Transmission electron microscopy investigation of the dislocation structure in a Ti-6Al-4V alloy subjected to an early stage of cyclic deformation

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Highlights

  • The α phase dislocation structure is similar to that formed during monotonic straining

  • Prismatic and basal slip are the dominant deformation modes

  • The prevalent <a>−type dislocations have a large screw component while rare <c+a> dislocations are of a mixed type

  • The β phase dislocations are associated with high Schmid factors and possess a large screw component

Abstract

A detailed investigation was conducted of the dislocation structure formed within both the α and β phase of a Ti-6Al-4V alloy subjected to an early stage of cyclic deformation. The cyclic loading-unloading was carried out in tension between applied stress levels of 1100 and 100 MPa, the macroscopic yield stress (σ02) being about 1205 MPa, at a strain rate of 2×10−3 s−1 to 300 cycles, as part of an in-situ synchrotron experiment. The dislocation type, arrangement and density were examined post-mortem using transmission electron microscopy, complemented by the automatic determination of individual crystallite orientations through precession-enhanced nanobeam diffraction. The deformation took place entirely through dislocation glide and no evidence of deformation twinning was found within either of the constituent phases. The dislocation structure within the α phase was fairly similar to that formed during monotonic straining and no presence of periodic planar dislocation networks suggested in the literature was detected. Prismatic and basal slip were the dominant observed deformation modes, based on the calculated Schmid factor and critical resolved shear stress values reported in the literature. The <a>−type dislocations present after straining typically displayed a large screw component, with a majority of them having a pure screw character. There was also limited presence of <c+a> dislocations within the α phase and these were largely of a mixed type. Burgers vectors of the dislocations remaining after deformation within the β phase were also largely associated with the potential slip systems having high Schmid factor values and these dislocations mainly displayed a large screw component.

Introduction

Deformation behaviour of a two-phase (α+β) Ti-6Al-4V alloy is mainly controlled by the hcp α phase as the bcc β phase content is comparatively low (typically 5–10%) [1,2]. Several deformation modes (slip and twinning) can operate during plastic deformation of hcp titanium and its alloys [3]. The slip may occur on the prismatic {101¯0}, basal {0001} and first-type pyramidal {101¯1} planes with the Burgers vector <a> = 13<112¯0>. Deformation along the c-axis might be facilitated by slip with the Burgers vector <c+a> = 13<112¯3>, typically taking place on the first-type pyramidal planes, or by several twinning modes [3]. In contrast to α titanium for which deformation twinning has been frequently reported to occur during monotonic loading [1,2,4], in commercial polycrystalline titanium alloys, such as Ti-6Al-4V, twin formation is extremely restricted [1,2,5,6]. The activation of a particular deformation mode is generally governed by the corresponding critical resolved shear stress (CRSS) in conjunction with the Schmid factor [7]. Although a favoured monotonic deformation mode in α titanium is prismatic slip [8,9], an increased aluminium content in titanium alloys causes a reduction in the CRSS for basal slip to a level comparable with that for prismatic slip [5,6,[10], [11], [12], [13], [14]]. This results in enhanced basal slip activity in these alloys. It has been widely accepted that the <c+a> glide is associated with a markedly higher CRSS than the pure <a> deformation systems [5,10] and, thus, its operation is comparatively difficult. The <a> type dislocations remaining after monotonic deformation of α titanium and its alloys are largely of a screw type, which indicates that it is the edge and mixed dislocations that move significantly faster during straining [5,6,10]. This has indeed been confirmed by the in-situ transmission electron microscopy (TEM) studies [9,[15], [16], [17]]. The low mobility of <a>−type screw dislocations is attributed to their three-dimensional core structure [9,15,16,18].

Fatigue deformation of α titanium takes place through wavy slip, due to its relatively high stacking fault energy, and the microstructure at low stress/strain amplitudes is characterised by homogeneous distributions of curved dislocations that gradually evolve into dislocation cell structures dominating the saturation state [19,20]. The formation of such dislocation structures is accompanied by profuse deformation twinning, similar to monotonic straining [1,2,4]. By contrast, in complex titanium-based alloys, slip planarity is markedly increased. This phenomenon is not only a consequence of lowering the stacking fault energy, but it is also influenced by the shear modulus, atomic size misfit, solute content, short-range order, and dispersion particles. Thus, different alloys might display significant differences in the fatigue behaviour. In addition, the crystal orientation has also an effect on the fatigue deformation mode and dislocation configuration in titanium alloys. Xiao and Umakoshi [[21], [22], [23]] have reported that the dislocation structure evolution during fatigue of Ti-5 at.% Al single crystals oriented for single or double prism slip starts with the homogeneously distributed screw dislocation arrays or localised planar dislocation bands on the primary glide plane and homogeneous dislocation distribution on the secondary glide plane, respectively. Such evolution continues in both cases with the gradual development of dense dislocation bundles aligned along the [0001] direction until a well-developed “saturation bundle structure” (SBS) is finally formed at the fatigue saturation stage. Dislocations in commercial titanium alloys subjected to fatigue deformation also tend to arrange in localised planar slip bands, predominantly formed on the prismatic and basal glide planes, which appear to evolve towards dislocation bundles at higher stress/strain amplitudes [[24], [25], [26], [27]]. It should be noted that the studied Ti6246 [24] and Ti6242 [26,27] alloys contained ordered Ti3Al precipitates the presence of which markedly enhances the planarity of dislocation slip. Both the homogeneously distributed dislocations and those located within the planar slip bands are largely screw in character [[21], [22], [23], [24], [25], [26], [27]], similar to monotonic straining [5,6,9,10,[15], [16], [17]]. As expected [1,2,5.6], deformation twinning is generally limited in commercial titanium alloys subjected to fatigue deformation, but some formation of {101¯2} extension twins has recently been observed within the grain boundary regions of hard (low Schmid factor) grains in a Ti-6Al-4V alloy [28,29]. It has been suggested that such twin formation was triggered by the stresses generated by the dislocation pileups in neighbouring soft (high Schmid factor) grains impinging on the grain boundaries.

Thus far, there have been few studies focused on the deformation processes occurring in the β phase in two-phase titanium alloys. It has been reported that, in these alloys, slip within the β phase can typically take place on {110} 〈111〉 and {112} <111> systems with possible cross-slip from {110} to {112} slip planes [30,31]. The β phase is generally considered softer than its α counterpart [6], as it possess a higher number of slip systems, and the ease of slip transfer across the interface from the α phase largely depends on a deviation from the exact Burgers orientation relationship between the two phases [[30], [31], [32]]. The presence of the β phase, which can take over deformations along the c-axis of the α lattice, may also be beneficial for the ductility of two-phase titanium alloys [1,2,6].

In spite of several publications on the α phase dislocation structures produced by fatigue deformation in commercial titanium alloys [[24], [25], [26], [27]], the published reports describing such structures within the constituent phases in a Ti-6Al-4V alloy, despite its high technological importance, have thus far been very limited. To the best of the authors' knowledge, the only currently available systematic investigation of the fatigue microstructures in a Ti-6Al-4V alloy at the dislocation level is that recently published by Jha et al. [33] that focuses solely on the α phase. In addition, some of the findings of the above study appear to be rather controversial, e.g. the formation of well-developed planar dislocation networks was surprisingly reported already at the very early stage of the fatigue deformation. Such profuse dislocation network formation has never been reported in other titanium alloys subjected to fatigue deformation at ambient temperature. In order to expand current understanding of the fatigue dislocation structures in the Ti-6Al-4V alloy as a function of loading conditions, the present work aims to investigate the dislocation character, density, and arrangement, as well as the corresponding slip systems, within both the α and β phase of a Ti-6Al-4V alloy subjected to cyclic tensile loading-unloading deformation. The straining is performed under conditions corresponding to an early stage of high cycle fatigue loading as part of an in-situ synchrotron experiment. The obtained dislocation characteristics will be used for the validation of the related outputs of the newly developed Bayesian analysis of the synchrotron high-energy X-ray data [34], with the aim to improve model predictions for crack initiation in aero-propulsion components.

Section snippets

Experimental procedures

The material used in the present work was a commercial (AMS 4928) Ti64 alloy supplied by the Airport Metals Australia having a nominal composition of 6 wt% Al, 4 wt% V and the balance Ti. The material was received in a form of a 12.7 mm diameter rod in an annealed condition. The crystallographic texture of both the α and β phase was determined by X-ray diffraction (XRD) using a PANalytical X'Pert PRO MRD texture diffractometer. Pole figures were measured for several major reflections of both

Starting material

Fig. 1 shows the crystallographic texture of both the constituent phases in the starting rod material represented by the inverse pole figures for the rod axis direction. The maximum intensity corresponding to the α phase was about 2.6 times random and it was concentrated around the [11¯00] crystal lattice direction. The orientation distribution for the β phase displayed two maxima centered at the [001] and [011] lattice directions and having intensities of about 2.1 and 2.3 times random,

α phase

It should be emphasised that all the cyclic deformation of the α phase occurred entirely through dislocation slip and no indication of deformation twinning was observed within this phase during the present work. This is in good agreement with the published literature, as the deformation twin formation has been reported in the Ti-6Al-4V alloy only very rarely for localised small areas, typically within grains having loading axes close to the [0001] lattice direction and for the late stages of

Conclusions

1. It has been observed that the α phase deformation took place entirely through dislocation slip and no evidence of deformation twinning was found within this phase.

2. Dislocation structure formed during cyclic loading was principally similar to that formed during monotonic deformation, apart from the local formation of dense dislocation bundles, which is consistent with the studied early stage of cyclic straining, and no presence of periodic planar dislocation networks suggested in the

Data Availability Statement

The raw/processed data required to reproduce these findings cannot be shared at this time as the data also forms part of an ongoing study.

Declaration of Competing Interest

The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.

Acknowledgements

The authors acknowledge funding from Defence Aviation Safety Authority (DASA), Australian Department of Defence, via Defence Science Technology, research agreement. The Deakin University's Advanced Characterisation Facility is acknowledged for provision of the research facilities used in the present work. The authors would also like to thank Ms. Madeleine Burchill (Maritime Division, DST) for her feedback regarding the analysis.

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    • Dislocation structures in a Ti–6Al–4V alloy subjected to cyclic tensile deformation

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      Citation Excerpt :

      Only very few dislocations were observed to be of an edge or a mixed type. The above findings are consistent with those made by the current authors during their recent study of the Ti–6Al–4V alloy subjected to the early stage of cyclic deformation [35]. The observed dominant presence of the screw-type dislocations within the β phase deformed microstructure might be expected, based on the general deformation behaviour of the bcc lattice materials [29–34].

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