In situ TEM observations of dislocation dynamics in α titanium: Effect of the oxygen content

Materials Science & Engineering A 703 (2017) 331–339

Contents lists available at ScienceDirect

Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea

In situ TEM observations of dislocation dynamics in α titanium: Effect of the oxygen content

MARK

⁎

B. Barkiaa,1, J.P. Couziniéa, , S. Lartigue-Korineka, I. Guillota, V. Doquetb a b

Université Paris Est, ICMPE (UMR 7182), CNRS, UPEC, F-94320 Thiais, France LMS, Ecole Polytechnique, CNRS UMR7649, Université Paris-Saclay, 91128 Palaiseau, France

A R T I C L E I N F O

A B S T R A C T

Keywords: Titanium Plasticity Transmission electron microscopy Oxygen In situ Dislocations

Plastic deformation micro-mechanisms, dislocation structures and glide kinetics in two titanium batches with moderate and high oxygen contents (450 and 3200 wppm, respectively) are investigated by in situ tensile tests in a transmission electron microscope, at room temperature. In both materials, plastic deformation is accommodated with < a > type screw dislocations gliding mainly in prismatic planes. The movement of screw segments is jerky, oxygen-dependent and strongly controlled by pinning on localized obstacles if the oxygen content is high. Dislocation multiplication is mainly controlled by the opening of loops produced by a double-cross slip mechanism at super-jogs. Evidences of composite glide between prismatic, first order pyramidal and basal planes are pointed out, proving the existence of intensive cross slip.

1. Introduction It is well established that the viscoplastic behaviour of α titanium (α-Ti) is strongly dependent on the interstitials’ content (oxygen, nitrogen, carbon, hydrogen), as a result of the interactions between these impurities and gliding dislocations [1]. Dislocations-interstitials interactions lead to different phenomena, i.e. hardening effect, dynamic strain aging [2,3] and have nasty practical consequences in titanium alloys such as cold dwell effect, loosening of bolted assemblies, subcritical crack growth. At microscopic scale, transmission electron microscopy (TEM) observations performed after deformation highlighted the effects of interstitial solutes on dislocations’ arrangement [4]. While dislocations are curved and uniformly distributed in high purity titanium, with a similar proportion of edge and screw segments, the increase in impurities content causes a transition from wavy to planar glide, and to a more heterogeneous microstructure made of long straight screw dislocations with b=a/3 < 1-210 > Burgers vector (so called < a > dislocations) along {10-10} prismatic planes [1,2,4]. The dislocations become more restricted to their glide plane, and their density increases with increasing oxygen content [4]. While the plastic deformation micro-mechanisms have been widely investigated through dynamic observations of dislocations glide in β and α/β Ti alloys, as well as in γ Ti-Al alloys [5–9], only few in situ TEM deformation experiments were conducted on α-Ti, either through

⁎

1

tensile tests [10–14] or nano-compression [15]. They emphasize high Peierls stresses on screw dislocations, a jerky motion of dislocations, and an intensive cross slip activity. Screw dislocations glide is controlled by a locking-unlocking mechanism [12], initially proposed to interpret slip features in beryllium [16]. A tomographic analysis of dislocations interactions in commercially pure titanium reveals that cross slip mainly occurs between the prismatic and pyramidal planes [14]. These interpretations are supported by atomistic computer simulations showing that the most stable configuration involves a sessile core of < a > dislocations spread in pyramidal planes in α-Ti [17]. A core transformation from a sessile to a glissile configuration is thus needed to allow the defects to move along the prismatic planes. The effect of solutes – especially oxygen – on the deformation behaviour of α-Ti has been addressed by in situ experimental studies on single crystals and micro-pillars [12,15]. The presence of oxygen impacts dislocation motion with strong pinning effects and also induces an increase of the critical resolved shear stress (CRSS) for prismatic slip. It has been proposed that the hardening effect results from a modification of the dislocations core structure [11,12,18,19]. High-angle annular dark field scanning (HAADF)-scanning TEM observations confirm this hypothesis, as they directly reveal a substantial change of the core structure with the presence of oxygen atoms close to screw dislocations cores in α-Ti [15]. Recent ab initio calculations shed a new light on the nature of the oxygen-dislocation interactions in hcp Ti and Zr [15,20,21]. These studies evidence a strong repulsion between oxygen

Corresponding author. E-mail address: [email protected] (J.P. Couzinié). Present address: MSSMAT, CentraleSupelec, UMR CNRS 8579, Université Paris-Saclay, 92295, France.

http://dx.doi.org/10.1016/j.msea.2017.07.040 Received 2 June 2017; Received in revised form 14 July 2017; Accepted 15 July 2017 Available online 22 July 2017 0921-5093/ © 2017 Elsevier B.V. All rights reserved.

Materials Science & Engineering A 703 (2017) 331–339

B. Barkia et al.

Table 1 Chemical composition of the two titanium batches. Material

H (wppm)

O (wppm)

C (wtppm)

N (wppm)

Fe (wppm)

Ti (wppm)

Oéq (at. % )

Ti0 T60

5±2 15 ± 3

450 ± 30 3200 ± 100

< 80 70 ± 20

< 30 60

< 200 1750 ± 50

Balance Balance

0.138 1.020

3. Results

atoms in the dislocation core, due to the distortion or destruction of interstitial sites leading to the cross slip of dislocations on adjacent prismatic planes [21]. These results demonstrate that there is a crucial need to bring a new experimental insight on the impact of oxygen solute on dislocations dynamics during plasticity at the appropriate scale. In order to gain a better understanding of the effect of oxygen on the dislocation kinetics and on their motion (slip, cross slip, overcoming of localized obstacles) in α-Ti, in situ straining experiments are run at room temperature in a TEM, on two batches of polycrystalline titanium with different oxygen contents. Slip features are analyzed in crystals with quite different orientations favouring either prismatic slip or basal slip.

3.1. Activated slip systems and dislocation motion In both titanium batches, plastic deformation is achieved mainly by < a > dislocations with a/3 < 11-20 > Burgers vectors. In all the testing conditions, no evidence of π1〈c+a〉 slip is found. In T60, the critical resolved shear stresses (CRSS) of the different slip systems have been determined previously [3] (see Table 2). The CRSS for π1 < a > pyramidal slip is 1.18 times larger than for prismatic P < a > slip. The absence of < c+a > dislocations is thus in agreement with Schmid's law, since the maximum SFs for π1 < c+a > slip are less than 1.80 times those for prismatic slip. Fig. 2 shows successive frames taken from a video (video 1-Ti0-G1) of < a > dislocations with b=a/3[11-20] Burgers vector gliding in grain G1 of the Ti0 sample with moderate oxygen content (Ti0-G1, Fig. 1). In this case, the highest SF is reached for the π1a (−110−1) system:

2. Materials and experimental procedure Two batches of titanium with different purities are investigated. The chemical compositions are given in Table 1. The first batch – denoted T60 – has a high impurity content and an "oxygen equivalent" Oeq=1.020 atomic (at.) % with 3200 ppm of oxygen in weight (wppm). The material, issued from a cold-rolled plate exhibits equiaxed grains with a mean size of 45 µm in the rolling-transverse plane (RD-TD) and a mean size of 25 µm in the normal direction (ND). The c axes are in the TD-ND plane and < 10-10 > directions are mostly parallel to RD [2,3]. The second batch – denoted Ti0 –has a low impurity content and an Oeq=0.138 at. % with 450 wppm of oxygen. The material, issued from a forged billet has a mean grain size of 220 µm and the c axes are normal to the billet axis [22]. Micro dog-bone tensile specimens, 5.5 mm long, and 2.3 mm wide, are machined out along the rolling direction of the T60 plate and along the axis of the Ti0 billet, and mechanically ground to a thickness of 100 µm. Final thinning of the central section to electron transparency is achieved using a Tenupol-5 twin jet polisher and a 7% perchloric acid, 27% butanol, 66% methanol solution, at a voltage of 15 V and a temperature of 243 K. In situ deformation experiments are performed at room temperature in a JEOL 2000EX transmission electron microscope operating at 200 kV. The samples are mounted on a displacementcontrolled micro tensile stage allowing ± 60° tilt and are stretched until dislocation motion is observed. Then the stage displacement is blocked and the system allowed to relax, while dislocations motion is recorded by a CDD camera able to capture 25 images per second. More than six hours of usable video sequences were recorded during the straining of four samples (two samples for each material). The results presented in this paper refer to four representative grains (two grains in T60 [T60G1 and T60-G2] and two grains in Ti0 [Ti0-G1 and Ti0-G2]), whose stereographic projections, obtained using Stereo-Proj [23] are presented in Fig. 1. Schmid factors (SFs) are computed for each system considering that the tensile axis is nearly parallel (within ± 10° [24]) to the TEM tilt axis. The active slip systems are determined by stereographic analysis of the slip traces left by dislocations on the foil surfaces [24]. In the following, a specific notation will be used to distinguish the two possible first order pyramidal planes for each < a > dislocation: π1a will be used for the system corresponding to the highest SF, π1b for the weakest. Otherwise, P and B will denote prismatic and basal slip systems, respectively.

SFπ1a (−110−1)=0.49 SF P (−1100)=0.44 SF π1b (−1101)=0.29 SF B (0001)=0.21 Supplementary material related to this article can be found online at http://dx.doi.org/10.1016/j.msea.2017.07.040. The dislocations are close to the screw orientation, and unambiguously lie in the (−1100) prismatic plane, as shown by the analysis of the slip traces left at the foil surfaces (Fig. 2a). In this grain, no deviation in cross slip planes (π1 especially) has been observed during all the experiment. Under stress, the dislocations show a jerky motion (video 1-Ti0-G1) already observed in the case of high purity titanium [12]. They execute a series of rapid jumps with duration less than 40 ms over variable distances (Fig. 2b and e), and then suddenly come at rest (Fig. 2a, c and f). In the oxygen-rich titanium (T60), the frame-by-frame analysis of the videos highlights the jerky motion of < a > dislocations. Fig. 3 illustrates the plastic activity in grain T60-G1. Two systems are active in this grain. The first one denoted S1 corresponds to the motion of a/3[-1210] defects in parallel (10−10) prismatic bands. In the observation conditions of Fig. 3, the projected Burgers vector of the dislocations is nearly parallel to the intersection of the slip plane and the foil surface. The motion sequence of a dislocation in the prismatic plane is displayed in Fig. 3b-g. Initially in the field of view – as pointed out by the white arrow – (Fig. 3), the dislocation first extends in the prismatic plane by the motion of the (mobile) non-screw parts (Fig. 3b and c). As the screw segment is straight and sometimes pinned (Fig. 3g), the motion is jerky, but quite different from what was observed in Ti0, since gliding is achieved segment-by-segment over short distances between pinning points. In the same grain, a second slip system (S2) is activated. It corresponds to a/3[11-20] dislocations. The SFs nearly have the same values as those for grain G1 in Ti0, for equivalent slip planes. The Schmid factors are: SFπ1a (1−101) = 0.50 SF P (1−100) = 0.45 SFπ1b (−1101) = 0.29 SF B (0001) = 0.22 332

Materials Science & Engineering A 703 (2017) 331–339

B. Barkia et al.

T

T

(a) grain T60-G1

(b) grain T60-G2

T

T

(c) grain Ti0-G1

(d) grain Ti0-G2

Fig. 1. Stereographic projections of the analyzed grains. T is the tensile axis. (a), (b) T60: high oxygen content titanium. (c), (d) Ti0: low oxygen content titanium.

straight and aligned along the screw orientation, suggesting an intense lattice friction in the glide plane (video 2-T60-G1, S2 system). Motion is always jerky, and achieved by short displacements along the dislocations, due to pinning by extrinsic obstacles. Consequently some debris are made visible by the creation of jogs and dipoles. Dislocation generation by double cross slip also occurs, as visible in Fig. 4. Two screw dislocations denoted d1 and d2 move along a plane near to the π1a (1101) plane, and undergo pinning (Fig. 4a-b, for d1 in G) and (Fig. 4b for d2 in J). The segments near the jogs bow out of d1 and create an alphashaped dislocation (Fig. 4c-g). With further straining, the initial dislocation moves forward, and the new loop expands, generating two new parallel screw dislocations (Fig. 4g), one (d1”) moving to the right and the other (d1’) to the left (video 2-T60-G1, S2 system). In the case where the jog is not sufficiently high to allow loop expansion, dipoles are rather dragged. Supplementary material related to this article can be found online at http://dx.doi.org/10.1016/j.msea.2017.07.040.

Table 2 Critical resolved shear stress (CRSS) normalized to prismatic slip in T60 (from [3]). Slip systems

Prismatic < a >

Basal < a >

π1 < a >

π1 < c+a >

1

1.48 ± 0.10

1.18 ± 0.04

1.82 ± 0.15

In Fig. 3a, the dislocations glide from the top of the field of view towards the left side. When they initially arrive in the field of view (top right of the field), the dislocations lie in prismatic (1−100) planes. However the analysis of the slip at surfaces reveals curved traces typical of a change in glide plane from P toward π1a, also suggested by the increase in their apparent thickness. The transitions, i.e. the multiple cross slip events, are smooth and continuous, and the dislocations actually glide (left middle) in an "undefined plane" between the two maximum SF slip planes P and π1a. The motion of defects in this configuration is highlighted in Fig. 4 by successive frames. Dislocations are

333

Materials Science & Engineering A 703 (2017) 331–339

B. Barkia et al.

Fig. 2. Ti0-G1. Locking-unlocking mechanism of b=a/3[11-20] dislocations in (−1100) prismatic plane. The dislocation A is blocked in (a) and moves rapidly in (b) before locking in (c). The arrow in (b) indicates the initial position of A before motion. The insert (i) detailed the movement of dislocation B initially in (d); (e) and (f): successive frames considering the glide of C. A remanent contrast appears in (e) due to the fast motion of C. Note that micrographs (i), (ii) and (iii) are differences images. They show the flying distances of dislocations A, B and C during the process. The flying time has not been accurately measured, as it was less than one frame (40 ms). T is the tensile direction.

Fig. 6 shows an operating intra-granular localized source anchored at point S (Fig. 6a). The dislocations have a b=a/3[2-1-10] Burgers vector. The basal system has the highest SF equal to 0.48, whereas the prismatic one is less favoured:

Deviations of the slip plane are also observed in other grains for the T60 sample. Fig. 5 displays an example from grain G2 in which SF(π1a) is higher than SF(P). In this grain, the dislocations have b=a/3[-2110] Burgers vector, and the SFs are: SF SF SF SF

π1a (01−1−1) = 0.49 P (01−10) = 0.45 π1b (01−11) = 0.29 B (0001) = 0.21

SF SF SF SF

The situation is close to the previous one, with a pyramidal system better oriented than the prismatic one. In Fig. 5a dislocations move from the top to the bottom of the micrograph, along parallel prismatic planes. They are not particularly straight in Fig. 5a but more rectilinear in Fig. 5b. The activation of the π1a (01-1-1) system seems effective, considering the progressive evolution of the slip traces from P to π1a, from the left to the right of Fig. 5c. The effect is more flagrant if we consider Fig. 5b and d. In the latter, defects glide from the grain boundary to the left. Even though the traces are not totally straight on the foil surfaces, some of them match with pyramidal traces. Moreover the presence of slip traces between pyramidal π1a and B strongly suggests the activation of the basal system, despite its unfavourable Schmid factor.

B (0001) = 0.48 π1a (01−11) = 0.32 π1b (01−1−1) = 0.13 P (01−10) = 0.11

The mechanism of dislocation generation by the source is conventional. Under stress and repeatedly, a dislocation segment rotates around the pinning point S (video 3-Ti0-G2), and emits two screw segments from each side of the source (Fig. 6b-g). However, the observation of surface traces clearly indicates multiple cross slip. In the first stages of source activity, dislocations are emitted in a plane situated between P and B and close to π1a (Fig. 6a″ and c′). Such a slip is clearly not stable, as traces re-orientate between P and π1a far from the anchor point. The emitted dislocations are not particularly straight (Fig. 6f), and seem locked during short intervals, before gliding over variable (but quite long) distances (Fig. 6f′ and g′). Supplementary material related to this article can be found online at http://dx.doi.org/10.1016/j.msea.2017.07.040. A similar pattern is observed in different zones of the strained grain. Fig. 7 illustrates the wavy nature of slip, attesting to a frequent and intensive cross slip. Apart from rare cases (Fig. 8), the slip traces correspond to both P and π1a systems, in spite of their relatively low SF. The activation of the basal system is also indirectly emphasized (Fig. 8). A screw dislocation, denoted A, with b=a/3[2-1-10] Burgers vector moves in an intermediate P+π1a plane (Fig. 8a) from the left to the right, before it cross slips to a plane whose trace lies between π1a and B (Fig. 8b and c). After some motion, the dislocation cross slips back, and

3.2. Wavy slip in titanium with moderate oxygen content In the cases reported above, most of grains in Ti0 or T60 were well oriented to favour P, π1 and even composite P+π1 glide (Fig. 5). Slip traces were generally rectilinear on surfaces, even when the slip plane was located between P and π1. The situation is totally different in some grains for which the SF is the highest for the basal plane. A typical example is given for a Ti0 sample and illustrated in Figs. 6 and 7. 334

Materials Science & Engineering A 703 (2017) 331–339

B. Barkia et al.

Fig. 3. T60-G1. Dislocation activity in oxygen-rich titanium. Two < a > systems are active: S1 with b1=a/3[-12-10] dislocations in P1 (10-10) prismatic plane and S2 with b2=a/3[1120] dislocations initially gliding from top right in P2 (1−100) plane. (a)-(f) glide in P1 planes. Dislocation A is initially shown by an arrow in (a). Micrographs (b) to (g) display the motion of A in the P1 (10−10) prismatic plane. Arrows in (b) and (e) indicate the sense of motion for the edge part of A; arrows in (f) highlight some pinning points on the A dislocation. The difference between (f) and (g) is displayed in (i). T is the tensile direction.

its trace corresponds to an intermediate plane between P and π1a (Fig. 8e) (video 4-Ti0-G2). It is worth noticing that pure basal slip is never observed, whatever the impurity content or grain orientation. Supplementary material related to this article can be found online at http://dx.doi.org/10.1016/j.msea.2017.07.040.

barrier in the prismatic planes makes glide of < a > dislocations easy, as confirmed by atomistic and experimental studies on single crystals or polycrystals [24,25]. The present observations are indeed consistent with these studies, especially in Ti0 samples, where many grains are well oriented for prismatic slip. In such a case, and in accordance with previous experiments or recent simulations [17], the present observations confirm planar slip, and the typical locking-unlocking mechanism [16], i.e. a succession of transitions between glissile and sessile core configurations. In glissile mode, the curvature of the dislocations highlights the modest friction stress along prismatic planes. The range of flying distances is large: from a few nanometres to more than 600 nm, as illustrated by Fig. 2iii. The situation is however more complex for higher oxygen content. Contrary to what occurs in micro-pillars [15], plastic deformation is never homogeneous in the oxygen-rich titanium batch. Dislocations are always confined within bands, in which many debris are present (Fig. 9). The homogeneous deformation in pillars could result from the high stresses attained, sufficient to activate slip of another type of dislocations such as b= ± [10-12] [15]. As for dislocations motion, the present study supports an oxygeninduced reduction in dislocations mobility. Probably pinned by extrinsic obstacles such as impurities, dislocations jump over shorter distances along the line between two stable (sessile) positions. As motion occurs segment by segment, the dislocations velocity is significantly reduced, and this results in a pronounced hardening effect. The presence of numerous debris (dipoles, loops) formed in the wake of

4. Discussion The above observations of deformation processes for two different α-Ti batches give the following main significant results: (1) plastic deformation is carried by < a > dislocations; (2) prismatic slip is always activated whatever the grain orientation or the impurity content; (3) when one of the two first-order pyramidal systems has a higher SF than the prismatic ones, dynamic observations suggest that slip of < a > dislocations takes place differently, depending on the oxygen content. In Ti0, dislocations move easily over large distances between pinning events, whereas they are quickly pinned after short distances in T60. (4) wavy slip due to profuse cross slip is observed when the maximum SF on the basal plane is the highest among all SFs. The first two results are expected, and will not be extensively discussed. Except when the tensile axis is parallel to the < c > axis, < a > dislocations generally accommodate deformation. The low Peierls 335

Materials Science & Engineering A 703 (2017) 331–339

B. Barkia et al.

Fig. 4. T60-G1. (a) Dislocations d1 and d2 of the S2 system (b2=a/3[11-20]). Traces at surfaces correspond to a glide close to π1a (1−101). d1 is initially pinned in G (full arrow). (b) formation of a jog in J after pinning; (b1) and (b2) a dipole loop is formed by punching off. Same mechanism is observed in (d) and (e) for d1 (full arrows). (c)-(g) mechanism of loop expansion from a jog at G for d1. Two dislocations d1’ and d1’’are thus formed. Dashed arrows represent the motion sense of dislocations. T is the tensile direction. Fig. 5. T60-G2 - motion of a/3[-2110] dislocations. (a) Near P (01−10) slip in the left corner. The deviation of the slip traces in the bottom right corner suggests the activation of π1a (01−1−1) system. (b) evidence of pyramidal glide of < a > dislocations close to a grain boundary. The black arrow indicates the presence of some slip traces π1a and B. (c)-(d) rotation of images (a) and (b), respectively, to highlight the curvature of the slip traces. Note the numerous debris in the wake of dislocations. T is the tensile direction.

1000 wppm O [15]. The core modification results in an increase of the recombination energy of screw dislocations with the increase of impurities content [11,26–28]. On the other hand, the most striking result for T60 is the relative instability of prismatic glide, especially when both prismatic and pyramidal SFs are high. Instability refers here to the numerous (smooth) slip trace deviations from pure prismatic to pyramidal orientations, i.e. a repeated cross slip phenomenon.

gliding dislocations is also emphasized (Fig. 9). This specific point will be discussed further. The effect of solute oxygen on defects motion had already been discussed in past pioneering studies. Oxygen atoms were assumed to induce a modification of dislocations core, thus hindering the sessileglissile transition [11]. This hypothesis has been confirmed recently by high-resolution HAADF-STEM projected images of screw dislocations whose cores seem narrower in Ti + 3000 wppm O than in Ti + 336

Materials Science & Engineering A 703 (2017) 331–339

B. Barkia et al.

Fig. 6. Ti0-G2 - details of a dislocation emission by a source initially in S (a). Dislocations have a a/3[2-1-10] Burgers vector. (a′) and (a″): first emission of a dislocation from S in a plane close to π1a and which escapes rapidly to the left in a P+ π1a glide as pointed out by the arrow in (b). A second emission is observed in (c), (c′) and (d). Dislocations motion is difficult between (e) and (f) before rapid displacement always in a composite P+ π1a slip. Differential micrographs in (f′) and (g′) show the jump length between (e) and (f) / (f) and (g). T is the tensile direction.

activation of pyramidal slip was reported above 150 K in pure titanium, in which it becomes more frequent with increasing temperature [17,18]. DFT calculations predict that dislocations would cross slip from primary P to π1, in which plane defects are straight, and move with difficulty due to the intense friction stress in these planes. At higher temperatures (300–570 K), the presence of edge dipoles and wavy slip at surfaces were interpreted as profuse cross slip. The way interstitial oxygen atoms influence the activation of π1 slip, is not so clear and several hypotheses have been put forward. A first possibility is that the critical resolved shear stress (CRSS) for π1 slip varies with the oxygen content. The values of normalized CRSS for π1 < a > /P < a > in polycrystalline titanium, measured during in situ tensile tests in a SEM, lead to ratios CRSS(π1 < a >)/CRSS (P < a >) close to 1.24 for T40 (1600 wppm O) and 1.18 for T60

The comparison between Ti0 and T60 shows significant differences from that respect, especially if we consider grains T60-G1/T60-G2 and Ti0-G1, which have similar orientations relative to the applied stress. In these conditions, the activation of pyramidal slip was unambiguously highlighted in T60 (Figs. 3, 4 and 5), but never in Ti0. Regardless of the impurity content, the π1 glide in α-Ti is often mentioned in the literature as a secondary slip system. At mesoscopic scale, it has been extensively reported by optical or scanning electron microscopy (SEM) analysis after mechanical testing. Wavy slip lines at the surface of Ti single crystals strained above 300 K have been attributed to the cross slip of < a > dislocations from P toward π1 [11], and deviations from P toward π1 occur with thin and sometimes long slip lines along pyramidal planes [22]. At microscopic scale, TEM studies reported this type of slip more or less clearly [10,11,17,29]. The

Fig. 7. Ti0-G2 - glide of a/3[2-1-10] dislocations. (a) g=10–11. Note the high number of debris. Dislocations are often pinned at jogs (arrow). The apparent double contrast for the dislocation A highlights the motion of the defect during the capture. (b) g=0002. Dislocations A are invisible. Surface traces are wavy and correspond to an intermediate slip between prismatic P and pyramidal π1a systems. T is the tensile direction.

337

Materials Science & Engineering A 703 (2017) 331–339

B. Barkia et al.

Fig. 8. Ti0-G2. Micrographs taken from the in situ tensile test of the Ti0 alloy. (a) gliding A dislocation with b=a/3[2-1-10] (arrow) between prismatic P (0−110) and the pyramidal π1a (01−11) planes. (b) and (c) cross-slip of the dislocation towards a plane between (01−11) and (0001). (d) Motion of the dislocation in the cross-slip plane evidenced by the difference between micrographs extracted from two successive frames: Δt = 40 ms between initial position (black contrast) and those after motion (white contrast). (e) Cross-slip from B+ π1a to π1a+P planes. T is the tensile direction.

stacking-fault associated to the < a > dislocation [21]. Oxygen seems to impose cross slip of the < a > dislocation into π1 plane, inducing the formation of jog-type defects along the line. In the particular case of titanium, for which the stable core is spread in pyramidal planes, oxygen locks < a > dislocations in these planes [20]. Consequently oxygen reduces the dislocations mobility in the prismatic planes, and favours cross slip in first order pyramidal ones. These simulations are supported by our observations for T60, in which pure prismatic glide was hardly observed. In the grains with maximum SF on a π1 < a > system, mixed traces correspond to a composite glide (multiple cross slip) due to oxygen atoms in the slip plane and/or in < a > dislocation cores [15,20]. The presence of debris in the wake of gliding dislocations reveals the cross slip phenomenon. In oxygen-rich titanium, the presence of extensive cross slip is indirectly proved by the creation of numerous dipoles and loops. Furthermore, the density of debris seems to increase with the oxygen content, as suggested by a comparison of strained Ti0 and T60 with similar SFs in P and π1a. According to recent studies on iron alloys [34], the formation of such debris is related to a cross slip mechanism, which supposes the interaction of moving kinks in different slip planes along screw dislocations. Cross-kinks may lead to the formation of loops [34,35] and such mechanism is effective in metals or alloys whose dislocations cores are non planar, as it is the case in α-Ti [20,32]. It is enhanced when small-size obstacles – such as solute impurities – are present [36]. In such a case, pinning of dislocations by impurities plays a key role in the formation of jogs or super-jogs. The rate of debris production with plastic straining is particularly high in T60, and many cross slip loops are continuously created (video 5-T60-G2). This might significantly contribute to the dislocations multiplication, but also to hardening [19]. Supplementary material related to this article can be found online at http://dx.doi.org/10.1016/j.msea.2017.07.040. However the situation is more complex when the CRSS on the basal system is high, as it is observed in Ti0 (Ti0-G2). As a matter of fact, the wavy glide observed in Ti0 has to be discussed in light of the active slip systems (point 4). Such phenomena have already been observed in α-Ti, but not really interpreted [10,16]. Analyses of the surfaces slip traces highlight extensive dislocations cross slip from prismatic to/from first order pyramidal and basal planes, but also from pyramidal to/from

T -1120

500nm

Fig. 9. T60-G2 - presence of debris in the wake of a/3[-1-120] dislocations during an in situ experiment. T is the tensile direction.

(3200 wppm O) [3]. The difference in CRSS seems to be slightly reduced with increasing oxygen content. At lower oxygen concentrations it is likely that this CRSS(π1 < a >)/CRSS (P < a >) ratio increases strongly as no first order π1 traces were observed for α-Ti with 100 ppm of oxygen [30]. The present study is in line with this tendency: for grains with similar SFs for a specific < a > system – (T60-G1 and G2 [O=3200 wppm] and Ti0-G1 [O = 450 wppm]: SF (π1a)≈0.50 > SF (P)=0.45) activation of first order pyramidal slip is only observed for T60, but never in Ti0. The evolution of the CRSS of π1 < a > with the oxygen content might be less pronounced than for the P < a > system, as it is the case for B < a > [31]. Consequently the ratio CRSS (π1 < a >)/CRSS(P < a >) could be larger for a lower oxygen content. In the HAADF atomic resolution image from Yu's study [15], the oxygen columns close to the screw dislocation core are aligned along the prismatic plane, in favour of an increase in the CRSS for this plane. A second possible way for oxygen to influence π1 slip is through interactions with dislocations cores. These interactions have been recently simulated using ab initio methods for Zr and Ti [15,17,20,21,32,33]. In both cases, strong repulsion is predicted between oxygen atoms and dislocations when interstitial atoms lie exactly within the dislocation core. In zirconium, the repulsion is explained by the destruction of octahedral sites containing oxygen atoms by the 338

Materials Science & Engineering A 703 (2017) 331–339

B. Barkia et al.

[6] P. Castany, E. Pettinari-Sturmel, J. Douin, A. Coujou, In situ transmission electron microscopy deformation of the titanium alloy Ti-6Al-4V: interface behaviour, Mater. Sci. Eng. A: Struct. Mater. Prop. Microstruct. Process. 483–484 (2008) 719–722, http://dx. doi.org/10.1016/j.msea.2006.10.183. [7] P. Castany, M. Besse, T. Gloriant, In situ TEM study of dislocation slip in a metastable β titanium alloy, Scr. Mater. 66 (2012) 371–373, http://dx.doi.org/10.1016/j.scriptamat. 2011.11.036. [8] M. Legros, D. Caillard, A. Couret, An in situ study at room temperature of deformation processes in a Ti23.7Al-9.4Nb alloy, Intermetallics 4 (1996) 387–401, http://dx.doi.org/ 10.1016/0966-9795(95)00055-0. [9] Q. Yu, J. Sun, J.W. Morris, A.M. Minor, Source mechanism of non-basal < c plus a > slip in Ti alloy, Scr. Mater. 69 (2013) 57–60, http://dx.doi.org/10.1016/j.scriptamat.2013. 03.009. [10] S. Naka, A. Lasalmonie, Cross-slip on the 1st order pyramidal plane (1011) of a-type dislocations [1210] in the plastic-deformation of alpha-titanium single-crystals, J. Mater. Sci. 18 (1983) 2613–2617, http://dx.doi.org/10.1007/BF00547577. [11] S. Naka, A. Lasalmonie, P. Costa, L. Kubin, The low-temperature plastic-deformation of alpha-titanium and the core structure of a-type screw dislocations, Philos. Mag. A: Phys. Condens. Matter Struct. Defects Mech. Prop. 57 (1988) 717–740, http://dx.doi.org/10. 1080/01418618808209916. [12] S. Farenc, D. Caillard, A. Couret, An in-situ study of prismatic glide in alpha-titanium at low-temperatures, Acta Metall. Mater. 41 (1993) 2701–2709, http://dx.doi.org/10.1016/ 0956-7151(93)90139-J. [13] S. Farenc, D. Caillard, A. Couret, A new model for the peak of activation area of alphatitanium, Acta Metall. Mater. 43 (1995) 3669–3678, http://dx.doi.org/10.1016/09567151(95)90150-7. [14] J. Kacher, I.M. Robertson, In situ TEM characterisation of dislocation interactions in alpha-titanium, Philos. Mag. 96 (2016) 1437–1447, http://dx.doi.org/10.1080/ 14786435.2016.1170222. [15] Q. Yu, L. Qi, T. Tsuru, R. Traylor, D. Rugg, J.W. Morris, M. Asta, D.C. Chrzan, A.M. Minor, Origin of dramatic oxygen solute strengthening effect in titanium, Science 347 (2015) 635–639, http://dx.doi.org/10.1126/science.1260485. [16] A. Couret, D. Caillard, Prismatic slip in beryllium. 1. The controlling mechanism at the peak temperature, Philos. Mag. A: Phys. Condens. Matter Struct. Defects Mech. Prop. 59 (1989) 783–800, http://dx.doi.org/10.1080/01418618908209820. [17] E. Clouet, D. Caillard, N. Chaari, F. Onimus, D. Rodney, Dislocation locking versus easy glide in titanium and zirconium, Nat. Mater. 14 (2015) 931–936, http://dx.doi.org/10. 1038/NMAT4340. [18] S. Naka, L. Kubin, C. Perrier, The plasticity of titanium at low and medium temperatures, Philos. Mag. A: Phys. Condens. Matter Struct. Defects Mech. Prop. 63 (1991) 1035–1043, http://dx.doi.org/10.1080/01418619108213935. [19] A. De Crecy, A. Bourret, S. Naka, A. Lasalmonie, High-resolution determination of the core structure of 1/3 (11-20)(10-10) edge dislocation in titanium, Philos. Mag. A: Phys. Condens. Matter Struct. Defects Mech. Prop. 47 (1983) 245–254, http://dx.doi.org/10. 1080/01418618308245221. [20] N. Chaari, Modélisation ab initio de la plasticité dans les métaux hexagonaux: zirconium et titane purs et effets de l′oxygène, Université Grenoble Alpes, 2015, 〈https://tel. archives-ouvertes.fr/tel-01269636〉. [21] N. Chaari, D. Rodney, E. Clouet, Oxygen - Dislocation interaction in zirconium from first principles, Acta Mater. 132 (2017) 416–424, http://dx.doi.org/10.1016/j.actamat.2017. 05.008. [22] B. Barkia, Viscoplasticité à l′ambiante du titane en relation avec ses teneurs en oxygène et en hydrogène, Éc. Polytech. (2014), 〈https://tel.archives-ouvertes.fr/tel-01137807〉. [23] F. Mompiou, Stereo-proj. 〈http://mompiou.free.fr/pages/stereo-fr.html〉. [24] D. Caillard, J.L. Martin, Thermally activated mechanisms in crystal plasticity, Permagon (2003). [25] N. Tarrat, M. Benoit, D. Caillard, L. Ventelon, N. Combe, J. Morillo, Screw dislocation in hcp Ti: dft dislocation excess energies and metastable core structures, Model. Simul. Mater. Sci. Eng. 22 (2014) 55016, http://dx.doi.org/10.1088/0965-0393/22/5/055016. [26] E. Levine, Deformation mechanisms in titanium at low temperatures, Trans. Metall. Soc. AIME 236 (1966) 1558–1564. [27] A. Akhtar, E. Teghtsoonian, Prismatic slip in alpha-titanium single-crystals, Metall. Trans. A: Phys. Metall. Mater. Sci. 6 (1975) 2201–2208, http://dx.doi.org/10.1007/ BF02818644. [28] M. Biget, G. Saada, Low-temperature plasticity of high-purity alpha-titanium singlecrystals, Philos. Mag. A: Phys. Condens. Matter Struct. Defects Mech. Prop. 59 (1989) 747–757, http://dx.doi.org/10.1080/01418618908209818. [29] D. Shechtman, D. Brandon, Orientation dependent slip in polycrystalline titanium, J. Mater. Sci. 8 (1973) 1233–1237, http://dx.doi.org/10.1007/BF00549337. [30] A. Churchman, The slip modes of titanium and the effect of purity on their occurrence during tensile deformation of single crystals, Proc. R. Soc. Lond. Ser. A: Math. Phys. Sci. 226 (1954) 216–226, http://dx.doi.org/10.1098/rspa.1954.0250. [31] H. Conrad, Rate controlling mechanism during yielding and flow of alpha-titanium at temperatures below 0.4 Tm, Acta Metall. 14 (1966) 1631, http://dx.doi.org/10.1016/ 0001-6160(66)90186-6. [32] D. Rodney, L. Ventelon, E. Clouet, L. Pizzagalli, F. Willaime, Ab initio modeling of dislocation core properties in metals and semiconductors, Acta Mater. 124 (2017) 633–659, http://dx.doi.org/10.1016/j.actamat.2016.09.049. [33] L. Liang, Ab Initio Simulation of Extended Defects of Alpha-Ti in Presence of Interstitial Atoms H & O, Université Paris-Saclay, 2016, 〈https://pastel.archives-ouvertes.fr/tel01355132〉. [34] D. Caillard, A. TEM, in situ study of alloying effects in iron. II—Solid solution hardening caused by high concentrations of Si and Cr, Acta Mater. 61 (2013) 2808–2827, http://dx. doi.org/10.1016/j.actamat.2013.01.049. [35] E. Furubayashi, Behavior of dislocations in Fe-3 percent Si under stress, J. Phys. Soc. Jpn. 27 (1969) 130–146, http://dx.doi.org/10.1143/JPSJ.27.130. [36] D. Caillard, M. Legros, A. Couret, Extrinsic obstacles and loop formation in deformed metals and alloys, Philos. Mag. 93 (2013) 203–221, http://dx.doi.org/10.1080/ 14786435.2012.705912.

basal planes in the material containing 450 wppm oxygen. These observations are correlated to the high SF on the basal plane, rather than to the oxygen content. Previous studies all report that basal slip is difficult at room temperature [26,27] in line with the dissociation instability for < a > dislocations cores in this plane [25]. Despite the low SFs on prismatic and first order pyramidal systems, the stress was high enough to activate it. In this specific case, extensive cross slip led to the formation of loops or dipoles, but also to a more homogenous dislocations structure. 5. Conclusions The effects of oxygen on the dynamic behaviour of dislocations in polycrystalline α-Ti were investigated by means of in situ tensile tests in a TEM. – Regardless of the impurity content, < a > screw dislocations control plastic deformation in α-Ti at room temperature. Prismatic slip is always observed. – Dislocations glide is characterized by a jerky movement. The motion is clearly affected by the oxygen content. Solute oxygen increases the pinning effect and reduces the jumping distance during the motion. In oxygen-rich titanium only – and depending on the grain orientation – intermediate slip traces of individual dislocations are found between prismatic and first order pyramidal planes, due to multiple cross slip. It is believed that oxygen strongly influences the activation of pyramidal slip, reducing the CRSS(π1 < a >)/CRSS (P < a >) ratio at high oxygen content. The continuous cross slip phenomenon and the presence of debris can also be related with recent atomistic simulations highlighting strong repulsion between oxygen atoms in octahedral sites and screw dislocation core, forcing local deviation of the screw defects. – When the grains are better oriented for basal slip – a condition encountered in the Ti0 sample with moderate oxygen content – screw dislocations leave wavy slip traces, sign of frequent changes in slip plane (from prismatic to pyramidal and basal and back, but also from pyramidal to/from basal). Such a phenomenon has not been observed in the oxygen-rich material that has a different texture, and a lack of grains suitably oriented for basal slip. However nothing suggests that a high oxygen content would change the mechanisms if the appropriate orientation conditions were met. Acknowledgements The authors would like to gratefully acknowledge Daniel Caillard (CEMES, CNRS) and Emmanuel Clouet (SRMP, CEA) for their interest on the subject and for the rich discussions and exchanges the last couple of months. This study was carried out in the framework of project FLUTI (ANR10-BLAN-915) funded by the Agence Nationale de la Recherche (ANR). References [1] H. Conrad, Effect of interstitial solutes on the strength and ductility of titanium, Progress. Mater. Sci. 26 (1981) 123–404, http://dx.doi.org/10.1016/0079-6425(81)90001-3. [2] B. Barkia, V. Doquet, J.P. Couzinie, I. Guillot, Room-temperature creep and stress relaxation in commercial purity titanium-Influence of the oxygen and hydrogen contents on incubation phenomena and aging-induced rejuvenation of the creep potential, Mater. Sci. Eng. A: Struct. Mater. Prop. Microstruct. Process. 624 (2015) 79–89, http://dx.doi.org/ 10.1016/j.msea.2014.11.073. [3] B. Barkia, V. Doquet, J.P. Couzinie, I. Guillot, E. Heripre, In situ monitoring of the deformation mechanisms in titanium with different oxygen contents, Mater. Sci. Eng. A: Struct. Mater. Prop. Microstruct. Process. 636 (2015) 91–102, http://dx.doi.org/10. 1016/j.msea.2015.03.044. [4] J. Williams, P. Tung, A. Sommer, Influence of oxygen concentration on internal stress and dislocation arrangements in alpha titanium, Metall. Trans. 3 (1972) 2979–2984, http:// dx.doi.org/10.1007/BF02652870. [5] P. Castany, F. Pettinari-Sturmel, J. Crestou, J. Douin, A. Coujou, Experimental study of dislocation mobility in a Ti–6Al–4V alloy, Acta Mater. 55 (2007) 6284–6291, http://dx. doi.org/10.1016/j.actamat.2007.07.032.

339