Tailoring microstructures and tensile properties of a precipitation-strengthened (FeCoNi)94Ti6 medium-entropy alloy

Journal Pre-proof Tailoring microstructures and tensile properties of a precipitation-strengthened (FeCoNi)94Ti6 medium-entropy alloy Y. Chen, H.W. Deng, Z.M. Xie, M.M. Wang, J.F. Yang, T. Zhang, Y. Xiong, R. Liu, X.P. Wang, Q.F. Fang, C.S. Liu PII:

S0925-8388(20)30820-3

DOI:

https://doi.org/10.1016/j.jallcom.2020.154457

Reference:

JALCOM 154457

To appear in:

Journal of Alloys and Compounds

Received Date: 8 November 2019 Revised Date:

19 February 2020

Accepted Date: 21 February 2020

Please cite this article as: Y. Chen, H.W. Deng, Z.M. Xie, M.M. Wang, J.F. Yang, T. Zhang, Y. Xiong, R. Liu, X.P. Wang, Q.F. Fang, C.S. Liu, Tailoring microstructures and tensile properties of a precipitationstrengthened (FeCoNi)94Ti6 medium-entropy alloy, Journal of Alloys and Compounds (2020), doi: https://doi.org/10.1016/j.jallcom.2020.154457. 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. © 2020 Published by Elsevier B.V.

Credit Author Statement: Yue

Chen:

Conceptualization, Methodology,

Investigation,

Data

Curation,

Writing-Original Draft, Haowei Deng: Investigation, Data Curation, Zhouming Xie: Methodology, Formal analysis, Data Curation, Writing-Review & Editing, Mingming Wang: Investigation, Junfeng Yang: Formal analysis, Data Curation, Tao Zhang: Project administration, Funding acquisition, Writing-Review & Editing, Ying Xiong: Funding acquisition, Rui Liu: Formal analysis, Xianping Wang: Supervision, Qianfeng

Fang:

Supervision,

Project

Changsong Liu: Project administration.

administration,

Funding

acquisition,

Tailoring microstructures and tensile properties of a precipitation-strengthened (FeCoNi)94Ti6 medium-entropy alloy Y. Chen a, b, H.W. Deng

a, b

, Z.M. Xie

,a

, M.M. Wang

a, b

, J.F. Yang a, T. Zhang

,c

, Y. Xiong d, R.

Liu a, X.P. Wang a, Q.F. Fang a, C.S. Liu a a)

Key Laboratory of Materials Physics, Institute of Solid State Physics, Chinese Academy of Sciences, Hefei, Anhui

230031, China b)

University of Science and Technology of China, Hefei, Anhui 230026, China

c)

School of Physics and Electronic Engineering, Guang Zhou University, Guangzhou, Guangdong, China

d)

State Key Laboratory of Environment-friendly Energy Materials, Southwest University of Science and Technology,

Mianyang 621010, China

Abstract The coherent precipitate L12-type Ni3Ti nanoclusters and nano-sized particles in a fcc-FeCoNi medium-entropy alloy (MEA) matrix were designed and realized through minor Ti addition. The mechanical properties were modulated by thermo-mechanical processing and various annealing treatments, which leads to different hardening mechanisms such as dislocation strengthening and precipitation strengthening. An excellent combination of yield strength (~ 893 MPa), ultimate tensile strength (UTS ~ 1263 MPa) and ductility (~ 24% elongation) can be achieved because of the precipitation strengthening from the L12-type Ni3Ti nano-sized precipitates. The present study is important not only for understanding the strengthening mechanisms of metallic materials, but also provides a generic way to obtain the desirable mechanical properties of alloys via microstructural regulation. Key words: Medium-entropy alloy; Precipitation strengthening; Coherent precipitates; Mechanical property controlling

Corresponding author: [email protected], Tel:+86-551-65596675 Corresponding author: [email protected], Tel:+86-551-65591175 1

1.

Introduction Metallic alloys are important structural materials. Conventional alloys are usually composed of

one principal element and the content of this solvent element is always more than 50% (at.%). The improvement of overall properties, especially the mechanical properties, relies on the addition of other minor solute elements to alter local stress fields and the corresponding dislocation movements. However, unexpected harmful phases such as the intermetallics and additional bulk-metallic glass may appear when the content of other solute elements exceeds solid solubility [1, 2]. In recent years, a new alloy design philosophy, i.e., medium-entropy alloys (MEAs) and high-entropy alloys (HEAs), has been proposed and attracted lots of attention due to their unique compositions, structures and properties [3-8]. High configurational entropy together with the decreased Gibbs free energy in these alloys retard formation of intermetallics and facilitate the formation of single solid solution phase, such as face-centered cubic (fcc), body-centered cubic (bcc) and hexagonal cubic phase (hcp) [9-11]. MEAs and HEAs have exhibited numerous interesting properties and unique characteristics, and attained combinations of physical and mechanical properties unattainable in conventional alloys [12-17]. Normally, single-phase fcc MEAs or HEAs show exceptional fracture toughness and outstanding ductility from room-temperature (RT) to cryogenic temperatures [18, 19]. For example, the classical fcc CrMnFeCoNi Cantor alloy exhibits exceptional tensile plasticity of ~80% and ~100% at RT and the cryogenic temperature, respectively [19]. However, these single-phase fcc MEAs always have poor strength especially the relatively low yield strength (YS), which will hinder their use in many structural applications [14]. It’s known that MEAs or HEAs have a large fluctuation of local lattice distortion which can impede dislocation glide and thereby lead to 2

strengthening [20, 21]. But this “severe lattice distortion” has been indicated just providing Labusch-type weak pinning on dislocation motion [21, 22]. Therefore, effective approaches to enhance the YS of fcc-structured MEAs with a simultaneous good ductility are still urgently needed. For the single-phase fcc MEAs or HEAs, a lot of methods have been proposed to improve the YS, such as the dislocation strengthening [23], introducing nano-twining [24], constructing hierarchical structures [25], gradient structures [26] and bimodal grain size microstructures [27]. Among these, mechanical treatments like severe plastic deformation and cold rolling induced dislocation strengthening are often applied to fcc-structured MEAs or HEAs to improve their strength [28, 29]. The strength of the CoCrNi MEA and FeCrNiCoMn HEA increased significantly after a multi-step equal channel angular pressing (ECAP) and cold rolling, respectively, but their ductility cannot retain the initial high level due to the shear localization [19, 28]. For multi-phase metals, the designed metastable phase [30, 31] and the precipitation particles [32] have been indicated effective methods for adjusting comprehensive mechanical properties in conventional alloys. Recently, precipitation hardening has been widely used to strengthen several fcc HEAs, such as the FeCoNiCr-based HEA, by adding precipitate-forming elements, like Mo, Nb, Al and Ti [33-35]. The strengthening effect is significantly relied on the precipitate type, size, and distribution [33]. For example, Mo has been demonstrated to facilitate formation of hard and brittle intermetallic sigma phase in fcc-CoCrFeMnNi and fcc-CoCrFeNiAl0.5 matrix [36, 37], which provides moderate strength increment but to some extent degrades the plasticity. For the FeCoNi-based HEA strengthened by nano-sized coherent L12-type Ni3(Al, Ti) precipitates through addition of 7 at.% Ti and 7 at.% Al, a combination of simultaneous high YS (1028 MPa) and good ductility (a total tensile elongation of 48%) at ambient temperature has been achieved [38]. These above findings trigger a 3

high level of interest about the synergistic effects of proper mechanical processing and precipitation hardening with coherent precipitates on both strength and ductility in fcc-MEAs and -HEAs. It’s known that most fcc-MEAs and -HEAs often contain the Cr element. However, Cr element would cause a much limited ductility effect on nanoparticles when compared with Fe and Co elements [38]. Moreover, harmful brittle phases such as the brittle σ phase may be introduced at grain boundaries because of too much Cr, which will degrade the mechanical properties of alloys [39]. In this work, a MEA alloy of FeCoNi (all elements in equal atomic ratio) with a single fcc solid solution phase has been chosen as the basic alloy. The FeCoNi base matrix has a large solid solubility of Ti. In addition, the enthalpy of formation (∆Hf) values between Fe, Co, Ni and Ti are -418, -386, -435 meV/atom, respectively. These strongly negative ∆Hf prefer forming intermetallic phases such as the L12-type Ni3Ti precipitates under certain heat treatments [40]. A (FeCoNi)94Ti6 (at.%) alloy was designed and prepared by vacuum arc melting, thermomechanical processing and post-deformation annealing or quenching. A systematic study of the alloying, mechanical processing and heat treatments induced precipitation strengthening effects on the microstructure, phase evolution and mechanical properties of the (FeCoNi)94Ti6 MEA has been carried out. This work would offer an analysis how the annealing temperature affects the precipitation process, tensile strength and ductility, and thus provide a guidance for regulating the strength and ductility of this (FeCoNi)94Ti6 MEA or other precipitation strengthened alloys. 2.

Experimental details

2.1. Preparing materials Ingots with nominal components of FeCoNi and (FeCoNi)94Ti6 (Fe31.3Co31.3Ni31.3Ti6 in at.%) were obtained by vacuum arc melting in a water-cooled copper crucible under a Ti-gettered 4

high-purity argon atmosphere. The starting materials of Fe, Co, Ni, Ti were all at least 99.9 wt% purity. All the samples were repeatedly melted for 6 times to insure the chemical homogeneity, and further homogenized at 1100 oC for ~2 h followed by water quenching. The synthesized MEAs are referred to as homogenized -FeCoNi and -(FeCoNi)94Ti6 MEAs. Then the homogenized MEA ingots are further hot-forged at 1100 oC to sheets with a thickness of 5 mm from the initial thickness of ~10 mm through 3-step forging. An intermediate soaking treatment was performed at 1100 oC for 5 min after the first and second forging passes in order to minimize the occurrence of plastic instabilities and edge cracks, and specimens were water quenched after the last forging. Finally, the forged MEAs were isochronally annealed at 200 °C, 300 °C, 400 °C, 500 °C and 1000 °C for 1 hour in argon atmosphere, respectively, and subsequently followed by furnace cooling or water quenching. 2.2. Tensile tests The dog-bone shaped specimens of both the homogenized, forged and annealed FeCoNi and (FeCoNi)94Ti6 MEAs for tensile tests were cut with a cross section of gage length of 5 mm, a width of 1.5 mm and a thickness of about 1 mm. The tensile specimen surfaces and sides were grinded to a 2000-grit SiC paper. Tensile tests were carried out using an Instron-5967 machine at RT with a constant speed of 0.3 mm/min, corresponding to an approximate strain rate of 10-3 s-1. 2.3. Phase identification and microstructural characterization Phase identification was examined by X-ray diffraction spectrometer (XRD, X’pert Pro MPD, Philips) with a Cu Kα radiation (0.1541 nm). The measurements were conducted over a scanning range from 30o to 100o using a step of 0.016o and a counting time of 5.08 s per step. The instrumental aberrations was corrected using a standard Si powder. The metallography was obtained using an optical microscope (ZEISS-AX10) after mirror 5

polishing and etching (25% H2SO4 aqueous solution). Specimens for transmission electron microscope (TEM) observations were initially ground to a thickness of about 50 µm, punched into 3 mm-diameter disks, and then electro-polished and thinned by twin-jet in Tenuple-5 at 25 V with an electrolyte of 20 vol % nitric acid and 80% methanol at -30 oC. The TEM observations were performed using a Tecnai G2 F20 TEM coupled with an energy-dispersive X-ray spectroscopy (EDS, INCA) analytical system operating at 200 kV. 3.

Results and discussion

3.1 Microstructures and tensile properties of the FeCoNi alloy Figs. 1a-1c display the metallographic images of the FeCoNi-H (homogenized), FeCoNi-H-F (homogenized + forged) and FeCoNi-H-F-1000A (homogenized + forged + 1000 oC annealed) FeCoNi MEA alloys, respectively. Equiaxed grains with average grain sizes of ~ 400 µm, 74 µm and 64 µm were observed in FeCoNi-H, FeCoNi-H-F and FeCoNi-H-F-1000A alloys, respectively, suggesting the grain refinement of FeCoNi alloys after the 3-step forging. The corresponding XRD pattern (Fig. 1d) shows that the phase constitution of the FeCoNi-H (homogenized) MEA alloy is simple and single solid solution with a typical fcc structure is indexed. After forging and 1000 oC annealing, the FeCoNi MEA alloys maintain the single phase fcc structure. Fig. 2a reveals the RT engineering stress - strain curves of the FeCoNi-H MEA alloys before and after post-deformation annealing at various temperatures. It can be seen from Fig. 2b and Table 1 that after the 3-step forging the yield stress (YS) and ultimate tensile strength (UTS) increase rapidly from 103 MPa and 320 MPa (for FeCoNi-H) to 454 MPa and 551 MPa (for FeCoNi-H-F), respectively. In this work, the YS is defined as the stress at 0.2% plastic strain. While after forging, the total elongation (TE) decreases from 49% to 26%. The annealing treatments at 400 oC or 500 oC 6

followed by furnace cooling have little influences on both the strength and ductility of the FeCoNi alloy as shown in the Fig. 2b. But after 1000 oC annealing followed by furnace cooling or water quenching, the YS deceases from the 551 MPa (for the forged one) to 155 MPa and the TE recovers from 26% to ~40%. This strength degradation and plastic recovery process corresponds to the typical microstructure evolution process which involves the recovery and recrystallization induced by the high temperature annealing. 3.2 Microstructures and tensile properties of the homogenized + forged + annealed (FeCoNi)94Ti6 alloy Fig. 3a displays the metallographic image of the homogenized (FeCoNi)94Ti6 (Ti6-H hereafter) alloy. Equiaxed grains with an average grain size of ~ 175 µm were observed in this Ti6-H MEA alloy, which is smaller than that of the FeCoNi-H alloy. This reduced grain size in the Ti6-H may come from the 6 at.% Ti addition. The corresponding XRD pattern shows that the phase constitution of this Ti6-H alloy is also a simple fcc structure. However, in the Ti6-H specimen, the fcc peaks are noted to slightly shift toward a smaller angle compared to the FeCoNi-H one as indicated by the locally amplified (111) diffraction peak in Fig. 3b inset. The shift of the diffraction peak to lower angle depends upon the dilated lattice constant as a result of Ti dissolution into the FeCoNi crystalline lattice, because the Ti has largest atomic radius of 1.45 Å to that of Fe (1.24 Å), Co (1.25 Å) and Ni (1.25 Å) [41]. After the 3-step hot forging, the Ti6-H-F still has equiaxed grains (from the forging surface) with grain sizes varying from 100 to 180 µm and an average grain size of 145 µm as shown in Fig. 4a. The following annealing at 200 oC ~ 1000 oC shows little effect on the grain shape and size. The average grain sizes of Ti6-H-F-200A, Ti6-H-F-500A and Ti6-H-F-1000A are 156 µm, 148 µm and 7

140 µm, respectively, as shown in Fig. 4b-4d. In addition, some micron-sized annealing twins are observed in these forged and annealed Ti6-H alloys, suggesting their low stacking fault energies. Fig. 5a reveals the RT engineering stress - strain curves of the Ti6-H MEA alloys before and after post-deformation annealing at various temperatures. The YS of the Ti6-H is 290 MPa, which is about 3 times that of the FeCoNi-H. Meanwhile, the UTS and TE of the Ti6-H are 592 MPa and 54%, respectively, as presented in Fig. 5b and Table 1, which are also larger than that of the FeCoNi-H. The increased strength is intuitively attributed to the solid solution strengthening from the Ti addition and the refined grains. After the 3-step forging, the YS of the Ti6-H-F increases significantly to 697 MPa which is about 400 MPa higher than that of its homogenized state, but on the other hand the TE decreases to 26%. The increase in strength and decrease in ductility come from the forging deformation which introduces a high density of defects like dislocations as discussed in the following microstructure analysis section. These pre-built dislocations accelerate an effective dislocation strengthening and could be additional nucleation sites for precipitation and hence increase the volume fraction of the precipitates, but on the other hand excess defects tend to cause stress concentration and limited ductility. After post-deformation annealing, there are abnormal increases in both YS and UTS especially in the Ti6-H-F-1000A, which are completely different from that in the annealed FeCoNi-H-F specimens. It’s know that high temperature annealing often produces recovered and/or recrystallized structures and thus weakens the deformation strengthening. After 1000 oC annealing and furnace cooling, the Ti6-H-F-1000A exhibits a super-high YS of 893 MPa and an UTS of 1263 MPa, respectively. The further increased strength in this Ti6-H-F-1000A alloy (compared to the Ti6-H-F) apparently comes from the other strengthening mechanism which is indicated the precipitation 8

strengthening as detailedly analyzed in the following section. It’s worth pointing out that in this work the annealing includes two main processes: i) 1h soaking process at the target temperatures and ii) slow cooling in the furnace. In order to clarify the temperature dependence of precipitation, the tensile curves of the water quenched and low temperature annealed Ti6-H-F specimens are also presented in Fig. 5a. For the 1000 oC water quenched (WQ) Ti6-H-F-1000WQ, the YS, UTS and TE are 290 MPa, 663 MPa and 54% as presented in Fig. 5b and Table 1, respectively, which are similar to that of the Ti6-H. That’s to say, the 1000 oC soaking just maintains the single fcc structured (FeCoNi)94Ti6 matrix and leads to a strength degradation and plastic recovery to the Ti6-H-F alloy. Thus the precipitation could be considered to occur during the slowly cooling process from 1000 oC to a certain temperature. The tensile properties of the Ti6-H-F-200A, Ti6-H-F-300A, Ti6-H-F-400A and Ti6-H-F-500A were evaluated as shown in Fig. 5. For the 200 oC annealed Ti6-H-F-200A, the YS, UTS and TE are 751 MPa, 1027 MPa and 28%, respectively, which are similar to that of the Ti6-H-F, indicating no precipitation strengthening under 200 oC annealing for just 1 hour. However， after 300 oC annealing for 1 hour, there is an observable improvement in the YS and UTS, resulting in 793 MPa and 1049 MPa, which are increased by 13.8% and 8.4%, respectively, compared to the Ti6-H-F. The Ti6-H-F-400A has the similar tensile performance to the case of the Ti6-H-F-300A. Further increasing the annealing temperature to 500 oC, the YS and UTS slightly increase to 801 MPa and 1097 MPa, respectively. It could be concluded that the precipitation strengthening in our current Ti6 alloys occurs after an annealing at 300 oC for 1 hour. The TEs of Ti6-H-F-200A, Ti6-H-F-300A, Ti6-H-F-400A and Ti6-H-F-500A are similar to that of the forged Ti6-H-F. It’s interesting to point out that even after 1000 oC annealing the TE of the Ti6-H-F-1000A is 24% which is slightly smaller 9

than that of the forged one. TEM analysis gives detailed information on the evolution of microstructures and precipitates in the forged and annealed Ti6-H alloys as shown in Fig. 6. In the Ti6-H, the recrystallized large grains with well-defined grain boundaries and almost no dislocations were observed as shown in Fig. 6a (left panel). The corresponding selected area electron diffraction (SAED) pattern with zone axis [011] in Fig. 6a (right panel) shows a typical fcc structure as indicated by the diffraction spots marked by the red circles, which is consistent with the above XRD results. After the 3-step forging, a high density of dislocations and dislocation tangles were introduced into the matrix as shown in Fig. 6b (left panel). The dislocation structure has been characterized by a configuration of planar glide along two non-coplanar slip systems, forming dislocation walls with a spacing distance between 400 and 800 nm. The high density of dislocations, dislocation tangles and intersecting dislocation walls play the dominant role in strengthening the Ti6-H-F alloy. The related SAED pattern in Fig. 6b (right panel) reveals a still single fcc structure in the Ti6-H-F. It is comprehensible that no precipitates were formed after the high temperature forging at 1100 oC because a subsequent water quenching was applied after the short-time high-speed hot forging. This rapid cooling cannot offer enough time for atomic diffusion and nucleation to form the precipitates. Similarly, it happens on the 1000 oC annealed and water quenched Ti6-H-F-1000WQ specimen that recovered and recrystallized large grains without any precipitates were observed as shown in Fig. 6c. While in the case of the furnace cooling (after the 1000 oC annealing), there are no dislocation structures but a high density of cuboidal nano precipitates in the Ti6-H-F-1000A as shown in Fig. 6d (left panel). The corresponding SAED pattern taken from the matrix and nano precipitates with zone axis [011] reveals not only the fcc matrix but also a newly formed crystal lattice with the L12 ordered structure as indicated by the 10

extra diffraction spots in the SAED pattern marked by the green circles in Fig. 6d (right panel). High-angle annular dark-field scanning transmission electron microscopy energy dispersive X-ray spectrometry (HAADF-STEM-EDS) was also conducted to qualitatively verify the elemental distribution in these precipitates. The EDS result (Fig. 7a and 7b) from one point on a precipitate particle shows that it contains 19at.% Ti, 61.9at.% Ni, 6.3at.% Fe and 13.8at.% Co, respectively, suggesting an atomic ratio of Ni : Ti = 3 : 1. Together with the SAED result, it could be concluded L12-type Ni3Ti precipitates, similar to the γ′ particles in Ni-based superalloys, were formed in the Ti6-H-F-1000A. The corresponding EDS elemental mapping images for different constituent elements are shown in Fig. 7c. These L12-type particles are apparently enriched in Ni and Ti, but depleted in Fe and Co. The size distributions of the precipitate particles shown in Figs. 8a and 8b indicate that the particles have an average size of 47 nm covering a range from 20 to 80 nm. The number density of these nano precipitates is about 3.9 × 10

21

/m3. To further study the interface structure between the

matrix and L12-type particles, the SAED pattern and high-resolution transmission electron microscopy (HRTEM, viewed along [011]) were employed, the results are shown in Figs. 8a and 8b. Fast Fourier transform (FFT) patterns from the matrix (see the SAED pattern highlighted by red frame in Fig. 8c inset) and the precipitate (see the SAED pattern highlighted by green frame in Fig. 8c inset) indicate that the L12-type particle exhibits a cube-on-cube orientation relationship with the fcc matrix. This orientation relationship has also been observed between the Ni3(Ti, Al) precipitate and the FeCoNiCr matrix in a (NiCoFeCr)94Ti2Al4 HEA alloy [42]. An enlarged image of the interface region between the precipitate and matrix (highlighted by the blue dotted square frame in Fig. 8c) is presented in Fig. 8d. It presents a perfect coherent interface between the precipitate and 11

matrix. The coherent interface could promote the precipitation strengthening by precipitate shearing mechanism [42]. So the super-high strength in the Ti6-H-F-1000A could be considered to be derived from the high density of coherent nano L12--type precipitates. However, due to its relative brittleness of the Ni3Ti phase, the presence of coarse (20 ~ 80 nm) nano precipitates would to some extent degrade the macroscopic plasticity of the alloy. This may be the main reason that the TE of Ti6-H-F-1000A is slightly lower than that of the Ti6-H-F even after 1000 oC annealing. However, considering the super-high YS (~ 893 MPa) and UTS (~ 1263 MPa), the simultaneously moderate TE of 24% is quite impressive. Fig. 9 presents the TEM results of the Ti6-H-F-200A, Ti6-H-F-300A, Ti6-H-F-400A and Ti6-H-F-500A alloys. After 200 oC annealing, a high density of dislocations, dislocation tangles and intersecting dislocation walls are remained in the Ti6-H-F-200A as shown in Fig. 9a (left panel), and there is no precipitate signal except for fcc matrix from the SAED pattern (Fig. 9a right panel). This result is consistent with the aforementioned tensile results that no precipitation strengthening appears after the 200 oC annealing. With further increasing the annealing temperature to 300 oC, 400 oC and 500 oC, the dislocation densities decrease, and the L12-type Ni3Ti precipitates are observed from the SAED patterns viewed through along [011] or [001] as shown in Fig. 9b-9d. It’s worth mentioning that it is hard to find the L12-type precipitates in the bright field (BF) TEM image, probably, due to the weak contrast and the too small size of the precipitates. Actually, nanocluster precipitates (~ 1 nm) can be seen in high-resolution TEM images of the Ti6-H-F-300A, Ti6-H-F-400A and Ti6-H-F-500A. To further indicate the L12-type Ni3Ti precipitation, a systematic XRD analysis were conducted. L12-type Ni3Ti precipitation is actually a process of diffusion of Ni and Ti solute atoms from the matrix to the precipitates. If the Ti dissolved in the FeCoNi matrix, the lattice parameter of the matrix 12

would be increased owing to the lattice expansion induced by the addition of Ti with the larger atomic radius. On the contrary, the Ti (and Ni) would subsequently be exsolved from the matrix to form the L12-type Ni3Ti precipitates because of the thermally-activated diffusion of Ti atoms, which would decrease the lattice parameter of the matrix. Fig. 10a shows comparison results of the normalized experimental XRD data of the FeCoNi-H, Ti6-H, Ti6-H-F-200A, Ti6-H-F-300A, Ti6-H-F-500A and Ti6-H-F-1000A. Intuitively, typical fcc phases of the matrix are indexed in all the specimens. While the diffraction peaks of the L12-type Ni3Ti phase are not discernible in the XRD patterns (even in Ti6-H-F-300A, Ti6-H-F-500A and Ti6-H-F-1000A specimens) because the peaks are overlapped with the primary fcc phases of the matrix. The peak positions of the indexed fcc phase in Ti6-H and Ti6-H-F-200A are similar and shift to relatively low angles in comparison with the FeCoNi-H as shown in Fig. 10a, indicating the dilated lattice in Ti6-H and Ti6-H-F-200A, as a result of Ti dissolution into the FeCoNi matrix. For example, the position of the (111) diffraction peaks shift from the 44.34o in FeCoNi-H to 43.41o in Ti6-H, as indicated by the magnified (111) diffraction peaks in Fig. 10b. While the peaks of the indexed fcc phase in Ti6-H-F-300A, Ti6-H-F-500A and Ti6-H-F-1000A shift to relatively high angles in comparison with the Ti6-H as shown in Fig. 10a and 10b. This peak shifting to relatively high angles indicates the constringent lattice of the matrix in both Ti6-H-F-300A, Ti6-H-F-500A and Ti6-H-F-1000A compared to Ti6-H, suggesting the exact precipitation of larger Ti (and/or Ni) atoms and further supporting the formation of L12-type Ni3Ti precipitates. 3.3 Microstructures and tensile properties of the homogenized + annealed (FeCoNi)94Ti6 alloy Fig. 11a shows the RT engineering stress - strain curves of the annealed Ti6-H MEA alloys without the pre-forging deformation, and the results of FeCoNi-H and Ti6-H are presented for 13

comparison. After 400 oC annealing and furnace cooling, the Ti6-H-400A has a high YS of 453 MPa and an UTS of 751 MPa, respectively, which are 56% and 27% higher than that of the Ti6-H. Specially, the TE of the Ti6-H-400A is as large as 49% which is slightly smaller than that of the Ti6-H (54%), as shown in Fig. 11b. For the Ti6-H-400A, the increase in strength is attributed to the precipitation strengthening from the precipitated L12-type Ni3Ti nanoclusters as indicated by the above TEM analysis and confirmed by the XRD result. For the pre-forged Ti6-H-F-400A, as mentioned above, it has an obvious increase in strength (YS ~ 750 MPa) but remarkable decrease in ductility (TE ~ 26%) compared to the Ti6-H. The additionally increased strength in Ti6-H-F-400A (compared to Ti6-H-400A) could be deduced from the pre-forging introduced high density of dislocations and dislocation walls which remain in the matrix after 400 oC annealing and produce significant dislocation strengthening. On the other hand, these excess pre-built dislocations tend to cause stress concentration, early necking and limited ductility. That’s to say, unlike the excess pre-built dislocations, precipitated L12-type Ni3Ti nanoclusters in the matrix would bring no disadvantageous effects on the RT ductility of Ti6-H-400A. However, further increasing annealing to 500 oC or 1000 oC, a perceptible decrease in ductility were observed in Ti6-H-500A (TE ~ 46%) and Ti6-H-1000A (TE ~ 21%), respectively. The decrease of ductility comes from the L12-type Ni3Ti precipitate coarsening from the nanoclusters to nano particles as indicated in the Ti6-H-F-1000A alloy. Five representative room-temperature tensile curves (Fig. 12a) of FeCoNi-H, Ti6-H, Ti6-H-400A, Ti6-H-F-400A and Ti6-H-F-1000A MEA alloys, as well as their corresponding UTS/TE values (Fig. 12b) and related microstructures (Fig. 12c) are summarized. Different mechanical properties of FeCoNi based MEA alloys are obtained by tuning the microstructures: i) the 14

homogenized FeCoNi equimolar alloy has high ductility (49%) but low strength (320 MPa), ii) solid solution strengthening induced by Ti addition and homogenizing treatments (> 1000 °C) results in moderate strength of 592 MPa with high ductility of 54%, iii) a combination of high strength (~ 751 MPa) and high ductility (49%) has been obtained in the Ti6-H-400A (without the pre-forging deformation) by annealing at moderate temperatures (i.e., 400 oC or 500 oC) because of the precipitation strengthening by the L12-type Ni3Ti nanoclusters, iv) a higher strength (~ 1060 MPa) but moderate ductility (26%) emerge after the post-deformation annealing at 400 oC in Ti6-H-F-400A due to the high density of pre-built dislocations and precipitated L12-type Ni3Ti nanoclusters, and v) an excellent combination of YS (~ 893 MPa), UTS (~ 1263 MPa) and ductility (TE ~ 24%) was achieved in the Ti6-H-F-1000A, resulting from the extensive nano-sized L12-type Ni3Ti precipitates. In this summarized chart, the room-temperature UTS and TE are used, and the UTS < 400 MPa is identified as “low strength”, 400 MPa < UTS < 700 MPa as “moderate strength”, 700 MPa < UTS < 1200 MPa as “high strength” and UTS > 1200 MPa as “super-high strength”, respectively. The TE > 40% is identified as “high ductility” and 20% < TE < 30% as “moderate ductility”. The data of some high performance HEA alloys such as the Co20Cr20Fe20Ni20Mo20 [7], (FeCoNiCr)93Al7 [43] (FeCoNiCr)93Mo7 [44], and (FeCoNiCr)94Al4Ti2 [14] are also presented for comparison as shown in Fig. 12b. 3.4 The strengthening mechanisms Fig. 13 presents the work-hardening behaviors (dσT/dε, where σT and ε are true stress and true strain, respectively) of the FeCoNi-H, Ti6-H, Ti6-H-F and Ti6-H-F-1000A alloys. All these sample exhibit a similar monotonic two-stage decrease in the strain hardening rate, and this two-stage strengthening commonly observed in alloys is related to dislocation configurations. At the first stage, 15

the sharp decrease in the strain hardening rate refers to a typical dislocation-strengthening feature described by the classical Taylor model [45]. Following this stage, the strain hardening originates from the formation of dislocation substructures, such as the Taylor lattice, high-density dislocation walls and dislocation cell [46, 47]. Intuitively, the lowest strain hardening rate has been observed in the FeCoNi-H, and it could be increased by solid solution strengthening, dislocation strengthening and precipitation strengthening, resulting in the highest strain hardening rate in the Ti6-H-F-1000A when the true strain is lower than 13% (before necking of Ti6-H-F-1000A). The main strengthening contributions are from solid solution strengthening, grain boundary strengthening, dislocation strengthening and precipitate strengthening. The yield strength (σy) can be estimated as the following equation [48]: σ y = σ 0 + σ A + ∆σ L12

(1)

Where σy is the yield stress, σ0 is the lattice friction stress (approximatively taken as 95 MPa, which is taken from the Al0.3FeCoCrNi HEA [43]), σA is the combined strengthening contributions from solid solution, grain boundaries and dislocations. ∆σL12 is the precipitation strengthening from the L12 precipitates. The Ti6-H-F is a single FCC structure. The grain boundary strengthening, solid solution strengthening and high density of dislocations play the dominant roles in strengthening mechanisms. The contribution of grain boundary strengthening from reducing the grain size is from the second term of the Hall-Petch relationship [48]: σ G = K y *d - 1 / 2

(2)

Where σG is grain boundary strengthening, Ky is the Hall-Petch coefficient and d is the average grain size (145 µm). The value Ky is approximatively taken as 824 MPa·µm1/2, which is also taken from 16

the Al0.3FeCoCrNi HEA [43]. As a result, the total contribution from the grain boundaries to the strength is approximatively taken as σG = 68 MPa. As the YS of the Ti6-H-F is 697 MPa, the remaining part of 534 MPa (697 MPa – 95 MPa – 68 MPa = 534 MPa) is intuitive from the strengthening contribution of the solid solution and dislocations [49]. In the case of Ti-free FeCoNi, grain boundary strengthening and dislocation strengthening play the dominant roles in strengthening mechanisms. The grain boundary strengthening is approximatively calculated as 96 MPa. As the FeCoNi-H-F exhibits a YS of 454  MPa, the corresponding dislocation strengthening is about 263 MPa (454 MPa– 95 MPa – 96 MPa = 263 MPa). The higher YS for the Ti6-H-F can be mainly attributed to the solid-solution strengthening effect of Ti. It can be roughly estimated that the solid solution strengthening in the Ti6-H-F is about 271 MPa (534 MPa – 263 MPa = 271 MPa). After 1000 oC annealing and furnace cooling, the Ti6-H-F-1000A exhibits a super-high YS and UTS, which is attributed to the formation of L12 precipitates in the matrix. The shearing mechanism is active for coherent precipitates [42]. For the shearing mechanism, the increase in yield strength results from contributions of coherency strengthening (∆σCS), modulus mismatch strengthening (∆σMS) and order strengthening (∆σOS) [50, 51]. The former two occurs before the dislocation shearing precipitates and the latter during shearing. Therefore, the total strength increment from shearing of precipitates is the larger of ∆σCS+∆σMS or ∆σOS [52, 53]. For the Ti6-H-F-1000A, there are almost no dislocation structures but a high density of cuboidal nano precipitates. Hence, the shearing mechanism for the contribution to yield strength increment is ∆σCS+∆σMS in the Ti6-H-F-1000A as described by the following equations [51]: ∆σ C S = M· α ε (G· ε) 3 / 2 (r· f/0.5G· b) 1 / 2

(3) 17

∆σ M S = M· 0.0055(∆G) 3 / 2 (2f/G) 1 / 2 (r/b) ( 3 / 2 m ) - 1

(4)

where M, G, and b are listed in Table 2, ε is the constrained lattice parameter mismatch, αε and m are the constant, ∆G is the modulus mismatch between the matrix and the precipitates, f is the volume fraction of the L12 precipitates (40.5%) and r is the mean radius of the precipitates (47 nm). The physical meaning and values of the symbols in the Eqs. (3) and (4) used are summarized in Table 2. The calculated values of ∆σCS and ∆σMS are 551 MPa and 57 MPa, respectively. Therefore, the contribution of L12 precipitation strengthening in Ti6-H-F-1000A is 551 + 57 = 608 MPa. The value of σ0 + σA in the Eq. (1) can be approximately taken as the YS (290 MPa) of Ti6-H-F-1000WQ. Consequently, the estimated YS of Ti6-H-F-1000A is 898 MPa, which is in reasonable consistent with the experimental values (893MPa). The excellent strength in Ti6-H-F-1000A is due to solid solution strengthening, grain boundary strengthening, and especially the precipitation strengthening from the coherent L12 precipitates. 4.

Conclusion A precipitation-strengthened (FeCoNi)94Ti6 (Ti6) MEA was fabricated. Microstructure

evolutions of dislocations and precipitates after forging deformation and annealing treatments were investigated by combining XRD and TEM techniques. In addition, the room temperature tensile properties of this Ti6 MEA have been successfully tailored by tuning the microstructures. Firstly, a moderate UTS value of 592 MPa with TE of 54% were obtained in the Ti6 MEA through solid solution strengthening by 6 at.% Ti dissolution into the FeCoNi equimolar alloy. Without the pre-forging deformation but with annealing at moderate temperatures (400 oC for 1h), the Ti6-H-400A achieved a combination of high strength (UTS ~ 751 MPa) and high ductility (TE ~ 49%) because of the precipitation strengthening from the L12-type Ni3Ti nanoclusters. These 18

nanocluster precipitates can effectively enhance the strength of the alloy without compromising its tensile ductility. With a post-deformation annealing at moderate temperatures, a higher strength (UTS ~ 1060 MPa) but moderate ductility (TE ~ 26%) emerged in Ti6-H-F-400A due to the effective deformation strengthening and precipitated coherent L12-type Ni3Ti nanoclusters. After a post-deformation annealing at 1000 oC, a super-high strength (UTS ~ 1263 MPa) and moderate ductility (TE ~ 24%) are achieved in the Ti6-H-F-1000A, resulting from the sufficient coherent nano-sized (20 ~ 80 nm) L12-type Ni3Ti precipitates. These coarse nano-precipitates can drastically increase the strength but to some extent degrade the tensile ductility. All these emphasized again microstructural optimization, such as the precipitate size and distribution, is a generic approach to tune the mechanical properties of precipitate strengthened MEAs and HEAs. Acknowledgement This work was financially supported by the National Natural Science Foundation of China (Grant Nos.: 51771184, 11735015, 11575241, 51801203, 11575231), the Natural Science Foundation of Anhui Province (Grant No. 1808085QE132), the Open Project of State Key Laboratory of Environment friendly Energy Materials (Grant No. 18KFHG02) and the Innovation Center of Nuclear Materials for National Defense Industry.

19

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Figure captions Fig. 1 Metallographic images of (a) the homogenized FeCoNi, (b) homogenized + forged FeCoNi, (c) homogenized + forged + 1000 oC annealed FeCoNi and (d) XRD patterns of these FeCoNi MEA alloys. Fig. 2 (a) Room-temperature tensile engineering stress-engineering strain curves of the FeCoNi alloys before and after forging and annealing for 1 h at different temperatures, and (b) the yield stress (YS), ultimate tensile strength (UTS) and total elongation (TE) of these FeCoNi alloys. Fig. 3 Metallographic image of (a) the homogenized (FeCoNi)94Ti6 and (b) XRD patterns of (FeCoNi)94Ti6 and FeCoNi alloys. Fig. 4 Metallographic images of (a) the homogenized + forged (FeCoNi)94Ti6 and post-deformation annealing at (b) 200 oC, (c) 500 oC and (d) 1000 oC. Fig. 5 (a) Room-temperature tensile engineering stress-engineering strain curves of the (FeCoNi)94Ti6 alloys before and after forging and post-deformation annealing at different temperatures, and (b) the yield stress (YS), ultimate tensile strength (UTS) and total elongation (TE) of these (FeCoNi)94Ti6 alloys. Fig. 6 The bright field TEM images and the corresponding SAED patterns of the specimens (a) Ti6-H, (b) Ti6-H-F, (c) Ti6-H-F-1000 WQ and (d) Ti6-H-F-1000A. Fig. 7 A HAADF-STEM image showing dispersion precipitates (a), EDS result showing the element content from one point on a precipitate (b) and EDS elemental mapping images for different constituent elements (c). Fig. 8 A bright field TEM image showing dispersion precipitates (a), a statistical result showing the corresponding particle size distribution (b), a HRTEM image and corresponding SAED patterns from the Ni3Ti precipitate and matrix (c), and an enlarged HRTEM image shown perfect coherent interfaces between the precipitate and matrix (d). Fig. 9 The bright field TEM images and the corresponding SAED patterns of the specimens (a) Ti6-H-F-200A, (b) Ti6-H-F-300A (c) Ti6-H-F-400A and (d) Ti6-H-F-500A. Fig. 10 (a) XRD patterns of FeCoNi-H, Ti6-H, Ti6-H-F-200A, Ti6-H-F-300A, Ti6-H-F-500A, Ti6-H-F-1000A alloys, and (b) the magnified 111 diffraction peaks. Fig. 11 (a) Room-temperature tensile engineering stress-engineering strain curves of the 24

(FeCoNi)94Ti6 alloys before and after annealing at different temperatures, and (b) the yield stress (YS), ultimate tensile strength (UTS) and total elongation (TE) of these (FeCoNi)94Ti6 alloys. Fig. 12 (a) Five representative room-temperature tensile curves of I: FeCoNi-H, II: Ti6-H, III: Ti6-H-400A, IV: Ti6-H-F-400A and V: Ti6-H-F-1000A MEA alloys, (b) their corresponding UTS/TE values, note that some high performance HEA alloys such as the Co20Cr20Fe20Ni20Mo20 [7], (FeCoNiCr)93Al7 [43] (FeCoNiCr)93Mo7 [44], and (FeCoNiCr)93Al4Ti2 [14] are also presented for comparison and (c) schematic diagrams showing the related microstructures. Fig. 13 Strain-hardening rate as a function of true strain.

25

Tables Table 1. Grain size, YS, UTS and TE of FeCoNi and (FeCoNi)94Ti6 MEA alloys. Grain size

YS

(µm)

(MPa) (MPa) (%)

Fcc matrix

400

103

320

49

FeCoNi-H-F

Fcc matrix

-

454

551

26

FeCoNi-H-F-1000A

Fcc matrix

-

155

461

40

Ti6-H

Fcc matrix

175

290

592

54

Ti6-H-400A

Fcc matrix + L12-type nanocluster

170

435

751

49

Ti6-H-500A

Fcc matrix + L12-type nanocluster

-

436

769

46

Ti6-H-1000A

Fcc matrix +

165

756

1054

21

Materials

Components

FeCoNi-H

UTS

TE

L12-type nanocluster Ti6-H-F

Fcc matrix

145

697

968

26

Ti6-H-F-200A

Fcc matrix

156

751

1027

28

Ti6-H-F-300A

Fcc matrix+ L12-type nanocluster

150

753

1049

27

Ti6-H-F-500A

Fcc matrix+ L12-type nanocluster

148

801

1097

26

Ti6-H-F-800A

-

-

895

1246

23

Ti6-H-F-1000A

Fcc matrix + L12-

140

893

1263

24

143

290

663

54

type nanoparticles Ti6-H-F-1000WQ

Fcc matrix

26

Table 2. Physical meaning and values of diffierent symbols used in the precipitation strengthening calculations [14, 43, 51, 54-56].

Symbol

Meaning

Values

Unit

G

Shear modulus

= 80 for FCC FeCoCrNiMn

GPa

∆G

Modulus mismatch

= 4 for Ni-based superalloys

GPa

between matrix and precipitates b

Burgers vector

= 0.255 for FCC FeCoCrNiMn

nm

ε

Constrained lattice

= 0.0017 for (FeCoCrNi)94Ti2Al4

Dimensionless

parameter misfit αε

Constant

= 2.6 for the FCC structure

Dimensionless

m

Constant

= 0.85

Dimensionless

M

Mean orientation

= 3.06 for the FCC polycrystalline

Dimensionless

factor

matrix

27

Fig. 1

28

Fig. 2

29

Fig. 3

30

Fig. 4

31

Fig. 5

32

Fig. 6

33

Fig. 7

34

Fig. 8

35

Fig. 9

36

Fig. 10

37

Fig. 11

38

Fig. 12

39

Fig. 13

40

The highlights are listed as following: 1.

Precipitation of coherent nanoclusters in a fcc-FeCoNi medium-entropy alloy (MEA) matrix was designed and realized.

2.

An excellent combination of yield strength (~ 893 MPa), ultimate tensile strength (~ 1263 MPa) and ductility (~ 24% elongation) has been achieved.

3.

This manuscript gives a detailed analysis of the relationships between the mechanical properties and microstructures.

Declaration of interest statement

I would like to declare on behalf of my co-authors that there is no conflict of interest in the submission of this manuscript, and manuscript is approved by all authors for publication.