Deformation twinning behaviors of the low stacking fault energy high-entropy alloy: An in-situ TEM study

Deformation twinning behaviors of the low stacking fault energy high-entropy alloy: An in-situ TEM study

Scripta Materialia 137 (2017) 9–12 Contents lists available at ScienceDirect Scripta Materialia journal homepage: www.elsevier.com/locate/scriptamat...

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Scripta Materialia 137 (2017) 9–12

Contents lists available at ScienceDirect

Scripta Materialia journal homepage: www.elsevier.com/locate/scriptamat

Regular article

Deformation twinning behaviors of the low stacking fault energy high-entropy alloy: An in-situ TEM study Jiabin Liu a, Chenxu Chen a, Yuqing Xu a, Shiwei Wu c, Gang Wang c, Hongtao Wang b,⁎, Youtong Fang d, Liang Meng a a

School of Materials Science and Engineering, Zhejiang University, Hangzhou 310027, China Institute of Applied Mechanics, Zhejiang University, Hangzhou 310027, China Department of Materials Science and Engineering, Shanghai University, Shanghai 200444, China d College of Electrical Engineering, Zhejiang University, Hangzhou 310027, China b c

a r t i c l e

i n f o

Article history: Received 6 February 2017 Received in revised form 14 April 2017 Accepted 1 May 2017 Available online 14 May 2017

a b s t r a c t The twinning mechanism in an FeCoNiCrAl0.1 high-entropy alloy was investigated by in-situ transmission electron microscopy. Deformation twins were commonly found on the conjugate plane and were produced by twinning dislocation gliding on every adjacent (111) plane layer by layer. The twinning dislocations were Shockley partials and were generated from the cross-slip. © 2017 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.

Keywords: Metals and alloys High-entropy alloys Transmission electron microscopy Twinning

High-entropy alloys (HEAs), with four or more equiatomic or nearequiatomic elements, have a single solid-solution phase with simple crystal structures such as face-centered cubic (FCC) or body-centered cubic, instead of generally complex phases and intermetallic compounds [1–4]. Highly distorted lattice structures in HEAs are expected because of the different atomic sizes and chemical bonds of the constituent elements. This may lead to a higher activation energy for dislocation nucleation in HEAs than for that in their conventional alloy counterparts [5]. This higher activation energy increases the slip resistance and thus affects the plastic deformation behaviors. However, the FeCoNiCrMn HEAs, a relatively extensively studied system, show that planar slips of 1/2 〈110〉 dislocations on {111} carry most of the plastic deformation thoughout a wide temperature range from room temperature to 1073 K [6,7]. Lowering the temperature to 77 K leads to simultaneous increases of the yield strength, ultimate tensile strength, and tensile ductility [8]. The fracture toughness values measured by compact tension tests at 77 K were over 200 MPa/m1/2 at crack initiation and 300 MPa/m1/2 in stable crack growth [8,9]. The proposed deformation mechanism was nanoscale twinning at low temperatures, as confirmed by post transmission electron microscopy (TEM) examination [8]. Although deformation twinning is important in determining the mechanical properties of HEAs, there are few direct experimental studies revealing the detailed behaviors and the underlying mechanisms. ⁎ Corresponding author. E-mail address: [email protected] (H. Wang).

http://dx.doi.org/10.1016/j.scriptamat.2017.05.001 1359-6462/© 2017 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.

In this work, in-situ TEM examination was carried out to investigate the twinning behavior of a model system, i.e., FeCoNiCrAl0.1, in term of its low stacking fault energy (SFE). The details of twin nucleation and thickening were uncovered in the FCC HEA. The in-situ study reveals that the deformation twinning behavior is largely determined by the interplay between twinning and dislocation activities. FeCoNiCrAl0.1 ingots were prepared by arc melting the constituent elements (N99.9% purity) in a water-cooled copper hearth under argon atmosphere. The ingots were re-melted five times to ensure chemical homogeneity. Then, FeCoNiCrAl0.1 sheets with dimensions of 2 mm × 10 mm × 100 mm were fabricated by suction casting using Cu molds. Slices with a thickness of 0.5 mm were cut from the as-cast FeCoNiCrAl0.1 sheets by electrical discharge machining. The slices were mechanically thinned to approximately 50 μm in thickness, and further electrolytically thinned by a twin-jet electropolisher. The electrolytical thinning process was carried out by using a mixture of 6 vol% perchloric acid, 30 vol% n-butanol and 64 vol% methanol at 12 V and −30 °C. Insitu TEM experiments were performed using a straining holder (Gatan 654) equipped in a JEM-2100 TEM operated at 200 kV. The specimens were strained by controlling the total elongation via a step motor in the straining holder. The minimum step was approximately 1 μm. The deformation process was recorded by a Gatan 831 charge-coupled device at a rate of 2 frames/s. SFE is one of the intrinsic parameters in determining the competition between the gliding of dislocations and twinning. The SFE value can be calculated by measuring the separation of partial dislocation pairs in

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Fig. 1. (a) A weak-beam dark-field TEM image of a partial dislocation pair in FeCoNiCrAl0.1; (b) The collected partial dislocation separations and dislocation character angles. The SFE was fitted to be 30 mJ/m2 (solid fitting curve), and the upper and lower boundaries (dashed fitting curves) were 34 mJ/m2 and 26 mJ/m2, respectively.

equilibrium according to [10] as following: d¼

  2 μb 2−v 2v cos2β 1− 2−v 8πγ 1−v

ð1Þ

where d is the separation distance between the 1/6 〈112〉 Shockley partials, μ the shear modulus, b the amplitude of the Burgers vector of the partials, ν the Poisson's ratio, β the dislocation character angle between the dislocation line and the diffraction vector. It is noted that the equilibrium condition is required for accurate measurements, and no external force, including imaging force, should have components in the gliding plane. To meet these conditions, the TEM samples were made from materials that were strained slightly over the yielding point. Fig. 1(a) displays a typical image of a partial dislocation pair obtained using the weak-beam dark-field TEM method. The partial dislocation pair is on the (111) habit plane when the incident beam is

parallel to the [111] direction imaged with a g = (220) vector. The separation distances between the partial dislocation pair and the corresponding character angles were directly measured from the weakbeam dark-field TEM images. Fig. 1(b) summarizes the d and β values from several partial dislocation pairs that met the equilibrium condition. According to Eq.(1), the SFE of the FeCoNiCrAl0.1 was calculated to be ~30 mJ/m2. Long stacking faults (SFs) are frequently observed in FeCoNiCrAl0.1 at low to medium strains. In the experimental setup in Fig. 2, the electron beam is aligned to [011] with the operation vector g = (002). The SFs are characterized by the typical banded contrast in the brightfield TEM images (Fig. 2(a)). Four gliding dislocations, indicated as D1–D4 in Fig. 2(a), were blocked by the SFs lying on the conjugate plane, which is inclined to the electron beam. As the straining holder only has the single tilting capability, the characterization of the dislocations on the same trace has been performed by transferring the sample

Fig. 2. Sequential snapshots of dislocation motion, extracted from the Supplemental Movie S1. (a)–(b) Bright-field TEM images of the dislocations cross-slipping from the primary plane onto the conjugate plane. (c) Bright-field TEM images of the cross-slipped dislocation split into a Frank partial and a Shockley partial. The gliding of the Shockley partial produced a stacking fault, (d) bright-field TEM image of the Frank partials on the conjugate plane as pointed out by dark arrows after tilt.

J. Liu et al. / Scripta Materialia 137 (2017) 9–12

to a double-tilt holder after the in-situ experiments. Series tilting under two-beam condition reveals the Burgers vector to be a/2 〈110〉 (Fig. S1.). It is noted that SFs can be effective barriers for intercrossed dislocation gliding [11]. Under straining, the dislocation pileup pushed D1 forward and led to cross-slip (Figs. 2(b–c), Supplementary Movie 1), leaving behind a Frank dislocation according to the following dislocation reaction [12]: i i i ah ah ah 110 þ 112 → 111 ð111Þ 2 3 ð111Þ 6 ð111Þ

ð2Þ

The Shockley partial moved away and formed a new SF parallel to the pre-existing one. The Shockley partial was identified because of the combined SFs, which showed banded contrast under the g = (002) condition. A similar process occured for subsequent dislocations and led to the formation of faulted zones. The Frank dislocations, as a byproduct, were observed by tilting the samples to the appropriate orientation (Fig. 2(d)). The Frank dislocations appeared as straight parallel dislocations associated with wide SFs under the g ¼ ð020Þ condition, which is consistent with the observation in Fe-Mn-C twinning induced plasticity steels [13]. The observation of Shockley partials and Frank dislocations on the conjugate plane supports the twinning mechanism proposed by Cohen and Weertman [12]. In their model, this mechanism would be activated when the full dislocation meets a Lomer-Cottrell barrier. Cross-slip has important implications for twin nucleation. One example is given in Fig. 3. To reveal the nucleation process, the conjugate sliding plane was tilted to the edge-on direction, as marked by slip plane B in Fig. 3(a). A narrow band with light gray contrast along the plane B is observed, which originated from the intersection of an array of dislocations on plane A in Fig. 3(a). The corresponding selected area electron diffraction (SAED) pattern in Fig. 3(b) is indexed to the matrix zone axis [01−1], showing no spots from a twin lamella. Because no other dislocation source can be found in the TEM view, this faulted zone is formed by the cross-slip of the dislocation array in plane A. Intensive gliding occurs under further in-situ straining, causing a final rupture along the slip trace (Fig. 3(c)). Twin lamella on the conjugate plane can be clearly identified by the SAED pattern in Fig. 3(d). The

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diffraction spots from the matrix and the twin are marked by the letters “M” and “T” respectively, showing the symmetry about the (1−1−1) plane. The twin thickness was measured to be approximately 60 nm in the dark-field TEM image (Fig. 3(e)). The intensive gliding process is shown in the online Supplementary Movie S2. The twin nucleation mechanism supports the models proposed by Cohen and Weertman [12] and by Fujita and Mori [14]. The thickening process of a twin band is shown in dark-field image sequences (online Supplementary Fig. S2 and Movie S3) by using the (111)T reflection. The initial twin band thickness increased from 30 nm to approximately 140 nm under continuous strain. It is noted that the band consisted of a twin bundle with a lamellar thickness less than 10 nm. The thickening and merging of twin lamellae are frequently observed during twin bundle growth, as shown in Fig. 4 and online Supplementary Movie S4. This process is characterized by slipping partial dislocations along twin boundaries. As shown in Fig. 4, a three-atomic-layer twin embryo increased the thickness to 15 layers and merged with the neighboring twin (Fig. 4(d)). The cross-slip mechanism may have important implications in the deformation of bulk materials as well. Ex-situ tensile tests have been performed on samples with a width of 5 mm and a gauge length of 25 mm. Intersection of deformation twins is frequently observed after tensile testing (Fig. S3), indicating activation of multiple slip systems as imposed by the deformation compatibility in polycrystals. Since the microstructural evolution information is missing in the post-TEM examination, it is not known how the cross-slip process affects the deformation behavior in bulk samples. However, the extended defects, such as long stacking faults or twin lamella, can easily block the motion of dislocations on the conjugate slip planes, as observed in both in-situ and ex-situ tensile experiments. Based on the in-situ experiments, the cross-slip process can be activated under such circumstance, leading to nucleation of deformation twins. The phenomenon is well studied by a recent large-scale molecular dynamic simulations on the formation of hierarchical twin structures in the low- SFE twin-induced-plasticity steels [11]. Since TBs are effective barriers for dislocation motion, a high work hardening rate is expected in FeCoNiCrAl0.1 and the crossslip mechanism will be further promoted [15,16].

Fig. 3. The full twinning process on the conjugate plane. (a) In the microstructure of the initial state, many dislocations on slip plane “A” are gliding and are blocked by slip plane “B”, (b) the corresponding SAED patterns of the dotted circle in (a) showing no twin spot, (c) the microstructure after deformation, (d) the corresponding SAED patterns of the dotted circle in (c) showing clear twin spots and (e) dark-field TEM image corresponding to (c), taken by using the (111)T spot. A deformation twin with a thickness of approximately 60 nm was generated on the conjugate plane.

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Fig. 4. Sequential snapshots of thickening of the deformation twins, extracted from the online Supplementary Movie 4. (a)–(d) High-resolution TEM images of the deformation twin growing to various layers.

In summary, the interplay of dislocations and twins in an FeCoNiCrAl0.1 HEA was investigated by in-situ TEM. The twin was produced by Shockley partials gliding on the adjacent (111) planes one by one. The Shockley partials were generated from the cross-slip of full dislocations on the primary plane. This work was financially supported by the National Natural Science Foundation of China (No. 11572281, 11,672,355 and u1434202) and the Fundamental Research Funds for the Central Universities (2016XZZX002–06). Supplementary data to this article can be found online at http://dx. doi.org/10.1016/j.scriptamat.2017.05.001. References [1] H.Y. Yasuda, K. Shigeno, T. Nagase, Scr. Mater. 108 (2015) 80–83. [2] J.W. Yeh, S.K. Chen, S.J. Lin, J.Y. Gan, T.S. Chin, T.T. Shun, C.H. Tsau, S.Y. Chang, Adv. Eng. Mater. 6 (5) (2004) 299–303.

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