In situ neutron diffraction study of a new type of stress-induced confined martensitic transformation in Fe22Co20Ni19Cr20Mn12Al7 high-entropy alloy

Materials Science & Engineering A 771 (2020) 138555

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In situ neutron diffraction study of a new type of stress-induced confined martensitic transformation in Fe22Co20Ni19Cr20Mn12Al7 high-entropy alloy Yajuan Shi a, Shilei Li a, Tung Lik Lee b, Xidong Hui a, Zhewei Zhang a, Runguang Li a, Minghe Zhang a, Saurabh Kabra b, Yan-Dong Wang a, * a

Beijing Advanced Innovation Center for Materials Genome Engineering, State Key Laboratory for Advanced Metals and Materials, University of Science and Technology Beijing, Beijing, 100083, China ISIS Facility, Rutherford Appleton Laboratory, Didcot, Oxfordshire, OX11 0QX, UK

b

A R T I C L E I N F O

A B S T R A C T

Keywords: Neutron scattering Stress High-entropy alloy Martensitic transformation Lattice strain

We successfully prepared an Fe22Co20Ni19Cr20Mn12Al7 alloy consisting of a face-center cubic (fcc) phase and body-center cubic (bcc) phase that exhibits an outstanding combination of true strength of 1430 MPa and ductility of 19.9% at room temperature (RT). The micromechanical behavior at RT and 77 K for the studied alloy during tensile deformation was investigated using in situ time-of-flight (TOF) neutron diffraction in combination with synchrotron-based high-energy X-ray diffraction (HE-XRD). Here, the striking finding of a large elastic strain of 7.0% and 5.6% is reported for the {200} bcc crystal plane, which was achieved at RT and 77 K, respectively. Such a large lattice distortion observed in the bcc phase was attributed to a new type of stressinduced confined martensitic transformation. We attributed the physical origin of this specific martensitic transformation to the intrinsic microstructural feature of a nano-scale continuous distribution of an ordered-todisordered crystal structure within the bcc phase, i.e., the disordered A2-structure was distributed continuously in the ordered B2-structure matrix. The stress-induced martensitic transformation from the metastable nanosized disordered A2-phase was confined by the stable B2-ordered matrix. The new findings in this study pro­ vide additional understanding of the deformation mechanisms of high-entropy alloys and insights into alloy design for further enhancement of the mechanical properties of high-performance structural materials used at cryogenic temperatures.

1. Introduction High-entropy alloys (HEAs) have demonstrated a high configuration entropy, sluggish atomic diffusion, the cocktail effect and special large lattice-distortion intrinsic characteristics [1,2]. Thus, they are being increasingly studied due to their excellent specific strength, exceptional ductility, superior mechanical performance at high temperatures, and outstanding cryogenic mechanical behavior [3–7]. In fact, a large number of HEAs can be designed via the so-called high configuration-entropy principle that tends to form either simple single-phase structures, such as face-centered cubic (fcc), body-centered cubic (bcc), and hexagonal close-packed (hcp), or a complex dual/multiple-phase structure, depending on the alloy composition, temperature and pressure parameters. However, only a tiny percent of HEAs may exist in the form of a single-phase solid solution, as the alloys consisting of multiple elements could be easily formed into intermetallic

phases if the phase equilibrium principle is considered. It can also be seen that some fcc single-phase alloys display a high ductility and poor strength, whereas the brittle nature of some bcc alloys with a high yield strength usually leads to difficulty during hot/cold processing; these factors limit their practical applications [8–11]. The deformation mechanisms generally found in HEAs include dislocation slip, mechanical twinning and phase transformation [12–14]. Gludovatz et al. [15] found that the fcc single-phase CoCrFe­ NiMn high entropy alloy displayed an exceptional damage tolerance characterized by a high tensile strength of 1280 MPa, ductility of 70%, and fracture toughness of 200 MPa� m1/2 at 77 K, which is superior to most metallic alloys. The deformation mechanism changed from planar-slip dislocation activity to mechanical nano-twinning, resulting in continuous strain hardening in the low temperature plastic flow regime that did not obey the well-known trade-off curves demonstrated by conventional alloys. It was suggested that the good mechanical

* Corresponding author. E-mail address: [email protected] (Y.-D. Wang). https://doi.org/10.1016/j.msea.2019.138555 Received 27 August 2019; Received in revised form 13 October 2019; Accepted 15 October 2019 Available online 15 October 2019 0921-5093/© 2019 Elsevier B.V. All rights reserved.

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properties in HEAs originate from not only the activation of mechanical twinning and/or stress-induced phase transformation due to low stacking-fault energy (SFE) but also the solid strengthening determined by the intrinsic high configuration entropy of chemical site occupation in the crystal lattice. The TaHfZrTi alloy [14], a typical single-phase bcc HEA, demonstrated irreversible phase transformation during unloading via tailoring the stability of the constituent phases, leading to a good combination of strength and ductility. Therefore, the strength-to-toughness balance could be achieved for dual-phase alloys having fcc and bcc phases, as the high plasticity in the fcc phase and high strength in the bcc phase could be effectively unified into an integrated microstructure. The dual- or multi-phase state for HEAs will thus be important in the alloy design for achieving excellent mechanical prop­ erties [16]. Great effort has been recently devoted to revealing the various changes in deformation mechanisms from the planar slip of dislocations to mechanical twinning to phase transformation in complex HEAs. The dual-phase Al0.6CoCrFeNi HEA [17] was investigated to observe the stress-induced martensitic transformation from the bcc phase to orthorhombic phase in the plastic deformation stage; the alloy exhibited a substantial yield strength of above 1 GPa with limited plastic deformation this behavior was caused by the existence of a hard and brittle matrix phase. Based on the specific atomic structure intrinsic to HEAs, two kinds of effects on the mechanical behavior can be summarized: (1) alloying with different sized elements in the unit cell may cause a large lattice distortion that enhances solution strengthening, and (2) the coexistence of various sublattices with different occupancies of element sites could give rise to a phase/twinning instability that enhances the plasticity in some chemically composed alloys. In fact, more severe distortions have been observed in HEAs than conventional alloys due to the uncertainty in atomic occupancies, which contributes to the excess configuration entropy in HEAs. The impact of intrinsic lattice distortion on solid so­ lution strengthening in a series of body-centered-cubic Al-containing HEAs has been experimentally investigated by Che et al. [18]. The first-principles calculations to investigate the lattice distortion and phase transition in chemically complex alloys were performed by Ye et al. [19]. Moreover, short-range order (SRO), which denotes a corre­ lation between the chemical occupancies of atomic sites, may exist in solid solution random alloys and has been found to affect the strength due to non-negligible solute/solute interactions under modified statis­ tics [20,21]. The SRO effect on strength has been demonstrated for or­ dered phase of NiAl in fcc solid solution alloys [22]. However, experimental-based investigations on the evolution of peak broad­ ening during loading and its related specific phase transformation behavior in HEAs have been rarely reported. The extremely challenging work facing to the multi-principal element alloy community is to develop approaches for quantifying peak broadening experimentally and to establish relationships among lattice strain, phase stability and mechanical properties. In situ neutron and high-energy X-ray diffraction (HE-XRD) tech­ niques have been used to quantitatively characterize deformation and phase transformation behaviors [23,24]. The above in situ techniques can provide an in-depth understanding of the excellent mechanical properties of HEAs from the microstructural aspects and deformation mechanisms. The in situ neutron diffraction experiments by Woo et al. [25] indicated that dislocation glide is the dominant deformation mode at 800 K and diffusion-controlled dislocation creep is dominant at 1000 K. The stress-induced martensitic transformation scenario from bcc to hcp phase for TaHfZrTi [14] and from bcc to bct for Al0.6CoCrFeNi HEAs [17] was traced by neutron diffraction and high-energy X-ray diffraction. Although many previous investigations on micromechanical behaviors of HEAs have been conducted, the quantitative relationship among lattice distortion, phase stability and mechanical properties, particularly on the specific intrinsic physical mechanisms of phase transformation, are still unclear. The CoCrFeNiAlx and CoCrFeNiMnAlx systems have been studied in recent years and shown transition from

single fcc to bcc phase with increasing the Al content [10,16,17]. With increasing the volume fraction of bcc phase, both the strength and hardness increased, while the alloy also became brittle. The reason of Fe22Co20Ni19Cr20Mn12Al7 alloy selected for this investigation is that a main fcc phase is retained as the matrix with a low volume fraction of nano-sized bcc phase with metastable state for achieving a good me­ chanical property in this HEA system. Due to the thermo-mechanical treatments, the order degree of nano-sized bcc phase may be easily tuned via the diffusion processing. Here, we prepared a hot-rolled and annealed non-equiatomic Fe22Co19Ni19Cr21Mn12Al7 (at.%) alloy that exhibits an outstanding combination of true strength of 1430 MPa and ductility of 19.9% at room temperature (RT). In situ time-of-flight (TOF) neutron diffraction and synchrotron-based HE-XRD techniques were employed to investigate the mechanical properties and deformation mechanisms of the studied alloy at RT and 77 K. A new type of stress-induced confined martensitic transformation was revealed that led to a large lattice distortion in the bcc phase of the studied alloy. Our results are expected to deepen the understanding of the deformation mechanisms and micromechanical behaviors in HEAs and be helpful in the design of new alloys. 2. Materials and experimental methods 2.1. Fabrication of materials An approximately 20 kg ingot of the Fe22Co19Ni19Cr21Mn12Al7 (at. %) HEA was fabricated in a vacuum induction furnace with high-purity (purity>99.9 wt%) constituent elements for five times and then hotforged to 50% thickness at 1373 K to ensure homogeneity and annealed at 1473 K for 2 h. Further grain refinement was achieved by hot-rolling to a 60% reduction in thickness and 1 h annealing at 1173 K in an Ar atmosphere followed by air cooling. The chemical composition of the HEAs was measured an energy dispersive spectroscope that was attached to the transmission electron microscope. The chemical composition of fcc and bcc phases is given in Table 1. It can be seen that the bcc phase contains much more Al and Ni and less Fe, Co and Cr than the fcc phase in matrix, and the Mn content remains equal in the two phases. 2.2. Neutron diffraction experiment The in situ time-of-flight (TOF) neutron diffraction measurements were performed during uniaxial tensile loading at room temperature and 77 K on a ENGIN-X diffractometer with a stress rig (ISIS spallation neutron source, the Rutherford Appleton Laboratory, UK). The experi­ mental setup is schematically shown in Fig. 1. The main advantage of time-of-flight (TOF) neutron diffraction is that the interplanar spacings for the {hkl} reflections in the fcc and bcc phases for the studied HEAs can be simultaneously measured with an elastic strain accuracy of 5 � 10 5. The neutron scattering gauge volume was 4 � 4 � 4 mm, which was defined by the 4 � 4 mm incident slit and the radial colli­ mators. The two detector banks at 2θ ¼ �90� (i.e., the axial detector and the radial detector) allowed simultaneous collection of diffraction pat­ terns with the scattering vector oriented parallel and perpendicular to the uniaxial tensile direction of the samples. The loading axis was par­ allel to the rolling direction and oriented 45� relative to the incident beam. The stress rig had a load range of �100 kN and was mounted on the diffractometer horizontally. The dog bone-shaped tensile rod sam­ ples were machined with a diameter of 8 mm and gauge length of 38 mm and with M12 screw-threaded ends. The samples were tested using a strain rate of 2 � 10 4 s 1. Stress control mode was used in the elastic region and was subsequently switched to displacement control after the sample yield point. The strain control was used instead because creep effects can occur in the material that is the macroscopic strain increases while the sample is being held at constant applied load above the yield point. Diffraction patterns were collected at each load step for an 2

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Table 1 Chemical compositions of fcc phase and bcc phase in the HEA (at.%). Phases

Chemical compositions (at.%)

fcc bcc

Fe 22.8 � 0.3 6.8 � 0.2

Co 20.3 � 0.3 12.9 � 0.3

Ni 16.1 � 0.2 36.4 � 0.4

Cr 21.5 � 0.2 3.2 � 0.1

Mn 11.6 � 0.2 11.6 � 0.3

Al 7.7 � 0.1 29.1 � 0.2

microscope (HRTEM) operated at 300 KV. 3. Results 3.1. Specific ordered-to-disordered coupling structure in bcc phase Fig. 2a–b show the bright-field TEM image and corresponding SAED patterns of the studied alloy. The phases having fcc and bcc-based structures can be identified well. The volume fractions of the fcc and bcc phases are 83% and 17%, respectively, which were determined by neutron diffraction and HE-XRD. The 100 superlattice diffraction spot can be observed in the SAED pattern from the selected region in the bcc phase, indicating that the bcc phase contained the ordered B2 phase. Interestingly, pairs of additional satellite-like spots distributed around 110, 200, 11ð Þ0 spots in Fig. 2b, which was attributed to a periodicity of anti-phase boundary (APB) domains in A2/B2 phase due to chemical or displacive deviations from the average atomic structure before deformation, being consistent with Ref. [27]. The 100 superlattice spot (circled by red in inset of Fig. 2c) was selected to take dark-field images (as shown in Fig. 2c), the random distribution of nano-sized B2 clus­ ters/domains (in ordered structure) along 110 and 1ð Þ10 direction within the bcc phase can be seen, which were both spatial and crystal­ lographic confined martensitic transformation. The confined martensitic transformation (CMT) is a new concept suggested by Khachaturyan et al. [28], describing that the martensitic transformation is totally or partially depressed during cooling due to the existence of various element defects (disordered occupation of elements or extra vacancy) in parent phase. It is possible to generate the stress- or strain-induced martensite during deformation, however, both transformation kinetics and crystallography in the CMT are different from that found in the traditional martensitic transformation (TMT). For the alloys having CMT, the stress-induced martensite is confined in both real sample space and reciprocal space to be the nanodomain-like morphology. This is due to that the nucleation of martensite dominates the whole transformation process, in contrast with a collective boundary motion between martensite and parent phase found in the TMT via domino reaction or avalanche transformation. The stress-induced CMT has been studied well in NiFeGaCo [29] and TiNbZrSn [30] shape memory alloys, but was rarely investigated in HEAs. The HRTEM image taken from the [001] zone shown in Fig. 2d, with the FFT patterns displayed in Fig. 2e–f, gives detailed information on the distribution of the disordered structure (confirmed a modulated feature) and B2-ordered structure in the bcc phase. A domain size of 2–5 nm is shown, which indicates that the disordered bcc phase (A2-structure) was distributed continuously in ordered B2-structure matrix and coherent with the B2 matrix. The modulated structure originating from the disordered bcc phase indicates that the disordered occupancy of constitutive elements was metastable in the studied HEA.

Fig. 1. Schematic illustration of the ENGIN-X in situ neutron diffraction measurement.

interval of ~30 min until the sample failed. The Rietveld method and the GSAS software package [26] were used to analyze the intensity and peak position for individual diffraction peaks. Lattice strain provides an important estimation for both macro­ scopic and intergranular stresses in the studied alloy during deforma­ tion, which can be calculated using equation (1):

εhkl ¼

dhkl d0hkl d0hkl

(1)

whereεhkl is the lattice strain of the {hkl} plane, dhkl andd0hkl are the interplanar spacings with and without an external load, respectively. 2.3. High-energy X-ray diffraction experiment The in situ HE-XRD uniaxial tensile experiments were further per­ formed at RT and 180 K at beamline 11-ID-C at the Advanced Photon Source, Argonne National Laboratory. A monochromatic X-ray beam (0.4 � 0.4 mm) with a wavelength of 0.117250 Å was used and the diffraction patterns were collected by a two-dimensional (2D) large area detector. Almost the same method with neutron diffraction was used for analyzing the evolution of lattice strains. As the detector used in the HEXRD covered a large reciprocal space, the texture evolution and minute changes in the crystal structure can be directly traced.

3.2. Mechanical properties

2.4. Transmission electron microscopy

In our case, different hot-rolled samples were further annealed at four different temperatures (1173 K, 1223 K, 1273 K and 1373 K) for 1 h to tailor the mechanical performance of the HEAs. The stress-strain curves for the above heat-treated samples are plotted in Fig. 3a, which were tested at 298 K at a strain rate of 1 � 10 3 s 1 using a tensile testing machine (with an extensometer). All the samples showed significant strain-hardening behavior. Specifically, the samples annealed at 1173 K

Transmission electron microscopy (TEM) was performed to examine the microstructure of the studied HEAs before and after deformation. The TEM samples were prepared by ion milling. The crystal structures for the undeformed and deformed samples and the orientation rela­ tionship between the bcc and fcc phases were investigated using a FEI Tecnai G2 F30 field-emission high-resolution transmission electron 3

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Fig. 2. (a) TEM image and the corresponding SAED pattern (b) along the [001] zone of the bcc and B2 phases in Fe22Co20Ni19Cr20Mn12Al7 HEA prior to deformation. (c) High magnification dark-field TEM images of the B2 phase in the undeformed sample which were taken by 100 superlattice spot (circled by red in inset of c). (d) HRTEM image of the bcc phase in the unde­ formed sample, showing the coherent relation­ ship between the disordered bcc precipitates and the B2 matrix. The FFT patterns (e) correspond­ ing to the red solid square region and (f) corre­ sponding to the yellow solid square region for the HRTEM image. (For interpretation of the refer­ ences to colour in this figure legend, the reader is referred to the Web version of this article.)

Fig. 3. (a) Representative tensile true strain-stress curves for the Fe22Co19Ni19Cr21Mn12Al7 (at.%) HEAs at room temperature. (b) Ultimate engineering stress tensile strength comparison and total ductility of the present work with various advanced steels, alloys [33], FeNiMnAl system HEAs [16,31,32].

for 1 h exhibited outstanding mechanical properties, i.e., a yield strength (YS) of 750 MPa, ultimate tensile strength (UTS) of 1435 MPa and ductility of 19.9%. The samples annealed at 1373 K for 1 h also possessed an excellent combination of UTS (1129 MPa) and ductility (41.2%). Fig. 3b shows a direct comparison of the engineering UTS and total ductility among the studied HEAs and the magnesium alloys, aluminum

alloys, advanced steels, and other reported FeNiMnAl system HEAs [16, 31,32]. (The data for other advanced alloys in the figure were taken from Ref. [33].) Our newly developed HEAs outperformed the most advanced steels and HEAs in similar systems and seemed to overcome the so-called strength-ductility trade-off principle such that they exhibited an exceptional strength-ductility combination.

4

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3.3. Micromechanical behavior investigated by in situ diffraction

101bct martensite, as discussed in detail in upcoming sections. The neutron diffraction spectra of the studied HEA at 77 K (not shown here) show a similar trend as at RT. The micromechanical and lattice distortion behavior of the alloy can be understood from the lattice strain evolution obtained from the results of in situ TOF neutron diffraction and HE-XRD analyses. The response of lattice strains to the applied stress for different {hkl} planes in the fcc and bcc phase along the loading direction (LD) and transverse direction (TD) at RT and 77 K are shown in Fig. 5. The whole deformation process for the studied HEA can be divided into three stages according to the lattice strain evolution versus the applied stress/strain: (1) the elastic deformation stage (Zone I), (2) the stress transfer from fcc to bcc stage (Zone II) and (3) the stress-induced martensitic transformation stage (Zone III). In Zone I, all lattice strains for both fcc and bcc phases respond proportionately to the applied stress with different slopes until 600 MPa at RT (Fig. 5a) and 700 MPa at 77 K (Fig. 5b), respectively, which is due to the elastic anisotropy in the two phases. The slope of the applied stress vs. lattice strains for {211}, {210} and {200} reflections in the bcc phase is slightly lower than that of {111}, {220} and {311} reflections in the fcc phase, indicating that the bcc phase has a lower elastic modulus than the fcc phase. However, the difference in elastic modulus between the two phases is not large. It indicates that a minimal load transfer existed between the bcc and fcc phases in the elastic deformation. In Zone II, the linear relationship between the lattice strain and stress disappeared as the strength yield begins, the lattice strains exhibit nonlinear responses with a further increase of the stress from 600 MPa to 800 MPa at RT (as shown in Fig. 5a) and from 700 MPa to 1000 MPa at

Fig. 4a shows the macroscopic stress-strain curve of the Fe22Co19 Ni19Cr21Mn12Al7 (at.%) HEA annealed at 1173 K for 1 h and tested at RT and 77 K at a strain rate of 2 � 10 4 s 1 with the dog bone-shaped tensile sample shown in the inset of Fig. 4a. The black line represents real-time measurements in Fig. 4a, whereas the pink and red points indicate where the neutron diffraction patterns were taken. The trends are clearly visible due to the load relaxation during the 30 min holding period under RT and 77 K. When the testing temperature decreased from RT to 77 K, the YS and UTS increased from 600 and 1187 MPa to 700 and 1415 MPa, respectively, with ductility decreasing from 19% to 14%, respectively. The alloy had a higher work-hardening ability at 77 K than RT. The tendency of work hardening and true strain behaviour was similar at both 298 K and 77 K (inset of Fig. 4a). The slight increase in work hardening rate in the initial stage of plastic deformation at 77 K is due to the easy generation of confined martensitic transformation, which accommodates the plastic deformation in fcc phase. The final UTS in bcc phase, achieved with a low strain at 77 K, controls the fracture as observed in the neutron experiment. Fig. 4b–c show the neutron diffraction spectra of the HEA along the axial and radial directions at different stress levels: undeformed, strained to 9% then unloaded, and reloaded to fracture, as circled in Fig. 4a. The diffraction spectra indicate that the alloy had a duplex (fcc þ bcc) microstructure before deformation. The diffraction peaks from the bcc phase became weak during deformation and the 111fcc peak almost overlapped the 110bcc peak. After fracture, new diffraction peaks appeared, which were considered to come from the stress-induced

Fig. 4. (a) The true stress-strain curves of Fe22Co20Ni19Cr20Mn12Al7 HEA obtained from in situ neutron diffraction of real-time and time averaged measurements at RT and 77 K. The inset in (a) shows the outer appearance of the pre-deformed samples. (b) and (c) show the neutron diffraction spectra of the HEA from the axial and radial detectors at different stress/strain levels. 5

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Fig. 5. Neutron diffraction obtained lattice evolution of the Fe22Co20Ni19Cr20Mn12Al7 HEA with the (a) applied stress (d) and strain at 298 K and (b) applied stress 77 K. (c) HE-XRD obtained lattice strain evolution with the applied stress at 298 K.

77 K (as shown in Fig. 5b). All lattice strains of the fcc phase begin to decrease after yield point, which indicates that fcc phase yields prefer­ entially to the bcc phase due to a strain relaxation in the fcc phase. More stress applied to the fcc phase is further transferred to the bcc phase. The lattice strains of bcc were increased from 1% to 3% as seen from Fig. 5a, depending on grain orientation. This indicates that the confined phase transformation has already been triggered. The clear increase in all lattice strains for the bcc phase is due to the stress-induced modulus softening, which should be related to the specific phase transformation characteristics [34]. In Zone III, the strain curve of the bcc planes deviates from linearity to a greater extent compared to the previous two stages. All lattice strains of fcc phase are less than 0.008, while all maximum lattice strains of the bcc phase are greater than 0.015. The large elastic strains of 7.0 % and 5.6% are observed for {200} crystal plane in the bcc phase at RT and 77 K, respectively. To our knowledge, it is the first experimental evi­ dence that so large lattice strain was found in two-phase HEAs during deformation. During the stage, the stress in the bcc phase reached a critical level, the confined martensitic transformation was initiated in some grains of the bcc phase (having about 2%–5% superelastic strain), as indicated from the large lattice strain. The superelastic strain devel­ oped in bcc phase (as almost as an elastic inclusion), with an additional small plastic strain, provides a mismatch with the plastic strain devel­ oped in fcc phase. We believe that the plastic deformation still occurred in the bcc phase, the change of intensity vs. azimuth angle for (211)bcc (not shown here), indicating that the texture component [110]//LD was

developed, but the plastic strain in bcc should be much smaller, in comparison with that of the fcc phase. Furthermore, the lattice strain of the bcc phase was larger than that of the fcc phase, suggesting that the bcc phase was subjected to greater stress than the fcc phase during the latter deformation stage. This confirms bcc phase in the alloy had an excellent strengthening effect during the plastic deformation stage. The evolution of the lattice strain and microstructure observed during deformation was further studied using the HE-XRD method. The changes in lattice strains as a function of applied stress (Fig. 5c) obtained from HE-XRD agreed well with those observed via neutron diffraction at RT (Fig. 5a–d). 3.4. Fine structural evolution during plastic deformation The 2-D in situ HE-XRD diffraction images covering the entire 360� azimuthal range recorded for different planes in the alloy that deformed at different selected strains 0, 1%, 2%, 7%, 17%, 26% at 298 K are shown in Fig. 6a–f. Note that 90� and 180� correspond to the loading direction (LD) and transverse direction (TD) respectively. All diffraction peaks for the bcc phase gradually became weak with increasing stress and plastic strain along the LD. Interestingly, the two theta of bcc200 peak deviated most severely from the initial angle among all the bcc peaks, which is good consistent with the large lattice strain of the 200 peak (Fig. 5). The 1-D diffraction profiles for the 110bcc and 211bcc at 950 MPa were obtained by integrating the 2-D diffraction pattern along η from 6

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Fig. 6. Selected 2-D X-ray diffraction images along the full azimuthal angle (η ¼ 0–360� ) for the alloy at different strains of (a) 0, (b) 1%, (c) 2%, (d) 7%, (e) 17%, (f) 26% at 298 K in situ HE-XRD. Note that 90� and 180� correspond to the loading direction (LD) and transverse direction (TD) respectively.

0� to 90� over a range of �5� . The structural evolution characteristics were studied by tracing 1-D diffraction profiles of different reflections under different stress levels, as plotted in Fig. 7. It can be seen that when the stress was higher than 772 MPa, new diffraction peaks 112bct and 101bct appeared, which were initiated in the 211bcc and 110bcc along SD (we defined η ¼ 45� as SD) and LD respectively, as shown in Fig. 7a–b, indicating that the stress-induced martensitic transformation occurred in the plastic deformation stage. This trend is consistent with that observed in an Al0.6CoCrFeNi alloy [17]. However, a colossal lattice distortion was observed in the bcc phase for our studied alloy system,

while both uniaxial UTS and elongation simultaneously improved greatly. The intensity and width the 110bcc and 211bcc diffraction peaks became weak after 925 MPa with further increasing stress, which in­ dicates that the confined martensitic transformation almost completed before fracture. A loading-unloading-reloading tensile test with TOF neutron diffraction at RT and 77 K was performed to further study the revers­ ibility of the martensitic transformation. The neutron diffraction pat­ terns derived from axial (LD) and radial banks (TD) at different stress levels are presented in Fig. 8. In contrast to the results by HE-XRD, not

Fig. 7. 1-D diffraction profiles of (a) 211bcc along SD (Azimuth ¼ 45� ), (b) 110bcc along TD and during loading by HE-XRD. 7

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only the martensitic transformation can be observed in the tensile di­ rection but also the latter stage of the plastic deformation from the TD, as circled in Fig. 8a–d. The peaks of the 101bct martensite appeared adjacent to the 110bcc peaks for the samples loaded at stress of approximately 762 MPa and 1030 MPa at RT and 77 K respectively, as seen clearly from the both the TD and LD neutron diffraction patterns (Fig. 8a–d). It indicted that the confined martensitic transformation can be triggered at a much lower stress (approximately 762 MPa) for the sample deformed at RT. The reason that a relatively lower stress is required to induce the martensitic transformation at a higher tempera­ ture can be explained by the characteristics of stress-induced martensitic transformation. Interestingly, the martensite almost reverted at 77 K after unloading to 20 MPa, while the more stable martensite remained during unloading at RT. This suggests that the martensitic trans­ formation is irreversible in the bcc phase after the large lattice distortion is induced at RT.

martensites had a short-range structure, similar structure was reported in Ni2FeGa alloy [35]. The multiple possible selections from the high-symmetry bcc phase resulted in the formation of preferred martensitic domains/variants, which could be related to the applied uniaxial loading mode. Gerold et al. [36] has already demonstrated that the short-range order or shearable precipitates can promote alloy slip planarity significantly, especially for multi-principle element systems. Thus, the nano-scale continuous distribution of an ordered-to-disordered coupling crystal structure within the bcc phase should play a crucial role in improving the mechanical properties [37]. The short-range order is considered to accelerate the martensitic trans­ formation more easily than long-range order [38]. The confined martensitic variants with short-range order can be an effective carrier for dislocation accumulation under tensile loading to enhance strain hardening and promote the strengthening effect [39]. Furthermore, Fig. 10a confirmed the existence of a modulated martensitic structure generated after plastic deformation, with satellite reflections spots distributed in a dimensional reciprocal space in the bcc/B2 domains. The 100 superlattice spot (circled by red in inset of Fig. 10a) was selected to take dark-field images for the samples prior to and after deformation, respectively, for distinguishing the distribution of ordered and disordered phase. It is can be seen that the nano-sized metastable A2 phase (displayed with bright contrast) was randomly distributed in the B2 phase before deformation, while large amounts of bright lamellar 101bct martensites existed in the A2/B2 domains after deformation, as presented in inset of Fig. 10b and b respectively. The similar modulated martensites were reported in Ni52.1Mn24.8Ga23.1 alloy [40] and Ni–34Mn–16Al alloy [41]. Nano-scale mechanical twinning can be shown as an additional

3.5. Stress-induced confined martensitic phase transformation Fig. 9 shows the TEM images and corresponding SAED patterns of bcc phase in the deformed HEA sample, which were taken from the [011], [1ð Þ11] and [001] zone axes. All the SAED patterns (Fig. 9b–d) correspond to the middle region of the selected microstructure images (Fig. 9a is a typical representation of them). The satellite spots appeared around each of the main spots in Fig. 9b–d. We attributed these satellite spots to the stress-induced confined phase-transformed variants of 101bct martensite. The different quantity and position of the spots might be caused by different martensitic variants. The satellite reflections around the main diffraction spots indicate that the stress-induced

Fig. 8. In situ neutron diffraction patterns at different stress levels under tensile loading-unloading-reloading for the Fe22Co20Ni19Cr20Mn12Al7 HEA: (a) 298 K and (c) 77 K along TD; (b) 298 K and (d) 77 K along LD. 8

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Fig. 9. (a) TEM images and (b, c and d) the corresponding SAED patterns of the bcc phase in the deformed HEA sample taken from the [011], [1ð zone axes, respectively.

deformation mechanism in the fcc phase for the samples deformed at RT by the representative bright-field (BF) and dark-field (DF) TEM images shown in Fig. 10c-d. Fig. 10e is the SAED of nanotwins. The critical shear stress for activating mechanical twinning should be related to the deformation temperature. It has been already reported in a single-phase FeCoNiCrMn HEA with the fcc structure that deformation twinning was not activated until the plastic deformation stage that occurred close to fracture at RT [15]. In contrast with the deformation at RT, the strength at 77 K remained a high level in the single-phase fcc alloy, which was attributed to the earlier activation of deformation twinning, i.e., that the progressively developing new interfaces in the microstructure provide an additional work hardening. In our case, the lattice strains for 311fcc and 200fcc decreased at the stress beginning from 530 MPa (in zone II) then increased from the stress greater than 710 MPa, as shown in Fig. 5c. The abovementioned tendency suggests that the mechanical twinning could be active in some grains having a specific preferred orientation in the fcc phase, which have large Schmid factor favorable for plastic deformation.

σ ¼ Vfcc σEfcc þ Vbcc σEbcc

(3)

where Ehkl is the diffraction elastic modulus for the specific {hkl} plane in each phase; Vfcc and Vbcc are volume fraction of fcc and bcc phase, respectively. Here, the uniaxial stress state for each phase was assumed, meanwhile the interaction of stress along the transverse direction was neglected. Moreover, we selected the lattice strains along the LD of 311fcc and 211bcc for estimating the phase stress because that stress is less affected by the intergranular stress, as demonstrated by the crystal plastic simulations [42]. The experimental stress-strain curves at RT and 77 K are shown in Fig. 11a b, respectively, in comparison with that calculated from the individual stress for both fcc and bcc phases. It can be seen that the calculated macrostresses match well with the experimental values observed over the whole strain region, which indicates that the phase stress should be accurately estimated. With increasing the applied stress, the stress subjected to the bcc phase is obviously increased and the maximum applied stress excesses 3.0 GPa at a strain of 18% for the RT deformation and at a strain of 14% for the 77 K deformation. The workhardening rate (WHR), representing by the derivative of the stress with respect to the strain for both fcc and bcc phases, is shown as insets in Fig. 11, respectively. The WHR drops rapidly and then decreases slowly down to an almost constant value at higher stress. The notable difference between the WHRs for bcc phase and fcc phase can be clearly evidenced at both RT and 77 K, with an obvious higher WHR in the bcc phase. It should be noted that the WHR observed for bcc phase at RT is slightly higher than that at 77 K. The intense work hardening observed in the bcc

4. Discussion 4.1. Strengthening mechanisms in bcc phase The hkl-specific phase stress, σhkl, for individual phase and the macrostress, σ, can be approximately estimated from their relative lat­ tice strains, εhkl, by the following equation,

σ hkl ¼ Ehkl � εhkl

Þ11] and [001]

(2) 9

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Fig. 10. (a) SAED pattern on the bcc phase (inset of (a) is corresponding image) of deformed sample taken from the [011] zone axes. Inset of (b) and (b) are dark field of bcc domain taken by 100 superlattice spot before and after deformed. (c) bright-field image, (d) dark-field image, (e) the SADP of nanotwins on the fcc phase in deformed HEA.

phase with the combination of high strength and high ductility was attributed to its confined martensitic transformation. The slight decrease in stress for the bcc phase near fracture could indicate that the bcc phase easily fractures [43]. Our analysis of the stress partition between the bcc and fcc phases indicates that, although the bcc phase comprised approximately 17% volume fraction in the studied alloy, it provided approximately 510 MPa to the stress (subjected to approximately 44.8% of UTS). To further understand the physical origin of the high WHR in bcc phase leading to a very high tensile fracture stress of 3.0 GPa, the strengthening mechanisms are depicted in detail as follows. In general, the four strengthening mechanisms that contribute to the flow strength have been found in HEAs, which consist of grain-boundary strength­ ening, solid-solution strengthening, precipitation strengthening, and dislocation strengthening. Here, the disordered bcc phase and B2 matrix were considered as a whole phase in the alloy, as they were completely coherent with each other. Thus, the precipitation strengthening contri­ bution to the UTS can be neglected. Primarily, we used a simplistic UTS strength model [44] to estimate the fracture stress, σUTS, of the bcc phase, which can be expressed as shown in equation (4):

σ UTS ¼ σ gb þ σ ss þ σ ds þ σ tr

According to a dislocation density model [45], the evolution of dislo­ cation density during deformation can be followed by analyzing the change in the slope of the modified W-H plot by equation (5): � qffiffiffiffiffiffiffiffiffiffiffiffiffi� ffiffiffiffiffi � pffiffi (5) ΔK ffi 0:9 D þ πM 2 b2 2 ρ KC0:5 where K ¼ 2sinθ/λ; △K ¼ 2cosθ(△θ)/λ; △θ is the full width at half maximum (FWHM) that is calibrated by extracting the instrumental broadening from the sample via the equation FWHM ¼ 1

½ðFWHMmeasured Þ2 ðFWHMinstrumental Þ2 �2 ; θ and λ are the diffraction angle and the wavelength of neutron diffraction or X-rays, respectively; D and ρ are average grain size and dislocation density, respectively; b is the Burgers vector of 0.230 nm; M is the Taylor factor (2.9 for both fcc and bcc materials [46]); C is the dislocation contrast factor. In the alloy, the values for Chkl for screw type {111} dislocations were calculated using the ANIZC software [47], with elastic constants inputs of C11 ¼ 194 GPa, C12 ¼ 124 GPa, and C44 ¼ 143 GPa [48]. The slope, obtained by modi­ fied W-H plots, is proportional to the square root of dislocation densityρ, and the intercept is inversely proportional to the average grain size D. The analysis of the modified W-H plots is shown in Fig. 11c. The dislo­ cation density was estimated to be 2.64 � 1015 m 2 (Fig. 11d), which was used to calculate the dislocation strengthening (σds) using the equation:

(4)

where σgb is the grain boundary strengthening, σss is solid solution strengthening, and σds is dislocation strengthening. For the materials subjected to severe plastic deformation, the dislocation density can be roughly estimated through the Williamson-Hall (W-H) method.

σ ds ¼ M � α � G � b � ρ1=2

10

(6)

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Fig. 11. Response of microstress for 211bcc and 311fcc to the applied strain during the deformation process from neutron diffraction at (a) 298 K and (b) 77 K. (c) The modified W-H plot of the bcc domain and (d) the curve of dislocation density and strain.

precipitates were dispersed in the B2 matrix, the specific nano-scale lattice modulation of A2/B2 and martensitic domains formed in the bcc phase might play a crucial role in the martensitic transformation by enabling an exceptionally high tensile fracture stress of 3.0 GPa at RT and 77 K. It is suggested that the stress-induced martensitic trans­ formation in the disordered bcc phase is triggered just by a slight change in lattice parameters from the approximate cubic structure to tetragonal structure, which accompanies with the broken symmetry to generate the lattice modulation (or atomic shuffling) during the transformation. Due to the stress-induced martensitic transformation for the disordered bcc domains is confined in both real sample space and reciprocal space by the stable ordered B2 matrix, both transformation kinetics and crystal­ lography are different from that found in the traditional martensitic transformation. For the bcc phase having confined martensitic trans­ formation, the nucleation or rearrangement of martensitic domains dominates the whole transformation process, in contrast with a collec­ tive boundary motion between martensite and parent phase found in the traditional martensitic transformation via domino reaction or avalanche transformation. Fig. 12 shows the schematic diagram denoting the change in microstructural features from the as-received sample to deformed sample based on the TEM observations. It is suggested that the random distribution of disordered A2 and ordered B2 domains in the undeformed sample was transformed to the preferred arrangement of B2 and 101bct orientated martensitic domains after deformation. That is to say, the martensitic domains with tetragonal structure having small c/a ratio, similar to the tweed structure observed in precursor martensite, are distributed randomly along different orientations in the cubic parent phase for the virgin sample. Under unified stress field, the martensitic domains should be rearranged with preferred selection and further be deformed with increasing c/a ratio. The stress-induced confined

where M, b and ρ are the same as the parameters in equation (5); α is an empirical constant of 0.23 [49]; G is the shear modulus (�65 GPa) for the AlCoCrCuFeNi system [50]. Parameter σds was calculated to be 513 MPa by equation (6) for the strength that arose from the deformed dislocation densities. The strengthening contributions from grain boundary strengthening (σgb) and solid solution strengthening (σss) can be approximately estimated as σ0.2, where σ0:2 ¼ σ ss þ σgb . Based on the phase stress at 0.2% plastic strain calculated from the lattice strains as shown in Fig. 11, the values of σ0.2 were approximately 900 MPa at RT and 1100 MPa at 77 K, respectively. Thus, the strength that resulted from the martensitic transformation (σtr) can be estimated as 1587 MPa and 1387 MPa, which occupied 52.9% and 46.2% of the total strength (3.0 GPa) at RT and 77 K, respectively. This indicates that martensitic transformation strengthening provided a significant contribution to the UTS of the bcc phase in the Fe22Co20Ni19Cr20Mn12Al7 HEA, especially at RT. The martensitic transformation may be the reason for outperforming the strength-ductility trade-off. This phase transformation mechanism strategy was previously modified by titanium alloys [51]. 4.2. The chemical origin of confined martensitic transformation mechanism It is well known that the pure B2-type intermetallic compound is brittle and hard, which can improve the strength but sacrifice the ductility of alloys. In the studied alloy, the disordered bcc phase as a precursor with a relatively low symmetry existed in the B2 phase. A modulated 101bct martensite were observed in bcc phase after defor­ mation. B2 phase decomposed and modulated plate structure observed in AlxCoCrCuFeNi system HEA [52,53]. Large amounts of nano-sized bcc 11

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3) The large lattice strain in the bcc phase was attributed to a stressinduced martensitic transformation having a nano-scale continuous order-to-disorder distribution in the bcc phase, where the martensitic transformation initiated in the metastable disordered bcc phase and was confined by the stable B2-ordered matrix. The confined martensitic variants with short range order were observed after deformation, where the satellite spots distributed in a dimensional reciprocal space. 4) The in situ TOF neutron diffraction experiments on the tensile loading-unloading-reloading process demonstrate that the martens­ itic transformation was triggered at a higher stress (1030 MPa) at a lower temperature (77 K), in comparison with a lower stress (762 MPa) at RT required to generate martensitic transformation. The martensitic transformation was irreversible at RT, but almost reversible at 77 K. 5) An additional strength enhancement in the duplex alloy was attrib­ uted to the specific nano-scale lattice modulation of A2/B2 and oriented martensitic domains formed in the bcc phase, which had an exceptionally high tensile fracture stress of 3.0 GPa at RT and 77 K. Declaration of competing interest The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. Acknowledgements This work was supported by the National Key Research and Devel­ opment Program of China (Grant No. 2017YFA0403804), the National Natural Science Foundation of China (NSFC) (Grant No.s 51471032 and 51527801), the Fundamental Research Funds for the Central Univer­ sities (Grant No.s 06111020) and the State Key Laboratory for Advanced Metals and Materials (Grant No. 2016Z-19). The use of the Advanced Photon Source was supported by the US Department of Energy, Office of Science, Office of Basic Energy Sciences, under Contract No. DE-AC0206CH11357. The neutron beam time at ENGIN-X of ISIS Neutron Source, Rutherford Appleton Laboratory, UK under experiment no. RB1720247 is acknowledged.

Fig. 12. Schematic disorder-order unit cell models of the (a) undeformed sample (containing ordered B2 phase and disordered bcc phase) and (b) deformed sample (containing ordered B2 phase and 101bct martensite) in accordance with the observed TEM image in deformed sample.

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martensitic transformation should further increase the nano-scale het­ erogeneity in atomistic distortion between the unstable disordered bcc domains and stable B2 matrix, which contributes to an obvious increase in the flow stress or causes an intense work hardening rate in bcc phase.

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5. Conclusions In the present investigation, in situ (TOF) neutron diffraction and high-energy X-ray diffraction techniques were used to reveal the intrinsic origin of a confined martensitic transformation and its related mechanical properties in a dual-phase Fe22Co20Ni19Cr20Mn12Al7 HEA fabricated by hot-forging, hot-rolling and annealing. Based on the experimental results, stress analyses and the origin of the confined martensitic transformation, the following conclusions can be drawn: 1) The investigated alloy designed with multiple non-equivalent highentropy compositional dual-phase microstructures overcame the strength-ductility trade-off by means of a stress-induced confined martensitic transformation in the bcc solid solution for enhancing the work-hardening rate. 2) Striking large elastic strains of 7.0% and 5.6% were found for the 200bcc crystal plane by in situ neutron diffraction in the studied alloy during tensile deformation at 298 K and 77 K, respectively.

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