HIGH ENTROPY ALLOY HAVING TWIP/TRIP PROPERTY AND MANUFACTURING METHOD FOR THE SAME

Abstract

The present invention relates to a high entropy alloy having more improved mechanical properties by controlling contents of additive elements in a NiCoFeMnCr 5-element alloy to control stacking fault energy, thereby controlling stability of a γ austenite phase to control a transformation mechanism, wherein the stacking fault energy is controlled in a composition range of Ni.sub.aCo.sub.bFe.sub.cMn.sub.dCr.sub.e (a+b+c+d+e=100, 1≦a≦50, 1≦b≦50, 1≦c≦50, 1≦d≦50, 10≦e≦25, and 77a−42b−22c+73d−100e+2186≦1500), and thus, the γ austenite phase exhibits a twin-induced plasticity (TWIP) property or a transformation induced-plasticity (TRIP) property in which the γ austenite phase is subjected to phase transformation into an ε martensite phase or an α′ martensite phase, under stress, thereby having improved strength and elongation at the same time to have excellent mechanical properties.

Claims

1. A high entropy alloy having TWIP (twin induced plasticity)/TRIP (transformation induced plasticity) property, which is represented by the following Chemical Formula:
Ni.sub.aCo.sub.bFe.sub.cMn.sub.dCr.sub.e  [Chemical Formula] (a+b+c+d+e=100, 1≦a≦50, 1≦b≦50, 1≦c≦50, 1≦d≦50, 10≦e≦25, and 77a−42b−22c+73d−100e+2186≦1500).

2. The high entropy alloy having TWIP/TRIP property of claim 1, wherein: the high entropy alloy includes 10 at. % or less of at least one element of C, N, Al, Ti, V, Cu, Zr, Nb, or Mo.

3. The high entropy alloy having TWIP/TRIP property of claim 1, wherein: in Chemical Formula above, 77a−42b−22c+73d−100e+2186≦500.

4. The high entropy alloy having TWIP/TRIP property of claim 1, wherein: in Chemical Formula above, 77a−42b−22c+73d−100e+2186≦200.

5. The high entropy alloy having TWIP/TRIP property of claim 1, wherein: the high entropy alloy includes a γ austenite single phase.

6. The high entropy alloy having TWIP/TRIP property of claim 1, wherein: the high entropy alloy simultaneously includes a γ austenite phase and an ε martensite phase.

7. The high entropy alloy having TWIP/TRIP property of claim 1, wherein: a γ austenite phase in the high entropy alloy is subjected to multi-stage phase transformation into an α′ martensite phase through an ε martensite phase during strain.

8. The high entropy alloy having TWIP/TRIP property of claim 1, wherein: a free energy change (ΔG.sub.hcp-fcc) when a γ austenite phase in the high entropy alloy is phase-transformed into an ε martensite phase during strain is 1500 J/mol or less (based on calculation of Thermo Calc, TCFE8).

9. The high entropy alloy having TWIP/TRIP property of claim 8, wherein: when the free energy change (ΔG.sub.hcp-fcc) is 1500 J/mol or less (based on calculation of Thermo Calc, TCFE8), the high entropy alloy exhibits the TWIP property, and when the free energy change (ΔG.sub.hcp-fcc) is 500 J/mol or less (based on calculation of Thermo Calc, TCFE8), the high entropy alloy exhibits the TRIP property.

10. The high entropy alloy having TWIP/TRIP property of claim 8, wherein: when the free energy change (ΔG.sub.hcp-fcc) is 200 J/mol or less (based on calculation of Thermo Calc, TCFE8), the high entropy alloy simultaneously includes the γ austenite phase and the ε martensite phase.

11. The high entropy alloy having TWIP/TRIP property of claim 8, wherein: the free energy change (ΔG.sub.hcp-fcc) is 500 J/mol or less (based on calculation of Thermo Calc, TCFE8), and a free energy change (ΔG.sub.bcc-fcc) at the time of phase transformation when the γ austenite phase in the high entropy alloy is phase-transformed into the α′ martensite phase is −2500 J/mol or less to −5000 J/mol or more (based on calculation of Thermo Calc, TCFE8).

12. The high entropy alloy having TWIP/TRIP property of claim 1, wherein: in Chemical Formula above, 1≦a≦7, 32≦b≦50, 32≦c≦50, 1≦d≦7, and 15≦e≦25.

13. A manufacturing method for a high entropy alloy having TWIP (twin induced plasticity)/TRIP (transformation induced plasticity) property comprising: preparing a raw material; and manufacturing the high entropy alloy by alloying the raw material, wherein in the preparing of the raw material, the raw material is prepared to satisfy the following Chemical Formula, and a free energy change (ΔG.sub.hcp-fcc) when a γ austenite phase (fcc) in the manufactured high entropy alloy is phase-transformed into an ε martensite phase (hcp) is 1500 J/mol or less (based on calculation of Thermo Calc, TCFE8):
Ni.sub.aCo.sub.bFe.sub.cMn.sub.dCr.sub.e  [Chemical Formula] (a+b+c+d+e=100, 1≦a≦50, 1≦b≦50, 1≦c≦50, 1≦d≦50, 10≦e≦25, and 77a−42b−22c+73d−100e+2186≦1500).

14. The manufacturing method of claim 13, further comprising: after the manufacturing of the high entropy alloy, performing homogenization treatment by hot rolling a manufactured ingot to 80% or less of an original thickness, and annealing in an Ar atmosphere at 1200±300° C. for 48 hours or less, followed by quenching.

15. The manufacturing method of claim 14, further comprising: controlling a microstructure size of the high entropy alloy by cold rolling the homogenized high entropy alloy to 10% or more of an original thickness, and annealing in an Ar atmosphere at 900±200° C. for 24 hours or less, followed by quenching.

16. The manufacturing method of claim 13, wherein: the free energy change (ΔG.sub.hcp-fcc) is 500 J/mol or less (based on calculation of Thermo Calc, TCFE8).

17. The manufacturing method of claim 13, wherein: the free energy change (ΔG.sub.hcp-fcc) is 200 J/mol or less (based on calculation of Thermo Calc, TCFE8).

18. The manufacturing method of claim 13, wherein: a free energy change (ΔG.sub.bcc-fcc) at the time of phase transformation when the γ austenite phase in the high entropy alloy is phase-transformed into an α′ martensite phase is −2500 J/mol or less to −5000 J/mol or more (based on calculation of Thermo Calc, TCFE8).

Description

BRIEF DESCRIPTION OF THE DRAWINGS

[0026] FIG. 1 is a graph showing a free energy change (ΔG.sub.hcp-fcc) at the time of phase transformation according to contents of respective elements in a NiCoFeMnCr 5-element alloy by using Calphad calculation.

[0027] FIG. 2 shows comparison between a free energy change (ΔG.sub.hcp-fcc) at the phase transformation predicted by Calphad calculation and a free energy change (ΔG′.sub.hcp-fcc) calculated by Equation derived from the present invention in Table 1.

[0028] FIG. 3 is a graph showing a free energy change (ΔG.sub.hcp-fcc) at the time of phase transformation by controlling a content ratio of (Fe, Co) to (Ni, Mn) in an increased direction in the NiCoFeMnCr 5-element alloy using Calphad calculation.

[0029] FIG. 4 shows results of an inverse pole figure map (IPF map) and a phase map obtained by electron backscattering diffraction (EBSD) measurement on high entropy alloy specimens of Comparative Example, and Examples 3 and 5 according to the present invention in Table 2.

[0030] FIG. 5 shows results of a tensile test at room temperature on the high entropy alloy specimens of Comparative Example, and Examples 3 and 5 according to the present invention in Table 2.

[0031] FIG. 6 shows results of a Kernal average misorientation map obtained by EBSD measurement of a 60% local strain region after the tensile test of the high entropic alloy of Comparative Example in Table 2.

[0032] FIG. 7 shows results of the IPF map and an electron channeling contrast image (ECCI) of a 15% local strain region (a1) and a 45% local strain region (b1) after the tensile test of the high entropy alloy of Example 3 according to the present invention in Table 2.

[0033] FIG. 8 shows results of the IPF map and the ECCI of the 15% local strain region (a1) and a 30% local strain region (b1) after the tensile test of the high entropy alloy of Example 5 according to the present invention in Table 2.

[0034] FIG. 9 is phase maps of (a) 15%, (b) 30%, (c) 60%, and (d) 75% local strain regions after the tensile test of the high entropy alloy of Example 5 according to the present invention in Table 2.

[0035] FIG. 10 shows a change in phase fraction according to a local strain rate in the tensile test of the high entropy alloy of Example 5 according to the present invention in Table 2.

[0036] FIG. 11 shows tensile test results at room temperature on a Ni.sub.20Co.sub.20Fe.sub.20Mn.sub.20Cr.sub.20 high entropy alloy specimen of Comparative Example 2 and a Ni.sub.8Co.sub.32Fe.sub.32Mn.sub.8Cr.sub.20 high entropy alloy specimen of Example 4 according to the present invention in Table 3.

[0037] FIG. 12 shows phase map results obtained by electron backscattering diffraction (EBSD) measurements before the tensile test (a) and after the tensile test (b) on the Ni.sub.20Co.sub.20Fe.sub.20Mn.sub.20Cr.sub.20 high entropy alloy specimen of Comparative Example 2 and the Ni.sub.8Co.sub.32Fe.sub.32Mn.sub.8Cr.sub.20 high entropy alloy specimen of Example 4 according to the present invention in Table 3.

[0038] FIG. 13 is a graph showing a free energy change (ΔG.sub.bcc-fcc) at the time of phase transformation by controlling a content ratio of (Fe, Co) to (Ni, Mn) in an increased direction in the NiCoFeMnCr 5-element alloy using Calphad calculation.

[0039] FIG. 14 shows X-ray diffraction analysis results showing phase transformation behaviors of (a) before cold rolling and (b) after cold rolling of the Ni.sub.5Co.sub.35Fe.sub.35Mn.sub.5Cr.sub.20 high entropy alloy specimen of Example 12 according to the present invention.

DETAILED DESCRIPTION OF THE EMBODIMENTS

[0040] Embodiments of the present invention are described in detail with reference to accompanying drawings.

[0041] The present invention is intended to further improve mechanical properties by controlling stacking fault energy of the above-described high entropy alloy to control a strain mechanism, and has an object of providing a high entropy alloy having excellent mechanical properties in which a γ austenite single phase microstructure or a dual-phase microstructure simultaneously having a γ austenite phase and an ε martensite phase is formed, and the γ austenite phase exhibits a twin-induced plasticity (TWIP) property or a transformation induced-plasticity (TRIP) property by an ε or α′ martensite phase under stress, thereby having improved strength and elongation at the same time.

[0042] To this end, the high entropy alloy of the present invention is composed of five elements of Cr, Mn, Fe, Co, and Ni, which are metal elements having a similar interatomic size of 10% or less and a mixed thermal relation close to almost 0. However, as compared to the conventional high entropy alloy, i.e., Cr.sub.20Mn.sub.20Fe.sub.20Co.sub.20Ni.sub.20 (so called a Cantor alloy) in which respective constituent elements are constituted at the same fraction, the high entropy alloy of the present invention has a state in which an entire composition is not in the same fraction by predicting an alloy system with a small free energy change (ΔG.sub.hcp-fcc) during the phase transformation from the γ austenite (FCC) into the ε martensite (HCP) through Calphad calculation and first principle calculation. In particular, a content ratio of (Fe, Co) to (Ni, Mn) may be controlled in an increased direction to reduce the stacking fault energy, and thus, the TWIP/TRIP behavior may be generated on the austenite phase at the time of stress application.

[0043] FIG. 1 is a graph showing a free energy change (ΔG.sub.hcp-fcc) at the time of phase transformation according to contents of respective elements in a NiCoFeMnCr 5-element alloy by using Calphad calculation. When one element changes, remaining elements are allowed to have the same fraction. At this time, compositions of the respective elements varied from 5 to 50%. Table 1 shows calculation results.

TABLE-US-00001 TABLE 1 Composition Cr Mn Fe Co Ni G.sub.fcc G.sub.hcp ΔG.sub.hcp-fcc ΔG′.sub.hcp-fcc Cr 1 5 23.75 23.75 23.75 23.75 −14615 −10883 3732 3755 2 10 22.5 22.5 22.5 22.5 −13858 −10683 3175 3146 3 15 21.25 21.25 21.25 21.25 −13031 −10476 2555 2537 4 20 20 20 20 20 −12162 −10234 1928 1928 5 25 18.75 18.75 18.75 18.75 −11267 −9965 1302 1319 6 30 17.5 17.5 17.5 17.5 −10360 −9674 686 710 7 35 16.25 16.25 16.25 16.25 −9363 −9450 −87 101 8 40 15 15 15 15 −8546 −9033 −487 −507 9 45 13.75 13.75 13.75 13.75 −7656 −8683 −1027 −1116 10 50 12.5 12.5 12.5 12.5 −6788 −8314 −1526 −1725 Mn 11 23.75 5 23.75 23.75 23.75 −8831 −8708 123 504 12 22.5 10 22.5 22.5 22.5 −10068 −9396 672 978 13 21.25 15 21.25 21.25 21.25 −11211 −9856 1355 1453 14 20 20 20 20 20 −12162 −10234 1928 1928 15 18.75 25 18.75 18.75 18.75 −12935 −10538 2397 2403 16 17.5 30 17.5 17.5 17.5 −13541 −10776 2765 2878 17 16.25 35 16.25 16.25 16.25 −13989 −10949 3040 3352 18 15 40 15 15 15 −14287 −11063 3224 3827 19 13.75 45 13.75 13.75 13.75 −14450 −11120 3330 4302 20 12.5 50 12.5 12.5 12.5 −14480 −11129 3351 4777 Fe 21 23.75 23.75 5 23.75 23.75 −13108 −10840 2268 2289 22 22.5 22.5 10 22.5 22.5 −12862 −10694 2168 2168 23 21.25 21.25 15 21.25 21.25 −12538 −10484 2054 2048 24 20 20 20 20 20 −12162 −10234 1928 1928 25 18.75 18.75 25 18.75 18.75 −11744 −9955 1789 1808 26 17.5 17.5 30 17.5 17.5 −11294 −9654 1640 1688 27 16.25 16.25 35 16.25 16.25 −10816 −9334 1482 1568 28 15 15 40 15 15 −10314 −9000 1314 1448 29 13.75 13.75 45 13.75 13.75 −9793 −8654 1139 1328 30 12.5 12.5 50 12.5 12.5 −9254 −8296 958 1208 Co 31 23.75 23.75 23.75 5 23.75 −12472 −9665 2807 2666 32 22.5 22.5 22.5 10 22.5 −12447 −9952 2495 2420 33 21.25 21.25 21.25 15 21.25 −12337 −10134 2203 2174 34 20 20 20 20 20 −12162 −10234 1928 1928 35 18.75 18.75 18.75 25 18.75 −11930 −10262 1668 1682 36 17.5 17.5 17.5 30 17.5 −11648 −10222 1426 1436 37 16.25 16.25 16.25 35 16.25 −11318 −10118 1200 1190 38 15 15 15 40 15 −10941 −9952 989 944 39 13.75 13.75 13.75 45 13.75 −10519 −9728 791 698 40 12.5 12.5 12.5 50 12.5 −10063 −9472 591 452 Ni 41 23.75 23.75 23.75 23.75 5 −9981 −9642 339 428 42 22.5 22.5 22.5 22.5 10 −10780 −9903 877 928 43 21.25 21.25 21.25 21.25 15 −11503 −10093 1410 1428 44 20 20 20 20 20 −12162 −10234 1928 1928 45 18.75 18.75 18.75  8.75 25 −12757 −10335 2422 2428 46 17.5 17.5 17.5 17.5 30 −13283 −10399 2884 2928 47 16.25 16.25 16.25 16.25 35 −13734 −10428 3306 3428 48 15 15 15 15 40 −14102 −10422 3680 3928 49 13.75 13.75 13.75 13.75 45 −14379 −10382 3997 4428 50 12.5 12.5 12.5 12.5 50 −14556 −10306 4250 4928

[0044] As shown in the drawing, it could be confirmed that the free energy change (ΔG.sub.hcp-fcc) was gradually decreased when contents of Cr, Fe and Co were increased, and the free energy change (ΔG.sub.hcp-fcc) was gradually decreased when contents of Mn and Ni were decreased. It could be confirmed that linearity was maintained when the contents of the respective elements were within 5 to 50%. Based on this, the free energy change (ΔG.sub.hcp-fcc) at the time of the phase transformation could be predicted through simple fitting equation in consideration of contribution degree of the respective elements at the time of phase transformation, and a result value was named ΔG′.sub.hcp-fcc.

[0045] The ΔG′.sub.hcp-fcc could be calculated by the following Equation 1.

ΔG′.sub.hcp-fcc=77a−42b−22c+73d−100e+2186  [Equation 1]

[0046] In Equation 1, a, b, c, d, and e are compositions (at %) of the respective elements. Calculated values of the ΔG′.sub.hcp-fcc according to the respective composition are shown in Table 1 below.

[0047] FIG. 2 shows comparison between a free energy change (ΔG.sub.hcp-fcc) at the phase transformation predicted by Calphad calculation and a free energy change (ΔG′.sub.hcp-fcc) calculated by Equation derived from the present invention in Table 1. As appreciated in the drawing, it could be confirmed that excellent linearity was exhibited in the entire composition range of the present invention, which indicated that the composition of the high entropy alloy could be efficiently controlled by using the ΔG′.sub.hcp-fcc Equation derived from the present invention.

[0048] FIG. 3 is a graph showing a free energy change (ΔG.sub.hcp-fcc) at the time of phase transformation by controlling a content ratio of (Fe, Co) to (Ni, Mn) in an increased direction in the NiCoFeMnCr 5-element alloy using Calphad calculation. As shown in the drawing, it could be confirmed that the free energy change (ΔG.sub.hcp-fcc) was gradually decreased at the time of the phase transformation by making a non-equi-atomic state in the Cantor alloy and decreasing the contents of Ni and Mn while increasing the contents of Fe and Co.

[0049] For example, the free energy change (ΔG.sub.hcp-fcc) value was greatly reduced to be 1500 J/mol or less (based on Calphad calculation) at the time of the phase transformation by simultaneously controlling the contents of Fe and Co to be 22 at. % or more and the contents of Ni and Mn to be 18 at. % or less which are the composition ranges of the present invention in the Cantor alloy so that the fraction of the constituent elements was not the same. This reduction of the free energy change (ΔG.sub.hcp-fcc) with respect to the phase transformation is closely related to a decrease in the stacking fault energy in materials, and thus, a TWIP strain mechanism due to activation of a twin may appear. In addition, when a free energy change (ΔG.sub.hcp-fcc) value is decreased to 500 J/mol or less (based on Calphad calculation), the twin strain is further activated by instability of the γ austenite phase, and thus, it is possible to generate transformation into the TRIP strain mechanism, i.e., phase transformation into the martensite phase during the strain. Further, when the free energy change (ΔG.sub.hcp-fcc) value at the time of the phase transformation is greatly decreased and becomes 200 J/mol or less (Calphad calculation standard), stability of the ε martensite phase is stabilized, and a dual-phase microstructure simultaneously having the γ austenite phase and the ε martensite phase at the time of solidification appears, and in this case, the γ phase in the dual-phase exhibits the TRIP strain mechanism.

[0050] To investigate effects of the decrease in the free energy change (ΔG.sub.hcp-fcc) on a material strain behavior at the time of the phase transformation, Cr, Mn, Fe, Co, and Ni constituting the alloy were prepared as parent elements with a purity of 99.9%, and casted by an induction melting method having stirring effect by an electromagnetic field, and homogenized by hot rolling at a temperature of 900° C. to 50% and treating the alloy in an Ar atmosphere at 1200° C. for 3 hours. In the present invention, since the casting is able to implementing a high temperature through an arc plasma in addition to the induction melting method, the alloy is able to be manufactured through a commercial casting process by utilizing an arc melting method capable of rapidly forming a bulk homogeneous solid solution and minimizing impurities such as oxides and pores, and a resistance heating method in which temperature is able to be precisely controlled. In addition to the commercial casting method capable of dissolving raw material metals, the alloy may be manufactured by preparing raw materials into powder and the like, and sintering the materials at high temperature/high pressure by spark plasma sintering or hot isostatic pressing using powder metallurgy, and in the case of the sintering method, there are advantages in that the microstructure may be more precisely controlled and components in desired shapes may be easily manufactured. The homogenization treatment of the manufactured specimen is preferably performed by hot rolling the manufactured ingot to 80% or less of an original thickness, and annealing in an Ar atmosphere at 1200±300° C. for 48 hours or less, followed by quenching.

[0051] Then, to obtain an appropriate grain size, cold rolling was performed to obtain a thickness reduction of 80% (single phase) or 50% (dual-phase), followed by annealing in an Ar atmosphere at 900° C. for 3 minutes (single phase) or 800° C. (dual-phase) for 1 hour, thereby refining the crystal grains. In order to control the size of the microstructure of the homogenized high entropy alloy specimen, it is preferred to perform cold rolling the homogenized high entropy alloy specimen to 10% or more of an original thickness, and annealing in an Ar atmosphere at about 900±200° C. for 24 hours or less, followed by quenching.

[0052] Mixing ratios of the elements of Examples of the present invention are shown in Table 2 below.

TABLE-US-00002 TABLE 2 Deformation ΔG.sub.hcp-fcc mechanism Specimen Composition(at %) (J/mol) (Phase species) Comparative Cr.sub.20Mn.sub.20Fe.sub.20Co.sub.20Ni.sub.20 1927.8 Dislocation (γ) Example 1 Example 1 Cr.sub.20Mn.sub.18Fe.sub.22Co.sub.22Ni.sub.18 1494.4 TWIP (γ) Example 2 Cr.sub.20Mn.sub.16Fe.sub.24Co.sub.24Ni.sub.16 1105.4 TWIP (γ) Example 3 Cr.sub.20Mn.sub.14Fe.sub.26Co.sub.26Ni.sub.14 771.0 TWIP (γ) Example 4 Cr.sub.20Mn.sub.12Fe.sub.28Co.sub.28Ni.sub.12 482.9 TRIP (γ) Example 5 Cr.sub.20Mn.sub.10Fe.sub.30Co.sub.30Ni.sub.10 245.3 TRIP (γ) Example 6 Cr.sub.20Mn.sub.8Fe.sub.32Co.sub.32Ni.sub.8 59.4 TRIP (γ + ε) Example 7 Cr.sub.20Mn.sub.6Fe.sub.34Co.sub.34Ni.sub.6 −74 TRIP (γ + ε) Example 8 Cr.sub.20Mn.sub.4Fe.sub.36Co.sub.36Ni.sub.4 −154 TRIP (γ + ε) Example 9 Cr.sub.20Mn.sub.2Fe.sub.38Co.sub.38Ni.sub.2 −178 TRIP (γ + ε) Comparative Cr.sub.20Mn.sub.0Fe.sub.40Co.sub.40Ni.sub.0 −147 Dislocation (BCC) Example 2

[0053] As shown in Table 2, a Cr.sub.20Mn.sub.20Fe.sub.20Co.sub.20Ni.sub.20 high entropy alloy including constituent elements in the same fraction was manufactured as Comparative Example 1 of the present invention.

[0054] FIG. 4 shows results of an inverse pole figure map (IPF map) and a phase map obtained by electron backscattering diffraction (EBSD) measurement on high entropy alloy specimens of Comparative Example 1, and Examples 3 and 5 according to the present invention.

[0055] As shown in the IPF map for the specimens of Comparative Example 1 (a), Example 3 (c), and Example 5 (e) in FIG. 4, it could be confirmed that the crystal grains were homogenized on all the specimens after the specimen manufacturing process, and had an average grain size of about 4 μm.

[0056] In addition, all the specimens were measured to be γ FCC single phases through the phase map results on the specimens of Comparative Example 1 (b), Example 3 (d), and Example 5 (f) in FIG. 4, and thus, it could be confirmed that the specimens of Examples 1 to 5 of the present invention as well as the specimens of Comparative Examples were also the γ FCC single phase high entropy alloys.

[0057] FIG. 5 shows results of a tensile test at room temperature on the high entropy alloy specimens of Comparative Example 1, and Examples 3 and 5 according to the present invention. It could be confirmed that the high entropy alloy (Cantor alloy) of Comparative Example 1, in which the constituent elements are included in the same fraction, had a tensile strength of 690 MPa and an elongation of 55%, and thus, mechanical properties were excellent due to basic properties of the high entropy alloy.

[0058] However, it could be confirmed that the high entropy alloy specimen of Example 3 (dot) and Example 5 (solid) according to the present invention had a similar yield strength and further improved mechanical properties as compared to those of the high entropy alloy of Comparative Example 1. Specifically, the high entropy alloy of Example 3 had an increased tensile strength of 740 MPa and an elongation of up to 65%, and the high entropy alloy of Example 5 had a more increased tensile strength of up to 860 MPa and an elongation of up to 70%.

[0059] As described above, the mechanical properties of the high entropy alloy according to the present embodiment were further improved as compared to those of the high entropy alloy having the same fraction according to Comparative Example 1, which is known to have excellent mechanical properties. Hereinafter, a reason that the mechanical properties of the high entropy alloy according to the present embodiment are improved is described.

[0060] FIG. 6 shows results of a Kernal average misorientation map obtained by EBSD measurement of a 60% local strain region after the tensile test of the high entropic alloy of Comparative Example 1. As shown in the drawing, even in a region with 60% large strain, a strain mechanism by dislocation in which strain occurs by activation of the dislocation could be confirmed as previously reported, and twin formation or phase transformation, etc., could be not confirmed.

[0061] FIG. 7 shows results of the IPF map and the ECCI of the 15% local strain region (a1) and the 45% local strain region (b1) after the tensile test of the high entropy alloy of Example 3 according to the present invention.

[0062] As shown in the case where the high entropy alloy of Example 3 was strained, it could be confirmed that the stacking fault band (SF band) was shown in some regions (a3) in the 15% local strain region (a1), but the strain by dislocation (a2) was mainly shown. On the contrary, when the strain rate was increased to 45%, the stacking fault band and the twin were found as shown in (b2) and (b3), and the strain mechanism exhibited the plastic strain properties of TWIP. The TWIP effect is a phenomenon that is generated in alloys having a FCC structure with sufficiently small stacking fault energy, and is a technique that improves mechanical properties of materials using the twin that is generated during the strain. Specifically, when the twin is generated during the strain, the movement of the dislocation and propagation of cracks, etc., are interrupted, and thus, work hardenability and elongation of the material are improved. It could be appreciated that mechanical properties of the high entropy alloy of Example 3 were improved due to the TWIP effect which was not exhibited in the high entropy alloy of Comparative Example. In particular, in the high entropy alloy of the present invention, the twin was activated even at room temperature, thereby promoting improvement of the mechanical properties at the time of strain.

[0063] FIG. 8 shows results of the IPF map and the ECCI of the 15% local strain region (a1) and the 30% local strain region (b1) after the tensile test of the high entropy alloy of Example 5 according to the present invention.

[0064] As shown in the case where the high entropy alloy of Example 5 was strained, it could be confirmed that the stacking fault band (a2), and twin (a3) were found in the 15% local strain region (a1), and thus, the strain mechanism exhibited plastic strain properties of TWIP. In particular, as shown in the phase map (b1) for the 30% local strain region (b1), it could be confirmed that the existing γ FCC phase (Black) region was converted to the ε HCP phase (Grey) region through the phase transformation. This phenomenon is also shown in alloys with FCC structure having sufficiently small stacking fault energy, and it may be confirmed that the plastic strain properties of TRIP in which the phase transformation between FCC phase and HCP phase is generated through twin strain mechanism during the strain, were shown. The TRIP effect is similar to the TWIP effect, but utilizes a stress-induced martensite phase which is generated during the strain. Specifically, when the phase transformation is generated during the strain, a twist effect is generated in the material due to the formed martensite phase, thereby increasing strength of the material, and movement of the dislocation is interrupted by a new interface, and thus, work hardenability and elongation of the material are further improved. It may be appreciated that mechanical properties of the high entropy alloy of Example 5 were further improved due to the TRIP effect which was not exhibited in the high entropy alloy of Comparative Example.

[0065] FIG. 9 is phase maps of (a) 15%, (b) 30%, (c) 60%, and (d) 75% local strain regions after the tensile test of the high entropy alloy of Example 5 according to the present invention. As shown in the drawing, it could be clearly confirmed that the phase transformation from the existing γ FCC phase (black) region to the ε HCP phase (gray) region was accelerated as the local strain was increased in the high entropy alloy of Example 5.

[0066] FIG. 10 shows a change in phase fraction according to a local strain rate in the tensile test of the high entropy alloy of Example 5 according to the present invention. It could be appreciated that in Example 5, the phase transformation started to be generated in a region having a local strain rate of about 18%, and about 80% of phase transformation from γ to ε was generated at 140% local strain.

[0067] Table 3 shows mixing ratios of elements of Comparative Examples 1 to 4 and Examples 1 to 17 of the present invention, and kinds of constitution phases thereof.

TABLE-US-00003 TABLE 3 Kinds of Specimen Composition (at %) main phase Comparative Example 1 Ni.sub.0Co.sub.40Fe.sub.40Mn.sub.0Cr.sub.20 α Comparative Example 2 Ni.sub.20Co.sub.20Fe.sub.20Mn.sub.20Cr.sub.20 γ Comparative Example 3 Ni.sub.26Co.sub.14Fe.sub.14Mn.sub.26Cr.sub.20 γ Comparative Example 4 Ni.sub.14Co.sub.21Fe.sub.21Mn.sub.14Cr.sub.30 γ + IC Example 1 Ni.sub.8Co.sub.19Fe.sub.45Mn.sub.8Cr.sub.20 γ + ε Example 2 Ni.sub.8Co.sub.34Fe.sub.34Mn.sub.8Cr.sub.16 γ + ε Example 3 Ni.sub.8Co.sub.30Fe.sub.30Mn.sub.8Cr.sub.24 γ + ε Example 4 Ni.sub.8Co.sub.32Fe.sub.32Mn.sub.8Cr.sub.20 γ + ε Example 5 Ni.sub.8Co.sub.32Fe.sub.32Mn.sub.8Cr.sub.18C.sub.2 γ + ε Example 6 Ni.sub.8Co.sub.32Fe.sub.32Mn.sub.8Cr.sub.18Al.sub.2 γ + ε Example 7 Ni.sub.8Co.sub.32Fe.sub.32Mn.sub.8Cr.sub.18Ti.sub.2 γ + ε Example 8 Ni.sub.8Co.sub.32Fe.sub.32Mn.sub.8Cr.sub.18Nb.sub.2 γ + ε Example 9 Ni.sub.8Co.sub.32Fe.sub.32Mn.sub.8Cr.sub.18Cu.sub.2 γ + ε Example 10 Ni.sub.6Co.sub.34Fe.sub.34Mn.sub.6Cr.sub.20 γ + ε(α′) Example 11 Ni.sub.5Co.sub.35Fe.sub.35Mn.sub.5Cr.sub.20 γ + ε(α′) Example 12 Ni.sub.4Co.sub.35Fe.sub.35Mn.sub.4Cr.sub.20N.sub.2 γ + ε(α′) Example 13 Ni.sub.4Co.sub.35Fe.sub.35Mn.sub.4Cr.sub.20V.sub.2 γ + ε(α′) Example 14 Ni.sub.4Co.sub.35Fe.sub.35Mn.sub.4Cr.sub.20Zr.sub.2 γ + ε(α′) Example 15 Ni.sub.4Co.sub.35Fe.sub.35Mn.sub.4Cr.sub.20Mo.sub.2 γ + ε(α′) Example 16 Ni.sub.2Co.sub.38Fe.sub.38Mn.sub.2Cr.sub.20 γ + ε(α′)

[0068] It was confirmed that in the alloys of Comparative Examples 1 to 4, the γ or a single phase (some intermetallic compounds (ICs) were able to be precipitated) was formed. On the other hand, in Examples 1 to 16 of the present invention, the γ austenite and the ε (α′) martensite phase were simultaneously precipitated to show a dual-phase microstructure.

[0069] Meanwhile, as shown in Table 3, when the high entropy alloy of the present invention further included 10 at. % or less of at least one element selected from additive elements such as C, N, Al, Ti, V, Cu, Zr, Nb, and Mo, etc., it was possible to improve properties by strengthening solid solution or by strengthening precipitation while maintaining the existing dual-phase matrix structure.

[0070] FIG. 11 shows tensile test results at room temperature on a Ni.sub.20Co.sub.20Fe.sub.20Mn.sub.20Cr.sub.20 high entropy alloy specimen of Comparative Example 2 and a Ni.sub.8Co.sub.32Fe.sub.32Mn.sub.8Cr.sub.20 high entropy alloy specimen of Example 4 according to the present invention. It could be confirmed that the high entropy alloy (dash dot) of Example 4 according to the present invention was improved both in strength and elongation at the same time, thereby further improving mechanical properties, as compared to those of the high entropy alloy (solid) of Comparative Example. Specifically, the high entropy alloy of Example 5 had greatly increased tensile strength to about 980 MPa and greatly increased elongation of up to about 60%.

[0071] As described above, the mechanical properties of the dual-phase high entropy alloy according to the present embodiment were further improved as compared to those of the high entropy alloy having the same fraction according to Comparative Example 2, which is known to have excellent mechanical properties. Hereinafter, a reason that the mechanical properties of the high entropy alloy according to the present embodiment are improved is described.

[0072] FIG. 12 shows phase map results obtained by electron backscattering diffraction (EBSD) measurements before the tensile test (a) and after the tensile test (b) on the Ni.sub.20Co.sub.20Fe.sub.20Mn.sub.20Cr.sub.20 high entropy alloy specimen of Comparative Example 2 and the Ni.sub.8Co.sub.32Fe.sub.32Mn.sub.8Cr.sub.20 high entropy alloy specimen of Example 4 according to the present invention in Table 3. As shown in the drawing, the twin generated during the strain was activated in the region subjected to 60% of large strain, and the γ austenite phase and the ε martensite phase had phase fractions of 87.5% and 12.5% before strain (FIG. 10(a)), respectively, which were changed to 41.6% and 58.4% after strain (FIG. 10(b)). Further, when the twin was generated during the strain in a metastable γ austenite phase having a low stacking fault energy, the movement of the dislocation, propagation of cracks, etc., were interrupted, and thus, work hardenability and elongation of the material were improved, and finally, phase instability was activated and the γ austenite phase was phase-transformed into the ε martensite phase.

[0073] FIG. 13 is a graph showing a free energy change (ΔG.sub.bcc-fcc) at the time of phase transformation into the α′ martensite phase by controlling a content ratio of (Fe, Co) to (Ni, Mn) in an increased direction in the DeletedTextsalloy using Calphad calculation. As shown in the drawing, when the free energy change (ΔG.sub.bcc-fcc) at the time of the phase transformation from the γ austenite phase to the α′ martensite phase was −2500 J/mol or less (based on Calphad calculation) among dual-phase high entropy alloys in which the free energy change (ΔG.sub.hcp-fcc) at the time of the phase transformation from the γ austenite phase to the ε martensite phase was 200 J/mol or less (based on Calphad calculation), it was possible to obtain a dual-phase high entropy alloy capable of being subjected to stress-induced multi-stage phase transformation from the metastable γ austenite through the ε martensite phase up to the α′ martensite phase due to shear deformation through intersections of ε martensite bands by the phase transformation during strain.

[0074] However, when the free energy change (ΔG.sub.bcc-fcc) at the time of the phase transformation from the γ austenite phase to the α′ martensite phase was extremely low to −5000 J/mol or less (based on Calphad calculation), since stability of the metastable γ austenite phase was significantly reduced, a pre-strain microstructure became an α phase having a BCC structure, and thus, the dual-phase high entropy alloy of the present invention could not be manufactured.

[0075] FIG. 14 shows X-ray diffraction analysis results showing phase transformation behaviors of (a) before cold rolling (upper graph in the drawing) and (b) after cold rolling (lower graph in the drawing) of the Ni.sub.5Co.sub.35Fe.sub.35Mn.sub.5Cr.sub.20 high entropy alloy specimen of Example 12 according to the present invention. As shown in the drawing, it could be confirmed that in the Ni.sub.5Co.sub.35Fe.sub.35Mn.sub.5Cr.sub.20 high entropy alloy, the microstructure in which the γ phase is a main phase in the dual-phase with γ phase and ε (α′) phase before cold rolling was changed to the dual-phase structure in which the fraction of the ε phase and the α′ phase was increased after cold rolling. These results indicate that as the stacking fault energy in the material is reduced, the metastable γ austenite phase undergoes the multi-stage phase transformation to the α′ martensite phase through the ε martensite phase during strain, which may further contribute to improvement of tensile strength and elongation. Further, when the phase transformation into the martensite phase is generated during the strain, a twist effect is generated in the material due to the formed martensite phase, thereby increasing strength of the material, and movement of the dislocation is interrupted by a new interface, and thus, work hardenability and elongation of the material are improved, and mechanical properties are more improved when the multi-stage phase transformation process is performed.

[0076] As reviewed above, according to the present invention, the high entropy alloy composed of five elements of Ni, Co, Fe, Mn, and Cr was developed, wherein the high entropy alloy had improved mechanical properties in which strength and elongation were simultaneously improved through the implementation of the single phase TWIP, TRIP or the dual-phase TRIP, the multi-stage phase TRIP strain mechanism, by reducing the stacking fault energy. In this case, the stacking fault energy was predictable by Calphad calculation, and was capable of being easily predicted by ΔG′.sub.hcp-fcc=77a−42b−22c+73d−100e+2186 (a, b, c, d, and e represent compositions (at %) of Ni, Co, Fe, Co, and Cr, respectively) which was derived through fitting of the calculation results. As a result, the high entropy alloy having the TWIP/TRIP property according to Examples may be applied not only as materials for offshore plants and structural materials for polar extreme environment which require excellent toughness and high strength at a low temperature, but also as structural materials for high-temperature extreme environment which require an excellent high-temperature creep property and high-temperature strength through the low stacking fault energy, such as projectile propulsion units, nuclear pressure vessels, cladding tubes, and high-efficiency next generation turbine blades for thermal power generation, etc.

[0077] Although the preferred embodiments of the present invention have been disclosed for illustrative purposes, those skilled in the art will appreciate that Examples described above are provided to exemplarily describe technical idea of the present invention, and various modifications, additions and substitutions are possible, without departing from the scope and spirit of the invention as disclosed in the accompanying claims.

[0078] The protection scope of the present invention must be analyzed by the appended claims and it should be analyzed that all spirits within a scope equivalent thereto are included in the appended claims of the present invention.

[0079] While this invention has been described in connection with what is presently considered to be practical exemplary embodiments, it is to be understood that the invention is not limited to the disclosed embodiments, but, on the contrary, is intended to cover various modifications and equivalent arrangements included within the spirit and scope of the appended claims.