Microstructures and mechanical properties of TixNbMoTaW refractory high-entropy alloys

Microstructures and mechanical properties of TixNbMoTaW refractory high-entropy alloys

Author’s Accepted Manuscript Microstructures and mechanical properties of TixNbMoTaW refractory high-entropy alloys Z.D. Han, H.W. Luan, X. Liu, N. Ch...

1MB Sizes 0 Downloads 63 Views

Author’s Accepted Manuscript Microstructures and mechanical properties of TixNbMoTaW refractory high-entropy alloys Z.D. Han, H.W. Luan, X. Liu, N. Chen, X.Y. Li, Y. Shao, K.F. Yao www.elsevier.com/locate/msea

PII: DOI: Reference:

S0921-5093(17)31601-5 https://doi.org/10.1016/j.msea.2017.12.004 MSA35846

To appear in: Materials Science & Engineering A Received date: 3 September 2017 Revised date: 1 December 2017 Accepted date: 2 December 2017 Cite this article as: Z.D. Han, H.W. Luan, X. Liu, N. Chen, X.Y. Li, Y. Shao and K.F. Yao, Microstructures and mechanical properties of Ti xNbMoTaW refractory high-entropy alloys, Materials Science & Engineering A, https://doi.org/10.1016/j.msea.2017.12.004 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.

Microstructures and mechanical properties of TixNbMoTaW refractory high-entropy alloys Z.D. Han a, H.W. Luan a, X. Liu a, b, N. Chen a, *, X.Y. Li a, Y. Shao a, K.F. Yao a, * a

School of Materials Science and Engineering, Tsinghua University, Beijing 100084, China Institute of Materials, China Academy of Engineering Physics, Mianyang 621907, China *Corresponding authors: N. Chen and K.F. Yao. E-mails: [email protected] mail.tsinghua.edu.cn; [email protected] b

Abstract Refractory high-entropy alloys (RHEAs) are newly developed candidate materials for high-temperature applications. Among the existing RHEAs, NbMoTaW RHEA possesses the best mechanical properties with combined high strength, excellent thermal stability and softening resistance at elevated temperatures. However, the NbMoTaW RHEA is quite brittle at room temperature, which would restrict its application as structural material. Here, TixNbMoTaW RHEAs were developed by alloying Ti in the NbMoTaW RHEA. It shows that the room temperature ductility of the RHEAs increases from 1.9% of the NbMoTaW RHEA to 11.5% of the TiNbMoTaW RHEA, and the yield strength increases from 996 MPa of the NbMoTaW RHEA to 1455 MPa of the TiNbMoTaW RHEA. In addition, the TixNbMoTaW RHEAs keep stable single BCC structure up to their melt points. The present result indicates that Ti addition could effectively enhance both the ductility and strength of the NbMoTaW RHEA. The combined performance of superior mechanical properties and high thermal stability of the TixNbMoTaW RHEAs promises them an important role in engineering applications. 1

Keywords: High entropy alloy; Alloy design; Microstructure; Mechanical property; Thermodynamic calculation.

1. Introduction Structural materials with superior mechanical properties at elevated temperatures are in high demand for the development of gas turbine engines. Since raising the working temperature in engines can effectively increase the efficiency of gas turbines [1], the requirement for high-temperature materials keeps increasing. Among the applied high-temperature alloys, nickel-based superalloys exhibit the best elevated temperature properties and have been extensively used. However, the service temperatures of the nickel-based superalloys are approaching their melt points [2], making the further improvement of turbine efficiency to be extremely difficult. Therefore, it is of great importance to develop new elevated temperature materials that could work at higher temperatures. The newly developed high-entropy alloys (HEAs), consisting of four or more principal metallic elements in equiatomic or near-equiatomic ratios, have received more and more attention due to their unique structures and excellent mechanical properties [3-5]. Considered as new types of superalloys, HEAs possess outstanding strength and hardness [4,6-13], high oxidation resistance [14-16], high thermal stability and high softening resistance at elevated temperatures [17,18]. By introducing refractory elements in HEAs, the developed refractory HEAs (RHEAs) are mainly composed of Ti, V, Zr, Nb, Mo, Ta, W and Hf, which may 2

exhibit better high-temperature performance. Among the reported RHEAs [19-25], NbMoTaW RHEA possesses the highest softening resistance at elevated temperatures [20]. Therefore, it has received the most attention and been extensively studied as the potential candidate in gas turbine engine materials [20,26-28]. However, the NbMoTaW RHEA is brittle at room temperature, posing a significant challenge for the manufacturing and assembling of workpieces made from this alloy. To tackle this issue, a critical step towards the application is to improve the room temperature ductility of the NbMoTaW RHEA. The brittle feature of the NbMoTaW RHEA is supposed to be caused by Mo and W elements, which are intrinsically brittle pure metals. Thus, decreasing their contents or removing them could effectively improve the ductility of the RHEAs. J. W. Qiao et.al [29] designed NbTaV-(Ti,W) RHEAs by removing Mo element. Among these RHEAs, the NbTaVTi alloy without W element possesses the highest ductility. However, removing the brittle refractory elements decreases the mechanical performance of the RHEAs at high temperatures. Decreasing average valence electron concentration (VEC) was reported to make Mo and W-containing alloys more ductile [30]. Meanwhile, this method has been applied in the RHEA design recently [31]. Based on this strategy, alloying method was employed to enhance the ductility of the NbMoTaW RHEA in present work. Ti alloys possess excellent mechanical properties, and good thermal stability. Moreover, Ti element possesses small heats of mixing with Nb, Mo, Ta and W elements and lowest VEC compare to other elements. Therefore, introducing Ti into the NbMoTaW 3

RHEA could improve the mechanical properties and besides sustain the microstructural stability. In this paper, TixNbMoTaW RHEAs were prepared with different Ti additions, based on which the effects of Ti additions on their structural stability and mechanical properties were investigated.

2. Experimental procedures TixNbMoTaW RHEAs (the molar ratio x=0, 0.25, 0.5, 0.75 and 1, denoted by Ti0, Ti0.25, Ti0.5, Ti0.75 and Ti1, respectively) were prepared using vacuum arc melting of the corresponding raw elements. The nominal chemical compositions of the TixNbMoTaW RHEAs are provided in Table 1. The purities of each element are higher than 99.96% (by weight). To make sure that the ingots are as homogeneous as possible, each ingot was re-melted at least five times. The ingots were cut into cuboid samples with a dimension of 3 mm × 3 mm × 6 mm for studying their microstructure and mechanical properties. Table 1. Nominal chemical compositions (in at.%) of the TixNbMoTaW RHEAs. Alloys/element

Nb

Mo

Ta

W

Ti

Ti0

25

25

25

25

0

Ti0.25

23.5

23.5

23.5

23.5

5.9

Ti0.5

22.2

22.2

22.2

22.2

11.1

Ti0.75

21.1

21.1

21.1

21.1

15.8

Ti1

20

20

20

20

20

Equilibrium phase diagrams of the TixNbMoTaW RHEAs were calculated by Thermal-Calc software with TTNI 8 database. The crystal structures of the RHEAs were characterized by X-ray diffraction (XRD, Rigaku D/max-RB) using a Cu Kα 4

radiation. The thermal parameters of the TixNbMoTaW RHEAs were analyzed by a simultaneous thermal analyzer (TG-DSC; Netzsch STA449F3) in a flowing argon atmosphere with a heating rate of 20 K min-1. The microstructures and chemical compositions were examined by Quanta 200 FEG scanning electron microscope (SEM) with energy dispersive spectrometer (EDS) accessory. The room temperature mechanical properties of the RHEAs were conducted by WDW-50 testing machine at a strain rate of 5.6×10-4 /s using the prepared cuboid samples. The hardness of the RHEAs was measured by MH-3 Vickers hardness tests.

3. Results and discussions 3.1. Microstructures

The calculated equilibrium phase diagrams of the TixNbMoTaW RHEAs are shown in Fig. 1. With increasing Ti content, the melting temperature of the TixNbMoTaW RHEAs decreases, which is due to that Ti has the lowest melting temperature compared to Nb, Mo, Ta and W elements. Below the melt point of the TixNbMoTaW RHEAs, the liquid is solidified in a single BCC phase. At temperatures above 1000 ˚C, the BCC phase is quite stable and no phase transformation occurs. Below 640 ˚C, a HCP phase appears in Ti0.25 alloy with Ti addition, which should mainly contain Ti element.

5

Fig. 1. Calculated equilibrium phase diagrams of the TixNbMoTaW RHEAs: a, Ti0 alloy; b, Ti0.25 alloy; c, Ti0.5 alloy; d, Ti0.75 alloy; e, Ti1 alloy.

Upon solidification, the formation temperature of the HCP phase increases with increasing Ti content. B2 phase mainly containing Mo and Ta elements appears at 480 ˚C in Ti0.25 alloy and the formation temperature of the B2 phase remains a constant with different Ti additions as shown in Fig. 1. The equilibrium phase diagrams indicate that the thermal stability of the BCC phase in the TixNbMoTaW RHEAs decreases with increasing Ti. Despite this, the BCC phase of all the Ti xNbMoTaW RHEAs maintains high thermal stability even at temperatures above 2500 ˚C, which is 6

much higher than the service temperatures for the present gas turbines.

Fig. 2. The XRD patterns of the as-cast TixNbMoTaW RHEAs.

Fig. 2 shows the XRD patterns of the as-cast TixNbMoTaW RHEAs. Single BCC phase is identified in all of the as-cast TixNbMoTaW RHEAs. With increasing Ti content, the peaks of the samples shift to lower 2θ. Therefore, the lattice constant of the as-cast TixNbMoTaW RHEAs increases from 3.226 Å of Ti0 alloy to 3.240 Å of Ti1 alloy. The increased lattice constant is resulted by the solute Ti with the largest radius compared with other elements in the TixNbMoTaW RHEAs.

Fig. 3. The DSC curves of the NbMoTaW and TiNbMoTaW RHEAs.

Fig. 3 shows the DSC results of NbMoTaW and TiNbMoTaW RHEAs. No 7

endothermic or exothermic peak is identified in the DSC curves of both NbMoTaW and TiNbMoTaW RHEAs, indicating that no phase transformation occurs and the TixNbMoTaW RHEAs possess high thermal stability in the range of 200 ˚C-1200 ˚C. Both the XRD and DSC results show that only single BCC phase is formed and no phase transformation occurs in the as-cast TixNbMoTaW RHEAs, which is different from the prediction of multi-phase formation in the TixNbMoTaW RHEAs by the calculated equilibrium phase diagrams. However, both the CALPHAD and DSC results imply that the TixNbMoTaW RHEAs keep stable crystal structure up to the melt points. The high thermal stability may be contributed by the high entropy effect [32], which could restrain the phase transformation or separation in the HEAs. Commonly, several parameters of the mixture entropy (ΔS) [33], mixture enthalpy (ΔH) [33], atomic packing parameter (γ) [34], valence electron concentration (VEC) [35] and thermodynamic parameter (Ω) [33] are employed to predict the phase selection and phase stability in HEAs. They are defined as following and are listed in Table 2: ∆𝑆 = −𝑅 ∑𝑛𝑖=1 𝑐𝑖 ln𝑐𝑖

(1)

∆𝐻 = 4 ∑𝑛𝑖=1,𝑖≠𝑗 ∆𝐻𝑖𝑗𝑚𝑖𝑥 𝑐𝑖 𝑐𝑗

(2)

(𝑟𝑆 +𝑟̅ )2 −𝑟̅ 2

𝛾 = (1 − √

(𝑟𝑆

+𝑟̅ )2

(𝑟𝐿 +𝑟̅ )2 −𝑟̅ 2

)⁄(1 − √

(𝑟𝐿 +𝑟̅ )2

); 𝑟̅ = ∑𝑛𝑖=1 𝑐𝑖 𝑟𝑖

𝑉𝐸𝐶 = ∑𝑛𝑖=1 𝑐𝑖 (𝑉𝐸𝐶)𝑖 𝛺=

𝑇𝑚 ∆𝑆𝑚𝑖𝑥 ; |∆𝐻𝑚𝑖𝑥 |

𝑇𝑚 = ∑𝑛𝑖=1 𝑐𝑖 (𝑇𝑚 )𝑖

(3) (4) (5)

Where R is gas constant, H ijmix is the mixture enthalpy of the ith and jth components, ci and cj are the atomic percentage of the ith and jth components, 8

respectively, rS and rL are the radius of the smallest and largest atoms, ri is the atomic radius of the ith component, (VEC)i is the VEC of the ith component, (Tm)i is the melting temperature of the ith element, respectively. Table 2. The ΔS, ΔH, γ, Ω, and VEC of the TixNbMoTaW RHEAs. Alloys

ΔS (J/K·mol)

ΔH (kJ/mol)

γ

Ω

VEC

Ti0

11.53

-3

1.053

3.84

5.5

Ti0.25

12.71

-3.04

1.079

4.17

5.41

Ti0.5

13.15

-3.06

1.079

4.29

5.33

Ti0.75

13.33

-3.06

1.079

4.36

5.26

Ti1

13.38

-3.04

1.079

4.40

5.2

With increasing Ti content, the mixture entropy of the TixNbMoTaW RHEAs increases from 11.53 J/K·mol of Ti0 alloy to 13.38 J/K·mol of Ti1 alloy. The high mixture entropy of the TixNbMoTaW RHEAs could make the solid solution structures more stable. The mixture enthalpy represents the interaction between different components. More negative ΔH means larger binding force between the elements, which makes it easier to form intermetallics. In contrast, positive ΔH promotes the separation of the components. The mixture enthalpy values of the TixNbMoTaW RHEAs are all close to 0 so that solid solution structures are preferred. The values of parameter γ of the TixNbMoTaW RHEAs are all smaller than the critical value of 1.175 while the values of Ω of the TixNbMoTaW RHEAs are larger than 1.1. These results indicate that the solid solution structures of TixNbMoTaW RHEAs are stable [33,34]. With increasing Ti content, the VEC value of the TixNbMoTaW RHEAs 9

decreases from 5.5 of Ti0 alloy to 5.2 of Ti1 alloy. All of them are much smaller than 6.87, indicating that the crystal structures of the TixNbMoTaW RHEAs are stable in single BCC solid solution structure [35].

Fig. 4. The BSE images of the as-cast TixNbMoTaW RHEAs: a, Ti0 alloy; b, Ti0.25 alloy; c, Ti0.75 alloy; d, Ti1 alloy.

The SEM backscatter electron (BSE) images of the as-cast TixNbMoTaW RHEAs are shown in Fig. 4. It can be seen that all of the as-cast TixNbMoTaW RHEAs possess dendrite structures. Element distributions of the as-cast TixNbMoTaW RHEAs are further identified by an EDS method. The EDS results of the as-cast TixNbMoTaW RHEAs are shown in Table 3. The bright regions in Fig. 4 are enriched with the heavy elements like W and Ta, while the dark regions are enriched with the light elements such as Ti, Nb and Mo. 10

Table 3. The EDS results of the as-cast TixNbMoTaW RHEAs in different regions. Concentrations (at.%)

Ti

Nb

Mo

Ta

W

Dendrite arm

-

19.4

23.1

25.9

31.6

Interdendrite region

-

28.2

29.6

22.7

19.5

Dendrite arm

0.5

23.4

25.0

24.5

26.6

Interdendrite region

2.3

28.9

27.8

22.3

18.7

Dendrite arm

4.4

21.2

22.2

23.7

28.5

Interdendrite region

12.8

29.7

25.2

18.5

13.8

Dendrite arm

7.1

19.0

22.7

24.6

26.6

Interdendrite region

19.7

29.9

26.1

16.2

8.1

Dendrite arm

8.6

19.0

20.2

23.0

29.2

Interdendrite region

29.2

24.2

23.4

14.9

8.3

Ti0

Ti0.25

Ti0.5

Ti0.75

Ti1

3.2 Mechanical properties

Fig. 5. a, The compressive stress-strain curves of the as-cast TixNbMoTaW RHEAs at room temperature; b, The yield strength and ductility of the as-cast TixNbMoTaW RHEAs with various Ti contents. 11

The engineering stress-strain curves of the as-cast TixNbMoTaW RHEAs at room temperature are shown in Fig. 5a, based on which the derived mechanical parameters of all the RHEAs are listed in Table 4. The yield strength value (σ0.2), the peak strength (σp) and the plastic strain (εp) of the Ti0 alloy at room temperature are 996 MPa, 1148 MPa and 1.9%, respectively (Fig. 5a). As the Ti content increases from Ti0 alloy to Ti1 alloy, the yield strength and plastic strain of the as-cast TixNbMoTaW RHEAs increase substantially. The σ0.2 of Ti1 alloy is 1455 MPa whereas εp increases to 11.5%. The yield strength and the plastic strain of the as-cast TixNbMoTaW RHEAs versus Ti content are shown in Fig. 5b. It shows that both the yield strength and plastic strain of the as-cast TixNbMoTaW HEAs increase nearly linearly with Ti content. Table 4. Room-temperature yield strength (σ0.2), peak strength (σp) and plastic strain (εp) of the as-cast TixNbMoTaW RHEAs. Alloys

σ0.2 (MPa)

σp (MPa)

εp (%)

Ti0

996

1148

1.9

Ti0.25

1109

1197

2.5

Ti0.5

1211

1578

5.9

Ti0.75

1304

1593

8.4

Ti1

1455

1910

11.5

12

The strength and ductility of the TixNbMoTaW RHEAs at room temperature increase with Ti additions. The mechanism may be complex due to a combination of multi-effects. Firstly, the addition of large radius Ti increases the lattice constant of the TixNbMoTaW RHEAs, which further enhances the effect from solid solution hardening. The effect of solid solution hardening contributed by the lattice distortion can be described [36,37] as below: 𝜅 = ∑𝑛𝑖=1 𝑐𝑖 𝛿𝑖 ⁄𝑟𝑖

(6)

Where ci is the atomic percentage of the ith component, δi is the atomic size misfit of the ith component, ri is the atomic radius of the ith component, respectively. The atomic size misfit δi can be evaluated approximately. For example, δTi in the TiNbMoTaW RHEA can be approximated by: 𝛿Ti = 𝑆 Ti − 𝑆 NbMoTaW

(7)

Where STi and SNbMoTaW are the average interatomic spacing of pure Ti and equiatomic NbMoTaW RHEA, respectively. The average interatomic spacing of the TixNbMoTaW RHEAs can be calculated with the lattice constants measured by the XRD methods. The strengthening coefficients κ of the Ti0, Ti0.25, Ti0.5, Ti0.75 and Ti1 RHEAs are 2.4×10-4, 1.3×10-3, 2.8×10-3, 4.9×10-3 and 6.9×10-3, respectively. Thus, the strength of the TixNbMoTaW RHEAs increases with Ti additions. Secondly, the grain boundary cohesion of the TixNbMoTaW RHEAs has been improved by Ti additions. The fracture images of the Ti0 and Ti1 alloys are shown in Fig. 6. The Ti0 alloy (Fig. 6a) shows an intergranular fracture mode, similar to the report in Ref [20], While Ti1 alloy (Fig. 6b) shows transgranular fracture morphology. 13

The transition of the fracture mode indicates that the grain boundary cohesion of TixNbMoTaW RHEAs is improved by Ti additions. The improvement of grain boundary cohesion could efficiently suppress the intergranular fracture [38,39]. As a result, the ductility of the TixNbMoTaW RHEAs was enhanced by Ti additions [40,41]. In addition, the compositional segregation becomes more significant with increasing Ti content (as seen in Table 3). The dendrite arms mainly containing W element are strong but brittle, whereas the inter-dendrite regions enriched in Ti, Nb and Mo elements are less strong but relatively ductile. In the compressive process, the ductile inter-dendrite regions enriched with Ti and Nb elements are more prone to deform with response to the applied stress and are able to bear more strains than the strong and brittle dendrite arms enriched with W element. Consequently, the ductility of the TixNbMoTaW RHEAs increases with Ti additions [37,42].

Fig. 6. The SEM images of fracture surfaces of TixNbMoTaW RHEAs: a, Ti0 alloy; b, Ti1 alloy.

Thirdly, L. Qi et.al found [30] that decreasing VEC can improve the ductility of brittle refractory alloys containing Mo/W by alloying transition metals such as Ti, Zr or Hf. The crystal orientation of <100> is known to be weakest for refractory alloys to bear external stress. Under the tensile direction <100>, lots of dislocations are 14

activated, provided that shear instability occurs before the ideal tensile stress is reached. In this case, the alloys could become intrinsically ductile [30]. Alloying brittle Mo/W alloys with Ti, Zr or Hf decreases the VEC values. As a result, the Fermi level of the system decreases and the shear instability indeed occurs earlier, therefore improving the ductility of the alloys. Recently, a new ductile RHEA, Hf0.5Nb0.5Ta0.5Ti1.5Zr has been produced [31] based on the electron theory above. Additionally, the data of other RHEAs also showed that lower VEC can make RHEAs more ductile [29,43,44]. In the TixNbMoTaW RHEAs, the VEC decreases from 5.5 of Ti0 alloy to 5.2 of Ti1 alloy with increasing Ti content (as seen in Table 2), which may efficiently improve the ductility of the TixNbMoTaW RHEAs.

4. Conclusions With Ti additions, TixNbMoTaW RHEAs (x=0, 0.25, 0.5, 0.75 and 1) have been developed. The TixNbMoTaW RHEAs possess single BCC crystal structure, which keeps stable up to their melt points. Alloying with Ti element effectively improves both the strength and ductility of the TixNbMoTaW RHEAs at room temperature. Specifically, the yield strength of the TiNbMoTaW alloy is 1455 MPa, which is 46% higher than that of NbMoTaW RHEA. While its ductility at room temperature is as high as 11.5%, which is about 5 times higher than that of NbMoTaW RHEA. The strength enhancement of the TixNbMoTaW RHEAs can be interpreted by the solid solution hardening effect. The improved ductility of the TixNbMoTaW RHEAs can be understood by the grain boundary cohesion improvement, ductile elemental 15

segregation and shear instability. The excellent mechanical performance and high thermal stability of these RHEAs make them hold potentials for applications as structural materials.

Acknowledgments This work was supported by the National Natural Science Foundation of China (NSFC, Grant Nos. 51571127 and 51271095).

Reference [1] J.H. Perepezko, Science 326 (2009) 1068-1069. [2] J.A. Lemberg, R.O. Ritchie, Adv. Mater. 24 (26) (2012) 3445-3480. [3] B. Cantor, I.T.H. Chang, P. Knight, A.J.B. Vincent, Mater. Sci. Eng. A 375-377 (2004) 213-218. [4] Y.J. Zhou, Y. Zhang, Y.L. Wang, G.L. Chen, Appl. Phys. Lett. 90 (18) (2007) 181904. [5] Y.F. Ye, Q. Wang, Y.L. Zhao, Q.F. He, J. Lu, Y. Yang, J. Alloy. Compd. 681 (2016) 167-174. [6] Y. Dong, Y.P. Lu, J.R. Kong, J.J. Zhang, T.J. Li, J. Alloy. Compd. 573 (2013) 96-101. [7] S. Liu, M.C. Gao, P.K. Liaw, Y. Zhang, J. Alloy. Compd. 619 (2015) 610-615. [8] J.M. Zhu, H.M. Fu, H.F. Zhang, A.M. Wang, H. Li, Z.Q. Hu, Mater. Sci. Eng. A 527 (26) (2010) 6975-6979. [9] H.Y. Ding, Y. Shao, P. Gong, J.F. Li, K.F. Yao, Mater. Lett. 125 (2014) 151-153. [10] H.Y. Ding, K.F. Yao, J. Non-Cryst. Solids 364 (2013) 9-12. [11] Z.D. Han, X. Liu, S.F. Zhao, Y. Shao, J.F. Li, K.F. Yao, Prog. Nat. Sci. Mater. Int. 25 (5) (2015) 365-369. [12] S.F. Zhao, Y. Shao, X. Liu, N. Chen, H.Y. Ding, K.F. Yao, Mater. Des. 87 (2015) 625-631. [13] S.F. Zhao, G.N. Yang, H.Y. Ding, K.F. Yao, Intermetallics 61 (2015) 47-50. [14] T.M. Butler, M.L. Weaver, J. Alloy. Compd. 674 (2016) 229-244. [15] B. Gorr, M. Azim, H.J. Christ, T. Mueller, D. Schliephake, M. Heilmaier, J. Alloy. Compd. 624 (2015) 270-278. 16

[16] C.M. Liu, H.M. Wang, S.Q. Zhang, H.B. Tang, A.L. Zhang, J. Alloy. Compd. 583 (2014) 162-169. [17] M.H. Chuang, M.H. Tsai, W.R. Wang, S.J. Lin, J.W. Yeh, Acta Mater. 59 (16) (2011) 6308-6317. [18] C.Y. Hsu, T.S. Sheu, J.W. Yeh, S.K. Chen, Wear 268 (5-6) (2010) 653-659. [19] O.N. Senkov, S.V. Senkova, C. Woodward, D.B. Miracle, Acta Mater. 61 (5) (2013) 1545-1557. [20] O.N. Senkov, G.B. Wilks, J.M. Scott, D.B. Miracle, Intermetallics 19 (5) (2011) 698-706. [21] O.N. Senkov, C. Woodward, D.B. Miracle, JOM 66 (10) (2014) 2030-2042. [22] C.C. Juan, M.H. Tsai, C.W. Tsai, C.M. Lin, W.R. Wang, C.C. Yang, S.K. Chen, S.J. Lin, J.W. Yeh, Intermetallics 62 (2015) 76-83. [23] N.N. Guo, L. Wang, L.S. Luo, X.Z. Li, Y.Q. Su, J.J. Guo, H.Z. Fu, Mater. Des. 81 (2015) 87-94. [24] N.D. Stepanov, D.G. Shaysultanov, G.A. Salishchev, M.A. Tikhonovsky, Mater. Lett. 142 (2015) 153-155. [25] M. Todai, T. Nagase, T. Hori, A. Matsugaki, A. Sekita, T. Nakano, Scripta Mater. 129 (2017) 65-68. [26] Y. Zou, S. Maiti, W. Steurer, R. Spolenak, Acta Mater. 65 (2014) 85-97. [27] O.N. Senkov, G.B. Wilks, D.B. Miracle, C.P. Chuang, P.K. Liaw, Intermetallics 18 (9) (2010) 1758-1765. [28] Y. Zou, H. Ma, R. Spolenak, Nat. Commun. 6 (2015) 7748. [29] H.W. Yao, J.W. Qiao, M.C. Gao, J.A. Hawk, S.G. Ma, H.F. Zhou, Y. Zhang, Mater. Sci. Eng. A 674 (2016) 203-211. [30] L. Qi, D.C. Chrzan, Phys. Rev. Lett. 112 (11) (2014) 115503. [31] S. Sheikh, S. Shafeie, Q. Hu, J. Ahlström, C. Persson, J. Veselý, J. Zýka, U. Klement, S. Guo, J. Appl. Phys. 120 (16) (2016) 164902. [32] J.W. Yeh, Ann. Chim-Sci. Mat. 6 (2006) 633-648. [33] X. Yang, Y. Zhang, Mater. Chem. Phys. 132 (2-3) (2012) 233-238. [34] Z.J. Wang, Y.H. Huang, Y. Yang, J.C. Wang, C.T. Liu, Scripta Mater. 94 (2015) 28-31. [35] S. Guo, C. Ng, J. Lu, C.T. Liu, J. Appl. Phys. 109 (10) (2011) 103505. 17

[36] I. Toda-Caraballo, P.E.J. Rivera-Díaz-del-Castillo, Acta Mater. 85 (2015) 14-23. [37] Z.D. Han, N. Chen, S.F. Zhao, L.W. Fan, G.N. Yang, Y. Shao, K.F. Yao, Intermetallics 84 (2017) 153-157. [38] C. T. Liu, J. H. Schneibel, P. J. Maziasz, J. L. Wright, D.S. Easton, Intermetallics 4 (1996) 429-440. [39] C.T. Liu, J.O. Stiegler, Science 226 (4675) (1984) 636-642. [40] J.H. Luan, Z.B. Jiao, G. Chen, C.T. Liu, J. Alloy. Compd. 602 (2014) 235-240. [41] J.H. Luan, Z.B. Jiao, W.H. Liu, Z.P. Lu, W.X. Zhao, C.T. Liu, Mater. Sci. Eng. A 704 (2017) 91-101. [42] P.H. Wu, N. Liu, W. Yang, Z.X. Zhu, Y.P. Lu, X.J. Wang, Mater. Sci. Eng. A 642 (2015) 142-149. [43] Y.D. Wu, Y.H. Cai, X.H. Chen, T. Wang, J.J. Si, L. Wang, Y.D. Wang, X.D. Hui, Mater. Des. 83 (2015) 651-660. [44] C.C. Juan, K.K. Tseng, W.L. Hsu, M.H. Tsai, C.W. Tsai, C.M. Lin, S.K. Chen, S.J. Lin, J.W. Yeh, Mater. Lett. 175 (2016) 284-287.

18