Microstructure evolution during deformation

After reheating the martensitic specimens at the deformation temperature and held for 180 s, no obvious microstructural changes like lath coarsening, as reported in [96], were observed in the Į' lath structure. However, the current thermomechanical processing resulted in remarkable microstructural changes in the martensitic Ti-6Al-4V alloy (i.e. grain refinement, ȕ phase formation and change in the misorientation angle/axis distribution), as shown in Figs. 4-6 to 4-8.

Figure 4-6: EBSD images of the martensitic Ti-6Al-4V alloy deformed at 700°C and 0.001 s^-1 to different strains: a) 0.2, b) 0.4, c) 0.6 and d) 0.8. Black, red, white and green lines represent Į-Į grain boundaries with a misorientation angle (T) greater than 15°, 5°<T<15°, 2°<T<5° and the intervariant interfaces having Ȉ13b=57.4q/<11ʹത0> twin characteristics, respectively. The ȕ phase is shown in blue.

The black arrows represent the compression direction.

As aforementioned, the initial martensitic structure mostly revealed high angle grain boundaries (HAGBs), with a misorientation angle distribution quantitatively

different from a randomly textured hexagonal close packed (hcp) material (i.e. the Mackenzie distribution [232]). Multiple peaks were observed across the misorientation angle distribution, mainly centering about the 10q, 60q, 63q, and 90q misorientation angles (Fig. 4-2c), which are similar to the boundary population in a martensitic structure formed from pure titanium through the crystallographic variant selection mechanism during the martensitic phase transformation [177]. It appeared that most intervariant interfaces had twin characteristics (i.e. 613b= ) with the {10 11} interface plane [230] similar to the mechanical twins formed in the primary laths (Fig. 4-3).

Figure 4-7: a) Misorientation distribution of the Ti-6Al-4V alloy deformed under different thermomechanical conditions. b-d) The characteristics of misorientation axis distribution corresponding to the misorientation angle distribution at different strains: b) 0.2, c) 0.4 and d) 0.8. The relative misorientation frequency was plotted using a bar width of 0.5°. The initial locations of different types (I-V) of intervariant misorientation axis in the non-deformed martensitic microstructure are marked by circles on the standard stereographic triangle in each figure.

The deformation revealed a significant effect on the misorientation characteristics, progressively altering the initial lamellar structure to ultrafine equiaxed grains with HAGBs (Figs. 4-6 and 4-7). At an early stage of deformation, there was a significant increase in the population of low angle grain boundaries (LAGBs), mainly

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in the middle of martensitic laths (Figs. 4-6a and 4-7a, b). Beyond a strain of 0.4, the peak intensity of LAGBs gradually decreased and became wider (Fig. 4.7c). The multiple peaks at about the 60q and 90q misorientation angles gradually merged with an increase in the strain forming a single peak. At the same time, the peaks broadened and their intensities were gradually being reduced (Fig. 4-7c). As a result, the misorientation angle distribution at a strain of 0.8 became nearly uniform for misorientation angles greater than 15q and hardly any peak was distinguishable (Fig.

4-7d). This suggests that most of the LABs were progressively converted to high angle boundaries with increasing strain.

Figure 4-8: Standard stereographic triangles of the rotation axes in crystal coordinate system for 613b=60q/<11ʹത0> intervariant boundaries at 700°C and 0.001 s^-1 at different strains: a) 0, b) 0.2, c) 0.4 and d) 0.8. The MRD is multiples of a random distribution.

Simultaneously, the misorientation axis of most martensitic intervariant interfaces/boundaries also lost their unique characteristics with strain (Fig. 4-7). As stated earlier, the undeformed martensitic microstructure was characterized by the Burgers orientation relationship (OR). As a result, the misorientation axis exhibited different clusters located at various positions on the stereographic triangle (Fig. 4-2c).

With increasing strain, these clusters became scattered continuously (Figs. 4-7b, c, d), indicating a deviation from the ideal Burgers OR. The spread of misorientation axis can be attributed to the activation of different slip systems in neighbouring Į' laths

through the interaction of dislocations with intervariant interfaces. In addition, the population of 613b corresponding to 60°/ gradually decreased (Fig. 4-6) as the misorientation axis deviated progressively from the axis with increasing strain (Fig. 4-8).

Corresponding to the dramatic changes in the misorientation profile, a progressive break-up of Į' laths into ultrafine equiaxed grains can be observed. In other words, the volume fraction of ultrafine equiaxed grains gradually increased with deformation at the expense of the lamella structure, leading to a nearly fully equiaxed grain structure at high strains (Figs. 4-6 and 4-9a). At a strain of 0.2, the newly formed refined grains had an average grain size of ~300 nm and their formation was found to be related to the alignment of Į' laths with respect to the compression axis (Fig. 4-6a).

It appears that the laths along the loading axis are preferable for fragmentation to take place. These types of laths kept experiencing shear strain when they were tilted towards the specimen radius direction. At a later stage of deformation (i.e. a strain of 0.4), a significant increase in the volume fraction of fine equiaxed grains was observed.

Only a small proportion of alpha laths were retained in the deformed microstructure, with their length direction perpendicular to the loading axis. Closer microstructural examination around the strained laths (Fig. 4-6b) revealed quite a number of low angle grain boundaries forming from the intervariant interfaces into the retained alpha laths.

This indicates a progressive fragmentation of alpha laths by interaction between dislocations and boundaries. At a deformation strain of 0.6, the initial lath structure was nearly fully broken up into globular segments ranging from 300 to 800 nm (Fig.

4-6c). Inside the few strained and highly kinked laths, numerous low angle grain boundaries were observed, as a result of rearrangement of dislocations. Finally, at a strain of 0.8, the initial lath microstructure was nearly fully replaced by fine equiaxed grains characterized by being dislocation free and mostly high angle boundaries (Fig.

4-6d).

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Figure 4-9: a) The volume fraction and average size of equiaxed grains as a function of strain at 700°C and 0.001 s^-1. The grains with an aspect ratio of less than 2.5 were considered as equiaxed using the grain detection function in HKL Channel 5 software. b) XRD patterns of Ti-6Al-4V alloy for starting martensitic structure (Martensite) and deformed structure (Deformed) at 700°C, strain of 0.8 and strain rate of 0.001 s^-1.

Along with warm compression to a strain of 0.8, the lath structure progressively became fragmented into very fine equiaxed grains with a size of ~150-800 nm depending on the thermomechanical condition (Fig. 4-6). It is noteworthy that E phase also appeared in the microstructure during the deformation. The presence of E phase was confirmed by both EBSD and XRD analysis [233] (Figs. 4-6 and 4-9b). This suggests that the microstructure may partially transform to E through adiabatic heating and/or phase transformation induced by deformation, which will be discussed in detail later.

4.3.4 The effect of thermomechanical parameters on UFG