Transmission electron microscopy investigation of the dislocation structure in a Ti-6Al-4V alloy subjected to an early stage of cyclic deformation
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 {}, basal {0001} and first-type pyramidal {} planes with the Burgers vector <a> = <>. Deformation along the c-axis might be facilitated by slip with the Burgers vector <c+a> = <>, 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 {} 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 [] 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.
References (47)
- et al.
Perspectives on titanium science and technology
Acta Mater.
(2013) - et al.
Stress-state dependence of slip in titanium-6Al-4V and other HCP metals
Acta Metall.
(1981) A study of active deformation systems in titanium alloys: dependence on alloy composition and correlation with deformation texture
Mater. Sci. Eng. A
(2003)- et al.
Analysis of the different slip systems activated by tension in an α/β titanium alloy in relation with local crystallographic orientation
Acta Mater.
(2005) - et al.
Combination of in-situ SEM tensile test and FFT-based crystal elasticity simulations of Ti-6Al-4V for an improved description of the onset of plastic slip
Mech. Mater.
(2017) - et al.
Crystal plasticity modeling of slip activity in Ti–6Al–4V under high cycle fatigue loading
Int. J. Plasticity
(2009) - et al.
Incipient slip and long range plastic strain localization in microtextured Ti-6Al-4V titanium
Acta Mater.
(2016) - et al.
An in situ study of prismatic glide in α titanium at low temperatures
Acta Metall. Mater.
(1993) - et al.
Experimental study of dislocation mobility in a Ti–6Al–4V alloy
Acta Mater.
(2007) - et al.
In situ transmission electron microscopy deformation of the titanium alloy Ti–6Al–4V: Interface behaviour
Mater. Sci. Eng. A
(2008)
Influence of simple metals on the stability of <a> basal screw dislocations in hexagonal titanium alloys
Acta Mater.
Orientation dependence of cyclic deformation behavior and dislocation structure in Ti-5at.% Al single crystals
Mater. Sci. Eng. A
Low cycle fatigue behavior of an α+β titanium alloy: Ti6246
Mater. Sci. Eng. A
Low-cycle fatigue behavior and deformation substructure of Ti–2Al–2.5Zr alloy at 298 and 673 K
Mater. Sci. Eng. A
Slip transfer and deformation structures resulting from the low cycle fatigue of near-alpha titanium alloy Ti-6242Si
Int. J. Plast.
Dislocation interactions and crack nucleation in a fatigued near-alpha titanium alloy
Int. J. Plast.
Concurrent operation of <c+a> slip and twinning under cyclic loading of Ti-6Al-4V
Scr. Mater.
Analysis of deformation mechanisms operating under fatigue and dwell-fatigue loadings in an α/β titanium alloy
Int. J. Fatigue
Room temperature deformation and mechanisms of slip transmission in oriented single-colony crystals of an α/β titanium alloy
Acta Mater.
On the process of transition of the cubic-body-centered modification into the hexagonal-close-packed modification of zirconium
Physica
Effect of strain amplitude on low cycle fatigue and microstructure evolution in Ti-6Al-4V: a TKD and TEM characterization
Mater. Charact.
Slip and fatigue crack formation processes in an α/β titanium alloy in relation to crystallographic texture on different scales
Acta Mater.
Dislocation structures representing individual slip systems within the α phase of a Ti–6Al–4V alloy deformed in tension
Mater. Sci. Eng. A
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2022, Materials Science and Engineering: ACitation 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].