Microstructure and mechanical properties of oxide dispersion strengthened FeCoNi concentrated solid solution alloys

Journal of Alloys and Compounds xxx (xxxx) xxx

Contents lists available at ScienceDirect

Journal of Alloys and Compounds journal homepage: http://www.elsevier.com/locate/jalcom

Microstructure and mechanical properties of oxide dispersion strengthened FeCoNi concentrated solid solution alloys Yuanhang Guo a, Mingyang Li a, Pei Li b, Cunguang Chen b, Qian Zhan a, Yongqin Chang a, *, Yanwen Zhang c, d a

School of Materials Science and Engineering, University of Science and Technology Beijing, Beijing, 100083, China Institute for Advanced Materials and Technology, University of Science and Technology Beijing, Beijing, 100083, China Materials Science and Technology Division, Oak Ridge National Laboratory, Oak Ridge, TN, 37831, USA d Department of Materials Science and Engineering, University of Tennessee, Knoxville, 37996, USA b c

a r t i c l e i n f o

a b s t r a c t

Article history: Received 22 July 2019 Received in revised form 16 November 2019 Accepted 18 November 2019 Available online xxx

Microstructure and mechanical properties of two kinds of oxide dispersion strengthened concentrated solid solution alloys (CSAs), FeCoNi-1.5Y2O3 (FCNY) and FeCoNi-1.2Ti-1.5Y2O3 (FCNTY), are studied through a comparing investigation with FeCoNi (FCN) CSAs. All these alloys are fabricated by mechanical alloying, spark plasma sintering, hot rolling and annealing treatment. For three kinds of alloys, both the as-milled powders and bulk materials are of single face-centered cubic structure. Electron backscattered diffraction results reveal that the texture transformation is suppressed during the hot rolling process because the movement of grain boundaries is hindered by the oxide particles. Compared with FCN CSAs, grains are refined by 43% and 47% for FCNY and FCNTY CSAs, respectively. Nano-sized Y2O3 (monoclinic structure) and Y2Ti2O7 (pyrochlore structure) particles are uniformly distributed in FCNY and FCNTY CSAs, respectively. Both Y2O3 and Y2Ti2O7 particles show a semi-coherent relationship with the matrix. Yield strength of FCN, FCNY and FCNTY CSAs is 559, 981 and 1050 MPa, respectively. Theoretical calculations illustrate that high strength of FCNY and FCNTY CSAs comes from refined grains and highdensity nano-sized oxide particles. © 2019 Elsevier B.V. All rights reserved.

Keywords: Metal matrix composites Mechanical alloying Microstructure Phase transitions Mechanical properties

1. Introduction Nowadays, the newly developed CSAs, including high entropy alloys (HEAs), have attracted increasing attention due to their unique structures and remarkable properties, such as high strength, high ductility, good oxidation and corrosion resistance, and outstanding irradiation-resistance [1e6]. Normally, CSAs tend to form single-phase structures, such as face-centered cubic (FCC), body-centered cubic (BBC), hexagonal close-packed (HCP) or a mixture of them [4,7,8]. Contrary to the traditional situation, FCC CSAs have shown better swelling resistance than that of BCC CSAs [9]. However, poor mechanical property is one of the prominent shortages for FCC structure alloys and restricts their applications. Generally, metal and alloys can be enhanced by refined grains, precipitate hardening, solid solution strengthening and highdensity dislocations [10e12]. In previous studies, oxide dispersion

* Corresponding author. E-mail address: [email protected] (Y. Chang).

strengthening (ODS) has been proven an effective way to improve the mechanical properties by refining grains and hindering dislocation movement [13]. In addition, large numbers of interfaces are introduced into the matrix by these high-density nano-sized particles (NPs). These interfaces can act as sinks for irradiation defects, such as dislocation loops, voids and helium bubbles, which is beneficial for the high resistance to the irradiation induced hardening and embrittlement [14,15]. It is, therefore, predicted that these ODS-CSAs may possess substantial performance improvement because it combines the advantages of both CSAs and ODS alloys. For ODS-CSAs, two kinds of NPs, Y2O3 and Al2O3, were mainly used in previous research and these NPs were introduced to CSAs with the mechanical alloying (MA) method. Hadraba et al. [16] reported that grain size of CoCrFeNiMn was refined 50% by adding 0.3 wt% oxide particles. The ultimate tensile strength and the yield strength increased by 30% and 70% at room temperature and 800
C, respectively. Praveen et al. [17] reported that the hardness of AlCoCrFe HEAs increased from 1050 ± 20 HV1 to 1070 ± 20 HV1

https://doi.org/10.1016/j.jallcom.2019.153104 0925-8388/© 2019 Elsevier B.V. All rights reserved.

Please cite this article as: Y. Guo et al., Microstructure and mechanical properties of oxide dispersion strengthened FeCoNi concentrated solid solution alloys, Journal of Alloys and Compounds, https://doi.org/10.1016/j.jallcom.2019.153104

2

Y. Guo et al. / Journal of Alloys and Compounds xxx (xxxx) xxx

because of the addition of oxide particles. Gwalani et al. [18] discovered that compressive yield strength of Al0.3CoCrFeMnNi HEAs increased from 979 MPa to 1759 MPa due to the addition of 3 vol% Y2O3. Liu et al. [19] found that the introduction of 0.25 wt% Y2O3 improved the tensile strength of CrMnFeCoNi HEAs from 868 MPa to 1001 MPa and ODS-CrMnFeCoNi HEAs also showed excellent wear-resistance. Jia et al. [20] reported that the yield strength increased from 654 MPa to 1754 MPa for FeCrCoNi HEAs strengthened by nano-Y2O3 particles. Rogal et al. [21] found that the compressive yield strength of ODS-CoCrFeMnNi increased from 1180 MPa to 1600 MPa with 5 wt% additions of a-Al2O3 NPs. Yang et al. [22] reported the yield strength, fracture strength and compressive ratio of Al0.4FeCrCo1.5NiTi0.3 HEAs reinforced with 8 wt % nano-Al2O3 particles reached 2.05 ± 0.01 GPa, 2.14 ± 0.01 GPa and 13.98 ± 0.25%, respectively. Prasad et al. [23] found that FeNiCoCrAlMn HEAs dispersed with 2 wt% Al2O3 possessed less thermal expansion coefficient and better oxidation resistance. These recent results demonstrate that the addition of nano-sized oxide particles is an effective and suitable way to design CSAs with superior mechanical properties. The previous studies focused on the quaternary or quinary CSAs and few work concentrated on ODS ternary alloys. Hightemperature and high-fluence irradiation experiments suggested that the improved radiation resistance depended critically on the type alloying elements and their concentration. Some medium entropy alloys (MEAs) with the unique electronic structure and atomic-level inhomogeneity [5,24] show better irradiation resistance than that of HEAs containing 4 or more elements [25e28]. Microstructural evolution of CSAs with and without nano-sized particles has not been fully analyzed, and in-depth understanding is lacking. In this work, FeCoNi ternary CSAs were chosen as the base metal because of its excellent swelling resistance and corrosion resistance [25,29]. Y2O3 and Ti powders were introduced into the matrix to form nano-sized particles. Microstructural evolution at different stages was investigated. Size distribution, crystal structure, chemical composition and orientation relationship of the NPs were systematically analyzed. The strengthening efficacies from different strengthening mechanisms were quantitatively calculated based on the microstructure observation. Furthermore, deformation mechanism during the hot rolling (HR) process and precipitation mechanism of the NPs were also discussed in detail. 2. Materials and methods High purity Fe (>99.9 wt%, <48 mm), Co (>99.9 wt%, 3e5 mm), Ni (>99.7 wt%,<48 mm), Ti (>99.5 wt%,<45 mm) and Y2O3 (>99.7 wt%, 30 nm) powders were blended with the nominal chemical composition FeCoNi (FCN), FeCoNi-1.5Y2O3 (FCNY) and FeCoNi1.2Ti-1.5Y2O3 (FCNTY) in a planetary ball-miller (TENCAN POWDER-XQM-4) for 60 h at 440 rpm in argon atmosphere. Stainless steel vials and balls were used as milling media, and the ball-to-powder mass ratio was 15:1. Ethyl alcohol was introduced as the processing controlling agent (PCA) to prevent metal oxidation and cold welding. As-milled powders were consolidated at 1050
C for 5 min with pressure of 45 MPa in vacuum with the SPS (SPS-201-10 IV) method to produce the cylindrical type specimens with the dimension of 30 mm  10 mm (d  h). Subsequently, cylindrical samples were rolled at 900
C from 10 mm to 2 mm in steps of 1 mm reduction per pass. Annealing treatment was conducted at 950
C for 5 h in argon atmosphere and cooled in the furnace. Average particle size of the as-milled powders was determined by laser particle size distribution analyzer (LPSDA, LMS-30). The metallographic analysis of the as-milled powders and annealed samples was carried out using scanning electron microscopy (SEM,

ZEISS-EVO 18) on the secondary electron (SE) mode. Density of the bulk materials was tested using Archimedes’ principle. Crystalline structure of the powders and bulk materials at different stages was detected using X-ray diffraction equipment (XRD, Rigaku D/maxRb) with Cu Ka radiation (l ¼ 0.15406 nm). Standard Silicon powders (seen in Supplementary Fig. 1) were used to measure the full width at half maximum (FWHM) curve and subtract the instrumental error. Grain size and texture observation were carried out by auger electron spectrometer (AES, ULVAC-PHI 710) operated at a voltage of 15 kV, a tilt angle of 70
and a step size of 0.10 mm. Specimens for EBSD characterization were mechanically ground, polished with standard metallurgical methods, and then electropolished with a solution of 10 ml perchloric acid and 90 ml ethanol at 30 V for 2 min. To insure the accuracy of electron backscattered diffraction (EBSD) data, collecting points with the absolute confidence index (CI) value less than 0.1 were filtered before clean-up and subsequent calculation. Tolerance specifying the maximum misorientation angle was set to 15
when calculating the volume fraction of fibers. Microstructures and distribution of the NPs were analyzed by using transmission electron microscopy (TEM, FEI Tecnai F20 and JEM 2010F) with energy dispersive spectrometer (EDS) system. EDS test for composition analysis was carried out on the scanning transmission electron microscopy (STEM) mode. Orientation relationship and coherency between the NPs and the matrix were measured by lattice images and their corresponding Fast-Fourier Transformation (FFT) images obtained from high resolution TEM (HRTEM). TEM samples were prepared by ion-beam thinning. Total content of yttrium and titanium elements in FCNY and FCNTY CSAs was tested using inductance coupling plasma (ICP) optical emission spectrometer (Varian 715-ES). Samples used for SEM, EBSD and TEM tests were in annealed state and sampling location was on the rolling direction (RD)-transverse direction (TD) section in the middle of thickness. Tensile test was performed in an electromechanical universal testing machine CMT4105 (100 kN load capacity) at a strain rate of 1  103 s1. The specimens (gauge length Lo, 10 mm and width, 3 mm) were taken along the rolling direction. Three samples for each kind of materials were tested to obtain the average values. Fracture morphology after tensile tests was observed by field emission scanning electron microscopy (FESEM, Zeiss GeminiSEM 500) on the InLens mode. Vickers hardness was measured using 401 MVD Wolpert Wilson Instruments™, and 15 points were tested. 3. Results 3.1. Microstructure of FCN, FCNY and FCNTY CSAs 3.1.1. SEM observation Fig. 1 shows the SEM images (SE mode) of the as-milled powders and the corresponding annealed samples of FCN, FCNY and FCNTY CSAs. The as-milled powders are in spherical or ellipsoidal shape with an average diameter of 13.8, 14.7 and 15.1 mm, respectively (seen in Supplementary Fig. 2). After SPS and hot rolling, the asmilled powders are solidified and deformed along the rolling direction, as shown in Figs. 1(d), (e) and (f). The bulk relative density is 96.5%, 94.4% and 93.9% of the theoretical value for FCN, FCNY and FCNTY CSAs, respectively. For three kinds of materials, only a few pores are observed, which indicates that SPS þ HR is an effective way to obtain high-density bulk materials. 3.1.2. XRD Fig. 2 shows the XRD spectra of FCN, FCNY and FCNTY CSAs at different processing stages. For three kinds of samples, diffraction peaks of the raw metals disappear after milling for 60 h and new

Please cite this article as: Y. Guo et al., Microstructure and mechanical properties of oxide dispersion strengthened FeCoNi concentrated solid solution alloys, Journal of Alloys and Compounds, https://doi.org/10.1016/j.jallcom.2019.153104

Y. Guo et al. / Journal of Alloys and Compounds xxx (xxxx) xxx

3

Fig. 1. SEM images (SE mode) of the as-milled powders (top images a-c) and annealed samples (bottom images d-f) of FCN, FCNY and FCNTY CSAs.

Fig. 2. XRD spectra of (a) FCN, (b) FCNY, and (c) FCNTY CSAs at different processing stages. (Raw powders, MA powders: mechanical alloying powders, SPS: spark plasma sintering, Hot rolling, and Heat treatment.)

Please cite this article as: Y. Guo et al., Microstructure and mechanical properties of oxide dispersion strengthened FeCoNi concentrated solid solution alloys, Journal of Alloys and Compounds, https://doi.org/10.1016/j.jallcom.2019.153104

4

Y. Guo et al. / Journal of Alloys and Compounds xxx (xxxx) xxx

peaks with weak intensity appear, indicating that the raw powders have been fully alloyed. After sintering, the diffraction intensity of all peaks increases rapidly and all the diffraction peaks belong to a single FCC structure. In addition, all peaks shift slightly to a higher Bragg angle (2q). During the MA process, a large number of defects and high quantity of stress are introduced into the powders, which causes the broad diffraction peak and poor crystallinity. After the SPS process, the internal stress releases, the lattice distortion recovers, the FWHM of all peaks also becomes narrow and the intensity increases, implying improved crystallinity and grain growth. After HR and heat treatment, diffraction peaks change differently for three kinds of materials. For FCN CSAs, the strongest diffraction peak transforms from (111) plane to (220) plane after HR. Peak intensity difference between two planes increases after annealing. For FCNY CSAs, (111) diffraction peak remains to be the strongest but the intensity of (220) plane also increases after HR. The peak intensity difference between (111) plane and (220) plane

shrinks after annealing. The above results illustrate that the orientation of partial grains changes in FCNY CSAs, but this tendency is impeded partly. For FCNTY CSAs, (111) plane remains the strongest peak and higher than other peaks both at HR and annealing stages. For three kinds of samples, change in peak positions is negligible during SPS, HR and heat treatment processes, and no new phases are detected, which means that the FCC phase is stable. The relative intensity change of the diffraction peaks corresponds to different preferred orientation. For three samples, experimental parameters at different stages are identical except for the chemical composition. Thus, it can be inferred that grain orientation change is suppressed by the addition of Y2O3 and further prohibited as titanium is introduced into the matrix. 3.1.3. EBSD analysis Fig. 3 shows the microstructure and grain size distribution of FCN, FCNY and FCNTY CSAs. Ultrafine grains with the average size of

Fig. 3. Inverse pole figure maps of (a) FCN, (b) FCNY, and (c) FCNTY CSAs and grain size distribution of three kinds of alloys (involved twin structures). (Note: in IPF maps, “ACI” is short for average confidence index.)

Please cite this article as: Y. Guo et al., Microstructure and mechanical properties of oxide dispersion strengthened FeCoNi concentrated solid solution alloys, Journal of Alloys and Compounds, https://doi.org/10.1016/j.jallcom.2019.153104

Y. Guo et al. / Journal of Alloys and Compounds xxx (xxxx) xxx

5

Fig. 4. (a) Schematic representation of deformation and annealing texture components and fibers in FCC alloys and the corresponding ODF figures of FCN, FCNY and FCNTY CSAs on sections of f2 ¼ 0
and 45
. (b) Volume fraction of fibers of FCN, FCNY and FCNTY CSAs.

854 ± 480 nm (including twin boundaries (TBs)) are obtained in FCN CSAs. The grain size of FCN fabricated by vacuum arc melting (VAM) method is generally several tens of micrometers [30], which means that MA is an effective way to refine grain for FeCoNi CSAs. The grain size is further refined by introducing Y2O3 and Ti into the matrix, as shown in Fig. 3(b) and (c). The mean grain sizes of FCNY and FCNTY CSAs are 483 ± 172 and 456 ± 187 nm, respectively. In Fig. 3, points with the absolute CI value less than 0.1 have been removed, which correspond to the black area in all the inverse pole figure (IPF) maps. Fig. 4 shows the typical deformation and annealing texture components in FCC structure alloys on the selected orientation distribution function (ODF) sections (f2 ¼ 0
and 45
) [31e33], and the corresponding ODF sections of FCN, FCNY and FCNTY CSAs. The corresponding Euler angles and Miller indices are listed in Table 1. In Fig. 4(a), strong a-fiber texture components, such as Brass-type, A-type, P-type and Rotated Goss-type texture, are identified. All of them belong to the {110} type texture except for different crystallographic directions. Another strong texture component is

Table 1 Deformation and annealing texture components and fibers in FCC alloys. Texture component

Cube (C) Brass (Bs) A P Rotated Goss (RtG) E D

Symbol

⋄ ▽ ◎ * Q ※

Miller indices

{001}<100> {110}<112> {110}<111> {011}<211> {110}<110> {111}<110> {112}<110>

Euler angles (
)

Fiber

f1

F

f2

0/90 35 55 60 90 0/60 0

0/90 45 45 45 45 55 35

0 0 0 0 0 45 45

e a, b

a a a g e

Cube-type texture, which corresponds to {001}<110>. In FCN CSAs, the volume fraction of a-fiber component is 39.1%, which is much higher than other fibers. According to the XRD results (Fig. 2(a)), the strongest diffraction peak is detected on (220) plane at the annealing state, which agrees well with the ODF results. For FCNY and FCNTY CSAs, strong g-fiber texture corresponding to the {111}

Please cite this article as: Y. Guo et al., Microstructure and mechanical properties of oxide dispersion strengthened FeCoNi concentrated solid solution alloys, Journal of Alloys and Compounds, https://doi.org/10.1016/j.jallcom.2019.153104

6

Y. Guo et al. / Journal of Alloys and Compounds xxx (xxxx) xxx

Fig. 5. TEM bright field images of (a) FCN, (c) FCNY, and (e) FCNTY CSAs and their corresponding SAED patterns: (b), (d) and (f). SAED patterns in (d) and (f) is along [011]M zone axis and the twin structure along [01 1]T zone axis.

<110> component is observed on the f2 ¼ 45
section, which corresponds to the highest peak on the (111) plane in Figs. 2(b) and (c). Weak Brass-type and D-type textures are also observed in FCNY and FCNTY CSAs. The differences of relative volume fraction between g-fiber and a-fiber components are -7.5% and 8.6% for FCNY and FCNTY CSAs, respectively, which means stronger g-fiber texture exists in FCNTY CSAs. According to the XRD results, the difference of the peak intensity between (111) plane and (220) plane in FCNY CSAs (Fig. 2(b)) is smaller than that in FCNTY CSAs (Fig. 2(c)), suggesting that the EBSD and XRD results agree with each other. The volume fraction of deformation texture components measured by EBSD is micro texture but the XRD peak intensity is proportional to the bulk texture, which results in the numerical deviation between two results. In this work, ODF results are supported by the XRD results, semi-quantitatively confirming that texture component (g-fiber / a-fiber texture) changes during the HR process in FCN CSAs. But this change could be well suppressed by the addition of Y2O3 and further prohibited with the introduction of titanium.

3.1.4. TEM observation Fig. 5 presents typical TEM images of FCN, FCNY and FCNTY CSAs and their corresponding selected area electron diffraction (SAED) patterns. Elongated grains are observed in Fig. 5(a). Diffraction rings in Fig. 5(b) reveal that only the FCC structure exists in the matrix. Figs. 5(c) and (e) indicate that grains of FCN CSAs are highly refined with the addition of Y2O3 and Ti. Nano-sized particles marked by red and blue arrows are uniformly distributed in the grains and at the grain boundaries. Twin structures marked by yellow solid lines are also observed in both FCNY and FCNTY CSAs. In Figs. 5(d) and (f), the diffraction patterns of the matrix are marked by white solid lines and the selected twin structures are marked by yellow dashed lines. Plane (111) is identified as the twinning plane in FeCoNi CSAs. 3.1.5. STEM observation Figs. 6(a) and (b) show the STEM images of FCNY and FCNTY CSAs. The changes in contrast feature come from differences in chemical composition. The nearly spherical NPs are uniformly

Please cite this article as: Y. Guo et al., Microstructure and mechanical properties of oxide dispersion strengthened FeCoNi concentrated solid solution alloys, Journal of Alloys and Compounds, https://doi.org/10.1016/j.jallcom.2019.153104

Y. Guo et al. / Journal of Alloys and Compounds xxx (xxxx) xxx

7

Fig. 6. STEM images of (a) FCNY and (b) FCNTY CSAs and (c) size distribution of nano-sized particles in two kinds of alloys.

Table 2 Chemical composition (at.%) at each point in FCNY and FCNTY CSAs. at.%

O

Ti

Cr

Mn

Fe

Co

Ni

Y

FCNY CSAs (Fig. 6(a))

Point Point Point Point Point

1 2 3 4 5

36.0 29.6 63.9 24.3 e

e e e e e

0.3 0.4 0.2 0.4 e

0.4 0.4 0.3 0.6 e

21.0 23.6 8.3 25.4 36.0

18.8 21.0 7.2 22.5 31.4

16.6 19.4 5.3 20.2 32.7

6.9 5.6 14.8 6.6 e

FCNTY CSAs (Fig. 6(b))

Point Point Point Point Point

6 7 8 9 10

45.1 35.9 48.4 68.8 e

5.1 10.0 11.0 20.2 e

e e e e e

e e e e e

13.7 14.7 9.7 1.0 34.6

16.2 16.3 11.4 1.0 32.2

14.6 14.5 9.3 0.7 33.2

5.3 8.7 10.2 11.3 e

distributed in these two kinds of alloys, which correspond to the particles marked in Figs. 5(c) and (e). Several representative particles are chosen to quantitatively analyze the chemical composition, and the values are summarized in Table 2. The results demonstrate that the NPs in FCNY CSAs mainly composed of yttrium and oxygen. Trace elements, such as Cr and Mn, are also detected, which may come from stainless steel vials and balls. During the long-time high energy ball milling process, it is very hard to avoid impurity element pollution [34]. Finer particles enriched with titanium, yttrium and oxygen elements are detected in FCNTY CSAs (Table 2). More than two hundred particles are

counted to calculate the average diameter and size distribution in these two kinds of samples, and the results are given in Fig. 6(c). For FCNY CSAs, the diameters of the NPs range from 2.5 to 60 nm with the average size of 13.1 nm. For FCNTY CSAs, diameter distribution is narrower and the average diameter reduces to 10.6 nm with the addition of Ti. The NPs with the diameter less than 10 nm in FCNTY CSAs account for around 53%, while the ratio of FCNY CSAs is only about 40%. It can be concluded that the addition of Ti is also beneficial to refine the NPs by the formation of Y-Ti-O particles in ODS FeCoNi CSAs, just like in other ODS materials [35e37]. In FCNY CSAs, high-density NPs prevent the movement of grain boundaries,

Please cite this article as: Y. Guo et al., Microstructure and mechanical properties of oxide dispersion strengthened FeCoNi concentrated solid solution alloys, Journal of Alloys and Compounds, https://doi.org/10.1016/j.jallcom.2019.153104

8

Y. Guo et al. / Journal of Alloys and Compounds xxx (xxxx) xxx

Fig. 7. HRTEM micrograph of Y2O3 particles in FCNY CSAs with different diameter: (a) ~5.0 nm, 7.1 nm, (b) ~15.4 nm and their corresponding FFT images, (c) ~56.7 nm and its SAED patterns (Note: in FFT image and SAED patterns, “P” is short for particle and “M” is short for matrix.).

Fig. 8. HRTEM micrograph of Y2Ti2O7 particles with diameters of (a) ~9.3, (b) ~21.9, (c) ~28.9 nm and its corresponding FFT image in FCNTY CSAs (Note: In FFT images, “P” is short for particle and “M” is short for matrix.).

which retard grain rotation during HR process and refines grains during annealing treatment. In FCNTY CSAs, the impediment effect of finer NPs is stronger, which result in more stable microstructures during the HR and annealing processes.

3.1.6. HRTEM Fig. 7 shows the HRTEM images of several representative NPs with different diameters in FCNY CSAs. Their corresponding FFT images or SAED patterns are given in the corner. The red dashed lines on the FFT or diffraction images represent the NPs and white arrows represent the matrix. Spherical or elliptical particles are

Please cite this article as: Y. Guo et al., Microstructure and mechanical properties of oxide dispersion strengthened FeCoNi concentrated solid solution alloys, Journal of Alloys and Compounds, https://doi.org/10.1016/j.jallcom.2019.153104

Y. Guo et al. / Journal of Alloys and Compounds xxx (xxxx) xxx

9

red dashed lines. Interplanar distances d(111)M of the matrix and d(601)P of Y2O3 are 2.060 Å and 2.303 Å. The misfit parameter d is 0.12, indicating that the semi-coherent interface remains even as the particle coarsens to 56.7 nm (Fig. 7(c)). Fig. 8 shows the HRTEM images of representative NPs with different diameters in FCNTY CSAs and their corresponding FFT images. The blue dashed lines on the FFT images represent the NPs and white solid lines represent the matrix. Hexagonal or elliptical particles are identified as Y2Ti2O7 with pyrochlore structure [40]. In Figs. 8(a) and (b), parallel relationship exists between the Y2Ti2O7 particles and the matrix: (040)P//(200)M, (222)P//(111)M and zone axis [101]P//[011]M. Interplanar distances d(111)M of the matrix and d(333)P of Y2Ti2O7 are 2.063 Å and 1.943 Å, respectively. The misfit parameter d is 0.06, indicating a semi-coherent interface between Y2Ti2O7 and the matrix. In Fig. 8(c), the diffraction patterns of (040)P, (200)M and (000) are not parallel, but the (22 2)P plane is still parallel to (111)M plane, which means the semi-coherent interface still exists. Fig. 9. Engineering stress-strain curves of FCN, FCNY and FCNTY CSAs.

3.2. Mechanical properties identified as Y2O3 with the monoclinic (M) structure. The misfit parameter (d) between the interface of particle and matrix is commonly calculated using the equation [35]:

d¼

dP  dM dM

(1)

where dP and dM are interplanar spacing of oxide particles and the matrix, respectively. Generally, d is calculated based on the lattice spacing in databases, and thus crystallographic data of Y2O3 (monoclinic [38]) is used as the theoretical values. The crystalline interplanar spacing of the matrix is determined from the XRD results in Fig. 2(b). Interplanar distances d(200)M of the matrix and d(204)P of Y2O3 are 1.784 Å and 1.934 Å, respectively. The misfit parameter d is 0.08, which is between 0.05 (coherent) and 0.25 (incoherent) [39] and indicates a semi-coherent interface between Y2O3 and the matrix. In Fig. 7(c), two independent sets of SAED patterns are obtained. One set is confirmed as the matrix, marked by white solid lines. The other set is identified as Y2O3, marked by

Engineering stress-strain curves of FCN, FCNY and FCNTY CSAs are shown in Fig. 9. Mechanical properties of three samples are given in Table 3 as well as other typical ODS alloys [41e50]. FCN CSAs show the high yield strength and ultimate tensile strength along with good ductility, and the corresponding values are 559 ± 18 MPa, 709 ± 6 MPa and 12.0%, respectively. For common FCN CSAs prepared by VAM, the corresponding values are 199 MPa, 495 MPa and 55%, respectively [30]. The enhanced tensile properties are mainly ascribed to the highly refined grains. Yield strength is further improved by introducing NPs into the matrix, which reduces the grain size (Figs. 3(b) and (c)) and interacts with dislocations during the deformation process. In Table 3, FCNY and FCNTY CSAs possess higher yield strength than ODS austenitic steels and partial ODS ferritic/martensitic (F/M) steels. However, ODS F/M steels, such as 14 YWT and MA957, exhibit higher tensile strength because of the BCC/BCT matrix structure and finer NPs. Moreover, the microhardness results also show that the hardness of FCN CSAs prepared by MA is higher than that prepared by VAM. The hardness

Table 3 Mechanical properties of FCN, FCNY and FCNTY CSAs and other typical ODS alloys. Chemical composition

Microstructurea Processa

Ultimate tensile strength (MPa)

Yield strength Elongation Hardness Reference (MPa) (%) (HV)

ODS-Co 14 YWT MA957 FCNTY

Coe20Cre5Al-2.4Hf-1.5Y2O3 Fee14Cre3W-0.4Ti-0.3Y2O3 Fe-13.87Cr-1.05Ti-0.3Mo-0.22Y FeCoNi-1.2Ti-1.5Y2O3

FCC þ HCP F F FCC

SPS þ HR WE HE SPS þ HR

2850 1564 1433 1140

e 1435 1158 1050

1.2 12.0 11.4 1.2 ± 0.8

e 499 e 367 ± 9

9Cr-ODS FCNY

Fe-9.08Cr-0.14C-1.97W-0.23Ti-0.29Y-0.16O FeCoNi-1.5Y2O3

F/M FCC

HE þ Forge SPS þ HR

1262 1117

1033 981

7.8 2.9 ± 0.8

405 349 ± 6

F/M

HIP þ HR

1085

966

11.7

e

[42] [43,44] [45] Present work [46] Present work [43]

A

HIP þ HR

981

e

15.3

397.5

[47]

F A A FCC

HR 925 Forge þ HR 852 HIP þ Forge þ HR 920 SPS þ HR 709

790 601 575 559

16 30 48.5 12.0 ± 2.0

345 e e 232 ± 5

A

HIP þ Forge þ HR 723

458

46

e

[48] [49] [50] Present work [51]

FCC

VAM

199

55

141

[30]

Fe-0.1Ce9Cre1W-0.2V-0.1Ta-0.3Y2O3 ODSEurofer 97 ODS steel Fe-16.7Cr-13.7Ni-0.2Ti-0.19Y-0.9W-0.45Zr0.38O PM2000 Fee20Cr-5.59Al-0.50 Y2O3-0.51Ti ODS-310 Fe-23.97Cr-18.33Ni-0.32Ti-1.93Mo-0.26Y-0.16O ODS-304 Fe-16.7Cr-7.62Ni-1.03Mo-0.49Ti-0.23Y-0.5 N FCN FeCoNi ODS-316 L Fe-16.82Cr-13.23Ni-0.72Si-0.28Ti-0.40Mn0.3Al-2.48Mo-0.2N-0.35Y2O3 FCN FeCoNi

495

a

The following acronyms are used: HCP (hexagonal closed-packed); F (ferritic); WE (warm extrusion); M (martensitic); HE (hot extrusion); HIP (hot isostatic pressing); A (austensitic); VAM (vacuum arc melting).

Please cite this article as: Y. Guo et al., Microstructure and mechanical properties of oxide dispersion strengthened FeCoNi concentrated solid solution alloys, Journal of Alloys and Compounds, https://doi.org/10.1016/j.jallcom.2019.153104

10

Y. Guo et al. / Journal of Alloys and Compounds xxx (xxxx) xxx

Fig. 10. SEM images (InLens mode) of the fracture morphology after tensile tests of (a) FCN, (c) FCNY, (e) FCNTY CSAs and their corresponding high magnification images (b), (d) and (f). (Note: Images (b), (d) and (f) correspond to the magnified white dashed rectangle areas in images (a), (c) and (e)).

is further improved with the addition of NPs, which agrees with the tensile strength. In order to further investigate the deformation mechanism of the ductility variation in three kinds of samples, fracture surfaces of the tensile tested samples are examined by SEM, and the results are shown in Fig. 10. Dimples with typical plastic fracture features are observed for both FCN and FCNY CSAs. Huge and deep dimples in FCN CSAs illustrate good plasticity compared with those in FCNY CSAs. At the bottom of dimples, nano-sized particles could be found which may act as the source of cracking during tensile test. Brittle surface with intergranular fracture is observed in FCNTY CSAs, which means that the addition of Ti reduces the ductility of FeCoNi CSAs. In all three kinds of samples, a small number of deformed pores exist, which may originate from unclosed pores at the annealed state (Figs. 1(d)e(f)) and are believed to be the main crack source during tensile test. In this paper, the lowest bulk density is measured in FCNTY CSAs, which means more unclosed pores at the powder boundary and poor bonding among the powders. This weakening effect coming from pores causes the failure of materials before grains undergo the plastic deformation, which results in the

brittle fracture of FCNTY CSAs. 4. Discussion 4.1. Formation of the monoclinic Y2O3 particles in FCNY CSAs In this work, M-Y2O3 particles with different diameters are observed in FCNY CSAs (Fig. 7), but Y2O3 particles with cubic (C) structure are blended as raw powders (Fig. 2(b)), which means that phase transition occurs during the manufacturing process. It has been reported that C-Y2O3 was a stable form of yttrium oxide at room temperature and atmospheric pressure [51]. It transformed to a more closely packed hexagonal (H) phase at 2326 K and atmospheric pressure [52]. Only when powders were pressed at and above 10 GPa, cubic structure was no longer stable and started to transform to the monoclinic structure at room temperature and this transition was reconstructive and irreversible. M-Y2O3 remained metastable when the pressure was lowered to atmospheric pressure [53]. The above results reveal that pressure is the key factor for the “C / M” phase transition. Furthermore, Zhang

Please cite this article as: Y. Guo et al., Microstructure and mechanical properties of oxide dispersion strengthened FeCoNi concentrated solid solution alloys, Journal of Alloys and Compounds, https://doi.org/10.1016/j.jallcom.2019.153104

Y. Guo et al. / Journal of Alloys and Compounds xxx (xxxx) xxx

11

Fig. 11. Schematic route maps of monoclinic Y2O3 formation.

et al. [37] reported that C-Y2O3 was observed in the Co-based ODS alloys with addition of 1.5 wt% Y2O3 prepared by MA and SPS at 1145
C. Oono et al. [35] reported that C-Y2O3 and M-Y2O3 coexisted in the Ni-based ODS alloys prepared by MA, with the following SPS and hot rolling at 1150
C. In this work, cylindrical samples are sintered at 1050
C and hot rolled at 900
C for 2 h with 80% thickness reduction, and only M-Y2O3 is observed in FCNY CSAs. All these studies indicate that HR may provide sufficient pressure and promote the “C / M” phase transition in ODS alloys strengthened by Y2O3 particles. Moreover, it has been widely reported that the crystalline structure of initial Y2O3 particles was destroyed and decomposed into yttrium and oxygen atoms in ODS steels during MA process and those atoms precipitated at the following heat working stages [54,55]. The precipitation temperature was 800
C and 690
C for Fe-based [56] and Co-based alloys [37], respectively. In our work, heat working temperatures during SPS and HR are higher than both temperatures, and thus yttrium and oxygen atoms are considered to precipitate continually. In general, different states of yttrium oxide and formation of M-Y2O3 are conjectured and drawn in the schematic route map shown in Fig. 11. The C-Y2O3 particles with diameter around 30 nm are decomposed and dissolved into the matrix in the form of yttrium and oxygen atoms during MA process (Figs. 11(a) and (b)). And then yttrium and oxygen atoms precipitate and form a certain number of C-Y2O3 particles at the sintering and pre-heating stages (Figs. 11(c) and (d)). After the first pass rolling, these C-Y2O3 particles transform to MY2O3 particles (Fig. 11(e)). M-Y2O3 particles coarsen and new C-Y2O3 particles precipitate during the intermediate annealing process (Fig. 11(f)). New formed C-Y2O3 particles transform to M-Y2O3 particles (Fig. 11(g)) after the second rolling. This “C / M” phase transition last seven times and finally high-density M-Y2O3

particles are distributed in FCNY CSAs (Fig. 11(h)). These M-Y2O3 particles coarsened and effectively hindered the grain growth during annealing process (Fig. 11(i)). 4.2. Strengthening mechanism Strengthening mechanisms in the ODS polycrystalline materials are generally ascribed to solid-solution strengthening, grain boundary strengthening, dislocation strengthening and dispersion strengthening. Yield strength can be a simple summation of these four individual contributions and expressed as [57]:

s0.2 ¼ s0 þ DsSS þ DsGB þ DsOro þ DsDis

(2)

where s0.2 is the yield strength of ODS-CSAs, s0 is the lattice friction strength of FeCoNi CSAs, which is adopted as 161 MPa [58,59]. DsSS, DsGB, DsDis and DsOr are strengthening contributions from solid solution, grain boundary, dislocations and NPs, respectively. 4.2.1. Solid-solution strengthening For the traditional ODS alloys, the effect of solid-solution strengthening is normally attributed to the dilute solution alloys, such as chromium or tungsten [60,61]. But for CSA systems, the terms “solute” and “solvent” lose their conventional meanings. How to evaluate, or define, the precise contribution of solidsolution strengthening in HEAs, remains a challenge. He et al. [57] simply treated (FeCoNiCr)94Ti2Al4 (at.%) as a FeCoNiCr solvent matrix containing Ti þ Al solutes, and then calculated the strength enhancement caused by solid-solution hardening. For two kinds of CSAs with same composition except for different thermomechanical procedures, the solid-solution strengthening contributions are

Please cite this article as: Y. Guo et al., Microstructure and mechanical properties of oxide dispersion strengthened FeCoNi concentrated solid solution alloys, Journal of Alloys and Compounds, https://doi.org/10.1016/j.jallcom.2019.153104

12

Y. Guo et al. / Journal of Alloys and Compounds xxx (xxxx) xxx

25.4 and 14.4 MPa, respectively. The value is too small to account for the strength difference, which suggests that solid-solution is not the dominant strengthening mechanism. In FCNY and FCNTY CSAs, no yttrium or titanium is detected in the matrix according to EDS results (Table 2), suggesting that most of the Y and Ti have precipitated from the matrix. Hence, the solid-solution hardening is treated as zero for three kinds of materials in this paper.

4.2.2. Grain boundaries strengthening It is well known that refined grains improve the strength of metals and alloys because smaller grain size provides a highvolume fraction of GBs. These GBs impede dislocation movement during the plastic deformation. The relationship between yield strength and grain size can be described by the Hall-Petch equation [62]:

DsGB ¼ ky/d1/2

(3)

where ky is the Hall-Petch coefficient with the value of 275 MPa∙mm1/2 [58,63], d is the diameter, and the values of FCN, FCNY and FCNTY CSAs are 854, 483 and 456 nm, respectively. Therefore, the contribution from GB strengthening is 298, 396 and 407 MPa for FCN, FCNY and FCNTY CSAs, respectively.

4.2.3. Oxide particle strengthening In FCNY and FCNTY CSAs, nano-sized M-Y2O3 and Y2Ti2O7 particles located on the grain boundaries can effectively retard the movement of grain boundaries and thus restrain the grain growth by the well-known Zener pinning effect at the recovery and recrystallization stages. The motion of dislocations can also be inhibited by these NPs even at high temperature, which results in stable microstructures and excellent creep strength in ODS alloys. According to the Orowan strengthening mechanism, the effect of oxide dispersion strengthening based on bypass of oxide particles by dislocation loop is calculated by the following equation [64]:

DsOro ¼ 0:84

 pr  MGb ! ln ffiffiffiffi ffi q 4b pffiffiffiffiffiffiffiffiffiffiffi 3p  p 2pr 1  v 4 2f

(4)

where M is Taylor factor (3.06 for FCC), G is shear modulus (60 GPa for FeCoNi [59]), b is burgers vector (0.2525 nm for FeCoNi [65]), v is Poisson’s ratio (0.35 for FeCoNi [59]), r and f are the average radius and volume fraction of oxide particles, respectively. According to the ICP results, yttrium and titanium contents (wt.%) in FCNY and FCNTY CSAs are 0.72%, less than 0.01% and 0.63%, 1.13%, respectively. Combined with TEM and STEM results (Figs. 5 and 6) and assuming sufficient precipitation of yttrium and titanium, r and f are 6.54 nm, 1.41% and 5.28 nm, 2.29% for FCNY and FCNTY CSAs, respectively. Thus, the contribution of oxide particle strengthening in FCNY and FCNTY CSAs are calculated to be 202 and 300 MPa, respectively. 4.2.4. Dislocation strengthening For the dislocation strengthening, it can be estimated by BaileyHirsch relationship:

DsDis ¼ MaGbr1/2

(5)

where a is a constant taken as 0.2 [57], r is the dislocation density. The dislocation density is roughly determined through WilliamsonHall method [66]. True XRD peak broadening b (the observed peak broadening minuses the instrument broadening) composes of crystallite size broadening bG and strain broadening bS. Based on the assumption of Cauchy-type function, they are expressed as:

b ¼ bG þbS

(6)

bG ¼ Kl/(D$cosq)

(7)

bS ¼ 4ε$tanq

(8)

where K is a constant taken as 0.9, l is the wavelength of Cu Ka radiation (0.15406 nm), D is crystallite size, ε is micro strain and q is

Fig. 12. Plots of bcosq as a function of 4sinq, and the slope of the linear fit shows the value of micro strain ε.

Please cite this article as: Y. Guo et al., Microstructure and mechanical properties of oxide dispersion strengthened FeCoNi concentrated solid solution alloys, Journal of Alloys and Compounds, https://doi.org/10.1016/j.jallcom.2019.153104

Y. Guo et al. / Journal of Alloys and Compounds xxx (xxxx) xxx

the Bragg angle of the corresponding peak. Taking the micro strain q as variable parameter, Equation (6) can be rewritten as:

bcosq ¼ Kl/D þ 4sinq$ε

(9)

The slope and intercept of the linear fit of the bcosq-4sinq plot determine the micro strain ε and the crystallite size D. The bcosq4sinq curves of FCN, FCNY and FCNTY CSAs are plotted in Fig. 12. The micro strain ε of FCN, FCNY and FCNTY CSAs is 0.013, 0.073 and 0.088, respectively, which means that annealing treatment results in low-density dislocations in FCN CSAs and high-density dislocations in FCNY and FCNTY CSAs. The dislocation density can be derived from the micros strain:

pffiffiffi .

r ¼ 2 3ε ðD , bÞ

(10)

According to Equations (10) and (5), dislocation strengthening contribution is evaluated as 123, 244 and 276 MPa for FCN, FCNY and FCNTY CSAs, respectively. Based on the above discussion, the theoretical yield strengths of FCN, FCNY and FCNTY CSAs are 582, 1003 and 1145 MPa, respectively. Fig. 13 shows the strength contributions from these four individual mechanisms (left column) as well as the experimental values (right column). It clearly shows that the theoretical calculations are in reasonable agreement with the experimental values. The improved strength in FCNY and FCNTY CSAs originates from the refined grains and high-density nano-sized oxide particles. The discrepancy may come from the following reasons. The main error results from unclose pores in the matrix (Figs. 1(d)e(f), which act as the weak part during the tensile test for all three kinds of samples. For FCNTY CSAs, the difference between calculated and experimental value is relatively large compared with that of FCN and FCNY CSAs. The large error may come from several aspects. Firstly, the lowest bulk relative density and more unclosed pores observed in FCNTY CSAs means that the probability of forming larger pores increases and these large pores could result in premature fracture. Secondly, the effect of oxide particle strengthening in FCNTY CSAs

13

may be overestimated due to the inhomogeneous distribution of the nano-sized particles compared with that of FCNY CSAs. The theoretical strengthening contribution is calculated based on TEM and STEM results. Although several different regions are chosen to calculate the average size and volume fraction of the NPs, the observation and measurements are still carried out at the nanometer scale. Thirdly, excessive titanium in FCNTY CSAs exists in the form of titanium oxide with the diameter larger than one hundred nanometers according to our previous studies [67]. These titanium oxides may act as the source of cracking during tensile test. The above reasons may explain the considerable errors between theoretical and experimental values for FCNTY CSAs compared with those for FCN and FCNY CSAs. 4.3. Microstructure evolution of three kinds of CSAs during HR process It is generally known that discontinuous dynamic recrystallization (DDRX) dominates the grain evolution for alloys with low to medium stacking fault energy (SFE) when deformation temperature (T/Tm) is larger than 0.5. However, the precipitation of second phase particles (e.g. M-Y2O3 and Y2Ti2O7 particles) has an opposite effect on DDRX. When the pinning force exceeds the driving force for boundary bulging, DDRX is suppressed and finer DDRX grains will be formed at larger strains [68]. Moreover, recent paper reported that continuous dynamic recrystallization (CDRX) during hot deformation has been unexpectedly discovered in CrCoNi medium entropy alloys with the appearance of necklace structure which has been broadly thought as typical indication of DDRX [69]. The above findings indicate that microstructure evolution analysis could be complex for three kinds of FeCoNi-based CSAs discussed in this paper. Normally, it has been widely recognized that grains are crushed, elongated and broken during the HR process, and subgrain rotation also happens at the same time [70], which correspond to the dynamic recrystallization and texture transformation in FCN CSAs. For FCNY CSAs, nano-sized CeY2O3 particles precipitate continuously and transform to M-Y2O3 particles under hot-rolled

Fig. 13. The calculated strength contribution from different strengthening mechanisms (on the left) and experimental values (on the right) for FCN, FCNY and FCNTY CSAs. (Note: “Cal.” is short for calculation and “Exp.” is short for experiment.)

Please cite this article as: Y. Guo et al., Microstructure and mechanical properties of oxide dispersion strengthened FeCoNi concentrated solid solution alloys, Journal of Alloys and Compounds, https://doi.org/10.1016/j.jallcom.2019.153104

14

Y. Guo et al. / Journal of Alloys and Compounds xxx (xxxx) xxx

pressure. H.R.Z. Sandim et al. reported that complete static recrystallization is intensively suppressed because of the strong interaction between the fine dispersion of stable Y2O3 particles and grain boundaries after severe cold rolling reduction and annealing at high temperature [71]. For FCNTY CSAs, nano-sized semicoherent Y2Ti2O7 particles precipitate and uniformly distributed in the matrix. These faceted-shape particles remain stable during HR process and ultra-fine grain structure was also stable annealed at 1150
C up to 1000 h. This superior thermal stability of grain structure was mainly ascribed to the strong pinning effect of the nano-sized semi-coherent Y2Ti2O7 particles on grain boundary migrations in austenitic ODS steels [72,73]. In this paper, grain boundaries pinned by the NPs are clearly observed in FCNY and FCNTY CSAs (Figs. 6(a) and (b)). The classical expression for the Zener pinning force Fz by taking into account the shape of the precipitates could be expressed as [74],

(6) The values of the yield strength for FCN, FCNY and FCNTY CSAs are 559, 981 and 1050 MPa, respectively. The enhanced strength in FCNY and FCNTY CSAs originates from the refined grains and high-density nano-sized oxide particles.

Author contributions section Y. Guo conceived the project, performed most experiments and analyzed results with guidance from Y. Chang and Y. Zhang. M. Li and P. Li helped to manufacture the mechanical alloyed powders and annealed samples. C. Chen contributed towards discussion and data interpretation. Q. Zhan gave guidance on the analysis of TEM results. Y. Guo wrote the paper with assistance from all other authors. Declaration of competing interest

Fz ¼

2ks gf kv r

(11)

where ks is a geometrical factor (p for spheres), kv depends on the shape of the NPs (4p/3 for spheres), g is the specific boundary energy, r is the particles radius, and f is the volume fraction of the NPs. It is obvious from this equation that, for the same matrix, higher number density and smaller the NPs the more efficient the pinning force will be. At the pre-heating stage, a certain amount of the NPs precipitate in both FCNY and FCNTY CSAs. During the HR process, new formed M-Y2O3 and stable Y2Ti2O7 particles constantly hinder the movement of dislocations, subgrains and grains, which retard g / a transformation and preserve prior gfiber in the matrix. As discussed in Section 4.2, r and f are 6.54 nm, 1.41% and 5.28 nm, 2.29% for FCNY and FCNTY CSAs, respectively. These values quantitatively indicate that strong pinning effect in FCNTY CSAs compared with FCNY CSAs, which cause higher volume fraction of preserved g-fiber in FCNTY CSAs.

5. Conclusion In this paper, FCN, FCNY and FCNTY CSAs are fabricated by MA, SPS, hot rolling and annealing treatments. These samples are characterized by XRD, EBSD, TEM, STEM and HRTEM techniques. Mechanical properties are evaluated and quantitatively calculated based on various strengthening mechanisms. The main conclusions can be drawn as follows. (1) Raw powders are fully alloyed after milling for 60 h for all the samples. During the SPS and following heat working processes, only FCC phase is detected for the bulk materials. (2) Average grain size is 854 ± 480 nm for FCN CSAs. Grains are refined to 483 ± 172 and 456 ± 187 nm for FCNY and FCNTY CSAs with the addition of nano-sized oxide particles. (3) For FCN CSAs, the texture component changes during the hot rolling process, while this phenomenon is partially suppressed by the presence of the Y2O3 particles in FCNY CSAs, and further inhibited by the finer and higher-volumefraction Y2Ti2O7 particles in FCNTY CSAs. (4) Monoclinic Y2O3 particles with diameter around 13.1 nm are uniformly distributed in FCNY CSAs and show a semicoherent relationship with the matrix. These monoclinic Y2O3 particles are transformed from precipitated cubic Y2O3 particles. (5) Pyrochlore Y2Ti2O7 particles with the diameter around 10.6 nm are uniformly distributed in FCNTY CSAs, and also show a parallel relationship with the matrix.

No conflict of interest exits in the submission of this manuscript, and this manuscript is approved by all authors for publication. Acknowledgements This work was supported by the National Natural Science Foundation of China (No. 11775017, 11175014), the Fundamental Research Funds for the Central Universities (FRF-GF-17-B5) and Beijing Municipal Natural Science Foundation (No. 2162023). Y. Zhang was supported as part of the Energy Dissipation to Defect Evolution (EDDE), an Energy Frontier Research Center funded by the US Department of Energy, Office of Science, Basic Energy Sciences under contract number DE-AC05-00OR22725. Appendix A. Supplementary data Supplementary data to this article can be found online at https://doi.org/10.1016/j.jallcom.2019.153104. References [1] Y.F. Ye, Q. Wang, J. Lu, C.T. Liu, Y. Yang, High-entropy alloy: challenges and prospects, Mater. Today 19 (2016) 349e362, https://doi.org/10.1016/ j.mattod.2015.11.026. [2] D.B. Miracle, O.N. Senkov, A critical review of high entropy alloys and related concepts, Acta Mater. 122 (2017) 448e511, https://doi.org/10.1016/ j.actamat.2016.08.081. [3] K. Jin, B.C. Sales, G.M. Stocks, G.D. Samolyuk, M. Daene, W.J. Weber, Y. Zhang, H.B. Bei, Tailoring the physical properties of Ni-based single-phase equiatomic alloys by modifying the chemical complexity, Sci. Rep. 6 (2016), https:// doi.org/10.1038/srep20159, 20159. [4] T. Nagase, P.D. Rack, J.H. Noh, T. Egami, 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 (2015) 32e42, https://doi.org/10.1016/ j.intermet.2014.12.007. [5] Y. Zhang, S.J. Zhao, W.J. Weber, K. Nordlund, F. Granberg, F. Durabekova, Atomic-level heterogeneity and defect dynamics in concentrated solidsolution alloys, Curr. Opin. Solid State Mater. Sci. 21 (2017) 221e237, https://doi.org/10.1016/j.cossms.2017.02.002. [6] Y. Zhang, G.M. Stocks, K. Jin, C.Y. Lu, H.B. Bei, B.C. Sales, L.M. Wang, L.K. Beland, R.E. Stoller, G.D. Samolyuk, M. Caro, A. Caro, W.J. Weber, Influence of chemical disorder on energy dissipation and defect evolution in concentrated solidsolution alloys, Nat. Commun. 6 (2015), https://doi.org/10.1038/ ncomms9736, 9736. [7] S.G. Ma, Y. Zhang, Effect of Nb addition on the microstructure and properties of AlCoCrFeNi high-entropy alloy, Mater. Sci. Eng. A 532 (2012) 480e486, https://doi.org/10.1016/j.msea.2011.10.110. [8] R.R. Chen, G. Qin, H.T. Zheng, L. Wang, Y.Q. Su, Y.L. Chiu, H.S. Ding, J.J. Guo, H.Z. Fu, Composition design of high entropy alloys using the valence electron concentration to balance strength and ductility, Acta Mater. 144 (2018) 129e137, https://doi.org/10.1016/j.actamat.2017.10.058. [9] S.Q. Xia, X. Yang, T.F. Yang, S. Liu, Y. Zhang, Irradiation resistance in AlxCoCrFeNi high entropy alloys, JOM 67 (2015) 2340e2344, https://doi.org/ 10.1007/s11837-015-1568-4.

Please cite this article as: Y. Guo et al., Microstructure and mechanical properties of oxide dispersion strengthened FeCoNi concentrated solid solution alloys, Journal of Alloys and Compounds, https://doi.org/10.1016/j.jallcom.2019.153104

Y. Guo et al. / Journal of Alloys and Compounds xxx (xxxx) xxx [10] N. Hansen, Hall-Petch relation and boundary strengthening, Scr. Mater. 51 (2004) 801e806, https://doi.org/10.1016/j.scriptamat.2004.06.002. [11] M. Mabuchi, K. Higashi, Strengthening mechanisms of Mg-Si alloys, Acta Mater. 44 (1996) 4611e4618, https://doi.org/10.1016/1359-6454(96)000729. [12] J.S. Benjamin, Dispersion strengthened superalloys by mechanical alloying, Metall. Trans. 1 (1970) 2943e2951, https://doi.org/10.1007/BF03037835. [13] M. Naginia, R. Vijay, K.V. Rajulapati, A.V. Reddy, G. Sundararajan, Microstructure-mechanical property correlation in oxide dispersion strengthened 18Cr ferritic steel, Mater. Sci. Eng. A 708 (2017) 451e459, https://doi.org/ 10.1016/j.msea.2017.10.023. [14] S. Yamashita, Y. Yano, S. Ohtsuka, T. Yoshitake, T. Kaito, S. Koyama, K. Tanaka, Irradiation behavior evaluation of oxide dispersion strengthened ferritic steel cladding tubes irradiated in JOYO, J. Nucl. Mater. 442 (2013) 417e424, https:// doi.org/10.1016/j.jnucmat.2013.04.051. [15] N. Oono, S. Ukai, S. Kondo, O. Hashitomi, A. Kimura, Irradiation effects in oxide dispersion strengthened (ODS) Ni-base alloys for Gen. IV nuclear reactors, J. Nucl. Mater. 465 (2015) 835e839, https://doi.org/10.1016/ j.jnucmat.2015.06.057. [16] H. Hadraba, Z. Chlup, A. Dlouhy, F. Dobes, P. Roupcova, M. Vilemova, J. Matejicek, Oxide dispersion strengthened CoCrFeNiMn high-entropy alloy, Mater. Sci. Eng. A 689 (2017), https://doi.org/10.1016/j.msea.2017.02.068, 262-256. [17] S. Praveen, A. Anupam, T. Sirasani, B.S. Murty, R.S. Kottada, Characterization of oxide dispersed AlCoCrFe high entropy alloy synthesized by mechanical alloying and spark plasma sintering, Trans. Indian Inst. Met. 66 (2013) 369e373, https://doi.org/10.1007/s12666-013-0268-4. [18] B. Gwalani, R.M. Pohan, J. Lee, B. Lee, R. Banerjee, H.J. Ryu, S.H. Hong, Highentropy alloy strengthened by in situ formation of entropy-stabilized nanodispersoids, Sci. Rep. 8 (2018), https://doi.org/10.1038/s41598-018-32552-6, 14085. [19] X.Y. Liu, H. Yin, Y. Xu, Microstructure, mechanical and tribological properties of oxide dispersion strengthened high-entropy alloys, Materials 10 (2017) 1312, https://doi.org/10.3390/ma10111312. [20] B. Jia, X.J. Liu, H. Wang, Y. Wu, Z.P. Lu, Microstructure and mechanical properties of FeCoNiCr high-entropy alloy strengthened by nano-Y2O3 dispersion, Sci. China Technol. Sci. 61 (2018) 179e183, https://doi.org/10.1007/s11431017-9115-5. [21] Ł. Rogal, D. Kalita, L. Litynska-Dobrzynska, CoCrFeMnNi high entropy alloy matrix nanocomposite with addition of Al2O3, Intermetallics 86 (2017) 104e109, https://doi.org/10.1016/j.intermet.2017.03.019. [22] S.F. Yang, Y. Zhang, X. Yan, H. Zhou, J.H. Pi, D.Z. Zhu, Deformation twins and interface characteristics of nano-Al2O3 reinforced Al0.4FeCrCo1.5NiTi0.3 high entropy alloy composites, Mater. Chem. Phys. 210 (2018) 240e244, https:// doi.org/10.1016/j.matchemphys.2017.11.037. [23] H. Prasad, S. Singh, B.B. Panigrahi, Mechanical activated synthesis of alumina dispersed FeNiCoCrAlMn high entropy alloy, J. Alloy. Comp. 692 (2017) 720e726, https://doi.org/10.1016/j.jallcom.2016.09.080. [24] Y. Zhang, T. Egami, W.J. Weber, Dissipation of radiation energy in concentrated alloys: unique defect properties and microstructural evolution, MRS Bull. 44 (2019) 798e811. https://doi.org/10.1557/mrs.2019.233. [25] C.Y. Lu, L.L. Niu, N.J. Chen, K. Jin, T.N. Yang, P.Y. Xiu, Y. Zhang, F. Gao, H. Bei, S. Shi, M.R. He, I.M. Robertson, W.J. Weber, L.M. Wang, Enhancing radiation tolerance by controlling defect mobility and migration pathways in multicomponent single-phase alloys, Nat. Commun. 7 (2016), https://doi.org/ 10.1038/ncomms13564, 13564. [26] K. Jin, C. Lu, L.M. Wang, J. Qu, W.J. Weber, Y. Zhang, H. Bei, Effects of compositional complexity on the ion-irradiation induced swelling and hardening in Ni-containing equiatomic alloys, Scr. Mater. 119 (2016) 65e70, https://doi.org/10.1016/j.scriptamat.2016.03.030. [27] T.N. Yang, C.Y. Lu, K. Jin, M.L. Crespillo, Y. Zhang, H.B. Bei, L.M. Wang, The effect of injected interstitials on voids formation in self-ion irradiated nickel containing concentrated solid solution alloys, J. Nucl. Mater. 488 (2017) 328e337, https://doi.org/10.1016/j.jnucmat.2017.02.026. [28] T.N. Yang, C.Y. Lu, G. Velisa, K. Jin, P.Y. Xiu, M.L. Crespillo, Y. Zhang, H.B. Bei, L.M. Wang, Effect of alloying elements on defect evolution in Ni-20X binary alloys, Acta Mater. 151 (2018) 159e168, https://doi.org/10.1016/ j.actamat.2018.03.054. [29] C.H. Tsau, S.X. Lin, C.H. Fang, Microstructures and corrosion behaviors of FeCoNi and CrFeCoNi equimolar alloys, Mater. Chem. Phys. 186 (2017) 534e540, https://doi.org/10.1016/j.matchemphys.2016.11.033. [30] P.P. Li, A.D. Wang, C.T. Liu, Composition dependence of structure, physical and mechanical properties of FeCoNi(MnAl)x high entropy alloys, Intermetallics 87 (2017) 21e26, https://doi.org/10.1016/j.intermet.2017.04.007. [31] G.D. Sathiaraj, W. Skrotzki, A. Pukenas, R. Schaarschuch, R.J. Immanuel, S.K. Panigrahi, J.A. Chelvane, S.S.S. Kumard, Effect of annealing on the microstructure and texture of cold rolled CrCoNi medium-entropy alloy, Intermetallics 101 (2018) 87e98, https://doi.org/10.1016/ j.intermet.2018.07.014. [32] A. Shabani, M.R. Toroghinejad, A. Shafyei, P. Cavaliere, Effect of cold-rolling on microstructure, texture and mechanical properties of an equiatomic FeCrCuMnNi high entropy alloy, Materialia 1 (2018) 175e184, https://doi.org/ 10.1016/j.mtla.2018.06.004. [33] L.A.I. Kestens, H. Pirgazi, Texture formation in metal alloys with cubic crystal structures, Mater. Sci. Technol. 32 (2016) 1303e1315, https://doi.org/10.1080/

15

02670836.2016.1231746. [34] Ł. Rogal, D. Kalita, A. Tarasek, P. Bobrowski, F. Czerwinski, Effect of SiC nanoparticles on microstructure and mechanical properties of the CoCrFeMnNi high entropy alloy, J. Alloy. Comp. 708 (2017) 344e352, https://doi.org/ 10.1016/j.jallcom.2017.02.274. [35] N. Oono, Q.X. Tang, S. Ukai, Oxide particle refinement in Ni-based ODS alloy, Mater. Sci. Eng. A 649 (2016) 250e253, https://doi.org/10.1016/ j.msea.2015.09.094. [36] T. Yamashiro, S. Ukai, N. Oono, S. Ohtsuka, T. Kaito, Microstructural stability of 11Cr ODS steel, J. Nucl. Mater. 472 (2016) 247e251, https://doi.org/10.1016/ j.jnucmat.2016.01.002. [37] L. Zhang, S. Ukai, T. Hoshino, S. Hayashi, X.H. Qu, Y2O3 evolution and dispersion refinement in Co-base ODS alloys, Acta Mater. 57 (2009) 3671e3682, https://doi.org/10.1016/j.actamat.2009.04.033. [38] T. Atou, K. Kusaba, K. Fukuoka, M. Kikuchi, Y. Syono, Shock-induced phase transition of M2O3 (M¼Sc, Y, Sm, Gd, and In)-type compounds, J. Solid State Chem. 89 (1990) 378e384, https://doi.org/10.1016/0022-4596(90)90280-B. [39] H.O. Zhuo, J.C. Tang, N. Ye, A novel approach for strengthening Cu-Y2O3 composites by in situ reaction at liquidus temperature, Mater. Sci. Eng. A 584 (2013) 1e6, https://doi.org/10.1016/j.msea.2013.07.007. [40] E. Chtoun, N. Hanebali, P. Garnier, J.M. Kiat, X-rays and neutrons Rietveld analysis of the solid solutions (1-x)A2Ti2O(7-x)MgTiO3 (A¼Y or Eu), Eur. J. Solid Inorg. Chem. 34 (1997) 553e561. [41] H. Yu, S. Ukai, N. Oono, Tensile properties of Co-based oxide dispersion strengthened superalloys, J. Alloy. Comp. 714 (2017) 715e724, https:// doi.org/10.1016/j.jallcom.2017.04.276. [42] D.A. McClintock, M.A. Sokolov, D.T. Hoelzer, R.K. Nanstad, Mechanical properties of irradiated ODS-EUROFER and nanocluster strengthened 14YWT, J. Nucl. Mater. 392 (2009) 353e359, https://doi.org/10.1016/ j.jnucmat.2009.03.024. [43] D.T. Hoelzer, K.A. Unocic, M.A. Sokolov, Z. Feng, Joining of 14YWT and F82H by friction stir welding, J. Nucl. Mater. 442 (2013) S529eS534, https://doi.org/ 10.1016/j.jnucmat.2013.04.027. [44] M.L. Hamilton, D.S. Genes, R.J. Lobsinger, G.D. Johnson, W.F. Brown, M.M. Paxton, R.J. Puigh, C.R. Eiholzer, C. Martinez, M.A. Blotter, Fabrication Technological Development of the Oxide Dispersion Strengthened Alloy MA957 for Fast Reactor Applications, PNNL-13168, PNL, Richland, WA, 2000. https://www.osti.gov/biblio/752621-Y5dmWn/webviewable/. [45] Y.F. Li, T. Nagasaka, T. Muroga, A. Kimura, S. Ukai, High-temperature mechanical properties and microstructure of 9Cr oxide dispersion strengthened steel compared with RAFMs, Fusion Eng. Des. 86 (2011) 2495e2499, https:// doi.org/10.1016/j.fusengdes.2011.03.004. €ning, M. Rieth, J. Hoffmann, A. Mo € slang, Production, microstructure and [46] T. Gra mechanical properties of two different austenitic ODS steels, J. Nucl. Mater. 487 (2017) 348e361, https://doi.org/10.1016/j.jnucmat.2017.02.034. lez-Carrasco, G. Ciapetti, M.A. Montealegre, S. Pagani, J. Chao, [47] J.L. Gonza N. Baldini, Evaluation of mechanical properties and biological response of an alumina-forming Ni-free ferritic alloy, Biomaterials 26 (2005) 3861e3871, https://doi.org/10.1016/j.biomaterials.2004.09.058. [48] Z.J. Zhou, S.Y. Sun, L. Zou, Y.L. Schneide, S. Schmauder, M. Wang, Enhanced strength and high temperature resistance of 25Cr20Ni ODS austenitic alloy through thermo-mechanical treatment and addition of Mo, Fusion Eng. Des. 138 (2019) 175e182, https://doi.org/10.1016/j.fusengdes.2018.11.020. [49] M. Wang, Z.J. Zhou, H.Y. Sun, H.L. Hu, S.F. Li, Microstructural observation and tensile properties of ODS-304 austenitic steel, Mater. Sci. Eng. A 559 (2013) 287e292, https://doi.org/10.1016/j.msea.2012.08.099. [50] J.R.O. Leo, S. Pirfo Barroso, M.E. Fitzpatrick, M. Wang, Z. Zhou, Microstructure, tensile and creep properties of an austenitic ODS 316L steel, Mater. Sci. Eng. A 749 (2019) 158e165, https://doi.org/10.1016/j.msea.2019.02.014. [51] M. Foex, J.P. Traverse, Investigation about crystalline transformation in rare earths sesquioxides at hifh temperature, Rev. Int. Hautes Temp. Refract. 3 (1966) 429e453. re, C. Landron, P. Melin, D.L. Pricea, J.P. Coutures, [52] L. Hennet, D. Thiaudie rar, M.L. Saboungi, Melting behavior of levitated Y2O3, Appl. Phys. Lett. J.F. Be 83 (2003) 3305e3307, https://doi.org/10.1063/1.1621090. , Phase transitions in yttrium oxide at [53] E. Husson, C. Proust, P. Gillet, J.P. Itie high pressure studied by Raman spectroscopy, Mater. Res. Bull. 34 (1999) 2085e2092, https://doi.org/10.1016/S0025-5408(99)00205-6. [54] S. Yamashita, S. Ohtsuka, N. Akasaka, S. Ukai, S. Ohnuki, Formation of nanoscale complex oxide particles in mechanically alloyed ferritic steel, Philos. Mag. Lett. 84 (2004) 525e529, https://doi.org/10.1080/ 09500830412331303609. [55] P. He, P.L. Gao, Q. Tian, J.M. Lu, W.Z. Yao, An in situ SANS study of nanoparticles formation in 9Cr ODS steel powders, Mater. Lett. 209 (2017) 535e538, https://doi.org/10.1016/j.matlet.2017.08.051. [56] H. Oka, M. Watanabe, S. Ohnuki, N. Hashimoto, S. Yamashita, S. Ohtsuka, Effects of milling process and alloying additions on oxide particle dispersion in austenitic stainless steel, J. Nucl. Mater. 447 (2014) 248e253, https://doi.org/ 10.1016/j.jnucmat.2014.01.025. [57] J.Y. He, H. Wang, H.L. Huang, X.D. Xu, M.W. Chen, Y. Wu, X.J. Liu, T.G. Nieh, K. An, Z.P. Lu, A precipitation-hardened high-entropy alloy with outstanding tensile properties, Acta Mater. 102 (2016) 187e196, https://doi.org/10.1016/ j.actamat.2015.08.076. [58] Y.Y. Zhao, T.G. Nieh, Correlation between lattice distortion and friction stress in Ni-based equiatomic alloys, Intermetallics 86 (2017) 45e50, https://

Please cite this article as: Y. Guo et al., Microstructure and mechanical properties of oxide dispersion strengthened FeCoNi concentrated solid solution alloys, Journal of Alloys and Compounds, https://doi.org/10.1016/j.jallcom.2019.153104

16

Y. Guo et al. / Journal of Alloys and Compounds xxx (xxxx) xxx

doi.org/10.1016/j.intermet.2017.03.011. [59] Z. Wu, H. Bei, G.M. Pharr, E.P. George, Temperature dependence of the mechanical properties of equiatomic solid solution alloys with face-centered cubic crystal structures, Acta Mater. 81 (2014) 428e441, https://doi.org/ 10.1016/j.actamat.2014.08.026. , J. Malaplate, J. Garnier, F. De Geuser, F. Barcelo, P. Wident, [60] M. Dade A. Deschamps, Influence of microstructural parameters on the mechanical properties of oxide dispersion strengthened Fe-14Cr steels, Acta Mater. 127 (2017) 165e177, https://doi.org/10.1016/j.actamat.2017.01.026. [61] J.J. Shen, Y.F. Li, F. Li, H.L. Yang, Z.S. Zhao, S. Kano, Y. Matsukawa, Y. Satoh, H. Abe, Microstructural characterization and strengthening mechanisms of a 12Cr-ODS steel, Mater. Sci. Eng. A 673 (2016) 624e632, https://doi.org/ 10.1016/j.msea.2016.07.030. [62] M. Praud, F. Mompiou, J. Malaplate, D. Caillard, J. Garnier, A. Steckmeyer, B. Fournier, Study of the deformation mechanisms in a Fe-14% Cr ODS alloy, J. Nucl. Mater. 428 (2012) 90e97, https://doi.org/10.1016/ j.jnucmat.2011.10.046. [63] Z.G. Wu, Temperature and Alloying Effects on the Mechanical Properties of Equiatomic FCC Solid Solution Alloys, PhD Diss, University of Tennessee, 2014. http://trace.tennessee.edu/utk_graddiss/2884. [64] S.M.S. Aghamiri, N. Oono, S. Ukai, R. Kasada, H. Noto, Y. Hishinuma, T. Muroga, Microstructure and mechanical properties of mechanically alloyed ODS copper alloy for fusion material application, Nucl. Mater. Energy. 15 (2018) 17e22, https://doi.org/10.1016/j.nme.2018.05.019. [65] C. Varvenne, A. Luque, W.A. Curtin, Theory of strengthening in fcc high entropy alloys, Acta Mater. 118 (2016) 164e176, https://doi.org/10.1016/ j.actamat.2016.07.040. [66] G.K. Williamson, W.H. Hall, X-ray line broadening from filed aluminium and wolfram, Acta Metall. 1 (1953) 22e31, https://doi.org/10.1016/0001-6160(53)

90006-6. [67] Y. Guo, M. Li, C. Chen, P. Li, W. Li, Q. Ji, Y. Zhang, Y. Chang, Oxide Dispersion Strengthened FeCoNi Concentrated Solid-Solution Alloys Synthesized by Mechanical Alloying, Intermetallics, 2018. [68] T. Sakai, A. Belyakov, R. Kaibyshev, H. Miura, J.J. Jonas, Dynamic and postdynamic recrystallization under hot, cold and severe plastic deformation conditions, Prog. Mater. Sci. 60 (2014) 130e207, https://doi.org/10.1016/ j.pmatsci.2013.09.002. [69] G. He, Y. Zhao, B. Gan, X. Sheng, Y. Liu, L. Tan, Mechanism of grain refinement in an equiatomic medium-entropy alloy CrCoNi during hot deformation, J. Alloy. Comp. 815 (2020), https://doi.org/10.1016/j.jallcom.2019.152382, 152382. [70] S.E. Ion, F.J. Humphreys, S.H. White, Dynamic recrystallisation and the development of microstructure during the high temperature deformation of magnesium, Acta Metall. 30 (1982) 1909e1919, https://doi.org/10.1016/ 0001-6160(82)90031-1. [71] H.R.Z. Sandim, R.A. Renzetti, A.F. Padilha, D. Raabe, M. Klimenkov, R. Lindau, €slang, Annealing behavior of ferritic-martensitic 9%Cr-ODS-Eurofer A. Mo steel, Mater. Sci. Eng. A 527 (2010) 3602e3608. [72] H. Oka, M. Watanabe, N. Hashimoto, S. Ohnuki, S. Yamashita, S. Ohtsuka, Morphology of oxide particles in ODS austenitic stainless steel, J. Nucl. Mater. 442 (2013) 164e168, https://doi.org/10.1016/j.jnucmat.2013.04.073. [73] X. Mao, K.H. Oh, J. Jang, Evolution of ultrafine grained microstructure and nano-sized semicoherent oxide particles in austenitic oxide dispersion strengthened steel, Mater. Char. 117 (2016) 91e98, https://doi.org/10.1016/ j.matchar.2016.04.022. chet, M. Militzer, A note on grain size dependent pinning, Scr. Mater. 52 [74] Y. Bre (2005) 1299e1303, https://doi.org/10.1016/j.scriptamat.2005.02.021.

Please cite this article as: Y. Guo et al., Microstructure and mechanical properties of oxide dispersion strengthened FeCoNi concentrated solid solution alloys, Journal of Alloys and Compounds, https://doi.org/10.1016/j.jallcom.2019.153104