Figure 1 shows the microstructure of Ti-6.6Al-4.6Sn-4.6Zr-0.9Nb-1.0Mo-0.32Si alloy. The average size of β grain is about 80 μm, and some primary α phases are visible at the β grain boundaries. The microstructure consists of coarse colonies of similarly aligned Widmanstätten α + β phase. Size of both α lath and α + β colonies increase with the growth of β grain size. To be specific, the average size of α + β colonies in Figure 1 is about 28 μm and that of α lath is about 7.2 μm.

Microstructural characteristics (such as morphology and volume fraction of α and β phases, β grain size, etc.) are important factors that significantly influence the mechanical properties of Ti alloys (Ding et al., 2002; Jia et al., 2014; Li et al., 2018). The mentioned microstructural characteristics are determined by chemical composition, processing, and heat treatment method (Ding et al., 2002). The bimodal, Widmanstätten, and basket microstructures are often used in near α and α+β Ti alloys (Pan et al., 2009). Widmanstätten microstructure in Ti alloys has been reported to provide a better creep resistance compared with bimodal microstructure (Balasundar et al., 2014).

The room and high temperature tensile properties of Ti-6.6Al-4.6Sn-4.6Zr-0.9Nb-1.0Mo-0.32Si alloys are listed in Table 1. At room temperature, yield strength, ultimate strength, elongation, and elastic modulus of Ti-6.6Al-4.6Sn-4.6Zr-0.9Nb-1.0Mo-0.32Si alloys are 1,010 MPa, 1,150 MPa, 12%, and 118 GPa, respectively. The yield strength is higher than that of near α Ti-60 (∼982 MPa) (Li et al., 2017), Ti-6242s (931 MPa) (Hosseini et al., 2017), and IMI834 (∼931 MPa) (Balasundar et al., 2014). The high temperature tensile properties are affected greatly by temperature increasing. Ultimate strength and elongation of Ti-6.6Al-4.6Sn-4.6Zr-0.9Nb- 1.0Mo-0.32Si alloys are 698 MPa and 14% at 873 K. In comparison with those of 873 K, ultimate strength decreases about 10 and 20%, while elongation increases about 43 and 143% when tested at 923 K. The ultimate strength is about 13% higher than that of near α Ti-60 (∼617 MPa) at 873 K (Li et al., 2017).

Figure 2A,B show the SEM fractographies of Ti-6.6Al-4.6Sn-4.6Zr-0.9Nb-1.0Mo-0.32Si alloy after room temperature and 973 K tensile, respectively. A large number of ductile dimples and a few stream-like patterns are observed in Figures 2A,B, which indicates that the fracture made of both room temperature and 973 K is ductility fracture. Dimples in Figure 2B are more common and deeper than those in Figure 2A, which would indicate better ductility of specimens tested at 973 K tensile than that of room temperature. The result is consistent with Table1.

The curves of logarithmic scale of creep strain vs. logarithmic scale of time are shown in Figure 3, which exhibit typical features of creep curves of Class-M solution alloy: the primary stage with a decreasing creep rate, the steady-state stage with steady creep rate, and the third stage with significantly increasing creep rate before creep rupture. The steady state creep stages last relatively longer at low temperatures or low stresses compared with those at high temperatures or high stresses. The steady stage lasts for a long time and the test is terminated after a constant creep rate obtained, when the temperature or stress is low. While, with a high the temperature or stress, the first and second stage are short, directly entering the third stage, and the minimum creep rate is determined as steady state creep rate.

In general, the creep rate of this Ti-6.6Al-4.6Sn-4.6Zr-0.9Nb-1.0Mo-0.32Si alloy can be obtained from the following Eq. 1 (Xiao et al., 2011): ε = A σ n   exp ( − Q R T ) (1) where ε • is steady state creep rates (creep rates), A is the material property constant, σ is the stress, n is the stress exponent, Q is the activation energy, R is the gas constant, and T is the temperature. Steady state creep rates of Ti-6.6Al-4.6Sn-4.6Zr-0.9Nb-1.0Mo-0.32Si alloy at different testing conditions are shown in Figure 4. The creep rate is found to increase with increasing temperature and stress. The creep strain of Ti-6.6Al-4.6Sn-4.6Zr-0.9Nb-1.0Mo-0.32Si alloy under 200 MPa is much smaller than that of IMI834 0.28% under 150 MPa at 873 K (Balasundar et al., 2014). Creep rate of near-α Ti alloy Ti-6Al-4Sn-4Zr-0.8Mo-1Nb-W-0.25Si is 3.35 × 10^–8 s⁻¹, 4.68 × 10^–7 s⁻¹, or 1.27 × 10^–6 s⁻¹ under 150, 200 MPa, or 250 MPa at 873 K (Zheng et al., 2021). Steady state creep rate of Ti-5Al-5Mo-5V-1Fe-1Cr alloy is 9.27 × 10^{– 8} s⁻¹ at 773 K, 200 MPa (Nie et al., 2014). The steady-state creep rate of TC18 samples with lamellar structure is 1.90 × 10^–8 s⁻¹ at 773 K, 200 MPa (Nie et al., 2016). The creep rate of near -α Ti alloy IMI834, IMI834 melt using high iron sponge, and IMI834 melt using low iron sponge with fully transformed β structure -lamellar α and retained β films is 5 × 10^–8∼1.5 × 10⁻⁷s⁻¹ at 923K, 220 MPa (Mishra et al., 2005). At 873K under an applied load of 210 MPa, the creep rate of Ti6242S is 5 × 10^–8 s⁻¹ after 60 h (Wimler et al., 2020). It can be concluded from the above report that the creep rate of Ti-6.6Al-4.6Sn-4.6Zr-0.9Nb-1.0Mo -0.32Si alloy is significantly reduced.

The stress exponent (n) could be calculated by logarithmic interpolation of Eq. 1. Figure 5 shows the curves of stress - creep strain rate of Ti-6.6Al-4.6Sn-4.6Zr-0.9Nb-1.0Mo-0.32Si alloy on a double logarithmic scale. At 873 K, 923 K, and 973 K, the stress exponents are about 3.5, 3.5, and 3.1 in Figure 5. All the stress exponents locate within the range of solid solutions 3.0–3.5, so the creep mechanism is diffusion viscous-drag-controlled (Pan et al., 2009). When the stress exponent is 3.5, creep is dominated by dislocation sliding at low stress levels, but the stress exponent is 4.6 when creep is governed by climb of dislocation at higher stress level. If Al content in Ti alloy is high and temperature of creep test is high, the transition from climb of dislocation into glide of dislocation may occur (Pan et al., 2009).

The stress exponents of Ti alloy decrease with temperature increasing. If the shear modulus is added to Eq. 1, the steady-state creep rate can be deduced as follows (Es-souni, 2001): ε ˙ = A G b R T ( σ G ) n D = A G b R T ( σ G ) n D 0 ⁡ exp ( − Q R T ) (2) Q [ kJ/mol ] = − R d   ln ( ε ˙ T G n - 1 ) d   ( 1000 T ) (3) where G is the shear modulus and D is the diffusion coefficient. The activation energy (Q) of Eq. 3 can be deduced from Eq. 2. According to Arrhenius equation, the activation energy can be calculated by the Arrhenius plot of logarithmic ε ˙ or ε ˙ T G ( n − 1 ) vs. the reciprocal Kelvin temperature 1,000/T that provided constant stress.

The activation energies of Ti-6.6Al-4.6Sn-4.6Zr-0.9Nb-1.0Mo-0.32Si alloy under 150–300 MPa are calculated based on Eq. 3. The curve of logarithmic ε ˙ of Ti-6.6Al-4.6Sn-4.6Zr-0.9Nb-1.0Mo-0.32Si vs. the reciprocal Kelvin temperature 1,000/T are shown in Figure 6. The activation energies of Ti-6.6Al-4.6Sn-4.6Zr-0.9Nb-1.0Mo-0.32Si alloy are 366 kJ/mol, 354 kJ/mol, and 337 kJ/mol under 150 MPa, 200 MPa, and 300 MPa, respectively, which is much higher than that of α Ti self-diffusion 240 kJ/mol (Malakondaiah and Rama Rao, 1981) and close to Al diffusion activation energy of α Ti (329 kJ/mol) (Köppers et al., 1997). The activation energy values fluctuate slightly with variation of stress.

Creep rupture lives of Ti-6.6Al-4.6Sn-4.6Zr-0.9Nb-1.0Mo-0.32Si alloy were tested at 873 K, 923 K, and 973 K, respectively. Figure 7 shows the stress-creep rupture life of Ti-6.6Al-4.6Sn-4.6Zr-0.9Nb-1.0Mo-0.32Si alloy on a double logarithmic scale. The creep rupture life deceases with temperature or applied stress increasing. There is a linear correspondence between the logarithmic applied stress and the logarithmic creep rupture lifetime. An inverse linear relationship is found between the creep rupture life and the creep rate by Monkman and Grant, which is named as the Monkman-Grant relationship (Ali and Tamin, 2013). A constant (M ∼ 0.09) is obtained. A constant of Ti-6.6Al-4.6Sn-4.6Zr-0.9Nb- 1.0Mo-0.32Si alloy calculated by Eq. 4 is 0.75∼0.89. Creep rupture lives depends on the creep deformation rate, indicating that creep performance failure is controlled by strain (Evans et al., 1996). tf × ε • = M ( ∼ 0.9 ) (4) Figure 8 shows the SEM microstructure near the fractography after tensile and after creep. A few circular micro cavities and intact continuous α phases on the grain boundaries are observed after tensile in Figure 8A. However, Figure 8B displays many circular micro cavities and few intact continuous α phases on the grain boundaries after creep. The fracture mechanism of Ti-6.6Al-4.6Sn-4.6Zr-0.9Nb-1.0Mo-0.32Si alloy after creep is different from that after tensile. The circular microcavities on the grain boundaries grow and combine with each other to form cracks during creep, which indicates a typical intergranular fracture mechanism. The grain boundary strength decreases during high temperature creep.

It is well known that creep mechanisms of near α titanium alloys include glide of dislocation and grain boundary sliding at high temperature (Pan et al., 2009). As the self and impurity diffusion in α Ti is lower than that in β Ti, deformation of α phase plays a dominant role in the process of creep. The dislocation motions are blocked in α phase. Figure 9 shows transmission electron microscopy (TEM) of dislocation motions in both grain boundaries (Figure 9A) and α lath (Figure 9B) after creep at 923 K with applied stress of 200 MPa. The microstructure of Ti-6.6Al-4.6Sn-4.6Zr-0.9Nb-1.0Mo-0.32Si alloy after heat treatment is Widmanstätten microstructure with several colonies in β grains as shown in Figure 1. The α+β laths present an almost identical crystalline orientation with each other in α+β colonies, which has a great effect on the creep properties in Ti alloys. The Burgers orientation relationship is followed between α + β laths in Widmanstätten and basket microstructure, i.e., the (110) [111] slip system of β phase is coplanar with the (0001) [1,120] slip system of α phase (Miller et al., 1987; Zheng et al., 2021). The dislocation slip is not effectively blocked by interfaces of α+β laths in each colony, so dislocations can easily occur across the interfaces of α+β laths and the entire colony. However, each various colony has a unique relative orientation and the interface between colony can effectively prevent sliding transport (Bhattacharyya et al., 2003). Thus, the size of α+β colonies has a great influence on the creep properties of Ti alloy. The α+β colony size will be reduced with the decrease of β grain size, therefore, the α+β colony size can be effectively controlled by controlling the β crystal size (Yang et al., 2015). Moreover, small β grains can introduce more boundaries to serve as the barriers for dislocation transmission, which could influence the mechanical properties of Ti alloy (He et al., 2019). The creep rate becomes higher when the sizes of average grains become smaller.

The precipitated particles can also hinder dislocation slip and increase creep resistance. Figure 10 shows the precipitation of Ti₃Al in Ti-6.6Al-4.6Sn-4.6Zr-0.9Nb-1.0Mo-0.32Si alloy after creep at 923 K with stress of 200 MPa. As shown in Figure 10, the Ti₃Al precipitates are small particles distributed uniformly in Ti-6.6Al-4.6Sn-4.6Zr-0.9Nb-1.0Mo-0.32Si alloy. The size of Ti₃Al particles becomes larger with increasing of creep temperature and time. The moving dislocations cut Ti₃Al particles in process of creep (Zheng et al., 2021), as shown in Figure 10. The shear stress along the active slip planes locally reduces as in process of deformation, which leads to planar slip. A large number of dislocations are concentrated at grain boundaries which will stop most of the dislocation from further transmission (Li et al., 2011a; Li et al., 2011b). Therefore, grain boundaries, interfaces, and precipitates are important influencing factors for creep properties.