3.1. Microstructures
Figure 2a–c is the microstructures of three different heat-treated samples. All microstructures contained α
2 phase particles, lamellar O phase, secondary acicular O phase, and B2 matrix. It can be seen that they are very similar, except for the differences in the secondary acicular O phase. All the aged microstructures contained secondary acicular O phase within the B2 matrix. That is because the O and B2 phases have specific orientation relationships (O.R.) in the low-energy configuration [
12]. The statistical data of microstructural characteristics in the three microstructures are listed in
Table 2. In
Table 2, the standard deviations have provided with the mean values. It can be seen that as aging temperature increased, the diameter of equiaxed α
2 phase showed nearly unvaried value, the volume fractions of total precipitates and secondary acicular O phase, and width of the secondary acicular O phase increased, while the length of acicular O decreased. Hence, the secondary acicular O phases were determined by aging temperature. The thickest size of secondary acicular O phase was presented at 840 °C, and the thinnest size was displayed at 760 °C. From our previous results in our group [
14], the equiaxed particles in these microstructures were determined by isothermal forging temperature. The lamellar O phases were decided by solution temperature, and the secondary acicular O phases were adjusted by aging temperature. As aging temperature increased, the thickness of secondary acicular O phase increased, while the length and volume fraction decreased. The major research objectives in this work were tensile and creep behaviors, and the deformation mechanisms of the three microstructures in the latter.
3.2. Room Temperature Tensile Properties and Deformation Mechanism
Tensile properties of the equiaxed microstructure at ambient temperature are presented in
Table 3. In
Table 3, the standard deviations have provided for these mean values of the four mechanical properties. In most cases the results represent the average of at least three tests. Thus, the tensile strength of HT-760 was better than that of HT-840. However, it exhibited worse elongation and reduction area.
Strengthening mechanisms of Ti
2AlNb-based alloys are similar to those of traditional titanium alloys. The various parameters that influence the yield strength are: (1) grain size of the B2 phase, (2) strengths of the individual phases influenced by changes in their composition with heat-treatment, (3) volume fractions of α
2/O and B2 phases, (4) lath size of O phase, and (5) dislocation substructure [
17]. It can be seen that grain boundary strengthening is the most important factor. The sizes of the secondary acicular O phase should be another important factor. In the three microstructures, aging temperatures are varied. Previous results showed that equiaxed particles were regulated by isothermal forging, lamellar O phases were controlled by varied solution temperature, and the secondary acicular O phases were adjusted by the aging temperature. The coarsening of secondary acicular O phases was caused by increasing aging temperature. Thus, the tensile mechanical properties of three microstructures were mainly decided by the secondary acicular O phases. This variation trend of strength is caused by noticeable coarsening and volume fraction reduction of secondary acicular O phase with increasing aging temperature [
18]. This result is in agreement with literature [
15]: that the yield strength is associated with the sizes and morphologies of constituent phases.
Equiaxed α
2 phase, B2 phase, and the region around the α
2 and O phases are the major factors influencing the variation in ductility. The differences in
E value are mainly ascribed to the volume fractions of the O phase and B2 phase. The O-phase has fewer available slip systems than the B2 structure [
16]. During tensile deformation, the ductility of the alloy was determined by the number of slip systems. Compared with B2 phase, O phase has fewer slip systems, thus the lower
E value was attributed to the limited slip systems in O-phase-dominated microstructures. The microstructure of HT-840 has more B2 phases and less acicular O phases, so the ductility of HT-840 at room temperature was superior to that of HT-760.
In the following discussion, room temperature (RT) tensile deformation mechanism of equiaxial microstructure was further analyzed by SEM and TEM techniques.
Figure 3 is the tensile fractographs of HT-760. The fractographs contain the fibrous zone (A in
Figure 3a), radiation area (B in
Figure 3a), and the shear lip (C in
Figure 3a). The shear lip of the fracture face was smallest. From
Figure 3, it can be seen that a crack started from the specimen’s interior (fibrous zone). The orientation of the crack growth was presented along the arrows of
Figure 3a. As the stress increased to critical value, the crack was propagated rapidly. The shear lip was formed on the rim of the fracture.
Figure 3c is the microstructure of the radiation area. The size of the raised particle is similar to that of the equiaxed α
2 phase. The tear ridge of the equiaxed α
2 particles showed the shear fracture characteristics. The explanation for this is that the deformed abilities of different phases were varied during tensile tests. The slip systems in equiaxed α
2 phase was less than that of B2 matrix, thus B2 matrix has better deformability. Due to different deformability of each phases, the microvoids were easily formed at phase interface between equiaxed α
2 phase and B2 matrix. The shear dimples were formed along the direction of shearing stress (
Figure 3d). From above analysis, the room temperature fracture mechanism of HT-760 is a mixed-rupture characteristic of quasi-cleavage and dimples.
The surface slip characteristics of the deformed samples could reflect the deformation mechanism of samples.
Figure 4a is the micrograph of surface slips in HT-760. Wavy slip lines were present in the B2 phase. Meanwhile, longer slip lines were passed through equiaxed α
2 particles and B2 matrix. This phenomenon showed that cross-slips and multiple slips operated in the B2 phase. There are longer slip lines in the equiaxed α
2 phase—this means that plane slippage characteristics were present in the α
2 phase. B2 phase has more slip planes than the α
2/O phase, thus, deformation was mainly focused in the B2 matrix [
19,
20]. To further explore the RT deformation mechanism of HT-760, the microstructures were studied by TEM. In the undeformed area of tensile samples (
Figure 4b), the dislocation density in the interior of the equiaxed α
2 phase was lower, while the dislocation density at phase interfaces between B2 and O was higher. The explanation for this is that during isothermal forging, the large deformation was occurred at phase interfaces between α
2/O and B2 phases. After tensile deformation (
Figure 4c), dislocation of the primary α
2 phase was straight, and formed a dislocation net. Dislocation pile-up and tangling around the lamellar O were also observed after a high degree of deformation. Multi-slips were activated in the equiaxed α
2 particles. In the region of large plastic strain, the dislocation net formed by dislocation slips was passed through all equiaxed particles. It could relieve high stress in the particles, and coordinate deformation between the equiaxed α
2 particles and B2 matrix (
Figure 4d).
In the equiaxed microstructure, besides the equiaxed α
2 particles, the influence of lath O phases on RT tensile properties is also important.
Figure 5 is the undeformed and deformed lath O phases. The undeformed lamellar O phase has lower dislocation density (
Figure 5a).
Figure 5b is the deformed lamellar O phase, it can be seen that slip bands originated at the interface between B2 and O phases, and then passed through the lamellar O phase and B2 matrix. There were many piled up dislocations at the phase interface between O and B2. Slip bands such as B2 → O → B2 were also observed, because the slip bands from the B2 phase were went through the lamellar O phase and B2 matrix. The continuous B2 phase promoted the development of slip systems. This is the key factor for improving the ductility of this alloy (
Figure 5b). Based on the above analysis about the deformation of the equiaxed microstructure, the mechanisms of nucleation and propagation of cracks could be concluded. In the equiaxed microstructure, slip bands were transferred via equiaxed particles, cracks were easily formed at the interface between the equiaxed α
2 phase and B2 phase.
3.3. High Temperature Tensile Deformation
Figure 6 shows the engineering stress-strain curves of HT-760 at different temperatures and strain rates. The curves present two characteristics. Firstly, the deformation temperature is sensitive to the flow stress. As the temperature increased, the strength of alloy decreased, while the ductility increased. At the strain rate of 0.001 s
−1, the deformation temperature increased from 600 to 700 °C, as the peak stress decreased from 778 to 712 MPa. Meanwhile, the flow stresses of the alloys also decreased. Secondly, at constant strain rate, the ductility of the alloy improved with increasing deformation temperature. The main reason is that (a) before tensile deformation, the samples needed to be heated to the deformation temperature. The temperature range of 600–700 °C belongs to the O + B2 phase region. As more secondary acicular O phases were precipitated from the B2 matrix, the relative content of B2 phase decreased, which is a process of precipitation strengthening. Another reason is that (b) the activation energy of the alloy also increased with the increasing deformation temperature. Hence, the flow stress of the alloy decreased while the ductility increased.
Figure 7 shows high-temperature fractographs under the different deformation temperatures, with strain rates of 0.01 s
−1. The fracture appearance shows that ductile fracture occurred during high temperature deformation. From the microscopic morphology, it can be seen that the dimple pattern gradually increased with increasing test temperature, meaning that the ductility of the alloy increased. The microstructure evolution under the different strain rates with deformation temperatures of 700 °C are presented in
Figure 8. Relative to the heat-treated microstructure, the undeformed zones were heat-treated again and had more acicular O phases. In the deformed area, as the strain rate increased, the volume fraction of O phases was nearly constant, owing to the attainment of phase equilibrium, while the lamellar O phase was not stable and inclined to the tensile direction. Therefore, the microstructural evolution during deformation was dominated by phase and morphology changes.
3.4. Creep Properties
The creep properties of three microstructures were also investigated. Creep curves at the stress of 150 MPa and temperature of 650 °C are showed in
Figure 9. As the aging temperature increased, the creep resistances of the microstructures also improved. The creep rate of the primary creep stage for HT-840 was lower than those of HT-800 and HT-760. HT-840 had the best creep resistance, at 650 °C/150 MPa. The reasons for different creep properties in three heat treatment schedules are the differentiation of microstructure before creep deformation. The microstructure after isothermal forging mainly contained equiaxed α
2 particles and the lamellar O phase formed by air cooling. After solution treatment, the acicular O phase was dissolved in B2 matrix. Some had already become globular during heat-treatment. During aging treatment processing of the O + B2 phase region (760–840 °C), very fine secondary acicular O phases were presented at 760 °C. This acicular O phase easily deformed during creep deformation; thus, the creep resistance of HT-760 was worst. The coarse lamellar phase formed at high aging treatment temperature (840 °C) had the function of strengthening during creep deformation.
Figure 10 and
Figure 11 are the microstructures of HT-760 and HT-840, respectively, after creep deformation. When the aging temperature increased, the dislocation density of the equiaxed α
2 particles reduced. In the HT-760 microstructure, there were several dislocation pile-ups at the interfaces of equiaxed α
2 particles and lamellar O phase. Some dislocations passed through the interiors of equiaxed particles, reducing the creep resistances of the equiaxed particles. For the HT-840 microstructure, dislocation hardly went around the coarser lamellar O phase, while the secondary acicular O phase was smaller and more easily deformed. Boehlert and Miracle [
21] overviewed the creep behaviors reported for α
2- and O-based Ti-Al-Nb alloys. They divided the creep behaviors into three regimes, and identified the creep mechanism in each regime according to the values of
n and
Q. When
n = 1 to 1.9 and
Q = 107 to 187 kJ/mol, the creep process is controlled by grain-boundary diffusion. When
n = 1.9 to 2.8 and
Q = 256 to 327 kJ/mol, the creep process is controlled by grain-boundary sliding. When
n = 3.5 to 7.2 and
Q = 241 to 376 kJ/mol, the creep process is controlled by dislocation climb. It is generally accepted that high-temperature dislocation climb is associated with lattice self-diffusion, while grain-boundary sliding is associated with either lattice self-diffusion or grain-boundary diffusion. Typically, the activation energy for grain-boundary diffusion is half that of lattice self-diffusion [
22,
23,
24]. It is no denying that diffusion has an important influence on creep properties of alloys. During current research, the creep mechanism of equiaxed microstructure is therefore dominated by a dislocation-controlled creep process. This conclusion is supported by observations of dislocation densities. The current observations are insufficient to determine whether a glide or climb mechanism is rate controlling.