Low-fault-energy high-entropy alloy and preparation method thereof

Abstract

The invention discloses a preparation method of a low-fault-energy high-entropy alloy, which comprises the following steps: (1) weighing spherical CoCrNiAl alloy, Co powder, Cr powder, Ni powder and Ti powder according to a stoichiometric equation AlaTibNicCodCre, (2) putting the weighed powder into a ball milling tank for ball milling, uniformly mixing the powder, and then drying, sieving and drying for storage; (3) 304L stainless steel is selected as a base material, and a sand mill is used for polishing the surface of the base material to remove surface oxide skin until a bright surface is exposed; and (4) a laser directional energy deposition method is adopted, the dried powder is placed in a powder feeding tank in a laser system, high-purity Ar gas is used for feeding the powder to a round spot laser head for melting and stacking, and layer-by-layer continuous deposition is conducted. The generation and expansion of cracks in laser additive manufacturing can be inhibited by reducing the stacking fault of a high-entropy alloy system, a large number of coherent L12 phases can be separated out in situ, and the excellent mechanical property is achieved at the normal temperature.

Description

Technical Field

[0001] This invention relates to the field of high-entropy alloy technology, specifically to a low-stacking-fault-energy high-entropy alloy and its preparation method. Background Technology

[0002] Laser additive manufacturing technology, with its unique advantages in manufacturing complex geometries and achieving ultrafine grain microstructures, is leading a revolution in the field of metal processing. Meanwhile, multi-principal element high-entropy alloys, as a new type of material breaking through traditional design concepts, have shown great potential in extreme environment applications such as aerospace and energy chemicals due to their excellent strength, corrosion resistance, and high / low temperature performance. Applying laser additive manufacturing technology to the forming of high-entropy alloys can theoretically combine the advantages of both, opening up a new path for manufacturing high-performance, structurally integrated key components. However, this ideal path faces a severe common challenge: microcracks are easily formed during additive manufacturing, seriously impairing the integrity, reliability, and mechanical properties of components. The crack problem has become one of the core bottlenecks restricting the engineering application of laser additive manufacturing of high-entropy alloys.

[0003] Laser additive manufacturing is essentially an extremely non-equilibrium process of rapid melting and solidification combined with complex thermal cycling. The rapid scanning of the laser beam causes the material to undergo intense and localized heating and cooling, generating extremely high temperature gradients and cooling rates (up to 10⁻⁶). 3 -10 8 (K / s). Uncoordinated thermal contraction (thermal stress) and phase transformation volume changes (phase transformation stress) between different regions of the material are "frozen" inside the component, forming significant residual tensile stress. When the local stress exceeds the material's strength limit or crack propagation resistance, cracks initiate and propagate. Furthermore, non-equilibrium solidification can create microstructural features such as fine cellular structures, high dislocation density, and elemental segregation. While these structures may provide strengthening effects, they can also become stress concentration points or brittle pathways, especially along columnar grain boundaries or molten pool boundaries, promoting the formation of hot cracks (solidification cracks) or solid-state cracks.

[0004] Studies have shown that many high-entropy alloy systems, including classic FeCoCrNi systems and refractory high-entropy alloys, exhibit widespread crack susceptibility during LPBF (Laser-to-Break) processes, with the number far exceeding publicly reported levels. These cracks not only reduce the compactness and mechanical properties of parts but also pose a fatal threat to their fatigue life and dynamic reliability. To address this challenge, researchers have proposed solutions from multiple perspectives, primarily including optimizing heat input and temperature field distribution by adjusting laser power, scanning speed, and scanning strategy to reduce residual stress. However, this method has limited effectiveness for materials with high intrinsic crack susceptibility and has a narrow optimization window, lacking broad applicability. Developing new alloys specifically suited for additive manufacturing conditions is another approach. For example, adding trace amounts of carbon (C) to strengthen grain boundaries and improve grain boundary bonding can suppress intergranular cracking; or utilizing the thermal cycling of additive manufacturing to induce second-phase dissolution can alter the matrix composition, reducing stacking fault energy and promoting twinning plasticity. In recent years, a strategy based on the physical nature of material deformation has received significant attention: controlling the stacking fault energy of alloys. Stacking fault energy is the energy required to generate stacking faults per unit area in a crystal. It profoundly affects plastic deformation mechanisms such as dislocation slip, cross-slip, twinning, and martensitic phase transformation.

[0005] In traditional high stacking fault energy (SFE) materials, dislocations readily undergo cross-slip and climb, forming entanglements or cellular structures. Under the thermal stress of additive manufacturing, stress tends to concentrate in localized areas such as the cell walls of these dislocations, ultimately releasing energy through the initiation of microcracks, leading to cracking. Conversely, reducing stacking fault energy can encourage materials to preferentially coordinate strain during deformation (including deformation caused by thermal stress) through plane slip, stacking fault formation, mechanical twinning, or the induction of ε-martensitic phase transformation. This principle provides a novel approach to crack suppression in additive manufacturing: by actively reducing SFE through alloy design, the strain energy generated by thermal cycling during printing is directed to form numerous nanoscale planar crystal defects (such as stacking faults and twins), rather than being used for crack initiation and propagation. This is equivalent to providing an additional, harmless dissipation channel for residual stress, thereby suppressing cracking at its source. The effectiveness of this strategy has been preliminarily verified. For example, a research team from Central South University and South China University of Technology successfully reduced the solid-state fracture (SFE) of an equiatomic FeCoCrNi high-entropy alloy to a lower level by introducing approximately 2.4 at.% Al. Comparative experiments showed that the original alloy without Al exhibited significant microcracks after LPBF, while the Al-doped alloy... 0.1 The CoCrFeNi alloy achieved completely crack-free forming. More importantly, while the tensile strength was slightly improved, the elongation after fracture was significantly increased by about 55%, achieving an excellent synergy between strength and plasticity. Research from Northwestern Polytechnical University also shows that local deformation behavior can be controlled through segregation and dislocation self-organization in the additive manufacturing process, achieving a combination of high dislocation density and high plasticity. Summary of the Invention

[0006] To address the aforementioned technical problems, the first objective of this invention is to provide a method for preparing a low-stack-fault-energy high-entropy alloy, and the second objective is to provide the low-stack-fault-energy high-entropy alloy prepared thereby. By reducing the stacking fault energy of the high-entropy alloy system, the initiation and propagation of cracks in laser additive manufacturing are suppressed, and a large amount of coherent L12 phase can be precipitated in situ, exhibiting excellent mechanical properties at room temperature.

[0007] To achieve the first objective mentioned above, the present invention provides the following technical solution: a method for preparing a low-stacking-fault energy high-entropy alloy, characterized by the following steps:

[0008] (1) According to the stoichiometric formula Al a Ti b Ni c Co d Cr e Weigh out spherical CoCrNiAl alloy, Co, Cr, Ni, and Ti powders, where a and b are both 6.25 (at.%), c is 25-50 (at.%), d is 25-50 (at.%), and e is 12.5 (at.%).

[0009] (2) Place the weighed powder into a ball mill jar for ball milling to make the powder mix evenly, then dry, sieve, and store in a dry place;

[0010] (3) Select 304L stainless steel as the base material, and use a sand mill to grind the surface of the base material to remove the surface oxide scale until a bright surface is exposed;

[0011] (4) Using laser directional energy deposition, dry powder is placed in the powder feeding tank of the laser system, and high-purity Ar gas is used to send the powder to the circular spot laser head for melting and deposition, and the deposition is carried out layer by layer.

[0012] Preferred: The stoichiometric formula is: Al 6.25 Ti 6.25 Ni 50 Co 25 Cr 12.5 Or Al 6.25 Ti 6.25 Ni 37.5 Co 37.5 Cr 12.5 Or Al 6.25 Ti 6.25 Ni 25 Co 50 Cr 12.5 .

[0013] This invention reduces the stacking fault energy of the alloy system by adjusting the Ni and Co contents, thereby suppressing cracks induced by high Al+Ti content during laser additive manufacturing. Simultaneously, the alloy composition can precipitate a large amount of L12 phase without heat treatment, exhibiting high strength and excellent plasticity at room temperature. This solves the cracking problem of high-entropy alloys with high L12 phase content in laser additive manufacturing.

[0014] In the above scheme, the CoCrNiAl, Ti, Ni, Co, and Cr powders have a particle size of 45-105 μm and a purity of ≥99.5%. To prevent the volatilization of low-melting-point Al, this invention uses Al-containing spherical alloy CoCrNiAl powder for formulation.

[0015] In the above scheme: ball milling parameters: ball-to-material ratio is 4:1, ball milling speed is 200-300 rpm / min, forward rotation for 1 hour, reverse rotation for 1 hour, with an interval of 300 seconds in between. After ball milling, the mixed powder is dried in a vacuum drying oven to remove surface moisture.

[0016] In the above scheme: sieve through a 100-200 mesh sieve.

[0017] In the above scheme: in step (3), after polishing, the oil stains on the surface of 304L stainless steel are cleaned with alcohol.

[0018] In the above scheme, the laser model is RC-LMS-6000-R fiber laser.

[0019] In the above scheme: laser energy density 18.75-35 J / mm 2 .

[0020] In the above scheme, the laser parameters are: laser power P=950 W, scanning rate v=780 mm / min, Z-axis lifting of 0.3 mm per layer, circular spot diameter of 3 mm, and powder filling speed of 2 r / min.

[0021] A method for preparing a low stacking fault energy high entropy alloy, as described above, yields a low stacking fault energy high entropy alloy.

[0022] Beneficial effects:

[0023] This invention reduces the stacking fault energy of the alloy, releases residual stress caused by rapid heating and cooling during the fabrication process through dislocation proliferation and movement, and inhibits crack propagation at grain boundaries by reducing the size of the coherent L12 phase. Because this alloy system has a high Al+Ti content, the L12 phase can precipitate in situ during the printing process, exhibiting excellent strength and ductility at room temperature without the need for subsequent heat treatment. Attached Figure Description

[0024] Figure 1 shows the powder morphology of the three alloys in Example 1 after ball milling.

[0025] Figure 2 An exploratory diagram of laser-directed energy deposition technology.

[0026] Figure 3 The macroscopic morphology of the three alloys in Example 1 is shown.

[0027] Figure 4 Here is the XRD image of the alloy.

[0028] Figure 5 This refers to the microstructure of the alloy.

[0029] Figure 6 This represents the elemental distribution in the alloy.

[0030] Figure 7 These are dislocations and stacking faults in the alloy.

[0031] Figure 8 This is the engineering stress-strain curve of the alloy.

[0032] Figure 9 The images show the fracture surfaces and side views of the three alloys. Detailed Implementation

[0033] The present invention will now be further described with reference to the accompanying drawings and embodiments.

[0034] Example 1

[0035] (1) Powder selection: according to Al 6.25 Ti 6.25 Ni 50 Co 25 Cr 12.5 Al 6.25 Ti 6.25 Ni 37.5 Co 37.5 Cr 12.5 Al 6.25 Ti 6.25 Ni 25 Co 50 Cr 12.5 (at.%) Weigh out CoCrNiAl alloy, Ti, Ni, Co, and Cr powders. All powders are spherical, with a size of 45-105 μm and a purity of ≥99.5%. Since Al is easily volatilized during the preparation process, spherical CoCrNiAl powder was chosen for the proportioning. A FA224 electronic balance with an accuracy of 0.0001 g was used for weighing.

[0036] (2) Powder ball milling: The prepared powder is placed in a stainless steel ball mill jar and ball milled in a planetary ball mill. The ball milling parameters are: ball-to-powder ratio of 4:1, ball milling speed of 200-300 rpm / min (240 rpm / min is selected in this embodiment), ball milling for 1 hour in the forward direction and 1 hour in the reverse direction, with a 300-second pause between the forward and reverse directions to ensure that the powder is fully mixed and uniform. After ball milling, the powder is passed through a 100-200 sieve using a stainless steel sieve and then placed in a vacuum drying oven at 60℃ for 12 hours for drying and later use.

[0037] (3) Substrate selection: Magnetic 304L stainless steel with dimensions of 100mm in length, 100mm in width, and 10mm in height was selected as the substrate. The substrate has a very small dilution effect on the high-entropy alloy components during the laser directional energy deposition process. The oxide scale on the surface of the substrate was removed using a grinding wheel to prevent the formation of holes in the lower part of the sample during the deposition process. The surface oil was cleaned with alcohol and dried with a hair dryer for later use.

[0038] (4) Sample preparation: Laser directed energy deposition (DED) was used. The dried powder was placed in the powder feeding container of the laser powder feeding system. High-purity argon gas (99.999%) was used to transport the powder onto the substrate. Simultaneously, the laser beam was focused to form a small molten pool, which was then deposited gradually into a wall-like sample via a zigzag path. High-purity argon gas was used throughout the deposition process to prevent oxidation of the sample. Laser power P = 950 W, scanning rate v = 780 mm / min, Z-axis lift per layer was 0.3 mm, spot diameter was 3 mm, and powder filling speed was 2 r / min. Laser energy density was 24.36 J / mm². 2 The deposited thin wall is 80 mm wide, 2 mm thick, and 48 mm high.

[0039] Figure 1 The powder morphology and elemental distribution of the alloy after ball milling are shown. (a) Al 6.25 Ti 6.25 Ni 50 Co 25 Cr 12.5 (b)Al 6.25 Ti 6.25 Ni 37.5 Co 37.5 Cr 12.5 (c) Al 6.25 Ti 6.25 Ni 25 Co 50 Cr 12.5 Spherical powder has a significant impact on the surface morphology of laser additive manufacturing. (The last sentence appears to be incomplete and possibly refers to Al.) 6.25 Ti 6.25 Ni 50 Co 25 Cr 12.5Al 6.25 Ti 6.25 Ni 37.5 Co 37.5 Cr 12.5 Al 6.25 Ti 6.25 Ni 25 Co 50 Cr 12.5 The three alloys were weighed according to their atomic ratios and then mixed thoroughly in a ball mill. The elemental distribution diagram shows that no elements are clustered; the elements are distributed randomly.

[0040] Laser energy density has a significant impact on the porosity of alloys. Laser energy densities of 18.75 J / mm² were used. 2 19.79 J / mm 2 20.83 J / mm 2 21.86 J / mm 2 23.08 J / mm 2 24.36 J / mm 2 25.64 J / mm 2 26.92 J / mm 2 30 J / mm 2 31.67 J / mm 2 33.33 J / mm 2 35 J / mm 2 Conduct experiments. For example... Figure 2 As shown, 24.36 J / mm 2 The lowest porosity was observed at the specified energy density, at 0.037%. Therefore, Example 1 used 24.36 J / mm². 2 Energy density is used in additive manufacturing.

[0041] from Figure 3 The macroscopic morphology diagram shows that (a) Al 6.25 Ti 6.25 Ni 50 Co 25 Cr 12.5 (b)Al 6.25 Ti 6.25 Ni 37.5 Co 37.5 Cr 12.5 (c) Al 6.25 Ti 6.25 Ni 25 Co 50 Cr 12.5 Al 6.25 Ti 6.25 Ni 50 Co25 Cr 12.5 The alloy exhibits numerous macroscopic cracks that extend along the structural direction, and the number of cracks is high. Al 6.25 Ti 6.25 Ni 37.5 Co 37.5 Cr 12.5 Macroscopic cracks are present at the bottom of the alloy, few in number and short in length. Al 6.25 Ti 6.25 Ni 25 Co 50 Cr 12.5 The alloy surface is free of cracks and has a good morphology.

[0042] from Figure 4 The XRD images of the alloys show that the same diffraction peaks are present in all three alloys of Example 1, indicating that the phases in the alloys are the same, all being face-centered cubic structures. The diffraction peaks in the (111) direction have the highest intensity and gradually shift to the left, with the larger Co atoms leading to an increase in interplanar spacing.

[0043] from Figure 5 The metallographic microstructure of the three alloys shows that:

[0044] (a) Al 6.25 Ti 6.25 Ni 50 Co 25 Cr 12.5 ;

[0045] (b) Al 6.25 Ti 6.25 Ni 37.5 Co 37.5 Cr 12.5 ;

[0046] (c) Al 6.25 Ti 6.25 Ni 25 Co 50 Cr 12.5 Al 6.25 Ti 6.25 Ni 50 Co 25 Cr 12.5 The alloy contains cracks that penetrate the deposited layer, and microcracks appear at the boundaries of dendrites in different directions along the long axis of the dendrites. 6.25 Ti 6.25 Ni 37.5 Co 37.5 Cr 12.5 Numerous microcracks exist, and they are present in the intergranular region. Al 6.25 Ti 6.25 Ni25 Co 50 Cr 12.5 The alloy is crack-free, and its microstructure is dendritic.

[0047] Figure 6 Chinese: (a) Al 6.25 Ti 6.25 Ni 50 Co 25 Cr 12.5 (b) Al 6.25 Ti 6.25 Ni 37.5 Co 37.5 Cr 12.5 (c)Al 6.25 Ti 6.25 Ni 25 Co 50 Cr 12.5 The elemental distribution diagrams of the alloys show that all three alloys contain a large number of nano-precipitates. The second phase is spherical and rich in Ni, Al, and Ti elements, surrounded by a network structure rich in Co and Cr. As Co replaces Ni, the size of the second phase gradually decreases.

[0048] Figure 7 For dislocations and stacking faults in laser additive manufacturing of printed alloys, including:

[0049] (a) Al 6.25 Ti 6.25 Ni 50 Co 25 Cr 12.5 ;

[0050] (b) Al 6.25 Ti 6.25 Ni 37.5 Co 37.5 Cr 12.5 ;

[0051] (c) Al 6.25 Ti 6.25 Ni 25 Co 50 Cr 12.5 All three alloys contain a large number of dislocations. As the Co content increases and the Ni content decreases, the number of stacking faults gradually increases. The alloys with low stacking fault energy release residual stress generated during the printing process through dislocation propagation and movement, thus avoiding alloy cracking.

[0052] Figure 8 The figures show the engineering stress-strain curves of the alloy. Due to the presence of a large amount of coherent L12 phase, the printed alloys all exhibit high yield strength and tensile strength. Al 6.25Ti 6.25 Ni 50 Co 25 Cr 12.5 The alloy exhibits the highest yield strength and tensile strength, reaching 842.5 MPa and 1157.2 MPa respectively, and its plasticity reaches 20.6%. Al 6.25 Ti 6.25 Ni 37.5 Co 37.5 Cr 12.5 The alloy exhibits a yield strength of 686.4 MPa and a tensile strength of 933.6 MPa, with a plasticity of 28.1%. Al 6.25 Ti 6.25 Ni 25 Co 50 Cr 12.5 The alloy has a yield strength of 573.2 MPa and a tensile strength of 908.7 MPa, with a plasticity as high as 34.4%.

[0053] Figure 9 The images show the fracture surface and fracture side morphology of three alloys, (a) Al 6.25 Ti 6.25 Ni 50 Co 25 Cr 12.5 (b)Al 6.25 Ti 6.25 Ni 37.5 Co 37.5 Cr 12.5 (c) Al 6.25 Ti 6.25 Ni 25 Co 50 Cr 12.5 All three alloys exhibited numerous dimples on their fracture surfaces, indicating that the fracture mode was ductile fracture. Al 6.25 Ti 6.25 Ni 25 Co 50 Cr 12.5 The large and deep dimples in the alloy indicate that its toughness is higher than that of the other two alloys. The fracture morphology shows that there are a large number of slip bands during the tensile process of the alloy.

[0054] Although embodiments of the invention have been shown and described, it will be understood by those skilled in the art that various changes, modifications, substitutions and alterations can be made to these embodiments without departing from the principles and spirit of the invention, the scope of which is defined by the appended claims and their equivalents.

Claims

1. A method for preparing a low-stacking-fault energy high-entropy alloy, characterized in that, Prepare according to the following steps: (1) According to the stoichiometric formula Al a Ti b Ni c Co d Cr e Weigh out spherical CoCrNiAl alloy, Co, Cr, Ni, and Ti powders, where a and b are both 6.25 (at.%), c is 25-50 (at.%), d is 25-50 (at.%), and e is 12.5 (at.%). (2) Place the weighed powder into a ball mill jar for ball milling to make the powder mix evenly, then dry, sieve, and store in a dry place; (3) Select 304L stainless steel as the base material, and polish the surface of the base material to remove the surface oxide scale until the bright surface is exposed; (4) Using laser directional energy deposition, dry powder is placed in the powder feeding tank of the laser system, and high-purity Ar gas is used to send the powder to the circular spot laser head for melting and deposition, and the deposition is carried out layer by layer.

2. The method for preparing the low-stacking-energy high-entropy alloy according to claim 1, characterized in that: The stoichiometric formula is: Al 6.25 Ti 6.25 Ni 50 Co 25 Cr 12.5 Or Al 6.25 Ti 6.25 Ni 37.5 Co 37.5 Cr 12.5 Or Al 6.25 Ti 6.25 Ni 25 Co 50 Cr 12.5 .

3. The method for preparing the low-stacking-energy high-entropy alloy according to claim 1, characterized in that: The CoCrNiAl, Ti, Ni, Co, and Cr powders have a particle size of 45-105 μm and a purity of ≥99.5%.

4. The method for preparing the low stacking fault energy high entropy alloy according to any one of claims 1-3, characterized in that: Ball milling parameters: ball-to-material ratio of 4:1, ball milling speed of 200-300 rpm / min, forward rotation for 1 hour, reverse rotation for 1 hour, with an interval of 300 seconds in between. After ball milling, the mixed powder is dried in a vacuum drying oven to remove surface moisture.

5. The method for preparing the low stacking fault energy high entropy alloy according to claim 4, characterized in that: Sieve through a 100-200 mesh sieve.

6. The method for preparing the low stacking fault energy high entropy alloy according to claim 5, characterized in that: In step (3), after polishing, the oil stains on the surface of 304L stainless steel are cleaned with alcohol.

7. The method for preparing the low stacking fault energy high entropy alloy according to claim 6, characterized in that: The laser used is an RC-LMS-6000-R fiber laser.

8. The method for preparing the low stacking fault energy high entropy alloy according to claim 7, characterized in that: Laser energy density 18.75-35 J / mm 2 .

9. The method for preparing the low stacking fault energy high entropy alloy according to claim 8, characterized in that: The laser parameters are as follows: laser power P=950 W, scanning rate v=780 mm / min, Z-axis lifting of 0.3 mm per layer, circular spot diameter of 3 mm, and powder filling speed of 2 r / min.

10. A low-stack fault energy high-entropy alloy prepared by the preparation method of any one of claims 1-9.

Images