Effect of High-Temperature Isothermal Annealing on the Structure and Properties of Multicomponent Compact Ti-Al(Nb,Mo,B)-Based Materials Fabricated via Free SHS-Compression

Abstract

This study investigates TNM-type titanium aluminide alloys, representing the third generation of β-stabilized γ-TiAl heat-resistant materials. The aim of this work is to study the combustion characteristics and to produce compact materials via the free SHS compaction method from initial powder reagents taken in the following ratio (wt%): 51.85Ti–43Al–4Nb–1Mo–0.15B, as well as to determine the effect of high-temperature isothermal annealing at 1000 °C on the structure and properties of the obtained materials. Using free SHS compression (self-propagating high-temperature synthesis), we synthesized compact materials from a 51.85Ti–43Al–4Nb–1Mo–0.15B (wt%) powder blend. Key combustion parameters were optimized to maximize the synthesis temperature, employing a chemical ignition system. The as-fabricated materials exhibit a layered macrostructure with wavy interfaces, aligned parallel to material flow during compression. Post-synthesis isothermal annealing at 1000 °C for 3 h promoted further phase transformations, enhancing mechanical properties including microhardness (up to 7.4 GPa), Young’s modulus (up to 200 GPa) and elastic recovery (up to 31.8%). X-ray powder diffraction, SEM, and EDS analyses confirmed solid-state diffusion as the primary mechanism for element interaction during synthesis and annealing. The developed materials show promise as PVD targets for depositing heat-resistant coatings.

1. Introduction

Modern heat-resistant alloys face increasingly stringent requirements each year, yet currently developed alloys do not fully meet these demands. Considerable research and application efforts are focused on titanium-aluminide-based materials, which offer a high melting point, low density (approximately 9 g/cm³ for nickel superalloys vs. about 4 g/cm³ for γ-TiAl alloys), high elastic moduli, an increasing yield strength (for TiAl) with temperature, oxidation and ignition resistance, and an excellent strength-to-density ratio [1,2,3]. However, their practical use below 700 °C is severely limited by low formability, fracture toughness, and insufficient elongation at fracture [4]. Notably, titanium aluminides undergo intensive wear above 600 °C, making TiAl-based alloys highly susceptible to cracking, reduced fatigue life, and low damage tolerance. Their room-temperature ductility and deformability are exceptionally poor, restricting their application as structural and engineering components. Consequently, enhancing the properties of titanium aluminide alloys for broader practical use remains a pressing and highly relevant challenge.

Research has demonstrated that adding β-stabilizers (Nb, Mo, Cr, V) to TiAl-based alloys significantly improves their processability, mechanical properties, and oxidation resistance [5,6,7]. These elements act by stabilizing the high-temperature β-phase at lower temperatures, which enhances workability and refines the microstructure. Researchers [8] identified an optimal Al/(Al + Ti) ratio (≥50 at%) and a β-stabilizer content (~5 at%) that synergistically enhance oxidation resistance and fracture toughness. This specific composition window promotes the formation of a protective Al₂O₃ scale while maintaining adequate microstructural stability. Known examples include TNB alloys such as Ti-46Al-8Nb and Ti-46Al-8Ta, along with their boron-containing derivatives, where boron is added primarily for grain refinement [9,10]. Recent years have seen the development of third-generation TNM alloys from the Ti-Al-(Nb,Mo) system, which are specifically designed as β-stabilized γ-alloys with an improved balance of properties [11,12]. The mechanisms behind these improvements are well understood. As shown in [13], the addition of Nb leads to substantial solution strengthening, an increased γ-phase volume fraction, and significant grain refinement, thereby simultaneously improving strength and ductility. The authors of [14] confirmed that Nb also enhances the oxidative properties of TiAl alloys by promoting the formation of a more continuous and adherent alumina layer. Similarly, studies [15,16] indicated that using Mo as a β-phase stabilizer in the TiAl matrix not only improves mechanical characteristics through solid solution strengthening but also enhances oxidation resistance by reducing oxygen diffusivity and stabilizing protective oxide scales. The synergistic effect of multiple β-stabilizers, particularly Nb and Mo, has become a focus of recent research, leading to advanced alloy systems with tailored microstructures and optimized high-temperature performance.

The introduction of new titanium aluminide alloys in aircraft turbines involves the progressive replacement of nickel-based superalloy blades. A significant challenge is substituting existing blades with lighter alternatives in more thermally stressed regions of gas turbine engines. According to General Electric, using gamma alloys in low-pressure turbine blades can provide a mass advantage of 100–180 kg compared to traditional nickel superalloys, which have twice the density. Another study [17] developed a new Ni-based superalloy, AMS-OR, with an optimal combination of printability, oxidation resistance, and mechanical properties. It was demonstrated that the strength and ductility of the AMS-OR alloy can significantly exceed those of existing commercial nickel superalloys. TiAl(Nb,Mo) intermetallic alloys exhibit superior properties, including specific strength and high creep resistance at elevated temperatures [18,19,20]. To further enhance physicomechanical properties, the current research is actively exploring intermetallic materials based on multicomponent systems modified with small boron additions (a structure-modifying agent). For instance, TiAl(Nb,Mo)B-based materials have been produced [21,22,23]. Research in this field is vigorously pursued worldwide and holds significant promise for practical applications.

TNB and TNM alloys are fabricated using various methods, including spark plasma sintering [24], casting [25], induction melting [26], and hot isostatic pressing [27], among others. Each technique possesses specific capabilities, application boundaries, and limitations. A promising and energy-efficient approach for producing these alloys is free SHS compression [28,29]. This method combines combustion in the mode of self-propagating high-temperature synthesis (SHS) [30,31] with high-temperature shear deformation [32]. The essence of this method lies in the consolidation and shaping of the synthesized material under constant, relatively low pressure (not exceeding 50 MPa) without the use of specialized molds. This constitutes a significant advantage of free SHS compression for producing large-scale plates and sheets compared to conventional compaction techniques. A further benefit is the highly favorable stress–strain state imposed on the deformable synthesized material, which undergoes not only densification but also shear deformations. This combination of deformations leads to a substantial reduction in defects and porosity in the resulting compact materials. All these specific features of the free SHS compression method, along with its technological and design parameters, significantly influence the combustion process, phase and structure formation, and, consequently, the physicomechanical properties of the final materials. This raises the possibility of fabricating finished materials and components with tailored microstructures and properties.

The aim of the present work is to investigate the combustion characteristics and fabricate compact materials via free SHS compression from initial powder reagents with the following composition (wt%): 51.85Ti–43Al–4Nb–1Mo–0.15B, and to determine the effect of high-temperature isothermal annealing at 1000 °C on the structure and properties of the obtained materials.

2. Materials and Methods

2.1. Objects

The initial reagents were commercial powders of Ti (45 µm, 99.1%), amorphous B (1 µm, 94.0%), Al (5–7 µm, 98.6%), Nb (40–63 µm, 99.9%), and Mo (20 µm, 99.1%). The powder mixture was prepared according to the nominal composition of 51.85Ti–43Al–4Nb–1Mo–0.15B (wt%). The powders were dried, thoroughly blended, and then uniaxially pressed into cylindrical green compacts with a diameter of 50 mm and a mass of 100 g. The compacts were formed with a specific relative density, which was optimized in the present study to ensure efficient combustion propagation.

Due to the relatively low exothermicity of the reaction between the initial components, a separate “chemical furnace” based on a highly exothermic mixture was employed to initiate the self-sustaining combustion wave. This igniter mixture, with a molar composition of 3Ti + Al + 2C, was chosen for its high chemical affinity with the Ti-Al(Nb,Mo,B) system and its ability to provide sufficient ignition temperature. A tungsten coil was used to initiate the reaction within the chemical furnace placed on top of the main powder compact. The heat released by this furnace then ignited the primary SHS reaction in the underlying sample.

2.2. Method of Preparation

Figure 1 presents a schematic diagram of the free SHS compaction process used to fabricate targets for the subsequent deposition of protective heat-resistant coatings via PVD methods. After the combustion wave had propagated through the entire volume of the compact, the synthesized product was immediately subjected to a high-temperature shear deformation process under conditions of free SHS compaction. As illustrated below, the directions of material flow during deformation (indicated by magenta arrows in Figure 1) were perpendicular to the applied axial pressure (shown by the blue arrow). This specific deformation mode is a hallmark of the free SHS compaction technique.

This process resulted in the formation of consolidated disc-shaped samples (pellets) with a final diameter of approximately 75 mm and a height of 10 mm. The surface of these compacts was subsequently ground and mechanically polished to achieve the required roughness (R_a < 0.1 µm) for subsequent materials science characterization and to qualify them as targets for PVD processes.

2.3. Research Techniques

The combustion temperature and velocity under free SHS compression conditions were determined using WRe5/WRe20 tungsten–rhenium thermocouples. These were strategically positioned at the center of the powder compact, at a distance of 10 mm from the center, and within the center of the chemical furnace to capture the ignition event. The thermocouple signals were recorded with a QMBox ADC data acquisition system at a high sampling rate to ensure temporal resolution. The combustion wave velocity was calculated as the ratio of the known distance between two thermocouples to the measured time difference of the wave passage between them, providing a direct measure of the reaction front-propagation kinetics.

Microstructural characterization of the synthesized materials was carried out using a Carl Zeiss UltraPlus field emission scanning electron microscope (SEM) operated in “compo” (compositional) mode for enhanced atomic number contrast. Chemical analysis was performed via energy-dispersive X-ray spectroscopy (EDS) using an INCA Energy 350 XT X-ray microanalysis system (Oxford Instruments, Abingdon, UK) to determine local phase composition and elemental distribution. Phase identification was conducted via X-ray powder diffraction (XRD) on an ARL X’tra diffractometer (ARL, Switzerland) using CuK_α radiation (λ = 1.5418 Å). Data were collected over a 2θ range of 34–67° with a step size of 0.02°, focusing on the primary intermetallic peaks for phase quantification.

Isothermal annealing was performed at 1000 °C for 3 h in a high-vacuum chamber of an HTK2000 attachment mounted on the ARL‘XTRA diffractometer. This in-situ setup allowed for subsequent phase analysis without exposing the sample to the atmosphere, minimizing surface oxidation.

Instrumented indentation testing was conducted in accordance with the ISO 14577-1:2002 standard to evaluate local mechanical properties on polished cross-sections. Tests were performed using a Nanoscan-4D nanoindentation system (NaukaSpetsPribor, Moscow, Russia) with a Berkovich diamond tip. The testing protocol involved an applied load of 0.5 N, a loading time of 10 s, a dwell time of 2 s at peak load for creep minimization, and an unloading time of 10 s. Mechanical characteristics, including indentation hardness and modulus, were derived from the load–displacement curves. The values presented in the article are averages, with measurements statistically validated across at least five independent samples to ensure reproducibility.

3. Results and Discussion

The chemical system under investigation exhibits low exothermicity, making direct ignition using a tungsten spiral impractical. Experimental measurements confirmed that applying a 60 V electrical current to the spiral achieved a maximum temperature of 800–1000 °C, which proved insufficient to initiate a self-sustaining combustion wave in the powder mixture. To overcome this limitation, a highly exothermic “chemical furnace” was employed. This igniter composition, consisting of titanium, aluminum, and carbon black powders, generated a combustion temperature of 1650 °C (Figure 2a) upon electrical initiation—sufficient to reliably ignite the primary powder compact. Thus, the effective ignition temperature for the studied system was established at 1650 °C, resulting in stable, self-propagating combustion behavior.

The relative density of the green compacts significantly influenced both combustion characteristics and propagation stability. Using our previously developed rheological methodology [33] for optimizing powder processing parameters, we fabricated cold uniaxially pressed compacts with relative densities ranging from 0.42 to 0.46. While higher densities were technically achievable, they complicated the necessary drilling of thermocouple holes. Conversely, compacts with densities below this range exhibited insufficient mechanical strength for handling and experimentation.

Figure 2a shows characteristic combustion thermograms for a compact with a relative density of 0.6. All profiles exhibit classical sharp temperature spikes upon passage of the combustion wave. The chemical furnace reached 1650 °C (thermocouple 1), while the synthesized material maintained temperatures of 1345–1360 °C (thermocouples 2 and 3). A clear density dependence was observed: increasing the relative density from 0.6 to 0.63 reduced combustion temperatures from 1360 to 1220 °C, with further density increases showing negligible additional effects. This temperature reduction can be attributed to enhanced interparticle contact at higher densities, which improves heat conduction away from the reaction zone. Combustion velocity, calculated from time differences (t1, t2) between thermocouples spaced 10 mm apart, remained constant at approximately 1.2 mm/s across all density variations. Consequently, compacts with relative density of 0.6 were selected for subsequent studies as they yielded maximum combustion temperatures. This optimization ensures extended dwell time in the plastic state during subsequent shear deformation, ultimately enhancing the quality and density of the final synthesized materials.

To obtain high-quality materials via free SHS compression, a systematic investigation was conducted to optimize the key processing parameter—applied pressure. Experiments revealed that a compaction pressure of 20 MPa yielded optimal results, producing materials with residual porosity not exceeding 1.5%. This low porosity indicates that excellent consolidation was achieved during the combined synthesis and deformation process. The resulting compacts were subsequently machined to produce specimens for comprehensive material characterization, including scanning electron microscopy (SEM), X-ray powder diffraction (XRD), and mechanical testing. Additionally, large-diameter targets (75 mm) were fabricated from the compacts for subsequent application as cathodes in physical vapor deposition (PVD) systems for depositing protective heat-resistant coatings [34,35,36], demonstrating the scalability and practical applicability of the developed method.

SEM analysis revealed a distinctive layered macrostructure in the synthesized materials (Figure 3). The microstructure consists of alternating macro-layers with a characteristic wavy morphology. Critically, the orientation of these layers aligns with the direction of material flow during compression (indicated by magenta arrows in Figure 3), while being perpendicular to the applied axial load (blue arrow in Figure 3). This specific texture is a direct consequence of the shear deformation inherent to the free SHS compression process. The microstructure prominently features irregularly shaped grains exhibiting bright contrast, with sizes reaching up to 80 μm (white regions). Given that the SEM imaging was performed in the “compo” mode, which emphasizes atomic number contrast, these brighter formations correspond to phases containing heavier elements. This suggests the presence of refractory-rich phases, particularly particles of niobium and molybdenum, which are known to form stable intermetallic compounds or solid solutions within the Ti-Al matrix and remain partially undissolved during the rapid synthesis process.

The microstructure of the synthesized materials consists of alternating intermetallic phases of TiAl (γ-phase) and Ti₃Al (α₂-phase), with a minor presence of TiAl₃. This multiphase architecture, commonly reported in literature for Ti-Al-based systems [37,38,39], is characterized by a distinctive lamellar morphology, which is also prominently featured in the obtained materials (Figure 4a). The combined processes of combustion synthesis and subsequent shear deformation resulted in the formation of a primary tetragonal γ-TiAl matrix (Spectra S1 in Figure 4c and S2 in Figure 4d,

Table 1. Quantitative energy dispersive analysis of chemical elements at the selected points indicated in Figure 4.

Spectrum		Content of Chemical Elements, wt%
	B	Al	Ti	Nb	Mo
S1	-	32.95	63.72	-	-
S2	-	35.97	64.03	-	-
S3	-	39.49	60.51	-	-
S4	-	20.36	79.64	-	-
S5	-	42.00	10.59	-	47.41
S6	-	5.61	-	-	94.39
S7	-	37.72	-	-	62.28
S8	-	19.60	22.34	-	58.06
S9	-	37.05	29.29	-	33.66
S10	-	37.03	2.31	-	60.66
S11	-	4.35	-	88.53	-

), along with hexagonal α₂-Ti₃Al phases (Spectrum S4 in Figure 4c, Table 1). These phases exhibit a well-defined lamellar structure, comprising parallel, alternating plates of the respective intermetallic. Notably, SEM-EDS analysis at specific points (Table 1) revealed the presence of partially dissolved niobium and molybdenum particles. At the periphery of larger particles, interdiffusion with titanium and aluminum has been initiated, while, for finer particles, this dissolution process was essentially complete (Spectrum S5 in Figure 4d, Table 1), indicating their incorporation into the intermetallic matrix.

A critical aspect of the process is the rapid quenching effect that occurs following combustion. After the passage of the combustion wave, the immediate compression by the press plunger induces significant heat loss from the sample, leading to abrupt cooling. This rapid quenching effectively arrests the phase transformation and microstructure evolution processes at the moment of compression. Ideally, these processes would continue towards equilibrium given sufficient time at high temperatures or slower cooling rates. However, a fundamental process dilemma exists: with increasing time after combustion, the synthesized material loses its plastic properties, which are essential for effective consolidation during the shear deformation step. Therefore, a key objective of this study was to identify the optimal temporal window—maximizing phase formation and microstructural development while retaining sufficient material plasticity for successful deformation and densification.

Figure 4c (Spectrum S6) corresponds to a central molybdenum particle (bright white region), surrounded by a distinct gray interfacial layer (Spectrum S7). Figure 4b presents the characteristic X-ray elemental map of the highlighted area. It is evident that the region characterized by Spectrum 6 is molybdenum.

EDS analysis of these points (Table 1) confirms the formation of a MoAl₂ phase at the interface. This provides direct evidence for a diffusion-controlled solid–liquid interaction mechanism between the molten aluminum (enabled by the combustion temperature of 1360 °C) and the solid molybdenum particle during the synthesis. A further concentric rim, approximately 2–3 µm thick, is observed encircling the MoAl₂ phase (Spectra S8 and S9, Table 1). Based on EDS results, this rim is hypothesized to be a ternary Ti-Al-Mo compound forming at the boundary between the MoAl₂ and the titanium aluminide matrix. The exact stoichiometry of this ternary phase is determined through the XRD analysis presented subsequently. Within the bulk material, smaller molybdenum particles (below 10 µm in one dimension) were found to have undergone complete interdiffusion with titanium and aluminum during combustion and deformation, resulting in the formation of a homogeneous ternary Ti-Al-Mo phase (Spectrum S10, Figure 4d, Table 1). In contrast, larger molybdenum particles exhibit incomplete diffusion gradients (Spectrum S10, Figure 4d), indicative of kinetic limitations.

The behavior of niobium particles differs significantly due to their fourfold higher abundance. Consequently, not all niobium particles interacted to the same extent. Spectrum S11 (Figure 4c) points to a fragmented niobium particle with no detectable diffusion halo, suggesting it remained largely inert. In the volume of material one can find niobium particles, however, they display rounded edges and are surrounded by a pronounced diffusion zone enriched with aluminum, titanium, and, notably, up to 14 wt% boron (Spectra S2 and S3, Figure 5). This suggests active dissolution and the formation of boride-containing phases. However, there are not many such particles in the entire volume; in most cases, Nb grains of a fragmented shape without clearly defined diffusion boundaries are found.

The elemental mapping results presented in Figure 4b reveal overlapping characteristic X-ray signals for molybdenum and niobium, indicating their co-localization and potential synergistic interaction within the matrix. Near the phase boundary (Spectrum S4, Figure 5,

Table 2. Quantitative energy-dispersive analysis of chemical elements at the selected points indicated in Figure 5.

Spectrum		Content of Chemical Elements, wt%
	B	Al	Ti	Nb	Mo
S1	6.85	-	-	82.11	-
S2	13.11	15.68	26.96	40.00	-
S3	14.42	18.82	32.02	30.91	-
S4	-	21.96	64.11	13.92	-

), zones are enriched in aluminum and titanium but are devoid of boron and exhibit reduced niobium content. This compositional profile strongly supports the in-situ formation of a ternary Ti-Al-Nb intermetallic compound during the synthesis process.

XRD analysis of the powdered material (particle size < 160 µm) confirmed that it primarily consists of titanium aluminide intermetallic phases: TiAl (γ-phase), Ti₃Al (α₂-phase), and a minor amount of TiAl₃ (Figure 6). Characteristic peaks for phases containing Mo and B were not distinctly resolved due to their low volumetric concentration and peak overlapping with the major phases. A notable observation is the angular shift of the Nb-related diffraction peaks from their standard positions (Figure 6), indicating lattice parameter changes consistent with the formation of niobium-containing compounds. Combined with the aforementioned SEM-EDS results, this confirms the formation of a solid solution beyond simple binary intermetallic. Its stoichiometry, estimated as (Al_0.86Nb_0.14)(Ti_2.85Nb_0.15), suggests significant niobium dissolution in the titanium aluminide matrix.

The crystal lattice parameters of this ternary compound are close to those of Ti₃Al (

Table 3. Quantitative energy-dispersive analysis of chemical elements at the selected points indicated in Figure 5.

Phase	Lattice Parameters, Å	Crystal System	Space Group	ICDD PDF2 No.
Nb	a = 3.30332	cubic	Im-3m	34-0370
TiAl	a = 2.844 c = 3.9447	tetragonal	P4/mmm	10-84-3907
TiAl3	a = 5.782 c = 4.629	hexagonal	P63/mmc	10-82-5277
Ti3Al	a = 5.77 c = 4.62	hexagonal	P63/mmc	65–7534
(Al0.86Nb0.14)(Ti2.85Nb0.15)	a = 5.77 c = 4.64	hexagonal	P63/mmc
Ti2AlMo	a = 3.17	cubic	Pm-3m	10-82-5224
Nb2Al	a = 9.943 c = 5.186	tetragonal	P42/mnm	12-0074

), leading to significant peak overlapping. According to SEM-EDS data, this solid solution formed via diffusion of aluminum and titanium into niobium particles and is predominantly located at their periphery.

This finding is consistent with the ternary Ti-Al-Nb phase diagram at the 700 °C isothermal section (Figure 7), which reveals a stability region for a ternary compound with stoichiometry Ti₄NbAl₃. Although the diagram represents the elemental composition without accounting for Mo and B, it substantiates the feasibility of forming a Ti-Al-Nb-based ternary compound. Thus, the phase composition can be approximated as 60 at% TiAl + 20 at% Ti₃Al + 20 at% Ti₄NbAl₃. Furthermore, SEM-EDS analysis indicated the formation of a ternary Ti-Al-Mo compound at the interface between the MoAl₂ phase and the titanium aluminide matrix. The positions of its characteristic XRD reflections show the closest match to the Ti₂AlMo compound (Table 3), suggesting this is the most likely phase formed under the given synthesis conditions.

Given the intended high-temperature application of the synthesized materials, they were subjected to isothermal annealing at 1000 °C for 3 h. In situ X-ray powder diffraction analysis was performed directly within the vacuum chamber after each hour of annealing, with a final diffractogram recorded after three hours of treatment and subsequent cooling. The presence of tungsten reflections in Figure 6 indicates the location of the tungsten heating element beneath the sample.

Based on the SEM and EDA results, it was established that the diffusion processes at the grain boundaries of Nb and Mo did not end as a result of the sharp heat removal during compression, and products were formed that were not in an equilibrium state; thus, high-temperature annealing should contribute to their continuation.

Despite the dynamic high-vacuum conditions (~5 × 10⁻⁵ bar), aluminum oxide (Al₂O₃) was detected on the material surface after one hour of thermal treatment. Characteristic Al₂O₃ peaks become clearly visible in the material after two hours of annealing (Figure 6c). The intensity of these peaks at approximately 43° and 57° further increases after the third hour, indicating progressive growth of the oxide layer on the surface. Additionally, these peaks shift toward higher angles after the third hour and subsequent cooling. The annealing treatment induced significant phase transformations: (i) a substantial increase in Nb₂Al phase intensity with a concurrent decrease in Nb reflection intensity; (ii) notable enhancement of Ti₃Al reflections accompanied by a shift of h0l peaks (101 and 202) toward larger interplanar spacing, suggesting increased formation of the (Al_0.86Nb_0.14)(Ti_2.85Nb_0.15) solid solution consistent with the ternary phase diagram; (iii) the complete disappearance of TiAl₃ reflections after 3 h, indicating its full transformation at high temperatures.

Figure 8 shows the microstructure of the material after annealing, and

Table 4. Quantitative energy-dispersive analysis of chemical elements at the selected points indicated in Figure 8.

Spectrum		Content of Chemical Elements, wt%
	B	Al	Ti	Nb	Mo
S1	8.40	2.16	1.41	88.03	-
S2	-	19.55	48.42	32.04	-
S3	14.42	26.18	42.83	30.98	-
S4	-	20.83	55.59	-	23.58
S5	-	29.86	70.14	-	-

lists the results of the EDS analysis over a 2 × 4 µm scanning area. The figure reveals that the intermetallic phases have a more distinct lamellar structure (region S5), which supports the XRD results indicating an increased content of the Ti₃Al phase. It can also be observed that the molybdenum grains have completely transformed into a ternary compound based on Mo-Ti-Al (region S4). No pure molybdenum was found within the material volume. Furthermore, diffusion processes continued into the core of the niobium grain during annealing. It is evident that, along the perimeter of the grain, which has decreased in size to less than 10 µm (region S1), a rounded area measuring up to 15–20 µm has formed. This area consists of two phases: Nb₂Al and (Al_0.86Nb_0.14)(Ti_2.85Nb_0.15), corresponding to regions S2 and S3, respectively.

The influence of heat treatment on the mechanical properties of the obtained samples was determined. Figure 9 presents the loading/unloading curves from indentation, based on which the mechanical characteristics were determined. As shown in

Table 5. Average values of mechanical properties of the material before and after heat treatment.

Type of Heat Treatment	Microhardness, GPa	Modulus of Elasticity, GPa	Elastic Recovery, %
Without annealing	5.4	175	21.3
After annealing at 1000 °C for 3 h	6.6 (up to 7.4)	185 (up to 200)	26.5 (up to 31.8)

, the average values of mechanical properties of the materials after heat treatment increased: microhardness and elastic recovery increased by 1.2 times, while Young’s modulus increased slightly by 10 GPa. The maximum microhardness for the heat-treated material reached values up to 7.4 GPa, elastic recovery up to 31.8%, and Young’s modulus up to 200 GPa. This confirms the formation of new phases or transformation of existing intermetallic phases in the material during heat treatment. Thus, if the initial material predominantly contained the TiAl phase (with a microhardness of 3.1 GPa for the single-phase material), after heat treatment, its amount decreased while the amount of Ti₃Al phase increased (with a microhardness of 4.5 GPa for the single-phase material). It was also found that the amount of free niobium, which has relatively low hardness (1.3 GPa), decreases after heat treatment. As previously shown by SEM, diffusion zones are formed at phase boundaries, which, according to XRD, can be identified as both the Nb₂Al binary phase and the (Al_0.86Nb_0.14)(Ti_2.85Nb_0.15) solid solution, both having significantly higher hardness than pure niobium. Consequently, the presence of these phases and the transformation of intermetallic phases enhance the mechanical characteristics of the material after annealing.

The comprehensive findings of this study regarding the phase composition, microstructure, and mechanical properties of the synthesized materials allow for optimistic predictions about their successful application as targets for depositing protective heat-resistant coatings onto various components. A key advantage is the potential use of these targets both in their as-synthesized state and after the implemented post-annealing treatment.

Materials obtained via free SHS compression offer a unique combination of a fine-grained, consolidated structure with a high degree of plasticity inherited from the shear deformation process. This may be preferable for certain PVD techniques where target ductility is a critical parameter to prevent cracking during operation. The annealed state provides a significant enhancement in hardness and elastic recovery, driven by the formation of a more stable and harder phase assemblage, including ternary solid solutions and intermetallic compounds such as Nb₂Al. This improved mechanical robustness is highly beneficial for extending target service life under intense energetic bombardment during sputtering processes, potentially leading to more stable and consistent coating deposition rates.

Thus, the developed approach offers valuable flexibility, allowing for the selection of a target state (as-synthesized or annealed) that is optimally suited to the specific requirements of the PVD coating technology and the intended operational conditions of the protected details.

4. Conclusions

• Optimal synthesis parameters for initiating combustion and achieving maximum performance for the 51.85Ti–43Al–4Nb–1Mo–0.15B (wt%) composition were established. These include fabricating powder compacts with a relative density of 0.42 and employing a highly exothermic 3Ti-Al-2C (molar ratio) chemical furnace. This configuration ensures a maximum combustion initiation temperature of 1650 °C, which is critical for sustaining the reaction in this low-exothermicity system.
• The materials synthesized via free SHS compression exhibit a characteristic layered macrostructure with a wavy morphology. The layers are aligned parallel to the material flow direction during compression and perpendicular to the applied load. The microstructure consists of alternating intermetallic phases (TiAl (γ) and Ti₃Al (α₂)) with a minor amount of TiAl₃, and uniformly distributed niobium particles throughout the volume. Interdiffusion of titanium and aluminum occurred at the boundaries of molybdenum and niobium grains, leading to the formation of binary (e.g., MoAl₂) and ternary (Ti-Al-Mo, Ti-Al-Nb) compounds.
• Isothermal annealing at 1000 °C for 3 h promotes significant phase evolution, resulting in an increased volume fraction of the (Al_0.86Nb_0.14)(Ti_2.85Nb_0.15) ternary solid solution and the formation of the Nb₂Al intermetallic phase. This phase transformation confirms a solid-state diffusion mechanism as the primary driver for the formation of these compounds during the annealing process.
• The phase and structural changes induced by isothermal annealing directly lead to an enhancement of mechanical properties: (i) microhardness and elastic recovery increased by a factor of 1.2, and (ii) Young’s modulus exhibited a modest increase of ~10 GPa. The maximum values achieved for the annealed material were: 7.4 GPa (microhardness), 31.8% (elastic recovery), and 200 GPa (Young’s modulus). This improvement is attributed to the formation of a harder and more stable phase assemblage.
• The authors first demonstrated the feasibility of synthesizing multicomponent materials based on the Ti-Al-(Nb,Mo) system under conditions that combine combustion processes and high-temperature shear deformation. The results of the measured physico-mechanical properties of the obtained materials indicate their promising potential for use both as structural and functional components, e.g., as targets for depositing protective coatings via PVD methods.