Journal of Materials Science & Technology, 2020, 52(0): 235-242 DOI: 10.1016/j.jmst.2020.02.064

Research Article

The influence of solute atom ordering on the deformation behavior of hexagonal close packed Ti-Al alloys

Hao Wu,a,*, Yunlei Xua, Zhihao Wangb, Zhenhua Liuc, Qinggang Lid, Jinkai Li,a,*, Junyan Wu,a,*

School of Materials Science and Engineering, University of Jinan, Jinan 250022, China

School of Material Science and Engineering, Qilu University of Technology (Shandong Academy of Sciences), Jinan 250353, China

Technical Division, Sinomach Foundry and Metal Forming Co., Ltd., Jinan 250306, China

Shandong Provincial Key Laboratory of Preparation and Measurement of Building Materials, Jinan 250022, China

Corresponding authors: *.E-mail addresses:msewuh@ujn.edu.cn(H. Wu),mselijk@ujn.edu.cn(J. Li),msewujunyan@ujn.edu.cn(J. Wu).

Received: 2020-02-2   Accepted: 2020-02-23   Online: 2020-09-15

Abstract

A framework of compositionally graded Ti-Al alloys was proposed to elucidate the Al alloying dependent deformation mechanism at room temperature. Slip trace analysis demonstrated low-alloyed Ti-Al model materials were plastically deformed by prismatic slip, pyramidal slip, stacking faults, and deformation twins. Increasing Al concentration promoted chemical ordering, thus suppressing the dislocation motion and deformation twinning. The mechanism behind such a kind of composition-mediated deformation physics was discussed. Our findings are expected to be applicable in the design of next-generation high-performance structural materials through the tailoring of the chemical ordering of solute atoms in substitutional solid solution.

Keywords: Titanium alloys ; Diffusion ; Alloying ; Electron microscopy

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Cite this article

Hao Wu, Yunlei Xu, Zhihao Wang, Zhenhua Liu, Qinggang Li, Jinkai Li, Junyan Wu. The influence of solute atom ordering on the deformation behavior of hexagonal close packed Ti-Al alloys. Journal of Materials Science & Technology[J], 2020, 52(0): 235-242 DOI:10.1016/j.jmst.2020.02.064

1. Introduction

Hexagonal close packed (hcp) metals such as titanium are one of major workhorse structural materials for engineering applications [1]. Strengthening titanium via alloying has been extensively investigated in the past few decades [2,3]. Understanding the impact of alloying elements on the mechanism and mechanics of dislocation slip and deformation twinning is of great importance to guide the design and manufacturing of next-generation high-performance metallic alloys [4,5]. As a commonly used α-phase stabilizer, the addition of Al element disturbs the lattice stability and c/a ratio of hcp matrix, thus influencing the deformation mechanism and consequently mechanical properties [[6], [7], [8]]. Our recent work demonstrated (i) that Al alloying changed the structure of dislocation core through solute-dislocation interactions [9], and (ii) that preferred substitution of Al atoms onto {10-10} prismatic planes significantly weakened deformation anisotropy of hcp metals, making basal and prismatic slip contributing equally when the Al content is increased up to 6.6 wt% [10] and stimulating the oriented precipitation of ordered Ti3Al from supersaturated Ti(Al) solid solution [11]. In addition to dislocation slip, deformation twinning is another prevalent plasticity mechanism for hcp metals due to insufficient slip systems required for uniform deformation according to von Mises or Taylor criterion [[12], [13], [14]]. Except for temperature, strain rate, and loading condition [[15], [16], [17], [18]], the chemical constituent is also a key factor in determining the twin nucleation, kinetics, and variant selection [19,20]; however it is not immediately clear whether the alloying element facilitates twinning or not. On the one hand, deformation twinning is usually occurred through the motion of twinning dislocations, and its proceeding seems easy to be trapped by solute atoms such as Al [21]. On the other hand, the Al alloying decreases the stacking fault energy [22], in combination with elevated flow stress against dislocation sliding [23], both assisting deformation twinning [24]. These two roles compete with each other in a way that the former elevates the critical stress required for twin nucleation, while the latter creates favorable conditions for deformation twinning [25,26]. Additionally, the twinning behavior at extremely high Al concentration (e.g., >12 at.%, ordered Ti3Al was precipitated) remains unclear. Therefore the aim of the present work is to elucidate such a kind of competitive twinning mechanism and to investigate the influence of Al addition on deformation twinning and dislocation dynamics upon deformation at room temperature.

2. Experimental

Commercially pure titanium and aluminum foils were cut into several pieces with dimension of 50 mm × 50 mm (length × width), chemically etched via dilute HF and NaOH solution to a thickness of 100 μm and 700 μm, respectively, ultrasonically cleaned, and dried. They were then alternately stacked, hot pressed at 515 °C under 75 MPa for a period of 1 h, and diffusively annealed initially at 700 °C for 1 h followed by high-temperature annealing at 1100 °C for 4 h under a pressure of 30 MPa. The as-fabricated model alloys, with a gradient of Al content along depth, were plastically deformed under tension so as to elaborate the influence of Al concentration. The uniaxial tension was applied by employing an Instron-1186 Universal Testing Machine at a strain rate of 1 × 10-4 s-1. The tensile specimens are flat dog-bone-shaped with gauge sizes of 18 mm length by 2 mm thickness by 5 mm width. The microstructures before and after deformation were examined by a LEO 1530 scanning electron microscope (SEM). The step size, working distance, and acceleration voltage for electron backscatter diffraction (EBSD) scanning were 0.1 μm, 10 mm, and 20 kV, respectively. Commercial HKL Channel 5 software was used for data post-processing. For post-mortem characterization of deformed microstructure, a JEM-2100 transmission electron microscope (TEM) was utilized with a working voltage of 200 kV.

3. Results

A series of as-casted or forged Ti-Al alloys were routinely used over the past few decades to elaborate the influence of chemical constituents on the microstructure, deformation mechanism, and mechanical properties; however such a strategy may cause undesired deviations in grain size, components and metallurgical defects [27,28]. For evading the dilemma, a model material with a compositional gradient was designed in the present study and fabricated by a simple diffusion reaction of alternately stacked Ti foils and Al foils. Due to the large deviation in the melting points of raw material foils, they were first annealed at 700 °C for a period of 1 h to fully convert Al to intermetallic TiAl3 compounds. Diffusion-induced Kirkendall voids were formed and subsequently eliminated by applying a pressure of 30 MPa for 4 h at 1100 °C, accompanied by the evolution of a series of intermediate phases such as Ti2Al5, TiAl2, TiAl, and Ti3Al. The chain of diffusion reaction can be described as Ti + Al → Ti + TiAl3 (after annealing at 700 °C) → Ti + Ti2Al5 + TiAl2 + TiAl + Ti3Al (annealing at 1100 °C) → Ti(Al) with a gradient composition as shown in Fig. 1 after annealing. A more detailed description of phase transformation and diffusion kinetics during sample preparation can be found in our previous work [29].

Fig. 1.

Fig. 1.   (a) Back-scattered electron (BSE) images of annealed samples at 1100 °C for 4 h and (b) corresponding compositional profile analysis along the black arrow as indicated in panel (a), showing a gradient of chemical constituents along the diffusion direction. Scale bar, 30 μm.


Uniaxial tension testing was carried out to trigger the gliding of dislocations at room temperature and the elongation to fracture was measured as about 8.5%. We then performed post-mortem TEM characterization on tensile fractured samples for elucidating the influence of Al alloying on the deformation mechanism. As shown in Fig. 2, the left grain has an Al content of 12.2 ± 0.4 at.% (averaged by five independent measurements in the TEM chamber), surpassing the solubility limit and consequently inducing the precipitation of ordered α2-Ti3Al [30]. The grain size of α2-Ti3Al precipitates is very fine, merely no more than 10 nm. Electron diffraction revealed an orientation relationship between α-Ti matrix and α2-Ti3Al precipitates of (10-10)α//(10-10)α2, [2-423]α//[1-2132, which agrees well with our in situ TEM study on precipitation kinetics [11]. Within the grain interior, the transition from Al disordering to chemical ordering was clearly observed through high-resolution TEM imaging as shown in Fig. 3(a) (from bottom to top), as well as perfect coherent atomic array free of dislocations (Figs. 2(a) and 3 (c)).

Fig. 2.

Fig. 2.   TEM-based slip trace analysis. (a) Bright field (BF) image of two neighboring deformed grains with different Al concentrations of 12.2 at.% (left grain) and 10.5 at.% (right grain). (b-d) Selected area electron diffraction (SAED) patterns of (b) grain boundary, (c) left grain, and (d) right grain.


Fig. 3.

Fig. 3.   Dependence of Al ordering on the dislocation activity. (a) High-resolution TEM micrograph nearby the grain boundary. (b-f) Corresponding reconstructed lattice fringes through inverse fast Fourier transform (FFT). The selected diffraction frequencies were indicated at the left bottom of each panel. A considerable amount of dislocations were detected in the right grain, both on prismatic and pyramidal planes, in sharp contrast with lattice fringes on the left side.


The matrices of two grains of interest were indexed as hcp Ti according to pdf card # 65-3362 (left grain) and # 51-0631 (right grain). On the right side, only 10.5 ± 0.3 at.% Al was included without the precipitation of Ti3Al; instead, significant dislocation sliding on {10-10} prismatic and {10-11} pyramidal planes were observed in Figs. 2(a) and 3 (e). Compared to dislocation dynamics in the left grain, it seems plausible that increasing Al concentration dramatically elevated the critical resolved shear stress against dislocation gliding.

Fig. 4 shows the deformed microstructure in the interior of low-Al-containing right grain. Nanotwin lamella and stacking faults were both observed and experimentally supported by fast Fourier transform (FFT) and electron diffraction of the region where the high-resolution TEM image was taken. The twin lamella was about 3 nm in thickness and the twin boundary was indexed as {10-11} planes (Fig. 4(c)). By contrast, no twins and stacking faults was detected in the left grains, and we thus concluded that severe Al alloying may impede the generation of planer defects of stacking faults and deformation twins to some extent.

Fig. 4.

Fig. 4.   Deformation nanotwinning and stacking faults. The image was taken at the right grain of Fig. 2(a) with relatively low Al content. No twin was found in the opposite side. In panel (a), we show the atomic array of parent phase and twin variants, and reconfirmed the nanotwinning by (b) electron diffraction at the region where panel (a) was taken. (c) Inverse and (d) FFT images of the region of interest (red squares in panel a), revealing the {10-11} twin with a lamella thickness of merely 3 nm as well as stacking faults. (e) FFT patterns of the red-bordered region as plotted in panel (c).


4. Discussion

4.1. Competitive twinning mechanisms mediated by Al ordering

According to the classic twinning theory, the ability of the microstructure to nucleate a deformation twin is determined by atomic shearing mechanism as well as the motion of twinning dislocations onto the twin planes [31]. In this regard, the importing of Al solute atoms is expected to elevate the slip barrier of twinning dislocation thus impeding deformation twinning. On the other hand, Al alloying decreases the stacking fault energy (SFE), which facilitates the dissociation of dislocations and deformation twinning. These two roles played by Al alloying are mutually competitive; for example, deformation twins become increasingly important in relieving stresses particularly along the crystallographic <c> direction in severely alloyed Ti-Al alloys where pyramidal slip is strictly suppressed by hexagonal symmetry and excess Al addition [15].

The influence of Al alloying on SFE levels, γ, was evaluated using an empirical formula of [22]:

$\gamma =310/\exp (0.\text{133}{{N}_{\text{Al}}})$

where NAl denotes the atomic percent of Al concentration. Substitutions of NAl = 12.2% and NAl = 10.5% into Eq. (1) yield the SFE levels of 61 and 77 mJ/m2 for left and right grains, respectively. Such a low SFE helps with the dissociation of dislocations; however the motion of dislocations is dynamically governed by the degree of Al ordering, obeying the rule of “the larger the Al concentration, the harder the dislocation motion”. For instance, the gliding of partial dislocations seems relatively easier for low-alloyed right grains, leaving behind a ribbon of stacking faults and thereby constituting the first step to nucleate a deformation twin. In other words, the formation of stacking faults and twins are principally determined by Al alloying rather than SFE, because the latter is of the same order of magnitude for both grains, being assumed to be contributed equally. At room temperature, the solute atoms can easily “trap” the twinning dislocation [32], and further impede the atomic shearing and shuffling despite the decreased SFE favorable for twinning [33,34]. It is therefore concluded that Al alloying weakens the twinning activity, which is experimentally supported by a great deal of prismatic slip, pyramidal slip, stacking faults, and nanotwins within low-Al-containing grains (Fig. 2, Fig. 3, Fig. 4). Comparison of twin density and thickness in the present study with currently available data reported in the literatures (pure titanium) also highlights the role of Al alloying on the suppression of deformation twinning [24,30,35]. It should be mentioned that the factor of SFE becomes increasingly important with elevated temperatures and our finding may be reversed due to temperature-dependent thermally-assisted enhanced atomic activity and rapidly reduced dislocation slip barriers [5].

4.2. Grain-level EBSD analysis for excluding the influence of grain orientation

The competitive relationship between Al-alloying-dependent reduced SFE and elevated slip barrier was used for rationalizing the experimentally observed deviation in deformation mechanisms. It should be clarified here that the deformation behaviors of as-selected grains are influenced not only by Al concentration but also by grain orientation with respect to loading direction. The contribution of the latter was not sufficiently considered although similar phenomena were frequently detected through repeating TEM examinations on many grains with different crystallographic orientations (see Fig. 5 for an example).

Fig. 5.

Fig. 5.   The influence of Al ordering on deformation mechanism of hexagonal close packed Ti-Al alloys. The Al concentration is measured as (a) 5 at.% and (b) 12 at.%, respectively. Inset is the electron diffraction of panel (b) and shows the presence of ordered Ti3Al precipitates.


In order to exclude the interference of the grain orientation, statistically-relevant large-scale EBSD analysis was performed on as-received annealed sample and deformed sample as shown in Fig. 6, Fig. 7. Several of inspirational findings are summarized as follows: (i) The average value of kernel average misorientation (KAM) is almost zero for annealed sample, suggesting a neglected effect of the gradient distribution of Al concentration with depth on the cooling-induced residual stress/deformation (Fig. 1, Fig. 6); in other words, plastic deformation of the present compositionally graded Ti-Al alloys does not show significant differences compared to traditional as-casted homogeneous counterparts with the same Al concentration [36]. (ii) After applying a plastic strain of 5%, we found that low-alloyed regions are plastically deformed more severely, supported by a larger KAM value and a great deal of deformation substructures such as sub- and low-angle grain boundaries in Grains B and C (Fig. 7). This finding also agrees well with our TEM results that only a few dislocations were observed in severely alloyed grains. (iii) EBSD-based slip trace analysis demonstrated a considerable amount of mobile dislocations gliding on {10-10} or {11-20} prismatic planes in the Grains A-C corresponding to dense parallel slip lines as shown in Figs. 7(c) and (d). (iv) Intra-granular EBSD characterization on individual grains such as Grain C (or other finer grains) demonstrated that slip bands and highly localized deformation were still detected frequently in the upper region of Grain C with low Al concentration, rather than the severely-alloyed lower part although they both belong to an identical grain with the same texture and boundary condition (Fig. 7(d)). This finding thus excludes the influence of the grain orientation in TEM-based deformation mechanism analysis, and highlights the importance of Al ordering. Coupling EBSD and TEM observation also permit ones to address the mutually competitive role played by Al alloying, and demonstrate that increasing Al concentration promoted chemical ordering, thus suppressing the dislocation motion and deformation twinning at room temperature.

Fig. 6.

Fig. 6.   Electron backscatter diffraction (EBSD) characterization of annealed samples at 1100 °C for 4 h. (a) Inverse pole figure (IPF). (b) Kernel average misorientation (KAM). (c) Grain boundary and (d) its frequency distribution histogram. All the grain boundaries were colored according to their misorientation angle between two neighboring grains: blue for 2°-5°, red for 5° -15°, and black for > 15°. Scale bar, 50 μm.


Fig. 7.

Fig. 7.   EBSD characterization of deformed microstructure after a tensile strain of 5%. (a) Inverse pole figure (IPF). (b) The distribution of low- and high-angle grain boundaries. All the grain boundaries were colored according to their misorientation angle between two neighboring grains: blue for 2°-5°, red for 5°-15°, and black for > 15°. (c) The image quality (IQ, Kikuchi pattern), and (d) Kernel average misorientation (KAM). (e) Compositional profile analysis from the topmost to bottom surface of panel (a), indicating the middle Al-poor region. (f-i) EBSD-based slip trace analysis and pole figures of (f) Grain A, (g) Grain B, and (h-i) Grain C as plotted in panel (a). Deformation substructure as well as shear localization was clearly observed, but was component and orientation dependent. In the interior of Grain C, the influence of grain orientation is minimized and we found that the closer to the Al poor region, the larger the KAM, and the greater the degree of plastic deformation. Scale bar, 50 μm.


4.3. Dislocation-twin interaction and its influence on mechanical properties

As one of deformation modes in hcp metals, twinning is of great importance as a supplement to <c + a> slip to accommodate the accumulated stress along crystallographic <c> direction upon loading [37]. The addition of Al atoms strains the matrix and enhances the critical resolved shear stress required for both dislocation slip and twin nucleation, beyond which twinning dislocations were driven to move onto the rigorously “invariant” {10-11} plane. This behavior is distinct with “stepped” {10-11} twin boundary as observed in deformed and annealed Mg-Gd alloys with severe Gd segregation since only plastic deformation was imposed in the present study without subsequent annealing [38].

Fig. 8 shows the bright field image taken ahead and right-side of Fig. 2(a). The presence of nanotwins serves as high-angle grain boundaries in obstructing the dislocation motion (Fig. 8(a)). Confined dislocation gliding toward another adjacent twin lamella is observed in Fig. 8(b), and its contribution to strengthening is inversely proportional to the square root of lamella spacing [39]. As a twin-mediated soft mode, dislocations emitting from the junction of twin/grain boundaries seems to have the ability to glide onto the coherent twin boundary which suffer from a very limited slip resistance (Fig. 8(a)). These unique dislocation dynamics, on the one hand, harden the matrix through the refinement of internal microstructure scale, and on the other hand, provide sufficient rooms for dislocation nucleation, accumulation and multiplication, thus facilitating work hardening and prolonging uniform plastic deformation [25,40]. Future work should be made on simultaneously achieving high strength (solution strengthening) and desired ductility (maximizing the twin activity) through the tailoring of Al ordering.

Fig. 8.

Fig. 8.   Nanotwins and their interactions with dislocations. (a) Influence of Al alloying on the deformation nanotwinning. (b) Confined dislocation slip within twin lamella with the thickness of a few or tens of nanometers.


5. Conclusion

In summary, we fabricated a compositionally graded Ti-Al alloy and investigated the influence of Al concentration on deformation mechanism at room temperature. Upon uniaxial tension, the mechanisms of prismatic slip, pyramidal slip, stacking faults, and deformation twins were activated in low-Al-containing regions; however these deformation modes were progressively inhibited with chemical ordering. Our finding also suggests that tailoring Al ordering for the maximization of twin activity seems plausible to simultaneously elevate the strength and ductility.

Acknowledgements

This work was financially supported by the National Natural Science Foundation of China (No. 51701081), the Key Research and Development Program of Shandong Province (Nos. 2019RKB01018, 2019GGX104077 and 2017GGX20122), the Young-aged Talents Lifting Project from Shandong Association for Science & Technology (No. 301-1505001, recoded by University of Jinan), and the Shandong Provincial Natural Science Foundation, China (Nos. ZR2018PEM008 and ZR2018PEM005).

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