Dislocation interactions in near-alpha Titanium alloy Ti6242Si under LCF

Dislocation interactions were investigated in near-alpha titanium alloy Ti6242Si after low cycle fatigue. Samples from the gauge section and the crack initiation site on the fracture surface were studied. Grain pairs with different crystallographic orientations were analysed from the gauge section to understand the dislocation interactions near the boundary. Deformation was primarily caused by planar slip, localized into slip bands in primary alpha (αp) grains. Direct slip transfer was observed within grains having similar orientations. In contrast, slip transfer resulted in a different kind of dislocation nucleation in the neighboring hard grain where there was misorientation between the grains, with the observation of cross-slip. Strain transfer was observed between highly misoriented grain pairs. Crack nucleation occurred on an αp grain by basal plane splitting, due to the large tensile stress developed by a double ended pile-up. This kind of pile-up is suggested to result from the incomplete reversibility of dislocation motion during load reversals. The observation of superjogs on the basal dislocations in the crack nucleated αp grain provides a rationale for why cracks nucleate near, rather than on, the basal plane.


Introduction
Near-α titanium alloys such as Ti6242Si are employed in high temperature aero-engine compressors due to their combination of specific toughness, creep resistance and microstructural stability at temperatures around 400-600°C [1][2][3][4]. The component is exposed to a variety of loading regimes during service, for e.g., high stresses in the low cycle fatigue regime may be experienced at the thrust peak associated with take-off. This continues to compromise design [5,6] and hence is very important to understand. Further, the Ti alloys are generally notch sensitive and so crack initiation from the microstructure is often of greater concern than in other alloy systems where lifing can mostly be considered in terms of crack growth.
Furthermore, it is believed that hard/soft grain pairs act as fatigue initiation sites [12,13] and this occurs by the formation of near basal facets in the hard grains, due to load shedding by adjacent soft grains [16], [20][21][22][23]. The lack of slip transfer from the soft to hard grain will produce a sufficient stress concentration to result in the initiation of a fatigue crack in the hard grain; conversely, easy slip transfer in the microtextured regions [24] will result in homogenous plastic deformation.
In lamellar microstructures fatigue cracks initiate either within slip bands in α lamellae or along prior beta grain boundaries [25,26]. In bimodal microstructures, cracking can initiate either on the interface between the lamellar microstructure and the α p or within the α p itself [27]; initiation also depends on the fraction and size of the α p . In near-α alloys, such as Ti6242, with more than 50% α p , crack nucleation is expected to occur in the α p due to extensive grain boundary dislocation pile-up [27,28].
In the present work, the dislocation structures produced during low cycle fatigue in Ti6246Si are investigated in detail, both in the gauge section and the fracture surface of the sample, with the

Material and experiments
The near alpha titanium alloy Ti-6Al-2Sn-4Zr-2Mo-0.1Si was studied in this work. Mechanical tests were conducted using Mayes servohydraulic machines with Instron 8800 controllers. Static tensile tests were carried out on dog bone shaped samples at a strain rate of 10 -3 s -1 . Low cycle fatigue (LCF) tests were carried out on cylindrical samples with a 2.9 mm diameter and a 15 mm long gauge length. The tests were performed using a trapezoidal waveform with a ramp up/down time of 1 s, a 1 s hold at maximum stress of 95% of yield stress and an R ratio of 0.05.
Dislocation analysis on LCF samples was conducted using a JEOL JEM-2100F TEM/STEM with an accelerating voltage of 200 kV. 0.5 mm thick discs were cut from the gauge of the failed samples normal to the loading direction, ground to 100 to 150 mm using SiC paper and electropolished using 3% perchloric acid, 57% methanol and 40% butan 1-ol in a Tenupol at -40°C and 24V. TEM samples from the crack initiation site were prepared using the focused ion beam (FIB) lift-out technique in a dual beam FEI Helios NanoLab 600 using a 30 kV Ga ion beam. To protect the area of interest, a gas injection system was used to deposit a Pt-containing protective layer.
Samples were made electron transparent by thinning to 150 nm thickness. To obtain the α p grain orientations, Transmission Kikuchi Diffraction (TKD) was carried out at an accelerating voltage of 30 kV and working distance of 2mm.

Initial microstructure:
3 The alloy had a typical bimodal microstructure with α p in the transformed β, Fig. 1. Secondary alpha (α s ) platelets can be observed in the retained β. The volume fraction of α p was found to be ≈ 50%.  Fig. 3(b). This is due to the pile-up stress since grain 3 is unfavorably oriented for slip transfer with grain 2. This pile-up stress results in slip C in neighboring grain 3.
g.b analysis shows that the dislocations in slip bands A and B are of (a/3) [1 10] type gliding on prism plane (10 0), and that the dislocations in slip band C are of (a/3) [11 0] type, gliding on the basal plane. Direct slip transfer occurred across the boundary between grains 1/2, where 2> prism slip in grain 1 activated the same type of slip in grain 2. In contrast, strain transfer occurred across the boundary between grains 2/3, where one type of slip in grain 2 activated another type of slip in the neighboring grain 3. grain, having its c-axis oriented ≈55° to the loading axis, and a hard grain with its c-axis oriented ≈ 15° from the loading axis, Fig. 4 (a). The crystallographic orientation of the grain combination was obtained using Transmission Kikuchi Diffraction (TKD) of the TEM foil. The grain boundary is indicated. The orientation relationship of the pair was identified from selected area diffraction (SAD) patterns from these grains when one of the grains was tilted to a zone axis. The SAD patterns of the soft and hard grains are shown in Fig. 5

Fig. 5 SAD patterns from (a) soft and (b) hard grain when the hard grain was tilted to the [1 13] zone
The slip transfer across the soft/hard grain boundary can be seen in Fig. 4(a). Dislocations in the soft grain is piled-up against the grain boundary, Fig. 4(b). These are type dislocations with where ϕ is the angle between the normal to the slip planes and κ is the angle between the slip directions in the two grains. This m' value was calculated using Matlab toolbox STABiX [29],  Fig 3(d). This cross slip was observed between first and second order pyramidal planes. It is believed that cross slip is responsible for dislocation multiplication and patterning in fatigue [30]. Further, the mobility of dislocations on the cross slip planes is lower than on the primary planes [31].  Fig. 4

Dislocation interactions on the fracture surface:
We now turn to the dislocation interactions in the foil lifted-out near the crack initiation site. The foil contained an α p grain bounded by two-phase regions on either. Intense planar pile-ups (Fig.   6a) and dislocation loops (Fig. 6b) Fig. 6b. It can be noted that the mixed component of the loops are found to pile-up in the boundary between α s + β and α p , Fig.   6b. It was also observed that the slip bands in this grain did not transfer across the boundary between the α p grain and the two phase region, even though they are intense and large.
Based on these observations the following crack nucleation mechanism is proposed in this case.  [32]. Incomplete reversibility of dislocation motion could be realized if the lattice friction stress against dislocation motion is higher in the reverse direction than in the forward direction [32]. This could be rationalised as follows.  [32][33][34]. This kind of pile-up is possible only when the dislocation sources are located very close to each other, because for an isolated dislocation source, dislocation motion on a plane becomes reversible as no back stress is produced by the dislocations on a neighbouring plane. Therefore, the reversibility or otherwise of dislocation motion will depend on the statistical distribution of dislocation sources.
Then the embryonic crack might have nucleated from this double ended pile-up as follows. The stress field in this case would be the sum of applied stress and the internal stresses due to the dislocation pile-ups on successive slip planes. In addition, the mixed components of the basal dislocation loops were also found to pile up near the boundary, causing further stress intensification.
The stress distribution due to this double ended pile-up together with the observed dislocation structures in the α p grain is shown schematically in Fig. 6c σ xx = 3 a n(Δτ − 2κ)/ 2h ………….. (2) where 2a is the grain size, n is the number of stress cycles, Δτ is the cyclic shear stress applied to the grain, κ is the lattice friction stress and h is the slip plane interplanar spacing.
The tensile stress would split the atomic planes along the dislocation lines [35], inset of Fig. 6c.
Therefore the large tensile stress resulting from the double ended pile-up appears to be the mechanism responsible for crack nucleation in the α p grain. From the TEM foil, it was observed that the fracture plane is found to be oriented 5-8° from the basal plane. This near-basal plane cracking rather than true basal plane cracking might be a consequence of the observed jogging of the basal dislocations.

Conclusions
The dislocation interactions of near alpha titanium alloy Ti6242Si were studied after low cycle fatigue. The investigations were carried out on both gauge section and the crack initiation site of the sample and the following conclusions were made.
The alloy mainly deforms by planar slip in α p grains and the two phase region α s +β acting as a barrier for slip transmission.
Direct slip transfer was observed between similarly oriented grains and slip transfer resulted in nucleation of different kind of dislocations when there is misorientation between grains.
Strain transfer occurred across the boundary between grains when the misorientation is high, where one type of slip in grain activated another type of slip in the neighboring grain.
Double ended pile-up in an α p grain is responsible for crack nucleation in this alloy. Such structures can result from the incomplete reversibility of dislocation motion due to a slightly higher friction stress in the reverse direction than in the forward direction.
This kind of double ended pile-up results in a large tensile stress, providing a mechanism of crack initiation by plane splitting along the dislocation line. The observation of superjogs on the basal dislocations provides a rationale for why such cracks nucleate near rather than on the basal plane.