Experimental study of dislocation mobility in a Ti–6Al–4V alloy
Introduction
Titanium alloys are widely used because of their excellent mechanical properties combined with low density. The most used of titanium alloys in the aerospace industry is Ti–6Al–4V [1]. It is a two-phase alloy with an α phase (hexagonal close-packed) and a β phase (body-centered cubic). Its microstructure and thus its mechanical properties are very dependent on the thermomechanical processes performed during its production. The microstructure can be fully nodular, fully lamellar or duplex with nodules and lamellar colonies.
The deformation of pure α-titanium has already been widely studied [2], [3], [4], [5], [6], [7], [8]. Two types of Burgers vectors are possible for gliding dislocations: a-type , which can glide in prismatic, basal and first-order pyramidal planes, and c+a-type , which can glide in first- and second-order pyramidal planes. Some twinning systems can also be activated, but they are essentially observed at room temperature in low oxygen content alloys [8]. In α-titanium, prismatic glide of a-type dislocations is the main deformation mode and the deformation is controlled by the motion of the screw segments [2], [3], [4], [5], [6], [7]. Atomistic calculations show that the high lattice friction on screw segments is due to their three-dimensional core structure spread in several planes [9], [10], [11], [12], [13], [14]. Indeed, screw dislocations have a stable and sessile configuration which must recombine in the glide plane into a glissile and metastable configuration to be able to glide.
The deformation micromechanisms have been less studied in two-phase and polycrystalline alloys like Ti–6Al–4V principally because of their more complex microstructure. This complexity is due to the presence of two phases (eventually more) and the presence of different type of grains – nodules and lamellar colonies – in which the α phase can have different chemical composition and different dislocation behaviour. The contributions to the strength of the alloy are then more numerous, the main ones already studied being:
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the presence of short range order, especially observed in nodular alloys [15], [16], [17], [18];
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the presence of interfaces, particularly in lamellar alloys or colonies [19], [20], [21], [22].
In the present paper, we present a detailed study of the deformation micromechanisms in a Ti–6Al–4V alloy. We focus especially on the origin of the strength in lamellar colonies using in situ transmission electron microscopy (TEM) deformation experiments at room temperature, which is the only available technique for studying the dislocation dynamics under stress and temperature. This technique has been previously performed only on Ti–6Al–4V alloy in cryogenic conditions at 20 K [20] or on pure α-titanium [6], [7], but never on other titanium alloys.
Section snippets
Experimental
The alloy under investigation is an industrial Ti–6Al–4V compound with 6% aluminium by weight and 4% vanadium by weight. Other elements are present as impurities (mainly oxygen, nitrogen, carbon and iron).
It has a duplex microstructure (Fig. 1), with primary alpha nodules αP and lamellar colonies αS/β. The nodules and lamellar colonies have the same size of about 10 μm. Lamellar colonies consist of secondary alpha plates αS with an average thickness of about 500 nm separated by thin β plates (
Results
The main micromechanisms of the plastic deformation of this alloy were identified during in situ experiments. All observed gliding dislocations have an a-type Burgers vector. In lamellar colonies, as in nodules, we essentially found basal and prismatic glide and, less frequently, first-order pyramidal glide. In the three types of planes, dislocations generally had behaviour similar to that detailed here.
The first sequence (Fig. 3) shows the emission and propagation of a dislocation loop in an αS
Discussion
The TEM in situ technique allows information to be obtained on the dislocation dynamics. To avoid possible thin foil effects, it is useful to compare these results with post-mortem observations, exemplified in Fig. 8. We have observed several post-mortem samples and our in situ observations are in agreement with the bulk alloy deformation. Moreover, in situ experiments give results about the chronology of events and about the dynamics of dislocations contrary to post-mortem observations, which
Conclusion
In situ TEM deformation experiments have allowed the observation of the dislocation dynamics and the detailed study of the deformation micromechanisms in a titanium alloy with a complex microstructure.
All gliding dislocations have an a-type Burgers vector. Due to their core structure, they are preferentially aligned with their screw direction, resulting in a large density of rectilinear long screw dislocations. This core structure is almost certainly spread in the basal, prismatic and
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