Aluminum nitride (AlN) ceramics are widely used in various applications including microelectronic and optoelectronic devices and heat dissipation systems owing to their excellent properties, such as their high thermal conductivity and resistivity [1-4]. Moreover, AlN is well known for its extremely high theoretical thermal conductivity (approximately 320 W/mK). However, due to limitations in their production technology, it is not possible to entirely prevent impurities and microstructural defects in polycrystalline AlN. Thus, while the thermal conductivity of AlN is usually lower than its theoretical value, it still has high values [5-6], making it an excellent heat transfer component for practical use. In addition, AlN has a coefficient of thermal expansion (CTE) similar to that of Si [7], an important parameter with respect to semiconductor manufacturing. Moreover, because of its excellent mechanical properties and high corrosion resistance [8-9], AlN is a reliable protective material for use in machine manufacturing and heavy industry production.

Recently, various surface techniques, such as physical vapor deposition (PVD) and plasma spraying [10-12], have emerged for the facile fabrication of AlN thin films. Conversely, sintering is the most common approach for fabricating AlN bulk materials. However, AlN has a hexagonal crystal structure with extremely strong covalent bonding, and its self-diffusion coefficient is low, posing challenges to sintering processes. To obtain sintered AlN ceramics with high densities, higher sintering temperatures and longer sintering times are required when employing conventional sintering methods, such as pressureless sintering and hot-pressing sintering. This results in elevated energy consumption and, consequently, increased manufacturing costs. Spark plasma sintering (SPS), a relatively new sintering technology, has several advantages including a high heating rate, short sintering time, and nonpolluting nature [13-14], which can be used to prepare high-density ceramics at relatively low temperatures within short periods. During the SPS process, the temperature, pressure, and pulsed direct current are directly varied. For insulating powders such as AlN, the application of pulsed direct current on the surfaces of the powder particles can induce instantaneous discharge within the voids between the particles. This leads to Joule heating of the particles due to the effects of surface activation and purification induced by discharging, and resulting in an increase in the degree of densification.

By optimizing the sintering parameters, such as the heating rate, sintering temperature, and holding time, high-density sintered ceramics can be achieved. In this way, AlN ceramics with densities exceeding 95 % can be prepared by SPS without the need for sintering aids [15-16]. Research on AlNbased composites fabricated by SPS has attracted widespread attention in recent years to develop composite materials suitable for a range of applications. For example, Wang et al. [17] fabricated AlN-W composites for high-average-power electronic tubes. Gao et al. [18] reported that AlN-based composites containing Mo exhibited better overall performance than those without the additives. Also, AlN-SiC composite ceramics were fabricated by SPS for applications involving electromagnetic waves with wavelengths ranging from 1 kHz to 1 MHz [19]. Other composites such as AlNTiN, AlN-BN, and AlN–TiB₂ have also been investigated [20-22].

However, studies on AlN-MgO composites have been limited, and there is a shortage of reference data on these materials. The synthesis of MgAl₂O₄ and MgAlON ceramics with spinel structures can be achieved through various sintering methods using AlN, MgO, and Al₂O₃ as the raw materials. The resulting ceramics exhibit desirable properties, including high transparency, good slag corrosion resistance, good mechanical properties, and moderately high CTE values [23-26]. However, the feasibility of obtaining spinel-structured ceramics using AlN and MgO as the raw materials is still challenging. Moreover, the performance of AlN-MgO composites needs to be investigated to assess their practical applicability.

Therefore, in this study, we prepared AlN-MgO composites with different compositions by SPS using AlN and MgO as the raw materials. The phase compositions and microstructures of the composites were investigated, and their thermal and mechanical properties were compared with those of pure AlN and MgO samples sintered under the same conditions.

The XRD patterns of the sintered samples are shown in Fig 2. Only AlN and MgO phases were detected in the sintered AlN-MgO composites, regardless of the weight ratio of the raw materials. This indicates that there were no phase transitions or reactions between MgO and AlN during the sintering process. A diffraction peak related to Al₂O₃ was observed in the pattern of the sintered pure AlN ceramic, suggesting the presence of oxide impurities. Magnified diffraction peaks related to AlN (110) and MgO (200) are shown in Fig 2 (b) and (c). As can be seen from Fig 2(b), with increasing MgO contents (MgO contents lower than 60 wt%) the AlN (110) peak of the sintered AlN-MgO composites is shifted to a lower 2θ value compared with the peak of the pure AlN sample. Similarly, Fig 2(c) shows that the MgO (200) peak of the sintered AlN-MgO composites is shifted to a higher 2θ value compared with the peak of the pure MgO sample with increasing MgO contents, (MgO content lower than 80 wt%). These peak shifts may be caused by the formation of a solid solution phase, indicating that solid solutions were formed in both the AlN and MgO lattices of the sintered samples with different compositions.

Fig 3 shows the XPS spectra of the sintered AlN-MgO composites with different compositions. As can be seen from Fig 3(a)–(d), the Mg 2p and O 1s peaks of the sintered AlN-MgO composites exhibit significantly higher bonding energies (BEs) compared with those of peaks of the sintered pure MgO sample, indicating that the chemical bonding states between Mg and O were different in the composites. However, only a minor shift in the BEs was observed for the Al 2p and N 1s peaks. Based on these results, it can be concluded that a solid solution was primarily formed in the MgO lattice. Fig 3(e) and (f) show the fitted O 1s peaks of the sintered pure AlN and MgO-60 samples. For the sintered pure AlN sample, in addition to the peak related to the adsorbed oxygen (BE = 532.1 eV [27]), a peak was also observed at 531.6 eV; this corresponded to the Al-O bond [27]. This confirms the presence of Al₂O₃ as an impurity in the sample, consistent with the XRD analysis results. With the sintered MgO-60 sample, the peaks with BE values of 530.8 eV and 531.5 eV could be ascribed to MgO [28] and magnesium–aluminum composite oxides (MgAl_xO_y), respectively [27], indicating the formation of a (Mg, Al)O solid solution. The fitted N 1s peaks of the sintered pure AlN and MgO-40 samples are shown in Fig 3(g) and (h), respectively. In addition to a peak related to the Al-N bond (BE = 396.7 eV [30]), a peak related to the Al-N-O bond was also observed (BE = 396.3 eV) [31] in the spectrum of the sintered MgO-40 samples, indicating the formation of an Al(N, O) solid solution phase. The oxygen in the AlN solid solution was derived either from the MgO lattice, the Al₂O₃ formed on the AlN surface, or the residual oxygen in the sintering chamber. Thus, the XPS results were consistent with those of the XRD analysis, further confirming that the different types of solid solutions were formed within the MgO and AlN lattices with changes in the composition of the sintered AlN-MgO composites.

The calculated relative densities of the sintered samples are summarized in Table 1. Except for pure MgO, all of the sintered samples exhibited relative densities higher than 95%. As shown in Fig 4(a), SEM surface images showed that the sintered pure MgO sample had a grain size ranging from 10 μm to more than 100 μm. However, residual pores were present at the grain junctions and coarse grain boundaries in the sintered pure MgO samples (Fig 4(b) and (c)). It has been reported that MgO ceramics can be sintered readily at temperatures lower than 1000 °C [32] and MgO undergoes rapid grain growth at sintering temperatures higher than 1200 °C [33]. Hence, it can be concluded that, in this study, the extremely rapid growth of the MgO grains was induced by the high sintering temperature used during the short sintering process. This resulted in the formation of pores and the weakened grain boundaries, leading to a low relative density in the sintered pure MgO sample.

The detailed microstructures of the sintered samples were analyzed using TEM. Fig 5 shows the low-magnification cross-sectional TEM images of the sintered pure MgO, AlN, and MgO-60 samples. As can be seen from Fig 5(a), no grain boundaries were observed in the sintered TEM sample of pure MgO, indicating that it had an extremely large grain size (from 10 μm to more than 100 μm), as observed in Fig 4 (a). Since the average size of the pristine MgO powder particles was approximately 1 μm, the MgO grains grew rapidly during the sintering. Similarly, significant grain growth was observed with the sintered pure AlN sample, which also showed a significantly larger grain size (2–3 μm) after sintering (Fig 5(b)).

However, as shown in Fig 5(c), the MgO-60 sample exhibited a finer-grained microstructure. The MgO and AlN grains of the sample could be distinguished readily based on the results of EDS mapping results (Fig 5 (d) and (e)): the grain sizes were determined to be 1.0–1.5 μm and 0.5–1.0 μm, respectively. The average grain sizes of the MgO and AlN in the AlN-MgO composites were determined from their cross-sectional TEM images, and the results are shown in Fig 6. Neither the MgO and AlN grains showed significant growth during the sintering process, probably because MgO and AlN can be regarded as secondary phases with respect to each other. Therefore, it is considered that their grain growth was mutually inhibited during sintering. In addition, except for the MgO-80 sample, the size of the MgO grains in the composites decreased with increasing MgO content, while the variations in the size difference between the MgO and AlN grains was the same as that seen in the MgO grains. This suggests that the inhibition of grain growth was markedly more pronounced in the MgO-60 sample, which had a relatively smaller grain size. Relatively larger differences between the MgO and AlN grains were present in the MgO-80 sample, indicating that the growth of the MgO grains was impeded with a low AlN content.

The state of grain boundaries has a critical effect on the characteristics of sintered materials. Two types of grain boundaries were observed in the sintered composites, and their TEM images are shown in Fig 7. Fig 7 (a) shows a low-density grain boundary in the sintered MgO-20 sample. The corresponding EDS maps are presented in Fig 7(b) and (c), and show that the low-density grain boundaries were composed of a great part of oxygen and the low Mg content. Hence, the oxygen-rich interface was more likely attributed to the residual oxygen remaining between the grains, rather than to the diffusion of MgO. Fig 7(d) and (e) show TEM images of the grain boundary observed in the sintered MgO-60 sample. A clean interface was formed between the MgO and AlN grains without the aggregation of oxygen. The inverse fast Fourier transform (IFFT) result for the marked area, shown in Fig 7(f), reveals an interplanar spacing of 0.473 nm, corresponding to the MgAl₂O₄ (111) plane. MgAl₂O₄ is a spinel phase with a cubic structure and is formed by a reaction between MgO and Al₂O₃ under certain conditions [34]. Therefore, the formation of the MgAl₂O₄ spinel phase at the clean interface can probably be attributed to a reaction between MgO and the Al₂O₃ layer from the AlN surface.

However, given the presence of AlN grains, it is likely that a MgAlON phase was also formed at the MgO and AlN grain boundaries. MgAlON has the same structure as the MgAl₂O₄ spinel phase, and is considered to be a solid solution formed by MgO/MgAl₂O₄ when they enter the γ-AlON lattice, or because of the dissolution of Al₂O₃ and AlN in MgAl₂O₄ [35]. Dai et al. [36] also reported that a MgAlON phase is formed during SPS at 1700 °C when MgO, AlN, and Al₂O₃ were used as the raw materials. Therefore, it is likely that, in this study, MgAlON was formed at the grain boundaries of MgO and AlN. Thus, despite the difficulty in distinguishing MgAl₂O₄ from MgAlON, it can be concluded that a spinel phase was formed at the clean interface between the MgO and AlN grains.

Fig 8 shows the measured thermal conductivities and CTE values of the sintered samples. The corresponding thermal diffusivity and heat capacity values are listed in Table 2. Sintered pure AlN exhibited the highest thermal conductivity (53 W/mK) of all the samples, followed by sintered pure MgO (39.7 W/mK). The AlN-MgO composites showed low thermal conductivities, which varied from 14.9 to 33.3 W/mK. While AlN ceramic typically shows high thermal conductivity, the AlN ceramic prepared by SPS in this study exhibited a thermal conductivity of less than 55 W/mK. This could be attributed to its fine-grained microstructure, meaning it contained more grain boundaries compared with single-crystal AlN. Grain boundaries are a type of planar defect that scatter phonons, thereby decreasing thermal conductivity [37].

In addition, oxygen-related defects, including the residual oxygen present at the interface, the oxide layer on the AlN surface, and the lattice defects induced by oxygen (substitutions and vacancies), contribute to the reduction of the thermal conductivity through phonon scattering [38-39]. As a result, the AlN ceramic did not exhibit high thermal conductivity in this study. The thermal conductivity of the AlN-MgO composites strongly depended on the MgO/AlN mass ratio and decreased with increasing MgO content. Because MgO has lower thermal conductivity than AlN, the AlN-MgO composites showed reduced thermal conductivities. A previous analysis showed that AlN-MgO composites exhibited a fine-grained microstructure, with solid solutions and an oxygen-rich interface that formed between the grains. All these factors enhanced phonon scattering and thus contributed to the lowering of the thermal conductivity of the AlN-MgO composites. Therefore, the low thermal conductivity of the AlN-MgO composites was a result of both their composition and microstructure. Considering that the thermal conductivity of the MgO-20 sample was lower than that of pure MgO, the effect of the microstructure on the thermal conductivity of the composites was more pronounced in this study.

In the case of the CTE, the pure AlN ceramic showed the lowest CTE value at 4.47 × 10^-6 /K, whereas the pure MgO sample exhibited the highest value, which was 13.06 × 10^-6 /K. In general, the microstructure also influences the CTE value, in addition to the composition. For instance, for the same material system, a sample with finer grains generally shows a higher CTE value than that of a sample with a coarser microstructure. This is because the CTE of the grain boundaries is significantly greater than that of the grains themselves [40]. In the present study, the CTE of the AlN-MgO composites increased from approximately 7.20 ×10^-6 /K to 10.73 ×10^-6 /K with increasing MgO content. However, the opposite was observed in the case of the MgO-40 sample, which showed a slightly lower CTE (approximately 6.49 ×10^-6 /K) than that of the MgO-20 sample (approximately 7.20 × 10^-6 /K). Since the relative densities and microstructures of the sintered composites were similar, the exact reason for the abnormal results observed in the case of MgO-20 and MgO-40 remains unclear and requires further study. Overall, however, the effect of the composition on the CTE was more significant than that of microstructure.

The mechanical properties of the sintered ceramics were evaluated using the Vickers hardness test. Fig 9 shows the calculated hardness and fracture toughness values of the sintered pure AlN, MgO, and AlN-MgO composites. The pure AlN sample exhibited the highest hardness value at approximately 14 GPa. In contrast, the MgO sample exhibited an extremely low hardness at approximately 4 GPa. It was previously reported that MgO ceramics prepared by SPS exhibit a wide range of hardness values, from 8–13 GPa, based on the microstructure of the material [41]. Therefore, the low hardness of the MgO ceramic fabricated in this study may have been caused by its coarse grains. The hardness values of the AlN-MgO composites decreased with increasing MgO content; however, the opposite was observed in the case of the MgO-60 sample, which exhibited the highest hardness (12.86 GPa) of all the composites, close to that of the pure AlN sample. Thus, the hardness of the AlN-MgO composites depended on factors other than their composition.

The pure AlN sample also exhibited a low fracture toughness of approximately 1.9 MPa·m^0.5. With increasing MgO content, the toughness of the AlN-MgO composites first increased, reaching the maximum at a MgO content of 60 wt. %, and then decreased. The indentation morphology of the sintered pure MgO sample was irregular (see Fig 10), which made determining its fracture toughness difficult. The AlN and MgO-60 samples exhibited clear Vickers indentations, with cracks initiating from the diagonal ends of the indentations. In the case of the sintered pure MgO sample, however, extensive cracks were observed originating from the edges of the indentations, indicating poor crack-extension resistance. In Fig 10 (d) it can be seen that the cracks have mainly propagated along the grain boundaries of the sintered pure MgO sample, suggesting the grain boundary cohesion of MgO was weak and resulted in the intergranular fracture. These results indicate the large grain size and coarse grain boundaries of the sintered pure MgO sample were responsible for its poor mechanical properties. Therefore, the primary hardening and toughening mechanisms for the MgO-60 sample are grain refinement and strengthening. In addition, the formation of a solid solution also helps improve mechanical properties. In summary, the MgO-60 sample exhibited the best overall mechanical properties in this study.

In this study, AlN-MgO composites with different compositions were fabricated using SPS. The effects of the composition on their structure, thermal properties, and mechanical properties were investigated. Pure AlN and MgO were also sintered under the same conditions for reference. The results indicated that the AlN-MgO composites were composed of AlN and MgO phases, and different solid solutions were formed within the MgO and AlN lattices. Compared with the sintered pure samples, the AlN-MgO composites showed finer grain microstructures owing to the inhibition of grain growth between the two phases. TEM analysis indicated that two types of grain boundaries were formed in the composites: oxygen-rich low-density grain boundaries and clean boundaries with a spinel phase. The sintered pure AlN sample showed the highest thermal conductivity of 53 W/mK and lowest CTE of 4.47 × 10^-6 /K at 100 °C. Conversely, the thermal conductivity and CTE of the AlN-MgO composites decreased from 33.3 to 14.9 W/mK and generally increased from 6.49 ×10^-6 to 10.73 ×10^-6 /K with increasing MgO content. The composite with an MgO content of 60 wt. % exhibited the best overall mechanical properties, including a hardness of 12.86 GPa and toughness of 3.15 MPa·m^0.5, attributed to its fine microstructure. It also demonstrated a thermal conductivity of 16.8 W/mK and a CTE of 8.14 ×10^-6 /K. Thus, the variations in the thermal and mechanical properties of the AlN-MgO composites can be attributed to the combined effect of their composition and microstructure.