected area electron diffraction pattern as the inset. |
Above 1400 °C, Fig. 3 shows that the weight underwent a significant and continuous change up to 1700 °C. The weight loss rate then decreased in the temperature range 1700–1800 °C. Such changes are highlighted by solid-state NMR spectroscopy. The 29Si MAS NMR spectrum of the PSZTi5_18 sample completely changed (Fig. 5a): it is only composed of SiC4 units; thus, indicating a total decomposition of SiCxN4−x (0 ≤ x ≤ 4) units identified in the PSZTi5_10 and PSZTi5_14 samples. This corroborates the XRD results (Fig. 5b): the XRD pattern of the PSZTi5_18 sample identifies β- and α-SiC phases along with a TiCxN1−x (0 ≤ x ≤ 1) which can be assimilated to a TiC0.7N0.3 phase (Vegard law) based on the peak positions at 2θ = 36.4°, 42.2°, 61.1°, 73.2° and 77.0° that shifted towards TiC peak positions. The Raman spectrum of the PSZTi5_18 sample (Fig. 5c) highlights this decomposition since it identifies the TO bands of TiC and β-SiC phases while the absence of the sp2 carbon bands can be noticed. This indicates the complete carbothermal reduction of the Si–N bonds identified in the PSZTi5_14 sample forming Si–C bonds in the PSZTi5_18 sample. Carbothermal reduction reactions induce the evolution of nitrogen (See Table 2) which is responsible of the weight loss recorded by HT-TGA (Fig. 3): a weight loss of 30% has been measured in the temperature range 1400–1800 °C for the PSZTi5_18 sample which fits relatively well the difference in terms of chemical formula between the PSZTi5_10/PSZTi5_14 samples and the PSZTi5_18 sample. Therefore, the material generated at 1800 °C can be considered as a TiC/SiC composite.
The increase of the heat-treatment temperature to 1800 °C (PSZTi5_18 sample) fully changes the local phase micro-/nanostructure of the materials (Fig. 7). The material appears to be highly crystallized with relatively big crystals. Based on the associated EDX spectra in Fig. 8, these crystals are made of Ti and Si elements (top image) and of Si as identified in the down image. Based on EDX results, we can assume that the SAED patterns (inset in Fig. 8) identify SiC (zinc blende type, space group F3m, a = 0.436 nm) and TiC (rock-salt type, space group Fmm; a = 0.433 nm) which display close crystalline forms and lattice parameters.40 Both SAED patterns are indeed composed of distinct spots that are ascribed to the (111), (200), (220), (311), (222), (400) and (331) planes of the cubic structures of SiC and TiC. In the following section, we investigate the effect of the structural evolution occurring from 1000 to 1800 °C on the optical properties.
Fig. 8 TEM images of the PSZTi5_18 sample with the corresponding EDX spectra selected area electron diffraction patterns identifying (top) TiC and (down) SiC crystals. |
Samples | Powder | Bulk analogs | ||
---|---|---|---|---|
ρ He (g cm−3) | ρ a (g cm−3) | ρ b (g cm−3) | O.P. (%) | |
PSZTi5_10 | 2.38 | 2.40 | 2.59 | 8 |
PSZTi5_14 | 2.88 | 2.67 | 2.97 | 10 |
PSZTi5_18 | 3.67 | 1.78 | 3.69 | 52 |
Table 3 gathers the results obtained from the hydrostatic weight of bulk analogs (ρa apparent density, ρb bulk density and O.P. Open Porosity).
As expected, the bulk density increases with the temperature of preparation according to the extended crystallization of the materials. However, data also show that the carbothermal reduction reactions previously discussed involve a very high level of open porosity in the bulk PSZTi5_18 samples indicating that they form foams upon heat-treatment to this temperature. The effect of the temperature of synthesis of these materials on their selectivity which is defined by the absorptance:emittance ratio41 has been investigated. Absorptance in the solar spectrum wavelength and emittance in the infrared one are both related to hemispherical reflectance by Kirchhoff law. Fig. 9 plots the evolution of the hemispherical reflectance versus the wavelength for the PSZTi_T samples (T = 10, 14 and 18) at room temperature (RT). As expected, the reflectance is very low below λ = 2500 nm and is independent of the temperature of the preparation of the samples. It is close to expectation for an ideal SSA in the solar spectrum. However, the poorly crystallized samples (PSZTi5_10 and PSZTi5_14) have a common non-selective behaviour with a reflectance less than 35% above 2500 nm.
The increase of the temperature of synthesis of these materials at 1800 °C provides enormous benefits. Below 2500 nm wavelength, the reflectance of the PSZTi5_18 sample is decreased to around 10% and from 2500 to 10000 nm, a major increase of the reflectance close to 50% is identified. At 10000 nm, the reflectance evolves through the Christiansen point which is due to absorption mechanisms in the far-infrared range (lattice vibration), causing the material to behave like a blackbody (low reflectance = high emissivity, assuming Kirchhoff's law of radiation).42 Thus, above 10000 nm, the reflectance profile is characteristic of a highly crystallized SiC phase43 in the PSZTi5_18 sample.
The highly crystallized SiC phase in the PSZTi5_18 sample leads to a higher hemispherical reflectance above 2500 nm and therefore a better general selectivity value (Table 4) as calculated from Fig. 9 according to eqn (1) and (2) described in the experimental part. In Table 4, we can clearly observe the contribution of the highly crystalline TiCxN1−x phase which can be assimilated to a TiC phase present in the PSZTi5_18 sample. A RT selectivity of 1.917 has been calculated for the PSZTi5_18 sample.
Samples | α, absorptance | ε, emittance | α/ε, selectivity |
---|---|---|---|
PSZTi5_10 | 0.888 | 0.712 | 1.246 |
PSZTi5_14 | 0.882 | 0.684 | 1.289 |
PSZTi5_18 | 0.893 | 0.466 | 1.917 |
Ideal solar absorbers | 1 | 0 | ∞ |
To our knowledge, this is the highest selectivity reported so far for SiC-based materials.43–45 We should keep in mind that – in addition to the high crystal quality of the nanocomposite – one further parameter – which is not discussed here – can affect the selectivity: the porosity. Although, it is assumed that a porous structure could maximize the absorption of the solar radiation (see introduction), it has been reported that the total hemispherical emissivity can notably increase with the increase of the surface roughness/porosity.46 For instance, an increase of the porosity from 5 to 30 vol% in HfC-based materials led to significant gain of emissivity from 0.4 to 0.55;46 thereby a decrease of the selectivity. As a consequence, this is a property to be further investigated and we can reasonably expect an increase of the RT selectivity if dense materials (as pellets or coatings) can be prepared.
Footnote |
† Electronic supplementary information (ESI) available. See DOI: https://doi.org/10.1039/d2ma00886f |
This journal is © The Royal Society of Chemistry 2023 |