Medium-entropy alloy and preparation method and application thereof

Abstract

The invention relates to the technical field of metal alloy materials, in particular to a medium-entropy alloy and a preparation method and application thereof. The chemical formula of the medium-entropy alloy is Co < 54 > Cr < 25 > Ni < 10 > Al < 3 > Ti < 3 > W < 5 >. The Co-rich low-fault-energy medium-entropy alloy is designed, so that the deformation mechanism of the alloy is dominated by fault, and the alloy can keep high strength and excellent elongation at the same time. Through proper component design, under the combined action of Cr, Ti and W elements, the alloy keeps certain corrosion resistance while having excellent mechanical properties, has high corrosion potential in 3.5 wt.% NaCl, and has low corrosion current density in a 0.5 M H2SO4 solution.

Description

Technical Field

[0001] The present invention relates to the technical field of metal alloy materials, and in particular to a medium-entropy alloy and a preparation method and application thereof. Background Art

[0002] In recent years, multi-principal alloys, as a new type of metal alloy material, have attracted widespread attention due to their unique four core effects, excellent mechanical properties and physical and chemical properties. This new type of metal alloy has shown great application potential in many fields and has quickly become a hot spot in materials science research. Among the many multi-principal alloy systems, CoCrNi medium-entropy alloy stands out due to its excellent room temperature tensile plasticity and high fracture toughness. However, its yield strength and tensile strength are relatively low, which to some extent limits its wider application. Summary of the invention

[0003] Based on the above, the present invention provides a medium entropy alloy and a preparation method and application thereof. The present invention precisely controls the formation of precipitation phase by adding Al, Ti and W trace elements to the alloy matrix (CoCrNi). In order to further optimize the microstructure and performance of the alloy, the present invention also combines hot extrusion, cold rolling and heat treatment process means to control the grain size and precipitation phase distribution of the alloy, thereby optimizing and improving the performance of the alloy.

[0004] To achieve the above object, the present invention provides the following solutions:

[0005] One of the technical solutions of the present invention is a medium entropy alloy having a chemical formula of Co 54 Cr 25 Ni 10 Al3Ti3W5(at.%).

[0006] A second technical solution of the present invention is a method for preparing the above-mentioned medium-entropy alloy, comprising the following steps:

[0007] By Co 54 Cr 25 Ni 10 Al3Ti3W5 atomic percentage Co, Cr, Ni, Al, Ti, W metal single substances are mixed and vacuum arc melted to obtain button ingots;

[0008] The button ingot is subjected to induction melting, followed by casting to obtain a medium entropy alloy ingot;

[0009] The medium-entropy alloy ingot is subjected to homogenization annealing, then heat-insulated, hot-extruded, and cooled to obtain the medium-entropy alloy.

[0010] A third technical solution of the present invention is the use of the above-mentioned medium-entropy alloy in the preparation of structural materials for service in seawater environments or acidic environments.

[0011] A fourth technical solution of the present invention is a structural material, the raw materials of which include the above-mentioned medium-entropy alloy.

[0012] The present invention discloses the following technical effects:

[0013] (1) The present invention designs a Co-rich low stacking fault energy medium entropy alloy, so that the deformation mechanism of the alloy is dominated by stacking faults, thereby maintaining high strength while maintaining excellent elongation.

[0014] (2) Through appropriate composition design, under the combined action of Cr, Ti and W elements, the alloy has excellent mechanical properties while maintaining a certain corrosion resistance, with a high pitting potential in 3.5wt.% NaCl and a low corrosion current density in 0.5M H2SO4 solution.

[0015] (3) The present invention applies the hot extrusion preparation process commonly used in traditional alloys to the preparation of cast medium-entropy alloys. The present invention obtains a medium-entropy alloy with high strength and high plasticity while maintaining certain corrosion resistance by adjusting the preparation and post-treatment processes, which can be used as a structural material in seawater environment or acidic environment. BRIEF DESCRIPTION OF THE DRAWINGS

[0016] In order to more clearly illustrate the embodiments of the present invention or the technical solutions in the prior art, the drawings required for use in the embodiments will be briefly introduced below. Obviously, the drawings described below are only some embodiments of the present invention. For ordinary technicians in this field, other drawings can be obtained based on these drawings without paying creative work.

[0017] Figure 1 (a) Vickers hardness, (b) engineering stress-strain curve, and (c) strength and elongation curves of HE, HE-873-1, HE-973-1, HE-1073-1, and HE-1173-1 samples.

[0018] Figure 2 (a) Vickers hardness, (b) engineering tensile stress-strain curve, (c) strength and elongation curves of HE, HE-973-1, HE-973-8, HE-973-12 and HE-973-16 samples, and (d) comparison of strength-elongation product and yield strength with data in the literature.

[0019] Figure 3(a) XRD results of HE-873-1, HE-973-1, HE-1073-1 and HE-1173-1 samples, (b) XRD results of HE-973-8, HE-973-12 and HE-973-16 samples.

[0020] Figure 4 SEM images of (a) HE, (b) HE-873-1, (c) HE-973-1, (d) HE-1073-1, and (e) HE-1173-1 samples, (f) magnified SEM image of HE-1173-1 sample; EDS results of the area indicated by the red box in (f): (g) Co, (h) Cr, (i) Ni, (j) Al, (k) Ti, and (l) W.

[0021] Figure 5 SEM images of (a) HE-973-8, (b) HE-973-12, and (c) HE-973-16 samples.

[0022] Figure 6 (a) TEM bright field image of HE-973-12 sample, (b) TEM dark field image of R1 region in (a), and (c) SAED result; (d) TEM bright field image of annealed twins and (e) SAED result of R2 region in (d); (f) TEM bright field image of stacking fault.

[0023] Figure 7 The Fourier transform results of (a) high-angle annular dark field, (b) R3 area and (c) R4 area of ​​HE-973-12 sample, and the EDS surface scan results of the corresponding area of ​​(a): (d) Co, (e) Cr, (f) Ni, (g) Al, (h) Ti, and (i) W.

[0024] Figure 8 (a) True stress-strain and strain hardening rate curves of HE-973-12 specimen, (b) schematic diagram of microstructural changes during deformation process, (c)-(e) KAM diagrams and (f)-(h) BC diagrams of HE-973-12 specimens at different strains.

[0025] Fig. 9 (a) Engineering stress-strain curves of HE, HE+CR, HE+CR+1273K 1 min, and HE+CR+1273K 1 min+973K 12 h samples; (b) Vickers hardness.

[0026] Fig.10 These are the XRD results of HE+CR, HE+CR+1273K 1min and HE+CR+1273K 1min+973K 12h samples.

[0027] Fig.11 SEM images of (a) HE+CR, (b) HE+CR+1273K 1 min and (c) HE+CR+1273K 1 min+973K 12 h samples; EDS image of the area indicated by the red box in (c): (d) Co, (e) Cr, (f) Ni, (g) Al, (h) Ti, (i) W.

[0028] Fig.12 BC diagram, IPF diagram and GND diagram of (a)(d)(g) HE+CR, (b)(e)(h) HE+CR+1273K 1min and (c)(f)(i) HE+CR+1273K1min+973K 12h samples.

[0029] Fig.13 (a) TEM image of the HE+CR+1273K 1min+973K 12h sample crystal, (b) enlarged view of the S1 area in (a), and (c) dark field image of (b); (b) corresponding EDS results: (d) Co, (e) Cr, (f) Ni, (g) Al, (h) Ti, (i) W; (j) SAED result of the S2 area in (b); (k) high-resolution image of the W-rich phase, and (l) FFT result of the position shown in the S3 area in (k).

[0030] Fig.14 (a) TEM image of the grain boundary of the HE+CR+1273K 1min+973K 12h sample, (b) an enlarged view of the S4 area in (a), (c) SAED result of the S6 area in (b); (d) SAED result of the S5 area in (b); (b) corresponding EDS results: (e) Co, (f) Cr, (g) Ni, (h) Al, (i) Ti, (j) W.

[0031] Fig.15 (a) E of HE-973-12 and HE+CR+1273K 1min+973K 12h samples and samples in the literature pit Comparison of (b) 1 / E pit -1 / (E pit -E corr )curve.

[0032] Fig.16 (a) Nyquist diagram, (b) Bode impedance diagram, (c) Bode phase angle diagram, and (d) equivalent circuit diagram of As cast (cast sample), HE-973-12, and HE+CR+1273K 1min+973K 12h samples.

[0033] Fig.17SEM and corresponding EDS results of (a)(b)(c) As cast (cast specimen), (d)(e)(f) HE-973-12 and (g)(h)(i) HE+CR+1273K1min+973K 12h samples.

[0034] Fig.18 XPS fitting results of passivation films generated in 3.5wt.% NaCl solution for (a)-(e) As cast (cast specimen), (f)-(j) HE-973-12, and (k)-(o) HE+CR+1273K1min+973K12h samples.

[0035] Fig.19 (a) open circuit potential versus time curve, (b) Tafel curve, (c) comparison of corrosion current density and corrosion potential of the samples in the present invention with those of the alloys in the literature.

[0036] Fig. 20 SEM and corresponding EDS results of (a)(b)(c) As cast (cast specimen), (d)(e)(f) HE-973-12 and (g)(h)(i) HE+CR+1273K1min+973K 12h samples. DETAILED DESCRIPTION

[0037] Various exemplary embodiments of the present invention will now be described in detail. This detailed description should not be considered as limiting the present invention, but should be understood as a more detailed description of certain aspects, features, and embodiments of the present invention.

[0038] It should be understood that the terms described in the present invention are only for describing a particular embodiment and are not intended to limit the present invention. In addition, for the numerical range in the present invention, it should be understood that each intermediate value between the upper and lower limits of the scope is also specifically disclosed. The intermediate value in any stated value or stated range, and each smaller range between any other stated value or intermediate value in the described range is also included in the present invention. The upper and lower limits of these smaller ranges can be independently included or excluded in the scope.

[0039] Unless otherwise indicated, all technical and scientific terms used herein have the same meanings as those generally understood by those skilled in the art. Although the present invention describes only preferred methods and materials, any methods and materials similar or equivalent to those described herein may also be used in the implementation or testing of the present invention. All documents mentioned in this specification are incorporated by reference to disclose and describe the methods and / or materials associated with the documents. In the event of a conflict with any incorporated document, the content of this specification shall prevail.

[0040] It will be apparent to those skilled in the art that various modifications and variations may be made to the specific embodiments of the present invention description without departing from the scope or spirit of the present invention. Other embodiments derived from the present invention description will be apparent to the skilled artisan. The present invention description and examples are exemplary only.

[0041] The words “include,” “including,” “have,” “contain,” etc. used in this document are open-ended terms, meaning including but not limited to.

[0042] The first aspect of the present invention provides a medium entropy alloy having a chemical formula of Co 54 Cr 25 Ni 10 Al3Ti3W5(at.%).

[0043] A second aspect of the present invention provides a method for preparing the above-mentioned medium entropy alloy, comprising the following steps:

[0044] By Co 54 Cr 25 Ni 10 Al3Ti3W5 atomic percentage Co, Cr, Ni, Al, Ti, W metal single substances are mixed and vacuum arc melted to obtain button ingots (40g / piece);

[0045] Ten of the above button ingots were induction melted and then cast to obtain medium entropy alloy ingots (cylindrical Ø29.3×50 mm);

[0046] The medium-entropy alloy ingot is subjected to homogenization annealing, then heat-insulated, hot-extruded, and cooled to obtain the medium-entropy alloy.

[0047] In some embodiments of the present invention, the temperature of the vacuum arc melting is 2300-3800K, the time is 45s, and the number of times is 5 times; the temperature of the induction melting is 1700-2000K, and the time is 1min; the temperature of the homogenization annealing is 1473±1K, and the time is 24h; the temperature of the insulation is 1523±1K, and the time is 1h.

[0048] In some embodiments of the present invention, the extrusion speed of the hot extrusion is 22±5 mm / s, and the extrusion ratio is 11.88. The cooling is specifically air cooling to room temperature.

[0049] In some embodiments of the present invention, the cooling step further includes cold rolling and / or heat treatment.

[0050] In some embodiments of the present invention, the cold rolling parameters are set as follows: the roller diameter is 175 mm, the maximum rolling force is 66838 N, the rolling speed is 915.8 mm / s, the rolling temperature is room temperature (293 K), the reduction is 0.2 mm / pass, the rolling passes are 18-21 times according to the thickness of the prepared sample before rolling, the rolling thinning is 60%, and there is no lubrication during the rolling process. In some embodiments of the present invention, the heat treatment is treated at a temperature of 873 K-1073 K for 1-12 hours.

[0051] In some embodiments of the present invention, the heat treatment is first treated at 1273K for 1 min, and then treated at 873K-1073K for 1-12 hours.

[0052] After the heat treatment is completed, a water quenching cooling step is also included.

[0053] A third aspect of the present invention provides an application of the above-mentioned medium-entropy alloy in the preparation of structural materials for service in seawater environments or acidic environments.

[0054] A fourth aspect of the present invention provides a structural material, the raw materials of which include the above-mentioned medium-entropy alloy.

[0055] The technical solutions described in the present invention, unless otherwise specified, are all conventional solutions in the art, and the reagents or raw materials used, unless otherwise specified, are purchased from commercial channels or have been disclosed.

[0056] The technical solutions provided by the present invention are described in detail below in conjunction with the embodiments, but they should not be construed as limiting the protection scope of the present invention.

[0057] Example 1

[0058] 1. Experimental Materials and Procedures (Co 54 Cr 25 Ni 10 Preparation of Al3Ti3W5 (at.%) medium entropy alloy

[0059] Hot extrusion: The medium entropy alloy button ingots are obtained by arc melting of metal single substances of Co, Cr, Ni, Al, Ti, and W (purity>99.5wt.%). The oxide film on the surface of each raw material must be removed before preparation, and the sample is weighed using an electronic balance with an accuracy of 0.0002g, and the mass error is controlled within the range of ±0.002g. The weighed samples are ultrasonically cleaned with ethanol, air-dried, and then packaged in a vacuum bag to prevent oxidation. Button ingots are prepared by vacuum arc melting (the equipment used is a vacuum arc melting suction casting equipment, model WK-Ⅱ). During melting, the metal materials are arranged in order from high to low melting points from top to bottom to ensure that the metal materials are completely melted. In order to ensure that the elements of the alloy are fully fused, the arc is gently shaken during melting, and then the sample is turned over and melted again after cooling. Each button ingot is turned over and melted at least 5 times (the temperature of each melting is 3800±100K and the time is 45s). Afterwards, the button ingots were cast in an induction melting furnace (KZG-2, Kusite, induction melting temperature of 1800±100K, time of 1min) and copper mold to obtain cylindrical medium entropy alloy ingots (cast, Ф29.3×50mm). In order to ensure uniform composition, the obtained samples were homogenized and annealed at 1473K for 24h in an argon atmosphere, and then wrapped with a steel sleeve for the medium entropy alloy ingot. After being kept at 1523K for 1h, they were extruded using a four-column hydraulic press (HP32-315, Jiangsu Tiankai) at an extrusion speed of 22±5mm / s, and then air-cooled to room temperature to obtain a long rod with a diameter of 16mm and a length of about 600mm. The diameter of the medium entropy alloy part is about 8.5mm, and the extrusion ratio calculated is about 11.88.

[0060] Post-processing: Post-processing is divided into three types: heat treatment (i.e. annealing), cold rolling (i.e. rolling), and cold rolling + heat treatment; the equipment used for post-processing of materials includes annealing furnaces and cold rolling mills.

[0061] Before heat treatment, a vacuum tube sealing device is used to seal the cut samples in a glass tube in an argon atmosphere to prevent oxidation of the samples during heat treatment. The samples obtained by the above hot extrusion are named HE. For samples that are post-treated after hot extrusion, they are named according to their preparation process. For samples that are only heat treated, the naming rule is HE-heat treatment temperature-heat treatment time, that is: HE-873-1, HE-973-1, HE-1073-1, HE-1173-1, HE-973-4, HE-973-8, HE-973-12, HE-973-16 (for example, HE-873-1 means that the sample obtained by hot extrusion is directly heat treated, the heat treatment temperature is 873K, the heat treatment time is 1h; the heat treatment time unit is h). For samples that were only cold rolled, they were named HE+CR; for samples that were first cold rolled (the specific cold rolling process was: roller diameter was 175 mm, maximum rolling force was 66838 N, rolling speed was 915.8 mm / s, rolling temperature was room temperature (293 K), reduction was 0.2 mm / pass, and the number of rolling passes was 18-21 times according to the thickness of the prepared sample before rolling, cold rolled to 40% of the original thickness of the sample, and no lubrication was performed during the rolling process) and then heat treated, the naming rule was HE+CR+heat treatment process, i.e., HE+CR+973K 1h, HE+CR+1073K 1h, HE+CR+1173K 1h, HE+CR+1273K 1min, HE+CR+1173K 1h+973K 12h, HE+CR+1273K 1min+973K 12h (for example, HE+CR+973K 1h means that the sample obtained by hot extrusion is first cold rolled and then heat treated at a temperature of 973K for 1h).

[0062] 2. Annealing regulation of Co prepared by hot extrusion 54 Cr 25 Ni 10 Properties and structure of Al3Ti3W5 (i.e. post-treatment is only heat treatment)

[0063] 2.1 Effect of heat treatment on Co 54 Cr 25 Ni 10 Effect of mechanical properties of Al3Ti3W5 alloy

[0064] Figure 1 (a) shows the Vickers hardness of HE, HE-873-1, HE-973-1, HE-1073-1 and HE-1173-1 samples. The results show that the hardness of the alloy increases first and then decreases with the increase of annealing temperature. At an annealing temperature of 1073K, the hardness of the alloy increases to about 327HV, and as the annealing temperature further increases to 1173K, the hardness decreases to about 282HV. Figure 1 (b) and (c) show the engineering stress-strain curves and the strength and elongation of HE, HE-873-1, HE-973-1, HE-1073-1 and HE-1173-1 samples, respectively. The samples show good comprehensive mechanical properties, with high yield strength, ultimate tensile strength and good ductility. Similar to the change in hardness, with the increase of annealing temperature, the strength of the sample increases first and then decreases. After annealing at 973K for 1 hour, the sample has the best combination of strength and plasticity, with a yield strength of about 800MPa, a tensile strength of about 1279MPa and an elongation of about 46.6%. The research results show that the mechanical properties of the extrusion heat treated samples are related to the annealing temperature.

[0065] Since the HE-973-1 sample showed the best comprehensive mechanical properties, the annealing time dependence of the samples at 973 K was further studied. The annealing time was set to 8 h, 12 h, and 16 h. Figure 2 (a) shows the Vickers hardness of HE, HE-973-1, HE-973-8, HE-973-12 and HE-973-16 samples. As the holding time increases, the Vickers hardness reaches a maximum value of about 366HV at a holding time of 8 hours. The engineering stress-strain curve shows that annealing at 973K for 12 hours exhibits the best mechanical properties, with the yield strength and ultimate tensile strength continuing to increase to about 960MPa and about 1439MPa, while the ductility drops to about 36%. Overall, compared with the reported CoCrNi ternary alloy, the samples after hot extrusion heating treatment have better mechanical properties, and also show a good combination of strength and ductility in the alloy strengthened by precipitation phase ( Figure 2 (d)).

[0066] 2.2 Effect of heat treatment on Co 54 Cr 25 Ni 10 Effect of Al3Ti3W5 Alloy Microstructure

[0067] 2.2.1 XRD and SEM analysis

[0068] like Figure 3 As shown in (a), the XRD results of HE-873-1, HE-973-1, HE-1073-1 and HE-1173-1 samples show a single FCC structure. Similarly, the samples heat treated at 973K for different times also show a single FCC structure ( Figure 3 (b)).

[0069] In order to characterize the microstructural characteristics of the samples, SEM tests were also carried out on the heat treated samples ( Figure 4 and Figure 5).like Figure 4 As shown in (a)-(e), annealing significantly promotes precipitation at the grain boundaries, which becomes more serious at elevated temperatures. The precipitates show two different element enrichments: the larger black precipitates are rich in Al and Ti, while the smaller white precipitates are rich in W ( Figure 4 (f)-(l)).

[0070] Among the samples heat treated at different temperatures after extrusion, the sample annealed at 973K for 1 hour showed the best mechanical properties, which may be attributed to the precipitation of L12 phase. Phase diagram calculation shows that a large amount of L12 phase precipitates in the crystal at the intermediate temperature to provide second phase strengthening. On this basis, heat treatment experiments were carried out at 973K for different heat treatment times. Figure 5 The SEM images of HE-973-8, HE-973-12 and HE-973-16 samples are shown. The precipitated phase at the grain boundary does not change much with the increase of annealing time. The grain size fluctuates in the range of 9 to 13 μm (excluding twins). This degree of change does not bring about significant strength differences, but the content and size of the L12 phase will change with the annealing time. The fluctuation of strength should be attributed to the change in the size and amount of the L12 phase.

[0071] 2.2.2 TEM analysis of HE-973-12 sample

[0072] In order to further determine the microstructure of the precipitated phases within the grains and at the grain boundaries, TEM, HRTEM, SAED and EDS measurements were performed on the HE-973-12 sample ( Figure 6 and Figure 7 ). Figure 6 (a) shows the bright field TEM image of the crystal, and the enlarged dark field TEM image and corresponding SAED pattern of the R1 region are shown in Figure 1. Figure 6 In the crystal, a large number of spherical nano-precipitates were found in the FCC matrix, which were identified as ordered L12 phases (see Figure 6 The superlattice diffraction spots are marked by green dashed circles in (c). Figure 6 As shown in (d) and (e), annealing twins still exist inside the matrix. Figure 6(f) shows that there are a large number of stacking faults in the crystal. In Ni-based single crystal high-temperature alloys with low stacking fault energy, the deformation mechanism of the alloy changes with the increase of hot compression temperature. At 873K, the deformation mechanism changes from antiphase boundary shear to stacking fault shear; at 1073K, the deformation mechanism is mainly controlled by stacking fault shear, accompanied by the reaction of stacking faults and Lomer-Cottrell (LC) locks in different directions; above 1173K, although there are still some stacking faults, the main deformation mechanism is the bypass mechanism of dislocations. In the present invention, the sample to be hot extruded is first kept at a temperature of 1523K for one hour, and then the sample is taken out of the annealing furnace and hot extruded in the atmosphere. Due to the rapid cooling of the sample in air, this treatment can be regarded as a medium-temperature compression deformation, which will lead to the formation of a large number of stacking faults.

[0073] like Figure 7 As shown in the figure, the results of TEM, FFT and EDS at the grain boundary are shown. It can be seen that there are two element-enriched regions, one is Al-rich and Ti-rich region (Co2AlTi), and the other is W-rich region (Co3W). In the Co-Al-Ti-W system, discontinuous precipitation of Co3W and Co2AlTi phases at the grain boundary was found at 1023K.

[0074] The strengthening mechanism and deformation mechanism of the sample were analyzed. The deformation process of the HE-973-12 sample is as follows: Figure 8 As shown. Before deformation, the HE-973-12 sample contained precipitation phases within the grains and at the grain boundaries, and stacking faults existed before plastic deformation. At the beginning of deformation, due to the small stacking fault energy, a large number of stacking faults were widely generated, resulting in the formation of a large number of structurally stable LC locks. These locks act as obstacles to dislocation movement, resulting in a significant increase in the work hardening rate during the early plastic deformation process. In addition, LC locks can act as Frank-Reed dislocation sources, providing more dislocations to accommodate subsequent plastic deformation. This in turn causes more dislocations to accumulate and entangle around the LC locks, slowing down the rate of decrease in the later work hardening rate ( Figure 8 (a)). Since many nanoscale L12 particles exist on the dislocation slip path, it becomes more difficult for dislocations to move within the grain during tensile deformation until the shear stress is large enough for the dislocations to cut through the precipitated phase. When dislocation slip encounters large-sized particles, the dislocations gather around the grain boundaries ( Figure 8 (c) and (f)). As the plastic deformation progresses, the volume fraction of dislocations gradually increases ( Figure 8 Near fracture, the volume fraction of dislocations becomes higher, leading to rapid propagation of slip bands and the appearance of microcracks at grain boundaries ( Figure 8(e), (h)), and begin to connect with each other. Therefore, the fracture should occur along the grain boundaries and a few slip bands, resulting in a mixed transgranular-transgranular fracture mode ( Figure 8 (h)). In general, maintaining work hardening during plastic deformation can induce excellent tensile properties, which should be attributed to the synergistic activation of dislocations and stacking faults. The variation trend of work hardening rate in other samples is similar to that of HE-973-12 sample. Due to the different ratios among dislocation density, precipitate phase, grain boundary and stacking fault, the work hardening effect is affected, resulting in different plastic deformation behaviors.

[0075] This section studies the effect of heat treatment on the Co 54 Cr 25 Ni 10 The influence of entropy on the properties and microstructure of Al3Ti3W5 medium-entropy alloys. Orthogonal experiments on heat treatment temperature and time were carried out on the alloys, and the relationship between microstructure and mechanical properties, as well as the corresponding deformation mechanism, were studied. The following conclusions were drawn:

[0076] (1) Three types of precipitates were present in the hot extruded and subsequently heat treated samples, namely, L12 nanoparticles within the FCC grains, W-rich, Al-rich, and Ti-rich precipitates at the grain boundaries. Hot extrusion refines the grains and promotes the accumulation of stacking faults. The content of various precipitates can be adjusted by adjusting the heat treatment temperature and heat treatment time.

[0077] (2) The optimized heat treatment parameters were set at 973K for 12 hours, and the yield strength and tensile strength of the HE-973-12 sample were increased to about 960MPa and 1439MPa, respectively, accompanied by a relatively high plastic strain of about 36%. A good combination of high strength and high plasticity was achieved. Calculations of the strengthening mechanism show that the L12-type second phase strengthening mechanism provides the main contribution to the improvement of yield strength. The collective strengthening effects from L12 precipitation strengthening, grain boundary precipitation strengthening, grain boundary strengthening, and solid solution strengthening increase the yield strength of the HE-973-12 sample to about 960MPa.

[0078] (3) During deformation, a high density of stacking faults in different directions was activated in the HE-973-12 sample, leading to the formation of a stacking fault network and immovable LC locks, which provided an effective source of work hardening capacity, thereby achieving a significant strength-plasticity synergy. The L12 precipitates were sheared by dislocations, facilitating long-distance dislocation slip through the matrix. In contrast, dislocations were piled up at the interface between the grain boundary precipitates and the matrix, which increased the work hardening capacity in the early stages of plastic deformation but also led to stress concentration.

[0079] 3. Cold rolling and heat treatment to regulate the Co prepared by hot extrusion 54 Cr25 Ni 10 Properties and structure of Al3Ti3W5

[0080] 3.1 Effect of cold rolling and heat treatment on Co 54 Cr 25 Ni 10 Effect of mechanical properties of Al3Ti3W5 alloy

[0081] The samples prepared by hot extrusion were first rolled and thinned by 60% (the cold rolling parameters were set as follows: the roller diameter was 175 mm, the maximum rolling force was 66838 N, the rolling speed was 915.8 mm / s, the rolling temperature was room temperature (293 K), the reduction was 0.2 mm / pass, and the rolling passes were 18-21 times according to the thickness of the prepared samples before rolling, and there was no lubrication during the rolling process), and then heat treated (the temperature and time of the heat treatment refer to the previous text). The heat-treated samples were cooled by water quenching and then cut into dog-bone-shaped tensile pieces for quasi-static tensile tests at room temperature. The heat treatment process and the corresponding yield strength, tensile strength and elongation are shown in Table 1.

[0082] Table 1: Yield strength, tensile strength and elongation of samples with different treatment processes

[0083]

[0084] Note: “\” in the table means it cannot be stretched.

[0085] For samples that were heat treated in one step, the samples that were heat treated at 973K and 1073K for one hour were very brittle and could not be tested for tensile mechanical properties. This was because the heat treatment temperature was too low and the samples were not completely recrystallized. There were still a large number of slip bands left by rolling in the samples. This caused the samples to lose plasticity. After trying higher heat treatment temperatures and longer heat treatment times, it was found that although the plasticity of the alloy was improved, which should be attributed to the occurrence of recrystallization, the strength decreased significantly. Therefore, a heat treatment method of high-temperature heat treatment followed by low-temperature heat treatment was adopted. First, heat treatment was performed at high temperature for recovery and recrystallization, which was also helpful for the precipitation of the Co3W phase. Then, heat treatment at medium temperature promoted the precipitation of the precipitated phase, achieving the purpose of strengthening the second phase. The results show that the sample of HE+CR+1173K 1h+973K 12h cannot be stretched. The first step is annealing at 1173K for 1h. The heat treatment temperature is not enough, resulting in incomplete recrystallization. In the second step of annealing, the process of annealing at 973K for 12h causes a large amount of L12 phase precipitation, resulting in a sharp drop in plasticity. The sample of HE+CR+1273K 1min+973K12h showed a tensile strength of 1774MPa and an elongation of 10.8%. Compared with the HE-973-12 sample, the elongation decreased by 25%, but the yield strength and tensile strength increased by 464MPa and 335MPa, respectively.

[0086] like Fig. 9 The engineering stress-strain curves and hardness of the selected samples and the hot extruded samples (HE) are shown. 54 Cr 25 Ni 10 After rolling, the Al3Ti3W5 alloy showed a high strength of 2000MPa. However, there was an obvious softening phenomenon during processing. The dislocations in the sample after cold rolling reached a saturated state, causing the sample to lose the ability to process harden. After the rolled sample was heat treated at 1273K for 1min, the recrystallization was complete, the tensile strength dropped from 2000MPa of the rolled sample to 1675MPa, and the elongation increased to 22.3%. After 12 hours of precipitation heat treatment at 973K, the tensile strength was increased to 1774MPa, with an elongation of 10.8%.

[0087] 3.2 Effect of cold rolling and heat treatment on Co 54 Cr 25 Ni 10 Effect of Al3Ti3W5 Alloy Microstructure

[0088] In order to clarify the relationship between alloy microstructure and mechanical properties, it is necessary to characterize the alloy microstructure. First, analyze the phase composition of the sample, such as Fig.10Shown are the XRD test results of HE+CR, HE+CR+1273K 1min and HE+CR+1273K1min+973K 12h samples, indicating that the samples after cold rolling and annealing after cold rolling both exhibit a single-phase FCC structure.

[0089] like Fig.11 Shown is the SEM image of the sample. After rolling, the hot extruded sample showed densely interlaced slip bands. After annealing, the HE+CR+1273K 1min and HE+CR+1273K 1min+973K 12h samples showed dense white precipitates. The EDS results corresponding to the HE+CR+1273K 1min+973K 12h samples showed that these dense white precipitates were W-rich precipitates. Compared with the samples that were only annealed, the W-rich phase in the annealed samples after rolling not only precipitated at the grain boundaries, but also precipitated within the grains along the slip bands formed by rolling (such as Fig.10 (b) Indicated by the yellow dashed line). The high dislocation density and local stress concentration in the slip band can reduce the energy barrier for nucleation, thereby promoting the nucleation of the precipitate phase. In addition, the dislocation network may also provide diffusion channels, accelerate the migration and aggregation of atoms in the slip band region, and further promote the formation of the precipitate phase. Compared with the W-rich phase (299.1nm) in the HE-973-12 sample that was only annealed, the W-rich phase (106.652nm) in HE+CR+1273K1min+973K12h became smaller. This is because recrystallization and the precipitation of a large amount of precipitates occur simultaneously. The higher nucleation rate causes the precipitate phase to contact other boundaries (such as grains or another precipitate phase) very early during growth and stop growing.

[0090] like Fig.12The EBSD images of the samples are shown, showing the BC map, IPF map and GND map of the samples after rolling and annealing after rolling. The samples after rolling have dense slip bands, the grains are elongated and broken. At the same time, they have high dislocation density. Generally speaking, based on the evaluation of misorientation, the grain boundaries of 2°<θ<15° and θ≥15° are called low-angle grain boundaries (LAGB) and high-angle grain boundaries (HAGB), respectively. The annealed samples all showed very little low-angle grain boundary content (HE+CR+1273K 1min: 3.74%; HE+CR+1273K1min+973K 12h: 2.83%), which indicates that they have achieved complete recrystallization. The annealed samples showed a non-uniform structure, i.e., a structure of large grains plus small grains, and the average grain size was refined from ~13μm in the HE sample to 6.64μm in the HE+CR+1273K1min+973K12h sample. The high dislocation density generated during rolling provides stored energy for subsequent heat treatment, which becomes the driving force for the nucleation and growth of new grains at the recrystallization temperature. The energy provided by the high dislocation density helps to overcome the stability of the atomic arrangement inside the original grains and promote the formation of new grains. During the recrystallization process, a high nucleation rate will cause a large number of grains to grow simultaneously. Due to the presence of a large number of nucleation points, the new grains have already contacted each other before growing up and stopped growing, resulting in an overall reduction in grain size. At the same time, a large number of annealing twins appeared in the samples of HE+CR+1273K 1min and HE+CR+1273K1min+973K12h, and the proportion of twin boundaries was 30.8% and 36.8%.

[0091] In order to further characterize the microstructure of the HE+CR+1273K 1min+973K 12h sample, TEM, HRTEM, SAED and EDS tests were performed on it ( Fig.13 ).like Fig.13 As shown in (a), the sample after annealing shows a large number of annealing twins, and there are a large number of black submicron precipitates in the grains and grain boundaries. The S1 region inside the grain is selected for enlargement ( Fig.13 In (b), the region shows black precipitates and twin morphology. The corresponding EDS test results ( Fig.13 The black precipitated phase in (d)-(i) is W-rich phase, and the corresponding high-resolution image Fig.13 (k) and its Fourier transform result Fig.13 (l) shows that the phase is Co3W phase. This is consistent with the results obtained by SEM test ( Fig.11) phase, the Co3W phase is not only precipitated at the grain boundary, but also found in the grain. In addition, there are annealing twins and nano-L12 precipitates ( Fig.13 (c) and (j)).

[0092] The present invention also conducted TEM test on the grain boundary part of the HE+CR+1273K 1min+973K 12h sample, and the results are as follows: Fig.14 As shown in Figure 2, two precipitated phases appeared at the grain boundaries. The EDS results showed that they were W-rich phase and Al-rich Ti phase ( Fig.14 In (e)-(j)), the SAED results of the Al-rich and Ti-rich phase show that it is a Co2AlTi phase ( Fig.14 In (c), the SAED results of the W-rich phase and the diffraction spots of the L12 phase are superimposed, but it can still be distinguished as the Co3W phase ( Fig.14 In summary, the microstructural test results of the HE+CR+1273K 1min+973K 12h sample showed a microstructure including fine grains, annealing twins and precipitated phases (intragranular: nano-L12+submicron Co3W; grain boundary: submicron Co3W+Co2AlTi).

[0093] After analyzing various strengthening mechanisms and plastic deformation mechanisms, the present invention can summarize a series of events that occur during the deformation process of the sample. Before deformation, the HE+CR+1273K 1min+973K 12h sample contains precipitation phases in the crystal (L12 phase and Co3W) and at the grain boundary (Co3W and Co2AlTi), and dislocations exist before plastic deformation. At the beginning of deformation, due to the small stacking fault energy, a large number of stacking faults are widely generated to form a stacking fault network. The LC locks and the interaction between the stacking faults provide an efficient source of work hardening ability. They can fix dislocations and prevent dislocations from moving to grain boundaries or phase boundaries to cause catastrophic stress concentration. At the same time, they can act as Frank-Reed dislocation sources to generate new dislocations to adapt to subsequent deformation, resulting in a significant strength-ductility synergistic effect. Since many nano-scale L12 particles exist on the dislocation slip path, it becomes more difficult for dislocations to move within the crystal during tensile deformation until the shear stress is large enough for the dislocations to cut through the precipitation phase. When dislocation slip encounters submicron-sized precipitates Co3W and Co2AlTi, dislocations and stacking faults accumulate at the interface between the FCC matrix and the precipitates. When dislocation motion encounters grain boundaries or twin boundaries, they accumulate at the interface. Approaching fracture, the volume fraction of dislocations and stacking faults becomes higher, resulting in rapid propagation of slip bands, while some microcracks appear at grain boundaries and begin to connect with each other. In general, maintaining work hardening during plastic deformation can induce excellent tensile properties, which should be attributed to the synergistic activation of dislocations and stacking faults.

[0094] This section focuses on the effect of rolling heating treatment on the Co prepared by hot extrusion. 54 Cr 25 Ni 10 The influence of entropy on the properties and microstructure of Al3Ti3W5 medium-entropy alloy. The rolled alloy was heat treated with different processes, and the relationship between microstructure and mechanical properties, as well as the corresponding deformation mechanism, were studied. The following conclusions were drawn:

[0095] (1) The hot extruded samples were first rolled and then heat treated, which can effectively refine the grains and promote the precipitation of submicron Co3W phase. Finally, the preferred HE+CR+1273K 1min+973K 12h sample showed a microstructure containing fine grains, annealing twins and precipitation phases (intragranular: nanoscale L12+submicron Co3W; grain boundary: submicron Co3W+Co2AlTi).

[0096] (2) The optimized process parameters are annealing at 1273K for 1 min after rolling 60%, and then annealing at 973K for 12 h. Compared with the HE-973-12 sample, the elongation of the HE+CR+1273K 1 min+973K 12 h sample decreased by 25%, but the yield strength and tensile strength increased by 464 MPa and 335 MPa, respectively. A good combination of high strength and toughness was achieved. The calculation of the strengthening mechanism shows that the second phase strengthening and dislocation strengthening provide the main contribution to the improvement of yield strength.

[0097] (3) During the deformation process of the HE+CR+1273K 1min+973K 12h sample, a high density of stacking faults was activated in different directions, leading to the formation of stacking fault networks and immovable LC locks, which provided an effective source for work hardening ability, thereby achieving a significant strength-toughness synergy. The L12 precipitates were sheared by dislocations, which facilitated long-distance dislocation slip through the matrix. In contrast, dislocations were piled up at the interface between the submicron-sized Co2AlTi and Co3W precipitates and the matrix, which increased the work hardening ability in the early stages of plastic deformation, but also led to stress concentration.

[0098] 4. High strength and high toughness Co 54 Cr 25 Ni 10 Corrosion Resistance of Al3Ti3W5 Medium Entropy Alloy

[0099] 4.1 Electrochemical Corrosion Behavior in 3.5wt.%NaCl Solution

[0100] 4.1.1 Electrochemical analysis

[0101] In order to more accurately evaluate the corrosion resistance of the samples, a penalty function was set according to the evaluation index for evaluating the corrosion resistance of alloys in NaCl solution proposed by Nascimento et al.

[0102]

[0103] Where Z is the value of the penalty function, α1 and α2 are exchange constants whose values ​​are arbitrary and depend on the relative weights given to each performance indicator by the designer. pit and E pit -E corr ) have the same weight, so α1 and α2 are equal. pit and E corr The smaller the penalty function Z is, the better its corrosion resistance in NaCl solution is. Fig.15 As shown in (b), the horizontal and vertical axes are plotted as 1 / E pit and 1 / (E pit -E corr ), in which the closer the position of the alloy is to the lower left corner, the smaller the value of the penalty function Z is, and the better its corrosion resistance is. Therefore, As cast (cast sample), HE+973-12 and HE+CR+1273K 1min+973K 12h samples have good corrosion resistance in NaCl solution. In summary, Ascast, HE-973-12 and HE+CR+1273K 1min+973K 12h perform better in pitting potential, showing good corrosion resistance in NaCl solution.

[0104] In order to better analyze the corrosion behavior of the medium entropy alloy in the present invention in NaCl solution, electrochemical impedance spectroscopy analysis was performed on the sample to quantitatively analyze the passivation behavior of the alloy. Fig.16 Shown are the fitting results of the EIS data. Fig.16 (a) is the Nyquist plot. Generally speaking, the larger the arc radius of a capacitor, the better its corrosion resistance. The HE-973-12 sample shows the largest arc radius and the best corrosion resistance. Fig.16 (b) is the Bode impedance diagram, and it is found that the HE-973-12 sample exhibits the highest impedance at both high and low frequencies. Fig.16 (c) is the Bode phase angle diagram. The larger peak width shows that HE-973-12 and As cast samples have better corrosion resistance. Fig.16(d) is the equivalent circuit diagram used in fitting the EIS circuit. Since the HE+CR+1273K 1min+973K 12h curve in the Nyquist diagram shows a combination of arcs and approximate straight lines, different equivalent circuits are used to fit the three samples. For As cast and HE-973-12 samples, the following equivalent circuits are used: Fig.16 The equivalent circuit shown in (d) is shown in ①, and the equivalent circuit shown in ② is used for the HE+CR+1273K 1min+973K 12h sample. In the present invention, R s Represents the resistance of the electrolyte solution, which reflects the conductivity characteristics of the solution itself. c Refers to the resistance of the passivation layer, which indicates the ability of the passivation layer to hinder current transfer. ct is the charge transfer resistance, which represents the degree of hindrance to the charge transfer process on the electrode surface. c and Q dl are the phase angle elements of the capacitance of the passivation layer and the electric double layer, respectively, which describe the capacitance properties and response behaviors of the passivation layer and the electric double layer under the AC electric field. In order to obtain better fitting results, in the actual fitting process, a constant phase element (CPE) is used to adjust Q c and Q dl The specific parameters of CPE are simulated. The impedance of CPE is defined as follows:

[0105] Z CPE =[Y0(jω) n ] -1

[0106] Where ω is the angular frequency, Y0 is the pseudo capacitance of the CPE, j is the imaginary unit, and the value of n represents the degree to which the CPE deviates from the ideal capacitance characteristic and is between 0 and 1.

[0107] Table 2 Equivalent fitting circuit electrochemical parameters

[0108]

[0109] Use as Fig.16 The fitting results of the equivalent circuit shown in (d) are shown in Table 2. The variance χ of the fitting data is used 2 The accuracy of the fitting results was tested and it was found that 2The values ​​of are all small, and the fitting is good. Since the value of Rct depends on the Faraday process of oxidation or reduction reaction caused by electron transfer, it can indicate the speed of the corrosion process. The larger the Rct, the slower the corrosion reaction. The HE-973-12 sample shows a larger Rct than the HE+CR+1273K 1min+973K 12h sample, and the Rct value of the HE+CR+1273K 1min+973K 12h sample is greater than that of the As cast sample, which is consistent with the Tafel test. corr The results are consistent. Rc represents the resistance of the corrosion product layer. The Rc values ​​of the three samples are not high, indicating that the corrosion layer does not significantly hinder the current. However, the Rc of the HE+CR+1273K 1min+973K 12h sample is slightly higher, which may indicate that its corrosion layer is thicker or denser.

[0110] 4.1.2 Microstructure after corrosion

[0111] In corrosive environments including air, corrosion-resistant alloys can automatically form a protective or passivation film. This passivation film acts as an electronic barrier for charge transfer by isolating the alloy from direct contact with the surrounding environment, while preventing the diffusion of corrosive ions to the alloy surface, effectively slowing down the corrosion process. The specific composition and structure of the passivation film directly affect the corrosion resistance of the alloy. Based on this, the corrosion mechanism of the alloy can be deeply understood by analyzing the passivation film on the surface of the sample after corrosion. Therefore, SEM and EDS characterization were performed on the samples after Tafel testing. Fig.17 As shown in Figure 2, the three samples showed different degrees of corrosion morphology. Only a small part of the surface of the As cast sample showed a passivation film, and there were a large number of pitting pits ( Fig.17 As shown by the yellow arrow in (a), the passive film of the HE-973-12 sample becomes more, and the passive film of the HE+CR+1273K1min+973K 12h sample is almost entirely passive film. The passive films of the three samples show obvious enrichment of Al, Ti, and W.

[0112] In order to further characterize the characteristics of the passive film produced by the sample during the corrosion process, an XPS test was performed on it. The XPS test results were analyzed using Advantage software and its built-in database. After subtracting the Shirley baseline, the high-resolution XPS spectrum was decomposed and convoluted. Fig.18 The characteristic spectra of Co 2p, Cr 2p, Al 2p, Ti 2p, and W 4f and their fitting results are shown. Since the energy level corresponding to the Ni element could not be detected in the passivation film, it indicates that there is no Ni compound in the passivation film of these samples. The results show that the passivation films on the three samples are composed of compounds of elements such as Co, Cr, Al, Ti, and W.

[0113] 4.2Co 54 Cr 25 Ni 10 Electrochemical Corrosion Behavior of Al3Ti3W5 Medium Entropy Alloy in 0.5M H2SO4 Solution

[0114] 4.2.1 Electrochemical analysis

[0115] like Fig.19 Polarization curves of As cast, HE-973-12 and HE+12731min+97312h samples in 0.5M H2SO4 solution. Analysis of polarization curves Fig.19 The corrosion current density (i corr ), corrosion potential (E corr ) and pitting potential (E pit ), as shown in Table 3.

[0116] Table 3 Tafel curve electrochemical experimental results

[0117]

[0118] like Fig.19 As shown in (a), at room temperature, the As cast sample exhibits the most positive open circuit potential, which means it is least likely to start corrosion. As shown in Table 3, among the three samples, the maximum corrosion current density is 41.6uA / cm 2 , minimum is 22.06uA / cm 2 This is at a lower level compared to alloys of similar composition (e.g. Fig.19 As shown in (c), the maximum corrosion potential is -0.296 mV and the minimum is -0.307 mV, which is at a medium level, indicating that the alloy is not prone to corrosion in the H2SO4 solution environment.

[0119] 4.2.2 Microstructure after corrosion

[0120] In order to further explore the corrosion mechanism of the alloy in 0.5M H2SO4 solution, the present invention conducted SEM and EDS tests on three samples of As cast, HE-973-12 and HE+12731min+97312h that completed the Tafel test. Fig. 20 As shown in Figure 2, the three samples showed different degrees of corrosion morphology. The As cast sample had the shallowest corrosion, with only a small part of the sample surface showing a passivation film, but also had multiple pitting pits (such as Fig. 20 The passive film of the annealed sample becomes more, and the sample annealed after rolling is almost entirely passive film. The composition of the passive film of the three samples is the same, which is mainly composed of Al, Ti, and W oxides.

[0121] This section studies the 54 Cr 25 Ni 10 The corrosion resistance of Al3Ti3W5 medium entropy alloy, especially the performance of samples under different heat treatment processes. Through the electrochemical corrosion behavior study in 3.5wt.% NaCl and 0.5M H2SO4 solution, the present invention found that the cast sample (As cast), the annealed sample only (HE-973-12) and the rolled and heat treated sample (HE+CR+1273K1min+973K 12h) all showed excellent corrosion resistance.

[0122] (1) These samples showed high pitting potential in 3.5wt.% NaCl solution. In 0.5M H2SO4 solution, they showed low corrosion current density and medium corrosion potential. This shows that the alloy has good corrosion resistance in NaCl solution and sulfuric acid environment. By comparing the corrosion performance and mechanical properties of the alloy, it can be found that the HE-973-12 and HE+CR+1273K 1min+973K 12h samples obtained in the present invention have both good mechanical properties and corrosion resistance.

[0123] (2) By analyzing the microstructure after corrosion, we further confirmed the protective effect of the passive film and the key role of alloying elements in corrosion resistance. In particular, the enrichment of Cr, Ti and W helps to form a more stable passive film, thereby providing better protection. At the same time, the microstructure of the alloy, such as grain size, dislocation density and precipitation phase, also has a significant effect on its corrosion resistance.

[0124] In summary, the present invention designs a composition of Co 54 Cr 25 Ni 10 The medium entropy alloy of Al3Ti3W5 is prepared by hot extrusion, and then the microstructure of the alloy is adjusted by heat treatment or rolling and heat treatment process, including grain size, precipitation phase, defects, etc. Through these means, the alloy has excellent mechanical properties while maintaining good corrosion resistance. The main conclusions are as follows:

[0125] (1) For the hot extruded samples that were heat treated. Three types of precipitation phases were presented in the hot extruded and subsequently annealed samples, namely, the coherent L12 nano-precipitates within the FCC grains, and the W-rich and Al- and Ti-rich precipitation phases at the grain boundaries. Hot extrusion refines the grains and promotes the accumulation of stacking faults. The content of various precipitation phases can be adjusted by adjusting the annealing temperature and annealing time. The yield strength and tensile strength of the optimal HE-973-12 sample increased to 960 MPa and 1439 MPa, respectively, accompanied by a relatively high plastic strain of about 36%. A good combination of high strength and plasticity was achieved. L12 second phase strengthening provided the main contribution to the increase in yield strength. During the deformation process, the HE-973-12 sample activated a high density of stacking faults in different directions, resulting in the formation of stacking fault networks and immovable LC locks, which provided an effective source of work hardening ability, thereby achieving a significant strength-toughness synergistic effect. The L12 precipitates were sheared by dislocations, which facilitated long-distance dislocation slip through the matrix. Instead, dislocations are accumulated at the interface between the grain boundary precipitates and the matrix, which increases the work hardening capacity in the early stage of plastic deformation but also leads to stress concentration.

[0126] (2) The post-rolling heat treatment process can further strengthen the Co 54 Cr 25 Ni 10 Mechanical properties of Al3Ti3W5 medium entropy alloy. The results show that the process can effectively refine the grains and promote the precipitation of submicron Co3W phase. Ultimately, the preferred HE+CR+1273K 1min+973K 12h sample exhibits a microstructure containing fine grains, annealing twins and precipitated phases (intracrystalline: nanoscale L12 phase + submicron Co3W phase; grain boundary: submicron Co3W phase + Co2AlTi phase). Compared with the HE-973-12 sample, the elongation of the HE+CR+1273K1min+973K 12h sample decreased by 25%, but the yield strength and tensile strength increased by 464MPa and 335MPa, respectively. Calculations of the strengthening mechanism show that second phase strengthening and dislocation strengthening provide the main contribution to the improvement of yield strength. During the deformation process of the HE+CR+1273K 1min+973K 12h sample, high-density stacking faults were activated in different directions, ensuring that the alloy maintained a certain elongation under high strength.

[0127] (3) The samples under different treatment processes all showed corrosion resistance. Through the study of electrochemical corrosion behavior in 3.5wt.% NaCl and 0.5MH2SO4 solutions, the present invention found that the cast sample (As cast), the sample annealed only after hot extrusion (HE-973-12), and the sample rolled and heat treated after hot extrusion (HE+CR+1273K 1min+97312h) all showed excellent corrosion resistance. In NaCl solution, a higher pitting potential was shown, showing good pitting resistance. In 0.5MH2SO4 solution, the low corrosion current density and moderate corrosion potential of the sample indicate that the alloy has good corrosion resistance in a sulfuric acid environment. By analyzing the microstructure after corrosion, the present invention further confirmed the protective effect of the passivation film and the key role of alloying elements in corrosion resistance. In particular, the oxides and hydroxides of Cr, Ti and W help to form a more stable passivation film, thereby providing better protection. At the same time, the microstructure of the alloy, such as grain size, dislocation density and precipitation phase, also has a significant impact on its corrosion resistance.

[0128] The embodiments described above are only descriptions of the preferred modes of the present invention, and are not intended to limit the scope of the present invention. Without departing from the design spirit of the present invention, various modifications and improvements made to the technical solutions of the present invention by ordinary technicians in this field should all fall within the protection scope determined by the claims of the present invention.

Claims

1. A medium entropy alloy, characterized in that: The chemical formula is Co 54 Cr 25 Ni 10 Al3Ti3W5.

2. A method for preparing the medium entropy alloy according to claim 1, characterized in that: The following steps are involved: By Co 54 Cr 25 Ni 10 Al3Ti3W5 atomic percentage Co, Cr, Ni, Al, Ti, W metal single substances are mixed and vacuum arc melted to obtain button ingots; The button ingot is subjected to induction melting, followed by casting to obtain a medium entropy alloy ingot; The medium-entropy alloy ingot is subjected to homogenization annealing, then heat-insulated, hot-extruded, and cooled to obtain the medium-entropy alloy.

3. The method for preparing a medium entropy alloy according to claim 2, characterized in that: The temperature of the vacuum arc melting is 2300-3800K, the time is 45s, and the number of times is 5 times; the temperature of the induction melting is 1700-2000K, and the time is 1min; the temperature of the homogenization annealing is 1473±1K, and the time is 24h; the temperature of the insulation is 1523±1K, and the time is 1h.

4. The method for preparing a medium entropy alloy according to claim 2, characterized in that: The extrusion speed of the hot extrusion is 22±5 mm / s, and the extrusion ratio is 11.

88.

5. The method for preparing a medium entropy alloy according to claim 2, characterized in that: After the cooling is completed, the steps of cold rolling and / or heat treatment are also included.

6. The method for preparing a medium entropy alloy according to claim 2, characterized in that: The cold rolling parameters are set as follows: roll diameter is 175 mm, maximum rolling force is 66838 N, rolling speed is 915.8 mm / s, rolling temperature is room temperature, reduction is 0.2 mm / pass, rolling passes are 18-21 times, and rolling reduction is 60%.

7. The method for preparing a medium entropy alloy according to claim 2, characterized in that: The heat treatment is performed at a temperature of 873K-1073K for 1-12 hours.

8. The method for preparing a medium entropy alloy according to claim 2, characterized in that: The heat treatment is firstly treated at 1273K for 1 minute, and then treated at 873K-1073K for 1-12 hours.

9. Use of the medium-entropy alloy as claimed in claim 1 in preparing structural materials for service in seawater environment or acidic environment.

10. A structural material, characterized in that: The raw material includes the medium-entropy alloy described in claim 1.

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