The annealing treatment of deformed β-Ti alloys leads to the evolution of α and ω phases when applied temperatures are below the β-transus, whereas temperatures above the β-transus lead to nucleations and grain growth of the single β microstructure [60]. The effect of heating rate and temperature (500–600 °C) on the aging behavior of TIMETAL-LCB, VT22, and Ti-15333 β-Ti alloys were reported in [61]. TIMETAL-LCB and VT22 formed fine plate-like α at slow heating rates due to the precipitation of isothermal ω at low temperatures, which serves as nucleation sites for α [61]. However, at high heating rates, the formation of isothermal ω was avoided, leading to coarse, plate-like α microstructures with less desirable properties. Ti-15333, on the other hand, exhibited β phase separation (β′+βmatrix) during isothermal aging rather than isothermal ω formation. It has been reported that like isothermal ω, β′ can also act as a nucleation site for α [61,62]. It is crucial to note that there are three categories of ω phase, depending on the process of formation: (1) a deformation-induced ω phase, (2) athermal ω phase, and (3) isothermal ω phase. The athermal ω phase, resulting from rapid quenching [63], is a diffusionless transformation but is not related to martensitic transformation, while the isothermal ω phase forms during the low-temperature aging of β-Ti alloys. The deformation-induced ω-phase forms under applied stress/strain [63]. Evolving phases during the heat treatment can also affect the mechanical strength of the β-Ti alloys. Precipitation of α or ω precipitates in solution-treated or deformed β-type Ti alloys is useful to obtain improved static and fatigue strength values. On the other hand, α and ω phases have significantly higher intrinsic E (E(ω) ≈ 153 GPa, E(α) ≈ 115 GPa) than the β phase (E(β) ≈ 60–65 GPa) and hence are not useful for bio-implants [64]. A heat-treated microstructure was observed in the hot-compressed Ti-13V-11Cr-3Al alloy samples. GB maps were reported for samples subjected to hot compression tests at 930 and 1030 °C and a strain rate of 0.1 s−1 [65]. EBSD analyses showed that continuous dynamic recrystallization (CDRX) leads to considerable grain refinement through the dissociation of coarse deformed grains [65,66]. It was observed that well-developed subgrains formed by extended dynamic recovery (DRV) were responsible for the grain dissociation. Inhomogeneous compressive deformation, due to the faster evolution of the substructure at regions adjacent to the GBs, can be observed in. Severe deformation causes the evolution of a large number of dislocations near the GBs. The existence of these dislocations at the GB regions favors DRV and leads to accelerating substructure formation. The micrographs also showed the formation of some small substructure-free volumes of the old grains via CDRX, which were surrounded by LAGBs + HAGBs [65]. Grain boundary serrations and nucleations were typical of the propensity for discontinuous dynamic recrystallization (DDRX). However, the decreased fraction of subgrains and the absence of GB serrations as well as DDRX were observed in samples processed at 1030 °C [65]. Since the processing temperature (930 and 1030 °C) was above the β-transus, this caused the presence of a single β phase during compression deformation. In contrast, a low-temperature compression test (700 °C) of Ti-7Mo-3Al-3Nb-Cr (Ti-7333) caused the precipitation of the α phase [67]. A fine α phase was precipitated during isothermal deformation at 700 °C-10⁻³ s⁻¹ [67]. The evolution of intragranular α (spherical precipitates) and grain boundary α (αGB) with a combination of HAGBs+LAGBs was observed [67].

The annealing treatment of a cold-rolled Ti-5Al-5Mo-5V-3Cr (Ti-5553) alloy was performed [68]. The as-received Ti-5553 alloy showed a weak texture and its intensity got strengthened near a partial α- and complete γ-fiber texture after 40% RR. Annealing treatment (860 °C–5 min) of a cold-rolled Ti-5553 alloy led to decreased intensity of the deformation texture [68]. Further, aging treatment of the as-received sample (below β-transus) at 670 and 770 °C caused the evolution of a fine α-phase inside the coarse β-grains. The evolution of the α-phase diminishes the texture of the aged samples. In a previous investigation on the annealing treatment of a cold-rolled Ti-15333 alloy, microstructure and crystallographic texture evolution have been discussed [69]. In a study, the effects of cooling rate following β- and α/β-region (920–1000 °C) heat treatment on microstructure and phase transformation were investigated for a Ti-6.5Al-2Sn-4Zr-4Mo-1W-0.2Si (BT25y) alloy [70]. The BT25y alloy was soaked at 920–1000 °C for 10 min, and then cooled at a rate of either 0.15 °C/s–150 °C/s to RT [70]. Microstructure observations indicated that the microstructure of the BT25y alloy was significantly influenced by the cooling rate. When the material was cooled from the β phase field at a lower rate, the grain boundary α (α_GB) and Widmanstatten α (α_WGB) phases were precipitated [70]. However, increasing the cooling rate greatly restrained the precipitations of α_GB and α_WGB phases. In this case, acicular martensite (α′) was precipitated inside the β grain. The primary equiaxed-α was retained when the material was cooled down from the α/β phase field [70]. The content and size of equiaxed-α decreased with the increasing solution temperature but were independent of the cooling rate [70]. Texture evolution in a heat-treated Ti-22Nb-6Ta alloy was investigated in [71]. A well-developed {001}<1–10> texture was obtained in the cold-rolled sample and after heat treatment at 600 °C for 10 min. Moreover, a recrystallization texture of {112}<1–10> was developed at 900 °C for 30 min [71].