# Effects of Carbonization Temperature on Mechanical and Thermal Insulation Properties of Carbon Aerogel Composites using Phenolic Fibers as Reinforcement

Authors: Longlong Li; Fengqi Liu; Junzong Feng; Yonggang Jiang; Liangjun Li; Jian Feng
Journal: Journal of Nanomaterials
 (2023)

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## Abstract

A series of carbon fiber-reinforced carbon aerogel composites (C/CAs) were prepared via carbonizing phenolic fibers impregnated organic aerogel at temperatures ranging from 1,000 to 1,600°C. Phenolic fiber as soft reinforcement shrinks synchronously with the aerogel matrix during the preparation process, which effectively avoided microcracks and achieved an excellent reinforcement effect. The effects of carbonization temperatures on the mechanical and thermal insulation properties of the C/CAs were investigated via pore structural and morphological analysis, as well as the characterization of mechanical strength and thermal conductivity. The results show that the compressive strength of C/CA is 1.26–2.14 MPa inxy-direction and 0.55–1.20 MPa in z-direction. The obtained bending strength range from 1.92 to 3.62 MPa with the carbonization temperature increase from 1,000 to 1,600°C. The thermal conductivity of C/CA-1000 at 1,800°C is 0.1637 W·m−1·K−1 while that of reached 0.2713 W·m−1·K−1 of C/CA-1600. Further study found that the change of porosity and pore size caused by the closure of micropores at high temperatures should be responsible for performance evolution.


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## Body

## 1. Introduction

Carbon aerogels are defined as a class of nanoporous materials with high porosity, high specific surface area, and controllable microstructure [1]. Initially derived from organic hydrogel, carbon aerogels are obtained by replacing the liquid phase of the hydrogel with the gas and then carbonize at high temperatures. Porous carbon skeleton endows carbon aerogels with high thermal stability [2], low density, and low thermal conductivity [3], which is considered as great candidates for ultrahigh temperature thermal protection systems, especially for aerospace applications [4]. However, the weak connection between the skeleton particles of pure carbon aerogel leads to poor mechanical properties and limits its practical application. A large number of studies has been devoted to strengthening the structure of aerogel to achieve interconnected colloidal 3D carbon network with desirable thermal and mechanical properties, which can be summarized into two design methods: (1) introducing the mechanism of energy absorption by adding reinforcements (e.g., fiber [5, 6] or carbon nanomaterials [7]) and (2) improving the skeleton particle neck through controlling the reaction parameters (e.g., the catalyst ratio [8], the water content [9], and the solution pH [10]) or coating the organic polymer (e.g., isocyanate [11] or epoxy resin [12]). According to these studies, the use of fibers to manufacture fiber/carbon aerogel composite is an economical, simple, and effective strategy. But most of the related studies based on inorganic hard fibers, such as carbon fibers [13], ceramic fibers [14], and glass fibers [15] as reinforcements, may cause the cracking of composites owing to shrinkage mismatch with the aerogel matrix during the carbonization process [16]. In our previous work, oxidized polyacrylonitrile (PAN) fibers are used as reinforcements compounded with organic aerogel, the fibers shrink with aerogel matrix during carbonization, reducing the difference of shrinkage rates between them, and finally transformed into carbon fiber-reinforced carbon aerogel composites (C/CAs) with noncracking and strong mechanical properties [17, 18]. In addition, phenolic fibers also as precursors for carbon fibers have been characterized with respect to their high carbon yield, extraordinary flame corrosion resistance [19], instantaneous ultra-high temperature resistance [20], as well as their outstanding thermal insulation, for example, the thermal conductivity of a 3 mm thick phenolic fiber only 0.026 W·m−1·K−1 [21]. Phenolic fibers have been used in a wide variety of flame-resistant textiles, such as aerocomposites and thermal insulation.In this work, carbon fiber-reinforced carbon aerogel composites (C/CAs) were prepared by carbonizing phenolic fiber-reinforced resorcinol–formaldehyde aerogels. During the carbonization process, stabilized phenolic fibers transform into inorganic carbon fibers, and organic networks also evolved into carbon skeletons, so the effect of carbonization temperatures (Tc) on the structural, mechanical, and thermal properties of C/CAs was systematically investigated.


## 2. Experiments

### 2.1. Preparation of C/CAs

A series of C/CAs are obtained through the carbonization of phenolic fibers-reinforced resorcinol–formaldehyde aerogels at different temperatures. The preparation process of C/CAs is shown in Figure1(a). First of all, organic precursor sol is synthesized by sol–gel polymerization of resorcinol (R) and formaldehyde (F) with deionized water (W) as the solvent and a small amount of sodium carbonate (C) as the catalyst. Referring to previous studies [18], the values of R/C and R/W affect the morphology of carbon aerogels by changing the pH of the solution and then controlling the nucleation and growth of the sol particles. For the purposes of appropriate shrinkage, low density, and narrow pore size, the molar ratio of R : F : C : W is fixed at 1 : 2 : 0.002 : 64. The phenolic fiber felts with the size of 238 × 221 × 13 mm3 and the apparent density of 0.13 g·cm−3 are used as reinforcement. The fiber orientation in the felt is anisotropic, most of the fibers are along with the felt plane (in-plane direction), and only a few needled fibers are along with the perpendicular direction of the felt plane (through-plane direction). After keeping the RF sols at room temperature for 1 day, the phenolic fiber felts were introduced into the RF sol and repeatedly impregnated in vacuum, further gel for 1 day at 50°C, 2 days at 80°C in the water bath. The resulted PF/RF gel is replaced in ethanol for 7 days, then dried under the supercritical condition with CO2. The obtained PF/RF aerogel composites are subsequently carbonized in vacuum at 1,000, 1,200, 1,400, and 1,600°C for 1 hr to formed C/CAs. These C/CAs are referred to as C/CA-1000, -1200, -1400, and -1600, respectively. Figure 1(b) shows the photographs of a sample at different preparation stages, displaying obvious dimensional shrinkage. Figure 1(c) demonstrates the good machinability of C/CA composites, and the large-scale and complex productions can be realized by adjusting the size and shape of the fiber preforms.Figure 1
Preparation process and physical photos: (a) preparation process of the C/CA; (b) photographs of C/CA-1000 at different preparation stages; (c) photograph of a large-scale C/CA.
(a)(b)(c)


### 2.2. Characterization

The linear shrinkage (Ls) is calculated from the differences of the sample dimensions (D) before and after supercritical drying and carbonization stage, according to (1 − Dafter/Dbefore) × 100%. Bulk density (ρ) of samples is obtained by measuring the volume and the mass of the samples. N2 adsorption/desorption measurements (QuadraSorb SI, USA) and mercury porosimeter (AutoPore IV 9500, USA) are performed to obtain pore properties, such as specific surface area, total pore volume, and average pore diameter using a Brunauer–Emmett–Teller method. SEM (Hitachi S4800, Japan) investigated the microstructure and morphology of the samples with 20 kV accelerated voltage. Mechanical properties of C/CAs are conducted on the electronic universal machine (WDW model 100, China) with a crosshead speed of 0.5 mm·min−1 using at least five specimens with dimensions of 10 × 10 × 10 mm3 for compressive strength measurements, 50 × 5 × 4 mm3 for flexural strength measurements. The thermal conductivities (λ) at 25°C are measured by a FOX-200 apparatus based on the steady-state heat flow method, and that between 400 and 1,800°C in argon can be calculated according to Equation (1):(1)λ=αTρCpT,where α(T) is thermal diffusivity at measured temperature tested on Netzsch LFA427 using the small wafer (Φ12.7 mm × 1.8 mm) based on laser flash method, ρ is the sample density, and Cp(T) is the specific heat capacity, which value can be referred from Wiener et al. [4].


## 2.1. Preparation of C/CAs

A series of C/CAs are obtained through the carbonization of phenolic fibers-reinforced resorcinol–formaldehyde aerogels at different temperatures. The preparation process of C/CAs is shown in Figure1(a). First of all, organic precursor sol is synthesized by sol–gel polymerization of resorcinol (R) and formaldehyde (F) with deionized water (W) as the solvent and a small amount of sodium carbonate (C) as the catalyst. Referring to previous studies [18], the values of R/C and R/W affect the morphology of carbon aerogels by changing the pH of the solution and then controlling the nucleation and growth of the sol particles. For the purposes of appropriate shrinkage, low density, and narrow pore size, the molar ratio of R : F : C : W is fixed at 1 : 2 : 0.002 : 64. The phenolic fiber felts with the size of 238 × 221 × 13 mm3 and the apparent density of 0.13 g·cm−3 are used as reinforcement. The fiber orientation in the felt is anisotropic, most of the fibers are along with the felt plane (in-plane direction), and only a few needled fibers are along with the perpendicular direction of the felt plane (through-plane direction). After keeping the RF sols at room temperature for 1 day, the phenolic fiber felts were introduced into the RF sol and repeatedly impregnated in vacuum, further gel for 1 day at 50°C, 2 days at 80°C in the water bath. The resulted PF/RF gel is replaced in ethanol for 7 days, then dried under the supercritical condition with CO2. The obtained PF/RF aerogel composites are subsequently carbonized in vacuum at 1,000, 1,200, 1,400, and 1,600°C for 1 hr to formed C/CAs. These C/CAs are referred to as C/CA-1000, -1200, -1400, and -1600, respectively. Figure 1(b) shows the photographs of a sample at different preparation stages, displaying obvious dimensional shrinkage. Figure 1(c) demonstrates the good machinability of C/CA composites, and the large-scale and complex productions can be realized by adjusting the size and shape of the fiber preforms.Figure 1
Preparation process and physical photos: (a) preparation process of the C/CA; (b) photographs of C/CA-1000 at different preparation stages; (c) photograph of a large-scale C/CA.
(a)(b)(c)


## 2.2. Characterization

The linear shrinkage (Ls) is calculated from the differences of the sample dimensions (D) before and after supercritical drying and carbonization stage, according to (1 − Dafter/Dbefore) × 100%. Bulk density (ρ) of samples is obtained by measuring the volume and the mass of the samples. N2 adsorption/desorption measurements (QuadraSorb SI, USA) and mercury porosimeter (AutoPore IV 9500, USA) are performed to obtain pore properties, such as specific surface area, total pore volume, and average pore diameter using a Brunauer–Emmett–Teller method. SEM (Hitachi S4800, Japan) investigated the microstructure and morphology of the samples with 20 kV accelerated voltage. Mechanical properties of C/CAs are conducted on the electronic universal machine (WDW model 100, China) with a crosshead speed of 0.5 mm·min−1 using at least five specimens with dimensions of 10 × 10 × 10 mm3 for compressive strength measurements, 50 × 5 × 4 mm3 for flexural strength measurements. The thermal conductivities (λ) at 25°C are measured by a FOX-200 apparatus based on the steady-state heat flow method, and that between 400 and 1,800°C in argon can be calculated according to Equation (1):(1)λ=αTρCpT,where α(T) is thermal diffusivity at measured temperature tested on Netzsch LFA427 using the small wafer (Φ12.7 mm × 1.8 mm) based on laser flash method, ρ is the sample density, and Cp(T) is the specific heat capacity, which value can be referred from Wiener et al. [4].


## 3. Results and Discussion

### 3.1. Physical Properties, Pore Structural, and Morphological Analysis

The variations in size and mass are accompanied by the whole preparation process of C/CAs. Table1 presents the linear shrinkage, density, and carbon yield of the samples before and after supercritical drying and carbonization stages, respectively. It can be seen that the shrinkages in the drying stage are much lower than that in the carbonization stage, but there is no significant change in densities. In addition, these properties are similar after drying attributed to the same pretreatment and seem to be insensitive to the carbonization temperature, because the mass loss happened mainly under 600°C, and there is only a small loss beyond this temperature. Shrinkage of samples occurred in the first stage at the carbonization process below 1,000°C and the second at higher temperatures above 2,000°C [2].Table 1
Basic physical properties of samples at different stages of preparation.
SampleSupercritical dryingCarbonizationShrinkage (%)ρPF/RF (g·cm−3)Shrinkage (%)ρC/CA (g·cm−3)Carbon yield (%)C/CA-10002.220.2519.060.2652.40C/CA-12002.170.2620.470.2551.34C/CA-14002.080.2521.730.2550.17C/CA-16002.230.2521.350.2650.08Pore structural characteristics for C/CAs are shown in Figure2. Parameters such as specific surface area, total pore volume, and average pore diameter determined from these characteristics are shown in Table 2. The adsorption/desorption isotherms (Figure 2(a)) of the C/CAs start with a linear increase and follow by a sharp escalation at high relative pressures (0.9 < P/P0 < 1.0), which indicates that the capillarity is enhanced. According to the classification of International Union of Pure and Applied Chemistry [22], the isotherms belong to the compound types II and IV, indicating the existence of mesopores (2 nm < d < 50 nm) and macropores (d > 50 nm). The amount of adsorption decreased with the increase of Tc. The uptake of C/CA-1600 decreased remarkably to half the amount of C/CA-1000, especially, at the region of capillary condensation. These results indicate that mesopore in carbon aerogel decreased and pore size distribution became broad. Further, the pore size distribution curves of the C/CAs are obtained according to Barret–Joyner–Halenda theory (Figure 2(c)), concluding that the pore size distribution is wide, but still dominated by mesopores. From the analytical data of nitrogen adsorption, we can see that increasing the carbonization temperature results in a decrease in the specific surface area and pore volume, and an increase in the average pore diameter as considered. Compared with C/CA-1000, the BET surface area of C/CA-1600 sample is halved but the average pore diameter is doubled. These variations are caused by the increase in carbonization temperature, and there is a possibility that the fusion of skeleton particles induces micropores to be closed [2]. The mercury intrusion porosimetry (MIP) is adopted to characterize a wider range of pore size, which makes up for the deficiency nitrogen sorption measurements. As shown in Figure 2(b), the mercury intrusion curves of C/CAs show the typical large hysteresis between intrusion and extrusion branch. The filling starts with the largest external surface connected pores as mercury is nonwetting and only with increasing pressure smaller pores are filled. The three-stage increase of cumulative intrusion with pressure indicates that the graded pores exist in the samples. After removing the pressure, no decrease in cumulative intrusion due to a large amount of mercury remains in the pore system. Pore size distribution for samples based on MIP is shown in Figure 2(d). The multipeak structures show that the nanopores from the aerogel matrix and the remaining micropores between the fibers are contained in the sample. Otherwise, the peak position shifts with the increase of carbonization temperature indicating that the pore size increases gradually. The most probable pore size concentered in around 20 and 1,000 nm for the C/CA-1000 sample, and 34 and 2,000 nm for the C/CA-1600 sample. Furthermore, according to the dates from mercury porosimetry, the porosity changes from 60.21% to 82.40% with the increase of carbonization temperature, which is due to the fact that mercury injection cannot be detected for the tiny micropores (<3 nm). The percentage of micropores in C/CA-1000 is relatively large, so the porosity is underestimated. With the increase of pyrolysis temperature, the number of undetectable micropores decreases, and the measured value of porosity tends to the actual value and becomes larger.Figure 2
Pore structure characterization of C/CAs: (a) nitrogen adsorption/desorption isotherms; (b) mercury intrusion curves; (c) pore size distribution derived from nitrogen adsorption/desorption; (d) pore size distribution derived from mercury intrusion.
(a)(b)(c)(d)Table 2
Pore structure parameters of C/CAs.
SampleC/CA-1000C/CA-1200C/CA-1400C/CA-1600Date from nitrogen absorption–desorption isothermsBET surface area (m2·g−1)646.55449.61403.32297.10Total pore volume (cm3·g−1)1.100.800.740.60Average pore diameter (nm)33.9147.7952.7770.43Date from mercury porosimetryTotal pore volume (cm3·g−1)3.063.593.754.40Porosity (%)60.6167.4273.2282.40Figure3 shows the SEM images of C/CAs. It is revealed that the composites do not have any obvious cracks where carbon aerogel is uniformly filled in the fiber network (Figure 3(a)–3(d)). However, the cracks can be found in our previous study [18], in which carbon fibers and mullite fibers as hard reinforcement to strengthen carbon aerogels. In comparison, phenolic fiber as soft reinforcement shrinks synchronously with aerogel matrix during the preparation process, which effectively avoided microcracks and achieved an excellent reinforcement effect. The introduction of carbon aerogel can also prevent the aggregation of fiber bundles, which can effectively weaken the heat transfer caused by fiber overlap and reduces the thermal conductivity of composites [5]. Further observation found that the carbon aerogel matrix maintained a randomly interconnected 3D porous network composed of nanometer-sized strands or spheroidal particles, during which there are randomly distributed with pores of tens to hundreds of nanometers (Figure 3(e)–3(h)). With the increase of carbonization temperature from 1,000 to 1,600°C, the nanoparticles begin to agglomerate and the pore size expands, which is helping to improve mechanical properties because the sturdy carbon skeleton can withstand greater loads but enlarge the heat channel.Figure 3
Morphological characterization of C/CAs: (a–d) SEM images of composite surfaces of C/CA-1000, C/CA-1200, C/CA-1400, and C/CA-1600, respectively. (e–h) SEM images of the corresponding carbon aerogel matrix.
(a)(b)(c)(d)(e)(f)(g)(h)


### 3.2. Effects of Carbonization Temperature on Mechanical Properties

The typical deformation behavior of C/CAs compressing alongxy- and z-directions is shown in Figures 4(a) and 4(b). It can be observed that the compressive response in both directions shows obvious anisotropy. In xy-direction where the fibers arrangement is parallel to the direction of the compression load, the slower stage of stress growth occurs within 2% of the strain attribution to the fact that the load has not yet been transferred to fibers. Subsequently, after increases rapidly to the maximum, the stress decreases only gradually with increasing strain, indicating of fracture toughness. This behavior is characteristic of an elastic deformation formed through elastic bending and rotation of carbon fibers, followed by a multilayer complex split arising from permanent damage [23]. There are some differences in the z-direction where the fibers arrangement is perpendicular to the direction of the compression load, the stress staring with a linear increase with strain up to ∼4%–5%, then the stress–strain curves appear an upward concavity, and the stress rises at a low-rate path to strain exceeding 30%. The whole process is accompanied that the carbon aerogel matrix and fibers inclined to be compressed to absorb more external energy [24]. The compressive strength and modulus of C/CAs in xy- and z-directions are shown in Figures 4(d) and 4(e), with the increase of carbonization temperature from 1,000 to 1,600°C, and the compression properties in both directions increase. Further study found that the better compressive strength belongs to the xy-direction (1.26–2.14 MPa) rather than the z-direction (0.55–1.20 MPa), which is easy to understand that most fibers in xy-direction are distributed with a wide variety of angles to the compressive loading. In contrast, in the z-direction, most fibers are perpendicular to the load, and only a small bundle of needling fibers are used as the load-bearing structure. Similar to the deformation behavior in the xy-direction, the bending strain–stress curves also show a linear increase stage and then follow an erratic, and somewhat horizontal path to strains exceeding 10% (Figure 4(c)). Similarly, with the increase of the carbonization temperature, the bending strength increase from 1.92 to 3.62 MPa, and Young’s modulus incerase from 117 to 198 MPa (Figure 4(f)). One of the reasons for this is the aggregation of nanoparticles at higher carbonization temperature strengthens the ability to resist deformation, and the other is that the appearance of large hexagonal carbon layers closes the pores left by pyrolytic gases H2O, CH4, H2, and CO evolved from the carbon fibers, resulting in a gradual reduction of microdefects with increasing carbonization temperature. Additionally, the structure of the carbon fibers slowly evolved toward ideal graphite which ultimately favors improving the fiber’s mechanical strength [25]. Furthermore, compared with the reported composite system, where nanoparticles [7], carbon fibers [13], or ceramic fibers [15] were used as reinforcements, using phenolic fiber as soft reinforcement can achieve the integration of lightweight and high strength.Figure 4
Mechanical performance of the C/CAs: (a, b) representative compressive stress–strain curves and (d, e) compressive strength and Young’s modulus of C/CAs inxy- and z-directions, respectively; (c) bending stress–strain curves and (f) bending strength and Young’s modulus of C/CAs.
(a)(b)(c)(d)(e)(f)


### 3.3. Effect of Carbonization Temperature on Thermal Insulation Properties

Figures5(a) and 5(b) demonstrate the thermal conductivities at 25°C of pure carbon fibers and the corresponding carbon aerogel composites prepared at different carbonization temperatures, respectively. The results show that both of them increase with carbonization temperature from 1,000 to 1,600°C and lead to higher thermal conductivity when carbon aerogel is filled with fibers, which are attributed to the more solid heat conduction caused by the increase of density. The high-temperature thermal conductivity of C/CAs between 400 and 1,800°C is measured under an argon atmosphere, as shown in Figure 5(c). Obviously, the carbonization temperature has a great impact on the thermal conductivity especially at higher temperatures. For example, the thermal conductivity for the C/CA-1000 sample is 0.1637 W·m−1·K−1 at 1,800°C, and that of the C/CA-1200 sample is 0.2024 W·m−1·K−1 but increases rapidly with the treatment temperature to 0.2713 W·m−1·K−1 for the C/CA-1600 sample. For the carbon aerogel composites, the effective thermal conductivity can be described with a superposition of the solid, gaseous, and radiative thermal conductivity. A negligible contribution from radiative heat transfer to the total thermal conductivity is suppressed by the specific extinction coefficient (more than 1,000 m2·g−1 for carbon aerogels) [3]. The main part of the increased thermal conductivity, therefore, has to be attributed to the changes in the carbon fibers and the solid carbon matrix. The carbon fibers derived from phenolic fibers above 1,500°C show evolution toward ideal graphite occurred, increasing the amount of heat transported through the fibers [3]. The carbon skeleton aggregates and exposes more macropores, creating a more spacious heat channel, which is conductive to heat transfer. Additionally, the contribution of gaseous heat transfer should also be partly responsible for the increase of the total thermal conductivity, which is proportional to the porosity and related to the pore size. The proportion of macropores in the samples at high carbonization temperature increased significantly, while the inhibition of these macropores on gaseous heat conduction was weakened, so the dominant role of gaseous heat transfer was enhanced [4].Figure 5
Thermal insulation performance of C/CAs: (a) thermal conductivity at 25°C of C/CAs; (b) thermal conductivity at 25°C of carbon fibers obtained by carbonized phenolic fibers; (c) high-temperature thermal conductivity between 400 and 1,800°C of C/CAs; (d) comparison of thermal conductivity of different carbon aerogel composites.
(a)(b)(c)(d)The C/CAs present a unique combination of strong mechanical properties and low thermal conductivity, which offer a far superior high-temperature insulation performance to that of current carbon aerogel-based thermal insulation materials [13–15]. As shown in Figure 5(d), carbon aerogel-filled carbon foam composites shows poor heat resistance to 1,800°C with high thermal conductivity even high than 0.675 W·m−1·K−1. It has also been shown that the heat resistance of carbon aerogels can be improved by the addition of carbon fibers, but the widespread use of these has been limited by safety concerns related to cracks. Meanwhile, although the mechanical properties of carbon aerogel composites reinforced with polyacrylonitrile carbon precursor fiber are enhanced, the thermal insulation performance is obviously at a disadvantage. Significantly, the C/CAs reinforced by phenolic fibers can combine excellent thermal insulation and good mechanical properties. Considering that insulting materials can easily encounter squeezing or shaking in their working conditions [27], this performance is very promising from the perspective of practical insulation applications.


## 3.1. Physical Properties, Pore Structural, and Morphological Analysis

The variations in size and mass are accompanied by the whole preparation process of C/CAs. Table1 presents the linear shrinkage, density, and carbon yield of the samples before and after supercritical drying and carbonization stages, respectively. It can be seen that the shrinkages in the drying stage are much lower than that in the carbonization stage, but there is no significant change in densities. In addition, these properties are similar after drying attributed to the same pretreatment and seem to be insensitive to the carbonization temperature, because the mass loss happened mainly under 600°C, and there is only a small loss beyond this temperature. Shrinkage of samples occurred in the first stage at the carbonization process below 1,000°C and the second at higher temperatures above 2,000°C [2].Table 1
Basic physical properties of samples at different stages of preparation.
SampleSupercritical dryingCarbonizationShrinkage (%)ρPF/RF (g·cm−3)Shrinkage (%)ρC/CA (g·cm−3)Carbon yield (%)C/CA-10002.220.2519.060.2652.40C/CA-12002.170.2620.470.2551.34C/CA-14002.080.2521.730.2550.17C/CA-16002.230.2521.350.2650.08Pore structural characteristics for C/CAs are shown in Figure2. Parameters such as specific surface area, total pore volume, and average pore diameter determined from these characteristics are shown in Table 2. The adsorption/desorption isotherms (Figure 2(a)) of the C/CAs start with a linear increase and follow by a sharp escalation at high relative pressures (0.9 < P/P0 < 1.0), which indicates that the capillarity is enhanced. According to the classification of International Union of Pure and Applied Chemistry [22], the isotherms belong to the compound types II and IV, indicating the existence of mesopores (2 nm < d < 50 nm) and macropores (d > 50 nm). The amount of adsorption decreased with the increase of Tc. The uptake of C/CA-1600 decreased remarkably to half the amount of C/CA-1000, especially, at the region of capillary condensation. These results indicate that mesopore in carbon aerogel decreased and pore size distribution became broad. Further, the pore size distribution curves of the C/CAs are obtained according to Barret–Joyner–Halenda theory (Figure 2(c)), concluding that the pore size distribution is wide, but still dominated by mesopores. From the analytical data of nitrogen adsorption, we can see that increasing the carbonization temperature results in a decrease in the specific surface area and pore volume, and an increase in the average pore diameter as considered. Compared with C/CA-1000, the BET surface area of C/CA-1600 sample is halved but the average pore diameter is doubled. These variations are caused by the increase in carbonization temperature, and there is a possibility that the fusion of skeleton particles induces micropores to be closed [2]. The mercury intrusion porosimetry (MIP) is adopted to characterize a wider range of pore size, which makes up for the deficiency nitrogen sorption measurements. As shown in Figure 2(b), the mercury intrusion curves of C/CAs show the typical large hysteresis between intrusion and extrusion branch. The filling starts with the largest external surface connected pores as mercury is nonwetting and only with increasing pressure smaller pores are filled. The three-stage increase of cumulative intrusion with pressure indicates that the graded pores exist in the samples. After removing the pressure, no decrease in cumulative intrusion due to a large amount of mercury remains in the pore system. Pore size distribution for samples based on MIP is shown in Figure 2(d). The multipeak structures show that the nanopores from the aerogel matrix and the remaining micropores between the fibers are contained in the sample. Otherwise, the peak position shifts with the increase of carbonization temperature indicating that the pore size increases gradually. The most probable pore size concentered in around 20 and 1,000 nm for the C/CA-1000 sample, and 34 and 2,000 nm for the C/CA-1600 sample. Furthermore, according to the dates from mercury porosimetry, the porosity changes from 60.21% to 82.40% with the increase of carbonization temperature, which is due to the fact that mercury injection cannot be detected for the tiny micropores (<3 nm). The percentage of micropores in C/CA-1000 is relatively large, so the porosity is underestimated. With the increase of pyrolysis temperature, the number of undetectable micropores decreases, and the measured value of porosity tends to the actual value and becomes larger.Figure 2
Pore structure characterization of C/CAs: (a) nitrogen adsorption/desorption isotherms; (b) mercury intrusion curves; (c) pore size distribution derived from nitrogen adsorption/desorption; (d) pore size distribution derived from mercury intrusion.
(a)(b)(c)(d)Table 2
Pore structure parameters of C/CAs.
SampleC/CA-1000C/CA-1200C/CA-1400C/CA-1600Date from nitrogen absorption–desorption isothermsBET surface area (m2·g−1)646.55449.61403.32297.10Total pore volume (cm3·g−1)1.100.800.740.60Average pore diameter (nm)33.9147.7952.7770.43Date from mercury porosimetryTotal pore volume (cm3·g−1)3.063.593.754.40Porosity (%)60.6167.4273.2282.40Figure3 shows the SEM images of C/CAs. It is revealed that the composites do not have any obvious cracks where carbon aerogel is uniformly filled in the fiber network (Figure 3(a)–3(d)). However, the cracks can be found in our previous study [18], in which carbon fibers and mullite fibers as hard reinforcement to strengthen carbon aerogels. In comparison, phenolic fiber as soft reinforcement shrinks synchronously with aerogel matrix during the preparation process, which effectively avoided microcracks and achieved an excellent reinforcement effect. The introduction of carbon aerogel can also prevent the aggregation of fiber bundles, which can effectively weaken the heat transfer caused by fiber overlap and reduces the thermal conductivity of composites [5]. Further observation found that the carbon aerogel matrix maintained a randomly interconnected 3D porous network composed of nanometer-sized strands or spheroidal particles, during which there are randomly distributed with pores of tens to hundreds of nanometers (Figure 3(e)–3(h)). With the increase of carbonization temperature from 1,000 to 1,600°C, the nanoparticles begin to agglomerate and the pore size expands, which is helping to improve mechanical properties because the sturdy carbon skeleton can withstand greater loads but enlarge the heat channel.Figure 3
Morphological characterization of C/CAs: (a–d) SEM images of composite surfaces of C/CA-1000, C/CA-1200, C/CA-1400, and C/CA-1600, respectively. (e–h) SEM images of the corresponding carbon aerogel matrix.
(a)(b)(c)(d)(e)(f)(g)(h)


## 3.2. Effects of Carbonization Temperature on Mechanical Properties

The typical deformation behavior of C/CAs compressing alongxy- and z-directions is shown in Figures 4(a) and 4(b). It can be observed that the compressive response in both directions shows obvious anisotropy. In xy-direction where the fibers arrangement is parallel to the direction of the compression load, the slower stage of stress growth occurs within 2% of the strain attribution to the fact that the load has not yet been transferred to fibers. Subsequently, after increases rapidly to the maximum, the stress decreases only gradually with increasing strain, indicating of fracture toughness. This behavior is characteristic of an elastic deformation formed through elastic bending and rotation of carbon fibers, followed by a multilayer complex split arising from permanent damage [23]. There are some differences in the z-direction where the fibers arrangement is perpendicular to the direction of the compression load, the stress staring with a linear increase with strain up to ∼4%–5%, then the stress–strain curves appear an upward concavity, and the stress rises at a low-rate path to strain exceeding 30%. The whole process is accompanied that the carbon aerogel matrix and fibers inclined to be compressed to absorb more external energy [24]. The compressive strength and modulus of C/CAs in xy- and z-directions are shown in Figures 4(d) and 4(e), with the increase of carbonization temperature from 1,000 to 1,600°C, and the compression properties in both directions increase. Further study found that the better compressive strength belongs to the xy-direction (1.26–2.14 MPa) rather than the z-direction (0.55–1.20 MPa), which is easy to understand that most fibers in xy-direction are distributed with a wide variety of angles to the compressive loading. In contrast, in the z-direction, most fibers are perpendicular to the load, and only a small bundle of needling fibers are used as the load-bearing structure. Similar to the deformation behavior in the xy-direction, the bending strain–stress curves also show a linear increase stage and then follow an erratic, and somewhat horizontal path to strains exceeding 10% (Figure 4(c)). Similarly, with the increase of the carbonization temperature, the bending strength increase from 1.92 to 3.62 MPa, and Young’s modulus incerase from 117 to 198 MPa (Figure 4(f)). One of the reasons for this is the aggregation of nanoparticles at higher carbonization temperature strengthens the ability to resist deformation, and the other is that the appearance of large hexagonal carbon layers closes the pores left by pyrolytic gases H2O, CH4, H2, and CO evolved from the carbon fibers, resulting in a gradual reduction of microdefects with increasing carbonization temperature. Additionally, the structure of the carbon fibers slowly evolved toward ideal graphite which ultimately favors improving the fiber’s mechanical strength [25]. Furthermore, compared with the reported composite system, where nanoparticles [7], carbon fibers [13], or ceramic fibers [15] were used as reinforcements, using phenolic fiber as soft reinforcement can achieve the integration of lightweight and high strength.Figure 4
Mechanical performance of the C/CAs: (a, b) representative compressive stress–strain curves and (d, e) compressive strength and Young’s modulus of C/CAs inxy- and z-directions, respectively; (c) bending stress–strain curves and (f) bending strength and Young’s modulus of C/CAs.
(a)(b)(c)(d)(e)(f)


## 3.3. Effect of Carbonization Temperature on Thermal Insulation Properties

Figures5(a) and 5(b) demonstrate the thermal conductivities at 25°C of pure carbon fibers and the corresponding carbon aerogel composites prepared at different carbonization temperatures, respectively. The results show that both of them increase with carbonization temperature from 1,000 to 1,600°C and lead to higher thermal conductivity when carbon aerogel is filled with fibers, which are attributed to the more solid heat conduction caused by the increase of density. The high-temperature thermal conductivity of C/CAs between 400 and 1,800°C is measured under an argon atmosphere, as shown in Figure 5(c). Obviously, the carbonization temperature has a great impact on the thermal conductivity especially at higher temperatures. For example, the thermal conductivity for the C/CA-1000 sample is 0.1637 W·m−1·K−1 at 1,800°C, and that of the C/CA-1200 sample is 0.2024 W·m−1·K−1 but increases rapidly with the treatment temperature to 0.2713 W·m−1·K−1 for the C/CA-1600 sample. For the carbon aerogel composites, the effective thermal conductivity can be described with a superposition of the solid, gaseous, and radiative thermal conductivity. A negligible contribution from radiative heat transfer to the total thermal conductivity is suppressed by the specific extinction coefficient (more than 1,000 m2·g−1 for carbon aerogels) [3]. The main part of the increased thermal conductivity, therefore, has to be attributed to the changes in the carbon fibers and the solid carbon matrix. The carbon fibers derived from phenolic fibers above 1,500°C show evolution toward ideal graphite occurred, increasing the amount of heat transported through the fibers [3]. The carbon skeleton aggregates and exposes more macropores, creating a more spacious heat channel, which is conductive to heat transfer. Additionally, the contribution of gaseous heat transfer should also be partly responsible for the increase of the total thermal conductivity, which is proportional to the porosity and related to the pore size. The proportion of macropores in the samples at high carbonization temperature increased significantly, while the inhibition of these macropores on gaseous heat conduction was weakened, so the dominant role of gaseous heat transfer was enhanced [4].Figure 5
Thermal insulation performance of C/CAs: (a) thermal conductivity at 25°C of C/CAs; (b) thermal conductivity at 25°C of carbon fibers obtained by carbonized phenolic fibers; (c) high-temperature thermal conductivity between 400 and 1,800°C of C/CAs; (d) comparison of thermal conductivity of different carbon aerogel composites.
(a)(b)(c)(d)The C/CAs present a unique combination of strong mechanical properties and low thermal conductivity, which offer a far superior high-temperature insulation performance to that of current carbon aerogel-based thermal insulation materials [13–15]. As shown in Figure 5(d), carbon aerogel-filled carbon foam composites shows poor heat resistance to 1,800°C with high thermal conductivity even high than 0.675 W·m−1·K−1. It has also been shown that the heat resistance of carbon aerogels can be improved by the addition of carbon fibers, but the widespread use of these has been limited by safety concerns related to cracks. Meanwhile, although the mechanical properties of carbon aerogel composites reinforced with polyacrylonitrile carbon precursor fiber are enhanced, the thermal insulation performance is obviously at a disadvantage. Significantly, the C/CAs reinforced by phenolic fibers can combine excellent thermal insulation and good mechanical properties. Considering that insulting materials can easily encounter squeezing or shaking in their working conditions [27], this performance is very promising from the perspective of practical insulation applications.


## 4. Conclusion

In summary, the C/CAs were prepared by carbonized phenolic fiber-reinforced organic aerogel composites, and the effect of carbonization temperature on the mechanical and thermal properties was studied in detail. The synchronous shrinkage of phenolic fibers and aerogel matrix gives a good reinforcement effect to the composite system. The mechanical properties of C/CAs inxy- and z-directions show anisotropy, and both of them increase with the increase of carbonization temperature, which is mainly attributed to the carbon skeleton strengthened caused by the aggregation of nanoparticles. As a function of carbonization temperature, the thermal conductivity at 1,800°C almost doubled, compared to that of the samples with a carbonization temperature between 1,000 and 1,600°C. The main cause for the increase can be identified as the graphitization of carbon fibers and pore size enlargement. Nevertheless, the prepared C/CAs combined with the robust mechanical strength and excellent thermal resistance superior to that of other carbon-based thermal insulation composites especially at high temperatures, which is extremely promising from the perspective of ultra-high temperature thermal insulation applications.



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*Source: 1113343-2023-02-12.xml*