Advanced Multicomponent Alloys: From Fundamentals to Applications (Materials Horizons: From Nature to Nanomaterials) [1st ed. 2022] 9811947422, 9789811947421

Citation preview

Materials Horizons: From Nature to Nanomaterials

Zengbao Jiao Tao Yang   Editors

Advanced Multicomponent Alloys From Fundamentals to Applications

Materials Horizons: From Nature to Nanomaterials Series Editor Vijay Kumar Thakur, School of Aerospace, Transport and Manufacturing, Cranfield University, Cranfield, UK

Materials are an indispensable part of human civilization since the inception of life on earth. With the passage of time, innumerable new materials have been explored as well as developed and the search for new innovative materials continues briskly. Keeping in mind the immense perspectives of various classes of materials, this series aims at providing a comprehensive collection of works across the breadth of materials research at cutting-edge interface of materials science with physics, chemistry, biology and engineering. This series covers a galaxy of materials ranging from natural materials to nanomaterials. Some of the topics include but not limited to: biological materials, biomimetic materials, ceramics, composites, coatings, functional materials, glasses, inorganic materials, inorganic-organic hybrids, metals, membranes, magnetic materials, manufacturing of materials, nanomaterials, organic materials and pigments to name a few. The series provides most timely and comprehensive information on advanced synthesis, processing, characterization, manufacturing and applications in a broad range of interdisciplinary fields in science, engineering and technology. This series accepts both authored and edited works, including textbooks, monographs, reference works, and professional books. The books in this series will provide a deep insight into the state-of-art of Materials Horizons and serve students, academic, government and industrial scientists involved in all aspects of materials research. Review Process The proposal for each volume is reviewed by the following: 1. Responsible (in-house) editor 2. One external subject expert 3. One of the editorial board members. The chapters in each volume are individually reviewed single blind by expert reviewers and the volume editor.

Zengbao Jiao · Tao Yang Editors

Advanced Multicomponent Alloys From Fundamentals to Applications

Editors Zengbao Jiao The Hong Kong Polytechnic University Hong Kong, China

Tao Yang City University of Hong Kong Hong Kong, China

ISSN 2524-5384 ISSN 2524-5392 (electronic) Materials Horizons: From Nature to Nanomaterials ISBN 978-981-19-4742-1 ISBN 978-981-19-4743-8 (eBook) https://doi.org/10.1007/978-981-19-4743-8 © The Editor(s) (if applicable) and The Author(s), under exclusive license to Springer Nature Singapore Pte Ltd. 2022 This work is subject to copyright. All rights are solely and exclusively licensed by the Publisher, whether the whole or part of the material is concerned, specifically the rights of translation, reprinting, reuse of illustrations, recitation, broadcasting, reproduction on microfilms or in any other physical way, and transmission or information storage and retrieval, electronic adaptation, computer software, or by similar or dissimilar methodology now known or hereafter developed. The use of general descriptive names, registered names, trademarks, service marks, etc. in this publication does not imply, even in the absence of a specific statement, that such names are exempt from the relevant protective laws and regulations and therefore free for general use. The publisher, the authors, and the editors are safe to assume that the advice and information in this book are believed to be true and accurate at the date of publication. Neither the publisher nor the authors or the editors give a warranty, expressed or implied, with respect to the material contained herein or for any errors or omissions that may have been made. The publisher remains neutral with regard to jurisdictional claims in published maps and institutional affiliations. This Springer imprint is published by the registered company Springer Nature Singapore Pte Ltd. The registered company address is: 152 Beach Road, #21-01/04 Gateway East, Singapore 189721, Singapore

Contents

Part I

High-Entropy Alloys

1

Body-Centered Cubic High-Entropy Alloys . . . . . . . . . . . . . . . . . . . . . . Yuan Wu, Xiaoyuan Yuan, Xiaocan Wen, and Meiyuan Jiao

3

2

Face-Centered Cubic High-Entropy Alloys . . . . . . . . . . . . . . . . . . . . . . Weihong Liu and Boxuan Cao

35

3

Eutectic High-Entropy Alloys . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Wenna Jiao, Zhijun Wang, Sheng Guo, and Yiping Lu

53

4

Cubic Ordered Intermetallic Alloys . . . . . . . . . . . . . . . . . . . . . . . . . . . . . W. C. Xiao, Y. L. Zhao, and T. Yang

91

Part II

High-Temperature Superalloys

5

Fe-Based Heat-Resistant Steels . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 107 Wei Yan, Shenhu Chen, Ye Liang, Yanfen Li, and Lijian Rong

6

Ni-Base Superalloys: Alloying and Microstructural Control . . . . . . . 133 Shi Qiu and Zengbao Jiao

7

Overview of the Development of L12 γ' -Strengthened Cobalt-Base Superalloys . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 155 Wei-Wei Xu

Part III Advanced High-Strength Steels 8

Advanced High-/Medium-Mn Steels . . . . . . . . . . . . . . . . . . . . . . . . . . . . 179 Xiao Shen and Wenwen Song

9

G-Phase Strengthened Steels . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 225 Wenwen Sun and Xulong An

v

vi

Contents

10 Intermetallic-Precipitation-Strengthened Steels . . . . . . . . . . . . . . . . . . 247 Mengchao Niu, Haojie Kong, Bingchen Zhou, Wei Wang, and Zengbao Jiao Part IV Shape Memory Alloys 11 Abnormal Grain Growth and Single Crystals in Multicomponent Shape-Memory Alloys . . . . . . . . . . . . . . . . . . . . . . . 269 Jixun Zhang, Tao Yang, and Shuiyuan Yang 12 Polycrystalline Shape-Memory Alloy and Strain Glass . . . . . . . . . . . . 287 Aleksandr Shuitcev, Yunxiang Tong, Yu Wang, and Daoyong Cong

Part I

High-Entropy Alloys

Chapter 1

Body-Centered Cubic High-Entropy Alloys Yuan Wu, Xiaoyuan Yuan, Xiaocan Wen, and Meiyuan Jiao

1 Design of BCC Refractory High-Entropy Alloys (RHEAs) 1.1 Design of Single-Phase Solid Solution It is generally accepted that the mole fraction of each constituent element in HEAs is from 5 to 35%, which greatly expands the composition range of HEAs compared to the design concept in conventional alloys. Due to the unique design concept of HEAs, some unintended second phases may form even though the high-entropy effect still exists, which may influence the related performance of HEAs. Thus, how to design single-phase HEAs with good performance is a goal pursued by materials researchers. According to previous reports, it was found that the high mixing entropy is a pivotal factor but not an essential factor for the formation of single-phase solid solution. To improve the accuracy of prediction for designing new single-phase HEAs, the stability of single FCC or BCC HEAs can be predicted based on the atomic-size difference (δ) and valence electron concentration (VEC, Ω) as follows [1, 2]:

Y. Wu (B) · X. Yuan · X. Wen · M. Jiao State Key Laboratory for Advanced Metals and Materials, University of Science and Technology, Beijing, China e-mail: [email protected] © The Author(s), under exclusive license to Springer Nature Singapore Pte Ltd. 2022 Z. Jiao and T. Yang (eds.), Advanced Multicomponent Alloys, Materials Horizons: From Nature to Nanomaterials, https://doi.org/10.1007/978-981-19-4743-8_1

3

4

Y. Wu et al.

⎡ | N | ci (1 − ri /r )2 δ=| i=1

r=

n 

Ω=

ci ri

i=1 n 

(1)

ci Ωi

i=1

where ci , ri , and Ωi are atomic percentage, atomic radius, and VEC of the ith component, respectively. Combined with δ and Ω according to the abundant experimental data [3, 4], the single-phase solid solution can be formed when Ω > 1.1 and δ < 6.6%. FCC phase is stable when VEC > 8 and BCC phase is stable when VEC < 6.87. In spite of the fact that these parameters are based on statistical experimental data, it is significant and instructive for designing novel single-phase HEAs. Moreover, researchers propose the thermodynamic criterion combined with the inherent physical and chemical properties of the constituent elements, such as Gibbs free energy [5], mixing enthalpy and entropy [6], and vacancy exchange potential [7], which can further improve the accuracy of the prediction for designing single-phase HEAs combined with the empirical equation. The development of new high-performance materials is getting faster and faster. Conventional single-phase solid solution cannot meet the demand for high performance under different conditions, especially the harsh environments. Thus, the development of novel high-performance HEAs is not limited to single-phase structures. Because the BCC HEAs mainly consist of refractory elements, we will focus on the performance-oriented composition design of the refractory HEAs (RHEAs) in the next parts.

1.2 Development of RHEAs with Good Wear Resistance Due to the high solid-solution strengthening effect which is the intrinsic property of HEAs, the wear resistance in HEAs is inevitably better than that in conventional alloys, such as Ni-based superalloys and SS304 stainless steel [8–11], as shown in Fig. 1. The increment of the wear resistance can be dominated by the wear film formed on the surface, which is related to the oxides combined with the constituent elements. For example, some easily oxidized elements, such as Ti and Zr, can form oxide film more easily on the surface [9], which can improve the wear resistance resulting from the lubrication effect of the film. Thus, the component design for improving wear resistance should consider some elements which can form oxides [9, 11], or silicides [12] with a continuous and dense structure under corresponding friction conditions. Moreover, the temperature is another key factor to the wear resistance. Different

1 Body-Centered Cubic High-Entropy Alloys

5

Fig. 1 a and b are the comparison of the volume loss as a function of the sliding distance and the wear rates for Mo20 Ta20 W20 Nb20 V20 and Inconel 718, respectively [10]. c and d are the steadystate coefficient of friction and wear rate for HfTaTiVZr, TaTiVWZr, and SS304 as a function of temperature, respectively [11]

films exhibit different characteristics under different temperatures. Al2 O3 film is still continuous and dense at elevated temperatures, while the nitride film becomes incompact [13]. The silicide film has a good wear resistance from ambient to intermediate temperatures, but its wear properties are not clear at high temperatures [12]. Thus, the component design for the wear resistance should consider the application temperature. Yet, there are slight researches on friction and wear properties for RHEAs, especially for their tribological behavior at the elevated temperatures. More investigations are needed for the guidance to develop novel wear-resistance RHEAs.

1.3 Development of RHEAs with Good Oxidation Resistance The service temperature of conventional Ni-based superalloys is higher than 1100 °C, which has approached the ultimate service temperature of the metal material. Such service conditions impose some special peculiarities on the material, i.e., excellent high-temperature stability, strength, and oxidation resistance. Compared to the Ni-based superalloys, generally, the constituent elements in BCC RHEAs are the refractory elements, such as Nb, Mo, and Ta, which make the BCC RHEAs withstand higher temperatures without melting [13]. Besides, the BCC RHEAs have higher high-temperature stability, strength, and softening resistance compared to the conventional superalloys in the previous research [14]. Thus, good oxidation resistance is inevitable and necessary for the BCC RHEAs applying and working in the high-temperature environment.

6

Y. Wu et al.

Fig. 2 Isothermal oxidation curves for the HEAs at 1300 °C, showing the effects of Ti, V, and Si addition [15]

Outstanding oxidation resistance is dominated by the oxide film formed on the metal surface with stable chemical properties, denseness, good combination with matrix, and fewer defects during the oxidation process to hinder the further oxidation of the alloy and protect the matrix. But according to the characteristics of the primary refractory metal elements in the BCC RHEAs, such as Nb, Mo, Ta and W, the oxidation resistance of these refractory elements under high temperature is unsatisfactory. To improve the oxidation resistance of BCC RHEAs, the oxidationresistant film constituted with the component elements must have a good antioxidant capacity. Thus, some alloy elements, such as metallic elements Al and Cr, and nonmetal element Si, are added to the BCC RHEAs [15–19]. For example, Ti, V, and Si were added, and the content of these elements was adjusted in the NbCrMoAl0.5 RHEA [15], as shown in Fig. 2. The results show that the compact Al2 O3 film was generated due to Al in the matrix, and the addition of Ti and Si can also produce compact (TiCrNb)O2 film at 1300 °C. Both the two films were dense, and the oxidation resistance in this RHEA was improved significantly. However, porous VOx film with a low melting point [13] was formed and further deteriorated the oxidation resistance. Therefore, from the viewpoint of alloy design, we should consider not only whether the film is dense but also the melting point of the formed oxide film under high temperatures. Until now, Ti, Cr, Si, and Al are preferred to be added to the BCC RHEAs to improve the oxidation resistance on one hand and lower the density on the other hand.

1.4 Development of RHEAs with Good Corrosion Resistance Similar to the wear and oxidation resistance mechanism in BCC HEAs, the corrosion resistance of alloys is also related to the film formed on the surface in different solutions. The film usually is a passivation layer which can reduce the reaction rate with the surrounding solution and further protect the matrix from corrosion [20].

1 Body-Centered Cubic High-Entropy Alloys

7

From the previous research about the corrosion resistance of BCC RHEAs, the BCC RHEAs have a more excellent corrosion resistance in the high concentration of sulfuric acid, hydrochloric acid, nitric acid, etc. compared to conventional stainless steels, Ti-alloys and Cu-alloys [21, 22], as shown in Fig. 3. This phenomenon is mainly due to the fact that the elements used in conventional alloys to improve the corrosion resistance, such as Ta, Mo, Al, and Cr, have become the primary constituent elements in BCC RHEAs [20]. Besides, due to the unique characteristics of high entropy, the BCC RHEAs have some superiorities compared to the conventional alloys [23]. First, the high-entropy effect improves the forming ability of single phase, which can reduce the possibility of element segregation and thus can promote a uniform film formed on the surface to improve the passivation ability and pitting resistance. Second, the sluggish diffusion effect lowers the diffusion rate of the element from the matrix to surface, which can retard the alloy elements diffusing to the surface when the alloy is corroded. As a result, the corrosion and cauterization of the RHEAs are delayed. Last, the lattice distortion effect may impede the movement of the atoms, which may improve the stress corrosion. Therefore, the BCC RHEAs have enormous potential as corrosion-resistant materials. The BCC HEAs usually consist of elements which can form a passivation film. But different passivation films formed by different elements existing different characteristics in different solutions. For example, Zhang et al. studied the corrosion resistance of the base alloy TiZr0.5 Nb containing Cr, V, and Mo [24]. The result shows that all the alloys exhibit excellent corrosion resistance in the NaCl and H2 SO4 solutions. Among the three alloys, V and Mo decrease the resistance to general corrosion and increase the pitting corrosion resistance for the series of alloys in the NaCl solution slightly, while they greatly improve the corrosion resistance in the H2 SO4 solution. TaNbHfZrTi can generate a Ta2 O5 passivation film, which is very stable in the HNO3 environment [24]. However, in fluorine ions containing HNO3 solution, the corrosion Fig. 3 Comparison of pitting potential and corrosion current density for the Hf0.5 Nb0.5 Ta0.5 Ti1.5 Zr RHEA, previously reported HEAs, and some conventional passive alloys in the 3.5 wt% NaCl solution [22]

8

Y. Wu et al.

rate is much higher compared with the alloy in the HNO3 solution. The reason is that the Zr and Hf can form fluoride with fluorine ions more readily. By the way, the corrosion resistance of Ti1.5 ZrHf0.5 Nb0.5 Ta0.5 RHEA is much better than that of 316L stainless steel in the NaCl solution due to the high pitting potential [22]. Thus, the alloy design for the corrosion resistance should consider the chemical reaction between the constituent elements and the solution environment and lower the content of the disadvantageous element. Now, recent researchers come to notice that not only the chemical component but also the special structure and physical characteristics lead to the excellent corrosion resistance of BCC RHEAs [20], and it is expected to evolve RHEAs with better corrosion resistance than the existing alloys in the future.

1.5 Development of RHEAs with Good Mechanical Properties To obtain better mechanical properties no matter at ambient temperature or high temperatures, the single-phase BCC RHEAs are not limited to investigations from scholars nowadays. The abundant composition variability of BCC RHEAs provides plenty of opportunities to process and generate the desired phase structure to control the property of ultimate alloys. However, that will lead to a predicament, i.e., how to choose an appropriate chemical constituent from a mass of possible constituents. Different from chemical compositions, accurately grasping the commonness and characteristics of elements and determining the relationship between the structure and properties of the alloys can better and more accurately guide the design and optimization of BCC RHEAs, which will greatly reduce the amount of trial and error experiments. The component elements in BCC RHEAs can be classified into BCC stabilizer elements (W, Mo, Nb, Ta, and V), HCP stabilizer elements (Hf, Zr, and Ti), high mixing entropy elements (Cr, Co, and Al), and non-metallic elements (C, N, O, B, Si, etc.) according to their features. The stabilizer elements are conducive to forming a single BCC structure, which can improve the phase stability of the alloy significantly at high temperatures [14]. At ambient temperature, the stabilizer elements also can inhibit the formation of second phases and enhance the mechanical properties. The primary mechanical properties of RHEAs are focused on the compression properties [25–27]. The RHEAs with good tensile properties usually are the series of TiZrHfNbTa HEAs [28–30]. The TRIP effect was introduced into a TiZrHfTa RHEA via regulating the concentration of the BCC stabilizer element Ta [29], resulting in a good combination of high strength and plasticity due to the transformation of the metastable BCC phase during tension. The HCP stabilizer elements have an HCP structure at low temperatures and BCC structure at high temperatures [31, 32]. The inclusion of these elements in the BCC HEAs is beneficial to promoting the equilibrium of strength and toughness of the alloys and improving the plasticity of the alloys at the same time [33]. Besides, taking advantage of the instability of HCP

1 Body-Centered Cubic High-Entropy Alloys

9

elements at the intermediate temperatures in the BCC RHEAs, the size and distribution of the precipitates enriched with HCP elements can be controlled by appropriate heat treatment processes, and the mechanical properties can be improved [34]. The elements with high mixing entropy have negative mixing enthalpy with most refractory elements, resulting in a strong interaction during alloying. High mixing enthalpy can easily destroy the formation of single-phase disordered solid solution during solidification and promote the formation of new intermetallic compounds or lead to the formation of ordered phases. The BCC RHEAs containing Cr will form Laves phase when solidification [35–37], and the second phase has a significant reinforcement effect, showing excellent strength but low plasticity at the ambient temperature. Only a single-phase structure containing Al can improve the strength and plasticity of the alloys [38–40]. Increasing the Al content promotes the formation of ordered phases, such as the B2 phase, which makes the strength and plasticity decrease [41]. By the way, the high mixing entropy elements usually have a lower density than that of the primary constituent elements and thus reduce the density of the BCC HEAs on the one hand, and can improve the high-temperature oxidation resistance on the other hand, which plays an important role for high-temperature applications. The non-metallic elements added to the BCC HEAs can be classified into two types, i.e., Si with a large atom size which can form the silicide with the refractory elements readily, and B, C, N, O with a small atom size which usually occupy the interstitial lattice site. With the increase of Si additions in the NbMoTaW RHEA, the yield strength increases while the plasticity decreases. The appearance and increase of the silicide phase have a strong constraint on the matrix and hinder the motion of the dislocations, thus improving the strength of the alloy [12]. Lei et al. added 2% (atomic fraction) of oxygen to the TiZrHfNb BCC RHEA to generate the ordered oxygen complexes (OOCs) [30], as shown in Fig. 4. The OOCs act on the dislocations and cause the dislocation pinning, making the plastic flow more uniform. The addition of carbon in a TiZrHf0.5 NbMo0.5 RHEA improves the fracture strength and plasticity at ambient temperature, which results from the decrease of the solid-solution strengthening [42]. In general, the non-metallic atoms with a small size lead to the lattice distortion in alloys to hinder the dislocation motion, thus obviously playing a role in strong toughness. The RHEA is considered to be the most promising alloy in the next-generation superalloys. However, the cost of preparing RHEAs is high, and the scholars hope to improve the high-temperature performance via precipitation strengthening, which is similar to the Ni-based superalloys. Laves phase is an important reinforced phase in conventional alloys. The formation of Laves phase is mainly related to Cr, Mo, and Zr, which are beneficial to improve high-temperature properties [35, 37, 43, 44]. Thus, the Laves phase can be easily obtained in RHEAs, which can significantly improve the high-temperature oxidation resistance and mechanical properties. Another hightemperature reinforcement phase is the ordered B2 phase [45, 46]. The B2 phase can exist in the matrix of as-cast alloys or in the alloys after annealing at 1200 °C, and the common elements in these alloys are Al, Ti, Zr, and Nb [47, 48]. Meanwhile, Senkov found that the B2 phase can be formed by the decomposition of high-temperature BCC phase [45]. The coherent precipitation of the disordered plate-like BCC particles

10

Y. Wu et al.

Fig. 4 a Room-temperature tensile stress–strain curves for the as-cast TiZrHfNb, (TiZrHfNb)98 O2 and (TiZrHfNb)98 N2 HEAs. b O composition profile as a function of the distance to the interface for a selection of particles (left axis) and evolution of the composition of the main constituents relative to their respective matrix composition (right axis) [30]

forms on the ordered B2 matrix, which is similar to the coherent precipitation of L12 in the FCC matrix in Ni-based superalloys, resulting in excellent high-temperature oxidation resistance and strength.

1.6 Development of RHEAs with Good Irradiation Resistance The first report about the irradiation resistance in HEA was in 2012 [49]. The researchers developed a novel equimolar ZrHfNb alloy with a single BCC phase according to which the elements are frequently used for nuclear materials. Several irradiation-resistant HEAs have been reported over the past decade, including FCC HEAs, such as FeCoNiCrCu [50] and FeNiCrMn [51], and BCC HEAs, such as Ti2 ZrHfV0.5 Mo0.2 [52] and W38 Ta36 Cr15 V11 [53]. It is found that the HEAs had a more excellent irradiation resistance than the conventional irradiation-resistant alloys, such as stainless steels whether under the electron, helium ion, and neutron irradiation, as shown in Fig. 5. It may make the HEAs the potential radiation-resistant material. In BCC RHEAs, the mechanism of irradiation resistance originates from the high entropy. The high-entropy effect brings a unique structure and physical characterizations in the irradiation-resistant BCC RHEAs. In contrast to conventional alloy systems, the high entropy of mixing effect tends to stabilize the single solid-solution structure, which makes the RHEAs a low activation property under irradiation [54]. The severe lattice distortion in RHEAs makes the overlap region between the migration energy of vacancies and interstitials broad, which is beneficial in delaying the irradiation-induced dislocation loop growth [55, 56]. The sluggish diffusion effect may hinder the migration of the atoms and thus delay the irradiation damage [57].

1 Body-Centered Cubic High-Entropy Alloys

11

Fig. 5 Hardness versus depth for irradiated and un-irradiated samples of a HfTaTiVZr HEA and b 304 stainless steels [54]

Another radiation-resistant mechanism in RHEAs may come from the “selfhealing” process [58, 59]. RHEAs containing several primary elements have different local chemical environment [60], and thus many types of atomic vacancies and interstitials form under irradiation, exhibiting extremely high formation energy and high atomic-level stress as well as low migration energy [59]. These unstable vacancies and interstitials promote amorphization and follow by recrystallization. Due to the fact that Ti, Zr, and Hf are usually added to the RHEAs which is common in amorphous, it is easy for the cyclic process between amorphization and recrystallization, resulting in the high recombination rate for the annihilation of the defects and thus improving the irradiation resistance. Until now, however, the effect of different elements on the irradiation resistance as well as the mechanical properties before and after irradiation in RHEAs is less studied. More investigations need to be conducted, and the development of novel RHEAs with good irradiation resistance is a long way to go.

2 Microstructures of BCC RHEAs 2.1 Single-Phase Solid Solution The crystal structure is one of the basic factors that determine the physical, chemical, and mechanical properties of solid metals. Studies on RHEAs have found that the high mixing entropy enhances the phase stability of the solid solution, which hinders the formation of intermetallic compounds and promotes the formation of simple solid solutions in the alloys, a phenomenon that is more pronounced at high temperatures.

12

Y. Wu et al.

One of the most famous BCC RHEAs is the Senkov alloy [61], which contains Nb (BCC), Mo (BCC), Ta (BCC), and W (BCC) in equal molar ratios, and the as-cast sample has a single-phase BCC crystal structure, showing typical dendritic features. The quaternary alloys were then expanded to quintuple alloys by adding V (BCC) or Ti (HCP), and the as-cast alloys with different lattice parameters exhibited simple BCC solid-solution structures [62]. These two alloys exhibit outstanding softening resistance and thermal stability at ultra-high temperatures, but suffer from roomtemperature brittleness, which severely limits their processability and thus practical applications. In addition, Al acts as a strong BCC stabilizer [63–65] and tends to form strong bonds with the constituent elements within specific RHEA systems, such as TiZrNbHfTa. Thus, Al additions to TiZrNbHfTa are expected to improve the oxidation resistance, high temperature strength, and maintain a BCC solid-solution structure [63]. Due to the inherent properties of BCC RHEAs and the large atomic-radius difference between the interstitial atoms and basic elements of HEAs, interstitial C, N, or O atoms can also play a huge role in the crystal structure of alloys. In particular, there are some interesting research results. Recently, Wang et al. [66] successfully achieved a large plasticity of >10%, along with a high strength of >1750 MPa in the NbMoTaW RHEAs via grain boundary engineering with the addition of either metalloid B or C. Microalloying of B or C atoms which preferentially segregate to grain boundaries suppresses the O segregation at grain boundaries. A strong electronic interaction between B/C and the host metals ensures a strong bonding to adjacent metallic atoms, which contributes to the enhanced cohesion. Meanwhile, the enhanced grain boundary cohesion successfully suppressed the early intergranular crack and changed the deformation mode from intergranular to intragranular, eventually contributing to the plasticity enhancement. Moreover, the boron doping induced a high efficiency in grain refinement from 96.0 to 16.2 μm of the TiZrNb MEA, which is the main factor for strengthening. Dislocation-dominated deformation mechanism involving cross-slip and dislocation pining in the TiZrNb containing 500 ppm boron serves to enhance the strain-hardening capacity, leading to the enhancement of ductility [67]. It has been empirically established that the competition between the configuration entropy and enthalpy, the difference between the atomic radius and electronegativity of constituent elements [31, 41], and also the overall valence electron concentration [2] are a few key thermodynamic parameters for the formation of HEAs. All these parameters are very susceptible to pressure tuning. Actually, pressure is a very powerful tool to tune the atomic/electronic structure of various materials and has been employed to understand materials and search for novel materials through rich pressure-induced phase transitions in diverse systems. Therefore, in contrast to the seeming ultra-stability during heating or cooling, BCC HEAs might exhibit rich tunable behavior under high pressures. Ahmad et al. investigated the structural stability of a TiZrHfNb alloy with a disordered BCC structure during compression up to 50.8 GPa, and no phase transition was found [68]. Yusenko et al. explored another BCC-structured Al2 CoCrFeNi HEA which transitioned up to 60 GPa, and no phase was found as well [69]. Moreover, the structural evolution of a (TaNb)0.67 (HfZrTi)0.33 HEA during compression up to ~100 GPa seemed very

1 Body-Centered Cubic High-Entropy Alloys

13

robust without any detectable structural transition [70]. No phase transition was ever observed in the BCC-structured HEAs. Therefore, it seems that the BCC-structured HEAs are incredibly stable, and much higher pressure may be needed to induce phase transition. To lower the transition pressure of possible polymorphic phase transitions, Cheng et al. [71] chose an equiatomic AlCoCrFeNi HEA and monitored its structural evolution during compression up to 42 GPa. The AlCoCrFeNi alloy had an ordered BCC structure (B2 phase) and was reported to sit in the transition zone between the FCC and BCC phases, as x varies in the Alx CoCrFeNi HEA system (0 < x < 2) [72]. Indeed, they discovered a phase transition from the initial B2 phase to a highly distorted form starting at relatively low pressure of ~17.6 GPa. Besides the XRD peak splitting, severe peak weakening and broadening occurred during compression, which may have been caused by the significant lattice distortion developed in the sample. Therefore, their work was unable to resolve the atomic structure of the high-pressure phase. Nevertheless, it is the first time that a pressure-induced polymorphism was suggested in a BCC-structured HEA [71].

2.2 Short-Range Orderings (SROs) Recently, a large number of studies reveal that the arrangement of atoms in HEAs is not an ideal disordered state due to the diversity in atomic radii and the complex interactions between the constituent elements usually result in SRO structures during solidification or/and heat treatment processes [73–77]. Furthermore, the SROs observed in HEAs affect the mechanical and physical properties of alloys. It seems that the existence of topological and/or chemical SROs is one of the common yet key features of HEAs, and the manipulation of SROs may provide an effective way for tuning the properties of HEAs. In fact, SROs have been widely reported in different types of BCC RHEAs. For instance, Bu et al. [77] studied the local chemical fluctuation in a HfNbTiZr HEA and employed in situ TEM straining to investigate the deformation behaviors of the ductile BCC HfNbTiZr HEA. It was demonstrated that local chemical fluctuations (LCFs) induced dislocation pinning (Fig. 6), multiplications and cross-slips in the HfNbTiZr HEA. The pinning caused by the LCFs not only increases the dislocation motion barrier, but also stimulates dislocation multiplication. The local double cross-slips make the dislocations multiplication occur onto various planes, which effectively alleviates strain localization. Furthermore, Lei et al. [30] reported that the addition of 2% oxygen in the TiZrHfNb HEA promotes the formation of the OOCs since oxygen prefers to stay with the Ti/Zr atoms, as clearly shown in Fig. 7. The unique SROs of interstitial solutions play a decisive role in the mechanical properties. The SROs in (TiZrHfNb)98 O2 HEA act as precipitates to interact with dislocations by pinning and facilitate frequent dislocation cross-slip, leading to the change of the plastic deformation mode from planar slip to wavy slip.

14

Y. Wu et al.

Fig. 6 LCF-induced pinning. a A low-magnification bright field (BF)-scanning transmission electron microscopy (STEM) image of a pinned dislocation. b The high-resolution HAADF-STEM image corresponding to the square area in (a). c Intensity line profiles of the blue squared region in (b), inset is the corresponding enlarged HAADF-STEM image. d The geometric phase analysis (GPA) of (b) shows the strain fluctuation around the pinning point [78]

2.3 Intermediate and Complex Phases It was found that there are many factors affecting the formation of RHEAs [79], including mixing entropy, mixing enthalpy, atomic-size difference, and valence electron concentration. When the enthalpy of formation of intermetallic compounds is large enough to overcome the effect brought about by the high-entropy effect, intermediate or complex phases are generated. The HCP phase only appears as a second phase in some RHEAs containing Hf, Ti, and Zr. As a strong BCC stabilizer, Ta and Nb play an important role in phase formation in these HEAs [29, 80]. Huang et al. [29] used the TaHfZrTi high-entropy alloy as a model material and successfully designed the TaxHfZrTi two-phase RHEA with excellent mechanical properties by regulating the thermodynamic and mechanical stability of the phases. As shown in Fig. 8 [29], the transformation-induced plasticity (TRIP) effect induces the formation of the HCP phase, which exists as a second phase. Phase transformation yields an intensive strain-hardening effect by dynamic

1 Body-Centered Cubic High-Entropy Alloys

15

Fig. 7 Chemical characterizations of ordered oxygen complexes (OOCs) in the (TiZrHfNb)98 O2 HEA. a The HAADF-STEM image for the [011] BCC crystal axis with a differently adjusted contrast to reveal the existence of chemical short-range ordering. b The atomic number contrast analysis of the HAADF-STEM image reveals the OOCs. Red squares represent the Zr/Ti-rich regions, and yellow squares indicate the Hf/Nb-rich regions. c The aberration-corrected STEM-annular bright field image. The inset in (c) is an enlarged view of the OOCs, with the white arrows indicating the positions of the oxygen atom columns. d APT reconstruction of the O-doped HEA [30]

strain–stress partitioning between the BCC and HCP phases and promotes plastic deformation inside grains, which effectively suppresses early cracking and eventually gives rise to an outstanding combination of strength and ductility. Similarly, with the increase of Nb content in HfNbx Ta0.2 TiZr (x = 0, 0.15, 0.20 and 0.25) HEAs, phase constitution exhibits a transition from a nearly single HCP phase (i.e., Nb0) to dual phases of BCC and HCP, i.e., Nb0.15 and Nb0.2 [80]. It was found that proper addition of Nb could not only enhance the TRIP effect due to the reduced phase stability of the prior BCC phase, but also facilitate twinning in the product HCP (hexagonal close packing) phase at the late stage of deformation [80]. Moreover, Stepanov et al. [36] investigated the organization and hardness of HfNbTaTiZr alloy after annealing treatment. HfNbTaTiZr alloy precipitated nanoscale HCP phase in BCC matrix after annealing at 600 and 800 °C. The nanoscale HCP phase particles are distributed in the grain boundaries, which substantially increase the hardness of the alloy. Some of the RHEAs have a two-phase microstructure containing either two BCC phases, BCC and B2, or BCC and Laves phases [81]. Reports on the formation of B2 structure of RHEAs which contain Al, Ti, Nb, and Zr. The B2 phase can be present as a matrix phase in as-cast or 1473 K annealed HEAs [45, 46, 82]. Interestingly, most of the Al-containing RHEAs have a superalloy-like two-phase microstructure containing BCC and B2 phases [47, 83–86]. And the disordered BCC phase is usually present in the form of cuboidal nanoscale precipitates arranged along directions in a coherent B2 matrix. Jensen et al. [87] studied the AlMo0.5 NbTa0.5 TiZr RHEA that the arrangement of BCC cuboidal nanoprecipitates is along directions with thin continuous B2 channels around them. This alloy shows a compressive yield strength of ~750 MPa at 1000 °C, which exceeds the capability of any superalloys. While the ductility is very low at room temperature, which is attributed to the continuous ordered B2 matrix of this alloy [83]. Wang et al. [41] found that Al addition in the HEAs results in a notable strengthening effect, which is attributed

16

Y. Wu et al.

Fig. 8 XRD patterns and EBSD images of the as-cast Tax HEAs. The Ta concentration significantly influences the phase constitution of this alloy system, rendering either single (BCC) or dual-phase (BCC + HCP) structure. The decreasing of Ta content destabilizes the BCC matrix and promotes formation of HCP [29]

to solid-solution hardening and ordered phase hardening. When the amount of Al exceeds the solubility in a random solid solution, an ordered BCC phase with the lattice constant close to random one starts to form. The high electron density and Fermi level of Al contribute to the formation of covalent bonds. VEC and Ω-δ criteria seem to more precisely predict the formation of solid solutions. A low VEC value may be in favor of ordering structural transition of BCC phase [41]. Laves phase, which is related to Cr, Mo, and Zr, is an important strengthening phase in conventional alloys, leading to excellent high-temperature properties. It was found that introducing the Laves phase into refractory high-entropy alloys can significantly improve the oxidation resistance [43, 44], creep resistance, and mechanical properties [35–37], while inevitably reducing the room-temperature plasticity. Considering the low density of Cr and the high solubility with refractory elements [88], it can be used to reduce the density of HEAs and the Laves phase is present in the alloy in the form of large particles or fine precipitates. It was found that Al can be used as an alternative to Cr to prepare RHEAs resistant to high-temperature oxidation [18], which can not only reduce the density of the alloy, but also improve the strength.

3 Mechanical Properties of BCC HEAs In recent years, the emergence of HEAs, which is consisting of multiple principal elements in equimolar or near-equimolar ratios, has attracted extensive attention due to their unique phase constitution, structure characteristic, and outstanding mechanical behavior, such as high hardness, strength, and fracture toughness [89, 90].

1 Body-Centered Cubic High-Entropy Alloys

17

Overall, it has been reported that BCC-structured HEAs show higher strength and lower plasticity compared with FCC-structured HEAs, which indicates that the structure types are the dominant factor in controlling the strength or hardness of HEAs. In this section, we review some of the BCC HEAs concerning mechanical properties.

3.1 Mechanical Behavior at Room Temperature The room-temperature mechanical properties of BCC HEAs can be varied over a wide range depending on the phase constitution, heating treatment processing, and loading method (i.e., tensile or compression). The BCC-structured HEAs typically exhibit high strengths although the plasticity is poor. Therefore, the present reported mechanical property is mostly focused on compression tests. For instance, Senkov et al. [14] conducted the room-temperature compressive test of the NbMoTaW and VNbMoTaW HEAs (Fig. 9), where the yield strengths were 1058 and 1246 MPa, respectively, while the elongation was 1.5% and 1.7%, respectively. Limited compressive plasticity at room temperature suggests that the ductile-to-brittle transition for these alloys occurs above room temperature. After compression deformation, scanning electron microscopy imaging revealed that the fracture morphologies of the alloys contained a brittle quasi-cleavage fracture, suggesting that in this alloy, the primary failure mode is tensile rather than shear. The TaNbHfZrTi RHEA has been observed to exhibit relatively high strength and good plasticity [91]. For instance, it was found that this alloy exhibited a yield strength approaching 929 MPa, and a compression plasticity surpassing 50% at room temperature (Fig. 10). The excellent comprehensive mechanical properties were attributed to the synergy of solid-solution strengthening effect, dislocation slip, and twinning. Wang et al. [66] recently reported that doping small-sized metalloids boron or carbon in NbMoTaW RHEAs could preferentially replace oxygen at grain boundaries and promote stronger electronic interaction with the host metals, thus effectively

(a)

(b)

Fig. 9 Compressive engineering stress–strain curve of the NbMoTaW and VNbMoTaW HEA at room temperature [14]

18

Y. Wu et al.

Fig. 10 Engineering stress versus engineering strain compression curves of the Ta20 Nb20 Hf20 Zr20 Ti20 alloy and Ta20 Nb20 W20 Mo20 V20 and Ta25 Nb25 W25 Mo25 alloys [91]

alleviating the grain boundary brittleness and changing the fracture morphology from intergranular fracture to transgranular fracture. Consequently, both large plasticities of >10%, along with high strength of >1750 MPa at room temperature, are successfully achieved (Fig. 11). In addition, some BCC HEAs also display certain tensile plasticity behavior, such as the TaNbHfZrTi HEA and its derivative [28]. J. Cízek et al. [92] reported that the HfNbTaTiZr exhibits a tensile yield strength of 1000 MPa, an ultimate strength of 1010 MPa and a total elongation to failure of 13% (Fig. 12). In addition, heat treatment processing also has an influence on the mechanical property of alloys. Senkov et al. [93] examined the tensile properties of the TaNbHfZrTi HEA after being cold-rolled and then annealed at 1073 and 1273 K, respectively. As shown in Fig. 13, the 86.4% cold-rolled sheet had true tensile stress of 1295 MPa Fig. 11 Stress–strain curves of the as-cast base RHEA (black) and doped RHEAs with different contents of B at room temperature under compression [66]

1 Body-Centered Cubic High-Entropy Alloys

19

Fig. 12 Stress–strain curve for tensile tests of standard size HfNbTaTiZr alloy in the as-cast state [93]

and tensile ductility of 4.7%. After annealing at 1000 °C, the true tensile stress and ductility of the sheet were 1262 MPa and 9.7%, respectively. Wu et al. [94] reported a novel refractory HfNbTiZr HEA with a single BCC structure, which still holds the single BCC phase even after furnace cooling from homogenization at 1573 K for 6 h. The fracture strength and plastic strain of this HEA were about 969 MPa and 14.9%, respectively (Fig. 14). The significant work hardening behavior is considered as resulting from the movement and multiplication of dislocations. Then, Bu et al. [77] further confirmed that the observed local double cross-slips caused by the LCFs distribute dislocations onto various atomic planes homogenously, which is thought to be beneficial for ductility in HfNbTiZr HEAs. Lei et al. [30] further reported that proper additions of oxygen e to NbHfZrTi could promote the formation of ordered nanoscale regions within the HEA characterized by (O, Zr, Ti)-rich atomic complexes. The ordered interstitial complexes change the Fig. 13 True stress–true strain curves of as-cold-rolled and rolled-and-annealed sheet of HfNbTaTiZr [8]

20

Y. Wu et al.

Fig. 14 Tensile true stress–strain curve of HfNbTiZr alloy [95]

dislocation shear mode from planar slip to wavy slip and promote double cross-slip and thus dislocation multiplication through the formation of Frank–Read sources during deformation. Therefore, the yield strength increases from 0.75 GPa for the base HEA to 1.11 GPa for the doped O-2 HEAs, while the elongation has nearly doubled, increasing from 14.21% for the base HEA to 27.66% for the O-2 HEA, accompanied by a substantial work-hardening effect (Figs. 14 and 15). Though BCC HEAs have recently attracted extensive attention for their excellent high-temperature mechanical property, the brittleness and low work-hardening rate at ambient temperature limit their practical uses. Serious efforts have been devoted, and a few solutions have been proposed [79, 95, 95]. Huang et al. [29] reported that via tailoring the stability of the constituent phases, transformation-induced ductility and work-hardening capability are successfully achieved in a brittle TaHfZrTi alloy. Fig. 15 Room-temperature tensile stress–strain curves for the as-cast TiZrHfNb (denoted as base alloy), (TiZrHfNb)98 O2 (denoted as O-2) and (TiZrHfNb)98 N2 (denoted as N-2) HEAs. The inset shows the corresponding strain-hardening response [30]

a

1 Body-Centered Cubic High-Entropy Alloys

21

Phase transformation yielded an intensive strain-hardening effect by dynamic strain– stress partitioning between the BCC and HCP phases and promote plastic deformation inside grains, which effectively suppressed early cracking and eventually gave rise to an outstanding combination of strength and ductility. As shown in Fig. 16, the plasticity of BCC HEAs was greatly improved from 4% (i.e., Ta1) to 27% (i.e., Ta0.4), although the fracture strength is slightly lowered to 1100 MPa with decreasing of Ta content. Similarly, Wen et al. [80] further found that a proper addition of Nb to Ta0.2HfZrTi could not only enhance the TRIP effect due to the reduced phase stability of the prior BCC phase, but also facilitate twinning in the product HCP (hexagonal close packing) Fig. 16 Mechanical behavior of the as-cast Tax HEAs at room temperature. a Representative tensile true stress–strain curves. b The corresponding strain-hardening rate curves [29]

22

Y. Wu et al.

Fig. 17 a Room-temperature tensile mechanical properties of the Nbx alloys, b the corresponding strain-hardening rate curves, and c curves of true stress/work-hardening rate against true strain [80]

phase at the late stage of deformation. As a result, the HfNb0.2 Ta0.2 TiZr HEA shows much enhanced mechanical properties, i.e., pronounced work-hardening behavior and large uniform ductility of up to 26% (Fig. 17).

3.2 Mechanical Behavior at Elevated Temperatures BCC HEAs based on refractory elements have shown interesting properties at elevated temperatures, such as outstanding high-temperature strength, high oxidation, and corrosion resistance, which ensure them promising materials for hightemperature applications [44, 96]. Senkov et al. [14] found that at 600 °C and above, NbMoTaW and VNbMoTaW alloys showed extensive compressive plastic strain accompanied by the yield stress of both alloys dropped by 30–40% between room temperature and 600 °C and was relatively insensitive to temperature above 600 °C (Fig. 18). The high-temperature mechanical properties of these BCC HEAs compare favorably with conventional superalloys (Fig. 19). Senkov et al. [96] researched the compression properties of a refractory TaNbHfZrTi alloy in the temperature range of 296–1473 K (Fig. 20) and found that the properties were correlated with the microstructure developed during compression

1 Body-Centered Cubic High-Entropy Alloys

(a)

23

(b)

Fig. 18 Compressive engineering stress–strain curve of the NbMoTaW and VNbMoTaW HEA at elevated temperatures [14]

Fig. 19 Temperature dependence of the yield V20 Nb20 Mo20 Ta20 W20 HEAs and two superalloys [14]

stress

of

Nb25 Mo25 Ta25 W25

and

testing. They reported the alloy had a yield strength of 929 MPa at 296 K, 790 MPa at 673 K, 675 MPa at 873 K, 535 MPa at 1073 K, 295 MPa at 1273 K and 92 MPa at 1473 K with a strain rate of 10−3 s−1 [96]. Continuous strain hardening and good ductility were observed in the temperature range from 296 to 873 K. Partial dynamic recrystallization, leading to the formation of fine equiaxed grains near grain boundaries, was observed in the specimens deformed at 1073, and 1273 K and completed dynamic recrystallization was observed at 1473 K. Guo et al. [97] reported that at elevated temperatures, the MoNbHfZrTi alloy has a compression yield strength of 825 MPa at 1073 K, 728 MPa at 1173 K, 635 MPa

24

Y. Wu et al.

Fig. 20 Engineering stress versus engineering strain compression curves of the TaNbHfZrTi alloy at different temperatures [96]

at 1273 K, 397 MPa at 1373 K, and 187 MPa at 1473 K, and some fine grains form at grain boundaries due to partial dynamic recrystallization (Fig. 21). Feng et al. [98] used intrinsic material characteristics as the alloy-design principles to design CrMoNbV RHEA, which has high-temperature strengths (beyond 1000 MPa at 1273 K), superior to other reported RHEAs as well as conventional superalloys, as shown in Fig. 22. And the outstanding elevated-temperature mechanical properties were attributed to its large atomic-size and elastic-modulus mismatches, the insensitive temperature dependence of elastic constants, and the dominance of non-screw character dislocations caused by the strong solute pinning, thus leading to pronounced solid-solution strengthening effect.

Fig. 21 Compressive curves of MoNbHfZrTi alloy: a at room temperature and b at different elevated temperatures [97]

1 Body-Centered Cubic High-Entropy Alloys Fig. 22 Stress–strain curves of the CrMoNbV HEA at room temperature (RT) and elevated temperatures (573–1273 K) [98]

25

a

3.3 Mechanical Behavior at Cryogenic Temperatures Recently, the mechanical behavior of BCC HEAs at cryogenic temperatures is drawing increasing attention to extending their applications to cryogenic temperatures in the extreme service conditions. Wang et al. [99] reported the mechanical instability and tensile properties of a TiZrHfNbTa HEA at cryogenic temperatures. As the temperature is decreased from 277 to 77 K, the yield strength, ultimate tensile strength, and uniform elongation increase from 875 to 1547 MPa, 994 to 1762 MPa, and 14.1% to 15.2%, respectively, suggesting that the TaNbHfZrTi RHEA has great potential to be used in lowtemperature extreme conditions (Fig. 23). The synergistic effect of screw dislocation glide, {112} mechanical twinning and BCC → ω phase transformation guaranteed the alloy an excellent combination of strength and ductility at the cryogenic temperature. In addition, Qiao et al. [100] studied the microstructural characteristics and mechanical behaviors of BCC AlCoCrFeNi HEAs. They found that the yielding strengths and fracture strengths increase distinguishingly from 1450 to 1880 MPa and 2960 to 3550 MPa, respectively with temperature decreasing from 298 to 77 K (Fig. 24). However, the plasticity changes very gently with the value of 15.5% at 298 K and 14.3% at 77 K, while the fracture modes are intergranular and transgranular, respectively.

3.4 Hardness The Vickers hardness of HEAs is strongly dependent on the crystal structure, and the hardness of the BCC structure is generally higher than that of other structures. For

26

Y. Wu et al.

Fig. 23 Microstructure and tensile properties of the equiatomic TiZrHfNbTa alloy, a engineering stress–strain curves at 277 and 77 K; b true stress–strain curves at 277 and 77 K [99]

Fig. 24 Compressive true stress–strain curves of the AlCoCrFeNi HEA at a 298 and b 77 K [100]

example, He et al. [101] found that the hardness values of the (FeCoNiCrMn)100−x Alx HEA increase significantly with the increase of the Al content, where the structure of HEAs turns from fully FCC-structured to FCC + BCC phases and finally fully BCC structure. And they surmised that the BCC-structured B2 phase contributes significantly to the hardness. Senkov et al. [102] found that the Vickers microhardness results of the single BCC-phased TaNbWMo and TaNbWMoV HEAs were approximately 4455 and 5250 MPa, respectively. The hardness values of these two HEAs are greater than those of their constituents, which indicates that there is an apparent solid-solution strengthening effect in HEAs. Later, Senkov et al. [47, 91] further showed that the hardness value of the TaNbHfZrTi increases from 3.8 to 4.9 GPa by replacing Hf with Al in Al0.4 Hf0.6 NbTaTiZr HEA without changing its phase composition. While the addition of Al was also conductive to reduce alloy density, increase high-temperature strength and specific strength, and improve oxidation resistance.

1 Body-Centered Cubic High-Entropy Alloys

27

3.5 Wear Behavior HEAs show more competitive and potential for use in tools, molds, and structural components, so wear is a fundamental phenomenon in these applications. Poulia et al. [10, 103] investigated the wear behavior of Mo20 Ta20 W20 Nb20 V20 HEA and a commercial Ni superalloy, Inconel 718. The results show that in all cases of sliding distances and counter bodies, the HEA presented an improved wear response in comparison with Inconel 718 (Fig. 25). Ye et al. [104] investigated the friction and wear of a BCC equiatomic TiZrHfNb HEA using a nanoscratch method. The results revealed that the wear rate of the TiZrHfNb HEA is proportional to the applied load, and the wear-resistance scales linearly with the harness of the alloy. TiZrHfNb exhibited an improved wear resistance and lower coefficient of friction as compared to the pure Nb and C103 (Nb5.4Hf-2Ti), (Fig. 26), which is attributed to the higher hardness/strength and small plowing and adhesion friction coefficients, respectively. This suggests that the current HEA may have the potential for tribological applications.

4 Summary In summary, RHEA is a unique idea and relatively new field, which shows exceptionally bright mechanical properties, especially at elevated temperatures. Here, an overview of mechanical behavior was provided on the BCC HEAs, including hardness, compressive behaviors, tensile behaviors, and wear behaviors. It is worth noting that though the mechanical behavior of BCC HEAs was highlighted and even outperform conventional alloys in some aspects, continued efforts are needed to further improve the comprehensive mechanical properties of BCC HEAs to extend their application prospects in future industrial endeavors.

28

Y. Wu et al.

Fig. 25 Comparative diagrams of the volume loss (left) and the wear rate (right) of Mo20 Ta20 W20 Nb20 V20 versus Inconel 718, tested with both an alumina and a steel ball for sliding distances of a 400 m, b 1000 m, and c 2000 m, respectively [10]

1 Body-Centered Cubic High-Entropy Alloys Fig. 26 a Volume removed (or wear volume) against scratch distance under various normal loads in TiZrHfNb HE alloy. b Wear rate versus normal force and c wear-resistance coefficient versus hardness for pure Nb, C103, and TiZrHfNb HE alloys [104]

29

30

Y. Wu et al.

References 1. Zhang Y, Zhou YJ, Lin JP et al (2008) Solid-solution phase formation rules for multicomponent alloys. Adv Eng Mater 10(6):534–538 2. Guo S, Ng C, Lu J et al (2011) Effect of valence electron concentration on stability of fcc or bcc phase in high entropy alloys. J Appl Phys 109(10):103505 3. Yang X, Zhang Y (2012) Prediction of high-entropy stabilized solid-solution in multicomponent alloys. Mater Chem Phys 132(2–3):233–238 4. Chen Y, Li Y, Cheng X et al (2018) The microstructure and mechanical properties of refractory high-entropy alloys with high plasticity. Materials 11(2):208 5. Melnick AB, Soolshenko VK (2017) Thermodynamic design of high-entropy refractory alloys. J Alloy Compd 694:223–227 6. Senkov ON, Miracle DB (2016) A new thermodynamic parameter to predict formation of solid solution or intermetallic phases in high entropy alloys. J Alloy Compd 658:603–607 7. McAlpine SW, Logan JV, Short MP (2021) Predicting single phase stability and segregation in the NbMoTaTi-(W, V) high entropy alloy system with the vacancy exchange potential. Scripta Mater 191:29–33 8. Mathiou C, Poulia A, Georgatis E et al (2018) Microstructural features and dry-Sliding wear response of MoTaNbZrTi high entropy alloy. Mater Chem Phys 210:126–135 9. Poulia A, Georgatis E, Karantzalis A (2019) Evaluation of the microstructural aspects, mechanical properties and dry sliding wear response of MoTaNbVTi refractory high entropy alloy. Met Mater Int 25(6):1529–1540 10. Poulia A, Georgatis E, Lekatou A et al (2016) Microstructure and wear behavior of a refractory high entropy alloy. Int J Refract Metal Hard Mater 57:50–63 11. Pole M, Sadeghilaridjani M, Shittu J et al (2020) High temperature wear behavior of refractory high entropy alloys based on 4-5-6 elemental palette. J Alloy Compd 843:156004 12. Guo Z, Zhang A, Han J et al (2019) Effect of Si additions on microstructure and mechanical properties of refractory NbTaWMo high-entropy alloys. J Mater Sci 54(7):5844–5851 13. Speight JG (2005) Lange’s handbook of chemistry. McGraw-hill, New York 14. Senkov ON, Wilks GB, Scott JM et al (2011) Mechanical properties of Nb25 Mo25 Ta25 W25 and V20 Nb20 Mo20 Ta20 W20 refractory high entropy alloys. Intermetallics 19(5):698–706 15. Liu CM, Wang HM, Zhang SQ et al (2014) Microstructure and oxidation behavior of new refractory high entropy alloys. J Alloy Compd 583:162–169 16. Gorr B, Mueller F, Christ HJ et al (2016) High temperature oxidation behavior of an equimolar refractory metal-based alloy 20Nb20Mo20Cr20Ti20Al with and without Si addition. J Alloy Compd 688:468–477 17. Müller F, Gorr B, Christ HJ et al (2018) Effect of microalloying with silicon on high temperature oxidation resistance of novel refractory high-entropy alloy Ta–Mo–Cr–Ti–Al. Mater High Temp 35(1–3):168–176 18. Chang CH, Titus MS, Yeh JW (2018) Oxidation behavior between 700 and 1300 °C of refractory TiZrNbHfTa high-entropy alloys containing aluminum. Adv Eng Mater 20(6):1700948 19. Lo KC, Yeh AC, Murakami H (2018) Microstructural investigation of oxidized complex refractory high entropy alloys. Mater Trans MJ201611 20. Fu Y, Li J, Luo H et al (2021) Recent advances on environmental corrosion behavior and mechanism of high-entropy alloys. J Mater Sci Technol 80:217–233 21. Li M, Chen Q, Cui X et al (2021) Evaluation of corrosion resistance of the single-phase light refractory high entropy alloy TiCrVNb0.5 Al0.5 in chloride environment. J Alloys Compd 857:158278 22. Zhou Q, Sheikh S, Ou P et al (2019) Corrosion behavior of Hf0.5 Nb0.5 Ta0.5 Ti1.5 Zr refractory high-entropy in aqueous chloride solutions. Electrochem Commun 98:63–68 23. Hua XJ, Hu P, Xing HR et al (2022) Development and property tuning of refractory highentropy alloys: a review. Acta Metallurgica Sinica (English Letters) 1–35

1 Body-Centered Cubic High-Entropy Alloys

31

24. Li J, Yang X, Zhu R et al (2014) Corrosion and serration behaviors of TiZr0.5 NbCr0.5 Vx Moy high entropy alloys in aqueous environments. Metals 4(4):597–608 25. Stepanov ND, Shaysultanov DG, Salishchev GA et al (2015) Structure and mechanical properties of a light-weight AlNbTiV high entropy alloy. Mater Lett 142:153–155 26. Zhang Y, Liu Y, Li Y et al (2016) Microstructure and mechanical properties of a refractory HfNbTiVSi0.5 high-entropy alloy composite. Mater Lett 174:82–85 27. Tseng KK, Juan CC, Tso S et al (2018) Effects of Mo, Nb, Ta, Ti, and Zr on mechanical properties of equiatomic Hf–Mo–Nb–Ta–Ti–Zr alloys. Entropy 21(1):15 28. Zýka J, Málek J, Veselý J et al (2019) Microstructure and room temperature mechanical properties of different 3 and 4 element medium entropy alloys from HfNbTaTiZr system. Entropy 21(2):114 29. Huang H, Wu Y, He J et al (2017) Phase-transformation ductilization of brittle high-entropy alloys via metastability engineering. Adv Mater 29(30):1701678 30. Lei Z, Liu X, Wu Y et al (2018) Enhanced strength and ductility in a high-entropy alloy via ordered oxygen complexes. Nature 563(7732):546–550 31. Chen SY, Tong Y, Tseng KK et al (2019) Phase transformations of HfNbTaTiZr high-entropy alloy at intermediate temperatures. Scripta Mater 158:50–56 32. Tu CH, Wu SK, Lin C (2020) A study on severely cold-rolled and intermediate temperature aged HfNbTiZr refractory high-entropy alloy. Intermetallics 126:106935 33. Zherebtsov S, Yurchenko N, Panina E et al (2020) Microband-induced plasticity in a Ti-rich high-entropy alloy. J Alloy Compd 842:155868 34. Ma Y, Wu S, Jia Y et al (2019) Hexagonal closed-packed precipitation enhancement in a NbTiHfZr refractory high-entropy alloy. Metals 9(5):485 35. Senkov ON, Senkova SV, Miracle DB et al (2013) Mechanical properties of low-density, refractory multi-principal element alloys of the Cr–Nb–Ti–V–Zr system. Mater Sci Eng, A 565:51–62 36. Senkov ON, Senkova SV, Woodward C et al (2013) Low-density, refractory multi-principal element alloys of the Cr–Nb–Ti–V–Zr system: Microstructure and phase analysis. Acta Mater 61(5):1545–1557 37. Long Y, Liang X, Su K et al (2019) A fine-grained NbMoTaWVCr refractory high-entropy alloy with ultra-high strength: microstructural evolution and mechanical properties. J Alloy Compd 780:607–617 38. Lee K, Jung Y, Han J et al (2021) Development of precipitation-strengthened Al0.8NbTiVM (M=Co,Ni) light-weight refractory high-entropy alloys. Materials 14(8):2085 39. Liu XW, Bai ZC, Ding XF et al (2021) A novel light-weight refractory high-entropy alloy with high specific strength and intrinsic deformability. Mater Lett 287:129255 40. Pang J, Zhang H, Zhang L et al (2021) Ductile Ti1.5 ZrNbAl0.3 refractory high entropy alloy with high specific strength. Mater Lett 290:129428 41. Wang W, Zhang Z, Niu J et al (2018) Effect of Al addition on structural evolution and mechanical properties of the AlxHfNbTiZr high-entropy alloys. Mater Today Commun 16:242–249 42. Guo NN, Wang L, Luo LS et al (2016) Effect of composing element on microstructure and mechanical properties in Mo–Nb–Hf–Zr–Ti multi-principle component alloys. Intermetallics 69:13–20 43. Senkov ON, Senkova SV, Dimiduk DM et al (2012) Oxidation behavior of a refractory NbCrMo0.5 Ta0.5 TiZr alloy. J Mater Sci 47(18):6522–6534 44. Butler TM, Chaput KJ, Dietrich JR et al (2017) High temperature oxidation behaviors of equimolar NbTiZrV and NbTiZrCr refractory complex concentrated alloys (RCCAs). J Alloy Compd 729:1004–1019 45. Senkov ON, Jensen JK, Pilchak AL et al (2018) Compositional variation effects on the microstructure and properties of a refractory high-entropy superalloy AlMo0.5 NbTa0.5 TiZr. Mater & Design 139:498–511 46. Yurchenko NY, Stepanov ND, Zherebtsov SV et al (2017) Structure and mechanical properties of B2 ordered refractory AlNbTiVZrx (x = 0–1.5) high-entropy alloys. Mater Sci Eng: A 704:82–90

32

Y. Wu et al.

47. Senkov ON, Woodward C, Miracle DB (2014) Microstructure and properties of aluminumcontaining refractory high-entropy alloys. Jom 66(10):2030–2042 48. Senkov ON, Senkova SV, Woodward C (2014) Effect of aluminum on the microstructure and properties of two refractory high-entropy alloys. Acta Mater 68:214–228 49. Nagase T, Anada S, Rack PD et al (2012) Electron-irradiation-induced structural change in Zr–Hf–Nb alloy. Intermetallics 26:122–130 50. Nagase T, Rack PD, Noh JH et al (2015) In-situ TEM observation of structural changes in nano-crystalline CoCrCuFeNi multicomponent high-entropy alloy (HEA) under fast electron irradiation by high voltage electron microscopy (HVEM). Intermetallics 59:32–42 51. Kumar NAPK, Li C, Leonard KJ et al (2016) Microstructural stability and mechanical behavior of FeNiMnCr high entropy alloy under ion irradiation. Acta Mater 113:230–244 52. Lu Y, Huang H, Gao X et al (2019) A promising new class of irradiation tolerant materials: Ti2 ZrHfV0.5 Mo0.2 high-entropy alloy. J Mater Sci Technol 35(3):369–373 53. El-Atwani O, Li N, Li M et al (2019) Outstanding radiation resistance of tungsten-based high-entropy alloys. Sci Adv 5(3):eaav2002 54. Sadeghilaridjani M, Ayyagari A, Muskeri S et al (2020) Ion irradiation response and mechanical behavior of reduced activity high entropy alloy. J Nucl Mater 529:151955 55. Zhao S (2020) Defect properties in a VTaCrW equiatomic high entropy alloy (HEA) with the body centered cubic (bcc) structure. J Mater Sci Technol 44:133–139 56. Tong Y, Velisa G, Zhao S et al (2018) Evolution of local lattice distortion under irradiation in medium-and high-entropy alloys. Materialia 2:73–81 57. Lu C, Niu L, Chen N et al (2016) Enhancing radiation tolerance by controlling defect mobility and migration pathways in multicomponent single-phase alloys. Nat Commun 7(1):1–8 58. Egami T, Guo W, Rack PD et al (2014) Irradiation resistance of multicomponent alloys. Metall and Mater Trans A 45(1):180–183 59. Xia S, Zhen W, Yang T et al (2015) Irradiation behavior in high entropy alloys. J Iron Steel Res Int 22(10):879–884 60. Wu Y, Zhang F, Yuan X et al (2021) Short-range ordering and its effects on mechanical properties of high-entropy alloys. J Mater Sci Technol 62:214–220 61. George EP, Raabe D, Ritchie RO (2019) High-entropy alloys. Nat Rev Mater 4(8):515–534 62. Yao HW, Qiao JW, Gao MC, Hawk JA, Ma SG, Zhou HF, Zhang Y (2016) NbTaV-(Ti, W) refractory high-entropy alloys: experiments and modeling. Mater Sci Eng, A 674:203–211 63. Lin C-M, Juan C-C, Chang C-H, Tsai C-W, Yeh J-W (2015) Effect of Al addition on mechanical properties and microstructure of refractory AlxHfNbTaTiZr alloys. J Alloy Compd 624:100–107 64. Tang Z, Gao MC, Diao H, Yang T, Liu J, Zuo T, Zhang Y, Lu Z, Cheng Y, Zhang Y, Dahmen KA, Liaw PK, Egami T (2013) Aluminum alloying effects on lattice types, microstructures, and mechanical behavior of high-entropy alloys systems. JOM 65(12):1848–1858 65. Zhang C, Zhang F, Chen S, Cao W (2012) Computational thermodynamics aided high-entropy alloy design. JOM 64(7):839–845 66. Wang Z, Wu H, Wu Y, Huang H, Zhu X, Zhang Y, Zhu H, Yuan X, Chen Q, Wang S, Liu X, Wang H, Jiang S, Kim MJ, Lu Z (2022) Solving oxygen embrittlement of refractory high-entropy alloy via grain boundary engineering. Mater Today 67. Pang J, Zhang H, Zhang L, Zhu Z, Fu H, Li H, Wang A, Li Z, Zhang H (2020) Simultaneous enhancement of strength and ductility of body-centered cubic TiZrNb multi-principal element alloys via boron-doping. J Mater Sci Technol 68. Ahmad AS, Su Y, Liu SY, Ståhl K, Wu YD, Hui XD, Ruett U, Gutowski O, Glazyrin K, Liermann HP, Franz H, Wang H, Wang XD, Cao QP, Zhang DX, Jiang JZ (2017) Structural stability of high entropy alloys under pressure and temperature. J Appl Phys 121(23):235901 69. Yusenko KV, Riva S, Crichton WA, Spektor K, Bykova E, Pakhomova A, Tudball A, Kupenko I, Rohrbach A, Klemme S, Mazzali F, Margadonna S, Lavery NP, Brown SGR (2018) Highpressure high-temperature tailoring of high entropy alloys for extreme environments. J Alloy Compd 738:491–500

1 Body-Centered Cubic High-Entropy Alloys

33

70. Guo J, Wang H, Rohr Fv, Wang Z, Cai S, Zhou Y, Yang K, Li A, Jiang S, Wu Q, Cava RJ, Sun L (2017) Robust zero resistance in a superconducting high-entropy alloy at pressures up to 190 GPa. Proceed National Acad Sci 114(50):13144–13147 71. Cheng B, Zhang F, Lou H, Chen X, Liaw PK, Yan J, Zeng Z, Ding Y, Zeng Q (2019) Pressureinduced phase transition in the AlCoCrFeNi high-entropy alloy. Scripta Mater 161:88–92 72. Wang W-R, Wang W-L, Wang S-C, Tsai Y-C, Lai C-H, Yeh J-W (2012) Effects of Al addition on the microstructure and mechanical property of AlxCoCrFeNi high-entropy alloys. Intermetallics 26:44–51 73. Zhang R, Zhao S, Ding J, Chong Y, Jia T, Ophus C, Asta M, Ritchie RO, Minor AM (2020) Short-range order and its impact on the CrCoNi medium-entropy alloy. Nature 581(7808):283–287 74. Dasari S, Jagetia A, Sharma A, Nartu MSKKY, Soni V, Gwalani B, Gorsse S, Banerjee R (2021) Tuning the degree of chemical ordering in the solid solution of a complex concentrated alloy and its impact on mechanical properties. Acta Materialia 212 75. Ding Q, Zhang Y, Chen X, Fu X, Chen D, Chen S, Gu L, Wei F, Bei H, Gao Y, Wen M, Li J, Zhang Z, Zhu T, Ritchie RO, Yu Q (2019) Tuning element distribution, structure and properties by composition in high-entropy alloys. Nature 574(7777):223–227 76. Chen X, Wang Q, Cheng Z, Zhu M, Zhou H, Jiang P, Zhou L, Xue Q, Yuan F, Zhu J, Wu X, Ma E (2021) Direct observation of chemical short-range order in a medium-entropy alloy. Nature 592(7856):712–716 77. Bu YQ, Wu Y, Lei ZF, Yuan XY, Wu HH, Feng XB, Liu JB, Ding J, Lu Y, Wang HT, Lu ZP, Yang W (2021) Local chemical fluctuation mediated ductility in body-centered-cubic high-entropy alloys. Mater Today 46:28–34 78. Du Q, Liu XJ, Fan HY, Zeng QS, Wu Y, Wang H, Chatterjee D, Ren Y, Ke YB, Voyles PM, Lu ZP, Ma E (2020) Reentrant glass transition leading to ultrastable metallic glass. Mater Today 34:66–77 79. Zhang W, Liaw PK, Zhang Y (2018) Science and technology in high-entropy alloys. Sci China Mater 61(1):2–22 80. Wen XC, Wu Y, Huang HL, Jiang SH, Wang H, Liu XJ, Zhang Y, Wang XZ, Lu ZP (2021) Effects of Nb on deformation-induced transformation and mechanical properties of HfNbx Ta0.2 TiZr high entropy alloys. Mat Sci Eng A-Struct 805 81. Senkov ON, Miracle DB, Chaput KJ, Couzinie J-P (2018) Development and exploration of refractory high entropy alloys—a review. J Mater Res 33(19):3092–3128 82. Chen W, Tang QH, Wang H, Xie YC, Yan XH, Dai PQ (2018) Microstructure and mechanical properties of a novel refractory AlNbTiZr high-entropy alloy. Mater Sci Technol 34(11):1309– 1315 83. Senkov ON, Isheim D, Seidman DN, Pilchak AL (2016) Development of a refractory high entropy superalloy. Entropy 18(3):102 84. Senkov ON, Rao S, Chaput KJ, Woodward C (2018) Compositional effect on microstructure and properties of NbTiZr-based complex concentrated alloys. Acta Mater 151:201–215 85. Soni V, Senkov ON, Gwalani B, Miracle DB, Banerjee R (2018) Microstructural design for improving ductility of an initially brittle refractory high entropy alloy. Sci Rep 8(1):8816 86. Soni V, Gwalani B, Senkov ON, Viswanathan B, Alam T, Miracle DB, Banerjee R (2018) Phase stability as a function of temperature in a refractory high-entropy alloy. J Mater Res 33(19):3235–3246 87. Jensen JK, Welk BA, Williams REA, Sosa JM, Huber DE, Senkov ON, Viswanathan GB, Fraser HL (2016) Characterization of the microstructure of the compositionally complex alloy Al1 Mo0.5 Nb1 Ta0.5 Ti1 Zr1 . Scripta Materialia 121:1–4 88. Takeuchi A, Inoue A (2005) Classification of bulk metallic glasses by atomic size difference, heat of mixing and period of constituent elements and its application to characterization of the main alloying element. Mater Trans 46(12):2817–2829 89. Yeh SKCJW, Lin SJ, Gan JY, Chin TS, Shun TT, Tsau CH, Chang SY (2004) Nanostructured high-entropy alloys with multiple principal elements: novel alloy design concepts and outcomes. Adv Eng Mater 6:299

34

Y. Wu et al.

90. Cantor B, Chang ITH, Knight P, Vincent AJB (2004) Microstructural development in equiatomic multicomponent alloys. Mater Sci Eng, A 375–377:213–218 91. Senkov ON, Scott JM, Senkova SV, Miracle DB, Woodward CF (2011) Microstructure and room temperature properties of a high-entropy TaNbHfZrTi alloy. J Alloy Compd 509(20):6043–6048 ˇ 92. Cížek J, Haušild P, Cieslar M, Melikhova O, Vlasák T, Janeˇcek M, Král R, Harcuba P, Lukáˇc F, Zýka J, Málek J, Moon J, Kim HS (2018) Strength enhancement of high entropy alloy HfNbTaTiZr by severe plastic deformation. J Alloy Compd 768:924–937 93. Senkov ON, Semiatin SL (2015) Microstructure and properties of a refractory high-entropy alloy after cold working. J Alloy Compd 649:1110–1123 94. Wu YD, Cai YH, Wang T, Si JJ, Zhu J, Wang YD, Hui XD (2014) A refractory Hf25 Nb25 Ti25 Zr25 high-entropy alloy with excellent structural stability and tensile properties. Mater Lett 130:277–280 95. Lilensten L, Couzinié J-P, Bourgon J, Perrière L, Dirras G, Prima F, Guillot I (2016) Design and tensile properties of a bcc Ti-rich high-entropy alloy with transformation-induced plasticity. Mater Res Lett 5(2):110–116 96. Senkov ON, Scott JM, Senkova SV, Meisenkothen F, Miracle DB, Woodward CF (2012) Microstructure and elevated temperature properties of a refractory TaNbHfZrTi alloy. J Mater Sci 47(9):4062–4074 97. Guo NN, Wang L, Luo LS, Li XZ, Su YQ, Guo JJ, Fu HZ (2015) Microstructure and mechanical properties of refractory MoNbHfZrTi high-entropy alloy. Mater Des 81:87–94 98. Feng R, Feng B, Gao MC, Zhang C, Neuefeind JC, Poplawsky JD, Ren Y, An K, Widom M, Liaw PK (2021) Superior high-temperature strength in a supersaturated refractory highentropy alloy. Adv Mater e2102401 99. Wang S, Wu M, Shu D, Zhu G, Wang D, Sun B (2020) Mechanical instability and tensile properties of TiZrHfNbTa high entropy alloy at cryogenic temperatures. Acta Mater 201:517– 527 100. Qiao JW, Ma SG, Huang EW, Chuang CP, Liaw PK, Zhang Y (2011) Microstructural characteristics and mechanical behaviors of AlCoCrFeNi high-entropy alloys at ambient and cryogenic temperatures. Mater Sci Forum 688:419–425 101. He JY, Liu WH, Wang H, Wu Y, Liu XJ, Nieh TG, Lu ZP (2014) Effects of Al addition on structural evolution and tensile properties of the FeCoNiCrMn high-entropy alloy system. Acta Mater 62:105–113 102. Senkov ON, Wilks GB, Miracle DB, Chuang CP, Liaw PK (2010) Refractory high-entropy alloys. Intermetallics 18(9):1758–1765 103. Poulia A, Georgatis E, Lekatou A, Karantzalis A (2017) Dry-sliding wear response of MoTaWNbV high entropy alloy. Adv Eng Mater 19(2):1600535 104. Ye YX, Liu CZ, Wang H, Nieh TG (2018) Friction and wear behavior of a single-phase equiatomic TiZrHfNb high-entropy alloy studied using a nanoscratch technique. Acta Mater 147:78–89

Chapter 2

Face-Centered Cubic High-Entropy Alloys Weihong Liu and Boxuan Cao

1 Introduction Since the Bronze Age, the alloy design strategy is based on one or two primary metals and minor other alloying elements are added to obtain optimum properties. For instance, steels, the most widely used structural materials, consist primarily of iron, and minor secondary elements such as carbon and chromium are selectively alloyed to improve strength and corrosion resistance, respectively. However, after thousandyear exploration, the compositional space of traditional alloys has been extensively excavated and approaches their property limits, such as the trade-off between strength and plasticity. Currently, there is a pressing need for high-performance structural alloys in energy efficiency, industrial production, national defense technology and people’s livelihood. Therefore, new approaches are needed to design and develop high-performance structural alloys. In 2004, Cantor et al. [1] and Yeh et al. [2] from two independent research groups proposed a multi-principal component alloy design strategy, i.e., high-entropy alloys (HEAs). HEAs are usually made up of five or more metallic elements in relatively high concentrations (5–35 at.%) and hence also termed multi-component alloys or complex concentrated alloys. This equal proportions chemical composition in HEAs would maximize the mixing configurational entropic contribution to overcome the enthalpies of compound formation and then stabilize simple solid-solution phases [2]. These entropy-stabilized complex concentrated solid-solution alloys exhibit four unique features [3]: high-entropy effect, severe lattice distortion, sluggish diffusion W. Liu (B) School of Materials Science and Engineering, Harbin Institute of Technology (Shenzhen), Shenzhen, China e-mail: [email protected] B. Cao Department of Materials Science and Engineering, City University of Hong Kong, Hong Kong, China © The Author(s), under exclusive license to Springer Nature Singapore Pte Ltd. 2022 Z. Jiao and T. Yang (eds.), Advanced Multicomponent Alloys, Materials Horizons: From Nature to Nanomaterials, https://doi.org/10.1007/978-981-19-4743-8_2

35

36

W. Liu and B. Cao

effect, and cocktail effect. As the understanding of HEAs evolves, the local chemical ordering effect was proposed. The compositional concentrated alloys are not ideal solids, in which the atoms of different elements are not randomly distributed at lattice positions, and local chemical order was identified, as shown in Fig. 1 [4]. Such synergies result in promising mechanical properties, such as the excellent combination of strength and plasticity [5], high resistance to low-temperature brittleness [6], radiation [7], and better hydrogen resistance as compared to austenitic stainless steels [8–10]. Fortunately, the HEA design strategy has overwhelming advantages over traditional alloys which are based on one-single principal element in seeking high-performance structural and functional materials, since they provide an infinite compositional design space for tailoring phase structures, stacking-fault energies, and associated deformation mechanisms. The first research effort was placed on the single-phase FeCoNiCrMn HEA with the face center cubic (FCC) structure, which exhibits outstanding tensile ductility and toughness at low temperatures down to 77 K [5]. With extensive research activities on HEAs in recent years, the definition and understanding of HEAs are developing continuously. One of the important findings is that in most cases the configuration entropy of HEAs cannot suppress the formation of intermetallic compounds and

Fig. 1 Composition inhomogeneity in a CrFeCoNiPd HEA [4]

2 Face-Centered Cubic High-Entropy Alloys

37

then does not completely determine the corresponding alloy’s structures and properties. Furthermore, the single-phase FCC HEA alloys alone were found to be insufficiently strong for structural applications. However, its good plasticity and outstanding strain-hardening capacity enable them excellent base alloys for further precipitation strengthening [11–13]. The precipitation-strengthening effects of various kinds of second-phase intermetallics, including σ, μ, Laves, B2-NiAl, body center cubic (BCC), carbides, and L12 phases, have been extensively studied [13–22]. Among all these precipitation-strengthening HEAs, the dual-phase HEAs with a high density of coherent precipitation with L12 -type nanoparticles in the FCC high-entropy matrix exhibit the most promising mechanical properties [17, 19, 21]. For the past ten years, considerable progress has been made in developing and exploring the property limit of advanced HEAs. Here, we aim to make a comprehensive summary of these advances and point out challenges needed to be overcomed in the future. This chapter reviews these advances via three sections. The introduction section is followed by a description of the characteristic of single-phase HEAs with the FCC structure, where the emphasis was placed on mechanical properties under extreme environments. The third section summarizes the mechanical properties of multiple-phase HEAs at room and elevated temperatures, including incoherent and coherent precipitation strengthened HEAs.

2 The Single-Phase FCC HEAs The FCC structure alloys. The five-element equiatomic CrMnFeCoNi alloy was the first HEA reported having the single-phase FCC solid solution [1]. However, after heat treatment below 800 °C, the solid-solution phase decomposes into metallic (BCC-Cr) and intermetallic (L10 -NiMn and B2-FeCo) phases [14, 23]. Regardless, after annealing at the high temperatures above 800 °C, the metastable FCC phase can be preserved at room temperature upon ‘normal’ cooling rates. Noted that no indication of clustering or short-range ordering was characterized in the quenched FCC solid solution. Theoretical calculations reveal that the FCC structure becomes thermodynamics unstable and transforms to the hexagonal close-packed (HCP) structure when the temperature is below 50 K [24]. Although no direct experimental evidences confirm this temperature-induced phase transformation, a high-pressureinduced transformation from FCC to HCP was reported [25, 26]. Such phase transformation would undoubtedly affect the alloy’s mechanical performance, since the HCP phase is much harder than the FCC phase, but less slip systems are available. The CrMnFeCoNi alloy exhibits decent strength, ductility, and fracture toughness over a wide temperature range. Gludovatz et al. [5] reported that the FCC CrMnFeCoNi HEA has a yield strength of 410 MPa and ultimate tensile strength of 763 MPa with a tensile elongation of 57% at room temperature. The alloy has strong temperature-dependent strength and ductility when the testing temperature decreases from 293 to 77 K, with the yield strength and ultimate tensile strength increased by ~85% to 759 MPa and ~70% to 1280 MPa, respectively. Meanwhile, the tensile

38

W. Liu and B. Cao

elongation increased by ~25–70%. The strain-hardening exponent n is as high as ~0.4, which contributes to such good plasticity and high fracture toughness. Furthermore, the alloy has good resistance to hydrogen embrittlement (HE). Zhao et al. [27] reported that although the hydrogen adsorption capacity of CrMnFeCoNi (76.5 wt.ppm) is higher than that of 304 (63 wt.ppm) and 316L (36.5 wt.ppm) stainless steels, the alloy only suffers a small ductility loss of ~5% after 15 MPa gas hydrogen charging (300 °C, 15 MPa, 72 h), as shown in Fig. 2a. It is much less than that of ~61% for 304 and of ~27% for 316L austenitic steels. The fracture mode of 304 and 316 stainless steels changed from micro-voids coalescence ductile fracture to brittle intergranular fracture. By contrast, the fracture mode of the CrMnFeCoNi alloy is still ductile, with the fracture surface covered by ductile dimples (Fig. 2a). However, Nygren et al. [8] charged the CoCrFeMnNi alloy in an autoclave at a hydrogen gas pressure of 120 MPa and a temperate of 200 °C for 160 h, which introduced 146.9 wt. ppm hydrogen to the alloy. The hydrogen-induced ductility loss was measured to be 33%, and the fracture mode has changed from the ductile transgranular mode to brittle intergranular mode. Studies by Nygren et al. [8] and Zhao et al. [27] together indicate that the response to HE of the CrMnFeCoNi alloy is directly related with the hydrogen concentration, where hydrogen-induced grain boundary embrittlement occurs in the hydrogen concentration of 146.9 wt. ppm but not in 76.5 mass ppm. Based on the equimolar quintuple CrMnFeCoNi alloy, several related variants by removing alloying elements or adjusting the element ratio have been developed. One of the rationales was to reduce the stacking-fault energy to promote twinning induced plasticity (TWIP) and/or transformation induced plasticity (TRIP), which could improve the strain-hardening ability of materials. For instance, Li et al. [20] reported a deformation-induced martensitic transformation from a metastable FCC to HCP phase in the 50Fe–30Mn–10Co–10Cr (at.%) alloy, which achieved a good combination of strength and plasticity and provided a valuable design direction to explore HEAs with superior mechanical properties. Another effective method was to simply remove elements from the alloy. Two notable examples are the CrFeCoNi HEA and the CrCoNi medium-entropy alloy. Liu et al. [29] reported that after removing Mn the resulted CoCrFeNi HEA also has the FCC structure and exhibits an outstanding mechanical property even down to 4 K, with a tensile strength of 1260 MPa and a tensile elongation of 62%. More importantly, the fine-grained alloy (1.9 μm) shows strong resistance to HE [9], with the tensile elongation to failure (51–49%) barely no loss after 100 MPa hydrogen gas charging, as shown in Fig. 2b. The fracture mode is still the ductile transgranular mode, without the occurrence of hydrogen-induced grain boundary embrittlement (Fig. 2b). After further removal of Fe, the resulted single-phase FCC CrCoNi medium-entropy alloy exhibits more attractive mechanical strength and ductility, as well as fracture toughness than both the quintuple CrMnFeCoNi and quaternary CrFeCoNi HEAs. Gludovatz et al. [6] showed that the CrCoNi alloy exhibited a tensile strength of ∼1 GPa, failure strain of approximately 70% and a K JIc fracture toughness above 200 MPa m1/2 at room temperature. With the temperature decreased to 77 K, the tensile strength increases to 1.3 GPa and elongation to failure reached ∼90%, with a K JIc fracture toughness of 275 MPa m1/2 at crack initiation and exceeding 400 MPa.m1/2 for crack growth,

2 Face-Centered Cubic High-Entropy Alloys

39

Fig. 2 Engineering stress–engineering strain curves and fracture surfaces of the a CrMnFeCoNi [27] and b CrFeCoNi [28] HEAs with and without presence of hydrogen (with data indicating ductility loss due to hydrogen charging)

making it one of the most damage-tolerant the exceptional mechanical properties [6]. The alloy’s low-temperature properties are impressive; however, Yi et al. [30] showed that the CrCoNi is susceptible to HE. After 100 MPa hydrogen gas charging at 200 °C for 160 h, a very high ductility loss of 85.3% accompanied with a fracture mode transition from transgranular to intergranular was observed. Fortunately, minor alloying of Mo was found to partially suppress such hydrogen-induced grain boundary embrittlement and reduced the ductility loss to 41.6% [30].

40

W. Liu and B. Cao

All the above studies on HEAs with the FCC structure show that they have outstanding mechanical properties under extreme conditions, mainly at low temperatures and in hydrogen-containing environments. However, their properties were strongly influenced by alloy compositions and microstructures, since the composition sets an alloy’s elastic properties and atomic interactions, which then dominate its dislocation behaviors under mechanical loading. For multi-component HEAs, adjusting alloy compositions and microstructures could have profound effects on their mechanical properties, and then, proper compositions and thermomechanical processing should be adopted for tailoring superior mechanical properties.

3 Precipitation-Strengthened HEAs HEAs are firstly designed to stabilize the solid-solution phase based on the concept that high configurational entropy could suppress phase separation and compound formation. However, only a few HEAs exhibiting the single solid-solution phase with the FCC or BCC structure were developed. In contrast, most HEAs tend to form dual-phase or even multi-phase structures. As we stated in the first section, the single-phase FCC Cr–Fe–Co–Ni–Mn alloy systems alone were insufficiently strong for structural applications, but their outstanding strain-hardening capacity enable them excellent base alloys for further precipitation strengthening [11–13]. Therefore, various intermetallic compounds were introduced to strengthen the FCC high-entropy matrix and form multi-phase HEAs. These multi-phase HEAs not only promote a fundamental investigation of phase separation, but also increase the achievable property spectrum for HEAs with superior properties. In multi-phase HEAs, mechanical properties strongly depend on configuration of the constituent phases, phase interface damage resistance, and coordination deformation ability. Understanding the micromechanical behavior of multi-phase alloys is so important that could shed light on the macro-mechanical behavior and guide alloy composition designs and applications.

3.1 Incoherent Precipitate-Strengthened HEAs The strengthening effect is strongly dependent on the type, size, and distribution of the precipitated particles, which interact with and thus retard moving dislocation, thereby determining how a material is strengthened. Let us talk about the strengthening effects of incoherent precipitates such as μ, Laves, B2-NiAl, and BCC phases. First, adding Al into the CoCrFeNi HEA would promote the formation of the BCC phase in the FCC matrix [31]. When the Al concentration increases, the crystal structure of the matrix would transform from FCC to BCC. Such transition would enhance the alloy’s strength with compromising ductility. Furthermore, the eutectic solidification HEAs with a hard intermetallic phase in the FCC matrix have the advantage of near-equilibrium microstructures and could be promising for the balanced strength

2 Face-Centered Cubic High-Entropy Alloys

41

and ductility. Guo et al. [32] designed the first as-cast AlCoCrFeNi2.1 eutectic HEA, which possessed a fine lamellar FCC/B2 microstructure and showed an good combination of high tensile ductility (25.6%) and large fracture strength (994 MPa) at room temperature. Such excellent mechanical properties could be retained up to 700 °C. He et al. [18] and Liu et al. [22] reported that Nb can provide an excellent solution strengthening effect due to its large atomic size compared to other principal elements in the CoCrFeNiNbx alloy system and more importantly, is able to facilitate the formation of eutectic structures. With the aid of the pseudo-binary (CoCrFeNi)– Nb diagram, Liu et al. [22] developed a eutectic Nb-modified (CoCrFeNi)–Nb alloy, which was composed of a ductile FCC phase and a hard Nb-enriched Laves phase. The Laves phase is a well-known size compound, which is much harder than the FCC matrix. Thus, the resulted fine laminar eutectic structure in the CoCrFeNiNb0.155 alloy leads to promising mechanical properties with the yield strength and tensile elongation as high as 321 MPa and 21.3%, respectively, as shown in Fig. 3a. Liu et al. [13] further reported that the precipitation of hard topologically close-packed (TCP) phases could effectively harden the CoCrFeNi matrix than the Nb-induced eutectic structure. Mo alloying promoted co-precipitation of fine and hard (Cr, Mo)-rich σ phases and (Mo, Cr)-enriched μ phases in the CoCrFeNiMo0.3 alloy. As shown in Fig. 3b, these σ and μ particles are much harder than the FCC matrix, having a hardness over 8 GPa. Interestingly, these precipitates greatly improved the yield and tensile strength of the CoCrFeNiMo0.3 alloy without causing serious embrittlement. For instance, aging at 850 °C for 1 h leads to an ultimate tensile strength of 1.2 GPa with an attractive ductility of ∼19% in the CoCrFeNiMo0.3 alloy, as shown in Fig. 3b. Fu et al. [33] found that a synergistic alloy of Nb and Mo induced the precipitation of a new kind of Laves phase with the HCP structure in the FCC matrix. The synergistic precipitation strengthening and solid-solution effect result in a high yield and fracture strength of 426 MPa and 714 MPa, while maintaining a high tensile elongation of 17.4%. All these results evidently show that the precipitation of incoherent precipitates could effectively strengthen the FCC matrix without or with little sacrificing ductility. To achieve superior mechanical properties, optimized alloying elements and proper heat treatment procedures should be carefully chosen.

3.2 Coherent Precipitate-Strengthened HEAs for Advanced Structural Applications In 2016, He et al. [17] first reported that a small volume fraction of coherent L12 nanoparticles (~20%) precipitated in the (FeCoNiCr)94 Ti2 Al4 alloy. The compositions of commonly utilized L12 particles were shown in Fig. 4. By tailoring the compositions and thermomechanical processing, the optimized FCC-L12 microstructure resulted in an excellent integrated tensile property in the (FeCoCrNi)96 Ti2 Al4 HEA at room temperature: a high tensile strength of 1273 MPa and a tensile ductility of 17%. In 2017, Zhao et al. [19] developed the (CoCrNi)94 Al3 Ti3 medium-entropy

42

W. Liu and B. Cao

Fig. 3 Microstructures and tensile strain–stress curves of the a CoCrFeNiNbx (x = 0, 0.155, 0.206 and 0.309) and b CoCrFeNiMox (x = 0, 0.2 and 0.3) alloy system [13, 22]

alloy, which also has dual FCC/L12 microstructures. Compared with He’s work, the tensile ductility improved to 45% while maintaining a similar strength level. Yang et al. [34] introduced a high-density ductile L12 multi-component nanoparticles in the complex (FeCoNi)86 Al7 Ti7 alloy. The nanoprecipitate-strengthened alloy exhibits a superior tensile strength of 1.5 GPa, which was five times higher than that of the single-phase FeCoNi-based alloy, accompanied by a ductility as high as 50% at the ambient temperature. This alloy has an extremely high value of σUTS × ELT (72 GPa %), which is impressive and exceeds many other high-performance alloys at room temperature. The rapid development of modern industries raises urgent requests for hightemperature materials with superior temperature capabilities. The microstructure of the HEAs strengthened by multi-component L12 intermetallic particles resembles that of the conventional Ni-based superalloys with ordered γ' precipitates embedded in the solid-solution γ matrix, which motivates researchers to explore the possibilities of developing L12 -strengthened HEAs toward elevated temperature environments [17, 36]. The following achievements and challenges in developing advanced high-temperature HEAs have been briefly reviewed here, including microstructural stability, particle coarsening resistance, high-temperature strength, and intermediate-temperature embrittlement issues.

2 Face-Centered Cubic High-Entropy Alloys

43

Fig. 4 APT results of the L12 -strengthened Fe–Co–Ni–Al–Ti-based HEAs [35]

It should be noted that most of the precipitation-hardened HEAs originally designed for cryogenic and ambient temperature applications possess a small volume fraction of the L12 phase with an inadequate heat resistance [19]. For example, the volume fraction of the L12 phase is only ~20% in the (NiCoCr)94 Al3 Ti3 (at.%) HEA (Fig. 5a), which vanished as the temperature increased above 900 °C [19]. Since a high L12 volume fraction is favorable for providing a sufficient deformation resistance at elevated temperatures, efforts have been devoted to stabilizing the L12 particles and increasing the γ' -solvus temperature through alloying additions [37]. Various elements have been identified as potent L12 -stabilizing elements, including Al, Ti, Cu, Nb, Ta, and V [38, 39]. It has been demonstrated that the γ' -solvus temperature can be stabilized up to 1100 °C by carefully controlling alloying additions (Fig. 5b), which is comparable with advanced Ni-based superalloys [38]. Cu additions have been shown to promote L12 phase formation in the Ni2 CrFeCuAl0.3 HEA by forming Cu-rich clusters and acting as heterogeneous sites for the nucleation of L12 precipitates (Fig. 5c) [40, 41]. It should be noted that the stabilization of L12 strengtheners should not be achieved at the expense of degrading the γ/γ' dual-phase region and associated brittle deleterious intermetallic phase formation. These brittle intermetallic phases might lead to severe stress concentration and associated embrittlement by acting as sites for crack initiation and propagation [36, 42]. Under certain circumstances, these harmful intermetallic phases also deplete refractory elements from the matrix phase, decreasing the effectiveness of solidsolution strengthening. Al, Ti, and Nb elements are commonly used to induce L12 phase formation by acting as L12 -stabilizing elements; however, their concentrations should be appropriately controlled to avoid undesirable intermetallic phase formation. For example, as a result of excess Al additions, the needle-shaped B2-NiAl phase formed in Ni36 Co25 Fe15 Cr8 Al10 Ti6 and Ni16 Co15 Fe15 Cr23 Cu8 Al23 HEAs [43, 44]. The formation of the ï-Ni3 Ti phase has also been identified due to the presence of excess Ti in the Ni1.5 Co1.5 FeCrAl0.1 Ti0.4 HEA [45]. Various other kinds of harmful

44

W. Liu and B. Cao

intermetallic phases have also been identified in L12 -strengthened HEAs, including the Nb-enriched Laves phase in the Ni22 Co30 Fe18 Cr18 Al8 Nb4 HEA (Fig. 5d) [46], the L21 -type Heusler phase in the Ni30 Co30 Fe13 Cr15 Al6 Ti6 HEA (Fig. 5e) [42], and the D019 phase in the Co-30Ni-7Al-4Ti-5Mo-1Nb-1Ta HEA (Fig. 5f) [47]. Such brittle intermetallic phase formation should be appropriately addressed in designing and developing novel HEAs with high performance. However, it is challenging to obtain desired microstructures with targeted phases in designing multi-component alloys due to the wide compositional space and numerous elemental combinations. The computational-aid Calculation of Phase Diagrams (CALPHAD) technique offers an effective method to rapidly screen compositions and accelerates the alloy design [48]. The CALPHAD method has been successfully applied to guide the design of HEAs recently [19, 49]. A stable γ/γ' dual-phase microstructue with adequate resistance toward particle coarsening is preferred in designing L12 -strengthened HEAs for elevated temperature applications. It is reported that the coarsening rate of the L12 particles in the (NiCoCr)94 Al3 Ti3 and (NiCoFeCr)94 Al4 Ti2 (at.%) HEAs at 800 °C are slower than those of commercial Ni-based superalloys [17, 50]. Moreover, it is interesting to note that the coarsening rate can be profoundly retarded via concentrated Coconcentrations. For example, particles coarsen rapidly in the binary Ni-13.6Al alloy [51], whereas alloying addition of Co-significantly slows down the particle growth in the Ni–21.7Co–13.4Al alloy [52]. Particle coarsening can be further inhibited in

Fig. 5 a Spherical L12 particles embedded in the FCC matrix of the (NiCoCr)94 Al3 Ti3 HEA after aging at 800 °C [17]. b Morphologies of the Co43.5 Ni29 Cr10 Al10 Ti2 Mo2 Ta2 Nb1.5 HEA after long-term thermal exposure at 1100 °C, showing superb L12 phase stability [38]. c EDX compositional mapping of the Ni2 CrFeCuAl0.3 HEA, showing Cu-rich clusters in assisting L12 precipitation [40]. d NiAl and Laves phase formation in the Ni22 Co30 Fe18 Cr18 Al8 Nb4 HEA [46]. e The precipitation of D019 phase at the grain boundaries in the Co–30Ni–7Al–4Ti–5Mo–1Nb–1Ta HEA [47]. f Preferential nucleation of L21 -type precipitates in the vicinity of grain boundaries in the Ni30 Co30 Fe13 Cr15 Al6 Ti6 HEA [42]

2 Face-Centered Cubic High-Entropy Alloys

45

HEAs with higher Co-concentrations [53–55]. As a result, Co-rich HEAs demonstrated exceptional particle coarsening resistance (Fig. 6), which is vital for hightemperature applications. It should also be noted that the excess Co-content tends to destabilize the L12 structure and promote the formation of the disordered FCC structure. Therefore, alloy compositions should be carefully adjusted to leverage phase stability and coarsening resistance. The reduced particle coarsening rate in HEAs is associated with the sluggish diffusion in the high-entropy matrix phase. By evaluating the diffusion parameters of constitutional elements in HEAs, it is reported that, as compared with conventional FCC-structured metals, the larger lattice potential energy fluctuation in the HEAs leads to a higher diffusion activation energy and the associated sluggish diffusion kinetics in the temperature range between 900 and 1100 °C [56]. Since the Lifshitz-Slyozov-Wagner model assumes that the solute diffusion in the matrix is the rate-limiting step for the particle coarsening, the sluggish solute diffusion in the multi-component matrix is thereby beneficial for the microstructural stability among HEAs [57]. Moreover, the activation energy for L12 particle coarsening has been determined to be 378 kJ/mol in the Ni30 Co30 Fe13 Cr15 Al6 Ti6 HEA [42] and 360 kJ/mol in the Co37.6 Ni35.4 Al9.9 Mo4.9 Cr5.9 Ta2.8 Ti3.5 HEA [53]. Such values are generally higher than most solutes in Ni- and Co-based alloys with values ranging from 254 to 289 kJ/mol [58]. The higher activation energy suggests a higher diffusion barrier for vacancy formation and migration. This conclusion is in line with the proposed sluggish diffusion effect in HEAs. Promising mechanical properties have been demonstrated in L12 -strengthened HEAs at elevated temperatures. The high-temperature hardness of the Ni48.6 Co17 Fe9 Cr7.5 Al10.3 Ti5.8 Ta0.6 Mo0.8 W0.4 HEA is even superior to the Ni-based superalloy CM247LC, even though the latter comprises a higher concentration of refractory elements [61]. Preliminary creep properties have been reported

Fig. 6 Comparison of the particle coarsening rate in multiple alloys. Coarsening resistance can be improved via the concentrated Co-content [51–54, 59, 60]

46

W. Liu and B. Cao

among L12 -strengthed HEAs. It has been found that the creep rupture life of the Ni47.9 Fe8.9 Co16.9 Cr7.4 Al10.2 Ti5.8 Mo0.9 Nb1.2 W0.4 C0.4 HEA at 982 °C/159 MPa is 114 h, which is comparable with first-generation Ni-based superalloys (118 h for Rene 80 under 982 °C/145 MPa) [62, 63]. Moreover, low creep rates have also been found in the same HEA, demonstrating low creep strains of 0.04% at 750 °C/159 MPa and 0.25% at 850 °C/159 MPa for 800 h. These preliminary results demonstrated the possibilities of developing HEAs in an elevated temperature environment. In particular, the yield anomaly behavior has been found in HEAs, including Ni36 Co25 Fe15 Cr8 Al10 Ti6 , various L12 -strengthened Ni48.6 Co17 Fe9 Cr7.5 Al10.3 Ti5.8 Ta0.6 Mo0.8 W0.4 , and Ni40.7 Co20.6 Fe11.5 Cr12.2 Al7.8 Ti7.2 HEAs [43, 61]. Such unique behavior is believed to be inherited from the ordered L12 intermetallic phase, which is a result of thermally activated dislocation cross-slip from the octahedral plane to the cube plane and associated Kear-Wilsdorf locks formation in the ordered intermetallic phase [64]. Another different scenario has recently been found in the Co53 Ni30 Al10 Nb3 Ti2 Ta2 HEA upon deformation at 700 °C, in which extensive superlattice stacking faults have been observed in the L12 precipitates [65]. The interactions between the superlattice stacking faults resulted in sessile stair-rod dislocation configurations, preventing further gliding of the partial dislocation. Such interactions between superlattice stacking faults and the associated hardening effect in the L12 precipitates also lead to anomalous yield behavior in the Co-rich HEAs [65]. Besides flow anomalies, the high-temperature strength can also benefit from the elevated planar fault energies via carefully controlled alloying additions. Particle shearing requires the matrix dislocations to cut into the L12 precipitates, and associated planar faults will be generated under such circumstances, e.g., an anti-phase boundary (APB) for the sheared a/2 dislocation and superlattice intrinsic stacking fault (SISF) for the sheared a/3 dislocation [66, 67]. The onset of the particle shearing requires critical shear stress to be reached in order to overcome the energy barrier. According to first-principle calculations, it has been well demonstrated that the alloying additions of Nb, Ti, and Ta can profoundly increase the APB and SISF energies of the L12 precipitates to 530 and 286 mJ/m2 , which significantly increased the shearing resistance from the multi-component L12 particles to the shearing of the matrix dislocations [65]. Therefore, following attempts in designing HEAs for elevated applications should focus on tailoring the planar fault energies to enhance their high-temperature deformation resistance. Apart from the strength at elevated temperatures, the embrittlement issue should also be addressed at intermediate temperatures (roughly from 600 to 800 °C) among L12 -strengthened HEAs, which severely limits their potential engineering applications at elevated temperatures. Cao et al. [68] demonstrated that such embrittlement is associated with grain boundary stress concentrations and environmental attacks upon tension in an oxidizing atmosphere (Fig. 7a). Pronounced dislocations accumulated in the vicinity of grain boundaries to accommodate the thermally activated grain boundary sliding at intermediate temperatures, leading to severe stress concentrations along the grain boundaries. Moreover, the presence of deleterious brittle grain boundary precipitates with incoherent interfaces between matrix leads to large lattice mismatch and associated severe stress localization [42]. It should be noted that due to

2 Face-Centered Cubic High-Entropy Alloys

47

the complex chemical compositions and elemental interplays HEAs are more prone to have localized composition variations and diffusional heterogeneities, particularly at grain boundaries. As a result, distinctly different compositional and structural interfacial states from those at grain interiors can be stabilized at grain boundaries. For example, the enrichment of Ni, Al, and Ti elements was found at grain boundaries in the Ni–30Co–13Fe–15Cr–6Al–6Ti–0.1B (at.%) HEA [42], leading to the formation of grain boundary L21 -type Heusler phase. The brittle Heusler phase acts as a preferential site for crack initiation and propagation upon plastic deformations, which intensifies the stress concentration along grain boundaries and associated severe intergranular embrittlement (~4% at 700 °C). Through a duplex-aging heat treatment, it has been demonstrated that the precipitation of the intergranular L12 particles helps to pin the grain boundaries and retard their migration. In this way, the formation of the Heusler phase can be effectively suppressed, contributing to the brittle-to-ductile transition and the recovered tensile ductility to ~10% at 700 °C [42]. The polycrystalline HEAs with the duplex-aging heat treatment still suffered from the severe brittleness with massive intergranular facets on the fracture surface as the temperature further increased to 775 °C, which is associated with environmental-assisted embrittlement [68]. The adverse environment embrittlement has been well proven in the protective argon environment [68]. In strong contrast with the rapid tensile failure in the laboratory air, the protective testing environment helps to recover the tensile ductility to 15% at this temperature. An innovative heterogeneous grain structure has been proposed to conquer such environmental embrittlement by partial recrystallization and subsequent intergranular precipitation for grain boundary pinning [69]. By tailoring grainboundary characters and structures via thermo-mechanical treatments, the heterogeneous columnar-grained structures reduce the fraction of high-angle random grain boundaries and disrupt their connectivity. The crack propagation can be inhibited by such heterogeneous structure, resulting in a distinct intergranular-to-transgranular transition and associated tensile ductility recovery to ~18% at 800 °C [69] (Fig. 7b– d). Although the heterogeneous columnar-grained structure can effectively inhibit grain-boundary embrittlement, it should be noted that the high-density dislocations are stored in the heterogeneous columnar-grained structure, and therefore, it might not be thermodynamically stable, which tends to transform into the polycrystalline structure with prolonged annealing [70]. The approaches to enhance the thermal stability of such heterogeneous columnar-grain structure to defeat the intermediate temperature embrittlement is worthy of investigation in future study.

4 Summary Critical issues in developing coherent and incoherent precipitate-strengthened HEAs for elevated temperature applications have been briefly covered here. A stable γ/γ' dual-phase microstructure is preferred without other brittle intermetallic phase formation, and meanwhile, microstructural stability should be carefully evaluated. Computational-aided thermodynamic calculations accelerate the exploration of the

48

W. Liu and B. Cao

Fig. 7 a Stress–strain tensile curves at various deformation temperatures (in the laboratory air). Severe embrittlement can be found at intermediate temperatures [68]. b Intermediate temperature embrittlement can be eliminated by the heterogeneous columnar grain structure [69]. c Distinct brittle intergranular fracture to dimpled transgranular fracture by HCG structure. d Comparison of the grain character in equiaxed grain structure and HCG structure

target microstructures and phases, which is effective in screening out preferred alloy compositions. Various alloying additions have been identified to be beneficial for the stability of the L12 phase, expanding their temperature capabilities. Due to the sluggish diffusion effect, HEAs have shown exceptional resistance toward particle coarsening, which justified the motivation to pursue high-temperature structural materials within the scope of HEAs. HEAs have demonstrated outstanding high-temperature strength, including high-temperature hardness, high-temperature tensile strength, and creep resistance. However, tensile embrittlement has been found at intermediate temperatures, severely restricting their engineering applications. Such embrittlement is closely related to the grain-boundary stress concentrations and environmental-assist intergranular damage. A novel heterogeneous columnargrain structure has been developed, demonstrating exceptional resistance toward the intermediate temperature embrittlement.

2 Face-Centered Cubic High-Entropy Alloys

49

References 1. Cantor B, Chang I, Knight P, Vincent A (2004) Microstructural development in equiatomic multicomponent alloys. Mater Sci Eng A 375:213–218 2. Yeh JW, Chen SK, Lin SJ, Gan JY, Chin TS, Shun TT, Tsau CH, Chang SY (2004) Nanostructured high-entropy alloys with multiple principal elements: novel alloy design concepts and outcomes. Adv Funct Mater 6(5):299–303 3. Jien-Wei Y (2006) Recent progress in high entropy alloys. Ann Chim Sci Mat 31(6):633–648 4. Ding Q, Zhang Y, Chen X, Fu X, Chen D, Chen S, Gu L, Wei F, Bei H, Gao Y (2019) Tuning element distribution, structure and properties by composition in high-entropy alloys. Nature 574(7777):223–227 5. Gludovatz B, Hohenwarter A, Catoor D, Chang EH, George EP, Ritchie RO (2014) A fractureresistant high-entropy alloy for cryogenic applications. Science 345(6201):1153–1158 6. Gludovatz B, Hohenwarter A, Thurston KV, Bei H, Wu Z, George EP, Ritchie RO (2016) Exceptional damage-tolerance of a medium-entropy alloy CrCoNi at cryogenic temperatures. Nat Commun 7(1):1–8 7. Kumar NK, Li C, Leonard K, Bei H, Zinkle S (2016) Microstructural stability and mechanical behavior of FeNiMnCr high entropy alloy under ion irradiation. Acta Mater 113:230–244 8. Nygren K, Bertsch K, Wang S, Bei H, Nagao A, Robertson I (2018) Hydrogen embrittlement in compositionally complex FeNiCoCrMn FCC solid solution alloy. Curr Opin Solid State Mater Sci 22(1):1–7 9. Nygren KE, Wang S, Bertsch KM, Bei H, Nagao A, Robertson IM (2018) Hydrogen embrittlement of the equi-molar FeNiCoCr alloy. Acta Mater 157:218–227 10. Pu Z, Chen Y, Dai L (2018) Strong resistance to hydrogen embrittlement of high-entropy alloy. Mater Sci Eng A 736:156–166 11. Wu Z, Bei H, Pharr GM, George EP (2014) Temperature dependence of the mechanical properties of equiatomic solid solution alloys with face-centered cubic crystal structures. Acta Mater 81:428–441 12. Liu W, Yang T, Liu C (2018) Precipitation hardening in CoCrFeNi-based high entropy alloys. Mater Chem Phys 210:2–11 13. Liu W, Lu Z, He J, Luan J, Wang Z, Liu B, Liu Y, Chen M, Liu C (2016) Ductile CoCrFeNiMox high entropy alloys strengthened by hard intermetallic phases. Acta Mater 116:332–342 14. Pickering E, Muñoz-Moreno R, Stone H, Jones N (2016) Precipitation in the equiatomic highentropy alloy CrMnFeCoNi. Scr Mater 113:106–109 15. Shun T-T, Chang L-Y, Shiu M-H (2013) Age-hardening of the CoCrFeNiMo0. 85 high-entropy alloy. Mater Charact 81:92–96 16. Shun T-T, Hung C-H, Lee C-F (2010) The effects of secondary elemental Mo or Ti addition in Al0. 3CoCrFeNi high-entropy alloy on age hardening at 700 C. J Alloys Compd 495(1):55–58 17. He J, Wang H, Huang H, Xu X, Chen M, Wu Y, Liu X, Nieh T, An K, Lu Z (2016) A precipitationhardened high-entropy alloy with outstanding tensile properties. Acta Mater 102:187–196 18. Ming K, Bi X, Wang J (2017) Precipitation strengthening of ductile Cr15Fe20Co35Ni20Mo10 alloys. Scr Mater 137:88–93 19. Zhao Y, Yang T, Tong Y, Wang J, Luan J, Jiao Z, Chen D, Yang Y, Hu A, Liu C (2017) Heterogeneous precipitation behavior and stacking-fault-mediated deformation in a CoCrNibased medium-entropy alloy. Acta Mater 138:72–82 20. Li Z, Pradeep KG, Deng Y, Raabe D, Tasan CC (2016) Metastable high-entropy dual-phase alloys overcome the strength–ductility trade-off. Nature 534(7606):227–230 21. He J, Wang H, Wu Y, Liu X, Mao H, Nieh T, Lu Z (2016) Precipitation behavior and its effects on tensile properties of FeCoNiCr high-entropy alloys. Intermetallics 79:41–52 22. Liu W, He J, Huang H, Wang H, Lu Z, Liu C (2015) Effects of Nb additions on the microstructure and mechanical property of CoCrFeNi high-entropy alloys. Intermetallics 60:1–8 23. Otto F, Dlouhý A, Pradeep KG, Kubˇenová M, Raabe D, Eggeler G, George EP (2016) Decomposition of the single-phase high-entropy alloy CrMnFeCoNi after prolonged anneals at intermediate temperatures. Acta Mater 112:40–52

50

W. Liu and B. Cao

24. Ma D, Grabowski B, Körmann F, Neugebauer J, Raabe D (2015) Ab initio thermodynamics of the CoCrFeMnNi high entropy alloy: Importance of entropy contributions beyond the configurational one. Acta Mater 100:90–97 25. Zhang F, Wu Y, Lou H, Zeng Z, Prakapenka VB, Greenberg E, Ren Y, Yan J, Okasinski JS, Liu X (2017) Polymorphism in a high-entropy alloy. Nat Commun 8(1):1–7 26. Tracy CL, Park S, Rittman DR, Zinkle SJ, Bei H, Lang M, Ewing RC, Mao WL (2017) High pressure synthesis of a hexagonal close-packed phase of the high-entropy alloy CrMnFeCoNi. Nat Commun 8(1):1–6 27. Zhao Y, Lee D-H, Seok M-Y, Lee J-A, Phaniraj MP, Suh J-Y, Ha H-Y, Kim J-Y, Ramamurty U, Jang J-I (2017) Resistance of CoCrFeMnNi high-entropy alloy to gaseous hydrogen embrittlement. Scr Mater 135:54–58 28. Koyama M, Wang H, Verma VK, Tsuzaki K, Akiyama E (2020) Effects of Mn content and grain size on hydrogen embrittlement susceptibility of face-centered cubic high-entropy alloys. Metall Mater Trans A 51(11):5612–5616 29. Liu J, Guo X, Lin Q, He Z, An X, Li L, Liaw PK, Liao X, Yu L, Lin J (2019) Excellent ductility and serration feature of metastable CoCrFeNi high-entropy alloy at extremely low temperatures. Sci China Mater 62(6):853–863 30. Yi J, Zhuang X, He J, He M, Liu W, Wang S (2021) Effect of Mo doping on the gaseous hydrogen embrittlement of a CoCrNi medium-entropy alloy. Corros Sci 189:109628 31. Gangireddy S, Gwalani B, Soni V, Banerjee R, Mishra RS (2019) Contrasting mechanical behavior in precipitation hardenable AlXCoCrFeNi high entropy alloy microstructures: Single phase FCC vs. dual phase FCC-BCC. Mater Sci Eng A 739:158–166 32. Lu Y, Dong Y, Guo S, Jiang L, Kang H, Wang T, Wen B, Wang Z, Jie J, Cao Z (2014) A promising new class of high-temperature alloys: eutectic high-entropy alloys. Sci Rep 4(1):1–5 33. Fan R, Wang L, Zhao L, Wang L, Zhao S, Zhang Y, Cui B (2022) Synergistic effect of Nb and Mo alloying on the microstructure and mechanical properties of CoCrFeNi high entropy alloy. Mater Sci Eng A 829:142–153 34. Yang T, Zhao Y, Tong Y, Jiao Z, Wei J, Cai J, Han X, Chen D, Hu A, Kai J (2018) Multicomponent intermetallic nanoparticles and superb mechanical behaviors of complex alloys. Science 362(6417):933–937 35. Yang T, Zhao Y, Liu W, Kai J, Liu C (2018) L12-strengthened high-entropy alloys for advanced structural applications. J Mater Res 33(19):2983–2997 36. Cao B, Yang T, Liu W-H, Liu C (2019) Precipitation-hardened high-entropy alloys for hightemperature applications: a critical review. MRS Bull 44(11):854–859 37. Cao B, Zhao Y, Yang T, Liu CT (2021) L12-strengthened Co-rich alloys for high-temperature structural applications: a critical review. Adv Eng Mater 23(10):2100453 38. Cao B, Kong H, Ding Z, Wu S, Luan J, Jiao Z, Lu J, Liu C, Yang T (2021) A novel L12strengthened multicomponent Co-rich high-entropy alloy with both high γ' -solvus temperature and superior high-temperature strength. Scr Mater 199:113826 39. Suzuki A, Inui H, Pollock TM (2015) L12-strengthened cobalt-base superalloys. Annu Rev Mater Res 45:345–368 40. Gwalani B, Choudhuri D, Soni V, Ren Y, Styles M, Hwang J, Nam S, Ryu H, Hong SH, Banerjee R (2017) Cu assisted stabilization and nucleation of L12 precipitates in Al0. 3CuFeCrNi2 fcc-based high entropy alloy. Acta Mater 129:170–182 41. Gwalani B, Soni V, Choudhuri D, Lee M, Hwang J, Nam S, Ryu H, Hong SH, Banerjee R (2016) Stability of ordered L12 and B2 precipitates in face centered cubic based high entropy alloys-Al0. 3CoFeCrNi and Al0. 3CuFeCrNi2. Scr Mater 123:130–134 42. Yang T, Zhao Y, Fan L, Wei J, Luan J, Liu W, Wang C, Jiao Z, Kai J, Liu C (2020) Control of nanoscale precipitation and elimination of intermediate-temperature embrittlement in multicomponent high-entropy alloys. Acta Mater 189:47–59 43. Daoud H, Manzoni A, Wanderka N, Glatzel U (2015) High-temperature tensile strength of Al10Co25Cr8Fe15Ni36Ti6 compositionally complex alloy (high-entropy alloy). Jom 67(10):2271–2277

2 Face-Centered Cubic High-Entropy Alloys

51

44. Haas S, Manzoni AM, Krieg F, Glatzel U (2019) Microstructure and mechanical properties of precipitate strengthened high entropy alloy Al10Co25Cr8Fe15Ni36Ti6 with additions of hafnium and molybdenum. Entropy 21(2):169 45. Chang Y-J, Yeh A-C (2015) The evolution of microstructures and high temperature properties of AlxCo1. 5CrFeNi1. 5Tiy high entropy alloys. J Alloys Compd 653:379–385 46. Antonov S, Detrois M, Tin S (2018) Design of novel precipitate-strengthened Al–Co–Cr–Fe– Nb–Ni high-entropy superalloys. Metall Mater Trans A 49(1):305–320 47. Lass EA, Sauza DJ, Dunand DC, Seidman DN (2018) Multicomponent γ’-strengthened Cobased superalloys with increased solvus temperatures and reduced mass densities. Acta Mater 147:284–295 48. Sundman B, Lukas H, Fries S (2007) Computational thermodynamics: the Calphad method. Cambridge university press, Cambridge 49. Zhang C, Gao MC (2016) CALPHAD modeling of high-entropy alloys. Springer, pp 399–444 50. Zhao Y, Yang T, Han B, Luan J, Chen D, Kai W, Liu CT, Kai J-J (2019) Exceptional nanostructure stability and its origins in the CoCrNi-based precipitation-strengthened medium-entropy alloy. Mater Res Lett 7(4):152–158 51. Irisarri A, Urcola J, Fuentes M (1985) Kinetics of growth of γ' -precipitates in Ni–6· 75AI alloy. Mater Sci Technol 1(7):516–519 52. Davies C, Nash P, Stevens R (1980) Precipitation in Ni-Co-Al alloys. J Mater Sci 15(6):1521– 1532 53. Pandey P, Kashyap S, Palanisamy D, Sharma A, Chattopadhyay K (2019) On the high temperature coarsening kinetics of γ' precipitates in a high strength Co37. 6Ni35. 4Al9. 9Mo4. 9Cr5. 9Ta2. 8Ti3. 5 fcc-based high entropy alloy. Acta Mater 177:82–95 54. Sauza DJ, Dunand DC, Noebe RD, Seidman DN (2019) γ’-(L12) precipitate evolution during isothermal aging of a CoAlWNi superalloy. Acta Mater 164:654–662 55. Lapin J, Gebura M, Pelachová T, Nazmy M (2008) Coarsening kinetics of cuboidal γ’precipitates in single crystal nickel base superalloy CMSX-4. Kovove Mater 46:313–322 56. Tsai K-Y, Tsai M-H, Yeh J-W (2013) Sluggish diffusion in Co–Cr–Fe–Mn–Ni high-entropy alloys. Acta Mater 61(13):4887–4897 57. Cao B, Yang T, Fan L, Luan J, Jiao Z, Liu C (2020) Refractory alloying additions on the thermal stability and mechanical properties of high-entropy alloys. Mater Sci Eng A 797:140020 58. Neumeier S, Rehman H, Neuner J, Zenk C, Michel S, Schuwalow S, Rogal J, Drautz R, Göken M (2016) Diffusion of solutes in fcc Cobalt investigated by diffusion couples and first principles kinetic Monte Carlo. Acta Mater 106:304–312 59. Li X, Saunders N, Miodownik A (2002) The coarsening kinetics of γ' particles in nickel-based alloys. Metall Mater Trans A 33(11):3367–3373 60. Gelles D, Garner F, Brager H (1981) Effects of radiation on materials. In: 10th international symposium, ASTM STP, p 735e753 61. Tsao TK, Yeh AC, Kuo CM, Murakami H (2017) On the superior high temperature hardness of precipitation strengthened high entropy Ni-based alloys. Adv Eng Mater 19(1):1600475 62. Tsao T-K, Yeh A-C, Kuo C-M, Kakehi K, Murakami H, Yeh J-W, Jian S-R (2017) The high temperature tensile and creep behaviors of high entropy superalloy. Sci Rep 7(1):1–9 63. Yg N, Ohtomo A, Saiga Y (1976) Directional solidification of Rene 80. Trans Jpn Inst Met 17(6):323–329 64. Suzuki A, Pollock TM (2008) High-temperature strength and deformation of γ/γ' two-phase Co–Al–W-base alloys. Acta Mater 56(6):1288–1297 65. Cao B, Xu W-W, Yu C, Wu S, Kong H, Ding Z, Zhang T, Luan J, Xiao B, Jiao Z (2022) L12strengthened multicomponent Co–Al–Nb-based alloys with high strength and matrix-confined stacking-fault-mediated plasticity. Acta Mater 229:117763 66. Beauchamp P, Douin J, Veyssière P (1987) Dependence of the antiphase boundary energy upon orientation in the L12 structure. Philos Mag A 55(5):565–581 67. Milligan WW, Antolovich SD (1991) The mechanisms and temperature dependence of superlattice stacking fault formation in the single-crystal superalloy PWA 1480. Metall Trans A 22(10):2309–2318

52

W. Liu and B. Cao

68. Cao B, Wei D, Zhang X, Kong H, Zhao Y, Hou J, Luan J, Jiao Z, Liu Y, Yang T (2022) Intermediate temperature embrittlement in a precipitation-hardened high-entropy alloy: the role of heterogeneous strain distribution and environmentally assisted intergranular damage. Mater Today Phys 24:100653 69. Cao B, Kong H, Fan L, Luan J, Jiao Z, Kai J, Yang T, Liu C (2021) Heterogenous columnargrained high-entropy alloys produce exceptional resistance to intermediate-temperature intergranular embrittlement. Scr Mater 194:113622 70. Wu S, Wang G, Wang Q, Jia Y, Yi J, Zhai Q, Liu J, Sun B, Chu H, Shen J (2019) Enhancement of strength-ductility trade-off in a high-entropy alloy through a heterogeneous structure. Acta Mater 165:444–458

Chapter 3

Eutectic High-Entropy Alloys Wenna Jiao, Zhijun Wang, Sheng Guo, and Yiping Lu

1 Introduction Recently, eutectic high-entropy alloys [1] (EHEAs), as an important branch of highentropy alloys [2] (HEAs) or multi-principal-element alloys [3] (MPEAs), have attracted widespread attention in the metallic materials community due to their unique microstructures and attractive properties [4–8]. In the infancy of the development of HEAs, researchers were focused on exploring the multicomponent alloys of simple solid solutions, typically in the form of face-centered cubic (FCC), body-centered cubic (BCC), and hexagonal closed-packed (HCP) solid solutions. Generally, HEAs comprising single FCC phase process high ductility [9], while HEAs comprising single BCC phase exhibit high strength [10]. However, it is not easy to leverage the combined merits of both ductility and strength in single-phase HEAs. In addition, poor castability and compositional segregation dramatically limit the large-scale industrial applications of single-solid-solution forming HEAs. The concept of eutectic alloys could be effective to combine the advantages of the FCC- and BCC-structured HEAs and to remedy their weaknesses. For example, the eutectic alloys usually exhibit excellent combination of strength and ductility as in-situ formed composites containing soft and hard phases. Some eutectic alloys are more stable at high temperatures due to the low energy of the phase boundaries W. Jiao · Y. Lu (B) Engineering Research Center of High Entropy Alloy Materials (Liaoning Province), School of Materials Science and Engineering, Dalian University of Technology, Dalian, China e-mail: [email protected] Z. Wang State Key Laboratory of Solidification Processing, Northwestern Polytechnical University, Xi’an, China S. Guo Department of Industrial and Materials Science, Chalmers University of Technology, Gothenburg, Sweden © The Author(s), under exclusive license to Springer Nature Singapore Pte Ltd. 2022 Z. Jiao and T. Yang (eds.), Advanced Multicomponent Alloys, Materials Horizons: From Nature to Nanomaterials, https://doi.org/10.1007/978-981-19-4743-8_3

53

54

W. Jiao et al.

and exhibit high rupture strength and good high-temperature creep resistance. Moreover, the eutectic alloys can avoid solidification defects such as hot cracks, internal shrinkage, and compositional segregation [11, 12]. Following this line of thinking, Lu et al. [1, 13] firstly proposed to use the eutectic alloy concept to design HEAs (namely EHEAs), aiming to achieve good castability and excellent combination of strength and ductility in HEAs. They designed an EHEA of AlCoCrFeNi2.1 consisted of a soft FCC phase and a hard BCC phase, which achieved the desirable balance between strength and ductility in a wide temperature range temperature as well as the good seawater corrosion resistance. Since then, more and more researchers have dedicated to the researches on EHEAs, encouraged by the advantages of their structural and functional properties. The goal of this chapter is to systematically summarize the advances in EHEAs, including various approaches to locate the EHEA composition, their solidification microstructures, and mechanical properties of EHEAs.

2 Alloy Design for EHEAs EHEAs have been a hot topic over the past few years due to their excellent properties and great potential for industrial applications. So far, plenty of new EHEAs have been developed and many design strategies of EHEAs have been proposed. There includes the simple mixing method, the mixing enthalpy method, the method based on thermophysical parameters, the calculation of phase diagram (CALPHAD) method, and the machine learning method.

2.1 A Simple Mixing Method Based on the mixing enthalpy and constituent binary eutectics, Jiang et al. [14] proposed a simple mixing method to design EHEAs. It was inspired by the observation that most reported EHEAs contain two phases, and the elements in EHEAs can be divided into two groups according to mixing enthalpy, other chemical/physical properties and binary-phase diagrams: Elements in group A possess similar chemical activity and atomic radius, and their mixing enthalpy being close to zero; elements in group B not only have very negative mixing enthalpy with the elements in group A, but also have eutectic reactions with the elements in group A in binary systems. When the selected elements meet the above conditions, a dual-phase EHEA may be obtained. Following this line of thinking, they found that the atomic pairs of Co– Cr, Co–Ni, Co–Fe, Cr–Ni, Cr–Fe, and Fe–Ni are compatible with the condition for group A elements, and indeed the CoCrFeNi alloy forms a simple FCC solid solution, thus elements Co, Cr, Fe, Ni can be put into group A. The mixing enthalpies of

3 Eutectic High-Entropy Alloys

55

atomic pairs of Nb/Ta/Zr/Hf–Co, Nb/Ta/Zr/Hf–Cr, Nb/Ta/Zr/Hf–Fe, Nb/Ta/Zr/Hf– Ni are quite negative, and all of them undergo a eutectic reaction in binary systems, so elements Nb, Ta, Zr, Hf can be assigned to group B. In order to locate the eutectic point accurately, a simple mixing method based on calculating the eutectic point from binary alloys is presented as follows: (1/4)Ni84.5 Nb15.5 + (1/4)Co86.1 Nb13.9 + (1/4)Cr88 Nb12 + (1/4)Fe89.4 Nb10.6 = Co21.525 Cr22 Fe22.35 Ni21.125 Nb13 = (CoCrFeNi)87 Nb13 ≈ CoCrFeNiNb0.6 (1) (1/4)Ni86.3 Ta13.7 + (1/4)Co92 Ta8 + (1/4)Cr87 Ta13 + (1/4)Fe92.5 Ta7.5 = Co23 Cr21.75 Fe23.125 Ni21.575 Ta10.55 = (CoCrFeNi)89.45 Ta10.55 ≈ CoCrFeNiTa0.47 (2) (1/4)Ni91.2 Zr8.8 + (1/4)Co90.5 Zr9.5 + (1/4)Cr82.8 Zr17.2 + (1/4)Fe90.2 Zr9.8 = Co22.625 Cr20.7 Fe22.55 Ni22.8 Zr11.325 = (CoCrFeNi)88.675 Zr11.325 ≈ CoCrFeNiZr0.51 (3) (1/4)Ni87.5 Hf12.5 + (1/4)Co89 Hf11 + (1/4)Cr87 Hf13 + (1/4)Fe92.1 Hf7.9 = Co22.25 Cr21.75 Fe23.025 Ni21.875 Hf11.1 = (CoCrFeNi)88.9 Hf11.1 ≈ CoCrFeNiHf0.49 (4) According to the above equations, four alloy compositions close to EHEAs can be obtained. The exact EHEAs, CoCrFeNiNb0.45 , CoCrFeNiTa0.4 , CoCrFeNiZr0.55 , and CoCrFeNiHf0.4 can be obtained by adjusting elements Nb/Ta/Zr/Hf slightly, as shown in Fig. 1. In the method proposed by Jiang et al., it is obvious that elements in group A are in equiatomic ratio, which greatly reduces the cocktail effect of HEAs [15]. For this reason, Xie et al. [16, 17] modified the simple mixing method. They set (CoCrFeNi)x as one group while another group is (Coa Crb Fec Nid Me ), in which M denotes Nb/Ta/Zr/Hf and Coa Crb Fec Nid Me is a Laves phase (a, b, c, d are randomly selected around the chemical compositions of the Laves phase). In binary-phase diagrams, the corresponding contents of M element in eutectic reactions with Co, Cr, Fe, Ni were defined as g1 , g2 , g3 , g4 . For the latter calculation, e can be regarded as the sum of g1 , g2 , g3 , g4 . As a result, an alloy of Coa+x Crb+x Fec+x Nid+x Me can be obtained. Based on the simple mixing method, the value of e can be determined, and in binary systems, the eutectic point of Co/Cr/Fe/Ni–M is also determined, so when x is known as an exact value such as h, the compositions of EHEAs can be calculated by solving equations. The schematic explanation of the modified simple mixture method is shown in Fig. 2. EHEAs with non-equal molar ratio of Co0.86 Cr0.67 Fe0.76 Ni0.91 Hf0.4 , Co1.01 Cr0.82 Fe0.91 Ni1.06 Ta0.4 and Co0.81 Cr0.62 Fe0.71 Ni0.86 Nb0.4 were thus designed by the modified simple mixing method.

56

W. Jiao et al.

Fig. 1 Microstructures of EHEAs obtained by the simple mixing method: a CoCrFeNiNb0.45 , b CoCrFeNiTa0.4 , c CoCrFeNiZr0.55 , d CoCrFeNiHf0.4 [14]

Fig. 2 Schematic explanation of the modified simple mixture method [16]

2.2 A Mixing Enthalpy Method In the typical AlCoCrFeNi2.1 EHEA, the mixing enthalpy of atomic pair Al–Ni is far more negative than that of Co–Cr–Fe, which is −22 kJ/mol and close to zero, respectively. Lu et al. [18] divided the elements into two groups according to the mixing enthalpy; i.e., one is composed of Al and Ni, and the other is composed of Co, Cr, and Fe. The eutectic phases in the AlCoCrFeNi2.1 EHEA, so the B2 phase

3 Eutectic High-Entropy Alloys

57

and the FCC phase, are enriched in elements belonging to the above-mentioned two groups, respectively. Along this line of thinking, Lu et al. [18] firstly used the mixing enthalpy method to locate new EHEAs based on AlCoCrFeNi2.1 by replacing Al with other elements that have negative mixing enthalpy with Ni. They found that the mixing enthalpy of Ni–Zr, Ni–Nb, Ni–Hf, and Ni–Ta is −49 kJ/mol, −30 kJ/mol, − 42 kJ/mol, and −29 kJ/mol, respectively, which is similar to NiAl. Therefore, the four elements can be used as substitutes for Al. Their relative contents in the target EHEAs could be determined by establishing a relationship between their mixing enthalpy with Ni and the content of element Al in AlCoCrFeNi2.1 . Using the equations that are given as follows: x 22 Nb = =− , x = 0.73 Al 1 −30

(5)

x 22 Ta = =− , x = 0.76 Al 1 −29

(6)

x 22 Zr = =− , x = 0.45 Al 1 −49

(7)

x 22 Hf = =− , x = 0.52 Al 1 −42

(8)

Consequently, four alloy compositions close to EHEAs, Zr0.45 CoCrFeNi2.1 , Nb0.73 CoCrFeNi2.1 , Hf0.52 CoCrFeNi2.1 and Ta0.76 CoCrFeNi2.1 were designed. Using simple trial and error experiments to tune the compositions, a series of exact EHEAs, Zr0.6 CoCrFeNi2.0 , Nb0.74 CoCrFeNi2.0 , Hf0.55 CoCrFeNi2.0 , and Ta0.65 CoCrFeNi2.0 can be obtained, and their microstructures are displayed in Fig. 3.

2.3 A Method Based on Thermo-Physical Parameters Till date, several thermo-physical parameters have been proposed to predict the phase formation in HEAs, e.g., mixing enthalpy (ΔH mix ), entropy of mixing (ΔS mix ), atomic size difference (δ r ), valence electron concentration (VEC), and Ω, which are defined as follows [19–21]: ΔH mix =

n 

ij

4ΔHmix ci c j

(9)

i=1,i/= j

ΔS mix = −R

n  i=0

ci lnci

(10)

58

W. Jiao et al.

Fig. 3 Microstructures of EHEAs obtained by the mixing entropy method: a CoCrFeNi2.0 Zr0.6 , b CoCrFeNi2.0 Nb0.74 , c CoCrFeNi2.0 Hf0.55 , d CoCrFeNi2.0 Ta0.65 [18]

⌈ | n | 2 δr = | ci (1 − ri /r¯)

(11)

i=1

VEC =



ci VECi

Ω = Tm ΔSmix /|ΔHmix |

(12) (13)

where ci and cj are the molar fraction of the ith and jth element, respectively, n is ij the number of elements in an alloy system, ΔHmix is the mixing enthalpy of binary alloys, R is the gas constant, r i is the atomic radius of the ith element, r is the average atomic radius, and VECi is the valence electron concentration of ith element. When predicting the formation of solid solutions in HEAs, the parameters of Ω ≥ 1.1 and δ r ≤ 6.6% are proved to be quite effective, while VEC is a good parameter to distinguish the formation of FCC and BCC solid solutions. Essentially, elements with a higher VEC tend to stabilize the FCC phase while elements with a lower VEC tend to stabilize the BCC phase, and the threshold values of 8.0 and 6.87 can be used as a quite valid guidance. Thus, the question is now that, for exact EHEAs, whether the above-mentioned parameters will fall in a certain region. Chanda et al. [22] investigated the relationship between the mixing enthalpy (ΔH mix ) and atomic size difference (δ r ) as well as the relationship between the atomic size difference (δ r ) and

3 Eutectic High-Entropy Alloys

59

Fig. 4 Relationship between a the ΔHmix and δ r and b VEC and δ r for the single-phase FCC, single-phase BCC and eutectic/near-eutectic HEAs. Note Stars represent single-phase FCC HEAs; squares represent single-phase BCC HEAs; and circles represent EHEAs [22]

VEC for the single-phase FCC, single-phase BCC, and eutectic/near-eutectic HEAs, as shown in Fig. 4. They proposed that at conditions −18 kJ/mol ≤ ΔH mix ≤ − 6 kJ/mol, 6 ≤ VEC ≤ 8.5, and δ r > 3%, the eutectic phases become stable. In the literature, most EHEAs are composed of solid solution/solid solution (FCC/BCC, BCC/HCP), solid solution/intermetallic compound (IMC), IMC/IMC, and only EHEAs containing a soft phase and a hard phase exhibit superior mechanical properties. Based on that, Jin et al. [23] proposed a pseudo-binary strategy to design EHEAs using parameters VEC and ΔH mix . Their core idea is to find a stable solid solution and a stable IMC according to the phase formation rules, to form the two components as in traditional binary alloy phase diagrams, and then to tune the content of the two pseudo-phases to locate the eutectic point. To verify their strategy, three HEAs, CoCrFeNi2 , Co2 CrFeNi, and CoCrFe2 Ni were selected as the stable FCC phase, and NiAl was chosen as the stable IMC. In this way, they successfully prepared three exact EHEAs, Al17 Co14.3 Cr14.3 Fe14.3 Ni40.1 , Al17 Co28.6 Cr14.3 Fe14.3 Ni25.8 and Al17 Co14.3 Cr14.3 Fe28.6 Ni25.8 . All three EHEAs exhibited excellent tensile mechanical properties, with a tensile fracture strength higher than 1000 MPa and a total elongation over 10%. The relationship between ΔH mix and VEC of most known FCC/IMC-structured EHEAs was summarized in Fig. 5. It is obvious that at conditions −5 kJ/mol < ΔH mix < 0 kJ/mol and VEC ≥ 8, EHEAs could be acquired [23]. In addition, inspired by the solidification behavior of the eutectic alloys, He et al. [24] separated the components of HEAs into two groups according to the mixing enthalpy. Besides simply considering the value of mixing enthalpy between elements, they also considered the standard deviation of mixing enthalpy. By calculating the standard deviation of mixing enthalpy of a large number of single-phase HEAs and EHEAs, they found that the alloys exhibit eutectic structure when the standard deviation is greater than 5, as shown in Fig. 6. This grouping strategy was supported by the discovery of a few new EHEAs, Ni2 CoCrFeNb0.74 , CoCrFeNiTa0.43 , Ni2 CoCrFeHf0.55 , and CoCrFeNiZr0.5 .

60

W. Jiao et al.

Fig. 5 Relationship between ΔH mix and VEC of some FCC/IMC-structured EHEAs [23]

Fig. 6 Standard deviation of mixing enthalpies of different pairs in single-phase HEAs and EHEAs [24]

2.4 A Semi-quantitative Method Based on CALPHAD Phase diagrams are the key to understanding the alloys. In binary and ternary phase diagrams, the eutectic compositions and reaction temperatures can be easily acquired. However, due to the large number of elements in HEAs, the establishment of phase diagrams is quite complicated. In this regard, CALPHAD method becomes an alternative way to locate the eutectic point. He et al. [25] firstly calculated the pseudo-binary CoCrFeNi–Nb phase diagram using the Thermo-Calc software, as shown in Fig. 7. In the pseudo-binary-phase diagram, CoCrFeNi alloy is assumed as one pseudo-element; therefore, the compositional space of the EHEAs is quickly narrowed down. It can be seen in Fig. 7, when the content of Nb reaches 10 at.%,

3 Eutectic High-Entropy Alloys

61

Fig. 7 Calculated pseudo-binary-phase diagram for CoCrFeNiNbx alloys [25]

the alloy will show the eutectic structure. Otherwise, the hypoeutectic structure and hypereutectic structure will be stable when the Nb content is lower or higher than 10 at.%. In order to verify the reliability of the calculated phase diagram, a series of CoCrFeNiNbx alloys were prepared, and it was experimentally proved that the eutectic composition lies at Nb = 14 at.% (the position marked by the star in Fig. 7), which is very close to the eutectic point in the calculated phase diagram. Therefore, the CALPHAD method predicted the eutectic compositions effectively. Accordingly, many EHEAs can be obtained using this CALPHAD-assisted design approach. For example, Chen et al. [26] located the eutectic composition in CoCrFeNiAlx and AlCoCrFeNiNbx alloys, and Wu et al. [27] designed a Al17.4 Co21.7 Cr21.7 Ni39.2 EHEA by calculating pseudo-binary and pseudo-ternary phase diagrams.

2.5 A Machine Learning Method As described above, based on thermodynamics, element grouping strategies, and the CALPHAD method, several design approaches for EHEAs have been proposed. As a result, FCC/Laves, FCC/BCC(B2) EHEAs have been successfully designed, in which FCC phases are usually enriched in Co, Cr, Fe, and Ni, while BCC(B2) phases are usually enriched in Ni and Al. However, it remains a great challenge to locate the eutectic points in the huge compositional space that the concept of HEAs can offer. Wu et al. [28] used the machine learning (ML) method to locate the eutectic point in the AlCoCrFeNi high-entropy system, in which sufficient data

62

W. Jiao et al.

is available in the literature and ML can hence efficiently predict material characteristics including compositions, microstructures, physical and chemical properties by constructing models and discovering correlations among input elements. In the Al–Co–Cr–Fe–Ni system, a database containing the alloy compositions and phase constitutions of 321 alloys was built up using the literature results and CALPHAD calculations. The artificial neural network model was then trained to predict a large number of near-eutectic compositions. By analyzing the predicted compositions, the association between different elements and eutectic compositions was established, as shown in Fig. 8. It was found that Al is the critical element for the eutectic compositions, the content of which should be between 15 and 20 at.%. Cr is a strongly associated element with Al; i.e., for eutectic compositions, there is a corresponding content range for Cr with the fixing Al content. As for the miscible Ni, Co, and Fe elements, their contents can be estimated using the average VEC. This research not only located the eutectic compositions in the Al–Co– Cr–Fe–Ni system, but also provided a generic three-step design approach to locate EHEAs in multicomponent systems: (1) divide the elements into critical elements, strongly associated elements, and miscible elements by machine learning; (2) select the combinations of critical and strongly associated elements; and (3) determine the contents of miscible elements. In the future, the machine learning method is expected to be a promising method to develop abundant EHEAs with the further development of material databases.

Fig. 8 Statistics on the 400 predicted near-eutectic compositions. a Content distributions of every single element. b Maps showing the contents of Co, Cr, Fe, Ni with a reference to the Al content [28]

3 Eutectic High-Entropy Alloys

63

3 Solidification Microstructures The morphology, grain size, and volume fraction of eutectic phases play important roles in determining the mechanical, physical, and chemical properties of the EHEAs. In conventional eutectic alloys, many attempts have been made to classify the solidification microstructures (e.g., lamellar, rod-like, fibrous, broken lamellar, flakes, complex regular, quasi-regular), which is closely related to the materials characteristics such as the faceted/non-faceted nature, entropy of fusion and the phase fraction, as well as the solidification conditions [29]. With the continuous development of EHEAs, some new eutectic morphologies have been reported. Table 1 gives a list of reported EHEAs with fully eutectic microstructure exhibiting different morphologies. It can be seen that besides the commonly seen lamellar and rod-like morphologies in the traditional eutectic alloys, sunflower-like, chrysanthemum-like [30], and seaweed [31] eutectic morphologies are found in EHEAs. According to the variation of the microstructure, EHEAs can be classified into two categories based on the morphology: lamellar morphology and anomalous morphology.

3.1 Lamellar Eutectic Morphology In traditional eutectic alloys, lamellar- or rod-like structure will evolve when both eutectic phases exhibit a low entropy of fusion, while the irregular structures will form when one phase has a high entropy of fusion and the other one has a low entropy of fusion. In EHEAs, the theory of eutectic solidification is limited because of the large number and high content of constituents. Most EHEAs show a lamellar structure, but different from traditional eutectic alloys, the lamellar structure in EHEAs is generally fine. For example, the interlamellar spacing of an industrial-scale AlCoCrFeNi2.1 [1] EHEA is about 2 μm; He et al. [36] reported a CoCrFeNiNb0.65 EHEA with the lamellar spacing of ~0.25 μm at the as-cast condition, and after annealing the sample at 600 °C and 750 °C, the lamellar spacing of the alloy became 0.22 μm and 0.44 μm, respectively; Jiang et al. [76] reported a CoFeNi2 V0.5 Nb0.75 EHEA with the lamellar spacing of ~0.25 μm, and after annealing at 600 °C, 700 °C and 800 °C, the lamellar spacing decreased to 0.2 μm, 0.16 μm, and 0.18 μm, respectively; Jiang et al. [43] reported a CoCrFeNiTa0.4 EHEA with the lamellar spacing of 0.15–0.20 μm; Liu et al. [45] reported an Al19.3 Co15 Cr15 Ni50.7 EHEA with the lamellar spacing of ~3 μm; Jiang et al. [77] reported a bulk CoCrFeNiNb0.45 EHEA with the lamellar spacing of ~0.20 μm; Jin et al. [49] reported a CrFeNi2.2 Al0.8 EHEA with the lamellae width of FCC and B2 phase of ~1 μm and 0.40 μm, respectively. The EHEAs mentioned above are directly cast by vacuum arc melting or vacuum induction melting, and such an ultrafine solidification structure is rarely seen in directly cast eutectic alloys, especially in industrial-scale ingots. The refinement of the interlamellar spacing in EHEAs may be resulted from the increasing resistance to form appropriate crystalline phases and hence to maintain the cooperative eutectic growth, due to the presence

64

W. Jiao et al.

Table 1 List of reported EHEAs with fully eutectic microstructure exhibiting different morphologies Alloys

Phases

Morphology

Ref.

AlCrFeNiMo0.2

FCC + B2

Lamellar eutectic

[32]

AlCrFeNi

BCC + FCC

Lamellar eutectic

[32]

Al1.2 CrCuFeNi2

FCC + B2

Rod-like eutectic

[30]

Al2.2 CrCuFeNi2

BCC + B2

Sunflower-like eutectic

[30]

Al2.5 CrCuFeNi2

BCC + B2

Chrysanthemum-like

[30]

AlCoCrFeNi2.1

FCC + B2

Lamellar eutectic

[1]

AlCoCrFeNiNbx

BCC + HCP

Lamellar eutectic

[33]

CoFeNiVMo0.6

FCC + IM

Lamellar eutectic

[34]

CoFeNi2 V0.5 Nb0.75

FCC + Laves

Lamellar eutectic

[35]

CoCrFeNiNb0.65

FCC + Laves

Lamellar eutectic

[36]

Alx CrFeNi

BCC + B2

Plate-like eutectic

[37]

CoCrFeNiZr0.5

FCC + Laves

Lamellar eutectic

[38]

CoCrFeNiNb0.45

FCC + Laves

Lamellar eutectic

[39]

CoCrFeNiMnPd

FCC + Mn7 Pd9 Seaweed eutectic

[31]

CoCrFeNiNb0.5

FCC + Laves

[40]

Lamellar eutectic

CoCrFeNiMo0.8

FCC + IM

Lamellar eutectic

[41]

CoCrFeNiTa0.395

FCC + Laves

Lamellar eutectic

[42]

CoCrFeNiTa0.4

FCC + Laves

Lamellar eutectic

[43]

Fe20 Co20 Ni41 Al19

B2 + L12

Lamellar eutectic

[44]

Al19.3 Co15 Cr15 Ni50.7

FCC + B2

Lamellar eutectic

[45]

(Co40 Cr10 Fe5 Mo5 Ni40 )82.2 Al17.8

FCC + B2

Lamellar eutectic

[46]

Al0.7 CoCrFeNi

FCC + B2

Lamellar eutectic

[47]

Co27.6 Cr22.6 Fe19.6 Ni23.4 Ta3.4 Zr2.5 Al0.9

FCC + Laves

Lamellar eutectic

[48]

CrFeNi2.2 Al0.8

FCC + B2

Lamellar eutectic

[49]

10V15Cr5Mn10Co2515Ni25.3Fe9.7Nb

FCC + Laves

Lamellar eutectic

[50]

Ni30 Co30 Cr10 Fe10 Al18 W2

FCC + B2

Lamellar eutectic

[51]

CoCu0.5 FeNiTa0.5

FCC + Laves

Lamellar eutectic

[52]

AlCox CrFeNi3-x

FCC + B2

Lamellar eutectic

[53]

Al1Co25 Cr18 Fe23 Ni23 Ta10

FCC + Laves

Lamellar eutectic

[54]

Co25.1 Cr18.8 Fe23.3 Ni22.6 Ta8.5 Al1.7

FCC + Laves

Lamellar eutectic

[55]

Fe32.5-x Co10 Ni25 Cr15 Mn5 V10 Al2.5 Nbx

FCC + Laves

Lamellar eutectic

[56]

Al0.9 CoCrNi2.1

L12 + B2

Lamellar eutectic

[57]

(CoFe2 NiV0.5 Mo0.2 )91 Nb9

FCC + Laves

Lamellar eutectic

[58]

Al0.8 CoCr0.6 Fe0.7 Ni1.5

FCC + B2

Lamellar eutectic

[59]

(CoCuFeNi)81 Mo19

FCC + μ

Lamellar eutectic

[60] (continued)

3 Eutectic High-Entropy Alloys

65

Table 1 (continued) Alloys

Phases

Morphology

Ref.

Fe20 Co20 Ni41 Al19

FCC + B2

Lamellar eutectic

[61]

Nb0.74 CoCrFeNi2.1

FCC + Laves

Lamellar eutectic

[62]

MoNbRe0.5 TaW(TiC)1.0

BCC + FCC

–

[63]

Fe35 Ni25 Cr25 Mo15

FCC + σ

Lamellar eutectic

[64]

Al0.9 CoFeNi2

L12 + B2

Lamellar eutectic

[65]

Fe2 Ni2 CrMo1.25

FCC + σ

Lamellar eutectic

[66]

CrFeNi2 Nb0.6

FCC + Laves

Lamellar eutectic

[67]

Fe2 NiCrNb0.34

FCC + Laves

Lamellar eutectic

[68]

AlCrFeNiCo1.9

FCC + B2

Lamellar eutectic

[69]

AlCr1.3 TiNi2

L21 + B2

Lamellar eutectic

[70]

Re0.5 MoNbW(TaC)0.8

BCC + FCC

Lamellar eutectic

[71]

Zr0.6 CoCrFeNi2.0

FCC + Ni7 Zr2

Lamellar eutectic

[18]

Nb0.74 CoCrFeNi2.0

FCC + (Co, Ni)2 Nb

Lamellar eutectic

[18]

Hf0.55 CoCrFeNi2.0

FCC + Ni7 Hf2

Lamellar eutectic

[18]

Ta0.65 CoCrFeNi2.0

FCC + (Co, Ni)2 Ta

Lamellar eutectic

[18]

CoCrFeNiZr0.55

FCC + Laves

Lamellar eutectic

[14]

CoCrFeNiHf0.4

FCC + Laves

Lamellar eutectic

[14]

Al17 Co14.3 Cr14.3 Fe14.3 Ni40.1

L12 + B2

Lamellar eutectic

[23]

Al17 Co28.6 Cr14.3 Fe14.3 Ni25.8

FCC + B2

Lamellar eutectic

[23]

Al17 Co14.3 Cr14.3 Fe28.6 Ni25.8

FCC + B2

Lamellar eutectic

[23]

Co0.96 Cr0.76 Fe0.85 Ni1.01 Hf0.40

FCC + Laves

Lamellar eutectic

[17]

Co0.99 Cr0.70 Fe0.94 Ni1.07 Hf0.40

FCC + Laves

Lamellar eutectic

[17]

Al10 Co18 Cr18 Fe18 Nb10 Ni26

FCC + B2 + Laves

Lamellar eutectic

[72]

(CoFe2 NiV0.5 Mo0.2 )91 Zr9

FCC + Laves

Lamellar eutectic

[73]

NiAl–20V–17Cr

BCC + B2

Lamellar eutectic

[74]

NiAl–30V–5Mo

BCC + B2

Lamellar eutectic

[74]

NiAl–30Cr–Mo

BCC + B2

Lamellar eutectic

[74]

NiAl–20V–10Cr–5Mo

BCC + B2

Lamellar eutectic

[74]

AlSi6 Cu4 Fe2

(Al) + Si + β + Rod-like eutectic Al2 Cu

[75]

66

W. Jiao et al.

of many different atomic species in the melt [5]. The refinement of the lamellar spacing, as well as the significant solid solution hardening and lattice distortion, render EHEAs to potentially show superior mechanical properties.

3.2 Anomalous Eutectic Morphology Anomalous eutectic morphology is rare in EHEAs. The solidification mechanism is also an important issue to be resolved. In the sunflower-like microstructure seen in the Al2.2 CrCuFeNi2 alloy, as shown in Fig. 9, the main phases are disordered BCC phase (A2) and ordered BCC phase (B2). A B2 phase formed primarily as the disk floret, and then grew alternatively with an A2 phase on the disk floret, with a radiating eutectic structure. Finally, sunflower-like morphology appeared. When the size of the primary phase is small or the shape deviates from the spheroid (ideal isotropic growth), a chrysanthemum-like or a distorted sunflower-like microstructure appeared. In order to further understand the formation mechanism of the solidification microstructure, Guo et al. and Borkar et al. investigated the solidification/decomposition pathway of the as-cast Al2 CuCrFeNi2 alloy. Guo et al. [78] believed that during solidification, the B2 phase first solidified from the liquid phase as the primary phase, forming the disk floret, as shown in Fig. 10. After that, the alternating B2 and A2 phases solidified from the liquid phase as the lamellae eutectic structure forming the petals of the sunflower. As the temperature continued to decrease, the primary B2 phase in disk floret decomposed to form plenty of discontinuous particles of A2 phase and the matrix B2 phases (α' and β' in Fig. 10) due to the spinodal decomposition. Borkar et al. [79] believed that the eutectic phases solidified from the parent liquid directly, forming sunflower-like structure with ordered BCC phases (B2e ) and disordered BCC phases (βe ) at the eutectic temperature (Fig. 11b). Subsequently, eutectic cells/colonies formulated when the solidification ended, or they collided with neighboring cells/colonies (Fig. 11c). When the temperature is low, more complex multi-scale decomposition occurs within the two phases.

4 Mechanical Properties The materials science usually consists of four components, including (1) compositions and structures, (2) processing, (3) properties, and (4) performance [15]. In this part, the mechanical properties of EHEAs are described as follows.

3 Eutectic High-Entropy Alloys

67

Fig. 9 a–c Sunflower-like and chrysanthemum-like microstructures in Al2.2 CrCuFeNi2 ; d–e seaweed microstructures in CoCrFeNiMnPd [30, 31]

Fig. 10 Sketch of the formation mechanism of the sunflower-like structure. (L: liquid phase, α: B2 phase, β: A2 phase, α' and β' : B2 and A2 phase out of the spinodal decomposition, respectively) [78]

68

W. Jiao et al.

Fig. 11 Schematic representation of phase evolution during solidification, forming the sunflowerlike structure [79]

4.1 Mechanical Properties of As-Cast EHEAs In binary and ternary eutectic alloy systems, cast iron, Al–Si, Sn–Ag–Cu, etc., are widely used in various industries such as electronics, automotive, aerospace, and many emerging industries. It is well known that eutectic alloys are the most widely used cast alloys in industries. There are several reasons for that. First, the eutectic structure forms at temperatures as high as their reaction temperature, which provides the eutectic alloys with excellent high-temperature performance. Second, as a natural or in-situ composite, the fine eutectic microstructure usually shows good properties. Last but not least, eutectic alloys natively have good castability or fluidity, which alleviates the segregation and shrinkage induced cavity. EHEAs inherit the characteristics

3 Eutectic High-Entropy Alloys

69

of both HEAs and eutectic alloys and show excellent mechanical properties. Here, the discussions of mechanical properties mainly focus on compressive properties and tensile properties of EHEAs. The strength of eutectic alloys primarily relates to (i) morphology of the eutectic microstructure, (ii) crystal structure and volume fraction of the constituent phases and their properties and (iii) the interlamellar spacing λe . Thereinto, λe plays an important role, and the Hall–Petch equation can be used to quantitatively express the relationship between strength and lamellar spacing: σ y = σ 0 + kλe −1/2 , where σ y is the yield strength of the material, σ 0 is the friction stress, λe is the interlamellar spacing. Therefore, the mechanical properties of the EHEAs can be controlled effectively by reducing the lamellar spacing and adjusting the phase composition and volume fraction of the phases. This section mainly summarizes the mechanical properties of EHEAs obtained by direct casting. So far, the casting methods of EHEAs mainly include arc melting, induction melting, magnetic levitation melting, and directional solidification, in which arc melting is the most commonly used technique. Generally, two types of ingots can be achieved by arc smelting; one is the button-like ingot by mixing all the constituent elements in a copper crucible, and the other is cylindrical or cuboid ingot by the method of copper mold suction casting. Induction melting and magnetic levitation melting techniques are usually used to prepare larger-scale ingots, so they are more suitable for industrial productions. Directional solidification process can be effectively used to adjust the microstructure and property of EHEAs by tuning processing parameters. The most classic EHEA composition, AlCoCrFeNi2.1 , of 2.5 kg in weight, was first prepared using vacuum induction melting by Lu et al. [1]. The bulk alloy exhibited an excellent castability with few casting defects, significantly differentiating with other bulk HEAs, as shown in Fig. 12a. The cast microstructure of the AlCoCrFeNi2.1 EHEA showed a uniform and fine lamellar morphology with the interlamellar spacing of ~2 μm as shown in Fig. 12b. The EHEA showed an unprecedented combination of high tensile ductility (elongation of 25.6%) and high fracture strength (944 MPa) at ambient temperature [1]. In addition, considering that the actual composition of EHEAs could potentially deviate from the eutectic composition in the practical large-scale casting, three industrial-scale HEAs ingots near the eutectic composition (hypoeutectic alloy AlCoCrFeNi2.0 , eutectic alloy AlCoCrFeNi2.1 and hypereutectic alloy AlCoCrFeNi2.2 ) were prepared [13]. The results showed that both the hypoeutectic alloy and hypereutectic alloy displayed fairly uniform and fine lamella (the interlamellar spacing of 1–3 μm), which is similar to the exact EHEA AlCoCrFeNi2.1 . Figure 13 depicts the engineering tensile stress–strain curves of the three as-cast HEAs, and the detailed mechanical properties of casting bulk alloys are summarized in Table 2. All the near-eutectic HEAs exhibited high strength and high ductility at both ambient and cryogenic temperatures. Subsequently, several bulk EHEAs were direct casted using induction melting [53, 65, 77, 80, 81], and all of them showed excellent mechanical properties. The CoCrFeNiNb0.45 with ultrafinelamellar microstructure (lamellar spacing ~200 nm) exhibited a high yield stress of

70

W. Jiao et al.

1475 MPa, fracture stress of 2558 MPa, and fracture strain of 27.9% from compressive tests [77]. The Al0.9 CoFeNi2 EHEA with the interlamellar spacing of ~1–2 μm showed superior strength of 1005 MPa and ductility of 6.2% in tension at room temperature [65]. The AlCox CrFeNi3-x (x = 0, 0.2, 0.4, 0.6, 0.8, and 1.0) EHEAs with the interlamellar spacing of ~2–10 μm exhibited excellent compressive properties, with the fracture strength being higher than 2000 MPa and the compressive strain exceeding 35% [53]. Arc melting is usually used to produce small-scale HEAs. The microstructure and properties of AlCoCrFeNi2.1 EHEA produced by arc melting have been reported in various reports. Using the method of suction casting, Wani et al. [82] prepared a cuboid-like AlCoCrFeNi2.1 EHEA of 15 mm (width) × 90 mm (length) × 3 mm (thickness) to investigate its microstructure and mechanical properties. The average thickness of L12 and B2 phases was 0.57 μm and 0.20 μm, respectively, which was finer than that of large-scale ingots, as shown in Fig. 14. The tensile test at room temperature showed that the yield strength, ultimate tensile strength, and elongation to failure of the alloy were 620 MPa, 1050 MPa, and 17%, respectively. When the testing temperature decreased to −90 °C, −150 °C, −196 °C, the yield strength, ultimate tensile strength and elongation to failure of the alloy were 675 MPa, 1183 MPa, 18.3%, 697 MPa, 1319 MPa, 19.5%, and 857 MPa, 1461 MPa, 16.6%, respectively

Fig. 12 a Bulk AlCoCrFeNi2.1 EHEA ingot, b a microscopy image showing the eutectic microstructure of the AlCoCrFeNi2.1 alloy [1]

Fig. 13 Engineering tensile stress–strain curves of AlCoCrFeNix (x = 2.0, 2.1 and 2.2) alloys, a at room temperature (RT), b at −70 °C, c at −196 °C [13]

3 Eutectic High-Entropy Alloys

71

Table 2 Yield strength (σ y ), fracture strength (σ UTS ), and elongation (EL%) to fracture of Ni2.0 , Ni2.1 and Ni2.2 alloys at RT, −70, and −196 °C Test temperature

Alloys

σ y (MPa)

σ UTS (MPa)

EL (%)

RT

Ni2.0

545.6

1076

16.6

Ni2.1

546.4

1046

17.7

Ni2.2

544.5

1120

20.5

Ni2.0

580

1034

10.5

Ni2.1

595

1168

15.8

Ni2.2

570

1143

18.0

Ni2.0

715

952

3.7

Ni2.1

690

1051

6.7

Ni2.2

705

1151

9.3

−70 °C

−196 °C

[83]. The small-scale ingot exhibited superior tensile property than that of largescale ingot at cryo-temperatures, especially at −196 °C. Several researchers [84, 85] cast button-like AlCoCrFeNi2.1 EHEA ingots by arc melting and tested their tensile properties. Their results showed that the mechanical properties of the alloy were comparable to those in previous studies provided by Wani et al. Therefore, in terms of the AlCoCrFeNi2.1 EHEA, the two casting methods have little effect on the microstructure and properties of the alloy. Since small-scale alloy ingots are easier to prepare and can save raw materials, most of the as-cast EHEAs are made by arc melting. A summary of the mechanical properties including yield strength (σ y ), fracture strength (σ max ), and plastic strain (εp ) of the reported EHEAs prepared by arc melting is given in Table 3.

Fig. 14 a EBSD phase map and b TEM micrograph of the as-cast EHEA of AlCoCrFeNi2.1 prepared by arc melting showing the lamellar structure [82]

– –

1200(C) 1150(C)

RT RT RT RT RT RT RT RT RT RT RT RT RT RT RT RT 600 °C

BCC + B2

BCC + B2

BCC + B2

BCC + B2

BCC + B2

BCC1 + BCC2

FCC + B2

FCC + B2

FCC + BCC

FCC + Laves

FCC + BCC

FCC + Laves

FCC + Laves

FCC + BCC

FCC + Laves

Al0.9 CrFeNi

AlCrFeNi

Al1.1 CrFeNi

Al1.2 CrFeNi

Al1.3 CrFeNi

AlCrFeNiMo0.2

(Co40 Cr10 Fe5 Mo5 Ni40 )82 Al17.8

CrFeNi2.2 Al0.8

Ni30 Co30 Cr10 Fe10 Al18 W2

CoCu0.5 FeNiTa0.5

Al0.9 CoCrNi2.1

(CoFe2 NiV0.5 Mo0.2 )91 Nb9

CoCrFeNiHf0.4

Al0.8 CoCr0.6 Fe0.7 Ni1.5

CoCrFeNiHf0.4

857 ±

−196 °C

FCC + Laves

20.4 ±

16.6 ± 2.8(T)

1461 ± 3002.4(C)

980(T)

2050(C)

1501(C) 490(T)

2050.8(C)

1118.2(C)

1033(T)

3160(C)

1595(C) 645(T)

1266.5(T)

–

956(T)

2813(C)

668(C) 479(T)

3222(C)

1487.4(C)

3607.6(C)

3513.0(C)

906.8(C) 1122.8(C)

3450.4(C)

1044.0(C)

3002.4(C)

2200 ± 10(C)

770.0(C)

>50(C)

5(C)

14.8(T)

4.5(C)

19.3(C)

6.9(T)

24(C)

20.3(T)

9.1(T)

33.5(C)

28.7(C)

24.7(C)

31.4(C)

27.1(C)

34.6(C)

37.1(C)

17.0 ± 0.5(C)

19.5 ± 0.6(T)

17(T)

1319 ± 14(T)

2060 ± 5(C) 843.2(C)

2.1(T)

18.3 ± 1.9(T)

1155 ± 1183 ± 18(T)

CoCrFeNiNb0.5

697 ± 14(T)

−150 °C 21(T)

620 ± 675 ± 15(T)

RT −90 °C

B2 + L12

εp (%)

25(T)

σ max (MPa)

AlCoCrFeNi2.1

7(T)

Temperature

Phases

Alloy

σ y (MPa)

(continued)

[87]

[87]

[59]

[86]

[58]

[57]

[52]

[51]

[49]

[46]

[32]

[37]

[37]

[37]

[37]

[37]

[40]

[83]

[83]

[83]

[83]

Ref.

Table 3 Mechanical properties (σ y , σ max , εp ) of as-cast EHEAs prepared by arc melting under tensile (T) and compression (C) at different temperatures (RT: Room temperature)

72 W. Jiao et al.

– –

7(C)

767 ± 592 ± 16(C)

RT RT RT RT RT 600 °C 700 °C RT RT RT RT RT RT

FCC + Laves

L12 + B2

FCC + Laves

L12 + B2

FCC + B2

BCC + MC

B2 + L12

FCC + B2

FCC + B2

FCC + B2

FCC + Laves

Al19.25 Co18.86 Fe18.36 Ni43.53

Fe2 NiCrNb0.34

AlCoCrFeNi2

AlCrFeNiCo1.9

Re0.5 MoNbW(TaC)0.8

Fe20 Co20 Ni41 Al19

Al17 Co14.3 Cr14.3 Fe14.3 Ni40.1

Al17 Co28.6 Cr14.3 Fe14.3 Ni25.8

Al17 Co14.3 Cr14.3 Fe28.6 Ni25.8

CoFeNi2 V0.5 Nb0.75

Notes T-tensile test; C-compressive test

–

903 ± 8(C)

RT

FCC + Laves

CrFeNiNb0.35

2073(C)

731 ± 60(T)

473 ± 32(T)

3.4(C)

10.3 ± 2232(C)

1.7(T)

1145 ±

38(T)

14.8 ± 2.8(T)

14.0 ± 2.1(T)

18.7(T)

8.90 ± 0.15(C)

–

–

–

~17.2(T)

30.8(C)

~10(T)

30.23 ± 1.32(C)

14.4(C)

9.2(C)

3.7(C)

>50(C)

εp (%)

1001 ± 26(T)

1067 ± 50(T)

1103(T)

577(T) 479 ± 35(T)

2347 ± 40(C)

1340 ± 21(C)

~881.7(T)

2267(C)

988.7(C) ~501.7(T)

~956(T)

2411.12 ± 17.21(C)

~486(T)

1204.67 ± 12.85(C)

2060.6(C)

1232.1(C)

CrFeNi2 Nb0.6

1745(C)

RT

FCC + σ –

RT

FCC + σ 1875(C)

1065(C)

Fe2 Ni2 CrMo1.25

–

536(C)

800 °C

Fe35 Ni25 Cr25 Mo15

σ max (MPa)

σ y (MPa)

Temperature

Phases

Alloy

Table 3 (continued)

[35]

[23]

[23]

[23]

[44]

[71]

[69]

[69]

[69]

[90]

[68]

[89]

[88]

[67]

[66]

[64]

[87]

Ref.

3 Eutectic High-Entropy Alloys 73

74

W. Jiao et al.

In addition, directional solidification process is also commonly used to fabricate EHEAs. During directional solidification, heating power and withdrawal velocities determine the temperature gradient and growth direction of the alloy, which assure the microstructure of the alloy. The AlCoCrFeNi2.1 alloy rods were directionally solidified at a constant growth rate of 6 mm/min and cast into samples of Φ16 × 255 mm by Zhang et al. [91, 92]. The engineering tensile stress–strain curves at room temperature and at 600–1050 °C are shown in Fig. 15. The yield strength, ultimate tensile strength, and elongation to failure of the alloy were about 1150 MPa, 1000 MPa, and 18%, respectively, at room temperature, which was superior to that of the alloy prepared by arc melting and induction melting. When increasing the testing temperature, the strength of the alloy decreased, and the plasticity increased. Note worthily, the alloy exhibited the best comprehensive mechanical properties during 600 and 700 °C, which may be related to the twining deformation mechanism. When the test temperature was over 900 °C, the alloy exhibited ultra-high ductility, which could be attributed to the simultaneous occurring of the precipitate hardening and dynamic recrystallization. The magnetic field is also effective to tailor the morphology of materials during the solidification process. The influence of magnetic field-assisted directional solidification on the microstructure and mechanical properties in the AlCoCrFeNi2.1 EHEA prepared by directional solidification at different withdrawal rates (v = 1–150 μm/s) has been investigated by Wang et al. [93]. Figure 16 shows the effect of the growth speed and the magnetic field on the eutectic lamellar spacing (λl ) and the eutectic cellular spacing (λc ) of the directionally solidified AlCoCrFeNi2.1 EHEAs. It can be seen that eutectic lamellar spacing and eutectic cellular spacing decreased as the growth speed increased, and the magnetic field imposed. Consequently, the refinement of microstructures greatly affected mechanical properties of the alloy. From Fig. 17, it is obvious to notice that the plasticity of the EHEA was greatly improved with the increase of the withdrawal rates and the application of the magnetic field,

Fig. 15 Engineering tensile stress–strain curves of as-cast AlCoCrFeNi2.1 prepared by directional solidification, at the strain rate of 1 × 10–3 /s and temperatures of room temperature, 600 °C, 700 °C, 800 °C, 900 °C, 1000 °C, and 1050 °C [91, 92]

3 Eutectic High-Entropy Alloys

75

Fig. 16 Effect of the growth speed and the magnetic field on the eutectic lamellar spacing (λl ) and the eutectic cellular spacing (λc ) of the directionally solidified AlCoCrFeNi2.1 EHEA: a and b λl and λc as a function of the growth speed in the case of no magnetic field; c and d λl and λc histograms as a function of the growth speed in both cases of 0 T and 6 T magnetic field [93]

almost without sacrificing strength. Therefore, directional solidification would be an effective pathway to achieve superior mechanical properties in EHEAs. Recently, Shi et al. [94] used directional solidification to fabricate an Al19 Fe20 Co20 Ni41 EHEA with outstanding mechanical properties. The mechanism for the superior properties was that the directional solidified alloy had a hierarchically organized herringbone structure. The special structure enabled bionic-inspired hierarchical crack buffering, which helped the cracks to avoid catastrophic growth and percolation. In this structure, the high-density crack did not harm the elongation, but instead acted as an effective ductilizing strategy to compensate the finite tensile ductility of the lamellar structure with poor deformability. Therefore, the EHEA exhibited an ultra-high uniform tensile elongation of ~50% without sacrificing strength. The yield strength and uniform elongation of the former were 1.3 times and 3 times higher than that of the latter, respectively.

76

W. Jiao et al.

Fig. 17 Mechanical properties of the AlCoCrFeNi2.1 EHEA prepared by suction casting and directional solidification at various growth speeds with and without a 6 T magnetic field: a engineering tensile stress–strain curves of the specimens prepared by suction-cast and directional solidification at various growth speeds without magnetic field; b engineering tensile stress–strain curves of the directionally solidified specimens at growth speeds of 50 μm/s, 100 μm/s and 150 μm/s with and without a 6 T magnetic field [93]

4.2 Influence of Thermo-Mechanical Processing (TMP) on Mechanical Properties of EHEAs Like many other materials, thermo-mechanical processing (TMP) is conducted to further optimize the mechanical properties of EHEAs. Tailoring the structural/mechanical properties by the application of TMP provides immense possibilities for the development of novel EHEAs. The effects of rolling and annealing at different temperatures on microstructure and properties of EHEA were investigated. Wani et al. [82, 95] firstly studied the influence of cold-rolling and annealing on microstructure and mechanical of the AlCoCrFeNi2.1 EHEA. When the EHEA was heavily cold-rolled to 90%, the yield strength (YS), ultimate tensile strength (UTS), and the elongation to failure (εf ) of the alloy were 1652 MPa, 1800 MPa, and 6%, respectively. Cold-rolling resulted a drastic increase in YS and UTS but at the expense of elongation. When further annealing the rolled EHEA at different temperature for 1 h, the comprehensive mechanical properties could be effectively improved. Figure 18 shows the engineering stress– strain curves of the AlCoCrFeNi2.1 EHEA in the as-cast, cold-rolled, and annealed conditions. It is apparent that a significant enhancement in mechanical properties compared with the as-cast EHEA was achieved after thermo-mechanical processing. Realizing that the EHEA could be successfully tailored using thermo-mechanical processing, Shi et al. [96] obtained a kind of fine-grained dual-phase heterogenous lamella structure (DPHL) by multi-pass rolling the AlCoCrFeNi2.1 EHEA to 84–86% thickness reduction and annealing it at different temperatures (660, 700, 740, and 900 °C) for 1 h. The engineering stress–strain curves of the three DPHL EHEAs at room

3 Eutectic High-Entropy Alloys

77

Fig. 18 Engineering stress–strain curves of the AlCoCrFeNi2.1 EHEA in the as-cast, cold-rolled, and annealed conditions [95]

temperature are depicted in Fig. 19. All the EHEAs exhibited superior strength– ductility combination than the as-cast EHEA, which was attributed to both the twohierarchical constraint effect and the self-generated microcrack-arresting mechanism. In addition, Xiong et al. [85] designed a new thermo-mechanical processing route to achieve high strength and high ductility of the AlCoCrFeNi2.1 EHEA via precipitation strengthening in a heterogeneous structure. The EHEA was cold-rolled to ~70% and then annealed at 1000 °C for 1 h (the sample denoted as CR-R), and then aged at 600 °C for 50 h (the sample denoted as CR-R-A). Figure 20 shows the engineering stress–strain curves of the EHEA AlCoCrFeNi2.1 at as-cast state and after thermo-mechanical processing. It can be seen that the CR-R-A exhibited best comprehensive mechanical property. The outstanding combination of high strength and excellent ductility was attributed to a dual-phase heterogeneous structure and nano-precipitates in both phases. Bhattacharjee et al. [97] investigated the influence of cryo-rolling and annealing on the microstructure and mechanical properties of the AlCoCrFeNi2.1 EHEA. They found that when cryo-rolling the alloy to ~90% and then annealing it at 800 °C for 1 h, the EHEA showed an outstanding increase in strength but the elongation to failure remained almost unchanged compared to the as-cast alloy. The engineering stress– strain plots of the EHEA in various TMP conditions are displayed in Fig. 21. They attributed the high strength without loss of ductility to the synergistic effect exerted by the hierarchical microstructure. Compared with cold/cryo-rolling and hot-rolling, warm-rolling combined the advantages of the reduced flow stress and an improved surface finish. Based on that, Reddy et al. [98] warm-rolled the AlCoCrFeNi2.1 to

78

W. Jiao et al.

Fig. 19 Engineering stress–strain curves of the three DPHL EHEAs and the as-cast EHEA at room temperature [96]

Fig. 20 Engineering stress–strain curves of the EHEA AlCoCrFeNi2.1 at the as-cast condition and after thermo-mechanical processing [85]

3 Eutectic High-Entropy Alloys

79

Fig. 21 Engineering stress–strain plots of the EHEA in various heat-treated conditions [97]

90% at 400, 600, and 750 °C and then annealed the alloy at temperatures ranging from 800 °C to 1200 °C for 1 h. The excellent strength–ductility balance was acquired in the EHEA after warm-rolling at 750 °C, as seen in Fig. 22. Furthermore, a novel hybrid-rolling process was conducted on the AlCoCrFeNi2.1 EHEA by Reddy et al. [99]. The hybrid-rolling process included cryo-rolling and warm-rolling at 600 °C. An extremely heterogeneous microstructure consisting of retained lamellar and transformed nanocrystalline regions were obtained by the novel processing strategy, which resulted in the high yield strength of ~1900 MPa, ultra-tensile strength of ~2000 MPa, and ductility of ~8%. The mechanical properties of typical EHEA AlCoCrFeNi2.1 were improved by different TMP processes. Inspired by the positive effect from TMP processing on the as-cast EHEA, many researchers designed various TMP processing to tailor the mechanical properties of the EHEAs. Table 4 collected the tensile properties of EHEAs processed by different thermo-mechanical processing methods.

4.3 Influence of Other Preparation Techniques on Mechanical Properties of EHEAs Preparation techniques determine the microstructure of an alloy, which further determine the mechanical properties. Although directly casting is the most convenient

80

W. Jiao et al.

Fig. 22 Tensile stress–strain plots of the EHEA material processed in different thermo-mechanical conditions [98]

method to produce EHEAs, the emergence of other preparation methods provides more possibilities for tailoring the microstructures and properties of EHEAs. Han et al. [55] reported an equiaxed Co25.1 Cr18.8 Fe23.3 Ni22.6 Ta8.5 Al1.7 EHEA consisting of FCC phase with nano-sized L12 precipitates and C14 Laves phase prepared by powder metallurgy. These microstructures were shown to be stable at a temperature of 1273 K for 100 h. The EHEA possessed well-balanced mechanical properties such as yield strength (~800 MPa) and tensile ductility (~16%) at a temperature up to 1073 K. John et al. [100] firstly combined mechanical alloying and spark plasma sintering (MA-SPS) to synthesize the typical AlCoCrFeNi2.1 EHEA. The microstructure of the alloy is shown in Fig. 23, which was rather different from that of the as-cast alloy. The bimodal (hierarchical) microstructure was observed in the MA-SPS alloy, in which apart from minor fraction of microcrystalline FCC containing nanocrystalline FCC and B2 phases, Al2 O3 and hexagonal Cr7 C3 were also seen. The strengthening mechanism including grain boundary strengthening, twin boundary strengthening, and dispersion strengthening caused by the complex structure render the alloy exhibit a high fracture strength of 1915 MPa, but it fractured before reaching the yield point. Vikram et al. [101] prepared an AlCoCrFeNi2.1 EHEA through the additive manufacturing (AM) route. Compared with the cast EHEA, phases in the AMEHEA were still L12 and B2, but the morphologies of which exhibited dendritic and eutectic structure, as shown in Fig. 24. It is worth noting that the morphology has anisotropy. The

3 Eutectic High-Entropy Alloys

81

Table 4 Tensile properties of EHEAs after TMP Alloys

Processing

Phases

σ y (MPa)

σ max (MPa)

εp (%)

Ref.

AlCoCrFeNi2.1

AC

L12 + B2

620

1050

17

[82]

CR 90%,

FCC + B2

1625

1800

6

[82]

CR 90%, 800 °C 1 h FCC + B2

1108

1200

12

[82]

CR 90%, 1000 °C 1h

FCC + B2

844

1175

23

[95]

CR 90%, 1200 °C 1h

FCC + B2

648

1075

27

[95]

CR 84–86%, 660 °C FCC + 1h B2

1490 ± 17 1638 ± 25

16.0 ± 0.5

[96]

CR 84–86%, 700 °C FCC + 1h B2

1263 ± 16 1442 ± 19

21.4 ± 0.9

[96]

CR 84–86%, 740 °C FCC + 1h B2

1154 ± 9

1340 ± 15

25.0 ± 1.1

[96]

CR 70%, 1000 °C 1h

FCC + B2

732.3

1147.7

24.6

[85]

CR 70%, 1000 °C 1 h, A-600 °C 50 h

FCC + B2

1008.8

1475.6

19.2

[85]

Cryo-R 90%, 800 °C 1 h

FCC + B2

1437 ± 26 1562 ± 33

14 ± 1

[97]

WR-400 °C 90%

FCC + B2

1390 ± 14 1940 ± 57

2.4 ± 0.7

[98]

WR-400 °C 90%, 800 °C 1 h

FCC + B2

1069 ± 96 1433 ± 37

17.1 ± 3.4

WR-400 °C 90%, 1000 °C 1 h

FCC + B2

679 ± 6

1208 ± 17

27.8 ± 0.3

WR-400 °C 90%, 1200 °C 1 h

FCC + B2

479 ± 16

1096 ± 2

28.9 ± 2.6

WR-600 °C 90%

FCC + B2

1358 ± 25 1880 ± 14

4.5 ± 0.8

WR-600 °C 90%, 800 °C 1 h

FCC + B2

1110 ± 78 1482 ± 32

16.8 ± 3.0

WR-600 °C 90%, 1000 °C 1 h

FCC + B2

619 ± 7

1170 ± 14

29.5 ± 3.5

WR-600 °C 90%, 1200 °C 1 h

FCC + B2

508 ± 5

1507 ± 17

15.6 ± 1.2

WR-750 °C 90%

FCC + B2

1192 ± 11 1635 ± 49

18.0 ± 1.4

WR-750 °C 90%, 800 °C 1 h

FCC + B2

1108 ± 38 1507 ± 17

15.6 ± 1.2 (continued)

82

W. Jiao et al.

Table 4 (continued) Alloys

Al0.7 CoCrFeNi

Al0.9 CoCrNi2.1

Fe20 Co20 Ni41 Al19

Processing

Phases

σ y (MPa)

σ max (MPa)

εp (%)

WR-750 °C 90%, 1000 °C 1 h

FCC + B2

705 ± 40

1194 ± 50

31.0 ± 0.1

WR-750 °C 90%, 1200 °C 1 h

FCC + B2

482 ± 11

1158 ± 66

31.0 ± 2.8

HR 90%,

FCC + B2

1900

2000

8

[99]

Homogenized-1150 °C 1 h, CR 30%, 1100 °C 5 min

FCC + B2

770

1090

17

[47]

Homogenized-1150 °C 1 h, CR 30%, 1100 °C 5 min, 580 °C 24 h

FCC + L12 & B2 + BCC

1080

1370

15

AC

FCC + BCC

645

1033

6.9

CR 65%

FCC + BCC

1800

1938

0.97

CR 65%, 1000 °C 1h

FCC + BCC

820

1400

11.6

AC

FCC + BCC

520

1103

16.17

1220

1520

24.24

CR 83%, 660 °C 1 h FCC + BCC

Ref.

[57]

[61]

Notes AC-as-cast; CR-cold-rolling; WR-warm-rolling; HR-hybrid-rolling; Cryo-R–cryo-rolling

Fig. 23 SEM BSE micrographs of the MA-SPS AlCoCrFeNi2.1 alloy at a low magnification and b high magnification. 1—white phase (FCC), 2—light gray phase (B2) and 3—dark gray phases (carbide), and 4—black dots (Al2 O3 ) [100]

3 Eutectic High-Entropy Alloys

83

compression tests were conducted at different temperatures to evaluate the mechanical properties of the EHEA. The engineering stress–strain plots of the AM EHEA along (Z) and across (X) build at room temperature and along (Z) build at high temperature are shown in Fig. 25. The EHEA exhibited yield anisotropy along and across the build direction at room temperature, which could be entirely attributed to the microstructural features. Increasing the testing temperature, the yield strength first increased up to 400 °C and then gradually decreased.

Fig. 24 a SEM micrograph of as-printed AM EHEA from the bottom zone, b, c SEM micrograph at the Z and X faces, respectively. The micrographs reveal dendritic and eutectic features [101]

84

W. Jiao et al.

Fig. 25 Engineering stress–strain plot of the AM EHEA (a), (Z) and across (X) build at room temperature and (b) along (Z) build at high temperature [101]

5 Future Trends and Potential Applications 5.1 Future Trends Combining the advantages of both HEAs and eutectic alloys, EHEAs are expected to be the most promising HEAs for the large-scale and extensive industrial applications. Future trends of the EHEAs may be still directed toward two very fundamental aspects, on the one hand, the development of new alloy systems, and on the other hand, the optimization of properties and the establishment of the relationship between properties and microstructure of EHEAs. Most reported EHEAs can be divided into two categories according to their eutectic phases. One is composed of two hard eutectic phases, such as BCC/B2 phases [8, 37]. Such EHEAs usually cannot be applied in a load-bearing structure especially when cyclic fatigue exists but could be employed as functional materials. The other is composed of a hard phase and a soft phase, such as FCC + BCC (B2) or FCC + IMC. These EHEAs generally show desirable balance of strength and ductility, which has great potential in industrial applications. Therefore, when exploring new EHEAs, more attentions should be paid to the second kind of EHEAs. Besides, since the majority of EHEAs contain expensive elements, how to reduce the cost of the alloy while still obtaining the excellent performance of EHEAs is an issue that worth considering. It has been known that most EHEAs possess excellent properties in the cast condition; however, suitable secondary processing and thermo-mechanical treatments may further enhance the mechanical and functional properties of these alloys. However, there is still much space for exploring subsequent processing routes to improve the properties of EHEAs. Moreover, many underlying deformation mechanisms of various new EHEAs have not been understood conclusively.

3 Eutectic High-Entropy Alloys

85

5.2 Potential Applications In the future, EHEAs hold great potential as structural and functional materials in industrial applications. As for fabrications, the advantages in casting and welding ensure EHEAs to be suitable for large-scale industrial applications. Moreover, with a high content of Al or Cr, most EHEAs possess good corrosion and oxidation resistance. One of the most promising applications lies in components used at low temperatures. As the temperature decreases, the tensile strength of EHEAs increases, while the high ductility does not degrade significantly [83]. The prototype AlCoCrFeNi2.1 EHEA with excellent corrosion resistance and low-temperature mechanical properties can be potentially applied in propeller for icebreakers, replacing copper alloys and stainless steels. At an even lower temperature of 77 K, EHEAs are also applicable without the ductile–brittle transition. The high-temperature properties of the AlCoCrFeNi2.1 EHEA promote the related high-temperature applications. The directionally solidified AlCoCrFeNi2.1 EHEA exhibits high strength at 700 °C without showing intermediate temperature embrittlement [92]. The AlCr1.3 TiNi2 EHEA shows outstanding oxidation resistance and specific yield strengths at high temperatures [70]. In power station and aerospace, the alloys can be possibly applied for extreme service conditions. In addition, EHEAs can be used as coating materials in engineering structural parts. The EHEAs containing Hf, Nb, Ta, and Zr have been shown superior tribological properties due to the combination of soft FCC and hard Laves phases, and their excellent wear-resistance was demonstrated from room temperature to a high temperature of 1100 K [102]. The exploration of function properties provides another promising application direction for EHEAs. Recently, Tsai et al. [103] reported that some EHEAs have good damping properties and can be used as shock absorbers. With in-situ formed composite structures leading to combined properties, EHEAs could be revolutionary metallic materials. In the future, if the lamellar spacing of EHEAs is reduced to about tens of nanometers, with one being the hard-magnetic phase and the other being the soft magnetic phase, a giant magnetoresistance effect could occur.

References 1. Lu YP, Dong Y, Guo S et al (2014) A promising new class of high-temperature alloys: eutectic high-entropy alloys. Sci Rep 4:6200 2. Yeh JW, Chen SK, Lin SJ et al (2004) Nanostructured high-entropy alloys with multiple principal elements novel alloy design concepts and outcomes. Adv Eng Mater 6:299–303 3. Cantor B, Chang ITH, Knight P et al (2004) Microstructural development in equiatomic multicomponent alloys. Mater Sci Eng A 375–377:213–218 4. Chanda B, Potnis G, Jana PP et al (2020) A review on nano-ultrafine advanced eutectic alloys. J Alloy Compd 827:154226

86

W. Jiao et al.

5. Chang I, Cai Q (2021) From simple binary to complex multicomponent eutectic alloys. Prog Mater Sci 123:100779 6. Tiwary CS, Pandey P, Sarkar S et al (2021) Five decades of research on the development of eutectic as engineering materials. Prog Mater Sci 123(13–29):100793 7. Lu Y, Dong Y, Jiang H et al (2018) Promising properties and future trend of eutectic high entropy alloys. Scripta Mater 187:202–209 8. Baker I, Wu M, Wang ZW (2019) Eutectic/eutectoid multi-principle component alloys: a review. Mater Charact 147:545–557 9. Otto F, Dlouhy A, Somsen C et al (2013) The influences of temperature and microstructure on the tensile properties of a CoCrFeMnNi high-entropy alloy. Acta Mater 61(15):5743–5755 10. Senkov NO, Wilks BG, Miracle BD et al (2010) Refractory high-entropy alloys. Intermetallics 18(9):1758–1765 11. Glicksman ME (2011) Principles of solidification: an introduction to modern casting and crystal growth concepts. Springer Science & Business Media, pp 1–520 12. Smallman R, Ngan A (2011) Physical metallurgy and advanced materials engineering, 7th ed 13. Lu YP, Gao XZ, Jiang L et al (2017) Directly cast bulk eutectic and near-eutectic high entropy alloys with balanced strength and ductility in a wide temperature range. Acta Mater 124:143–150 14. Jiang H, Han KM, Gao XX et al (2018) A new strategy to design eutectic high-entropy alloys using simple mixture method. Mater Des 142:101–105 15. Zhang Y, Zuo TT, Tang Z et al (2014) Microstructures and properties of high-entropy alloys. Prog Mater Sci 61:1–93 16. Xie T, Xiong Z, Xu Z et al (2021) Another eutectic point of Co–Cr–Fe–Ni–M (Hf, Ta, Nb) high-entropy system determined using a simple mixture method correlated with mixing enthalpy. Mater Sci Eng A 802:140634 17. Xie T, Xiong Z, Xu Z et al (2020) Varying the eutectic composition of Co–Cr–Fe–Ni–Hf high-entropy system using a modified simple mixture method. Mater Sci Eng A 786:139420 18. Lu Y, Jiang H, Guo S et al (2017) A new strategy to design eutectic high-entropy alloys using mixing enthalpy. Intermetallics 91:124–128 19. Guo S, Ng C, Lu J et al (2011) Effect of valence electron concentration on stability of fcc or bcc phase in high entropy alloys. J Appl Phys 109(10):213 20. Yang X, Zhang Y (2012) Prediction of high-entropy stabilized solid-solution in multicomponent alloys. Mater Chem Phys 132(2–3):233–238 21. Ye YF, Wang Q, Lu J et al (2016) High-entropy alloy: challenges and prospects. Mater Today 19(6):349–362 22. Chanda B, Das J (2019) An assessment on the stability of the eutectic phases in high entropy alloys. J Alloy Compd 798:167–173 23. Jin X, Zhou Y, Zhang L et al (2018) A new pseudo binary strategy to design eutectic high entropy alloys using mixing enthalpy and valence electron concentration. Mater Des 143:49– 55 24. He F, Wang Z, Ai C et al (2019) Grouping strategy in eutectic multi-principal-component alloys. Mater Chem Phys 221:138–143 25. He F, Wang Z, Cheng P et al (2016) Designing eutectic high entropy alloys of CoCrFeNiNbx. J Alloy Compd 656:284–289 26. Chen HL, Mao H, Chen Q (2017) Database development and Calphad calculations for high entropy alloys: challenges, strategies, and tips. Mater Chem Phys 210:279–290 27. Wu M, Wang S, Huang H et al (2020) CALPHAD aided eutectic high-entropy alloy design. Mater Lett 262 28. Wu Q, Wang Z, Hu X et al (2020) Uncovering the eutectics design by machine learning in the Al–Co–Cr–Fe–Ni high entropy system. Acta Mater 182:278–286 29. Elliott R (1983) Eutectic solidification processing: crystalline and glassy alloys. Butterworths 30. Guo S, Ng C, Liu CT (2013) Anomalous solidification microstructures in Co-free AlxCrCuFeNi2 high-entropy alloys. J Alloy Compd 557:77–81

3 Eutectic High-Entropy Alloys

87

31. Tan Y, Li J, Wang J et al (2017) Seaweed eutectic-dendritic solidification pattern in a CoCrFeNiMnPd eutectic high-entropy alloy. Intermetallics 85:74–79 32. Dong Y, Lu Y, Kong J et al (2013) Microstructure and mechanical properties of multicomponent AlCrFeNiMox high-entropy alloys. J Alloy Compd 573:96–101 33. Ma SG, Zhang Y (2012) Effect of Nb addition on the microstructure and properties of AlCoCrFeNi high-entropy alloy. Mater Sci Eng A 532:480–486 34. Jiang L, Cao ZQ, Jie JC et al (2015) Effect of Mo and Ni elements on microstructure evolution and mechanical properties of the CoFeNixVMoy high entropy alloys. J Alloy Compd 649:585–590 35. Jiang L, Lu Y, Dong Y et al (2015) Effects of Nb addition on structural evolution and properties of the CoFeNi2 V0.5 high-entropy alloy. Appl Phys A 119(1):291–297 36. He F, Wang Z, Shang X et al (2016) Stability of lamellar structures in CoCrFeNiNbx eutectic high entropy alloys at elevated temperatures. Mater Des 104:259–264 37. Chen X, Qi JQ, Sui YW et al (2017) Effects of aluminum on microstructure and compressive properties of Al–Cr–Fe–Ni eutectic multi-component alloys. Mater Sci Eng A 681:25–31 38. Huo W, Zhou H, Fang F et al (2017) Microstructure and mechanical properties of CoCrFeNiZrx eutectic high-entropy alloys. Mater Des 134:226–233 39. Jiang H, Jiang L, Qiao D et al (2017) Effect of niobium on microstructure and properties of the CoCrFeNb x Ni high entropy alloys. J Mater Sci Technol 33(7):712–717 40. Chanda B, Das J (2018) Composition dependence on the evolution of nanoeutectic in CoCrFeNiNbx (0.45 ≤ x≤ 0.65) high entropy alloys. Adv Eng Mater 20(4):1700908 41. Guo Y, Liu L, Zhang Y et al (2018) A superfine eutectic microstructure and the mechanical properties of CoCrFeNiMox high-entropy alloys. J Mater Res 33(19):3258–3265 42. Huo W, Zhou H, Fang F et al (2018) Microstructure and properties of novel CoCrFeNiTax eutectic high-entropy alloys. J Alloy Compd 735:897–904 43. Jiang H, Han K, Qiao D et al (2018) Effects of Ta addition on the microstructures and mechanical properties of CoCrFeNi high entropy alloy. Mater Chem Phys 210:43–48 44. Jin X, Zhou Y, Zhang L et al (2018) A novel Fe20 Co20 Ni41 Al19 eutectic high entropy alloy with excellent tensile properties. Mater Lett 216:144–146 45. Liu D, Yu P, Li G et al (2018) High-temperature high-entropy alloys Alx Co15 Cr15 Ni70-x based on the Al–Ni binary system. Mater Sci Eng A 724:283–288 46. Gasan H, Ozcan A (2019) New eutectic high-entropy alloys based on Co–Cr–Fe–Mo–Ni–Al: design, characterization and mechanical properties. Met Mater Int 26(8):1152–1167 47. Gwalani B, Gangireddy S, Zheng Y et al (2019) Influence of ordered L12 precipitation on strain-rate dependent mechanical behavior in a eutectic high entropy alloy. Sci Rep 9(1):6371 48. Han L, Xu X, Wang L et al (2019) A eutectic high-entropy alloy with good high-temperature strength-plasticity balance. Mater Res Lett 7(11):460–466 49. Jin X, Bi J, Zhang L et al (2019) A new CrFeNi2 Al eutectic high entropy alloy system with excellent mechanical properties. J Alloy Compd 770:655–661 50. Rahul MR, Phanikumar G (2019) Design of a seven-component eutectic high-entropy alloy. Metall Mater Trans A 50(6):2594–2598 51. Wu Q, Wang Z, Zheng T et al (2019) A casting eutectic high entropy alloy with superior strength-ductility combination. Mater Lett 253:268–271 52. Xu T, Lu Y, Cao Z et al (2019) Effects of Ta addition on the microstructure and mechanical properties of CoCu0.5 FeNi high-entropy alloy. J Mater Eng Perform 28(12):7642–7648 53. Dong Y, Yao Z, Huang X et al (2020) Microstructure and mechanical properties of AlCox CrFeNi3-x eutectic high-entropy-alloy system. J Alloy Compd 823:153886 54. Han L, Quan T, Liu B et al (2020) Revealing the excellent high-temperature oxidation resistance of a non-equimolar Al1 Co25 Cr18 Fe23 Ni23 Ta10 compositional complex eutectic alloy. J Alloy Compd 846:156265 55. Han L, Xu X, Li Z et al (2020) A novel equiaxed eutectic high-entropy alloy with excellent mechanical properties at elevated temperatures. Mater Res Lett 8(10):373–382 56. Jain R, Dewangan SK, Kumar V et al (2020) Artificial neural network approach for microhardness prediction of eight component FeCoNiCrMnVAlNb eutectic high entropy alloys. Mater Sci Eng A 797:140059

88

W. Jiao et al.

57. Jin X, Liang Y, Bi J et al (2020) Enhanced strength and ductility of Al0.9 CoCrNi2.1 eutectic high entropy alloy by thermomechanical processing. Materialia 10:100639 58. Li R, Ren J, Zhang GJ et al (2020) Novel (CoFe2 NiV0.5 Mo0.2 )100-x Nbx eutectic high-entropy alloys with excellent combination of mechanical and corrosion properties. Acta Metall Sinica (English Letters) 33:1046–1056 59. Ma L, Wang J, Jin P (2020) Microstructure and mechanical properties variation with Ni content in Al0.8 CoCr0.6 Fe0.7 Nix (x = 1.1, 1.5, 1.8, 2.0) eutectic high-entropy alloy system. Mater Res Exp 7(1):016566 60. Shao Y, Ma H, Wang Y (2020) Effect of Mo addition on the microstructure and mechanical properties of CoCuFeNi high entropy alloy. Metals 10(8):1017 61. Shi P, Zhong Y, Li Y et al (2020) Multistage work hardening assisted by multi-type twinning in ultrafine-grained heterostructural eutectic high-entropy alloys. Mater Today 41:62–71 62. Sunkari U, Reddy SR, Athira KS et al (2020) Effect of niobium alloying on the microstructure, phase stability and mechanical properties of CoCrFeNi2.1 Nbx high entropy alloys: experimentation and thermodynamic modeling. Mater Sci Eng A 793:139897 63. Wei Q, Luo G, Zhang J et al (2020) Designing high entropy alloy-ceramic eutectic composites of MoNbRe0.5 TaW(TiC)x with high compressive strength. J Alloys Compd 818:152846 64. Yin Y, Tan Q, Wang T et al (2020) Eutectic modification of Fe-enriched high-entropy alloys through minor addition of boron. J Mater Sci 55(29):14571–14587 65. Jiang H, Qiao D, Jiao W et al (2021) Tensile deformation behavior and mechanical properties of a bulk cast Al0.9 CoFeNi2 eutectic high-entropy alloy. J Mater Sci Technol 61:119–124 66. Jiao W, Jiang H, Qiao D et al (2021) Effects of Mo on microstructure and mechanical properties of Fe2 Ni2 CrMox eutectic high entropy alloys. Mater Chem Phys 260:124175 67. Lee W-H, Oh Y, Jo M-G et al (2021) Microstructures and mechanical properties of CrFeNi2 Nbx eutectic multicomponent alloys. J Alloy Compd 860:158502 68. Song J, Chai Z, Zheng J et al (2021) Design Fe-based eutectic medium-entropy alloys Fe2 NiCrNbx . Acta Metall Sinica (English Letters) 34(8):1103–1108 69. Vikram RJ, Gupta K, Suwas S (2021) Design of a new cobalt base nano-lamellar eutectic high entropy alloy. Scripta Mater 202:113993 70. Wang M, Lu Y, Wang T et al (2021) A novel bulk eutectic high-entropy alloy with outstanding as-cast specific yield strengths at elevated temperatures. Scripta Mater 204:114132 71. Wei QQ, Xu XD, Li GM et al (2021) A carbide-reinforced Re0.5 MoNbW(TaC)0.8 eutectic high-entropy composite with outstanding compressive properties. Scripta Mater 200:113909 72. Jin X, Bi J, Liang Y et al (2021) Triple-phase eutectic high-entropy alloy: Al10 Co18 Cr18 Fe18 Nb10 Ni26 . Metall Mater Trans A 52:1314–1320 73. Li T, Lu Y, Wang T et al (2021) Grouping strategy via d-orbit energy level to design eutectic high-entropy alloys. Appl Phys Lett 119(7):071905 74. Ye X, Xiong J, Wu X et al (2021) A new infinite solid solution strategy to design eutectic high entropy alloys with B2 and BCC structure. Scripta Mater 199:113886 75. Oquab D, Josse C, Proietti A et al (2019) Solidification sequence and four-phase eutectic in AlSi6 Cu4 Fe2 alloy. Mater Charact 156:109846 76. Jiang L, Lu Y, Wu W et al (2016) Microstructure and mechanical properties of a CoFeNi2 V0.5 Nb0.75 eutectic high entropy alloy in As-cast and heat-treated conditions. J Mater Sci Technol 32(3):245–250 77. Jiang H, Qiao D, Lu Y et al (2019) Direct solidification of bulk ultrafine-microstructure eutectic high-entropy alloys with outstanding thermal stability. Scripta Mater 165:145–149 78. Guo S, Ng C, Liu CT (2013) Sunflower-like solidification microstructure in a near-eutectic high-entropy alloy. Mater Res Lett 1(4):228–232 79. Borkar T, Gwalani B, Choudhuri D et al (2016) Hierarchical multi-scale microstructural evolution in an as-cast Al2 CuCrFeNi2 complex concentrated alloy. Intermetallics 71:31–42

3 Eutectic High-Entropy Alloys

89

80. Zhang J, Qiu H, Zhu H et al (2021) Effect of Al additions on the microstructures and tensile properties of Alx CoCr3 Fe5 Ni high entropy alloys. Mater Charact 175:111091 81. Zhang L, Zhang Y (2020) Tensile properties and impact toughness of AlCox CrFeNi3.1-x (x = 0.4, 1) high-entropy alloys. Front Mater 7:92 82. Wani TBIS, Sheikh S, Lu YP et al (2016) Ultrafine-grained AlCoCrFeNi2.1 eutectic highentropy alloy. Mater Res Lett 4(3):174–179 83. Bhattacharjee T, Zheng R, Chong Y et al (2018) Effect of low temperature on tensile properties of AlCoCrFeNi2.1 eutectic high entropy alloy. Mater Chem Phys 210:207–212 84. Chen X, Xie W, Zhu J et al (2021) Influences of Ti additions on the microstructure and tensile properties of AlCoCrFeNi2.1 eutectic high entropy alloy. Intermetallics 128:107024 85. Xiong T, Zheng S, Pang J et al (2020) High-strength and high-ductility AlCoCrFeNi2.1 eutectic high-entropy alloy achieved via precipitation strengthening in a heterogeneous structure. Scripta Mater 186:336–340 86. Ma H, Shek CH (2020) Effects of Hf on the microstructure and mechanical properties of CoCrFeNi high entropy alloy. J Alloy Compd 827:154159 87. Xu ZQ, Ma ZL, Xia GH et al (2020) Microstructures and mechanical properties of CoCrFeNiHfx high-entropy alloys. Mater Sci Eng A 792:139820 88. Lu Y, Zhang M, Zhang L et al (2021) Cobalt-element-free eutectic medium-entropy alloys with superior mechanical performance and processability. Mater Sci Eng A 801:140421 89. Shi P, Li Y, Wen Y et al (2021) A precipitate-free AlCoFeNi eutectic high-entropy alloy with strong strain hardening. J Mater Sci Technol 89:88–96 90. Tian QW, Zhang GJ, Yin KX et al (2021) Effect of Ni content on the phase formation, tensile properties and deformation mechanisms of the Ni-rich AlCoCrFeNix (x = 2, 3, 4) high entropy alloys. Mater Charact 176:111148 91. Zhang Y, Wang X, Li J, Huang Y et al (2018) Deformation mechanism during hightemperature tensile test in an eutectic high-entropy alloy AlCoCrFeNi2.1 . Mater Sci Eng A 724:148–155 92. Zhang Y, Li J, Wang X et al (2019) The interaction and migration of deformation twin in an eutectic high-entropy alloy AlCoCrFeNi2.1 . J Mater Sci Technol 35(5):902–906 93. Wang J, Long Z, Jiang P et al (2020) Microstructure, crystallographic orientation and mechanical property in AlCoCrFeNi2.1 eutectic high-entropy alloy under magnetic field-assisted directional solidification. Metall Mater Trans A 51(7):3504–3517 94. Shi P, Li RG, Li Y et al (2021) Hierarchical crack buffering triples ductility in eutectic herringbone high-entropy alloys. Science 373(6557):912–918 95. Wani IS, Bhattacharjee T, Sheikh S et al (2016) Tailoring nanostructures and mechanical properties of AlCoCrFeNi2.1 eutectic high entropy alloy using thermo-mechanical processing. Mater Sci Eng A 675:99–109 96. Shi P, Ren W, Zheng T et al (2019) Enhanced strength-ductility synergy in ultrafine-grained eutectic high-entropy alloys by inheriting microstructural lamellae. Nat Commun 10(1):489 97. Bhattacharjee T, Wani IS, Sheikh S et al (2018) Simultaneous strength-ductility enhancement of a nano-lamellar AlCoCrFeNi2.1 eutectic high entropy alloy by cryo-rolling and annealing. Sci Rep 8(1):3276 98. Reddy SR, Sunkari U, Lozinko A et al (2019) Microstructural design by severe warm-rolling for tuning mechanical properties of AlCoCrFeNi2.1 eutectic high entropy alloy. Intermetallics 114:106601 99. Reddy SR, Yoshida S, Bhattacharjee T et al (2019) Nanostructuring with structuralcompositional dual heterogeneities enhances strength-ductility synergy in eutectic high entropy alloy. Sci Rep 9(1):11505 100. John R, Karati A, Joseph J et al (2020) Microstructure and mechanical properties of a high entropy alloy with a eutectic composition (AlCoCrFeNi2.1 ) synthesized by mechanical alloying and spark plasma sintering. J Alloys Compd 835:155424 101. Vikram RJ, Murty BS, Fabijanic D et al (2020) Insights into micro-mechanical response and texture of the additively manufactured eutectic high entropy alloy AlCoCrFeNi2.1 . J Alloys Compd 827:154034

90

W. Jiao et al.

102. Yu Y, He F, Qiao Z, Wang Z, Yang J (2019) Effects of temperature and microstructure on the triblogical properties of CoCrFeNiNbx eutectic high entropy alloys. J Alloy Compd 775:1376– 1385 103. Lin LF, Tsai CW (2019) Study on the damping behaviour of eutectic high-entropy alloys with lamellar structures. Philos Mag Lett 99(6):226–234

Chapter 4

Cubic Ordered Intermetallic Alloys W. C. Xiao, Y. L. Zhao, and T. Yang

1 Introduction Ordered IMs with long-range periodically ordered crystal structures have long been attractive high-temperature structural materials due to their many unique thermal– mechanical properties including excellent high-temperature strength, good oxidation, and corrosion resistance [1–3]. Such a superior performance can be largely due to their strong chemical binding and the highly restricted atomic mobility of the ordered superlattices. However, most conventional IMs exhibit remarkable strengthductility conflict at ambient temperatures, which significantly limit their widespread applications as vital structural components in many industries. For instance, topologically close-packed (TCP) intermetallic alloys are often very strong but extremely brittle at room temperature, which largely limits their practical fabricability and potential applications [4, 5]. Such an embrittlement issue can be primarily attributed to their intrinsically low-symmetry crystal structures and limited numbers of slip systems [6–8]. By comparison, some cubic ordered intermetallic alloys with the ordered L12 or B2 structures appear respectable ductile by tuning the chemical compositions or phase structures, making them potentially competent for practical structural utilization, especially at some harsh environments [6, 9–11].

W. C. Xiao · T. Yang Department of Materials Science and Engineering, City University of Hong Kong, Hong Kong, China W. C. Xiao Department of Mechanical Engineering, City University of Hong Kong, Hong Kong, China Y. L. Zhao (B) School of Materials Science and Engineering, Harbin Institute of Technology, Shenzhen, China e-mail: [email protected] T. Yang Hong Kong Institute for Advanced Study, City University of Hong Kong, Hong Kong, China © The Author(s), under exclusive license to Springer Nature Singapore Pte Ltd. 2022 Z. Jiao and T. Yang (eds.), Advanced Multicomponent Alloys, Materials Horizons: From Nature to Nanomaterials, https://doi.org/10.1007/978-981-19-4743-8_4

91

92

W. C. Xiao et al.

Notably, the silicide materials represent an important class of ordered IMs, which are expected to possess excellent creep and oxidation resistance, exhibiting significant potential for high-temperature structural applications. However, most silicide IMs as with the case for the majority of intermetallic are found to have low grain boundary strength and serious intergranular brittleness at ambient and intermediate temperatures. The cubic L12 -type Ni3 Si IMs appears to be the only one that can be made ductile and easily fabricated. Oliver and White [12] first showed that a small ductility improvement can be achieved for the Ni3 Si alloyed with ~3.2 at.% Ti. In 1989, Takasugi et al. [13] further revealed the underlying ductilization mechanisms by a comprehensive study on the effect of Ti addition on the grain boundary cohesion of the Ni3 Si intermetallic alloys, eventually producing the ductile Ni3 (Si, Ti) intermetallic materials. As reported, with the addition of Ti element, both the tensile strength and the ductility of the ternary intermetallic Ni3 (Si, Ti) are greatly improved at both room and elevated temperatures. This is because that replacing Si with Ti can reduce the difference in the electrochemical bonding environment at grain boundary planes, thereby improving the grain boundary cohesion and preventing the intergranular brittle fractures. This kind of ductilization approach is also applicable to other IMs, such as Ni3 (Al, Fe) [14] and Ni3 (Al, Mn) [15]. Likewise, the Ni3 (Si, Ti) IMs show a positive temperature dependence of yield stress as other intermetallics, which could be attributed to the Kear–Wilsdorf (K–W) locking mechanism at elevated temperatures. More importantly, because of the replacement of Si, the ternary Ni3 (Si, Ti) IMs have a relatively low density ( Mn > Fe. The macroalloying elements could segregate to the grain boundary and change the occupation landscape and/or the diffusion path of hydrogen on grain boundaries. As a consequence, the grain boundary would be strengthened effectively, which thereby brings about

94

W. C. Xiao et al.

a beneficial effect on reducing the moisture-induced embrittlement of Ni3 (Si, Ti) alloys. Furthermore, it is reported that pre-straining like the shot peening and predeformation alloy would also be effective in alleviating environmental embrittlement of the Ni3 (Si, Ti)-type intermetallic alloys [31, 32]. The deformation defects, such as dislocations and vacancies, introduced by the pre-deformation are suggested to scavenge hydrogen atoms that decomposed from moisture in the air, reducing their negative effect on the grain boundary and improving their embrittlement susceptibility.

3 Macroalloying Effect on the Microstructure Despite the variable influence of alloying elements on the environmental brittleness, alloying with additional elements is often considered to be an effective way to improve the mechanical and chemical properties of the Ni3 (Si, Ti)-type alloys, which are primarily correlated to the microstructures of the metallic alloys. Thus, understanding the microstructural evolutions with the variation of alloying elements is of great significance. The substitution behaviors of some transitional elements and their corresponding maximum solubility limits in the Ni3 (Si, Ti)-type alloys were carefully studied by Takasugi et al. [33, 34]. Before we start to review the alloying behaviors of the Ni3 (Si, Ti) alloys, let us first recall the crystalline structure and sublattice occupancy of the base alloy. In the Ni3 (Si, Ti) compound with a L12 structure, the Si and Ti atoms randomly occupy the cube corners, while Ni occupies the face center of the cubic. As shown in Fig. 1, the Ni3 (Si, Ti) can be assumed to be the Ni6 SiTi ternary compound (Pt6 CuMn structure type, cP32, space group Fm33m), in which Si and Ti atoms form ordered structure in their sublattice sites. The heats of formation between Ni6 SiTi and Ni6 SiX (or Ni6 XTi or X6 SiTi) and the difference in atom size between matrix atoms and the quaternary element X were measured to predict the substitution behavior and maximum solubility limits [33], shown in Fig. 2. Figure 2 shows the calculated heat of formation for three hypothetical types of structure configuration based on an extended Miedema’s equation [33, 35, 36]. By comparing the values of heat of formation, Cr, W, Mn, Re, Fe, Os, Co, Ir, Pt, Cu, and Au tend to substitute for Ni atoms. Zr, Hf, V, Nb, Ta, Mo, and Al tend to substitute for Ti atoms. Ga is likely to substitute for Si or Ti atom. Compared with the value of the host compound (the horizontal dashed line, ΔH (Ni6 SiTi) = −36:02 kJ/mole of atoms), Zr, Hf, Nb, and Ta were more likely substitute Ti atoms to form the Ni6 SiX structure, while Ir and Pt tend to replace with Ni and form the X6 SiTi structure. Table 1 further displays the phase composition and solubility limit of the Ni3 (Si, Ti) IMs alloyed with different elements [33, 34, 37–44]. As can be seen, the solid solubility of most elements is less than 2 at.%. Noting that the solubility of Nb and Ta was higher than other elements, which was primarily attributed to the smaller atomic size misfit with Ti atoms compared to other elements. In general, the solubility

4 Cubic Ordered Intermetallic Alloys

95

Fig. 1 Crystal structures of a Ni3 (Si, Ti) with L12 structure and b Ni6 SiTi with Pt6 CuMn structure. In (a), open circles correspond to Ni atoms, and full circles correspond to Si or Ti atoms. In (b), open circles correspond to Ni atoms, full circles correspond to Si atoms, and shaded circles correspond to Ti atoms [33]

Fig. 2 Calculated heats of formation for three structure configurations [33]

limits of the quaternary elements substituted for the Ti site in Ni3 (Si, Ti) are ranked experimentally as Nb > V > Cr > Mn > Fe > Hf > Zr, which is correlated with the size misfit parameter between Si elements and the quaternary element X, and the difference in the formation enthalpy between Ni3 (Si, Ti) and Ni3 (Si, X) structures [34]. Since a single-phase L12 -type microstructure is desirable for obtaining superior high-temperature tensile properties, Ta, Al, and Nb were the most suitable substitution elements to form a continuous Ll2 -type solid solution within their solubility limit according to the Table 1. Second phases are expected to form when the element contents exceed their solute limitation, which were identified as the D019 -type Ni3 Zr, D024 -type Ni3 (Ti), D0a-type Ni3 (Ta), Ni3 Nb, and Ni3 Hf, HCP-type Re-solid solution. For example, as shown in Fig. 3, when the addition of Ta is above the solubility

96

W. C. Xiao et al.

Table 1 Phase composition and solubility limit of the Ni3 (Si, Ti)-type IMs alloyed with different elements [33, 34, 37–44] Alloy

Microstructure

Solubility limits

Alloy

Microstructure

Solubility limits

Mo(Ti)

L12 + Niss