Low-Temperature Rapid Sintering of Dense Cubic Phase ZrW2−xMoxO8 Ceramics by Spark Plasma Sintering and Evaluation of Its Thermal Properties

In this study, we report a low-temperature approach involving a combination of a sol–gel hydrothermal method and spark plasma sintering (SPS) for the fabrication of cubic phase ZrW2−xMoxO8 (0.00 ≤ x ≤ 2.00) bulk ceramics. The cubic-ZrW2−xMoxO8 (0.00 ≤ x ≤ 1.50) bulk ceramics were successfully synthesized within a temperature range of 623–923 K in a very short amount of time (6–7 min), which is several hundred degrees lower than the typical solid-state approach. Meanwhile, scanning electron microscopy and density measurements revealed that the cubic-ZrW2−xMoxO8 (0.00 ≤ x ≤ 1.50) bulk ceramics were densified to more than 90%. X-ray diffraction (XRD) results revealed that the cubic phase ZrW2−xMoxO8 (0.00 ≤ x ≤ 1.5) bulk ceramics, as well as the sol–gel-hydrothermally synthesized ZrW2−xMoxO7(OH)2·2H2O precursors correspond to their respective pure single phases. The bulk ceramics demonstrated negative thermal expansion characteristics, and the coefficients of negative thermal expansion were shown to be tunable in cubic-ZrW2−xMoxO8 bulk ceramics with respect to x value and sintering temperature. The cubic-ZrW2−xMoxO8 solid solution can thus have potential applications in electronic devices such as heat sinks that require regulation of thermal expansion.

Similar to many hard-to-sinter ceramics, the most challenging problem for Mo-doped ZrW 2 O 8 substitutional solid solution is to devise a method to sinter the dense bulk ceramics. The phase diagram of the ZrO 2 -WO 3 system indicates that ZrW 2 O 8 is thermodynamically stable over a limited temperature range; precisely from 1380 K to 1520 K, while it is metastable from absolute 0 K to 1050 K. As it slowly cools from its stable state, ZrW 2 O 8 will decompose into ZrO 2 -ss and WO 3 . Hence, it is difficult to generate the pure phase via the high-temperature solid-state route [26][27][28]. Thus, the investigation of alternative low-temperature liquid-phase methods is rather mandatory. It has been reported that ZrW 2 O 8 powders can be prepared by dehydrating its hydrate precursor ZrW 2 O 7 (OH) 2 ·2H 2 O, which can be obtained using liquid-phase methods such as solgel [1,2,[29][30][31], hydrothermal [3,5,[32][33][34][35][36] and co-precipitation [35]. Kanamori et. al. [32] first reported that the cubic phase undoped ZrW 2 O 8 bulk ceramic, with a relative density of 98.6%, could be prepared by using the combined sol-gel and a spark plasma sintering (SPS) technique at a low temperature; precisely (873 K). There are, however, no reports about the fabrication of dense cubic-ZrW 2−x Mo x O 8 ceramics.
In the present study, a sol-gel-hydrothermal method was used for synthesizing ZrW 2−x Mo x O 7 (OH) 2 ·2H 2 O (0.00 ≤ x ≤ 2.00) precursors. Subsequently, the cubic-ZrW 2−x Mo x O 8 ceramics were prepared by SPS using pre-calcination powders as a raw material. Thermogravimetric-differential thermal analyzer and high-temperature X-ray diffraction techniques were employed to investigate the sintering temperatures of cubic-ZrW 2−x Mo x O 8 ceramics. In addition, the NTE coefficients of cubic-ZrW 2−x Mo x O 8 bulk ceramics were measured using a thermodynamic analyzer. The impact of various sintering temperatures on the relative densities and NTE coefficients of cubic-ZrW 2−x Mo x O 8 bulk ceramics was also investigated.

Synthesis of ZrW
was prepared by a sol-gel-hydrothermal method. In the first stage (sol-gel stage), ZrOCl 2 ·8H 2 O (Kishida Chemical, Osaka, Japan, 99%) was dissolved in 4 mol·L −1 acetic acid and 2-butanol in the air for 3 h. At the same time, WCl 6 (High Purity Chemicals, Saitama, Japan, 99.99%) and MoCl 5 (Sigma Aldrich, Saint Louis, MO, USA, 95%) were separately dissolved in ethanol in a nitrogen atmosphere in a stoichiometric ratio of W:Mo = (2 − x): x (x = 0.0, 0.1, 0.2, 0.5, 0.6, 0.7, 1.0, 1.5, 2.0). Following 3 h of stirring, the tungsten and molybdenum solutions were poured into the zirconium solution and allowed to stir for 72 h. The mixed solution comprising tungsten, molybdenum, and zirconium ions was heated in a silicone oil bath at around 360 K. Sol-gel precursors were obtained after the solution had been completely evaporated. The resultant sol-gel precursors were dissolved in 60 mL of distilled water during the second stage (hydrothermal stage). It was then added to a 100 mL Teflon-lined Parr bomb placed in a mantle heater at 453 K and allowed to stir continuously for 18 h. The sol-gel-hydrothermal precursors were subsequently centrifuged, washed with distilled H 2 O, and dried in an oven at 333 K.

Fabrication of Cubic-ZrW 2−x Mo x O 8 Ceramics
In this study, an SPS apparatus (SPS-515S; Fuji Electronic Industrial, Saitama, Japan), with a pulse duration of 3.3 ms and on/off pulse intervals of 12:2 s, was applied for the preparation of dense cubic-ZrW 2−x Mo x O 8 bulk ceramics. The sol-gel-hydrothermal precursors were pre-calcined at 573-723 K in a drying oven (DX-41, Yamato Scientific, Tokyo, Japan). The pre-calcination powders were ground and sieved to 45-100 µm (typical batch of about 1.0 g), and then placed into a graphite die (10 mm in diameter and 20 mm in depth) with a layer of protective carbon sheet. The sintering temperature was set to the temperatures based on the high-temperature XRD results of the ZrW 2−x Mo x O 7 (OH) 2 · 2 H 2 O precursors. Each sample's heating rate and holding period were 100 K·min −1 and 10 min, respectively. During the SPS procedure, the applied pressure of 50 MPa was held constant in an argon atmosphere until the conclusion of the sintering period. The volume contraction between the two graphite punches, namely the Z-axis linear shrinkage value, was recorded every 30 s during SPS.

Evaluation Method
The phase identification of the ZrW 2−x Mo x O 7 (OH) 2 ·2H 2 O precursors and cubic-ZrW 2−x Mo x O 8 ceramics was carried out using X-ray powder diffraction (XRD; Rigaku, Ultima IV, Tokyo, Japan). The XRD data were collected at a scan speed of 6 • ·min −1 in the 2θ range from 10 • to 40 • under 40 kV/40 mA with Cu-Kα radiation (λ = 0.15406 nm) using a continuum scanning method. To investigate the temperature range in which -ZrW 2−x Mo x O 8 exists as a cubic structure, thermogravimetric-differential thermal analyses (TG-DTA; SHIMADZU, DTG-60/60H, Kyoto, Japan) and high-temperature X-ray diffraction (HT-XRD; Empyrean, PANalytical, EA Almelo, The Netherlands) were utilized. The TG-DTA experiments were carried out with airflow in the range between room temperature and 1073 K, at a heating rate of 10 K·min −1 . The HT-XRD data for ZrW 2−x Mo x O 7 (OH) 2 ·2H 2 O precursors were recorded at the temperature ranging from room temperature to 973 K in the 2θ range of 10 • to 40 • . Prior to the collection of XRD data, the samples were heated to the desired temperature at a rate of 10 K·min −1 and kept for 2 min to equilibrate. The relative densities of the ZrW 2−x Mo x O 8 bulk ceramics were measured using Archimedes' principle in distilled water. The density of distilled water is 0.99754 g·cm −3 at a room temperature of 296 K and the theoretical density of ZrW 2 O 8 has been reported to be 5.08 g·cm −3 [13,14]. The microstructural analysis of the fractured surface of each bulk ceramic was performed with a scanning electron microscope (SEM; JEOL, JSM-IT100, Tokyo, Japan) along with energy dispersive spectroscopic (EDS) mapping. A thermodynamic analyzer (TMA; Shimadzu, TMA-60/60H, Kyoto, Japan), in the temperature range of 323-673 K with a heating rate of 1 K·min −1 was used to measure the changes in length relative to temperature.  samples may be arising from the difference in atomic numbers between Mo (42) and W (74), resulting in a difference of X-ray scattering cross-section. The ZrW2−xMoxO7(OH)2·2H2O precursors were measured using TG-DTA to understand better the compositional and weight variations of ZrW2−xMoxO8 bulk ceramics during heating. TG-DTA curves of ZrW2−xMoxO7(OH)2·2H2O (x = 0.0, 1.0) precursors are depicted in Figure 2a,b. Up to 500 K, the behaviors of all samples are more or less similar. An endothermic peak in the DTA accounted for 7-10% of the initial weight loss in the TG curves. This may occur due to the dehydration and dehydrogenation of the precursors. This behavior was further verified by the observation of the FT-IR spectra of ZrW2−xMoxO7(OH)2·2H2O precursors reported in the literature [3,14,37]. The DTA curves of the x = 1.0 have a pair of exothermic peaks located near 669 K and 817 K, while only one exothermic peak is observed at 863 K for non-doped ZrW2O8. The first exothermic peak is assigned to the formation of the cubic phase and the second exothermic peak of x = 1.0 sample is identified as the cubic-trigonal phase transformation. It has been reported that due to the instability of cubic phase ZrW2−xMoxO8 (x ≥ 1.0), a mixed-phase (cubic + trigonal) is produced during the conversion of the cubic phase to the trigonal phase [4][5][6][7]25]. The occurrence of this phenomenon implies that the preparation of pure cubic phase samples requires accurate analysis.  Figure 2a,b. Up to 500 K, the behaviors of all samples are more or less similar. An endothermic peak in the DTA accounted for 7-10% of the initial weight loss in the TG curves. This may occur due to the dehydration and dehydrogenation of the precursors. This behavior was further verified by the observation of the FT-IR spectra of [3,14,37]. The DTA curves of the x = 1.0 have a pair of exothermic peaks located near 669 K and 817 K, while only one exothermic peak is observed at 863 K for non-doped ZrW 2 O 8 . The first exothermic peak is assigned to the formation of the cubic phase and the second exothermic peak of x = 1.0 sample is identified as the cubic-trigonal phase transformation. It has been reported that due to the instability of cubic phase ZrW 2−x Mo x O 8 (x ≥ 1.0), a mixed-phase (cubic + trigonal) is produced during the conversion of the cubic phase to the trigonal phase [4][5][6][7]25]. The occurrence of this phenomenon implies that the preparation of pure cubic phase samples requires accurate analysis. The ZrW2−xMoxO7(OH)2·2H2O precursors were measured using TG-DTA to und stand better the compositional and weight variations of ZrW2−xMoxO8 bulk ceramics du ing heating. TG-DTA curves of ZrW2−xMoxO7(OH)2·2H2O (x = 0.0, 1.0) precursors are d picted in Figure 2a,b. Up to 500 K, the behaviors of all samples are more or less simil An endothermic peak in the DTA accounted for 7-10% of the initial weight loss in the T curves. This may occur due to the dehydration and dehydrogenation of the precurso This behavior was further verified by the observation of the FT-IR spectra ZrW2−xMoxO7(OH)2·2H2O precursors reported in the literature [3,14,37]. The DTA curv of the x = 1.0 have a pair of exothermic peaks located near 669 K and 817 K, while only o exothermic peak is observed at 863 K for non-doped ZrW2O8. The first exothermic peak assigned to the formation of the cubic phase and the second exothermic peak of x = 1 sample is identified as the cubic-trigonal phase transformation. It has been reported th due to the instability of cubic phase ZrW2−xMoxO8 (x ≥ 1.0), a mixed-phase (cubic + trigon is produced during the conversion of the cubic phase to the trigonal phase [4][5][6][7]25]. T occurrence of this phenomenon implies that the preparation of pure cubic phase samp requires accurate analysis. The compositional and thermal dependence of the phase behaviors of ZrW 2−x Mo x O 7 (OH) 2 ·2H 2 O precursors were examined by HT-XRD measurements. These measurements are anticipated to indicate the underlying mechanism responsible for the phase transition detected by the TG-DTA measurements. Figure 3a,b show the HT-XRD patterns of the

Results and Discussion
It was found that, from room temperature to 373 K, the two precursors maintained their corresponding precursor phases. At 423 K, the ZrW 2 O 7 (OH) 2 ·2H 2 O precursors became amorphous, then transformed into the cubic phase ZrW 2 O 8 at 823 K and a certain amount of WO 3 decomposed at 903 K. The ZrWMoO 7 (OH) 2 ·2H 2 O precursors, on the other hand, changed into an orthorhombic phase, then transformed into cubic-ZrWMoO 8 at 753 K, and ultimately converted into a cubic + trigonal mixed-phase after 833 K. This observation indicates that the change in crystal structure during heating depends on the Mo concentration and synthesis temperature of the cubic phase. The HT-XRD results shown in Figure 4 shows  The compositional and thermal dependence of the phase behavior ZrW2−xMoxO7(OH)2·2H2O precursors were examined by HT-XRD measurements. T measurements are anticipated to indicate the underlying mechanism responsible fo phase transition detected by the TG-DTA measurements. Figure 3a,b show the HTpatterns of the ZrW2−xMoxO7(OH)2·2H2O (x = 0.0, 1.0) precursors. It was found that, room temperature to 373 K, the two precursors maintained their corresponding precu phases. At 423 K, the ZrW2O7(OH)2·2H2O precursors became amorphous, then t formed into the cubic phase ZrW2O8 at 823 K and a certain amount of WO3 decomp at 903 K. The ZrWMoO7(OH)2·2H2O precursors, on the other hand, changed into an o rhombic phase, then transformed into cubic-ZrWMoO8 at 753 K, and ultimately conv into a cubic + trigonal mixed-phase after 833 K. This observation indicates that the ch in crystal structure during heating depends on the Mo concentration and synthesis perature of the cubic phase. The HT-XRD results shown in Figure 4 shows the ov process of phase transformation for the ZrW2−xMoxO7(OH)2·2H2O (x = 0.0, 0.1, 0.2, 0.5 0.7, 1.0, 1.5, 2.0) system during the heating procedure. The synthesis temperature of c ZrW2-xMoxO8 and the width of the synthetic "window" for the cubic phase decreas the x value increases: a gradual shift from the interval of 80 K between 823-903 K (x = to an interval of 50 K between 753-803 K (x = 1.0) and fell to the narrow range only a K for x = 2.0. In this study, the sintering temperatures were set to 923 K and 873 K fo 0.0, x = 0.1, x = 0.2; 873 K and 823 K for x = 0.5, x = 0.6, x = 0.7; 773 K and 823 K for x 743 K and 723 K for x = 1.5; 623 K for x = 2.0, respectively.    Figure 5 shows the Z-axis linear shrinkage curve of the ZrW2O8 bulk ceramic at a sintering temperature of 923 K. In the temperature range of 300-430 K, a slight Z-axis shrinkage is observed, which may be due to the evaporation of H2O and CO3 2-on the surface of the pre-calcining powders. The remarkable shrinkage caused by the densification  Figure 5 shows the Z-axis linear shrinkage curve of the ZrW 2 O 8 bulk ceramic at a sintering temperature of 923 K. In the temperature range of 300-430 K, a slight Z-axis shrinkage is observed, which may be due to the evaporation of H 2 O and CO 3 2− on the surface of the pre-calcining powders. The remarkable shrinkage caused by the densification takes place at a temperature of 450 K and ends at a set temperature of 923 K. It means that the ZrW 2 O 8 bulk ceramics sintered at a higher temperature (923 K) are denser than those sintered at lower temperatures (873 K), where the densification change can be confirmed by SEM micrographs of the fracture surfaces of ZrW 2 O 8 bulk ceramic shown in Figure 6. A relatively loose microstructure consisting of some pores was observed in the 873 Ksintered ceramic (Figure 6a). In comparison, the 923 K-sintered ceramic was dense due to the decrease in the number of pores (Figure 6b). In addition, the Z-axis linear shrinkage curve at the time of holding temperature does not affect the compactness of the bulk ceramic, which means that it takes only 6-7 min to be prepared using the SPS technique. Figure 7 shows the relative densities of ZrW 2−x Mo x O 8 bulk ceramics prepared by SPS at various sintering temperatures. It was found that the relative density of ZrW 2−x Mo x O 8 bulk ceramics increased with the rise in sintering temperature, which is consistent with the results of the Z-axis linear shrinkage curve. The relative densities of cubic-ZrW 2−x Mo x O 8 bulk ceramics were 96.5% and 91.1% for the x = 0.0 sample sintered at 923 K and 873 K, respectively; 95.3% and 93.6% for the x = 0.5 sample sintered at 873 K and 823 K respectively; 96.3% and 91.2% for the x = 1.0 sample sintered at 823 K and 773 K respectively; and near 90% for the x = 1.5 sample sintered at 743 K. This observation implies that SPS can densely fabricate the hard-to-sinter ceramics. On the other side, the relative density of x = 1.5 sintered at 723 K and x = 2.0 samples reached approximately 80%, which can be attributed to its extremely low sintering temperature.      , while the ZrMo2O8 bulk ceramic was found to contain a small amount of trigonal phase, indicating that the synthesis of a singlephase α-ZrMo2O8 is rather difficult. The absence of some special reflections (221, 310) for x = 1.0 and x = 1.5 samples is related to the occurrence of an order-disorder phase transformation, which shows variation in the space group from the low-temperature phase (space group: P 21 3) to the high-temperature phase (space group: P a 3 . The most important finding of the current study is that low temperature is a requirement for sintering pure dense cubic ZrW2−xMoxO8 bulk ceramics. In general, the fabrication of densified ZrW2O8 polycrystals requires a very high sintering temperature of more than 1373 K using the solid-phase method. The current study revealed that the SPS technique could reduce the sintering temperature required for high densification to 773-923 K. This low-  ) for x = 1.0 and x = 1.5 samples is related to the occurrence of an order-disorder phase transformation, which shows variation in the space group from the low-temperature phase (space group: P 21 3) to the high-temperature phase (space group: P a 3). The most important finding of the current study is that low temperature is a requirement for sintering pure dense cubic ZrW 2−x Mo x O 8 bulk ceramics. In general, the fabrication of densified ZrW 2 O 8 polycrystals requires a very high sintering temperature of more than 1373 K using the solid-phase method. The current study revealed that the SPS technique could reduce the sintering temperature required for high densification to 773-923 K. This low-temperature densification is associated with sintering compression pressure and the SPS pulsed direct current. Figure 9 shows the calculated and theoretical values of the lattice parameters of cubic-ZrW 2−x Mo x O 8 bulk ceramics according to Vegard's law. It can be seen that the lattice parameters of all solid solution intervals decrease linearly with the increase of Mo content. Moreover, this linear decreasing trend of the lattice parameters is in general agreement with Vegard's law, which indicates that the Mo 6+ are a solid solution substituted to the W site. In addition, the lowering of the lattice parameters due to the addition of Mo is expected to induce distortion of the crystal structure or decrease the relaxation area due to lattice volume contraction, which may affect the thermal expansion and phase transition behavior. temperature densification is associated with sintering compression pressure and the SPS pulsed direct current. Figure 9 shows the calculated and theoretical values of the lattice parameters of cubic-ZrW2−xMoxO8 bulk ceramics according to Vegard's law. It can be seen that the lattice parameters of all solid solution intervals decrease linearly with the increase of Mo content. Moreover, this linear decreasing trend of the lattice parameters is in general agreement with Vegard's law, which indicates that the Mo 6+ are a solid solution substituted to the W site. In addition, the lowering of the lattice parameters due to the addition of Mo is expected to induce distortion of the crystal structure or decrease the relaxation area due to lattice volume contraction, which may affect the thermal expansion and phase transition behavior.

Evaluation of NTE Coefficients of Dense Cubic-ZrW2−xMoxO8 Bulk Ceramics
The relative thermal expansion of cubic-ZrW2−xMoxO8 dense ceramics was measured by thermo-mechanical analyzer equipment. Figure 10a,b shows the change in relative thermal expansion as a function of temperature for the samples (x = 0.0, 0.5, 1.0, 1.5) sintered at higher temperatures and lower temperatures, respectively. It was evident that the thermal expansion curves decreased with increasing temperature, which indicates the NTE characteristics of the obtained cubic-ZrW2−xMoxO8 ceramics. A change in the NTE

Evaluation of NTE Coefficients of Dense Cubic-ZrW 2−x Mo x O 8 Bulk Ceramics
The relative thermal expansion of cubic-ZrW 2−x Mo x O 8 dense ceramics was measured by thermo-mechanical analyzer equipment. Figure 10a,b shows the change in relative thermal expansion as a function of temperature for the samples (x = 0.0, 0.5, 1.0, 1.5) sintered at higher temperatures and lower temperatures, respectively. It was evident that the thermal expansion curves decreased with increasing temperature, which indicates the NTE characteristics of the obtained cubic-ZrW 2−x Mo x O 8 ceramics. A change in the NTE coefficient, which corresponds to the phase transformation, is attributed to the rapid shift in the slope of x = 0.0 and x = 0.5 samples. The phase transformation temperature shifts to lower temperatures, from about 440 K (x = 0.0) to 350 K (x = 0.5), and drops to below room temperature when x ≥ 1.0, which is consistent with many previous studies conducted on ZrW 2−x Mo x O 8 powders. The repeated changes between positive and negative of x = 1.5 (sintered at 723 K) thermal expansion curves may be related to the sample with a low relative density. It should also be noted that the Mo substitution does not only cause a decrease phase transformation temperature and changes the NTE coefficients. The NTE coefficien calculated from the change in dimensions of material concerning temperature have be shown in Table 1. For the non-doped ZrW2O8 sample (sintered at 923 K), the NTE coe cients are estimated to be −8.68 × 10 −6 K −1 for the α phase and −4.61 × 10 −6 K −1 for the phase. The measured coefficients have values similar to the values reported for ZrW2 powders in the literature (α-ZrW2O8: −8.8 × 10 −6 K −1 ; β-ZrW2O8: −4.9 × 10 −6 K −1 ). In additio the ZrW2O8 sample (sintered at 873 K) with slightly smaller NTE coefficients (−8.44 × 1 K −1 for α phase and −4.05 × 10 −6 K −1 for β phase) may be related to the pores in the bu ceramic. The majority of cubic-ZrW2−xMoxO8 bulk ceramics exhibit similar behavior. A It should also be noted that the Mo substitution does not only cause a decrease in phase transformation temperature and changes the NTE coefficients. The NTE coefficients calculated from the change in dimensions of material concerning temperature have been shown in Table 1. For the non-doped ZrW 2 O 8 sample (sintered at 923 K), the NTE coefficients are estimated to be −8.68 × 10 −6 K −1 for the α phase and −4.61 × 10 −6 K −1 for the β phase. The measured coefficients have values similar to the values reported for ZrW 2 O 8 powders in the literature (α-ZrW 2 O 8 : −8.8 × 10 −6 K −1 ; β-ZrW 2 O 8 : −4.9 × 10 −6 K −1 ). In addition, the ZrW 2 O 8 sample (sintered at 873 K) with slightly smaller NTE coefficients (−8.44 × 10 −6 K −1 for α phase and −4.05 × 10 −6 K −1 for β phase) may be related to the pores in the bulk ceramic. The majority of cubic-ZrW 2−x Mo x O 8 bulk ceramics exhibit similar behavior. Additionally, when the Mo content rises, the absolute value of the NTE coefficients of the phase gradually decreases. In contrast, the NTE coefficients of the β phase remain almost unchanged with changes in Mo doping and stay around 4.0 ppm·K −1 for all samples. Due to their simple processing, excellent heat conductivity, and electrical conductivity, aluminum matrix and copper matrix composites have been widely used in industrial production. Therefore, it is anticipated that the thermal expansion coefficient of the overall composite can be modified to match the ceramic substrate (CTE: 6-12 ppm) when cubic-ZrW 2−x Mo x O 8 is compounded with copper (CTE: 16.5 ppm·K −1 ) or aluminum (CTE: 23.1 ppm·K −1 ). Therefore, self-heating or heat from external temperature during equipment operation can be avoided, preventing deformation or fracture of the junction between the substrate and the composites.  [38] −4.77 [38] x = 0.6 −7.82 (873 K) −4.60 (873 K) −5.70 (823 K) −4.01 (823 K) -−3.5 [5] x = 0.7 -−3.52 (873 K) -−2.91 (823 K) -- x = 1.0 -−4.28 (823 K) -−2.50 (773 K) −9.00 [16] −5.50 [16] x = 1.5 -−2.79 (773 K) ---−3.56 [38]

Conclusions
In conclusion, the ZrW 2−x Mo x O 7 (OH) 2 · 2 H 2 O precursors obtained by the sol-gelhydrothermal technique were employed as the raw material for SPS sintering, and HT-XRD determined the phase change process and temperature interval of various x contents. Under pressurization, the SPS pressure was held at 50 MPa, and the sintering temperature was maintained at x = 0.0, x = 0.1, x = 0.2 at 923 K; x = 0.5, x = 0.6, x = 0.7 at 873 K; x = 1.0 at 823 K; and x = 1.5 at 773 K, which resulted in pure cubic phase ceramics with relative densities above 90%. The bulk ceramics demonstrate NTE characteristics, and the NTE coefficients were tunable in cubic-ZrW 2−x Mo x O 8 bulk ceramics for x value. As a result, cubic-ZrW 2−x Mo x O 8 with customizable NTE coefficients can operate as a thermal expansion inhibitor for composites, allowing them to match any required thermal expansion coefficient to overcome interfacial deformation or even breakage generated by inconsistent thermal expansion coefficients.
Author Contributions: H.W. conducted the overall project and composed the manuscript; J.M. provided the resources; Y.X. and X.Z. provided the investigation and visualization; J.L. provided the methodology; X.X. and Y.Z. provided the validation; X.W. assisted with investigation and data curation; M.L. provided writing-reviewing and editing and supervised. All authors have read and agreed to the published version of the manuscript.