In-situ high-resolution transmission electron

Journal of Alloys and Compounds 709 (2017) 802e807

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In-situ high-resolution transmission electron microscopy investigation of grain boundary dislocation activities in a nanocrystalline CrMnFeCoNi high-entropy alloy Qingyun Lin a, Xianghai An a, *, Hongwei Liu b, Qunhua Tang c, Pinqiang Dai d, e, Xiaozhou Liao a, ** a

School of Aerospace, Mechanical and Mechatronic Engineering, The University of Sydney, Sydney, NSW 2006, Australia Australian Centre for Microscopy & Microanalysis, The University of Sydney, Sydney, NSW 2006, Australia School of Mechanical & Electrical Engineering, Putian University, Putian 351100, China d College of Materials Science and Engineering, Fuzhou University, Fuzhou 350108, China e School of Materials Science and Engineering, Fujian University of Technology, Fuzhou 350108, China b c

a r t i c l e i n f o

a b s t r a c t

Article history: Received 27 November 2016 Received in revised form 16 March 2017 Accepted 18 March 2017 Available online 21 March 2017

Relaxation activities of grain boundary dislocations in a nanocrystalline CrMnFeCoNi high-entropy alloy were investigated using in-situ high-resolution transmission electron microscopy. 1/3 Frank dislocations at a low-angle grain boundary were relaxed by climbing or emitting Shockley partials that produced stacking faults slightly narrower than the theoretically predicted width. The extended dislocation structures were quite stable since no any change of the Shockley partial position or the stacking fault width was induced under further e-beam irradiation. These behaviours indicate easy activation of Shockley partials and strong barriers for dislocation motion in the alloy, which explains well the excellent deformability and strengthening effect of the alloy. In contrast, Frank dislocations at a twin boundary maintained a compact core structure, but were accompanied with a rotation process of the adjacent grain. © 2017 Published by Elsevier B.V.

Keywords: High-entropy alloy Nanocrystalline Grain boundary Dislocation Transmission electron microscopy

1. Introduction High entropy alloys (HEAs), which are solid solution alloys based on five or more elements with (near) equiatomic concentration, have attracted considerable attention in the materials research community [1e5]. Of these newly designed alloys, the equiatomic CrMnFeCoNi HEA alloy with a single face-centered cubic (FCC) crystal structure exhibits exceptional combination of mechanical properties with tensile strength of above 1 GPa, ductility of ~60e70%, and fracture toughness of exceeding 200 MPa m1/2 [6e9]. Its strength and damage tolerance can be enhanced with decreasing temperature due to the transition of the dominant deformation mechanism from dislocation slip to deformation twinning [6]. In-situ deformation transmission electron microscopy

* Corresponding author. ** Corresponding author. E-mail addresses: [email protected] (X. An), [email protected] (X. Liao). http://dx.doi.org/10.1016/j.jallcom.2017.03.194 0925-8388/© 2017 Published by Elsevier B.V.

(TEM) investigations revealed that the excellent mechanical properties are mainly attributed to the easy motion of Shockley partial dislocations and arrest of dislocations at slip bands [10]. It is well known that the microstructures of grain boundaries and grain boundary dislocations play crucial roles in the plastic deformation of materials, especially in materials with fine grain sizes. For the CrMnFeCoNi HEA alloy, although its strength obeys the classical Hall-Petch relationship, its hardening coefficient is much higher than that of conventional alloys [11]. Therefore, exploration of the boundary structures and dislocation activities will be beneficial for the deep understanding of the structureeproperty relationship in the HEA alloy. In the present study, we conducted in-situ high-resolution TEM (HRTEM) observations of grain boundary dislocation activities in a nanocrystalline (NC) CrMnFeCoNi alloy processed by high-pressure torsion (HPT). Boundary dislocations, verified as 1/3 Frank partial dislocations, display distinct relaxation behaviours at different types of boundaries when exposed to electron-beam (e-beam) irradiation. Moreover, the in-situ observation on the emission of a Shockley partial from one of the Frank dislocations and its pinning inside a

Q. Lin et al. / Journal of Alloys and Compounds 709 (2017) 802e807

nanograin lends support to the propositions of the easy motion of Shockley partials generating excellent deformability and the strong barrier for partial dislocation motion creating an outstanding strengthening effect in HEAs. 2. Material and methods The single phase FCC equiatomic CrMnFeCoNi alloy was produced by vacuum induction melting of the constituent elements with at least 99.9 wt% purity in a water-cooled copper crucible under a pure argon gas atmosphere. The alloy was repeatedly melted for five times to improve its chemical homogeneity and then solution-annealed at 1473 k for 5 h. For HPT processing, the alloy was cut into a disc sample with a diameter of 20 mm and a thickness of 1.5 mm. HPT processing was performed through 10 revolutions under a quasi-constrained condition with an imposed pressure of 5 GPa at room temperature [12]. Following the HPT processing, the sample was mechanically ground to a thickness of ~100 mm and then electropolished using a Struers TenuPol-5 jet electropolishing unit and a solution of 6 vol % perchloric acid in methanol under an operating voltage of 20 V at 30 C. Microstructural characterization by scanning electron microscopytransmission Kikuchi diffraction (SEM-TKD) [13], TEM and in-situ HRTEM was undertaken in regions at a distance of ~2 mm from the edge of the disc and all micrographs were taken from planview. The SEM-TKD experiment was performed in the Zeiss Ultra SEM with the operating voltage of 20 kV and the scanning step size of 10 nm, while TEM and in-situ HRTEM observations were conducted using a JEOL 3000F TEM operating at 300 kV.

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relatively random and no obvious texture was detected, careful inspection revealed that there was orientation fluctuation within grains with noticeable low-angle grain boundaries (LAGBs), as indicated by white lines. This is often associated with high internal stress or strain in materials processed by HPT [14]. The typical TEM image in Fig. 1(b) also presents the substantially refined microstructure. The inset indexed diffraction pattern substantiates the single FCC phase, which implies that no phase transformation occurred during the HPT process. Different from the SEM-TKD image in Fig. 1(a), grain boundaries are quite blurred and ill defined in TEM images due to the limited dynamic recovery [15,16]. Multiple types of grain boundaries provide an excellent opportunity for the investigation of the structures and behaviours of boundary dislocations. Fig. 2 (a) shows a small nanograin of less than 10 nm and two adjacent larger grains with sizes of ~55 nm and ~30 nm, respectively. An HRTEM image of the small nanograin and its surrounding grain boundaries was taken along the ½110 direction and displayed in Fig. 2(b). The left boundary noted with A was a low-angle tilt boundary, while the right one marked with B was a twin boundary. It is interesting to find that these two boundaries exhibited distinct behaviours under e-beam irradiation. For boundary A, dislocation dissociation and climb occurred, while boundary B accompanied with a rotation process of the right grain. There are always boundary dislocations at a LAGB. To understand the activities of these grain boundary dislocations, it is crucial to determine their Burgers vectors. As indicated in Fig. 3 (a), the LAGB with a symmetric tilting angle q of ~6 consists of an array of three dislocations with a constant inter-dislocation distance of 9 {111}layers. Therefore, the spacing between two neighboring

3. Results and discussion After HPT processing, the CrMnFeCoNi HEA alloy presented homogeneously distributed equiaxed nanograins with an average grain size of ~55 nm, as shown in the SEM-TKD image in Fig. 1(a). Although the crystallographic orientation of the nanograins was

Fig. 1. (a) An SEM-TKD image of CrMnFeCoNi after HPT processing. The black and white lines indicate high- and low-angle grain boundaries, respectively; (b) a brightfield TEM image and a corresponding electron diffraction pattern.

Fig. 2. (a) A bright-field TEM image showing a small nanograin of less than 10 nm and two adjacent larger grains with sizes of ~55 nm and ~30 nm, respectively; (b) an HRTEM image of the small nanograin and its neighboring areas, with a LAGB noted as A and a twin boundary marked as B.

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Fig. 3. ½110 HRTEM images of the dislocation array at boundary A during the e-beam irradiation at elapsing times of (a) 00 000 , (b) 00 5900 , (c) 104800 , and (d) 20 2200 .

dislocations Ddisl can be calculated as follows,

Ddisl ¼

pffiffiffi pffiffiffi 9*dð111Þ 9* 33 a 9* 33*0:359 nm ¼ ¼ ¼ 1:979 nm sin a sin a sin 70:53

Where a is the lattice parameter of FCC CrMnFeCoNi, known as 0.359 nm [17]; and a is the angle between (111) and ð111Þ, 70.53 . According to the Frank-Billy equation for LAGB structures [18,19], the estimated modulus of the Burgers vector is 0.2074 nm as follows,

q 6 jbj ¼ 2 tan Ddisl ¼ 2 tan *1:979 nm ¼ 0:2074 nm 2 2 where b is the Burgers vector of the dislocation; q is the tilt angle (6 ). There are three main types of dislocations in FCC structures, i.e. full dislocations with a Burgers vector of 〈110〉/2, and two types of partial dislocations with Burgers vectors of 〈111〉/3 and 〈112〉/6, respectively. The moduli of these three types of dislocations in CrMnFeCoNi can be calculated as follows,

pffiffiffi pffiffiffi < 110 > ¼ 2 a ¼ 2*0:359 nm ¼ 0:2539 nm 2 2 2

pffiffiffi pffiffiffi < 111 > ¼ 3 a ¼ 3*0:359 nm ¼ 0:2073 nm 3 3 3 pffiffiffi pffiffiffi < 112 > ¼ 6 a ¼ 6*0:359 nm ¼ 0:1466 nm 6 6 6 where a is the lattice parameter of FCC CrMnFeCoNi, known as 0.359 nm [17]. Considering the experimental error of ~ ±0.2 in measuring the tilt angle q along the tangent direction of the bright lattice points in the high-resolution micrograph, i.e., the tilt angle q should be 5.8 e6.2 , the calculated modulus of the Burgers vector should be in the range of 0.2005e0.2143 nm. Only the modulus of 0.2073 nm for Frank partial dislocations with Burgers vectors of 〈111〉/3 in the HEA alloy is close to the calculated modulus. Therefore, the dislocations in the boundary are of the 〈111〉/3 type. Since the left grain in Fig. 2 is not on an exact 〈110〉 zone axis and the lattice fringe is not very clear at the dislocation core area, it is not an easy task to draw Burgers circuits around the boundary dislocations. By recourse to the O-line method (details in the Supplementary Materials), it is verified that the LAGB dislocations are Frank partial dislocations. As shown in Fig. 3(bed), dissociation and climb of Frank dislocations occurred at boundary A during a 3-min e-beam irradiation


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Fig. 4. (a) An HRTEM image around boundary B at 00 0000 . (b) An HRTEM image around boundary B at 20 1100. Four dislocations at the boundary are marked by “T”. Inset at the left bottom is a corresponding fast Fourier transformation pattern indicating twin relationship between the two nanograins on each side of the boundary; (c) an enlarged image around P the top first dislocation and a Burgers circuit around the dislocation; (d) a Burgers circuit (FS/RH perfect crystal convention) transferred to ¼ 3 reference lattice, b is defined by the finish- (F) to-start (S) vector in a right-hand (RH) circuit that comes to closure around the dislocation but fails to close in the perfect crystal.

process. In the early stages of the e-beam irradiation process, there was no obvious microstructural evolution. Energy accumulation provided by the e-beam triggered dislocation activities [20]. For dislocation I, no significant change occurred except that a stacking fault tended to extend at the end of the e-beam irradiation process, which can be seen more clearly in the Supplementary movie. Dislocation II started to dissociate at 00 5900 and a Shockley partial dislocation was emitted through the following dislocation reaction:

1 1 1 ½111/ ½112 þ ½110 3 6 6 Supplementary video related to this article can be found at http://dx.doi.org/10.1016/j.jallcom.2017.03.194. The emitted Shockley partial dislocation subsequently glided away from the LAGB, resulting in a stacking fault that spanned 7 {111} atomic layers in the small nanograin, as revealed in Fig. 3(c) and (d). In contrast, dislocation III climbed along the grain boundary for 2 atomic layers. As such, the distance between dislocations II and III reduced from the initial 9 {111} atomic layers to 7 layers. Since dislocation climb is achieved by mass transport [21], the climb behaviour of dislocation III originated from the diffusion of HPT-induced interstitials under the e-beam irradiation [22]. Finally,

the reconstruction of the dislocation configurations at boundary A resulted in an increase of the tilting angle between the two adjacent grains from 6 to 7. These different dislocation activities can be attributed to their positions at the LAGB. Dislocations I and III were close to grain boundary triple junctions, which may significantly influence their behaviour under external stimuli. More investigations will be needed to understand this. In contrast, triple junctions had little effect on dislocation II. Therefore, the activity of dislocation II would be more representative of the behaviour of grain boundary dislocations relatively far away from triple junctions, which will be further discussed. Compared to the low-angle tilt grain boundary on the left, the small nanograin had a roughly twin relationship with the right grain, as revealed in the inset fast Fourier transformation pattern in Fig. 4(b). Although there was no apparent structural change at the twin boundary during the e-beam irradiation, the right grain underwent a rotation process. Fig. 4(a) exhibits that the original crystallographic direction of the right grain at 00 0000 was not on a low index zone axis. However, lattice fringes on ½110 zone axis in the top region of the right grain is quite clear after 201100 e-beam exposure as observed in Fig. 4(b). Comparison of Fig. 4(a) and (b) shows almost no change of ð111Þ and ð111ÞT planes before and after

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the e-beam irradiation. Since the crystallographic orientation of the small nanograin remained almost the same, the clear lattice image should stem from a rotation process of the right grain around an axis perpendicular to the grain boundary plane. In addition, it is noticeable that there are four dislocations at the boundary, as marked by “T”. Fig. 4(c) presents an enlarged image including the top first dislocation with a Burgers circuit around. The 90 angle is marked to show that the circuit turned from (111)T plane to the vertical plane ð211ÞT . The FS/RH perfect crystal convention was applied to define the dislocation character [23], as shown in Fig. 4(d), which defines b by the finish- (F) to-start (S) vector in a right-hand (RH) circuit that comes to closure around the dislocation but fails to close in the perfect crystal and identifies that the nature of the dislocation is also a Frank dislocation with a Burgers vector of 1/3 〈111〉. Based on the HRTEM observations and simulations using the embedded-atom potential (EAP) method [24], 1/3 〈111〉 dislocations at twin boundaries can present two distinct relaxed processes, which depends on the position of the extra half plane with respect to the twin boundary plane. When the extra plane is contained in the reflex angle defined by the two sets of inclined ð111Þ and ð111ÞT planes, the Frank dislocations can relax by emitting a 90 Shockley partial dislocation followed by an intrinsic stacking fault. Conversely, if the extra plane is contained in the obtuse angle defined by the same sets of {111} planes, the 1/3 〈111〉 dislocation presents a compact core structure. Herein, the geometrical relationship between the extra half (111) planes of the four dislocations and the twin boundary plane, as shown in Fig. 4 (b), conforms to the second case. Therefore, in lieu of dissociation of a Frank dislocation observed in the LAGB above, the 1/3 [111] dislocations in the twin boundary would present a compact core structure without any apparent change during the e-beam irradiation process, while the grain rotated around an axis perpendicular to the grain boundary plane as the relaxation mode. Previous investigations revealed that in materials with low stacking-fault energies, a Frank partial dislocation might relax by the emission of a Shockley partial dislocation, which creates a stacking fault and leaves behind a stair-rod dislocation [23,25]. Considering the CrMnFeCoNi alloy with a low stacking-fault energy value of ~21 mJ/m2 [26], the dissociation of the Frank dislocation at boundary A is quite reasonable. Moreover, the high internal stress induced by HPT may also facilitate the dissociation of the Frank dislocation to reduce the system energy. As mentioned above, the emitted Shockley dislocation glided along the ð111Þ plane and stopped at the grain interior, leaving a stacking fault with a width of 7 {111} layers. In this case, a repulsive elastic force exists between the two dissociated dislocations and an attractive force of the stacking fault acts on them. Theoretically, the stacking fault should extend until the two forces balance. The width of the stacking fault can be calculated from the following equation [27]:

mb1 $b2 ð2 nÞ 2n 1 cos 2 q 8pð1 nÞd 2n mjb1 jjb2 jcos 4ð2 nÞ 2n ¼ 1 cos 2 q 8pð1 nÞd 2n

gSF ¼

where, gSF is stacking fault energy of the material (~21 mJ/m2), m is Shear modulus (80 GPa [26]), n is Poisson's ratio (0.25625 [26]), b1 and b2 are the Burgers vectors of two partial dislocations 16 ½112 and pffiffiffi pffiffiffi 1 ½110 (jb j ¼ 6*0:359 nm ¼ 0:1465 nm and jb j ¼ 2*0:359 1 2 6 6 6 nm ¼ 0:0846 nm) respectively, 4 is the included angle between 1 ½112 and 1 ½110 (54.74 ), d is the separation width of the two 6 6 dislocations (the width of the stacking fault), q is the discrepancy angle between the line direction of the stacking fault and the

Burgers vector (0 ). By substituting the corresponding values, the calculated width of the stacking fault in this case is 1.79 nm, which corresponds to the width of 8 {111} layers. The width of the stacking fault in the present study was 7 {111} layers, which is slightly lower than the theoretical value. However, longer time e-beam irradiation after 20 2200 did not induce any change of the Shockley partial position or the stacking fault width. This implies that the Shockley partial was pinned in the grain and the present configuration of the dislocations and stacking fault was relatively stable. The relatively stable extended dislocation structure in the HEA alloy can be mainly attributed to its lattice nature. It is well known that the lattice structure of HEA alloys is highly distorted due to different atomic sizes and chemical bonds of the constituent elements. This induces local elastic stress fields, which would interact with the stress field of dislocations and remarkably hinder dislocation motion that consequently increases strength [28]. The above observation provides well evidence at the atomic scale for the excellent deformability and an outstanding strengthening effect of the HEA alloy due to the easy activation of Shockley partial dislocations and the strong barrier for partial dislocation motion, respectively. 4. Conclusions In summary, by recourse to in-situ HRTEM observations, we studied the activities of 1/3 〈111〉 Frank partial dislocations located at a LAGB and a twin boundary in a nanocrystalline CrMnFeCoNi HEA alloy under e-beam irradiation. At the LAGB, Frank dislocations relaxed via dissociation and climb, while at the twin boundary the dislocation array showed insignificant change but accompanied with some rotation of one of the two twinned grains during the irradiation. The dissociation of a Frank dislocation at a LAGB emitted a Shockley partial and the emitted partial glide into the grain and finally terminated at the interior, leaving a 7-layer stacking fault, which is slightly narrower than that predicted theoretically. These experimental observations deepen our understanding of dislocation behaviors of HEAs and may also provide guidelines for the future design of advanced metallic materials with superior mechanical properties. Acknowledgements The authors acknowledge the scientific and technical input and support from the Australian Microscopy & Microanalysis Research Facility node at the University of Sydney. This project was supported by the Australian Research Council (DP150101121). Appendix A. Supplementary data Supplementary data related to this article can be found at http:// dx.doi.org/10.1016/j.jallcom.2017.03.194. References [1] J.W. Yeh, S.K. Chen, S.J. Lin, J.Y. Gan, T.S. Chin, T.T. Shun, C.H. Tsau, S.Y. Chang, Nanostructured high-entropy alloys with multiple principal elements novel alloy design concepts and outcomes, Adv. Eng. Mater. 6 (2004) 299e303. [2] Y. Zhang, T.T. Zuo, Z. Tang, M.C. Gao, K.A. Dahmen, P.K. Liaw, Z.P. Lu, Microstructures and properties of high-entropy alloys, Prog. Mater. Sci. 61 (2014) 1e93. [3] Q.H. Tang, Y. Huang, Y.Y. Huang, X.Z. Liao, T.G. Langdon, P.Q. Dai, Hardening of an Al0.3CoCrFeNi high entropy alloy via high-pressure torsion and thermal annealing, Mater. Lett. 151 (2015) 126e129. [4] G. Laplanche, O. Horst, F. Otto, G. Eggeler, E.P. George, Microstructural evolution of a CoCrFeMnNi high-entropy alloy after swaging and annealing, J. Alloys Compd. 647 (2015) 548e557. [5] B. Gludovatz, E.P. George, R.O. Ritchie, Processing, microstructure and mechanical properties of the CrMnFeCoNi high-entropy alloy, JOM 67 (2015) 2262e2270.

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