Current status of synthesis and consolidation strategies for thermo-resistant nanoalloys and their general applications

2.1 PS process

This process involves the use of extended periods of sintering the compact powders at very high temperatures (1,500–2,000°C) based on the chemistry of the powder in the absence of any applied pressure. Additives that ensure the densification (Figure 5) and uniform grain structure of the compacts either before the onset of the liquid phase or solid state during the sintering process are also added under a controlled atmosphere (Figure 6). By judiciously selecting the components and raw materials the process aids in achieving the desired densities which are very close to the theoretical estimates [78,79]. Advantages of PS includes lower cost, offers better control over composition, reduced shrinkage of the samples with more retainment of its original dimensions, and improved density and strength of the final product. Whereas, the disadvantages of the process are longer processing time as it relies solely on heat to achieve the desired density, limited to certain materials such as ceramics, and may not work well with more complex materials, the process is more susceptible to defects such as cracks and voids due to heat employed to achieve densification and it is limited to certain shapes and sizes [80].

Figure 5

(a) Schematic of silicon nitride ceramic preparation and (b) densification process of silicon nitride from α → β phase transition. Obtained with permission from ref. [81].

Figure 6

Depicts the PS and reactive PS processes. (a) Step 1: Where a loose bound mixed powder of the refractory precursors is filled into the metal die which is vibrated magnetically to settle up; (b) step 2: the powder and die are sintered in a sintering furnace at a desired temperature to get the final products; the image shows reactions as observed during the reactive PS process of HfB₂–SiC–VSi₂ composite. (c) post 5 h milling shows dispersed particle arrangement; (d) reaction at densification phase; (e) formation of a solid solution of (Hf,W)-B due to tungsten (W) getting inter-substituted in HfB₂ structure; (f) the melting of VSi₂ at 1,700°C; and (g) the obtained composite (HfB₂–SiC–VSi₂) with its final microstructure, post its sintering at 2,150°C. The images (a and b) are reprinted with permission from OpenLearn content, Copyright holder: The Open University© [78]. Images (c–g) are reprinted with permission from Ghadami et al. [79], Springer Nature.

Luo et al. studied the effects of different loadings of MgO–LiF as a sintering additive on the thermo-mechanical properties of pressureless sintered Si₃N₄ ceramics obtained at 1,620°C. The ceramics had a relative density of 92% when 7 wt% MgO and 3 wt% LiF were used as additives. The thermo-mechanical properties exhibited by the ceramic included a high bending strength of 556.73 (±10.18) MPa, hardness of 9.07 (±0.07) GPa, fracture toughness of 6.02 (±0.29) MPa/m^1/2, and thermal conductivity of 47.59 (±0.03) W/(m K). The improvement in the thermo-mechanical properties is attributed to the phase transition from the α to β phase due to the addition of LiF to Si₃N₄ ceramics. The dissolution and re-precipitation of the liquid phase α-Si₃N₄ is initiated along the Ostwald ripening process due to reaching the saturation point in the dissolved state. In any given scenario, the viscosity of the liquid phase is crucial in the materials phase transition. The lower the viscosity, the more the mass transfer contributes to the phase transition of the material. The formation of a silica phase in the compact due to supportive sintering additives also helps to decrease the viscosity of the material [81].

Allemand et al. compared two processes to sinter hexagonal barium aluminosilicate (BAS) monoliths using SPS and a novel ultra-fast pressureless sintering (UFPS) technology. The monolith was sintered between 1,590°C and 1,760°C, which is the thermal stability range of the material. The researchers found that the SPS process could only fabricate small portions of the BAS monoliths due to the manifestation of a thermal gradient that resulted in the partial melting of the monoliths. However, UFPS was able to process and fabricate the complex and larger shapes of the monoliths with equal efficiency and offers improved temperature control during the sintering process concerning SPS. A 1:2:2 molar mixture of powders including BaCO₃, Al(OH)₃, and SiO₂ quartz was used to synthesize BAS monoliths, respectively. All the powders were mixed and homogenized ultrasonically in pure ethanol at pH = 11 for 4 min and later dried for 2 h at 120°C. A portion of the powdered mixture was sintered by the SPS process, where it was kept in a graphite pressing (mold) device with two pistons and a single matrix. Initially, a BN₃ layer was coated onto the carbon sheet (Papyex^®) within the mold, and this coating aided in protecting the powdered oxide samples. The SPS functions under the temperature control (optical pyrometer) centered on a small orifice on the graphite mold surface. The graphite mold was encapsulated with a graphite mat (felt) to mitigate the thermal gradient and any heat loss that may have occurred inside the sample. The other portion of the powdered sample was sintered using the UFPS process, which is very similar to SPS, but excludes any electrical current passing through the graphite die or sample. The obtained green body was thereafter fabricated from the powders using compaction at 100 MPa within the mold. Allemand et al. employed three step based thermal cycles for their samples using both SPS and UFPS. Initially, the researchers maintained a heating ramp rate of 100°C/min to reach a final temperature of 900°C, which was maintained for 1 min; the dehydration process starts at 900°C by converting Al(OH)₃ into alumina (Al₂O₃). Second, the increase in temperature from 900 to 1,150°C at a heating ramp rate of 80°C/min and a holding time of 10 min enables the formation of intermediate species such as barium aluminate/silicate. Third, the final step involves increasing the temperature to 1,600°C at a ramp rate of 100°C/min and a holding time of 10 min; it is at this stage that the BAS forms. UFPS does not involve any pressure sintering procedure. Whereas, SPS involves the use of a high pressure of 40 MPa under vacuum conditions in all situations and for all the steps [82]. Refining the grain microstructure of the as-sintered material from an ultra to nanocrystalline size with decreased porosity and uniform homogeneity is highly desired. The nanocrystalline size enhances the reliability of the material and improves the properties of the material, including malleability. There are some processes, such as doping or field assisted sintering that inculcate increased thermo-mechanical properties with good purity and shape molding options for specific applications such as 3D printing and sputtering targets. It is challenging to yield pure phased elemental refractories using the solid-state PS that is usually performed at temperatures higher than 1,500°C. The higher temperature induces a capillary driving force during the sintering process that results in a synchronal grain growth in size and coarseness. This is a major problem, especially when producing nano-powdered refractories. PS, in general, involves a two-step sintering process: initially, the green body is compacted and heated to a certain high temperature (T ₁) without the final temperature being maintained for any period. This causes the relative density ρ to be approximately greater than 70–80%, and the shrinkage of the relative pore size, ensuring that the channels alongside the grain size are at a high thermodynamically unstable mark. Later, the green body is cooled to a lower temperature T ₂ (e.g., 100–200°C lower than T ₁) and is maintained at this temperature for a period so that the full density is achieved. The second step allows sintering at T ₂ that helps in densifying and inhibiting the grain growth. A completely suspended grain size can only be achieved using best practices. This process has yielded a beneficial Lifshitz–Slyozov–Wagner–Hillert grain size uniformity with a homogenized micro to nano-structure and has successfully constrained the grain size of yttria to ∼30 nm; additionally, the yttria has enhanced properties [83]. Dehghani et al. have recently reviewed the effects on the properties of introducing AlN–Y₂O₃ into pressureless sintered SiC ceramics [84]. Nguyen et al. explored the anatomical and morphological structure of porous ZrB₂–SiC–AlN composites that were fabricated using PS at 1,900°C, using electron microscopy [85]. Li et al. synthesized dense ultrafine-grained W using a two-step pressureless sintered process and achieved a 99.6% theoretical density at 1,600°C when their samples were sintered for 3 h. The researchers also observed no changes in texture, second-phase impurities, no abnormal grain growth, or grain boundary (GB) liquids even after cooling the samples [86]. Jafari et al. studied the effects and behavior properties of ZrB₂–SiC nanocomposites by introducing HfB₂ into the microstructures produced via PS for 1 h at 2,100 and 2,150°C using MoSi₂ and B₄C as sintering catalysts, respectively. With the increasing HfB₂ content, there was an apparent reduction in porosities. Moreover, the presence of MoSi₂ engendered a partial molten phase during the sintering process that markedly improved HfB₂ diffusion into ZrB₂. The apparent change and reduction in the porosity and density occurred in the B₄C catalyzed samples. When the samples were compared with micron-sized SiC and SiC _n (nano) the researchers found that the SiC _n were much smaller and finer than the SiC [87]. Similarly various high-temperature materials such as silicon nitride–barium aluminosilicate (Si₃N₄/BAS) [88], nano (Ca,Sr,Ba)ZrO₃ [89], nano Ti₃SiC₂ interlayer phased SiC [90], Gd₂O₃/Yb₂O₃ infused O'- SiAlON/Si₃N₄ [91], Al₂O₃–mullite–ZrO₂–SiC [92], calcium magnesium aluminosilicate/Yb₂SiO₅–YSZ [93], nitride-based calcium hexaaluminate (CaAl₁₂O_(19–1.5x)N _x ) [94], nano TiB₂–CrB₂–WB₂ [95], ZrB₂ ceramics co-doped with graphite and TiC [96], liquid phase sintered B₄C infused SiC–Al₂O₃–Y₂O₃ composite systems [97], B₄C–Co and ZrC–Co cermets [98], ZrB₂–MoSi₂ [99], MgF₂ + Y₂O₃ + MgO/β-silicon aluminum oxynitride (SiAlON) ceramics [100], (Ni, Co, WC)/TiB₂ [101], α-β-YbF₃-Si₃N₄ [102], ZrB₂–TiB [103], ZrB₂–SiC–TiB₂–B₄C [104], Ta_0.8Hf_0.2C/MoSi₂ [105], α-SiC/(Al₂O₃, Lu₂O₃, Er₂O₃, and CeO₂) [106], AlTiB/diopside:(MgCa(SiO₃)₂ [107], Al₂O₃/barium-magnesium aluminosilicate [108], B₄C–TiB₂/TiO₂ [109], TiB₂–40 wt% TiN [110], Ni _x Al _y TiC [111], Zn/Ce_0.8Gd_0.2O_1.9 [112], and HfB₂–SiC–VSi₂ [79] have been produced and processed by PS. This list is not comprehensive.

2.2 RS

This process relies on the in situ synthesis of UHTC via a chemical-based reaction, where a hard, compact powder sample is converted into a UHTC material either by a solid-state reaction process or by vapor-phase deposition. In this way, the material obtained has a constrained density gradient. One of the benefits of RS is that it can lead to materials with enhanced mechanical properties like strength, hardness, and wear resistance, due to the high material density achieved. Additionally, the technique allows for the production of high purity materials with minimal impurities, as the reaction can take place in a controlled environment. However, RS also has its drawbacks, which must be considered. This process can produce complex shapes and intricate parts, as the starting materials can be shaped before the reaction occurs. Furthermore, this method of manufacturing is cost-effective, and results in high-quality materials with minimal waste [113]. Despite its advantages, RS also comes with certain drawbacks. The process necessitates high temperatures, which can be energy-intensive and may even cause material degradation. Additionally, it involves a long processing time and may have limited material options available for forming the desired product. Finally, there is limited control over the reaction, which can lead to the production of unwanted byproducts or incomplete reactions. Figure 7 depicts a general process occurring during the RS process.

Figure 7

Refractory SiC produced at high temperatures of 1426.85 °C (1700K) by reaction sintering. Reprinted with permission for ref Tsuno et al. (2004) [114] from SPIE Journal Publishers under a Creative Commons Attribution 4.0 Unported License (CC BY).

Tsuno et al. developed a novel, highly robust nano SiC using the RS method. The nano SiC was a lightweight, attractive optical mirror with more than twice the bending strength in comparison to other SiC materials. The SiC produced by the RS method is polished by a diamond slurry that is devoid of pores and is extremely compatible with the visible and infrared region even excluding the chemical vapor deposition (CVD) based expensive SiC coating. The material exhibited a mean bending strength of 854 MPa, Young’s modulus of 400 GPa, fracture toughness of 3.3 MPa m^1/2, CTE of 3.9 × 10⁻⁶/K (RT-799.85°C), specific heat capacity of 680 J/kg/K, and thermal conductivity of 130 (W/m/K). They state that the material is well suited for use in the manufacture of large-scale objects due to the novel fabrication process that ensures very little shrinkage and involves a lower temperature. The material was synthesized by mixing SiC and C and was spray dried; it was later cold pressed to form a green body. The green body was thereafter reaction sintered at 1426.85°C (1,700 K) under extreme vacuum conditions during which the molten Si reacted with the C in the green body to yield the desired SiC. Some traces of Si remained. The product was then mechanically finished to the desired shape and used in optical applications for astronomy [114]. Xiaoju et al. synthesized a SiC/B₄C composite using Si infiltration in the RS method which exhibited excellent mechanical attributes. The dynamic strength results of the SiC/B₄C composite shows a peak value of 1,000/s and its strain continuously increased with the increase in the strain rate. The authors observed three deformation regions in the composite which, including inelastic deformation, failure rapid loading regions during the analysis for dynamic loading [115]. Luo et al. synthesized single phased Ba₂Ti₉O₂₀ (BTO) using the RS method with TiO₂ and BaCO₃ as precursor materials after sintering for 10 h at 1,150°C. The overall reaction activation energy calculated was approximately 386.17 kJ/mol. The researchers observed that the formation sequence of the phases related to the RS process followed the sequence described in Eq. (1).

The composite had pores on the BTO grains because of the sintering which may be due to the absence of oxygen vacancies and elements. However, the surface of the ceramic surface was very rough, and the texture is attributed to the heterogeneous distribution of Ba on the exterior surface of the BTO grain which contributes to the Ti enrichment. In a typical process, the pure powders of TiO₂ and BaCO₃ were ball milled in distilled water using ZrO₂ grinding balls for 6 h to form a slurry which was dried at 60°C. The dried slurry was sieved and mixed with a binder solution of polypropylene glycol. The mixture was not calcined, but it was pulverized, mixed, and dried. The obtained uniform dried powders were subsequently compacted at 200 MPa to form a green body which was thereafter debinded for 2 h at 600°C and later sintered for different soaking times in the temperature range of 900–1,150°C. The optimized microwave dielectric property of the BTO that was treated for 10 h at 1,150°C was ε _r = 34.5, with a Q × f of 11,400 (at 5.7 GHz) and τ _f of 2.89 ppm [116]. Likewise, several researchers have synthesized other combinations of high-temperature materials using the RS method for example: Zr_1−x Sn _x O₂ reinforced alumina-mullite refractory [48], nano-powdered BaZr_0.8Y_0.2O_3−δ [117], (Ti,Zr)B₂–(Zr,Ti)C [118], mullite-Al₂TiO₅ [119], W-ZrC [120], and nano sized MgO–CaZrO₃–β-Ca₂SiO₄ [121].

2.3 SPS

This process involves the flow of an alternative electric current, through the material while the material is being hot pressed. This process is relatively new in the present era, and it allows for the faster sintering of a material. However, this process along with hot pressing (HP) can produce only a defined cast of simple shapes, which would necessarily involve the use of another costly machining process for the manufacture of intricate engineering components. SPS, similar to PS and RS techniques, has several advantages such as the ability to create high-density materials, shorter processing time, cost savings, and uniform heating throughout the sample. Furthermore, SPS allows for the development of more uniform microstructures, improved properties, and greater control over grain size. It is applicable to various materials, including metals, ceramics, and composites. However, the SPS method has a few drawbacks, such as high equipment costs, limited sample size, scalability challenges, and process control difficulties [80,122]. Paul et al. (2021) synthesized ZrB₂–MoSi₂ composites and studied the effect of adding silicon carbide whiskers (SiC_w) to the composite to understand the wear performance. The composite ZrB₂–MoSi₂–SiC_w (ZMS_w) was produced using the SPS method during which the powders of pure raw ZrB₂, MoSi₂, and SiC_w were mixed in acetone at different stoichiometric ratios in a planetary mono-ball mill using WC grinding balls for 6 h at ∼250 rpm. The milled powders were heated for 1 h at 300°C to remove the acetone and were later dried and ground to a fine consistency and packed in a graphite foil lined graphite die for sintering using the SP process with multi-step sintering in the range of 1,400–1,700°C under a vacuum. A uniaxial pressure of 60 MPa was applied throughout the sintering cycle. The SPS heating rate was 100°C/min. The tribological results indicated that the self-lubricating tribo-oxide layer that formed during the frictional heating significantly improved the composite resistance. The operational wear maps show that the tolerant workable load limit for the ZMS_w lies between 20–40 N [123]. Figure 8 depicts the general process occurring during the SPS process and the differences between the conventional sintering and SPS process following the different process stages and the passage of the current in the nano powders during the reaction [124].

Figure 8

(a) A generalized spark plasma process scheme; comparability between (b) SPS and (c) conventional sintering methods; (d) graph depicting different stages of a typical SPS process; and (e) the passage of direct current (DC) pulse through the powdered particles during an SPS reaction. Reprinted with permission from ref. [124] from Hindawi.

Chakravarty et al. (2021) fabricated highly dense, super strength nano W–TaC–Ta₂O₅ (similar to the W-MC system, where M = Ta, Zr, Hf, Ti) by the SPS method. The fabrication was performed under high vacuum conditions (3 × 10⁻⁴ mbar) and in the temperature range of 1,773–2,273 K (1,500–2,000°C) under an applied stress of 25–100 MPa. The synthesis involved mixing the raw powders of W and TaC under an Ar atmosphere in a high energy ball mill using WC balls for 2 h at a speed of 250 rpm. Later, the milled mixture was reduced in H₂ for 2 h at 1,500°C and thereafter subjected to the SPS process [125]. Azevêdo et al. used the SPS method to synthesize nano phased (300–200 nm sized) Al₂O₃–WC–Co composites for abrasive machining application. The raw materials (Al₂O₃, WC, and Co powders) were initially mixed in a high energy mill. Subsequently, they were wet milled in the presence of ethanol, using carbide balls for 50 h at 450 rpm. The slurry was thereafter spread to dry on Al foil to remove ethanol. After the mixture was dried, it was placed in a graphite die and SPS was performed at 1,550°C at a heating ramp rate of 65°C/min under 40 MPa of pressure and 100 Pa of vacuum, with a residence time of 5 min. The authors found that the apparent porosity of the composite sample was 1.15% with 26.41 GPa of micro-hardness, 4.51 g/cm³ of apparent density, a friction coefficient of 0.54, and wear rate of 1.978 × 10⁻⁶ mm³/Nm [126]. Feng et al. prepared dense (Hf, Zr, Ti, Ta, Nb)B₂ high-entropy nano sized (300–500 nm) ceramics using the SPS method, during which they used stoichiometrically weighed powders of HfO₂, Ta₂O₅, TiO₂, ZrO₂, Nb₂O₅, B₄C, and carbon black. Three different formulations were performed for the synthesis of high entropy boride (HEB) ceramics using different quantities of Nb. Oxides were clubbed with an excess of B₄C and carbon. The powdered mixtures of HEB were synthesized initially for 3 h under a moderate vacuum (∼3 Pa) at 1,650°C using the boro/carbothermal reduction process. Subsequently, the HEB ceramics were subjected to a two-step SPS method under ∼2 Pa of vacuum and at the highest temperature peak of 2,100°C. During the SPS process, an isothermal hold for 5 min, with a 15 MPa uniaxial pressure at 1,650°C was applied to remove the impurities which occurred in the form of remnant oxide [127]. Ensuring the certainty of the spatial uniformity of the sample properties is a major challenge usually encountered in an SPS process. Several studies have also shown that some differences and a non-uniform temperature gradient/distribution occur during SPS tooling. The non-uniform temperature gradient is observed between the average temperature of the sample and the outer surface of the die that is in contact with the thermocouple of the unit. It is primarily attributed to the non-uniform thermo-electric properties of the SPS tooling. The improvement in the uniformity of the temperature may be achieved via optimization. However, the application of SPS in sintering thermo-electric samples has uncovered an additional heat source and temperature gradient. Another point to consider is the loss or generation of the heat during the process by the Peltier effect. In certain instances, e.g., with samples such as nano sized boron carbide, it was found that when the sample was sintered at high temperatures of 2,000 K, a significant temperature gradient was observed between the top and bottom sides of the sample. The temperature gradient causes severe loss and heterogeneity in its microstructure and mechanical properties. The contact surface of the specimen, graphite tooling, induction, and pulse frequency, also influence (or contribute to) a potent Peltier effect. A p-type conductor such as boron carbide and an n-type conductor such as graphite experiences punch-through electron–hole recombination over the top surface of the specimen, thus generating an electron–hole pair at the bottom surface during the SPS process. This phenomenon causes the cooling of the bottom contacts and the heating of the top layer. Researchers have observed a decrease in polarity and thermal changes when AC is used instead of DC electric fields. Due to the Peltier effect, the nano boron carbide demonstrates erratic melting behavior. When the innate inhomogeneity is easily noticeable along with a lower relative density of the sample, the driving force and time eventually contribute to an increase in density by 100%. Additionally, the irregular grain size results from the non-uniform relative density, which can be avoided by introducing dielectric plates (such as alumina) between the sample and graphite punches to optimize the quality of the properties of the final product by removing the thermo-electric effects [128]. Nayebi et al. synthesized TiB₂–Ti₃AlC₂ ceramic composites at 1,900°C for 7 min using the SPS method at 30 MPa biaxial pressure [129]. Buinevich et al. fabricated ball milled and combustion synthesized non-stoichiometric nano structured hafnium carbonitrides (HfC _x N _y ) which were further consolidated using the SPS method at a temperature of 2,000°C for 10 min under a constant pressure of 50 MPa [130]. Park et al. synthesized and studied the behavior of nano sized TaNbHfZrTi high-entropy alloy composites by the hydrogenation–dehydrogenation reaction at 1,100°C using the SPS method [131]. Jiang et al. studied the sintering behavior of the Ce₂Zr₃(MoO₄)₉ ceramics processed by the RS method and analyzed their vibrational spectra and microwave dielectric properties [132]. Sun et al. synthesized ball milled CaZrTi₂O₇-zirconolite based ceramics that were further consolidated by reactive SPS (Figure 9) for plutonium disposition application. The SPS synthesis was performed in the range of 900–1,450°C from 0.1–10 h, between 15 and 50 MPa of uniaxial pressure along with a solid-state reaction [133]. Rominiyi et al. studied the oxidation, microstructure, and tribological behavior of nano Ti–Ni_−x TiCN ceramic composites by SPS [134].

Figure 9

A generalized reaction spark plasma process scheme for the synthesis of Zirconolite and the obtained crystallinity and microstructure as reported by Sun et al. [133]. Reprinted with permission from Elsevier.

2.4 HP and HIP

This process involves the application of a high-grade temperature (up to 2,200°C) and uniform directional pressure (200–500 MPa) on the material, simultaneously to elevate the densification and sintering of the UHTC material at a temperature that is much lower than the surrounding temperature. Due to the fast rate of the densification process, the grain size in the material is always maintained which improves the thermo-mechanical and thermal properties of the UHTC material. The process is commonly used in the production of ceramics, metals, and composite materials. Here are some advantages and disadvantages of HP: HP produces materials with high density, uniformity, reduced porosity, improved strength, and grain alignment. Few drawbacks of the process include high equipment and manufacturing cost and the process requires a significant amount of energy to heat the material. It also has limited shape complexity and material compatibility as it may not be suitable for all materials due to high pressure and temperature of the process. The process is also time-consuming which can limit its use in certain manufacturing applications. On the other hand, HIP involves utilizing high temperatures and pressures to condense metal or ceramic powders into solid materials. The advantages of HIP comprise better material properties, improved surface finish, customizable shapes, material variety, and reduced waste. Nevertheless, HIP has limitations such as high cost, time consumption, safety hazards to operators, as well as size and material thickness restrictions [135,136,137]. Kang et al., prepared an in situ micro and nano scaled W steel first by selective laser melting (SLM) followed by the HIP. The HIP process has a maximum limit of pressure (207 MPa) and temperature (2,000°C). The HIP was performed completely in a pure Ar atmosphere at 950°C with a cooling/heating rate of 10°C/min, holding time of 120 min, and at a pressure of 150 MPa. During the post HIP process, it was found that the SLM samples exhibited a higher density without any change. The sample finally demonstrated a high yield strength of 1,025 MPa with an ultimate tensile strength between 980 and 1,470 MPa [138]. Sun et al. prepared a WTaMoNb/Cu nano composite with a stable microstructure using dual-step mechanical alloying and HP sintering [139]. Figure 10 shows the schematic illustrations of the HP and HIP processes, a step wise example for the HP synthesis of micro/nano SiC, and the basic difference between the HP and HIP processes [140,141,136,142].

Figure 10

(a) Generalized schematics of HP by Moustafa et al. [140]; schematic diagrams demonstrating (b) sintering process for the synthesis of SiC ceramic by HP method and (c) the heating curve designed for the process by Li et al. [141]; (d) conventional HIP method [136]; schematic comparison between (e) HP and (f) HIP processes by Dobrzański et al. [142]. Reprinted with permission from (a) Scientific Research Publishers; (b and c) from Elsevier; (d) from the AZoNetwork; and (e and f) from IntechOpen Limited, United Kingdom.

Lozanov et al. prepared Ir-based ultrahigh-temperature intermetallic composites (MIr_3±x , where M = Hf, Ta) using the HP process. The powders were mixed and filled in a graphite die maintained at ∼1 kPa in an Ar atmosphere at a pressure 110–130 kPa. The dies were pressed uniaxially at 20 MPa followed by heating at 2,073 K at a rate of 100 K/min with 20 min of exposure time at the peak temperature. Later, the chamber was cooled for 1 h to reach temperatures between 323 and 343 K, respectively. The obtained samples demonstrated a thermal diffusivity between 11 and 13 mm²/c, thermal conductivity of approximately 39.4 W/m K at 1,700 K, heat capacity (C _p) of approximately 102.3 J/mol K at T = 298 K, CTE of approximately 5.2 × 10⁻⁶ K⁻¹ between the temperatures of 300 and 1,100 K [74]. It is acknowledged that the preparation process of SiC heavily influences the chemistry and microstructure of the overall product thereby affecting the properties of the final product. Taking this factor into consideration, Šajgalík et al. prepared an additive-free ultra-high creep resistant SiC ceramic using rapid hot pressing. The researchers suggested that the β–α SiC phase transformation and the GB sliding controlled by GB diffusion are responsible for the creep mechanism at higher temperatures of 1,500–1,750°C and at compressive loads of 200–400 MPa [143].

Potanin et al. synthesized micro and nano sized MoSi₂–ZrB₂–SiC and MoSi₂–HfB₂–SiC composites using the combined synthesis of ball milling, combustion synthesis, and hot pressing. According to the researchers, the obtained composites have the potential to be used as high-temperature materials for industrial heaters. Powders of Mo, Hf, Zr, Si, amorphous B, and amorphous carbon black were mixed in the desired proportions and were mechanically activated in a planetary ball mill for 5 min at a speed of 700 rpm. The mixtures underwent combustion synthesis (self-propagating high-temperature synthesis; SHS) under an Ar atmosphere at 2 atm of pressure. The powders were subsequently ball milled. Thereafter, the powders were consolidated using HP under a pressure of 35 MPa at 1,600°C at a heating rate of 10°C/min for 10 min of isothermal exposure in a vacuum to yield the final composite products. The obtained composites had a 97–98% relative density and an unreactive continuous oxide layer of crystalline α-quartz (a high-temperature alteration of SiO₂) formed which was strengthened with HfO₂, ZrSiO₄, ZrO₂, and HfSiO₄ precipitates. The composites had 2–3% of residual porosity, hardness of 12–14 GPa, strength of 350–400 MPa, and fracture toughness of 6.5–7.1 MPa m^1/2. The composites exhibited some weight gain when tested for their oxidation mechanism for 4 h at 1,650°C which is a characteristic of the in situ development of a dense protective layer [4].

Table 3 lists the modulus properties of the high-temperature silicide based nano refractory materials [51,60]. Vorotilo et al. synthesized both nano and micron sized ZrB₂–TaB₂–TaSi₂ mixed composite ceramics using both the SHS and HP methods. The powders of Zr, B, Ta, and Si were milled and later subjected to the SHS process and milled again to obtain a fine powder which was further consolidated by an HP process at 1,700°C for 10 min of exposure time under a pressure of 30 MPa. The obtained composite ceramic yielded a relative density of 95–98%, hardness (HV₁₀) of 19.2 GPa, and crack resistance (K _1c) of 3.5 MPa m^1/2 [57]. The HIP process, in contrast, has an advantage over the SHS process because the samples are sintered at low temperatures for a short holding time. During consolidation, this prepares defect free sub-micro/nano structured powdered materials with minimized grain growth. However, when the stress state is that of HIP conditions and a material has a lower shear viscosity, extremely small grain size, and a flattened pore, then the pore deforms, shrinks anisotropically, and finally disappears during the holding time. The pressure driven sintering usually involves several mechanisms including plastic yielding, and diffusional and power-law creep. When the stress levels are greater than the yield stress, there is a probability of macroscopic deformation, and the stress potential determines the power-law creep for the obtained porous material. However, when the stress levels are low, the diffusional creep becomes the prevailing deformation mechanism. Here the macroscopic strain rate is linear relative to the hydrostatic stress and thermodynamic operating force for pore shrinkage, which is the sintering stress. To predict shrinkage in the final product, the constitutive equations help to determine the densification mechanisms. Notably, during the densification process, the use of pressures higher than 100 MPa for HIP sintering becomes mandatory and it is crucial to completely remove the material defects such as cracks, large grain sizes, and growth, respectively. For decreasing the large pore size/voids during the densification process with HIP, it is found that the use of very high pressures of more than 100 MPa sometimes do not alter or shrink the pore size because the time it takes to reach the reference density decreases, and it is inversely proportional to pressure. This renders the time dimensionless and excludes its consideration, thus, high pressures may be ineffective at reducing the large pore size during the HIP process [144]. Kaplanskii et al. synthesized and optimized a nano phased Ni₄₁Al₄₁Cr₁₂Co₆ alloy powder using plasma-spheroidization in a plasma torch and layered the alloy on the turbine rotor blade models using laser synthesis (laser powder bed fusion-LPBF). The researchers investigated the HIP effect on the LPBF parts for their structural and thermo-mechanical properties. The LPBF parts were initially aged to ensure that the reinforced agents were well precipitated and well dispersed on the layers and have relaxed residual stress. Later the thermal treatment at 1,423 K was performed in a vacuum resistance furnace equipped with a W heating element under a vacuum of 1.3 × 10⁻⁵ Pa for 3 h of constant thermal exposure. Post thermal treatment, the blades were partially subjected to an HIP process at 145 MPa of pressure and at 1,523 K to decrease the porosity and remove the structural and property-based anisotropy. It was observed that HIP improved and significantly decreased plastic deformation, and the mechanical strength of the rotor blade at 85 K, developed by the combination of the LPBF, aging, and HIP process, was greater than that of the other fabricated (as-HIP and as-LPBF + aging) blades [145]. Lv et al. synthesized and fabricated a new CrMoNbWTi-C refractory high-entropy alloy composite which was co-strengthened using fine-grained intermetallics and UHTC carbides by mechanical alloying and the HP method between 1,400 and 1,450°C. The composite demonstrated a good elastic modulus (247.95 GPa), compressive fracture strength (3,094 MPa), and hardness (8.26 GPa) [146]. Recently, Herrmann and Räthel described the basic process and applications of HP and HIP in great detail [147].

2.5 Plasma/Physical vapor deposition (PVD) and Chemical vapor deposition (CVD)

The deposition techniques can efficiently coat good quantities of the UHTC material precisely onto the substrate or on any of the surfaces of the components. The plasma coating generally involves melting the precursor material in gas plasma before the material is transferred to and deposited under high velocity onto the targeted substrate, where the coating cools and solidifies. However, this process (plasma spraying) generates large amounts of defects in its coatings which could improve the thermal insulation resistance even when there are immediate changes in the temperature gradient. Additionally, the CVD process utilizes the benefit of the chemical reactions involving the gas and precursor materials to deposit the coating on the substrate directly (Figure 11) [148,149,150,151,152]. The coating material obtained by this process is always dense, flawless, and corrosion and oxidation resistant in comparison with the spray coating. In comparison, few merits and demerits are often found with the PVD/CVD, even though it is one of the preferred methods used for producing UHTCs which results in remarkable thermal stability and mechanical properties, making them suitable for high-temperature applications like aerospace and nuclear industries. PVD/CVD provides UHTC materials with high purity, customizable composition, coating uniformity, and improved wear and corrosion resistance. However, the process is complex, costly, and requires skilled personnel and specialized equipment. The production capacity of PVD/CVD is limited due to the slow deposition rate, and the process can produce a rough surface finish, which may require additional processing steps to achieve the desired quality. Safety hazards such as high temperatures, vacuum environment, and hazardous chemicals must be mitigated with appropriate safety protocols and equipment [153]. He et al. synthesized high-temperature carbon/carbon (C/C) composites with SiC/pyrographitic carbon (PyC) core–shell structured nanowires using chemical liquid-vapor deposition (CLVD). The researchers used polycarbosilane (PCS) and SiC as the precursors and the preforms were made by punching a 2D needle into the carbon felt. The SiC nanowires and preforms were immersed in liquid xylene/PCS solution and pyrolyzed in the CLVD system at 950°C for 6 h, followed by heat treatment in Ar at a temperature range of 1,350–1,900°C for 2 h at a 10°C/min heating rate to yield the composites. The SiC/PyC core–shell structured nanowire composites were prepared by soaking the as-prepared SiC nanowire composites in xylene and then deposited using CLVD for 2 h at 950°C. At this stage the xylene is pyrolyzed into PyC and the PyC is gradually deposited onto the surface of the SiC nanowires leading to the formation of SiC/PyC core–shell structure nanowire composites [154].

Figure 11

A generalized schematic of CVD process adapted and modified from Dhand et al. [149,150,151,152].

Ren et al. fabricated a C/C bi-phase coated with a HfC–ZrC composite using CVD to investigate the anti-ablative and anti-oxidative properties. The raw materials used for the synthesis were the C/C substrate and SiC sandpaper which were ultrasonically cleaned and dried, HfCl₄, ZrCl₄, CH₄, H₂, and Ar. The deposition was performed in a vertical tube furnace. The HfCl₄ and ZrCl₄ powders were initially mixed by ball milling, and loaded into the sublimation chamber and the C/C substrate was loaded into the reaction chamber. The deposition was performed between 1,200 and 1,350°C at 3 kPa in an inert Ar atmosphere under a gas flow for 2 h. Later, CH₄ and H₂ were cut down and cooled in Ar ambience to yield the C/C bi-phase coated HfC–ZrC composite. The following equations express the possible chemical reactions occurring during the deposition process (Eqs. (2)–(6)).

During deposition:

The composites were evaluated for their anti-ablation property for 60 s by oxyacetylene ablation, where it was discovered that the HfO₂ and ZrO₂ played an important role and there were sufficient grains on the surface for the formation of a strong oxide layer, which helped to ensure negligible oxygen diffusion. Eqs. (7)–(12) express the reactions occurring during the ablation process.

During ablation:

After the test it was found that the linear ablation rate (R _l) and mass ablation rate (R _m) of the composite coating was −0.16 µm/s and 0.34 mg/s, respectively [155]. Temperature plays an important role in the deposition and microstructure control of the coating.

1. The surface reaction controls the coating deposition at lower temperatures and crystal nucleation dominates during the deposition process that results in a smaller grain size of the coated material.

2. At high temperatures, the gas transmission controls the coating and deposition process, where the large grain crystal growth predominates in the coated material. With the increase in the surface temperature, the release of gases such as CO/CO₂, saturated gas vapor pressure, followed by gas vaporization, and the production of voids within the oxide layer occurs. Due to a higher CTE of the carbide coating and oxide layers, excessive thermal stress and shock cause the coating to crack. Thus, the crack channels the diffusion of the oxidizing gas leading to rapid substrate and coating oxidation. During the post ablative stage, the rate of oxidation of the coating and the oxygen diffusion becomes greater due to the increase in the size of the cracks/voids, respectively. Additionally, this causes the coated oxide or melt material to lose its weight and thickness due to the high pressure/speed and mechanical abrasion inflicted by the ablation torch [142,143,144,145,146,147,148,149,150,151,152,154,155]. Tanaka et al. synthesized a highly ordered nanocomposite of cubic (B1)-TaC(111) and rhombohedral-Ta₃C₂(0001) phases over the Al₂O₃(0001) substrate using ultra-high vacuum DC magnetron sputtering of a thick TaC target under a vacuum of 5 mTorr in the presence of an Ar/C₂H₄ gas mixture. The reaction/deposition was performed at substrate temperatures of 1,123 and 1,273 K, respectively. The process involves 10 min of cleaning the Al₂O₃(0001) substrate with acetone, isopropanol, and distilled water using ultra-sonication, followed by drying under a stream of N₂ and mounting on a heating stage, and transferring to the deposition chamber after evacuating the chamber below 10⁻⁸ Torr. In the deposition chamber, the sample is thermally degassed at 1,273 K until the base pressure is <6.0 × 10⁻⁹ Torr. When the temperature is set to 1,123 or 1,273 K within the DC magnetron sputtering unit, TaC films are deposited for 2 h on the substrate in the presence of Ar and C₂H₄ gas mixtures at 5 mTorr of vacuum and at a deposition rate of 0.03 nm/s at 50 W of constant target power. The other parameters, such as variable current (0.15–0.16 A) and voltage (297–307 V) range, is followed up for creating TaC film during the reaction. Once the deposition is finished and the desired parameters attained, the heater is switched off, the samples are cooled, and analyzed [156].

Figure 12 depicts the general process of physical vapor deposition (PVD) using different types of magnetron (RF and DC types) sputtering to yield good quality coatings of the nano structured refractory materials onto the targeted substrate surface [157,158,159].

Figure 12

(a) Generalized schematic of physical vapor deposition (PVD) reaction involving the deposition of the liquefied metal source by Yu et al. [157]; schematic PVD diagram demonstrating (b) DC magnetron deposition unit; (c) magnified portions of the magnetron and the substrate holder within the chamber by Avino et al. (2020) [158]; (d) schematic diagram of PVD unit involving radio frequency (RF) based magnetron deposition unit by Kulkarni et al. [159]. Images (a–d): Reprinted with permission from Elsevier.

The PVD technique is useful because a wide variety of substrate materials, including alloys, ceramics, metals, glass, and polymers may be utilized. Additionally, PVD utilizes a very large scope of coating materials including alloys, metals, metal oxides, nitrides, and carbides. The process ensures a high-quality coating adhesion on the substrate with a refined microstructure based on a selection of suitable coating parameters (precursor composition, temperature, deposition rate, and film thickness). PVD is highly useful in opto-micro-electronic devices for superior quality thin film production, improved tribo-performance, anti-corrosion layering, thermal barrier, heat dissipators (sinks), and aesthetic coatings. Among the different types of available PVD methods, thermal evaporation is the most opted method due to its ease of operation and its ability to evaporate large varieties and quantities of materials. The versatility of the method helps to easily achieve the deposition of the desired material in the vapor phase under vacuum conditions (pressure < 5–10 torr) assisted by a resistive heater as the heating source. Thus, the transportation of the vapor atoms toward the low pressure zone helps in the deposition of the material onto the substrate aligned linearly from the source [160]. The formation of the desired fine-structured morphologies is a highly complex phenomenon because there is a contest between the surface transport effects and the depositing vapor flux during the deposition process. Thus, there is a two-fold challenge to optimize the parameters governing the PVD processing to obtain the desired microstructures. The primary challenge is selecting the spatial properties and the deposition conditions for the target to be deposited onto the substrate. The magnetron sputtering deposition parameters include, the gas/target/substrate composition, background pressure, temperature, and sputtering power. Likewise, the important properties for the deposition of the relevant materials are the atomic stoichiometry of the target species and the bulk/surface diffusion onto the substrate [161]. Kameneva et al. investigated improving the physico-mechano-tribological properties, structural and phase transformations by using cathodic arc deposition (Arc-PVD) equipped with magnetron sputtered targets of Ti and Al for the synthesis of a nanostructured Ti_1−x Al _x N coating on a WC–Co substrate. The researchers also studied the effects on the composition and phase concentration of the nano-thick sized coating by controlling the temperature and deposition conditions on the surface of the cathode and substrate during the process [162]. Jokanović et al. fabricated multilayered nano thin films of titanium-oxy-nitride and Cu-doped titanium nitride on a glass substrate using different PVD techniques and investigated its physico-chemical properties [163]. Kumar et al. deposited uniform nano TaC and SiC layers onto a graphite tube using the CVD method. Initially the authors synthesized tantalum chloride (TaCl₅) in situ by the chlorination of Ta chips at 550°C for 30 min. Later, a graphite tube was placed in the CVD chamber and TaC was deposited onto the tube due to the reduction reaction between TaCl₅ and CH₄ in the presence of H₂ gas at a pressure of 50–100 mbar and temperatures between 1,050 and 1,150°C. Similarly, another set of tubes was placed in the CVD chamber and SiC was deposited by introducing methyl-tri-chloro-silane and H₂ gas at 1,000°C and 50 mbar. The coating thickness at 1,150°C was approximately 600 and 400 µm at 1,050°C for the 10 h duration of the additional deposition. The samples were tested for their anti-ablation property for 120 s using an oxyacetylene flame. The results indicate that the coatings were intact without any change in their integrity albeit with a negligible amount of erosion. The analysis revealed the presence of nano aggregates of the surface oxide phases (TaO₂, Ta₂O₅, and SiO₂) of the TaC and SiC coatings, which is believed to have protected the underlying substrate from the high temperature of 2,000°C. The researchers suggested that these uniform CVD coatings on other geometrical objects/propulsion systems, including throat inserts, nozzles, and chambers of rocket motors, has potential which should be fulfilled in the coming years [164]. Zhu et al. performed both thermodynamic calculations and the synthesis of micro and nano mixed phases of ZrB₂ from a ZrCl₄–BCl₃–H₂–Ar system using CVD between 1,050 and 1,600°C at very low pressures [165]. Similarly, in their previous work, Zhu et al. performed an experimental synthesis using CVD accompanied by thermodynamic calculations for the production of a ZrC–SiC bi-phase from the ZrCl₄–C₃H₆–methyltrichlorosilane–H₂–Ar system using the FactSage thermochemical software [166]. Moraes et al. fabricated the sub-stoichiometric α-structured W–Ta–B thin films by PVD equipped with an in-house developed magnetron sputtering system. The researchers observed that the WB₂ thin films demonstrated an α-AlB₂-prototype structure upon crystallization instead of their ω-W₂B_5−z -prototype structure, which is thermodynamically stable. By simulating the experiments using density functional theory (DFT), it was found that the stability of the α-WB₂ thin films is extremely regulated by point defects, such as vacancies, within the target PVD materials. The presence of α-TaB₂ and the phase transformation during the thermal decomposition reaction aids in the transformation of α-W_1−x Ta _x B_2−z into a ω-W₂B_5−z -type structure when the temperatures are above 1,200°C. Experimentally, W_1−x Ta _x B_2−z thin films were synthesized by the sputtering of TaB₂, WB₂, and carbon targets on the ultrasonically pre-cleaned steel foil or Si(100) substrates at a chamber temperature of 700°C (substrate temperature maintained at 400°C) and at a working pressure of 0.4 Pa. Throughout the experiments, the base pressure during the coating process was maintained below 3 × 10⁻⁴ Pa, target power varied between 0 and 11 W/cm² at 0.25 Hz of substrate holder rotation and bias voltage (−50 V) and the Ar etching of the targets in the deposition unit performed in an Ar atmosphere of 6 Pa with an applying voltage potential of −750 V for 10 min [167]. Startt et al. conducted a study to investigate the mechanical and thermal properties of MoNbTaTi quaternary refractory alloy across a broad compositional space using DFT simulations along with experimental validation. The study found that there was good agreement between the simulation results and experimental data, which supports the use of DFT tools to predict properties of complex alloys. The first column of Figure 13 shows the thermodynamic properties calculated as a function of temperature for the equiatomic and experimental Ti-heavy (Ti34.7) composition using DFT under the quasi-harmonic approximations (QHA). The results were compared with the experimentally measured values for the Ti34.7 composition, which were obtained at approximately room temperature. Although there was some deviation between the DFT predictions and the experimental measurements, the modeling results were still quite close to the experimental values, indicating that the predictions of thermodynamic properties through both the harmonic and QHA are reliable [168].

Figure 13

Thermodynamic properties of different compositions of the MoNbTaTi quaternary refractory alloy calculated using the QHA and density functional perturbation theory. The equiatomic composition is represented by a solid black line, while the Ti-heavy (Ti34.7) experimental composition is shown as purple dashed lines and cross points in the first column. Additionally, the thermal properties were calculated for one composition above (indicated by solid-colored lines) and below (represented by dashed-colored lines) the equiatomic composition for each element track. Reprinted from ref. [168], with permission from Elsevier.

In their study, Alvi et al. employed DFT simulations to determine the mechanical properties of the MoNbTaTi quaternary refractory alloy. The simulations revealed a Young’s modulus value of 229 GPa, which agreed well with the experimental nanoindentation measurements. The researchers also observed that the high hardness observed in the metallic film could be attributed to the increased grain boundary strengthening resulting from the presence of nanocrystalline structure [169]. Dan et al. synthesized nano thin films of TiB₂/TiB(N)/Si₃N₄ deposited onto stainless steel (SS) substrates maintained at approximately 40°C by DC/RF magnetron sputtering. The researchers used TiB₂ and Si targets for the synthesis of the thin films. The substrates were chemically pre-cleaned and polished metallo-graphically before deposition. The chamber was evacuated to a base pressure of 5.0 × 10⁻⁴ Pa. In the presence of Ar plasma, the TiB₂ layer developed and in an atmosphere of Ar and N₂, the Si₃N₄ and TiB(N) layers were deposited. The obtained TiB₂/TiB(N)/Si₃N₄ thin film was tested for its photo-thermal conversion activity and the films exhibited 0.18 and 0.964 of emittance (at 82°C), and absorptance in the infrared region of 2.5–25 µm, and 0.3–2.5 µm in the solar spectrum range [170]. Table 4 lists the possible high-temperature refractory coating materials which can be coated on the surface of the C/SiC composites along with their melting points [171,172]. Table 5 lists the summarized melting points of a few inorganic refractory compounds and elemental metals which can be employed as high-temperature materials [171,172].

In this section we observe that the main factor for the deposition synthesis of the micro and nano phased refractory materials and their specific applications is based on the composite fabrication of the refractory matrix. The focus is also on the maintenance of a homogeneous microstructure for enhancing the thermal shock observed during the anti-ablation and oxidation tests. Very few research reports have so far been published regarding the preparation of other multi-functional composites using this method. Nevertheless, the fabrication and formulation of multi-functional composites comprising different mixtures and quantities of additives can produce diverse properties, such as superior thermal protection at higher temperatures, achievable full density, pore distribution, conditioned microstructures, enhanced thermo-physical attributes, and composite strength in the UHTC composites for their specific application. We have also noted that researchers have been modeling and simulating the parameters using FactSage, DFT, and other software to understand the effect on the synthesis, microstructure, and properties during the processing of the material composites. At present, a lot of attention has been focused on the conservation of material architecture for the prospective growth and maturation of the high-end C/C-UHTC composites for a desired set of properties and applications.