Improving mechanical properties of an FCC high-entropy alloy by γ′ and B2 precipitates strengthening

Journal Pre-proof Improving mechanical properties of an FCC high-entropy alloy by γ′ and B2 precipitates strengthening Ziyong Li, Liming Fu, Jian Peng, Han Zheng, Xinbo Ji, Yanle Sun, Shuo Ma, Aidang Shan PII:

S1044-5803(19)32083-2

DOI:

https://doi.org/10.1016/j.matchar.2019.109989

Reference:

MTL 109989

To appear in:

Materials Characterization

Received Date: 6 August 2019 Revised Date:

10 October 2019

Accepted Date: 29 October 2019

Please cite this article as: Z. Li, L. Fu, J. Peng, H. Zheng, X. Ji, Y. Sun, S. Ma, A. Shan, Improving mechanical properties of an FCC high-entropy alloy by γ′ and B2 precipitates strengthening, Materials Characterization (2019), doi: https://doi.org/10.1016/j.matchar.2019.109989. This is a PDF file of an article that has undergone enhancements after acceptance, such as the addition of a cover page and metadata, and formatting for readability, but it is not yet the definitive version of record. This version will undergo additional copyediting, typesetting and review before it is published in its final form, but we are providing this version to give early visibility of the article. Please note that, during the production process, errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain. © 2019 Published by Elsevier Inc.

Improving mechanical properties of an FCC high-entropy alloy by γ′ and B2 precipitates strengthening Ziyong Li1,2,3, Liming Fu1,2,3*, Jian Peng1,2,3, Han Zheng1,2,3, Xinbo Ji1,2,3, Yanle Sun1,2,3, Shuo Ma1,2,3, Aidang Shan1,2,3* 1. School of Materials Science and Engineering, Shanghai Jiao Tong University, 800 Dongchuan Road, Shanghai, 200240, People’s Republic of China 2. Collaborative Innovation Center for Advanced Ship and Deep-Sea Exploration (CISSE), Shanghai, 200240, People’s Republic of China 3. Shanghai Key Laboratory of High Temperature Materials and Precision Forming, Shanghai Jiao Tong University, Shanghai, 200240, People’s Republic of China Corresponding authors Liming Fu, Research Fellow, Tel: 86-21-54748974, Fax: 86-21-54740825, Email: [email protected]

Aidang Shan, Professor, Tel: 86-21-54747556, Fax: 86-21-54740825, E-mail: [email protected]

Abstract The FCC-structured high entropy alloys (HEAs) possess exceptional ductility and fracture toughness, but they generally exhibit insufficient strength for engineering applications. In this work, a precipitation strengthening non-equimolar (FeCoNi)81Cr9Al8Ti1Nb1 HEA was fabricated via arc melting and subsequent two-step aging treatment. The microstructure, phase constitution and mechanical properties of this alloy during aging were investigated systematically. The results indicate that coherent γ′ and incoherent B2 are major precipitates for the alloy under the two-step aging. An excellent balanced tensile property is achieved at room temperature even with extensive B2 grain boundary coverage. Quantitative calculations of the individual strengthening effects demonstrate that particle (γ′) shearing mechanism is the predominant strengthening mechanism. The high work-hardening capability of the FCC matrix could greatly suppress the propagation of microcracks originated at these brittle B2 phases and promote a retention of ductility. In addition, this alloy exhibits outstanding high-temperature tensile properties up to 700 ℃. It is attributed to the high thermal stability of the γ′ precipitates as well as the pinning effect of the grain-boundary B2 phases on the grain boundary. Present work will focus on optimizing of the alloy design of 1

HEAs and the precipitation strengthening of HEAs for high-temperature structural applications.

Keywords: High entropy alloy; Precipitation strengthening; γ′ and B2 precipitate; Microstructures; Mechanical properties; Strengthening mechanism

1. Introduction High entropy alloys (HEAs) with new compositions design approach, exhibiting unique microstructures characteristics, and unusual properties, have been attracting extensive attention in the past decades and a half

[1-11]

. Despite containing multiple

principal elements, the HEAs can readily form solid solutions with simple crystal structures

[2, 3]

. Typically, the FeCoCrNiMn based HEA with a face-centered cubic

(FCC) structure has been widely investigated due to excellent ductility and exceptional fracture toughness from room temperature to cryogenic temperatures [3, 4]. However, compared with most existing structural metallic materials, the FCC HEAs usually exhibit an insufficient strength for engineering applications [3-5]. Consequently, extensive efforts have been dedicated to improving the strength of the FCC-structured HEAs, such as by introducing phase transformation-induced hardening twinning-induced plasticity

[7]

, adopting precipitates

[8]

, interstitial atom

[9]

[6]

,

, grain

refinement [10] and multi-scale grains [11]. Precipitation strengthening has been extensively used to increase the strength of commercial alloys, e.g., Ni-based superalloys

[12-14]

. Essentially, precipitation

strengthening can be achieved by producing a particulate dispersion of obstacles to dislocation movement. Such an effect is strongly dependent upon the morphology, volume fraction, size, and distribution of the precipitates

[14]

. He et al.

[15]

demonstrated that minor additions of Al and Ti to the FeCoCrNi HEA can induce a large amount of coherent nanoscale Ni3Al-type γ′ phases formation in the FCC matrix, which drastically improves the strength of the HEA without compromising its tensile ductility. Zhao et al. [16] reported that the γ′ precipitates in HEA exhibit a much slower coarsening kinetics than that in the conventional Ni-based alloys owing to the 2

sluggish diffusion. The feature of high thermal stability of γ′ precipitates makes γ′-strengthened HEA to be new promising structural materials for high-temperature applications. In addition, it was found that the increased Al addition to FeCoCrNi-based HEAs facilitates the formation of NiAl-type B2 phase in addition to the γ′ precipitates[17-24].The B2 phases in many HEAs are mainly distributed at the grain boundary or the intragranular region, and the grain-boundary B2 can effectively suppress grain growth during annealing[23, 24]. NiAl is a brittle intermetallic compound at room temperature, but it is of interest for the development of high-temperature alloys due to its high melting temperature and excellent oxidation resistance

[25]

Recently, the NiAl B2 phase also have been utilized to strengthen the FCC HEAs 22]

. Diao et al.

[22]

.

[18,

developed a NiAl-strengthened Al0.3FeCoCrNi HEA via aging

processing (700 ℃ for 500 h), which gives a yield strength of 321 MPa with an elongation of 29 % at room temperature. Lu et al.

[18]

reported that a eutectic

AlCoCrFeNi2.1 HEA containing a fine lamellar FCC/B2 microstructure exhibits an excellent mechanical property with the fracture strength and elongation of 944 MPa, 25.6 %, and 538 MPa, 22.9 %, at room-temperature and 700 ℃, respectively. Owing to the brittleness of NiAl phase can be relieved by combining it with FCC phase, the NiAl-strengthened HEAs still retain a good ductility

[22, 25]

. These observations

demonstrated that γ′ and NiAl B2 can effectively improve the mechanical properties of HEAs individually. A recent study by Gwalani et al.

[19]

suggested that the

Al0.3FeCoCrNi HEA can be strengthened by combining γ′ and B2 phases after annealing at 620 ℃ for 50 h. Compared to the single FCC Al0.3FeCoCrNi HEA, the alloy with γ′ and B2 phases exhibits an excellent combination of yield strength (~ 490 MPa), ultimate tensile strength (~ 850 MPa), and ductility (~ 45% elongation), at room temperature. This research shows a new precipitation strategy for HEAs that can take advantage of both types of ordered phases simultaneously. It can be predictable that the HEA including high thermal stability of γ′ phase and high solvus temperature of B2 phase could possess an outstanding high-temperature mechanical property. However, the γ′ precipitates in Al0.3FeCoCrNi HEA are only stable below 620 ℃ and 3

then are replaced by ordered B2 precipitates above 700 ℃ [19]. It means that γ′ and B2 phases in Al0.3FeCoCrNi HEA can coexist at a relatively low temperature. The high thermal stability of the precipitates in the matrix is crucial to keep its strong reinforcing role, especially at high temperature

[16]

. Therefore, it is necessary to

improve the thermal stability of both γ′ and B2 phases to keep them stable and useful at high temperatures. In this study, a novel FeCoCrNi based HEA with Al, Ti and Nb additions is presented, and a two-step aging process is conducted to this HEA. The objective of the study is to promote the sable γ′ and B2 phases precipitation in the FCC matrix. After two-step aging, high density of γ′ phases are formed in the FCC matrix, and the B2 phases are mainly precipitated along grain boundary. The two-step aged alloy exhibits an outstanding balanced strength and tensile ductility at room and elevated temperatures, especially, the high-temperature tensile properties superior to most of the FeCoCrNi-based HEAs reported previously. The design methodology in optimizing the HEA with γ-γ′-B2 three phases is discussed. In addition, particular emphasis is focused on the characterization of the alloy microstructures and the effects these features on the mechanical properties.

2. Material design The thermal stability of the precipitates in the alloys are sensitive to the composition and heat treatments

[16, 19, 26]

. It is generally believed that Al element

alone is insufficient to stabilize γ′ precipitates in HEA systems, while co-additions of Al and Ti can greatly promote the precipitation and stabilization of the γ′ precipitates [26]

. Compared to the Al0.3FeCoCrNi HEA, the thermal stability temperature of the γ′

phase in (FeCoCrNi)94Al4Ti2 HEA is increased to at least 800 ℃ [16, 19]. Besides the Ti element, Nb element also is preferentially toward to the γ′ phase by substituting the Al sublattice. Addition of Nb can stabilize the γ′ precipitates by increasing its solvus temperature as well as its volume fraction [12, 27]. It reported that the γ′ precipitates are formed in a FeCoCrNi based HEA with addition of Al and Nb, and the γ′ volume 4

fraction is increased with increasing of Nb content

[28]

. Many Ni based superalloys

often contain Al, Ti and Nb elements to produce an optimized volume fraction and distribution of the γ′ precipitates, and then improving the mechanical properties 29]

[12, 13,

. Therefore, co-additions of Al, Ti and Nb to the current HEA are adopted to

stabilize γ′ precipitation in the FCC matrix and to achieve high strength at room or elevated temperatures. The chemical composition of the alloy is provided in Table 1. Considering that Al element is B2 and γ′ phases forming element, so adequate content of Al is added to the HEA. High Ti and Nb concentrations in the matrix could promote the precipitation of detrimental topologically close packed (TCP) phases, e.g., η and δ phases

[12, 13]

. Thus, minor Ti and Nb are added to the current HEA. In

addition, many research works have demonstrated that the single-FCC HEAs occur the phase decomposition during annealing at intermediate temperatures, e.g., FeCoCrNiMn

[30]

and FeCoCrNi

[31]

, and the Cr-rich σ phases were detected in both

HEAs. In the present study, the concentrations of the σ phases forming element Cr is reduced to mitigate the formation of this detrimental phase. The HEA defining parameters of the present HEA are calculated, i.e., mixing enthalpy (∆Hmix), atomic size difference (δ), valence electron concentration (VEC), and mixing entropy (∆Smix). These parameters are used to characterize the collective behavior of the constituent elements in multicomponent alloy

[32-34]

. It is generally

accepted that the formation ability of the solid solution phase in HEAs is combined control by ∆Hmix, δ, VEC and ∆Smix, rather than by the only factor of ∆Smix. Basing on the reported empirical parameters in literature

[33]

, the formation of solid solution

phase requires that ∆Hmix, δ, and ∆Smix simultaneously satisfy -22 ≤ ∆Hmix ≤ 7 kJ / mol, 0 ≤ δ ≤ 8.5, 11 ≤ ∆Smix ≤ 19.5 kJ /mol. Moreover, the larger VEC (≥8) favors the formation of FCC-structured solid solution, while the smaller VEC (≤6.89) favors the formation of BCC (body-centered cubic)-structured solid solution [34]. Table 2 lists the calculation results of these parameters for the present HEA. The calculation method of the parameters can be found elsewhere

[32, 33]

. It is obvious that all parameters are

within the range of requirements for the formation of an FCC-structured solid 5

solution.

Table 1 Chemical composition of the homogenized HEA Elements

Fe

Co

Ni

Cr

Al

Nb

Ti

Nominal, At. %

27.00

27.00

27.00

9.00

8.00

1.00

1.00

Measured, At. %

27.23

26.64

26.83

9.25

8.13

0.89

1.03

Table 2 Calculation results for this HEA alloy ∆Hmix (kJ / mol)

∆Smix (JK / mol)

VEC

δ

-8.64

13.06

8.11

4.47

3. Experimental procedures The non-equimolar (FeCoNi)81Cr9Al8Ti1Nb1 HEA under investigation here was prepared using a vacuum arc melting furnace from pure metals (> 99.9 wt.%) under a Ti-gettered high-purity argon atmosphere. The button alloys were re-melted at least five times to ensure the alloying elements uniformly distribute throughout the alloy. Then, the arc-melted button alloys were casted to a rectangular ingot using a water-cooled copper mold. Following melting, the ingots were homogenized at 1190 ℃ for 6 h followed by water-cooling (WC). To achieve a uniform full-recrystallized FCC structure, the homogenized ingot was cold-rolled with 80 % reduction in thickness and then annealed at 1100 ℃ for 1 h followed by water-cooling. Subsequently, two-step aging process was employed to heat treatment of this HEA, that is aging at 845 ℃ for 4 h and air-cooling (845 ℃/4 h, AC), and then aging at 720 ℃ for 10 h and air-cooling (845 ℃/4 h, AC + 720 ℃/10 h, AC). The sample identification (ID) and the corresponding processing conditions are summarized in Table 3.

6

Table 3 Sample ID and the corresponding processing conditions of the HEAs Processing conditions

Samples ID

As-cast

A

Homogenization (1190 ℃ / 6 h, WC)

B

Annealing (1100 ℃ / 1 h, WC)

C

One-step aging (845 ℃/4 h, AC)

S1

Two-step aging (845 ℃/4 h, AC + 720 ℃/10 h, AC)

S2

Phase identification was conducted using a Rigaku D/max 2500 X-ray diffraction (XRD) instrument by Cu-Kα radiation, with a scanning 2θ range of 20 to 100 deg and a scanning rate of 2 deg/min. The microstructures were characterized using a ZEISS optical microscope (OM), a JSM-7600F field emission scanning electron microscopy (SEM) coupled with energy dispersive spectrometer (EDS), and a JEM-2100F transmission electron microscopy (TEM). The casted, homogenized and annealed specimens after polishing were etched in a solution consisting of 100 ml HCl, 100 ml C2H5OH, and 5 g CuCl2 for OM and SEM (secondary-electron mode) observation. The aged specimens after polishing were examined using SEM (black-scattered electron mode) to distinguish the B2-NiAl phase based on atomic number contrast. TEM specimens were prepared as follows: the alloys were carefully ground to foils with a thickness of about 80 um, punched to Ф3 mm circular sheets and then further ground to about 40 um, and finally followed by twin-jet electrochemical polishing using an electrolyte of 5 % perchloric acid and 95 % ethanol with a voltage of 30 V at -30 ℃.

For quantitative analysis of the γ′ and B2 phases, particle density and

average particle diameter were analyzed using Image-Pro Plus image processing 7

software. At least five SEM (black-scattered electron mode) or TEM (dark field) images and 300 particles were examined to evaluate particle diameter and density. Dog-bone tensile specimens were machined from the annealed plates by electro-discharge cutting machine, following the specimens were carefully ground to obtain a bright and clean surface. The tensile specimens with a gauge length of 15 mm and a cross-section of ~ 4 × 0.9 mm2. The room-temperature tensile tests were carried out on a Zwick/Roell universal testing machine equipped with a clip-on extensometer for the strain measurement at a strain rate of 5×10-4 s-1. The high-temperature tensile at 600, 630, 650 and 700 ℃ were performed on a Shimadzu AG-10kNA tension tester under an initial strain rate of 5×10-4 s-1, and the strain was measured using a non-contact laser extensometer. Three tensile samples for each condition were tested to ensure good reproducibility. For the high-temperature tensile, the samples were first heated to desired temperature and then held there for 10 min before the start of the test. After tensile testing, the fracture surfaces were examined by SEM, and the fracture mechanism of these alloys was determined.

4. Results and discussion 4.1 Microstructures The JMatPro® 8 software is used to predict the equilibrium phases of the (FeCoNi)81Cr9Al8Ti1Nb1 HEA, as shown in Fig. 1. It can be seen that the transition of this HEA from the liquid phase to the γ phase initiate at 1370 ℃ and finish at 1298 ℃. With decreasing temperature, B2 and γ′ phases are formed at the temperatures of 1020 and 904 ℃, respectively. The BCC and sigma (σ) phases are formed at relatively low temperatures of 625 and 570 ℃, respectively. XRD diffraction patterns of the HEA specimens for different processing states are presented in Fig.2. A single FCC structure is observed from the A alloy and B alloy. The C alloy still retains a single FCC structure. By contrast, for the S1 alloy, the BCC and two weak diffraction peaks are detected, which are identified as B2-NiAl and γ′-Ni3Al phases, respectively. The 8

S2 alloy shows that the diffraction intensity of B2 and γ′ phases is increased as compared with that of the S1 alloy, implying the content of both phases is increased. The phase constituents detected using XRD for this HEA alloy in difference processing states essentially meet the predicted results of the thermodynamic calculation.

Fig.1. Phase equilibrium calculated by JMatPro® 8.

9

Fig. 2. XRD diffraction patterns of the HEAs for different processing states. A: casted alloy; B: homogenized alloy; C: cold-rolled alloy annealed at 1110 ℃ for 1 h; S1: one-step aged alloy (845 ℃/4 h, AC); S2: two-step aged alloy (845 ℃/4 h, AC+720 ℃/10 h, AC).

Fig.3 shows micrographs of the HEAs for different processing states. The A alloy exhibits a dendritic structure, as shown in Fig.3a. The B alloy shows a significant coarser grain structure, as observed in Fig.3b. The C alloy displays the typical equiaxed grain structure with a grain size of about 35 µm and no second phases are observed, as shown in Fig.3c. For the S1 alloy, it is clearly seen that the second phases in the shape of rhombus/short-rod-like precipitates are formed in the matrix, they are largely distributed at grain boundary, with a few in the intragranular region, as presented in Fig.3d-f. According to the above XRD analysis results of the samples in Fig.2, these precipitates could be B2 phases with a BCC structure.

10

Fig.3. Micrographs of the HEAs for different processing states. (a) OM image of the A alloy; (b) SEM secondary-electron image of the B alloy;(c) SEM secondary-electron image of the C alloy; (d), (e) and (f) SEM back scattered-electron (BSE) images of the S1 alloy (unetched) in different magnifications.

Fig.4 shows SEM images and EDS analysis of the S2 alloy. Comparing the microstructure of the S1 alloy (Fig.3e), the content and size of B2 precipitates are slightly increased, as displayed in Fig.4a. Fig.4b is the high-magnification SEM image of the triple junctions, showing that B2 precipitates are distributed along the grain boundary. EDS mapping indicates that these B2 precipitates are significantly enriched in Al. The composition of these B2 phases is measured 11

using the spot EDS, see the inset in Fig.4a. Compared to the composition of the homogenized alloy (Table 1), the Al element in the B2 phase is enriched to 20 at. %, while other constituent elements are not significantly changed. It has proved that the Al content ratio has a significant effect on lattice types and microstructures for HEAs

[35, 36]

. The addition of Al in FeCoCrNi HEA usually

changes the primary solid-solution phase from FCC to FCC+BCC to BCC, it is generally attributed to the larger atomic size of Al destabilizes the FCC lattice and introduces BCC phase [35, 36].

Fig.4. SEM-BSE images and the energy dispersive spectrometer (EDS) analysis of the S2 alloy (unetched).

Fig.5 shows TEM images of the S1 alloy. Fig.5a is the BF image taken from 12

the grain boundary region, which shows the precipitates are located at the grain boundary. The inset in Fig.5a is the [011] BCC zone axis SAED pattern taken from a grain-boundary precipitate. The superlattice spots (marked by white circles) are observed in the diffraction pattern. It indicates that these grain-boundary precipitates are an ordered phase with B2-type ordering. Fig.5b is the DF image recorded from the selected superlattice spot (100) in the SAED pattern (inset in Fig.5a). It can be seen that the rod-like B2 precipitates are distributed in chain shapes at the grain boundary. The [110]

FCC

zone axis SAED pattern (inset in

Fig.5b) taken from a grain-boundary B2 precipitate together with the neighboring FCC matrix. Kurdjumov-Sachs (K-S) orientation relation is identified between the B2 phase and FCC matrix in agreement with the results of other researchers [19, 20]. The large lattice misfit (~ 19.9 %) between the FCC matrix and B2 phase was reported, indicating FCC and B2 are incoherent

[22]

. Fig.5c-d are the BF images

taken from the intragranular region, besides the rhombus-like B2 phases precipitate in the matrix, another fine spherical precipitate can be observed. Fig.5e is the [011]

FCC

zone axis SAED pattern taken from the matrix in Fig.5d.

Additional superlattice reflections are detected in the diffraction pattern, indicating the formation of coherent ordered L12 (γ′) precipitates in the matrix. Fig.5f shows the DF image taken from the superlattice spot ( 100 ) in the SADE pattern (Fig.5e). The average grain diameter of these spherical γ′ precipitates is about 39 nm. In relative to Al0.3FeCoCrNi HEA in which the γ′ precipitates are only sable blow 700℃

[17, 20]

, the current HEA exhibits a high precipitation

temperature and the γ′ precipitates can stable up to 845 ℃. Therefore, the addition of Nb and Ti to substitute partial Al can greatly improve the stability of γ′ phase.

13

Fig.5.TEM images of the S1 alloy. (a) bright field (BF) image from the grain boundary region showing the grain-boundary B2, the inset shows the SAED pattern taken from a grain-boundary precipitate marked as the red circle in Fig.5a; (b) dark field (DF) image from the selected superlattice spot (100) from the SAED pattern (inset in Fig.5a), the inset shows the SAED pattern taken from a grain-boundary precipitate together with the neighboring FCC matrix marked as the white circle in Fig.5a; (c) BF image from the intragranular region showing the intragranular B2; (d) BF image form the intragranular region showing the γ′ precipitates; (e) the corresponding SAED pattern from Fig.5d; (f) DF image from the selected superlattice spot ( 100 ) in the SAED pattern (Fig.5e).

Fig.6 displays TEM images of the S2 alloy, the rob-like B2 phases precipitate in a chain shape along the grain boundary, as shown in Fig.6a. Fig.6b shows the BF of the long-rob-like precipitate in the intragranular, and the inset is the [011] BCC zone axis SAED pattern taken from this precipitate. The superlattice spots appear in the diffraction pattern. Additionally, a DF TEM image (Fig.6c) of the precipitate can be taken from the superlattice reflection (010). It demonstrates that this precipitate also is the B2 phase. As compared with the S1 alloy (Fig.5a-c), the size of grain-boundary B2 precipitates is not apparently changed, while the size of intragranular B2 is increased. Fig.6d shows the BF image from the FCC matrix, in which high-density 14

precipitates are observed. Fig.6e is the [011]FCC zone axis SAED pattern taken from the corresponding region in the Fig.6d. The presence of the super-lattice reflections in the SAED pattern confirms the formation of L12 (γ′) precipitates. These γ′ precipitates are highlighted in the DF image as shown in Fig.6f, which are recorded from a selected superlattice spot (100) in the SAED pattern (Fig.6e). The DF image shows that the spherical γ′ precipitates are well dispersed in the FCC matrix and have two different sizes. The average grain size of the large precipitates is ~ 43 nm, while the smallest one is ~19 nm. According to above TEM observation of the S1 alloy (Fig.5f), it can be deduced that the large γ′ precipitate was primarily formed at 845℃, the smallest one was formed at 720 ℃. The γ′ precipitates in some superalloys usually grow from the initial spheroid to cuboidal shape

[37]

. Many researchers reported that the morphology of γ′ precipitates

in HEAs usually shows spheroidal shape

[8, 15, 16]

. Similarly, the spheroidal γ′

precipitates are observed in current HEAs after two-step aging. The equilibrium morphology of precipitates is determined by minimizing the sum of elastic strain energy and interfacial energy

[38]

. The former is originated from lattice mismatch at

the precipitate/matrix interface, while the latter is caused by the interfaces between precipitate and matrix. The fine precipitates possess low coherent strain energy; hence the shape is mainly controlled by minimizing the interfaces area of the precipitates, resulting in spheroidal precipitates. However, with the growing of the precipitates, the shape of precipitates is gradually dominated by the elastic strain energy

[39]

. It has

been experimentally demonstrated that the spheroidal γ′ precipitates (diameter ~70 nm) of the (FeCoCrNi)94Ti2Al4 HEA after aging at 800 ℃ for 503 h was mainly controlled by the interfacial energy

[16]

. These γ′ precipitates remain spheroidal and nanoscale

size even after being aged for a long time, indicating a relatively low driving force for coarsening. Comparing the activation energy of γ′ precipitate coarsening with the diffusion activation energy of individual elements in the (FeCoCrNi)94Ti2Al4 HEA, it is suggested that the precipitate coarsening behavior should be controlled by the diffusion of either Al or Ti in the alloy matrix 15

[16]

. In the present work, Al, Ti and Nb

elements are co-added to the FeCoCrNi based HEA, their atoms are all larger than other constituent elements. According to the traditional point of view, the larger atom produces a larger lattice strain during elementary atom jump, resulting in an increase in migration barrier and ultimately lower diffusion mobility [40]. For the two-step aged (S2) alloy, the large γ′ precipitates obtained at one-step aging stage (845 ℃ for 4 h) only show a small degree coarsening from initial grain size ~39 to ~43 nm after two-step aging stage (720 ℃ for 10 h). Meanwhile, the smallest precipitates with size ~19 nm are formed. Apparently, the γ′ precipitates in the current HEA also exhibit excellent thermal stability.

Fig.6. TEM images of the S2 alloy. (a) BF images from grain boundary regions showing the grain-boundary B2 ; (b) BF image from the intragranular region showing the intragranular B2, the inset shows the SAED pattern of the intragranular B2; (c) DF image form the selected super lattice spot (010) in the SAED pattern (inset in Fig.6b); (d) BF image form the intragranular region showing the γ′ precipitates; (e) the corresponding SAED pattern from Fig.6d; (f) DF image from the selected super lattice spot (100) in the SAED pattern (Fig.6e).

4.2 Room-temperature tensile properties Fig.7 presents the representative room-temperature tensile stress-strain curves of 16

the HEA alloy in different processing state, and the tensile properties are listed in Table 4. The yield strength (YS), ultimate tensile strength (UTS) and fracture elongation of the C alloy are 280, 701 MPa (σUTS - σYS = 420 MPa, σYS/σUTS = 0.40) and 47 %, respectively. This FCC-structured (C) alloy exhibits a high working hardening ability and an excellent uniform deformation. The YS, UTS and fracture elongation of the S1 alloy are 491, 863 MPa and 32 %, respectively. The one-step aged (S1) alloy shows a moderate improvement in strength but with a slight loss of elongation compared to the C alloy. The presence of B2 and γ′ precipitates in the FCC matrix can improve the alloy strength. However, the large size of B2 precipitates and low-density of γ′ precipitates have limited effectively impede dislocations motion; thus, the strength increase is not significant. For two-step aged (S2) alloy, the YS and UTS are 863 and 1285 MPa, respectively, whilst the elongation still remains at a respectable value of 28 %. As compared to the C alloy, the S2 alloy is increased by 117 % and 95 % in the YS and UTS, respectively, and still exhibits a strong work-hardening capability (σUTS - σYS = 422 MPa, σYS/σUTS = 0.67). It is interesting that even with extensive brittle NiAl-type (B2) phases grain boundary coverage, the alloy still exhibits a room temperature elongation of 28 %. The high work-hardening capability of the FCC matrix could greatly suppress the propagation of microcracks originated at these brittle phases and promote retention of ductility. It has been reported that the HEA containing brittle σ and µ phases still exhibits a high strength but without causing a serious embrittlement owing to the ductility FCC matrix

[41]

.

The pronounced increase in strength is mainly attributed to the formation of high-density γ′ precipitates, and though B2 phases play a role in strengthening, its contribution seemed to be very small. The representative fracture surfaces of the C and S2 alloy are presented in Fig.8a and Fig.8b, respectively, in which both of them show a feature with the intragranular ductile fracture. This suggests that the brittle B2 phases are not seriously weaken the strengthening of the grain boundaries. Therefore, this HEA with γ-γ′-B2 three phases achieves a combination of high strength and good ductility at room temperature. 17

Fig.7. Stress-strain curves of the HEAs in different processing states at room temperature. The insets show the sketches of microstructure in the C, S1 and S2 alloys.

Fig.8. Fracture surface of the C (a) and S2 (b) alloy at room temperature.

18

Table 4 Tensile property of the HEAs at room and high temperature. Samples ID

Tensile temperature

σYS (MPa)

σUTS (MPa)

εf (%)

C

Room

280.7

701.8

47.1

S1

Room

491.5

863.5

32.6

S2

Room

863.5

1285.9

28.2

S2

600 ℃

619.6

939.9

35.9

S2

630℃

605.3

880.5

38.3

S2

650℃

584.3

823.4

38.8

S2

700℃

557.2

714.3

40.3

To

provide insight

into

the strengthening

contributions

of different

microstructural features in present HEA. A simplistic strength model is proposed to estimate the yield strength ( σ y ) and can be expressed as the following equation:

σ y = σ A + ∆σ B 2 + ∆σ γ ′ (1) where σ y is the yield stress, σ A is the combined strengthening contributions from solid solution, grain size (Hall-Petch) and dislocations. ∆σ B 2 and ∆σ γ ′ are the precipitation strengthening contributions from B2 and γ′ phases, respectively. The C alloy is comprised of a single-phase FCC structure. Generally, the fully annealed structure from a relatively high temperature has few dislocations et al.

[15]

[15, 42]

. He

estimated the solid-solution hardening contribution to the strength in

FeCoCrNi HEA by Ti (2 at. %) and Al (4 at. %) additions, suggesting that a small value (25.4MPa) account for the strength enhancement. Hence, the solid-solution strengthening and dislocations strengthening may not be the dominant strengthening mechanism in the present alloy. A high Hall-Petch coefficient was reported for many HEAs [19, 43], thus grain-boundary strengthening could have a great contribution to an increase in strength. The yield strength can be described by the classical Hall-Petch equation [19]:

σG = σ0 + 19

Ky d

(2)

Where σ G is grain-boundary strengthening, d is the average grain size, σ 0 is the friction stress and K y is the Hall-Petch coefficient. In this work, the value of σ 0 and K y are adopted from the Al0.3FeCoCrNi HEA, that is 95 MPa and 824 MPa.um1/2, respectively. The measured average grain size of the C alloy is ~ 35 um, and consequently, the grain-boundary strengthening is estimated to be ~ 234 MPa. Hence, it can be assumed that 280 MPa (YS of the C alloy) – 234 MPa = 46 MPa is the strengthening contribution from the solid-solution and dislocations. For the one/two-step aged alloy, the strength increment is ascribed to the formation of B2 and γ′ precipitates in the FCC matrix. Generally, the precipitation strengthening is governed by either a dislocation bypassing mechanism (Orowan-type) or particle shearing mechanism

[15, 44]

. Orowan mechanism occurs when the radius of

particles exceeds a critical value or is incoherent with the matrix. The shearing mechanism dominates when precipitates are sufficiently small and coherent. As discussed earlier (Fig.5 and 6), the B2 and γ′ phases are incoherent and coherent with the FCC matrix, respectively. For the B2 phases, the Orowan dislocation bypassing mechanism ( ∆σ Orowan ) contribution to yield strength increment can be described as [44]

: ∆σ Orowan = M ⋅

0.4Gb ln(2r / b) λP π 1 −υ

(3)

Where M, G, b, and υ are defined in Table 5, r is the mean radius of the circular cross-section in a random plane for the precipitate, r = 2 r , where r is the mean 3 radius of the precipitates, λP is mean edge-to-edge interprecipitate spacing,



π

 − 1 , where f is the volume fraction of the B2 precipitates. In the  4f 

λP = 2r 

shearing mechanism of γ′ precipitates, three contributing factors to the increase in yield strength are used, i.e., coherency strengthening ( ∆σ CS ), modulus mismatch strengthening ( ∆σ MS ) and order strengthening( ∆σ OS ). The larger of ∆σ CS + ∆σ MS 20

or ∆σ OS is the total strength increment from the dislocation shearing mechanism[15, 44, 45]

. The shearing mechanism for the contribution to yield strength increasement are

calculated as following equations [44]: 1

∆σ CS = M ⋅ α ε ( G ⋅ ε )

3 2

 rf  2    0.5Gb  1

∆σ MS = M ⋅ 0.0055 ( ∆G )

3 2

(4) 3m

 2 f 2  r  2      G  b

−1

(5)

1

∆σ OS = M ⋅ 0.81

γ apb  3π f  2   2b  8 

(6)

where M, G, and b are listed in Table 5, Ɛ is the constrained lattice parameter mismatch, α ε and m are the constant, ∆ G is the modulus mismatch between the matrix and the precipitates, f is the volume fraction of the γ′ precipitates, and γ apb is the antiphase boundary free energy of the precipitates. The physical meaning and values of the symbols in the Eq. 3-6 used in the current study are summarized in Table 5. The fraction and size of precipitates in one/ two-step aged (S1 and S2) alloy are evaluated from the microstructural observation and are listed in Table 6. The calculated values of contributions to yield strength from B2 and γ′ precipitations strengthening are listed in Table 7. For the S1 alloy, ∆σ Orowan is 39.3 MPa, ∆σ CS , ∆σ MS and ∆σ OS are 141.1, 17.4, and 157.2 MPa, respectively. The measured strength increments from Orowan dislocation bypassing mechanism and particle shearing mechanism are 39.3 and 159.1 MPa, respectively. Consequently, the overall strengthening contributed by B2 and γ′ precipitates in S1 alloy is 39.3 + 159.1 = 198.4 MPa. For the S2 alloy, ∆σ Orowan is 48.6 MPa, ∆σ CS , ∆σ MS and ∆σ OS from the large γ′ are 142.2, 17.3, and 151.4 MPa, respectively, however, that are 227.1, 33.0, and 362.6 MPa from the smallest γ′, respectively. The total strengthening contributed by B2 and γ′ precipitates in the S2 alloy reach to 48.6 + 159.5 + 362.6 = 570.7 MPa. The value of σ A in the Eq. 1 can be approximately taken as the YS 21

(280.7 MPa) of the C alloy. Thus, the estimated YS of the S1 and S2 alloy are 479.1 and 851.4 MPa, respectively, which are in reasonable agreement with the experimental values (listed in Table 4). Noteworthy, in the S2 alloy, precipitation strengthening account for the largest strength improvement, and still maintain a good ductility. It could be attributed to various strengthening mechanisms: solid-solution strengthening and grain-boundary strengthening, and in particular, the precipitation strengthening from γ′ phases. Table 5 Physical meaning and values of different symbols used in the strengthening mechanism calculations. Symbol

Meaning

Values

Unit

σ0

Friction stress

= 95 for FCC Al0.3CoCrFeNi[19]

MPa

ky

Hall-Petch coefficient

= 824 for FCC Al0.3CoCrFeNi[19]

MPa.um1/2

b

Burgers vector

= 0.255 for FCC FeCoCrNiMn[46]

nm

G

Shear modulus

= 80 for FCC FeCoCrNiMn[47]

GPa

∆G

Modulus mismatch between

= 4 for Ni-based superalloys[48]

GPa

= 0.12 for Ni-based superalloys[48]

J.m-2

= 3.06 for the FCC polycrystalline

Dimensionless

matrix and precipitates Antiphase boundary energy

γapb

of the precipitates M

Mean orientation factor

matrix[44] αε

Constant

= 2.6 for the FCC structure[44]

Dimensionless

ε

Constrained lattice

= 0.0017 for (FeCoCrNi)94Ti2Al4[15]

Dimensionless

parameter misfit m

Constant

= 0.85[15]

Dimensionless

υ

Poisson ratio

= 0.265 for FeCoCrNiMn[47]

Dimensionless

Table 6 Fraction and size of precipitates in S1 and S2 alloy. Samples ID

Precipitates

Precipitates fraction f (%)

Precipitates radius r (nm)

S1

B2

2.85

241

γ′

6.17

19.5

22

S2

B2

4.56

287

Large γ′

5.72

21.5

Smallest γ′

31.84

9.5

Table 7 Strength contributions to yield strength from different strengthening mechanisms (MPa) Samples ID

∆σCS

∆σMS

∆σCS + ∆σMS

∆σOS

∆σorowan

S1

141.1

17.4

158.5

157.2

39.3

S2

144.2/227.1

17.3/33.0

161.5/260.1

151.4/362.6

48.6

(large γ′/smallest γ′)

4.3 High-temperature tensile properties Fig.9 shows the high-temperature tensile stress-strain curves of the S2 alloy, and the tensile properties are listed in Table 4. It can be found that the dropping trend of both YS and UTS is small, especially for the YS, as increasing temperature. At temperature of 600 and 700℃, the YS, UTS and fracture elongation are 619 MPa, 939 MPa, 35.9 %, and 557 MPa, 714 MPa, 40 %, respectively. A promising performance for this HEA can be maintained up to at least 700 ℃. Fig.10 presents the comparison of the UTS between the S2 alloy with some reported HEAs and superalloys at different temperatures. Comparing the HEAs with different structure, e.g., the single FCC

[49]

(FeCoCrNiMn

),

FCC

+

γ′

(Al10Co25Cr8Fe15Ni36Ti6

[50]

,

Al8Co17Cr17Cu8Fe17Ni33 [51]), FCC + B2 (Al0.3FeCoCrNi [22], AlCoCrFeNi2.1 [18]), BCC + B2 (Al0.7CoCrFe2Ni

[52]

), and FCC + γ′+B2 (Al0.3FeCoCrNi

[19]

), the current S2

HEA exhibits higher tensile strength at all temperatures. In addition, the strength of the S2 HEA is even higher than some commercial superalloys, e.g., 800H alloy

[53]

and Inconel 617 alloy [54]. The resulting outstanding performance of this HEA at high temperatures should be attributed to the combined precipitation of γ′ and B2 phases in the FCC matrix. The S2 alloy has high-density of nanoscale γ′ precipitates in the matrix. The γ′ phases have significant effect on improving both YS and UTS as increasing temperature in many 23

Ni-based alloys, while the strength starts to decrease when beyond a certain temperature

[55]

.

This

phenomenon

was

also

observed

in

HEA,

e.g.,

Al10Co25Cr8Fe15Ni36Ti6[50]. In current HEA, both YS and UTS are slightly decreased as increasing temperature. The remarkable thermal stability of the γ′ precipitates in the HEA exhibits strong resistance against plastic deformation. Moreover, the γ/γ′ lattice misfit parameters in many HEAs is positive and which is decreased with increasing temperature, such characteristic benefit to high temperature properties

[56,57]

.

Furthermore, the γ′ precipitates in HEAs contain more alloying elements than Ni-based superalloys, resulting in the stronger γ′ strengthening and enhance the anti-phase boundary (APB) energy [8, 56]. The higher APB energy means higher energy for dislocation pairs to cut through γ′ precipitates. Generally, it is believed that the grain boundaries are weak at elevated temperatures and thus deformation easily occurs in this area

[58-60]

. Formation of

precipitates in the grain boundaries would suppress the local deformation by restricting dislocation motion around the grain boundary, resulting in strengthening of the alloy.

It is reported that a small amount of δ phases in Inconel 718 alloy is

essential for grain-boundary pinning

[59]

. However, excessive levels of δ phases are

detrimental to properties due to solute depletion and easy crack initiation and growth. Koul et al. [60] reported that the presence of continuous networks of grain-boundary M23C6 carbides dramatically reduce the rupture life of Inconel 738LC alloy, whereas the discontinuous M23C6 carbides at the grain boundaries is considered necessary for obtaining optimum creep properties. Chen et al. [58] suggested that the creep life of the austenitic heat resistant steel (Fe-20Cr-30Ni-2Nb, at %) is extended without ductility loss with increasing grain boundary coverage of Fe2Nb laves phases, owing to the grain boundary covered phases can effectively decrease grain boundary deformation. For the present S2 HEA, the grain-boundary B2 phases mainly display discontinuous distribution, see above microstructure observation in Fig. 4. After high-temperature tensile testing at 600 ℃, it can be found that the grain boundary of the alloy appears not to show some elongated grain structure, as shown in Fig.11a. This indicates that 24

grain-boundary B2 phases suppress the deformation of grain boundary. However, the magnified image (Fig.11b) of the gran boundary structure in Fig.11a shows that the space between the two discontinuous grain-boundary B2 phase is increased comparing to the initial microstructure (Fig.4). The local deformation of the precipitate-free grain boundary, not pinned by particles, will relax the stress concentration originated from the interaction between the near grain-boundary phases and dislocation

[61]

. In addition, the discontinuous grain-boundary precipitates could

hinder the grain boundary sliding, thereby reducing nucleation and growth of cavities [60]

. When the alloy tensile testing at 700 ℃, the elongated grain structure and the

coarsening grain-boundary B2 phases can be observed in Fig.11c and Fig.11d. The high-temperature deformation facilitates the coarsening of B2 phase. The pinning effect of B2 phases on the grain boundaries is reduced, thus the grain boundaries become easily deformable. Fig.11e and Fig.11f show the fractography of the alloys at 600 and 700 ℃, respectively. Evidently, the fracture of both alloys occurs mainly in ductile mode with the presence of many small dimples, while some microvoids can be seen in the alloy tensile tested at 700 ℃. The coarsening precipitates tend to become brittle whereas the surrounding solute depleted matrix behaves as a soft phase [62]. The solute depleted matrix is easily deformable, which contributes in the nucleation of voids at the interphase of precipitates and matrix. On the other hand, serrations behavior can be observed on the stress-strain curves, as shown in Fig.9. At temperature of 600 and 630 ℃, the serrated flow on the stress-strain curves is not significant. With further increasing testing temperatures, the distinct serration feature can be seen in the temperature of 650 and 700 ℃. Generally, serrations on the stress-strain curves for crystalline alloy is attributed to the different flow units, such as dislocations, twins, grain boundaries, precipitates etc., and strongly rely on temperature and strain rate

[63]

. The critical strain to initiate the serration in

present alloy is decreased with increasing the temperature, see Fig.9. The similar change between the critical strain for serration and temperature was observed in Al0.5CoCrCuFeNi HEA, which can be rationalized with an instability theory based on 25

the coupled dynamics of moving dislocation and solutes [64]. It is also reported that the critical strain required for the serration is increased with increasing grain size

[63]

.

Thus, the serration behavior in the present alloy (S2) usually appear at the plastic instability stage of the stress-strain curves with a high strain, see Fig.9, which might be related with the coarser grain size (~35 um) of the alloy. The present alloy possesses γ′ and B2 precipitates in the FCC matrix. Many researches have demonstrated that the serration phenomenon is linked to the interaction between solute atoms and moving dislocations or the precipitate-dislocation interaction

[63-68]

.

Chen et al. [64] reported that the γ′ precipitates in the Al0.5CoCrCuFeNi HEA may provide an obstacle for the moving dislocations, resulting in serrated flow behavior. Precipitation of the B2 phases in Al0.5CoCrCuFeNi HEA can alter the characteristic behavior of the serrated flow [67]. In addition, it is found that the serration behavior of HEAs is usually sensitive to the alloy elements

[63, 68]

. The serrated flow on the

stress-strain curves appear in Al0.3CoCrFeNi single crystals at 600 and 800 ℃, but there is no serration in CoCrFeNi crystals [68]. In the literature [68], it is suggested that Al atoms play a crucial role in the dynamic strain aging (DSA) of the alloy, owing to the solute atmosphere of Al create near a moving dislocation core. The solute atmosphere increases the frictional stress of dislocations. Once the dislocations escape from the atmosphere at a sufficiently high stress, the frictional stress will decrease abruptly, resulting in the serrated flow on the stress-strain curve. There are many factors that affect the serration behaviors, thus more work needs to be performed on the present alloy in the future. In our study, the results show that novel (FeCoNi)81Cr9Al8Ti1Nb1 exhibits excellent high-temperature mechanical properties and even higher than some commercial superalloys. In addition, the present alloy with a high content of Fe and Al which could be a cost-effective HEA. Therefore, the present HEA containing γ-γ′-B2 three phases exhibits a promising potential for high-temperature structural material. It is expected that the properties could be further improved by optimizing the thermomechanical processing and adjusting the morphology, fraction, size and 26

distribution of the precipitates.

Fig.9. Stress-strain curves of S2 alloys at high temperature.

27

Fig.10. Comparison of high-temperature mechanical properties.

Fig.11. Micrographs of the S2 alloys after high-temperature tension testing. (a) and (b) tension testing at 600℃;(c) and (d) tension testing at 700 ℃; (e) and (f) fracture surface at 600 and 700 ℃, respectively.

5 Conclusions A new (FeCoNi)81Cr9Al8Ti1Nb1 HEA strengthened by combining γ′ with B2 phases was designed and fabricated by melting and aging treatment. Based on the results and observations presented in this study, the following conclusions can be 28

reported: 1. Co-addition of Al, Ti and Nb to the FeCoCrNi-based HEA can promote the stable γ′ and B2 phases precipitation in the FCC matrix. The coherent γ′ and incoherent B2 phases can be formed simultaneously at one-step aging stage (845℃/4 h). High density of γ′ phases are precipitated at two-step aging stage (720℃/10 h). The γ′ phases with two grain sizes (~19 or ~ 43 nm) are distributed in the FCC matrix, while the B2 phases are mainly located along the grain boundary. 2. The results of the γ-γ′-B2 three phases in the two-step aged alloy show an excellent balanced tensile property at room temperature. The YS, UTS and fracture elongation are 863 MPa, 1285 MPa, and 28 % respectively. 3. Quantitative calculations of the individual strengthening effects demonstrate that particle (γ′) shearing mechanism is the predominant strengthening mechanism. The strengthening contributions from γ′ and B2 phases to the yield strength increment are 552.1 and 48.6 MPa, respectively. The high work-hardening capability of the γ matrix could greatly suppress the propagation of microcracks originated at these brittle B2 phases and promote a retention of ductility. 4. The two-step aged HEA also exhibits an excellent high-temperature tensile property up to 700 ℃. At temperature of 600 and 700 ℃, the YS, UTS, and fracture elongation are 619 MPa, 939 MPa, 36%, and 557 MPa,714 MPa, 40%, respectively. 5. The excellent high-temperature strength of this alloy is attributed to the high thermal stability of the γ′ precipitates as well as the pinning effect of the grain-boundary B2 phases on the grain boundary.

Acknowledgments This research was financially supported by National Major Science and Technology Project of China (No. 2014ZX07214-002). The author (L. Fu) is grateful to the financial support from Young Researchers Startup Fund for Youngman Research 29

at SJTU (SFYR at SJTU: No. 18X100040023).

Data availability The raw/processed data required to reproduce these findings cannot be shared at this time as the data also forms part of an ongoing study.

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34

Highlights

A new (FeCoNi)81Cr9Al8Ti1Nb1 HEA with γ′ and B2 phases was fabricated, which exhibits an outstanding balanced strength and ductility at room and elevated temperatures. Combined precipitation of γ′ and B2 phases in γ matrix that effectively improve the strength, while the ductility γ matrix suppress the brittleness of B2 and promote a retention of ductility. High thermal stability of the γ′ precipitates and the pinning effect of the grain-boundary B2 precipitates on grain boundary result in excellent high-temperature strength.

Conflict of interest statement: We declare that we have no financial and personal relationships with other people or organizations that can inappropriately influence our work, there is no professional or other personal interest of any nature or kind in any product, service and/or company that could be construed as influencing the position presented in, or the review of, the manuscript entitled.