3.1. CUG-1A–TiO2 Ball-Milling Process
As previously mentioned, lunar regolith particles exhibit irregular shapes and a wide range of particle-size distributions. Sintering without pre-treatment adversely affects the performance of a sample [
10,
11]. During the sintering of fine particles, the atomic diffusion distance reduces, thus enhancing particle solubility in the liquid phase and facilitating faster and more efficient sintering; this contributes to the improved properties of the sample [
41]. Therefore, to eliminate this interfering factor, the CUG-1A–TiO
2 powder was initially subjected to ball milling for different durations to determine the optimal process. The particle sizes of the CUG-1A–TiO
2 powder with different ball-milling durations are presented in
Table 3. After 6 h of ball milling, the particle size of T0 (D10, D50, and D90) decreased by 131%, 233%, and 163%, respectively, compared with the initial state (0 h). However, after a milling duration of 8 h, no significant change was observed in the particle size of CUG-1A–TiO
2. The variation curve for the specific surface area of the T6 powder after ball milling was also examined (
Figure 1); it showed an increase with a prolonged ball-milling duration. When the milling duration exceeded 6 h, no substantial changes were observed in the specific surface area of the powder. These results indicate that after a ball-milling duration of 6 h, the effect of ball milling on the CUG-1A–TiO
2 powder particles becomes negligible. Considering both the post-milling properties and experimental efficiency factors, we determined a ball-milling duration of precisely 6 h to be optimal.
The SEM image in
Figure 2 shows the T6 powder morphology at the milling durations of 0, 2, and 6 h. An evident refinement and uniformity of the powder can be observed with increasing milling duration. As shown in
Figure 2c, the large particles depicted in
Figure 2a are eliminated, whereas the aspect ratio and specific surface area of the particles are enhanced. Thus, ball milling for 6 h to prepare and sinter the powder yields a sample surface that exhibits significantly improved performance compared with direct preparation and sintering without ball milling.
3.2. Phase Analysis
The phase evolution during the sintering process was analyzed to better understand the effect of varying the TiO
2 content on the thermodynamic behavior of CUG-1A–TiO
2. The XRD pattern (
Figure 3) of the CUG-1A–TiO
2 samples indicates the presence of pseudobrookite (PDF#41-1432), diopside (PDF#41-1370), anorthite (PDF#41-1481), forsterite (PDF#34-0189), augite (PDF#41-1483), and amorphous glass. Additionally, a weak diffraction peak corresponding to magnesium dititanate (MgTi
2O
5) (PDF#35-0792) is observed for T10. Compared with T0, the diffraction peak of pseudobrookite (Fe
2TiO
5) becomes more apparent as the TiO
2 content increases, whereas the intensity of the hematite (Fe
2O
3) (PDF#33-0664) peak weakens. All the PDF cards can be found in the
Supplementary Materials. When the TiO
2 content exceeds 6 wt%, some hematite peaks disappear. Furthermore, as the TiO
2 content increases, the intensity of the diffraction peaks for certain augite and forsterite minerals initially increases and then decreases. This may be due to the potential vitrification of silicate minerals at high temperatures, forming an amorphous phase. Alternatively, it could be attributed to the generation of volatiles with low melting points. The XRD patterns of T6 samples sintered at different temperatures are presented in
Figure 4. It is evident that with increasing temperature, the diffraction peaks of pseudobrookite, diopside, anorthite, forsterite, and augite initially intensify and then weaken; however, the peaks corresponding to hematite (Fe
2O
3) diminish and eventually disappear. Nevertheless, when the temperature reaches 1160 °C, prominent diffraction peaks cannot be observed in the XRD curve. The XRD (
Figure 4) shows the loss of peaks at an elevated sintering temperature, indicating the mineral underwent significant vitrification, resulting in the formation of a glass phase, and it also indicates the formation of a magnesium-rich phase. Some studies have demonstrated that the melting temperature plays a crucial role in silicate formation and that, in addition to other factors, the mineral composition, flux compounds, presence and abundance of volatiles, and disordered substances contribute to the variations in the melting temperature of CUG-1A–TiO
2 [
42]. In addition, as is known, the amorphous phases with disordered structures and low softening temperatures can be more unstable [
18,
35]. Furthermore, the alteration could be intensified by elevated temperatures.
The phase compositions of the air-sintered samples undergo significant changes, including the formation of Fe
2O
3, Fe
2TiO
5, and TiO
2. According to the research conducted by Zhang et al. [
43], the formation of ilmenite products during sintering is highly dependent on temperature. At temperatures ranging from 600 to 800 °C, ilmenite undergoes oxidation, converting ferrous iron to ferric iron, and new phases such as Fe
2Ti
3O
9, Fe
2O
3, and TiO
2 emerge. However, the Fe
2Ti
3O
9 phase is metastable and disappears at temperatures above approximately 1000 °C, decomposing into Fe
2TiO
5 and TiO
2. At 1000 °C, Fe
2TiO
5 begins to form due to the decomposition of Fe
2Ti
3O
9 and the combination reaction of Fe
2O
3 and TiO
2. The correlation equations are presented in (4), (5), and (6).
Chen et al. [
44] investigated the elevated-temperature behavior of ilmenite and proposed that the oxidation products of FeTiO
3 vary with increasing temperature. At 600–800 °C, ferrous iron oxidizes to form an intermediate Fe
2O
3·2TiO
2 phase (7). Chen et al. [
45] also noted that the Fe
2O
3·2TiO
2 and Fe
2Ti
3O
9 phases are metastable at high temperatures. Thus, at 600–1000 °C, the Fe
2O
3·2TiO
2 phase decomposes to form Fe
2O
3 and TiO
2 (8). Furthermore, Fe
2Ti
3O
9 decomposes into Fe
2O
3 and TiO
2 at above ~1000 °C (9). Fe
2TiO
5 is formed from the recombination of Fe
2O
3 with TiO
2 at 1000–1200 °C (10). Although the specific temperature for the high-temperature formation of Fe
2TiO
5 varies across different studies, a general correspondence can be noted between the formation and the temperature range within which Fe
2TiO
5 is formed.
By analyzing the T0 curve and cross-referencing
Table 2, we can see that the CUG-1A utilized in this experiment belongs to the category of low-titanium lunar regolith simulant, with a TiO
2 content of approximately 1.9 wt%. Consequently, even after the aforementioned reactions, the formation of Fe
2TiO
5 is inevitably limited. This observation is consistent with the subdued diffraction peak of Fe
2TiO
5 in the T0 curve of the XRD pattern. With increasing TiO
2 content, hematite (Fe
2O
3) can react sufficiently with abundant TiO
2 to generate Fe
2TiO
5 (6) or (10) at elevated temperatures. This phenomenon further indicates the weakening of the diffraction peak of Fe
2O
3, along with a gradual enhancement of the diffraction peak of Fe
2TiO
5. These findings are consistent with those of Zhang et al. [
43] and Chen et al. [
44].
3.3. Thermal Analysis
The synchronous thermal analysis (TGA-DSC) results of T0 and T6 are presented in
Figure 5, revealing distinct differences between the two sets of curves. Additionally, the weightlessness process for T0 can be roughly categorized into three stages (
Figure 5a): the first stage occurs from room temperature to approximately 380 °C, during which the primary transformation in the sample involves the volatilization of water, encompassing both adsorbed water molecules on particle surfaces and within the crystal lattice. The second stage takes place between around 380 °C and 550 °C, characterized by a discernible steepening of the curve slope, indicating an elevated rate of mass loss. Wilkerson et al. [
21] determined that the mass loss events occurring at different temperature ranges during the heat treatment of lunar regolith simulant JSC-1A have distinct origins and compositions. The mass loss observed from approximately 200 °C to 500 °C primarily consists of H
2O, CO, CO
2, SO
2, and SO
3, which is likely a result of either physisorbed gases or the decomposition of carbonates and sulfates with low stability. A significant mass loss, accompanied by a substantial evolution of CO
2, is observed in the temperature range of 500 °C to 600 °C. This mass loss is attributed to the decomposition of a non-lunar trace carbonate, such as CaCO
3. Notably different from the vacuum sintering investigated in that study, air sintering involving oxygen effectively increases the heat energy required for the fusion process of amorphous phases and expedites their molten state evaporation [
45,
46]. Thus, considering that the sintering was conducted in an ambient air environment in this study, and the CUG-1A composition contained only minimal amounts of sulfide, it is likely that the material loss at this stage resulted from hydration phase decomposition and carbonate decomposition. The third stage occurs above ~700 °C; the rate of weightlessness decreases. The evident heat absorption peak is observed at 1083 °C in the corresponding DSC curve. This phenomenon may be attributed to the depletion of specific silicate materials, and the amorphous phase is vitrifying these mineral phases.
The TGA-DSC results of T6 (
Figure 5b) show an unexpected rise in the TGA curve between 600 °C and 783 °C, indicating the formation of new substances during high-temperature sintering. In addition, from 765 °C to 773 °C, a simultaneous exothermic peak appears in the DSC curves, providing evidence for our inference in
Section 3.2: we speculate that TiO
2 inclusion causes an increase in the amount of Fe
2Ti
3O
9, which subsequently decomposes into pseudobrookite (Fe
2TiO
5) as the temperature further increases. Notably, at equivalent temperatures, T6 experiences a substance loss that is 2.4465% lower than that of T0, suggesting that TiO
2 can react with molten amorphous Fe to form Fe
2TiO
5 and reduce the generation of volatile amorphous phases during the high-temperature firing of the CUG-1A. The diffraction peaks of pyroxene and olivine in the experimental group are stronger than those for T0, further supporting the results of this analysis, which means that the addition of TiO
2 may inhibit the decomposition or evaporation of mineral components.
3.4. Microstructure and EDS Analysis
Figure 6 presents the SEM images of the fractured and polished microstructure surfaces of the samples sintered at 1100 °C. Noticeable differences can be observed among T10 (e), T6 (f), T4 (g), and T0 (h). T0 exhibits irregular large pores in patch-like formations, while T10, T6, and T4 display individual pores. Moreover, the number of pores within the same observation range initially decreases and then increases with increasing TiO
2 content. The evaporation process during sintering restricts material transfer and reduces the presence of pores, resulting in the formation of macropores [
18,
45]. I. P. Alekseeva et al. [
47,
48] have also suggested that TiO
2 induces liquid unmixing, causing metastable liquid phase separation and lowering the crystallization temperature of glass. Similarly, Lim et al. [
49] demonstrated that the addition of TiO
2 to the regolith simulant decreased the viscosity of molten slag and improved the wettability between molten iron and slag.
Combining the results of SEM and XRD and the aforementioned related investigations, we postulate that the underlying factors contributing to this phenomenon may be as follows: First, air sintering facilitates liquid-phase melting to promote sintering. Second, the addition of TiO
2 reduces the viscosity, enhances the wettability and fluidity of glass, and facilitates element migration within the glass system, thereby lowering the crystallization temperature of the glass. Third, an appropriate amount of TiO
2 not only increases the titanium content in CUG-1A but also reacts with Fe
2O
3 to form Fe
2TiO
5 and with Fe in the molten amorphous state to form Fe
2TiO
5. These reactions effectively prevent substance evaporation and an excessive increase in the amount of liquid-phase components caused by the formation of a low-melting-point iron-containing solid solution. These also accelerate the mass transfer rate during sintering, leading to closer particle bonding. Consequently, the large pores cluster together and shrink into independent pores, ensuring uniform sample shrinkage, which contributes to enhance the relative density of the samples, resulting in an increase in flexural strength. The fracture morphologies of the samples are shown in
Figure 6i–l. Compared with T0, which still exhibits a granular structure after sintering (l), the experimental groups ((i), (j), and (k)) clearly display an increased presence of a fuller liquid phase during the sintering process. This phenomenon further supports the aforementioned viewpoint that the addition of TiO
2 promotes liquid-phase formation and facilitates closer bonding between neighboring particles, which will enhance the The relative densities of T4 and T6 samples; thus, the performance characteristics of these samples are considerably enhanced [
18,
45,
46].
However, the excessive addition of TiO
2 may have the opposite effect. When the TiO
2 content increases from 6 wt% to 10 wt%, the amounts of the liquid-phase products decrease (
Figure 6e). At higher temperatures, TiO
2 can act as an intermediate oxide within the glass system, replacing Si
4+ in the silicon-oxygen tetrahedral structure and forming coordination bonds with O
2−. This process disrupts the intricate silicon–oxygen tetrahedral network. Consequently, excessive doping of TiO
2 can significantly decrease the viscosity of glass [
50]. However, when TiO
2 is appropriately doped, it can effectively reduce viscosity and improve the fluidity of the glass. Severely reduced viscosity of the glass results in macroscopic deformation of the sample. The sample will experience severe deformation following sintering at an elevated temperature, such as 1160 °C. Furthermore, although the reduction in glass viscosity and increase in the liquid-phase content promote element migration within the glass system [
51,
52] and enhance the mass transfer processes, an excessive amount of TiO
2 can react with the Mg
2+ present in the low-melting-point magnesia solid solution or pyroxene to form magnesium dititanate (MgTi
2O
5) [
53,
54] (as indicated by the XRD curve of T10 shown in
Figure 3). As mentioned previously, an excess of evaporating substances during sintering can impede mass transfer. Similarly, an excess of the molten state can hinder gas discharge and increase the trapped gas pressure within the closed pores. This results in a higher number of closed pores, which inevitably have a detrimental effect on the sintering pattern performance. A significant increase and expansion of the internal pores are evident in T10 (
Figure 5e). Consequently, the inclusion of excessive TiO
2 doping in T10 leads to a significant deterioration in its properties, including shrinkage, relative density, and mechanical characteristics when compared to T6.
Table 4 presents the EDS results for the marked regions in
Figure 6. Regions A–C exhibit high concentrations of Fe, Ti, and O, while region D displays lower levels of Fe and Ti. The predominant components in regions A–C are Fe
2TiO
5, Fe
2O
3, and TiO
2, which align with the chemical Equations (7)–(9) when combined with the XRD findings. On the other hand, region D primarily consists of hematite (Fe
2O
3). The low presence of Mg in region D, along with the TGA-DSC results (
Figure 5a) and its secondary analysis, indicate the evaporation of the molten amorphous state. The reduced Mg content is presumed to be the result of its consumption from the amorphous phase in order to form numerous Fe-Ti-O phases. Consequently, Ca is not being concentrated elsewhere for the formation of other mineral crystals. Naturally, the hindrance of the mass transfer process due to evaporation inevitably leads to the formation of macroscopic pores. Regions A–C have high Mg and Fe contents. Based on the inference from the thermal analysis (
Section 3.3), we can conclude that the TiO
2 content effectively mitigates the generation or volatilization of certain low-melting-point substances during the high-temperature sintering of CUG-1A. This confirms that TiO
2 plays a positive role in reducing the evaporation of Mg- or Fe-containing low-melting-point substances formed during sintering [
18,
45], thereby enhancing the properties of the samples.
3.5. Sintering Shrinkage and Relative Density of CUG-1A–TiO2 Samples
The appearance of the samples after sintering is shown in
Figure 7. Notably, T6 (d), subjected to a sintering temperature of 1160 °C, exhibits severe deformation, the observed results appear to be consistent with the findings from our previous XRD analysis (as indicated by the XRD curve of 1160 °C shown in
Figure 4). whereas the surface of T0 (b) displays unevenness. The other samples are well-sintered. Furthermore, T0 appears reddish-brown owing to hematite formation, as confirmed in the XRD analysis (
Figure 3). In contrast, T6 displays a yellowish-brown color.
As the samples exhibit a round-like shape in the X–Y plane, the X-axis and Y-axis are combined into a single X-axis during measurements.
Figure 8 shows the linear shrinkage rate of each group of sintered samples in the X- and Z-axes at varying sintering temperatures. The relative densities of the sintered samples are shown in
Figure 9, and the bulk density and real density of the sintered samples are shown in
Table 5 and
Table 6, respectively. The shrinkage rate of the samples initially increases and then decreases with increasing temperature (
Figure 8), and the relative density of the samples exhibits a similar pattern (
Figure 8). Additionally, the shrinkage rate of the samples shows the same trend as the content of TiO
2 at the same temperature. The shrinkage rate along the X-axis is generally lower than that along the Z-axis. At the sintering temperature of 1100 °C, all samples exhibit maximum shrinkage in all directions and achieve the highest relative density. At 1100 °C, among all samples, T6 exhibits the highest sintering shrinkage of 17.9 ± 0.2% and 21.7 ± 0.5% along the X- and Z-axes, respectively, which are 13.29% and 25.43% higher than those of T0, respectively. The relative density of T6 reaches 91.06%, indicating a significant enhancement of 12.28% compared with that of T0, which has a relative density of 81.10%. The observed trends in the shrinkage rate and relative density are consistent with those discussed in
Section 3.4.
3.6. Mechanical Properties
To assess the impact of TiO
2 content on the mechanical properties of the CUG-1A samples, the flexural strengths of the sintered samples were determined and are illustrated in
Figure 10. The experimental group generally exhibits higher flexural strength compared to T0, indicating the influence of TiO
2 content on the mechanical properties of CUG-1A. The flexural strength of all sample groups initially increases and then decreases with increasing sintering temperatures. At a sintering temperature of 1100 °C, all sample groups achieve optimal flexural strength. However, unlike T0, which exhibits only slight deformation (
Figure 7b), samples in the experimental group experience severe melting deformation (
Figure 7d) and cannot be tested for flexural strength at 1160 °C. The enhancement in the flexural strength of the experimental group samples is not proportional to increasing TiO
2 content. Instead, it initially increases and then decreases; this is consistent with the trends observed for the sample shrinkage rate and relative density. Microstructural analysis reveals that appropriate TiO
2 doping enhances the flexural strength of CUG-1A. However, excessive TiO
2 doping leads to an increase in liquid-phase content during the sintering process, thereby impeding gas discharge. With increasing sintering temperature, decomposed substances continuously generate gas, resulting in elevated trapped gas pressure within closed pores and, subsequently, leading to an augmentation in the number of closed pores. Consequently, both flexural strengths and relative densities of the samples decrease.
The flexural strength of T6 sintered at 1100 °C reaches the highest value of 136.66 ± 4.92 MPa, exhibiting a remarkable increase of 65% compared to that of T0. These results demonstrate the effectiveness of TiO2 as an additive component in enhancing both relative density and flexural strength for lunar regolith structures.