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Fine-tuning of mechanical properties in V10Cr15Mn5Fe35Co10Ni25 high-entropy alloy through high-pressure torsion and annealing
Journal Pre-proof Fine-tuning of mechanical properties in V10Cr15Mn5Fe35Co10Ni25 high-entropy alloy through high-pressure torsion and annealing Peyman Asghari-Rad, Praveen Sathiyamoorthi, Nhung Thi-Cam Nguyen, Jae Wung Bae, Hamed Shahmir, Hyoung Seop Kim PII:
S0921-5093(19)31390-5
DOI:
https://doi.org/10.1016/j.msea.2019.138604
Reference:
MSA 138604
To appear in:
Materials Science & Engineering A
Received Date: 29 August 2019 Revised Date:
25 October 2019
Accepted Date: 26 October 2019
Please cite this article as: P. Asghari-Rad, P. Sathiyamoorthi, N. Thi-Cam Nguyen, J.W. Bae, H. Shahmir, H.S. Kim, Fine-tuning of mechanical properties in V10Cr15Mn5Fe35Co10Ni25 high-entropy alloy through high-pressure torsion and annealing, Materials Science & Engineering A (2019), doi: https://doi.org/10.1016/j.msea.2019.138604. This is a PDF file of an article that has undergone enhancements after acceptance, such as the addition of a cover page and metadata, and formatting for readability, but it is not yet the definitive version of record. This version will undergo additional copyediting, typesetting and review before it is published in its final form, but we are providing this version to give early visibility of the article. Please note that, during the production process, errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain. © 2019 Published by Elsevier B.V.
Fine-tuning of mechanical properties in V10Cr15Mn5Fe35Co10Ni25 highentropy alloy through high-pressure torsion and annealing
Peyman Asghari-Rada,b, Praveen Sathiyamoorthia,b, Nhung Thi-Cam Nguyena,b, Jae Wung Baea,b, Hamed Shahmirc, Hyoung Seop Kima,b,d,* a
Department of Materials Science and Engineering, Pohang University of Science and
Technology (POSTECH), Pohang, 37673, South Korea b
Center for High Entropy Alloys, Pohang University of Science and Technology
(POSTECH), Pohang 37673, South Korea c
Department of Materials Science and Engineering, University of Sheffield, Sheffield, United
Kingdom d
Graduate Institute of Ferrous Technology (GIFT), Pohang University of Science and
Technology (POSTECH), Pohang 37673, South Korea *
Corresponding Author: [email protected]
Abstract A V10Cr15Mn5Fe35Co10Ni25 (at%) high-entropy alloy (HEA) was subjected to highpressure torsion (HPT) and subsequently annealed under different conditions to study its microstructural evolution and mechanical properties. The HPT-processed sample consisted of a nanocrystalline structure with an average grain size of 30 nm. Annealing at 600 °C for 2 min led to annealing induced hardening because of the formation of sigma phase, whereas the samples annealed at higher temperatures showed a monotonous decrease in hardness because of dislocation recovery and grain growth. The average number of geometrically necessary dislocations decreased with increase in the temperature and duration time, indicating recovery of HPT-induced dislocations. Tensile properties indicated that the yield strength and ductility are notably influenced by the microstructure after annealing. The post-HPT annealing at 700 °C for 10 min resulted in an outstanding synergy of high tensile strength (~1.54 GPa) and good elongation-to-failure (~11%).
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Keywords: Post-deformation annealing, High-entropy alloy, High-pressure torsion, Grain refinement, Strengthening mechanisms, Thermo-mechanical processing
1. Introduction Multi-principal element alloys (MPEAs) are established as advanced materials with a distinct alloy design strategy, which is based on the use of a large number of major elements to form a simple structure with remarkable properties [1-4]. Based on their configurational entropy, MPEAs are classified into medium entropy alloys (MEAs) and high entropy alloys (HEAs) [5-11]. Among the MPEAs, the Co-Fe-Cr-Mn-Ni alloy system with a single-phase structure has received widespread attention due to its remarkable mechanical properties at liquid nitrogen temperature [12-14]. However, most of the single-phase face-centered cubic (FCC) HEAs show a rather low yield strength at room temperature, which could limit their usage in structural applications [15]. Several strengthening strategies such as grain refinement [16], solid-solution strengthening [17, 18], and precipitation strengthening [19, 20] have been utilized to improve the yield strength of FCC HEAs. Among these strengthening mechanisms, severe plastic deformation (SPD) techniques such as high-pressure torsion (HPT) are effective for enhancing the mechanical properties because SPD leads to fine grains [21, 22]. Recently, we attempted
to
enhance
the
strength
of
single-phase
FCC-structure
HEA
(V10Cr15Mn5Fe35Co10Ni25 at%) by grain refinement through the HPT process and achieved a significant increase in its strength from ~230 MPa (un-deformed condition) to ~2 GPa after 5 turns of HPT processing [23]. However, the corresponding elongation-to-failure showed a noticeable reduction from ~48% in the un-deformed condition to ~6% after HPT processing. This has been a common trend in highly strained materials: the strength increases compared to the initial un-deformed condition, while the corresponding elongation-to-failure decreases drastically [24]. The post-deformation annealing has been effectively utilized in conventional alloys and MPEAs to fine-tune the tensile properties and improve the ductility of the deformed material [16, 25, 26]. A good synergy of yield strength (~830 MPa) and elongation-to-failure (~65%) has been demonstrated in an equiatomic HEA (CoCrFeMnNi) by post-HPT
2
annealing at 800 °C for 10 min [27]. Similarly, enhancement in hardness and tensile properties have been achieved in CoCrNi MEA through post-HPT annealing at 500 °C and 600
°C
[28].
Accordingly,
further
investigation
of
post-HPT
annealing
of
V10Cr15Mn5Fe35Co10Ni25 HEA is desirable. Thus, the aim of the present study is to obtain optimal tensile properties in V10Cr15Mn5Fe35Co10Ni25 HEA by post-HPT annealing under different annealing conditions. Consequently, the microstructural evolution during post-HPT annealing and its influence on tensile behavior were investigated.
2. Experimental A V10Cr15Mn5Fe35Co10Ni25 (at%) HEA was cast in a vacuum induction furnace using elements with purity greater than 99.9%. The ingot was homogenized at 1100 °C for 6 h under argon gas and subsequently quenched in water. The homogenized sample was coldrolled from 6.2 to 1.3 mm with a thickness reduction ratio of ~79%. Disk-shaped samples with a diameter of 10 mm were cut from the rolled specimen. The disc-shaped samples were recrystallized at 900 °C for 30 min under argon gas and then quenched in water. The HPT processing was carried out on the disc samples under quasi-constrained conditions at 25 °C with a pressure of 6 GPa and a rotation rate of 1 revolution per minute (rpm) through 5 turns. The 5 turn HPT-processed samples were annealed in the temperature range 600 to 800 °C for a holding time from 2 to 60 min. X-ray diffraction (XRD) analysis was conducted on the post-HPT annealed samples to identify their constituent phases. Rigaku D/MAX-2500 XRD equipment with an incident beam of Cu Kα radiation was used. All scans were carried out with 2θ from 30 to 100°, a scan speed of 1°/min, and a step size of 0.02°. Electron backscatter diffraction (EBSD) analysis was conducted using a field emission scanning electron microscope (FE-SEM) with a step size of 50 nm (Model: Quanta 3D FEG, FEI Company, USA). The EBSD samples were prepared by mechanical polishing of samples with SiC paper up to 1200 grit, followed by diamond (3 µm and 1 µm paste), and then finished with colloidal silica polishing (0.04 µm). Orientation imaging microscopy (OIM) analysis software (TSL OIM analysis 7) was used to analyze the EBSD data. In addition, the average grain sizes of the annealed samples
3
were estimated using EBSD data by considering high angle grain boundaries (HAGBs) without considering annealing twin boundaries (Σ3, TBs). Electron-dispersive spectroscopy (EDS) analysis was performed using an FE-SEM (Model: JSM-7800F PRIME, JEOL Ltd., Japan). Transmission electron microscope (TEM) analysis (model: JEM-2100F, JEOL, Japan) was carried out on the HPT-processed sample using an acceleration voltage of 200 kV. The TEM samples were prepared by mechanical grinding to the thickness of ~60 µm and then thinning using electro-polishing with a solution of 10% HClO4 and 90% CH3COOH at an applied voltage of 25 V at room temperature. Vickers microhardness tests were performed on samples after post-HPT annealing with a load of 500 g and a dwelling time of 10 s. Tensile tests were carried out on the platetype miniature tensile specimens (gauge length: 1.5 mm, gauge width: 1 mm, thickness: 0.7 mm) with a quasi-static strain rate of 1 10−3 s−1 at room temperature (Instron 1361, Instron Crop., USA). A digital image correlation (DIC) method using optical 3D deformation analysis with white and black speckles on the surface of the tensile samples was conducted to measure the tensile strain (ARAMIS v6.1, GOM Optical Tech., Germany). The tensile specimens were cut from the 10 mm disc such that the position of the gage was located 2.5 mm from the center.
3. Results 3.1. Initial microstructure and microstructural evolution during HPT The EBSD inverse pole figure (IPF) map of the V10Cr15Mn5Fe35Co10Ni25 HEA before-HPT processing is illustrated in Fig. 1a, where the annealing twin boundaries (Σ3, TBs) and high-angle grain boundaries (HAGBs) are delineated by red and black colors, respectively. The microstructure of the undeformed sample (before-HPT) reveals equiaxed grains with an average size of ~7 µm. Moreover, a large fraction of annealing twins is apparent in the microstructure. The XRD patterns of V10Cr15Mn5Fe35Co10Ni25 HEA before and after 5 turns of HPT processing are illustrated in Fig. 1b. The XRD results indicate that both samples consisted of a single-phase FCC structure. Peak broadening was clear in the XRD pattern of the HPT-
4
processed sample because of the accumulated strain in the lattice and grain refinement during the HPT processing [23]. Figure 1c presents a bright-field TEM image of the HPT-processed sample after 5 turns. The deformed microstructure reveals intensive grain refinement with semi-equiaxed nano-structured grains having an average size of ~30 nm with vague grain boundaries. This microstructure is commonly observed in SPD processed materials. The vague grain boundaries indicate the presence of a high fraction of non-equilibrium boundaries [29]. In addition, the selected area electron diffraction (SAED) pattern taken along the [011] zone axis displays only the FCC phase with concentric ring patterns, indicating the nanocrystalline nature of the HPT-processed sample.
3.2. Microstructural evolution during post-HPT annealing The EBSD-IPF maps of V10Cr15Mn5Fe35Co10Ni25 HEA after post-HPT annealing at 600 to 900 °C and different holding times are illustrated in Fig. 2. The IPF map of the sample annealed at 600 °C for 10 min presents an annealed microstructure consisting of ultrafine grains with an average grain size of ~170 nm (Fig. 2a). The black area in this figure is related to a nano-structured matrix with very high density of accumulated strain, which is not indexed by the EBSD detector. An increase in the holding time from 10 to 60 min at 600 °C leads to grain growth, and the average grain size reaches ~370 nm (Fig. 2b). Furthermore, the results indicate a monotonous grain growth during annealing at 700 °C with an increase in annealing time. After annealing at 800 °C for 60 min, the average grain size increases drastically and reaches ~1.4 µm. The strain distribution based on the dislocation density can be studied using geometrically necessary dislocation (GND) by measuring the changes in the local orientation. The GNDs are generally induced at the interfaces to accommodate the strain compatibility between two grains during the plastic deformation. In the present study, the GND maps in Fig. 3 and GND density distribution curves in Fig. 4 were measured up to the third nearest neighbor with a maximum misorientation of 5°. The GND maps in Fig. 3 illustrate a higher strain distribution in the sample annealed at 600 °C for 10 min than the samples annealed under other conditions. The results demonstrate that
5
with increasing holding time (to 60 min at 600 °C), the average GND density decreases drastically to ~24.4 × 10
. By increasing the annealing temperature to 700 and 800 °C
at a similar holding time (60 min), the average GND density decreases to ~15.8× 10 ~10.5× 10
and
, respectively.
In Fig. 4, the GND density distribution curves of the post-HPT annealed samples show that the GND distribution is wide, and the maximum GND value is located at a higher dislocation density in the sample annealed at 600 °C for 10 min. As the annealing temperature and holding time increase, the distribution becomes sharp and narrow, and the maximum GND value shifts to lower dislocation density. To understand the microstructural evolution during post-HPT annealing, XRD analysis was conducted and the XRD patterns are presented in Fig. 5. Close inspection of the XRD results reveals the formation of new peaks with lower intensity as compared with the FCC peaks after annealing at temperatures in the range 600 to 800 °C. However, the microstructure is again a single FCC phase after annealing at 900 °C for 60 min. The additional peaks observed after treatment at the temperature range of 600 to 800 °C are attributed to the formation of sigma phase. Figure 6a illustrates the BSE image of the sample annealed at 800 °C for 60 min with corresponding EDS maps of its constituent elements. In the BSE image, it is apparent that the sigma phase precipitates at the triple junctions and the grain boundaries (indicated by the yellow arrow). The EDS maps confirm that the sigma phase precipitates are enriched in Cr and V, and depleted in Ni and Fe elements as compared to the matrix. In contrast, the elements Co and Mn show a homogeneous distribution in the matrix and precipitates. Similarly, V-Cr-rich sigma phase precipitation was also reported in V-containing HEAs like V20Cr15Fe20Ni45 and FeCrCoNiMnxVx [30, 31]. 3.3. Mechanical properties after post-HPT annealing Figure 7 shows the microhardness values after different processing conditions. The hardness value at the initial annealed condition was ~220 HV, and it reached a saturation value of ~505 HV after 5 turns of HPT processing. After post-HPT annealing at 600 °C for a short duration of 2 min, the hardness value increased initially and reached ~537 HV. Then, it decreased slightly with an increase in holding time and reached ~487 HV after 60 min.
6
However, by increasing the annealing temperature to 700 and 800 °C, the hardness evolution shows a monotonous trend of decrease. The engineering stress-strain plots and a summary of tensile properties after various processing conditions are illustrated in Fig. 8a and 8b, respectively. The before-HPT sample presents a remarkable elongation-to-failure of ~52% but with low yield and tensile strength. The tensile strength increased substantially to ~2 GPa after HPT processing; however, the elongation-to-failure declined to ~6%. The increase in the tensile strength with a decrease in elongation-to-failure is typical behavior in SPD-processed materials [24, 32]. Inspection of the tensile properties shows two different mechanical behaviors after post-HPT annealing. The samples annealed at 600 °C show a reduction in their ductility as compared with the HPT-processed sample. Whereas, the samples annealed at 700 and 800 °C present good ductility; however, this is at the cost of lower strength.
4. Discussion 4.1. Microstructural evolution during post-HPT annealing The GND maps illustrated in Fig. 3 represents a higher strain distribution in the sample annealed at 600 °C for 10 min as compared with the samples annealed at higher temperature and holding time. The high average GND density of ~47.3 × 10
implies
that the stored dislocation density during HPT processing is not fully recovered by annealing at 600 °C for 10 min. However, the reduced average GND density value in the samples annealed at 700 to 800 °C is attributed to the recovery of stored dislocations during the postHPT annealing. Furthermore, the GND density distribution curve of the sample annealed at 600 °C for 10 min demonstrates a wide GND distribution in which the maximum GND value is located at a higher dislocation density. Such GND distribution indicates that there is a large misorientation in the lattice due to high dislocation density generated during HPT processing. By increasing the annealing temperature and holding time, the GND density distribution becomes sharper and the maximum GND value shifts to lower dislocation density. This implies that the density of induced dislocations decreases during the post-HPT annealing, which leads to a decrease in misorientation in the lattice.
7
The XRD analysis reveals the formation of sigma phase after annealing at temperatures in the range 600–800 °C. The formation of sigma phase was previously reported in the HEAs processed by HPT followed by annealing. The sigma phase was formed in the HPT-processed Al0.5CoCrFeMnNi HEA after annealing at 800 °C for 60 min [16]. In addition, sigma phase was observed in CrFe2NiMnV0.25C0.125 HEA during annealing at 550 °C [26]. In general, the formation of sigma phase is substantially dependent on the heat treatment conditions, the density of lattice defects, and the chemical composition. In the present study, the formation of sigma phase was detected after a short-time annealing in the temperature range of 600–800 °C (Fig. 5). However, the sigma precipitates were observed after prolonged annealing at 700 °C for 1000 h in as-cast CrMnFeCoNi HEA with initial coarse-grained structure [33]. It is suggested that the slow kinetics of the sigma phase formation in equiatomic CrMnFeCoNi HEA is attributable to the sluggish diffusion in HEAs [34]. While Schuh et al. [35] observed that by subjecting CrMnFeCoNi HEA to the HPT process, the sigma phase can appear quickly after just a short time annealing. It is suggested that the fast formation of sigma phase is attributed to the high density of lattice defects which accelerates the diffusion process in the deformed material. Similarly, the sigma phase in the present HPT-processed HEA precipitates within a short time during the annealing process. In addition to the pre-deformation factor, the presence of certain elements in the alloying system also affects the kinetics of sigma phase formation. Thermodynamic calculations show that the addition of V to the Co-Fe-Cr-Mn-Ni alloy system increases the thermodynamic stability of the sigma phase [36]. Moreover, it was observed that the V element contributes to sigma phase formation and makes V-Cr-rich precipitates in V-containing HEAs [30, 31]. Hence, the addition of V can also influence the kinetics of sigma phase formation.
4.2. Mechanical properties after post-HPT annealing The hardness evolutions illustrated in Fig. 7 display that the hardness value initially increases in the sample annealed at 600 °C for 2 min and then decreases by increase in annealing time. Whereas, the samples annealed at 700 and 800 °C show a continuos decreasing trend. The increment of hardness during post-deformation annealing is observed in alloys due to precipitation of secondary phase, segregation, the formation of annealing
8
nano-twins, and grain boundary relaxation [27, 28, 37]. In CoCrMnFeNi HEA, it has been proposed that the formation of sigma phase and other nanostructured precipitates during annealing at an intermediate temperature lead to the hardness increment [34]. In CoCrNi MEA, annealing induced hardening has been observed in the absence of precipitation or secondary phase formation. Praveen et al. [28] suggested that the reduction of dislocation density due to the recovery (aside from grain boundary relaxation: the formation of equilibrium and well-defined grain boundaries) during annealing at 500 and 600 °C for 2 min, resulted in hardness increment in CoCrNi MEA. In the present study, the formation of sigma phase and grain boundary relaxation could be the main reasons for the hardness increment after annealing at 600 °C for 2 min. However, a detailed TEM analysis is required to confirm the hardening due to the grain boundary relaxation. Moreover, reduction in the hardness value by increasing the holding time and annealing temperature might be related to the softening due to coarsening of the recovered grains and decrease in the stored dislocation density (Figs. 2 and 3). The evolution of tensile properties shows good compliance with the microstructural evolution during post-HPT annealing. As mentioned in the previous section, by subjecting the HPT-processed sample to the annealing treatment, the sigma phase is formed in grain boundaries, leading to brittle mechanical properties [30]. On the other hand, the recovery of stored dislocations followed by grain growth during post-HPT annealing results in a softening effect that improves the ductility. The correlation between sigma phase formation and softening from recovery and grain growth controls the mechanical properties after post-HPT annealing. According to the XRD results presented in Fig. 5, the sigma phase could form very rapidly even after annealing at 600 °C for 2 min. However, the average GND values of samples annealed at 600 °C still present a high density of accumulated dislocations. Accordingly, the reduced ductility after annealing at 600 °C as compared with the HPTprocessed sample could be due to the formation of sigma phase. However, by increasing the annealing temperature to 700 and 800 °C, the GND density decreases and the grain size increases, which leads to an improvement in total elongation-to-failure. Furthermore, because the sigma phase particles are located in grain boundaries, they cannot directly contribute to strengthening by impeding the dislocations. However, these precipitates can reduce the rate of grain growth during annealing by the Zener pinning effect [38, 39].
9
It is important to note that after annealing at 700 °C for different holding times, only the total elongation-to-failure was increased and the uniform elongation was not changed noticeably (Fig. 8b). Similar mechanical behavior with low uniform elongation and high total elongation was reported earlier in conventional alloys and MPEAs [40, 41]. It has been proposed that the presence of pre-strain in the lattice results in high yield strength with no work hardening after yielding [42]. In the present study, the post-HPT annealed samples, which contained a large density of stored dislocations, exhibited similar mechanical behavior with pre-strained materials. Therefore, the samples annealed at 700 °C with high average GND density values, presented lower uniform elongation. However, the sample annealed at 800 °C showed an increment in uniform elongation as compared with the samples annealed at 700 °C. The increase in uniform elongation after annealing at 800 °C could be due to the reduction in the dislocation density, which improves the work hardening ability. As shown in Fig. 8b, the yield strength decreases gradually with increases in annealing time and temperature. The reduction in yield strength during the post-HPT annealing can be ascribed to the recovery of stored dislocations followed by grain growth. The effect of dislocation density on the yield strength can be described by the Bailey-Hirsch relationship, in which an increase in the dislocation density increases the yield strength [43]. Additionally, the correlation between the grain size and the yield strength can be depicted by the Hall-Petch relationship, where the reduction in the grain size improves the yield strength [44, 45]. Therefore, reduction of the dislocation density through the recovery process, in addition to increase in the grain size, causes the yield drop during post-HPT annealing. Figure 9 demonstrates the mechanical properties of V10Cr15Mn5Fe35Co10Ni25 HEA in comparison with conventional and multi-component alloys [27, 46]. It is apparent that the strength of the present HEA after post-HPT annealing is higher than that of high-strength conventional alloys. For instance, the sample annealed at 700 °C for 10 min presents a high tensile strength of 1.54 GPa with a total elongation-to-failure of ~11%. The outstanding mechanical properties obtained after post-HPT processing in the present study can be related to the contribution of different strengthening mechanisms. The presence of intensive dislocation density, ultrafine-grained matrix, and sigma-phase precipitations at grain boundaries that suppress grain growth, resulted in improvement of the tensile properties in post-HPT annealed V10Cr15Mn5Fe35Co10Ni25 HEA.
10
5. Conclusions In
the
current
work,
optimization
of
the
mechanical
properties
of
V10Cr15Mn5Fe35Co10Ni25 (at%) HEA was studied by HPT processing followed by annealing. The post-HPT annealing treatments were carried out under different annealing conditions that led to a variety of microstructures. A summary of the results from the present study follows. 1- The HPT-processed V10Cr15Mn5Fe35Co10Ni25 HEA exhibited a high tensile strength (~2 GPa) with ductility of ~6%. Nanoscale microstructure with an average grain size of ~30 nm was obtained after 5-turn HPT processing. 2- The average GND value of the post-HPT annealed samples presented a monotonous trend of decrease with increase in the annealing temperature and time due to recovery of HPT-induced dislocations and grain growth that occurred during subsequent annealing. 3- The samples annealed above 700 °C exhibited higher ductility as compared to the HPT-processed sample, but at the cost of losing strength. The optimal mechanical properties were achieved with annealing at 700 °C for 10 min, by which a remarkable synergy of the high tensile strength (~1.54 GPa) and good elongation-to-failure (~11%) was achieved.
Conflicts of interest No conflict of interest exists. We confirm that the article is original, and the article has not been published previously. The article has been written by the stated authors who are all aware of its content and approve its submission, and the article is not under consideration for publication elsewhere. All authors have approved the manuscript and agree with its submission to Materials Science and Engineering: A.
Acknowledgments We would like to acknowledge financial supports from POSCO TJ Park Foundation through POSCO Asia Fellowship, National Research Foundation of Korea (NRF) (2016M3D1A1023384),
and
Korea
Research
11
Fellowship
program
of
NRF
(2017H1D3A1A01013666). We would like to appreciate Prof. Alireza Zargaran for helping us with TEM analysis.
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Figures:
Figure 1. (a) EBSD inverse pole figure and grain boundary maps of V10Cr15Mn5Fe35Co10Ni25 (at%) HEA after annealing at 900 °C for 30 min, (b) XRD patterns before HPT and after 5 turn HPT processing, and (c) Bright-field TEM image and corresponding SAED pattern after 5 turn HPT processing.
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Figure 2. Electron backscatter diffraction (EBSD) inverse pole figure and boundary maps of post-deformation annealed samples. The color scale on the right side corresponds to the [001] inverse pole figure.
Figure 3. Geometrically necessary dislocation (GND) map of post-deformation annealed samples with the corresponding color scale on the right side.
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Figure 4. Geometrically necessary dislocation (GND) density distribution under different post-deformation annealing conditions.
Figure 5. X-ray diffraction (XRD) patterns of V10Cr15Mn5Fe35Co10Ni25 (at%) HEA after postdeformation annealing under different conditions.
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Figure 6. SEM-BSE image of V10Cr15Mn5Fe35Co10Ni25 (at%) HEA after post-deformation annealing at 800 °C for 60 min, and corresponding EDX mapping of its constituent elements.
Figure 7. Microhardness of V10Cr15Mn5Fe35Co10Ni25 (at%) HEA after different processing conditions.
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Figure 8. (a) Engineering stress-strain curves and (b) Evolution of tensile properties of V10Cr15Mn5Fe35Co10Ni25 (at%) HEA after different processing conditions.
Figure 9. Comparison of the tensile properties of HPT-processed and post-deformation annealed V10Cr15Mn5Fe35Co10Ni25 (at%) HEA with other high-strength conventional and multi-component alloys [27, 46] (This graph was adapted from reference [47]).
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Conflicts of interest No conflict of interest exists. We confirm that the article is original, and the article has not been published previously. The article has been written by the stated authors who are all aware of its content and approve its submission, and the article is not under consideration for publication elsewhere. All authors have approved the manuscript and agree with its submission to Materials Science and Engineering: A.
1
S0921-5093(19)31390-5
DOI:
https://doi.org/10.1016/j.msea.2019.138604
Reference:
MSA 138604
To appear in:
Materials Science & Engineering A
Received Date: 29 August 2019 Revised Date:
25 October 2019
Accepted Date: 26 October 2019
Please cite this article as: P. Asghari-Rad, P. Sathiyamoorthi, N. Thi-Cam Nguyen, J.W. Bae, H. Shahmir, H.S. Kim, Fine-tuning of mechanical properties in V10Cr15Mn5Fe35Co10Ni25 high-entropy alloy through high-pressure torsion and annealing, Materials Science & Engineering A (2019), doi: https://doi.org/10.1016/j.msea.2019.138604. This is a PDF file of an article that has undergone enhancements after acceptance, such as the addition of a cover page and metadata, and formatting for readability, but it is not yet the definitive version of record. This version will undergo additional copyediting, typesetting and review before it is published in its final form, but we are providing this version to give early visibility of the article. Please note that, during the production process, errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain. © 2019 Published by Elsevier B.V.
Fine-tuning of mechanical properties in V10Cr15Mn5Fe35Co10Ni25 highentropy alloy through high-pressure torsion and annealing
Peyman Asghari-Rada,b, Praveen Sathiyamoorthia,b, Nhung Thi-Cam Nguyena,b, Jae Wung Baea,b, Hamed Shahmirc, Hyoung Seop Kima,b,d,* a
Department of Materials Science and Engineering, Pohang University of Science and
Technology (POSTECH), Pohang, 37673, South Korea b
Center for High Entropy Alloys, Pohang University of Science and Technology
(POSTECH), Pohang 37673, South Korea c
Department of Materials Science and Engineering, University of Sheffield, Sheffield, United
Kingdom d
Graduate Institute of Ferrous Technology (GIFT), Pohang University of Science and
Technology (POSTECH), Pohang 37673, South Korea *
Corresponding Author: [email protected]
Abstract A V10Cr15Mn5Fe35Co10Ni25 (at%) high-entropy alloy (HEA) was subjected to highpressure torsion (HPT) and subsequently annealed under different conditions to study its microstructural evolution and mechanical properties. The HPT-processed sample consisted of a nanocrystalline structure with an average grain size of 30 nm. Annealing at 600 °C for 2 min led to annealing induced hardening because of the formation of sigma phase, whereas the samples annealed at higher temperatures showed a monotonous decrease in hardness because of dislocation recovery and grain growth. The average number of geometrically necessary dislocations decreased with increase in the temperature and duration time, indicating recovery of HPT-induced dislocations. Tensile properties indicated that the yield strength and ductility are notably influenced by the microstructure after annealing. The post-HPT annealing at 700 °C for 10 min resulted in an outstanding synergy of high tensile strength (~1.54 GPa) and good elongation-to-failure (~11%).
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Keywords: Post-deformation annealing, High-entropy alloy, High-pressure torsion, Grain refinement, Strengthening mechanisms, Thermo-mechanical processing
1. Introduction Multi-principal element alloys (MPEAs) are established as advanced materials with a distinct alloy design strategy, which is based on the use of a large number of major elements to form a simple structure with remarkable properties [1-4]. Based on their configurational entropy, MPEAs are classified into medium entropy alloys (MEAs) and high entropy alloys (HEAs) [5-11]. Among the MPEAs, the Co-Fe-Cr-Mn-Ni alloy system with a single-phase structure has received widespread attention due to its remarkable mechanical properties at liquid nitrogen temperature [12-14]. However, most of the single-phase face-centered cubic (FCC) HEAs show a rather low yield strength at room temperature, which could limit their usage in structural applications [15]. Several strengthening strategies such as grain refinement [16], solid-solution strengthening [17, 18], and precipitation strengthening [19, 20] have been utilized to improve the yield strength of FCC HEAs. Among these strengthening mechanisms, severe plastic deformation (SPD) techniques such as high-pressure torsion (HPT) are effective for enhancing the mechanical properties because SPD leads to fine grains [21, 22]. Recently, we attempted
to
enhance
the
strength
of
single-phase
FCC-structure
HEA
(V10Cr15Mn5Fe35Co10Ni25 at%) by grain refinement through the HPT process and achieved a significant increase in its strength from ~230 MPa (un-deformed condition) to ~2 GPa after 5 turns of HPT processing [23]. However, the corresponding elongation-to-failure showed a noticeable reduction from ~48% in the un-deformed condition to ~6% after HPT processing. This has been a common trend in highly strained materials: the strength increases compared to the initial un-deformed condition, while the corresponding elongation-to-failure decreases drastically [24]. The post-deformation annealing has been effectively utilized in conventional alloys and MPEAs to fine-tune the tensile properties and improve the ductility of the deformed material [16, 25, 26]. A good synergy of yield strength (~830 MPa) and elongation-to-failure (~65%) has been demonstrated in an equiatomic HEA (CoCrFeMnNi) by post-HPT
2
annealing at 800 °C for 10 min [27]. Similarly, enhancement in hardness and tensile properties have been achieved in CoCrNi MEA through post-HPT annealing at 500 °C and 600
°C
[28].
Accordingly,
further
investigation
of
post-HPT
annealing
of
V10Cr15Mn5Fe35Co10Ni25 HEA is desirable. Thus, the aim of the present study is to obtain optimal tensile properties in V10Cr15Mn5Fe35Co10Ni25 HEA by post-HPT annealing under different annealing conditions. Consequently, the microstructural evolution during post-HPT annealing and its influence on tensile behavior were investigated.
2. Experimental A V10Cr15Mn5Fe35Co10Ni25 (at%) HEA was cast in a vacuum induction furnace using elements with purity greater than 99.9%. The ingot was homogenized at 1100 °C for 6 h under argon gas and subsequently quenched in water. The homogenized sample was coldrolled from 6.2 to 1.3 mm with a thickness reduction ratio of ~79%. Disk-shaped samples with a diameter of 10 mm were cut from the rolled specimen. The disc-shaped samples were recrystallized at 900 °C for 30 min under argon gas and then quenched in water. The HPT processing was carried out on the disc samples under quasi-constrained conditions at 25 °C with a pressure of 6 GPa and a rotation rate of 1 revolution per minute (rpm) through 5 turns. The 5 turn HPT-processed samples were annealed in the temperature range 600 to 800 °C for a holding time from 2 to 60 min. X-ray diffraction (XRD) analysis was conducted on the post-HPT annealed samples to identify their constituent phases. Rigaku D/MAX-2500 XRD equipment with an incident beam of Cu Kα radiation was used. All scans were carried out with 2θ from 30 to 100°, a scan speed of 1°/min, and a step size of 0.02°. Electron backscatter diffraction (EBSD) analysis was conducted using a field emission scanning electron microscope (FE-SEM) with a step size of 50 nm (Model: Quanta 3D FEG, FEI Company, USA). The EBSD samples were prepared by mechanical polishing of samples with SiC paper up to 1200 grit, followed by diamond (3 µm and 1 µm paste), and then finished with colloidal silica polishing (0.04 µm). Orientation imaging microscopy (OIM) analysis software (TSL OIM analysis 7) was used to analyze the EBSD data. In addition, the average grain sizes of the annealed samples
3
were estimated using EBSD data by considering high angle grain boundaries (HAGBs) without considering annealing twin boundaries (Σ3, TBs). Electron-dispersive spectroscopy (EDS) analysis was performed using an FE-SEM (Model: JSM-7800F PRIME, JEOL Ltd., Japan). Transmission electron microscope (TEM) analysis (model: JEM-2100F, JEOL, Japan) was carried out on the HPT-processed sample using an acceleration voltage of 200 kV. The TEM samples were prepared by mechanical grinding to the thickness of ~60 µm and then thinning using electro-polishing with a solution of 10% HClO4 and 90% CH3COOH at an applied voltage of 25 V at room temperature. Vickers microhardness tests were performed on samples after post-HPT annealing with a load of 500 g and a dwelling time of 10 s. Tensile tests were carried out on the platetype miniature tensile specimens (gauge length: 1.5 mm, gauge width: 1 mm, thickness: 0.7 mm) with a quasi-static strain rate of 1 10−3 s−1 at room temperature (Instron 1361, Instron Crop., USA). A digital image correlation (DIC) method using optical 3D deformation analysis with white and black speckles on the surface of the tensile samples was conducted to measure the tensile strain (ARAMIS v6.1, GOM Optical Tech., Germany). The tensile specimens were cut from the 10 mm disc such that the position of the gage was located 2.5 mm from the center.
3. Results 3.1. Initial microstructure and microstructural evolution during HPT The EBSD inverse pole figure (IPF) map of the V10Cr15Mn5Fe35Co10Ni25 HEA before-HPT processing is illustrated in Fig. 1a, where the annealing twin boundaries (Σ3, TBs) and high-angle grain boundaries (HAGBs) are delineated by red and black colors, respectively. The microstructure of the undeformed sample (before-HPT) reveals equiaxed grains with an average size of ~7 µm. Moreover, a large fraction of annealing twins is apparent in the microstructure. The XRD patterns of V10Cr15Mn5Fe35Co10Ni25 HEA before and after 5 turns of HPT processing are illustrated in Fig. 1b. The XRD results indicate that both samples consisted of a single-phase FCC structure. Peak broadening was clear in the XRD pattern of the HPT-
4
processed sample because of the accumulated strain in the lattice and grain refinement during the HPT processing [23]. Figure 1c presents a bright-field TEM image of the HPT-processed sample after 5 turns. The deformed microstructure reveals intensive grain refinement with semi-equiaxed nano-structured grains having an average size of ~30 nm with vague grain boundaries. This microstructure is commonly observed in SPD processed materials. The vague grain boundaries indicate the presence of a high fraction of non-equilibrium boundaries [29]. In addition, the selected area electron diffraction (SAED) pattern taken along the [011] zone axis displays only the FCC phase with concentric ring patterns, indicating the nanocrystalline nature of the HPT-processed sample.
3.2. Microstructural evolution during post-HPT annealing The EBSD-IPF maps of V10Cr15Mn5Fe35Co10Ni25 HEA after post-HPT annealing at 600 to 900 °C and different holding times are illustrated in Fig. 2. The IPF map of the sample annealed at 600 °C for 10 min presents an annealed microstructure consisting of ultrafine grains with an average grain size of ~170 nm (Fig. 2a). The black area in this figure is related to a nano-structured matrix with very high density of accumulated strain, which is not indexed by the EBSD detector. An increase in the holding time from 10 to 60 min at 600 °C leads to grain growth, and the average grain size reaches ~370 nm (Fig. 2b). Furthermore, the results indicate a monotonous grain growth during annealing at 700 °C with an increase in annealing time. After annealing at 800 °C for 60 min, the average grain size increases drastically and reaches ~1.4 µm. The strain distribution based on the dislocation density can be studied using geometrically necessary dislocation (GND) by measuring the changes in the local orientation. The GNDs are generally induced at the interfaces to accommodate the strain compatibility between two grains during the plastic deformation. In the present study, the GND maps in Fig. 3 and GND density distribution curves in Fig. 4 were measured up to the third nearest neighbor with a maximum misorientation of 5°. The GND maps in Fig. 3 illustrate a higher strain distribution in the sample annealed at 600 °C for 10 min than the samples annealed under other conditions. The results demonstrate that
5
with increasing holding time (to 60 min at 600 °C), the average GND density decreases drastically to ~24.4 × 10
. By increasing the annealing temperature to 700 and 800 °C
at a similar holding time (60 min), the average GND density decreases to ~15.8× 10 ~10.5× 10
and
, respectively.
In Fig. 4, the GND density distribution curves of the post-HPT annealed samples show that the GND distribution is wide, and the maximum GND value is located at a higher dislocation density in the sample annealed at 600 °C for 10 min. As the annealing temperature and holding time increase, the distribution becomes sharp and narrow, and the maximum GND value shifts to lower dislocation density. To understand the microstructural evolution during post-HPT annealing, XRD analysis was conducted and the XRD patterns are presented in Fig. 5. Close inspection of the XRD results reveals the formation of new peaks with lower intensity as compared with the FCC peaks after annealing at temperatures in the range 600 to 800 °C. However, the microstructure is again a single FCC phase after annealing at 900 °C for 60 min. The additional peaks observed after treatment at the temperature range of 600 to 800 °C are attributed to the formation of sigma phase. Figure 6a illustrates the BSE image of the sample annealed at 800 °C for 60 min with corresponding EDS maps of its constituent elements. In the BSE image, it is apparent that the sigma phase precipitates at the triple junctions and the grain boundaries (indicated by the yellow arrow). The EDS maps confirm that the sigma phase precipitates are enriched in Cr and V, and depleted in Ni and Fe elements as compared to the matrix. In contrast, the elements Co and Mn show a homogeneous distribution in the matrix and precipitates. Similarly, V-Cr-rich sigma phase precipitation was also reported in V-containing HEAs like V20Cr15Fe20Ni45 and FeCrCoNiMnxVx [30, 31]. 3.3. Mechanical properties after post-HPT annealing Figure 7 shows the microhardness values after different processing conditions. The hardness value at the initial annealed condition was ~220 HV, and it reached a saturation value of ~505 HV after 5 turns of HPT processing. After post-HPT annealing at 600 °C for a short duration of 2 min, the hardness value increased initially and reached ~537 HV. Then, it decreased slightly with an increase in holding time and reached ~487 HV after 60 min.
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However, by increasing the annealing temperature to 700 and 800 °C, the hardness evolution shows a monotonous trend of decrease. The engineering stress-strain plots and a summary of tensile properties after various processing conditions are illustrated in Fig. 8a and 8b, respectively. The before-HPT sample presents a remarkable elongation-to-failure of ~52% but with low yield and tensile strength. The tensile strength increased substantially to ~2 GPa after HPT processing; however, the elongation-to-failure declined to ~6%. The increase in the tensile strength with a decrease in elongation-to-failure is typical behavior in SPD-processed materials [24, 32]. Inspection of the tensile properties shows two different mechanical behaviors after post-HPT annealing. The samples annealed at 600 °C show a reduction in their ductility as compared with the HPT-processed sample. Whereas, the samples annealed at 700 and 800 °C present good ductility; however, this is at the cost of lower strength.
4. Discussion 4.1. Microstructural evolution during post-HPT annealing The GND maps illustrated in Fig. 3 represents a higher strain distribution in the sample annealed at 600 °C for 10 min as compared with the samples annealed at higher temperature and holding time. The high average GND density of ~47.3 × 10
implies
that the stored dislocation density during HPT processing is not fully recovered by annealing at 600 °C for 10 min. However, the reduced average GND density value in the samples annealed at 700 to 800 °C is attributed to the recovery of stored dislocations during the postHPT annealing. Furthermore, the GND density distribution curve of the sample annealed at 600 °C for 10 min demonstrates a wide GND distribution in which the maximum GND value is located at a higher dislocation density. Such GND distribution indicates that there is a large misorientation in the lattice due to high dislocation density generated during HPT processing. By increasing the annealing temperature and holding time, the GND density distribution becomes sharper and the maximum GND value shifts to lower dislocation density. This implies that the density of induced dislocations decreases during the post-HPT annealing, which leads to a decrease in misorientation in the lattice.
7
The XRD analysis reveals the formation of sigma phase after annealing at temperatures in the range 600–800 °C. The formation of sigma phase was previously reported in the HEAs processed by HPT followed by annealing. The sigma phase was formed in the HPT-processed Al0.5CoCrFeMnNi HEA after annealing at 800 °C for 60 min [16]. In addition, sigma phase was observed in CrFe2NiMnV0.25C0.125 HEA during annealing at 550 °C [26]. In general, the formation of sigma phase is substantially dependent on the heat treatment conditions, the density of lattice defects, and the chemical composition. In the present study, the formation of sigma phase was detected after a short-time annealing in the temperature range of 600–800 °C (Fig. 5). However, the sigma precipitates were observed after prolonged annealing at 700 °C for 1000 h in as-cast CrMnFeCoNi HEA with initial coarse-grained structure [33]. It is suggested that the slow kinetics of the sigma phase formation in equiatomic CrMnFeCoNi HEA is attributable to the sluggish diffusion in HEAs [34]. While Schuh et al. [35] observed that by subjecting CrMnFeCoNi HEA to the HPT process, the sigma phase can appear quickly after just a short time annealing. It is suggested that the fast formation of sigma phase is attributed to the high density of lattice defects which accelerates the diffusion process in the deformed material. Similarly, the sigma phase in the present HPT-processed HEA precipitates within a short time during the annealing process. In addition to the pre-deformation factor, the presence of certain elements in the alloying system also affects the kinetics of sigma phase formation. Thermodynamic calculations show that the addition of V to the Co-Fe-Cr-Mn-Ni alloy system increases the thermodynamic stability of the sigma phase [36]. Moreover, it was observed that the V element contributes to sigma phase formation and makes V-Cr-rich precipitates in V-containing HEAs [30, 31]. Hence, the addition of V can also influence the kinetics of sigma phase formation.
4.2. Mechanical properties after post-HPT annealing The hardness evolutions illustrated in Fig. 7 display that the hardness value initially increases in the sample annealed at 600 °C for 2 min and then decreases by increase in annealing time. Whereas, the samples annealed at 700 and 800 °C show a continuos decreasing trend. The increment of hardness during post-deformation annealing is observed in alloys due to precipitation of secondary phase, segregation, the formation of annealing
8
nano-twins, and grain boundary relaxation [27, 28, 37]. In CoCrMnFeNi HEA, it has been proposed that the formation of sigma phase and other nanostructured precipitates during annealing at an intermediate temperature lead to the hardness increment [34]. In CoCrNi MEA, annealing induced hardening has been observed in the absence of precipitation or secondary phase formation. Praveen et al. [28] suggested that the reduction of dislocation density due to the recovery (aside from grain boundary relaxation: the formation of equilibrium and well-defined grain boundaries) during annealing at 500 and 600 °C for 2 min, resulted in hardness increment in CoCrNi MEA. In the present study, the formation of sigma phase and grain boundary relaxation could be the main reasons for the hardness increment after annealing at 600 °C for 2 min. However, a detailed TEM analysis is required to confirm the hardening due to the grain boundary relaxation. Moreover, reduction in the hardness value by increasing the holding time and annealing temperature might be related to the softening due to coarsening of the recovered grains and decrease in the stored dislocation density (Figs. 2 and 3). The evolution of tensile properties shows good compliance with the microstructural evolution during post-HPT annealing. As mentioned in the previous section, by subjecting the HPT-processed sample to the annealing treatment, the sigma phase is formed in grain boundaries, leading to brittle mechanical properties [30]. On the other hand, the recovery of stored dislocations followed by grain growth during post-HPT annealing results in a softening effect that improves the ductility. The correlation between sigma phase formation and softening from recovery and grain growth controls the mechanical properties after post-HPT annealing. According to the XRD results presented in Fig. 5, the sigma phase could form very rapidly even after annealing at 600 °C for 2 min. However, the average GND values of samples annealed at 600 °C still present a high density of accumulated dislocations. Accordingly, the reduced ductility after annealing at 600 °C as compared with the HPTprocessed sample could be due to the formation of sigma phase. However, by increasing the annealing temperature to 700 and 800 °C, the GND density decreases and the grain size increases, which leads to an improvement in total elongation-to-failure. Furthermore, because the sigma phase particles are located in grain boundaries, they cannot directly contribute to strengthening by impeding the dislocations. However, these precipitates can reduce the rate of grain growth during annealing by the Zener pinning effect [38, 39].
9
It is important to note that after annealing at 700 °C for different holding times, only the total elongation-to-failure was increased and the uniform elongation was not changed noticeably (Fig. 8b). Similar mechanical behavior with low uniform elongation and high total elongation was reported earlier in conventional alloys and MPEAs [40, 41]. It has been proposed that the presence of pre-strain in the lattice results in high yield strength with no work hardening after yielding [42]. In the present study, the post-HPT annealed samples, which contained a large density of stored dislocations, exhibited similar mechanical behavior with pre-strained materials. Therefore, the samples annealed at 700 °C with high average GND density values, presented lower uniform elongation. However, the sample annealed at 800 °C showed an increment in uniform elongation as compared with the samples annealed at 700 °C. The increase in uniform elongation after annealing at 800 °C could be due to the reduction in the dislocation density, which improves the work hardening ability. As shown in Fig. 8b, the yield strength decreases gradually with increases in annealing time and temperature. The reduction in yield strength during the post-HPT annealing can be ascribed to the recovery of stored dislocations followed by grain growth. The effect of dislocation density on the yield strength can be described by the Bailey-Hirsch relationship, in which an increase in the dislocation density increases the yield strength [43]. Additionally, the correlation between the grain size and the yield strength can be depicted by the Hall-Petch relationship, where the reduction in the grain size improves the yield strength [44, 45]. Therefore, reduction of the dislocation density through the recovery process, in addition to increase in the grain size, causes the yield drop during post-HPT annealing. Figure 9 demonstrates the mechanical properties of V10Cr15Mn5Fe35Co10Ni25 HEA in comparison with conventional and multi-component alloys [27, 46]. It is apparent that the strength of the present HEA after post-HPT annealing is higher than that of high-strength conventional alloys. For instance, the sample annealed at 700 °C for 10 min presents a high tensile strength of 1.54 GPa with a total elongation-to-failure of ~11%. The outstanding mechanical properties obtained after post-HPT processing in the present study can be related to the contribution of different strengthening mechanisms. The presence of intensive dislocation density, ultrafine-grained matrix, and sigma-phase precipitations at grain boundaries that suppress grain growth, resulted in improvement of the tensile properties in post-HPT annealed V10Cr15Mn5Fe35Co10Ni25 HEA.
10
5. Conclusions In
the
current
work,
optimization
of
the
mechanical
properties
of
V10Cr15Mn5Fe35Co10Ni25 (at%) HEA was studied by HPT processing followed by annealing. The post-HPT annealing treatments were carried out under different annealing conditions that led to a variety of microstructures. A summary of the results from the present study follows. 1- The HPT-processed V10Cr15Mn5Fe35Co10Ni25 HEA exhibited a high tensile strength (~2 GPa) with ductility of ~6%. Nanoscale microstructure with an average grain size of ~30 nm was obtained after 5-turn HPT processing. 2- The average GND value of the post-HPT annealed samples presented a monotonous trend of decrease with increase in the annealing temperature and time due to recovery of HPT-induced dislocations and grain growth that occurred during subsequent annealing. 3- The samples annealed above 700 °C exhibited higher ductility as compared to the HPT-processed sample, but at the cost of losing strength. The optimal mechanical properties were achieved with annealing at 700 °C for 10 min, by which a remarkable synergy of the high tensile strength (~1.54 GPa) and good elongation-to-failure (~11%) was achieved.
Conflicts of interest No conflict of interest exists. We confirm that the article is original, and the article has not been published previously. The article has been written by the stated authors who are all aware of its content and approve its submission, and the article is not under consideration for publication elsewhere. All authors have approved the manuscript and agree with its submission to Materials Science and Engineering: A.
Acknowledgments We would like to acknowledge financial supports from POSCO TJ Park Foundation through POSCO Asia Fellowship, National Research Foundation of Korea (NRF) (2016M3D1A1023384),
and
Korea
Research
11
Fellowship
program
of
NRF
(2017H1D3A1A01013666). We would like to appreciate Prof. Alireza Zargaran for helping us with TEM analysis.
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Figures:
Figure 1. (a) EBSD inverse pole figure and grain boundary maps of V10Cr15Mn5Fe35Co10Ni25 (at%) HEA after annealing at 900 °C for 30 min, (b) XRD patterns before HPT and after 5 turn HPT processing, and (c) Bright-field TEM image and corresponding SAED pattern after 5 turn HPT processing.
15
Figure 2. Electron backscatter diffraction (EBSD) inverse pole figure and boundary maps of post-deformation annealed samples. The color scale on the right side corresponds to the [001] inverse pole figure.
Figure 3. Geometrically necessary dislocation (GND) map of post-deformation annealed samples with the corresponding color scale on the right side.
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Figure 4. Geometrically necessary dislocation (GND) density distribution under different post-deformation annealing conditions.
Figure 5. X-ray diffraction (XRD) patterns of V10Cr15Mn5Fe35Co10Ni25 (at%) HEA after postdeformation annealing under different conditions.
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Figure 6. SEM-BSE image of V10Cr15Mn5Fe35Co10Ni25 (at%) HEA after post-deformation annealing at 800 °C for 60 min, and corresponding EDX mapping of its constituent elements.
Figure 7. Microhardness of V10Cr15Mn5Fe35Co10Ni25 (at%) HEA after different processing conditions.
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Figure 8. (a) Engineering stress-strain curves and (b) Evolution of tensile properties of V10Cr15Mn5Fe35Co10Ni25 (at%) HEA after different processing conditions.
Figure 9. Comparison of the tensile properties of HPT-processed and post-deformation annealed V10Cr15Mn5Fe35Co10Ni25 (at%) HEA with other high-strength conventional and multi-component alloys [27, 46] (This graph was adapted from reference [47]).
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Conflicts of interest No conflict of interest exists. We confirm that the article is original, and the article has not been published previously. The article has been written by the stated authors who are all aware of its content and approve its submission, and the article is not under consideration for publication elsewhere. All authors have approved the manuscript and agree with its submission to Materials Science and Engineering: A.
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