Characterization of Microporous Ceramics Based on Calcium Hexaluminate “In Situ” Produced by Direct Molding

Abstract

Calcium hexaluminate (CA₆) presents a wide application in high-temperature thermal insulation. Despite the high porosity levels achieved, the use of carbonated precursors in its synthesis inevitably produces CO₂ as a by-product. CA₆ was produced by combining different sources of alumina (α-Al₂O₃ and ρ-Al₂O₃) and lime (CaCO₃, Ca(OH)_2, and CaO) in aqueous suspensions that were cast and sintered to evaluate these routes on its physical properties. The products attained after sintering at 1550 and 1600 °C were characterized for crystal phases, real density, particle morphology, uniaxial compressive strength, apparent porosity, and pore size distribution. Part of the samples sintered at 1600 °C was subjected to a thermal shock test and was then evaluated for residual strength under diametrical compression, apparent porosity, pore size distribution, and flexural elastic modulus. The CA₆ samples produced from α-Al₂O₃ presented lower pore fraction and higher mechanical strength and modulus of elasticity. The superior properties of the materials produced with α-alumina were maintained after thermal shock. The acicular geometry of the CA₆ particles is related to their excellent thermal shock resistance and mechanical performance. The results indicated a more environmentally friendly system produced from α-Al₂O₃-CaO for industrial applications of high-temperature thermal insulation resistant to thermal shock damage.

Keywords:
Porous ceramics; Calcium hexaluminate; Direct casting; Alumina sources; Calcium oxide sources

1. Introduction

Calcium hexaluminate (CaAl₁₂O₁₉, CaO.6Al₂O₃, or CA₆) is an alumina-rich compound that contains 92 wt% of Al₂O₃ and 7-8 wt% of CaO and presents high thermal stability and refractoriness (melting point higher than 1800°C)¹-⁶. Due to this, microporous CA₆-based materials are highly stable in reducing atmospheres and present low solubility in alkaline slag, allowing its contact with molten steel⁷-⁹. Besides this, its low thermal conductivity (1-4 W.m^-1∙K^-1 at 1300 °C)⁸-¹⁴ and densification⁶,¹⁵-¹⁹ and thermal shock resistance²,⁵,⁸,⁹,²⁰-²² point out such porous structures as ideal candidates for application as high-temperature thermal insulators¹⁰-¹⁹ and hot-air filters²³-²⁵.

Due to their technological importance, during the past 30 years, several works studied the mechanism of porogenesis and solid-state reactions involved in such structure production employing different processing methods. In general terms, CA₆ can be produced from the combination of several sources of Al₂O_3, for instance, calcined alumina (α-Al₂O₃)²,³,⁵,⁶,¹⁶,¹⁷,¹⁹,²⁶-³², hydratable alumina (ρ-Al₂O₃)³³,³⁴,, and CaO, such as plain calcium oxide (CaO or lime)¹-⁴, calcium carbonate (α-CaCO₃, or calcite¹⁷,¹⁹,²¹,²²,³²,³⁴,³⁵, and μ-CaCO₃, or vaterite¹⁰), calcium hydroxide (Ca(OH)₂)³⁴, calcium chlorite (CaCl₂)³⁶,³⁷, and calcium aluminate cement (CAC)¹⁴,¹⁸,³⁶-³⁹. In these studies, frequently the raw materials were dry-mixed and compacted by uniaxial pressing⁵,⁶,¹⁶,¹⁷,¹⁹,²⁶-³², mixed with an organic compound and extruded³⁵,⁴⁰,⁴¹, or prepared by chemical routes using sol-gel⁵,²³ or gelcasting techniques³⁷,⁴⁰-⁴³ to produce monolith structures⁸,⁹,¹³,¹⁴,³³,³⁴, spherical beads³⁵, or porous aggregates⁸,⁹,¹¹,¹³,⁴⁴-⁴⁶. After the consolidation of these particles, and during the initial heating up to 900-1000ºC, the decomposition of hydroxylated and carbonated compounds generated empty spaces due to the particles’ density mismatch⁶,¹⁷,¹⁹. Alternatively, foaming agents⁴⁷-⁴⁹ and freeze-casting-based techniques⁴⁰ were introduced to increase initial pore content even further. In the 1000-1350ºC temperature range, the solid-state reactions involved in the formation of CA₆ begin at the interface between the Al₂O₃ matrix and CaO domains, where the concentration of Ca²⁺ ions is higher¹,⁶,¹⁰,¹⁷,¹⁹,⁵⁰. Consequently, low-melting point CaO-rich phases (Ca₃Al₂O₆ or C₃A, Ca₁₂Al₁₄O₃₃ or C₁₂A₇, and CaAl₂O₄ or CA) form firstly at these contact points¹⁷,¹⁹. As heating continues, these low-melting phases dissolve the neighboring particles, and as the liquid is progressively enriched in alumina, compounds with higher melting points (CaAl₄O₇ or CA₂ and CA₆) crystallize. Because typical CA₆ crystals are flat hexagonal plates of high aspect ratio, their formation is expansive and the porosity generated during the decomposition of precursors is partially preserved during sintering¹⁰,¹⁶,¹⁹.

Once produced and installed in steelmaking equipment, CA₆ structures must withstand thermomechanical stresses resulting from thermal cycling and temperature gradient between the hot and cold faces⁸,⁹,¹²,¹³,⁵¹. As the material undergoes successive cycles of thermal shock, its structure accumulates damage, which is a determining factor in its useful life. The increase in size and coalescence of cracks in the material drastically reduces its mechanical strength so failure can occur at stress levels much lower than that supported by the undamaged material⁵,⁹,¹⁸,²¹,²²,²⁶,³⁹,⁵²,⁵³. To minimize thermal shock damage, mineralizer agents, such as SiO₂, MgO, B₂O₃, and SiC were added to the compositions to tune the geometry of CA₆ crystals⁵⁴-⁶¹. For instance, whereas SiO₂ additions improved the compacting of the grains and reduced their aspect ratio and inter-particle porosity⁵⁴,⁵⁷,⁵⁸, the presence of MgO and in situ formed magnesium aluminate spinel (MgAl₂O₄) generate elongated crystals and compression forces amongst particles that behaved as toughening mechanisms, preventing the formation of long cracks and making this insulating material resistant to thermal shock damage⁶¹,⁶².

Although the use of calcium carbonate (CaCO₃) as a CaO precursor contributes to increasing the porosity, this compound inevitably produces CO₂ as a by-product (between 590 and 753°C)⁶,¹⁷,¹⁹. Studies show that CaO and Ca(OH)₂ are strategic sources of calcium for the production of CA₆, as they release a lower percentage of toxic gases in the final system. In previous works, combinations of α-Al₂O₃ and different calcium oxide sources have been evaluated mainly in processing routes based on uniaxial pressing¹⁰,¹⁷,¹⁹. However, they can also be formed by liquid-based routes, such as direct casting. To produce “in situ” CA₆, direct casting requires the production of a stable aqueous co-dispersion of Al₂O₃ and CaO sources’ particles after pH adjustment (above 9, to prevent CaCO₃ dissolution and avoid aluminas’ isoelectric point) and proper dispersion (using poly(sodium-ammonium acrylate)-based dispersants)⁶,³⁵. The high flowability of such suspensions facilitates the casting of parts into complex shapes with large volumes, or even direct application on hot linings¹²,¹³. Due to this, such a process would present fast application and energy saving¹²,¹³,⁶³. Because α-Al₂O₃ does not react with water, the addition of a hydraulic binder, such as calcium aluminate cement and hydratable alumina (ρ-Al₂O₃), is necessary to set the suspension before sintering. In both cases, the Al₂O₃-CaO ratio in their particles affects the amount of CA₆ formed¹⁴,³³,³⁴.

In comparison, to other processing routes, direct casting was less explored in literature despite its technological importance. In this context, the objective of this work was to produce microporous in situ CA₆ by direct casting combinations of alumina (α-Al₂O₃ and ρ-Al₂O₃) and lime sources (CaCO₃, Ca(OH)₂ and CaO) and to investigate the combined effect of different raw materials and sintering temperature on physical properties (density, crystal phases, microstructure, compressive strength, porosity, pore size distribution, thermal shock resistance).

2. Materials and Methods

2.1. Raw materials and samples preparation

The raw materials used in this work were: calcined alumina (α-Al₂O₃, 3.9 g/cm³, A1000SG, Almatis, USA), hydraulic alumina (ρ-Al₂O₃, 3.2 g/cm³, Alphabond 300, Almatis, USA), calcium carbonate (CaCO₃, 2.7 g/cm³, Vetec, Brazil), calcium oxide (CaO, 3.5 g/cm³, Vetec, Brazil) and calcium hydroxide (Ca(OH)₂, 2.2 g/cm³, Synth, Brazil). The binder used was calcium aluminate cement (CAC, Secar 71, Kerneos). The dispersant additives used were polyglycol) derivative (Castament FS60, Basf) and sodium polymethacrylate (Darvan 7NS, Vanderbilt Minerals). The raw materials were mixed in stoichiometric proportions to produce calcium hexaluminate (CaAl₁₂O₁₉, CA₆), represented by Equations 1-3 and defined in Table 1:

Table 1
Conditions for preparing aqueous suspensions.

Compositions	Binder (wt%)	Dispersant (wt%)	Solid content (%)
α-Al2O3-CaCO3	25	Polyglycol (0.2)	78.0
α-Al2O3-CaO	25	Sodium polymethacrylate (2.0)	75.8
α-Al2O3-Ca(OH)2	25	Sodium polymethacrylate (1.0)	74.0
ρ-Al2O3-CaCO3	-	Polyglycol (1.4)	66.0
ρ-Al2O3-CaO	-	Polyglycol (1.6)	57.0
ρ-Al2O3-Ca(OH)2	-	Polyglycol (2.8)	54.0

$$6A{l}_2{O}_3+CaC{O}_3\to 6A{l}_2{O}_3+CaO+C{O}_2\to CaA{l}_{12}{O}_{19}$$(1)

$$6A{l}_2{O}_3+CaO\to CaA{l}_{12}{O}_{19}$$(2)

$$\begin{array}{l}6A{l}_2{O}_3+Ca{\left(OH\right)}_2\to 6A{l}_2{O}_3+\\ CaO+H_{2}O\to CaA{l}_{12}{O}_{19}\end{array}$$(3)

The suspensions were prepared and poured under vibration into cylindrical (16 mm diameter × 18 mm height) and prismatic (75 mm length × 12.5 mm width × 12.5 mm height) molds. After casting, samples were subjected to curing processes (50 ^oC, 24 h in a saturated environment), drying (110 ^oC, 24 h), and sintering (1 ^oC/min up to 1000 ^oC, 1 h plateau, 5 ^oC/min. up to 1550 ºC or 1600 ^oC, 2 h plateau, 10 ºC/min cooling rate). Such conditions were employed to maximize the binding effect of hydratable alumina and reduce the likelihood of explosive spalling⁶⁴.

Samples sintered at 1600 ^oC also underwent a thermal shock procedure using the thermal fatigue by multiple cycle method, in which the specimens are subjected to several cycles of sudden heating and cooling, keeping the temperature difference fixed at values higher than that necessary to nucleate cracks (∆T_est ≥ ∆T_c). Each sample was inserted for 15 min into an oven at 1025 ^oC and quickly cooled to room temperature, with a ∆T of approximately 1000 °C. For each sample, 10 cycles were carried out, with a complete cycle being the process of heating and cooling the specimen⁶⁵.

2.2. Samples’ characterization

The products attained after sintering were characterized for crystal phases, real density, particle morphology, uniaxial compressive strength, apparent porosity, and pore size distribution. Part of the samples sintered at 1600 °C after the thermal shock test were then evaluated for residual strength under diametrical compression, apparent porosity, pore size distribution, and flexural elastic modulus. All tests were carried out on cylindrical samples except for the elastic modulus (in rectangular bars).

X-ray diffraction (XRD-6000, Shimadzu; Cu-Kα radiation = 1.54439 Å) was performed to identify the phases formed and the real density was assessed using an Upyc - 1200, v5.04 (Quantachrome) helium pycnometer. The microstructure at the external surface of the sintered samples was investigated using scanning electron microscopy (SEM, EVO MA10, Zeiss, previously covered by thin gold coating).

The compressive strength (σ_U, MPa) was calculated as the average of 5 specimens of the different CA₆ samples under uniaxial compression, according to ABNT NBR 6224:2001, using a universal testing machine (EMIC DL10.000) with a displacement speed of 0.15 mm/min. The rupture stress of each sample was calculated according to Equation 4:

$${\sigma }_U=4P/\pi D^2$$(4)

where P (N) is the load recorded at the fracture of each sample and D is the average diameter of each sample. The diametral compression (σ_D, MPa) test was carried out according to ABNT NBR 7222:2011 using the same equipment and the displacement speed was 0.05 mm/min. The reported value was an average of 5 samples, which residual strength was calculated according to Equation 5:

$${\sigma }_D=2P/\pi hD$$(5)

where h (mm) is the average height of each sample.

The apparent porosity (AP) and volumetric density (VD) were evaluated following the method based on the Archimedes principle (ASTM C 830). The samples (both green and sintered) in triplicate were initially dry-weighed (M_s). The green and sintered samples were then soaked in kerosene (ρ=0.8 g/cm³) or water (ρ=1.0 g/cm³), respectively, under vacuum for 1 h. The immersed (M_i) and wet (M_u) weights were then recorded. The apparent porosity and volumetric density were calculated according to Equations 6 and 7.

$$AP=100\%\times \left[\left(M_U–M_S\right)/\left(M_U–M_I\right)\right]$$(6)

$$VD=100\%\times \left[M_S/\left(M_U–M_I\right)\right]$$(7)

The total porosity (TP) was calculated from the ratio between volumetric density (VD) and real density (RD), according to Equation 8, where VD is the ratio between the mass and the external volume and RD was measured by He picnometry, as previously described.

$$TP=100\%\times \left[1-\left(VD/RD\right)\right]$$(8)

Pore size distribution was evaluated using the mercury porosimetry technique (Autopore IV 9500, Micromeritics, USA). The diameter of the pore into which mercury penetrates relates to the applied pressure through the Washburn equation (Equation 9), valid for cylindrical pores, where D is the pore diameter, P is the applied pressure, γ is the surface tension of the mercury and ϕ is the contact angle (130^o) between the mercury and the sample. The volume of mercury (V) that penetrates the pores is measured directly, as a function of the applied pressure (P). From the curve of the cumulative volume of mercury intruded into the sample after successive pressure increases, the average pore size and pore size distribution can be extracted.

$$D=(-4\gamma cos(\varphi ))/P]$$(9)

The flexural elastic modulus was obtained by the non-destructive impulse excitation method (Sonelastic, ATCP, ASTMC 747, Brazil) using bar-shaped samples.

3. Results and Discussion

The solid-state reaction that forms CA₆ is a double-diffusion mechanism, where Ca²⁺, Al³⁺, and O^2- ions are mutually exchanged between particles. Since the diffusion rates of Ca²⁺ in Al₂O₃ and Al³⁺ in CaO are significantly different, the overall reaction kinetics is slow, and several intermediate compounds may be formed and decomposed before reaching the equilibrium composition¹,⁶,¹⁷,¹⁹. The density of the intermediate compounds (C₃A = 3.15 g.cm^-3, CA = 2.9 g.cm^-3, and C₁₂A₇ = 2.6 g.cm^-3) have lower values than the reactants (Al₂O₃ = 3.9 g.cm^-3, CaO = 3.5 g.cm^-3 and CA₆ = 3.2 g.cm^-3) and the volumetric expansion resulting from its formation tends to separate the grains and make densification slower. Due to this aspect and to the large plate-like morphology of CA₆ crystals, such structures are intrinsically densification-resistant¹⁶-¹⁹.

In the samples produced with α-Al₂O₃, dehydration of the CAC hydrates occurs first (up to 400 °C), causing initial shrinkage¹⁴. As the temperature increases to 1500 °C, there is gradual volume expansion, as the formation of the CaAl₄O₇ (CA₂) phase is completed. Eventually, some shrinkage occurs between 1500 and 1650 °C due to densification. This result suggests that all the samples were well crystallized, and their formation process was completed, as shown in Figure 1

Figure 1
DRX pattern for samples CA₆ obtained by direct molding from compositions of α-Al₂O₃ and ρ-Al₂O₃ with different calcium sources sintered at 1600 °C. Identified phases CaAl₁₂O₁₉ (JCPDS 38-0470) and CaAl₄O₇ (JCPDS 36-0367).

The CA₆ samples produced in this work showed very similar real density values (~3.8 g.cm^‑3), as shown in Table 2

Table 2
Real density of CA₆ samples obtained by direct molding from α-Al₂O₃ and ρ-Al₂O₃ compositions with different calcium sources sintered at 1550°C and 1600°C.

Al2O3 Precursors	CaO Precursors		Ca(OH)2 (g.cm-3)
	CaCO3 (g.cm-3)	CaO (g.cm-3)
α-Al2O3 (1550/1600)	3.79/3.79	3.80/3.80	3.78/3.79
ρ-Al2O3 (1550/1600)	3.78/3.79	3.78/3.79	3.78/3.78

3.1. Uniaxial compression resistance and apparent porosity

The uniaxial compressive strength and apparent porosity results (Figure 2) show that the samples produced with ρ-Al₂O₃ presented higher apparent porosity, both for green or sintered samples and, consequently, a lower uniaxial compressive strength than those prepared with α-Al₂O₃.

Figure 2
(a) Uniaxial compression strength and (b) apparent porosity of samples CA₆ obtained by direct molding from compositions of α-Al₂O₃ and ρ-Al₂O₃ with different calcium sources sintered at 1550 °C and 1600 °C.

Regarding the alumina source, all samples prepared with ρ-Al₂O₃ presented higher apparent porosity (approximately 65%) than those prepared with α-Al₂O₃ (approximately 25%, Figure 2). The ρ-Al₂O₃ composition has some peculiar characteristics that affect its density, such as a significantly lower packing efficiency than α-Al₂O₃, lower density, and the presence of 10-15 wt.% of chemically bound water, generating green bodies with higher porosity than those obtained with α-Al₂O₃. This behavior can be explained by the greater water consumption in the preparation of ρ-Al₂O₃ suspensions (Table 1) as well as the dehydroxylation of the binder because pores formed spontaneously as the mixing water is vaporized and the hydrated phases decompose during calcination. It is worth mentioning that the transformation from ρ-Al₂O₃ to α-Al₂O₃ at higher temperatures involves the formation of other transition phases. For example, above 250 °C ρ-Al₂O₃ transforms firstly into boehmite and other metastable phases as temperature rises. Similar results were observed for other CaO-free ρ-Al₂O₃-bonded castable systems. During such phase transitions, the increase in density (from 3.2 g.cm^-3 to 3.8 g.cm^-3) behaves as a porogenic mechanism, whereas the decomposition of alumina hydrates creates an additional porosity in the matrix. It has been reported that the open porosity of CA₆ produced from ρ-Al₂O₃ compositions is approximately 60%⁶⁶,⁶⁷. Specifically, for the ρ-Al₂O₃-CaCO₃ composition sintered at 1600 °C, the open porosity reached 62.7%.

Some reactions occur in an α-Al₂O₃-CaCO₃ matrix with increasing temperature, leading to the formation of pores and intermediate phases, following a few steps. At room temperature (T = 25 °C), the suspension consists of the main matrix (α-Al₂O₃) and CaCO₃ aggregates that contain closed pores. In the temperature range of 500 - 900 °C, CaCO₃ decomposes, releasing carbon dioxide (CO₂), forming the first open pores¹⁷,¹⁹. Between 1100 and 1362 °C, the C₁₂A₇ and CA phases are formed, along with a small amount of liquid phase at the interface with the alumina matrix due to the melting of the C₁₂A₇ phase. As the temperature rises to 1400 °C, more alumina dissolves into the CA phase and the formation of the CA₂ phase begins. Between 1400 and 1500 °C, CA₂ reacts with the alumina excess, and the CA₆ phase is formed. At temperatures above 1500 °C, the CA₆ phase stabilizes in terms of microstructure and pore volume.

According to the literature, the total porosity (TP) of CA₆ produced by the α-Al₂O₃–CaCO₃ route was 50.5 ± 0.5%. For the samples sintered at 1600 ^oC produced in this work, the results of volumetric density (VD), and total porosity (TP) are presented in Table 3. The samples produced with α-Al₂O₃ presented porosity of around 40%, while the samples produced with α-bond reached around 67% of porosity, except for the sample produced with CaO, which displayed 62% porosity.

Table 3
Volumetric density (VD), and total porosity (TP) of the CA₆ samples sintered at 1600 ^oC.

Compositions	VD (g.cm-3)	TP (%)
α-Al2O3-CaCO3	2.285	41.2
α-Al2O3-CaO	2.367	39.1
α-Al2O3-Ca(OH)2	2.382	38.8
ρ-Al2O3-CaCO3	1.295	66.7
ρ-Al2O3-CaO	1.462	62.4
ρ-Al2O3-Ca(OH)2	1.261	67.6

3.2. Pore size distribution

Figure 3 shows the pore size distribution of samples produced with the compositions containing α-Al₂O₃ and ρ-Al₂O₃.

Figure 3
Pore size distribution of CA₆ samples obtained by direct molding from α-Al₂O₃ and ρ-Al₂O₃ compositions with different calcium sources after sintering at 1550°C and 1600°C.

CA₆, when produced with compositions α-Al₂O₃, presented a smaller average pore diameter (APD), in the range of 0.75 and 1.15 μm, when compared to ρ-Al₂O₃, in the range of 2-3 μm. This occurs due to the packing of the particles, which is facilitated by the size of the grains and the intertwining between them in the crystalline structure, resulting in a smaller percentage of pore volume compared to ρ-Al₂O₃. This can be explained by the greater difference in particle size between ρ-Al₂O₃ and CaO when compared to the significantly smaller particle size difference between α-Al₂O₃ and CaO. Because hydratable alumina presents a larger average particle size compared to the CaO, its packing efficiency is significantly lower which generates higher pore size in the samples. Another factor that contributes to the quantity and size of pores is the percentage of water in the composition during sample preparation, which is released as vapor during the sintering process leaving voids and small pores. Samples α-Al₂O₃-CaCO₃ presented a platelet structure, whereas for the ρ-Al₂O₃-CaCO₃ ones, a more equiaxial morphology in grains with curved boundaries and higher porosity and consequently lower mechanical properties.

3.3. Diametral compression strength, apparent porosity, and elastic modulus after thermal shock

Samples’ diametric compression strength, apparent porosity, average pore size, and elastic modulus measured before and after 10 cycles of heating and cooling are presented in Figures 4 and 5. Samples showed very similar levels of total porosity, apparent density, and average pore diameters before and after the procedure. The porous structure of the material proved to be very stable and was not significantly affected by thermal shock. The small diameter pores contributed to the material's low thermal conductivity and its resistance to thermal shock¹⁰,¹²,¹³,⁵³,⁶³.

Figure 4
(a) Diametral compression strength, (b) apparent porosity, and (c) elastic modulus of samples CA₆ obtained by direct molding from compositions of α-Al₂O₃ and ρ-Al₂O₃ with different calcium sources, sintered at 1600 °C before and after thermal shock procedure.

Figure 5
Pore size distribution of samples CA₆ obtained by direct molding from α-Al₂O₃ and ρ-Al₂O₃ compositions with different calcium sources, sintered at 1600 °C before and after thermal shock procedure.

The lamellar shape of crystals CA₆ can behave as a toughening mechanism, promoting crack deflection and grain interlocking to reduce damage by thermal shock²,³,⁹,²¹,²². The results attained for α-Al₂O₃-containing compositions corroborate this information, as there was no significant strength variation before and after the thermal shock cycles. On the other hand, samples produced with ρ-Al₂O₃ had significant thermal shock damage, which is characterized as a strength reduction after thermal cycling. This behavior may be associated with the arrangement of the crystals in a house of cards, a structure that is not favorable to heat transport by radiation and results in low thermal conductivity at high temperatures (> 1000 °C), and the higher porosity and the reduced pore dimensions, as verified through Hg porosimetry tests. Scanning electron microscopy for samples sintered at 1600 ^oC is shown in Figure 6.

Figure 6
Scanning electron microscopy of samples CA₆ obtained by direct molding from α-Al₂O₃ and ρ-Al₂O₃ compositions with different calcium sources, sintered at 1600 °C before thermal shock procedure.

Assessment of the elastic modulus of ceramics using non-destructive tests is useful in determining its structural integrity⁵²,⁵³. This property is significantly affected by the presence of cracks or laminations and by the degree of connection between them. The employed technique involves the determination of resonance frequencies of harmonic waves induced in a bar employing a quick mechanical excitation. According to the literature, the reduction in elastic modulus after a thermal shock cycle is associated with the spontaneous formation of microscopic cracks in ceramic samples. However, when the pores are randomly distributed, the effects of porosity interfere as little as possible with elastic modulus, which explains why the samples did not show changes in modulus after thermal shock⁶⁵,⁶⁶,⁶⁸,⁶⁹.

Thermal shock damage, when not catastrophic, is more intense at the surface of the material. This explains the reduction of mechanical strength after the thermal shock cycles since surface cracks control this property. On the other hand, porosity and elastic modulus are associated with the bulk sample, in which microstructure remained practically unchanged after thermal shock cycles.

4. Conclusions

Calcium hexaluminate (CA₆) samples prepared with ρ-Al₂O₃ present much lower uniaxial and diametral compression strength and elastic modulus when compared to those prepared with α-Al₂O₃, regardless of the calcium source, which is by the higher apparent porosity (approximately 65%) verified for these samples when comparing the α-Al₂O₃ compositions (approximately 25%), which can be correlated to the higher water consumption in the preparation of ρ-Al₂O₃ suspensions. Such samples CA₆ also have larger pores (in the range of 2 to 3 μm) when compared to α-Al₂O₃ compositions (in the range of 0.8 to 1.2 μm). The same behavior about the raw materials used was verified after the thermal shock procedure and it was verified that the samples produced were resistant to the procedure. The lamellar shape of particles CA₆ (mainly in α-Al₂O₃) favors anchoring, improving mechanical performance, with the route with CaO being the one with the highest values, showing that it is possible to synthesize CA₆ using a more environmentally friendly system.

5. Acknowledgments

The authors would like to acknowledge the Brazilian Research Foundations: FAPESP #2015/01541-0 and #2010/19274-5, CNPq and Finep for supporting this research and Multi-User Analytical Center of Univap.

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Publication Dates

• Publication in this collection
  11 Nov 2024
• Date of issue
  2024

History

• Received
  05 May 2024
• Reviewed
  20 Aug 2024
• Accepted
  20 Sept 2024