A method for inhibiting high-temperature brittle fracture of laser additive manufacturing high-strength nickel-based alloy

Abstract

The application discloses a method for inhibiting high-temperature brittle fracture of laser additive manufacturing high-strength nickel-based alloy, and belongs to the technical field of additive manufacturing engineering. The application first proposes that the total content of B reaches 0.1-0.8 wt.% by introducing a grain boundary element addition amount exceeding the upper limit of the existing nickel-based alloy composition threshold value, so as to promote the precipitation of grain boundary strengthening phase during heat treatment to strengthen the grain boundary, thereby avoiding the high-temperature brittle fracture of the high-strength nickel-based alloy formed by LPBF, and finally obtaining a laser powder bed fusion forming component which is free of forming cracks in the structure and has excellent high-temperature mechanical properties.

Description

Technical Field

[0001] This invention belongs to the field of additive manufacturing engineering technology, and in particular relates to a method for suppressing high-temperature brittle fracture of high-strength nickel-based alloys manufactured by laser additive manufacturing. Background Technology

[0002] Laser powder bed fusion (LPBF), as a high-precision laser additive manufacturing technology, exhibits unparalleled advantages over traditional processes in the integrated forming, lightweight design, and rapid manufacturing of complex parts, providing a novel solution for the fabrication of complex nickel-based alloy components in the aerospace field. Precipitation-strengthened nickel-based alloys, reinforced by dispersed γ′ phases (Ni3(Al,Ti,Ta)), are widely used in hot-end components in the aerospace field due to their excellent high-temperature strength, ductility, and corrosion resistance. With the increasing demands of the aerospace industry for ultra-high load-bearing capacity, extreme heat resistance, and high reliability, more stringent requirements are being placed on the high-temperature strength, ductility, and service stability of nickel-based alloys. High-strength nickel-based alloys for high-temperature service (above 850℃) typically contain a high proportion of Al, Ti, and Ta elements to obtain a high volume fraction of γ′ phase to maintain their high-temperature strength. However, for the manufacture of LPBF (Limited-Body Fabrication) for high-strength nickel-based alloys, high levels of γ′ phase-forming elements can reduce the alloy's weldability, leading to severe cracking during LPBF forming. Furthermore, as the γ′ phase content in the microstructure increases, the alloy achieves high strength but also suffers a loss of plasticity, potentially even resulting in high-temperature brittle fracture. This severely limits the application of LPBF technology in high-strength nickel-based alloy components.

[0003] Current research generally focuses on reducing / removing grain boundary elements such as B, C, Zr, and Si in nickel-based alloys to narrow the solidification temperature range and decrease the content of low-melting-point liquid films between grains, thereby reducing the alloy's susceptibility to hot cracking and preventing cracking during LPBF forming. However, grain boundary elements such as B, C, Zr, and Si can strengthen grain boundaries, reduce grain boundary diffusion rates, and improve microstructure stability. While reducing / removing these elements can suppress cracking in nickel-based alloys during LPBF forming, it also leads to a loss of high-temperature plasticity. For example, removing Zr and B from IN738 alloy can effectively suppress cracking during LPBF forming, but the alloy loses plasticity above 800℃, exhibiting brittle fracture. On the other hand, compared to casting (10... 0 -10 3 The low cooling rate of the process (K / s) results in an LPBF as high as 10. 8The high cooling rate (K / s) significantly refines the size of alloy dendrites and grains. Simultaneously, the layer-by-layer manufacturing technique of LPBF results in columnar grains that grow epitaxially along the forming direction. Compared to the equiaxed grains of cast nickel-based alloys, the fine columnar grains in the microstructure contribute to the high strength and low plasticity of LPBF-formed nickel-based alloys. Furthermore, when the temperature exceeds 0.5Tm (Tm is the alloy melting point), grain boundary strength is lower than intragranular strength. The high proportion of grain boundaries in LPBF-formed nickel-based alloys becomes the preferred failure site, and the columnar grains' inability to coordinate plastic deformation further exacerbates grain boundary stress concentration. Therefore, LPBF-formed nickel-based alloys typically exhibit insufficient plasticity at high temperatures. For high-strength nickel-based alloys containing a high proportion of γ′ phase, the difference between intragranular and grain boundary strength further increases, leading to severe intergranular brittle fracture at high temperatures. Therefore, in order to promote the application of LPBF technology in high-strength nickel-based alloy hot-end components, it is urgent to improve the composition of existing high-strength nickel-based alloys according to the forming characteristics of LPBF process in order to suppress their high-temperature brittle fracture.

[0004] Studies have found that after removing grain boundary elements such as B, C, Zr, and Si, the content of reinforcing phases such as carbides and borides precipitated along the grain boundaries in high-strength nickel-based alloys formed by LPBF is significantly reduced after heat treatment. This leads to a weakening of the high-temperature strength of the grain boundaries, premature failure of the grain boundaries during plastic deformation, and thus brittle fracture. Summary of the Invention

[0005] To address the aforementioned technical problems, this invention proposes a method for suppressing high-temperature brittle fracture of high-strength nickel-based alloys manufactured by laser additive manufacturing.

[0006] To achieve the above objectives, the present invention provides the following technical solution:

[0007] A method for suppressing high-temperature brittle fracture of high-strength nickel-based alloys manufactured by laser additive manufacturing involves adding boron (B) to the high-strength nickel-based alloy to achieve a total B content of 0.1-0.8 wt.% in the resulting pre-alloyed powder.

[0008] This invention is the first to propose that by introducing grain boundary element additions exceeding the upper limit of existing nickel-based alloy composition thresholds, the precipitation of grain boundary strengthening phases during heat treatment can be promoted to strengthen grain boundaries, thereby avoiding high-temperature brittle fracture of high-strength nickel-based alloys formed by LPBF.

[0009] Furthermore, the method specifically includes the following steps:

[0010] Adding boron (B) to a high-strength nickel-based alloy yields a pre-alloyed powder.

[0011] The pre-alloyed powder was prepared into alloy specimens using a laser powder bed melting process;

[0012] The alloy specimen was subjected to heat treatment.

[0013] This invention promotes the segregation of B and Cr elements in high-strength nickel-based alloys by introducing a B element content exceeding the upper limit of existing nickel-based alloy composition thresholds (the B content in existing nickel-based alloys is generally below 0.1 wt.%). Subsequently, heat treatment is used to induce the precipitation of a high proportion of Cr5B3 phase at the grain boundaries, thereby strengthening the grain boundaries and preventing premature failure during plastic deformation. Adding an appropriate amount of B not only inhibits cracking in LPBF-formed high-strength nickel-based alloys, but also gives the alloy excellent high-temperature strength and plasticity.

[0014] Furthermore, the high-strength nickel-based alloy is a precipitation-strengthened nickel-based alloy with a total mass fraction of γ′ phase-forming elements Al+Ti+Ta greater than 7wt.%.

[0015] Furthermore, the method for adding element B is selected from either of the following two methods:

[0016] Method 1: Add as an alloying element during the base material smelting process;

[0017] Method 2: Directly mix high-strength nickel-based alloy powder with element B.

[0018] Furthermore, in the second method, the direct mixing method is ball milling.

[0019] Furthermore, the heat treatment method is selected from one of atmospheric pressure heat treatment, hot isostatic pressing, and heat treatment under an argon protective atmosphere.

[0020] Furthermore, the specific steps of the atmospheric pressure heat treatment include: holding at 1050-1200℃ for 1-3 hours, and then aging heat treatment at 800-900℃ for 8-24 hours.

[0021] Furthermore, the specific steps of the hot isostatic pressing treatment include: applying an isostatic pressure of 100-200 MPa at 1050-1170℃ and holding it at that temperature for 1-3 hours, followed by aging heat treatment at 800-900℃ under normal pressure for 8-24 hours.

[0022] Furthermore, the specific operation steps of the heat treatment under the argon protective atmosphere include: holding at 1050-1200℃ for 1-3 hours in a flowing argon atmosphere, and then aging heat treatment at 800-900℃ for 8-24 hours.

[0023] Compared with the prior art, the present invention has the following advantages and technical effects:

[0024] While LPBF (Liquid-Layer Basis Forming) technology overcomes the limitations of traditional manufacturing in areas such as integrated part structures and complex shapes and functions, providing a novel solution for the forming of nickel-based alloys, the extremely high cooling rate and spatially varying temperature gradients of LPBF make it difficult to apply to most existing nickel-based alloys, resulting in problems such as forming cracks and insufficient high-temperature plasticity. To avoid cracking of nickel-based alloys during LPBF forming, existing methods generally reduce the content of grain boundary elements such as boron (B), carbon (C), and silicon (Si). Although this approach can prevent cracking during LPBF forming, the absence of grain boundary elements significantly deteriorates the alloy's high-temperature performance. In contrast, this invention simply requires the addition of boron exceeding the upper limit of existing compositional thresholds. This not only gives the alloy excellent formability but also effectively prevents high-strength nickel-based alloys from high-temperature brittle fracture, providing a new approach for the compositional design of nickel-based high-temperature alloys specifically for laser additive manufacturing.

[0025] This invention has significant engineering implications. In the aerospace field, high-strength nickel-based alloys are commonly used to manufacture hot-section components of engines. High-temperature brittle fracture significantly reduces the service life and reliability of these components, threatening flight safety. Effectively suppressing this problem can ensure the stable operation of aero-engines under extreme conditions, promoting performance improvements and technological innovation, and contributing to the development of new aircraft. In the energy and power industry, nickel-based alloy components are used in gas turbines, nuclear power equipment, etc. Overcoming high-temperature brittle fracture enhances the equipment's resistance to high temperatures and pressures, improves energy conversion efficiency, reduces equipment maintenance costs and downtime frequency, and promotes efficient and sustainable development in the energy sector. From a materials processing perspective, solving the high-temperature brittle fracture problem can expand the application scope of laser additive manufacturing technology in the field of high-strength nickel-based alloys, promote near-net-shape manufacturing of complex structures and high-performance parts, accelerate the industrialization of advanced manufacturing technologies, and enhance the core competitiveness of the nation's high-end equipment manufacturing. Attached Figure Description

[0026] The accompanying drawings, which form part of this invention, are used to provide a further understanding of the invention. The illustrative embodiments of the invention and their descriptions are used to explain the invention and do not constitute an undue limitation of the invention. In the drawings:

[0027] Figure 1 The images are optical microscope images, where (a) is an optical microscope image of IN738 alloy formed by LPBF in Comparative Example 1; and (b) is an optical microscope image of IN738 alloy formed by LPBF with 0.5 wt.% B element in Example 1.

[0028] Figure 2The images are transmission electron microscope (TEM) images, where (a) is a TEM image of the IN738 alloy in Comparative Example 1, with the inset showing the energy distribution spectrum of Cr; and (b) is a TEM image of the IN738 alloy with 0.5 wt.% B added in Example 1, with the inset showing the energy distribution spectrum of Cr.

[0029] Figure 3 Images are obtained using scanning electron microscopy (SEM) and transmission electron microscopy (TEM), where (a) is a SEM image of the IN738 alloy after heat treatment in Comparative Example 1; (b) is a SEM image of the IN738 alloy with 0.5 wt.% B added after heat treatment in Example 1; and (c) is a TEM image of the second phase in the IN738 alloy with 0.5 wt.% B added after heat treatment in Example 1, with the inset showing the selected area electron diffraction pattern of the second phase particles.

[0030] Figure 4 A comparison of the tensile curves of samples prepared in Comparative Example 1 and Example 1 at 900°C;

[0031] Figure 5 The images are optical microscope images, where (a) is an optical microscope image of the novel nickel-based alloy formed by LPBF in Comparative Example 2; and (b) is an optical microscope image of the novel nickel-based alloy formed by LPBF in Example 2 after adding 0.4 wt.% of B element.

[0032] Figure 6 The images are scanning electron microscope images. (a) is a scanning electron microscope image of the novel nickel-based alloy after heat treatment in Comparative Example 2; (b) is a scanning electron microscope image of the novel nickel-based alloy after heat treatment in Example 2 after adding 0.4 wt.% B.

[0033] Figure 7 A comparison of the tensile curves of the samples prepared in Comparative Example 2 and Example 2 at 900°C;

[0034] Figure 8 The images are optical microscope images, of which (a) is an optical microscope image of IN738 alloy with 1 wt.% B element added after LPBF forming in Comparative Example 3; (b) is an optical microscope image of IN738 alloy with 0.5 wt.% B element added after heat treatment in Example 1; and (c) is an optical microscope image of IN738 alloy with 1 wt.% B element added after heat treatment in Comparative Example 3. Detailed Implementation

[0035] Various exemplary embodiments of the present invention will now be described in detail. This detailed description should not be considered as a limitation of the present invention, but rather as a more detailed description of certain aspects, features, and embodiments of the present invention.

[0036] It should be understood that the terminology used in this invention is merely for describing particular embodiments and is not intended to limit the invention. Furthermore, with respect to numerical ranges in this invention, it should be understood that each intermediate value between the upper and lower limits of the range is also specifically disclosed. Every smaller range between any stated value or intermediate value within a stated range, and any other stated value or intermediate value within said range, is also included in this invention. The upper and lower limits of these smaller ranges may be independently included or excluded from the range.

[0037] Unless otherwise stated, all technical and scientific terms used herein have the same meaning as commonly understood by one of ordinary skill in the art. While only preferred methods and materials have been described herein, any methods and materials similar or equivalent to those described herein may be used in the implementation or testing of this invention. All references to this specification are incorporated by way of citation to disclose and describe methods and / or materials associated with those references. In the event of any conflict with any incorporated reference, the content of this specification shall prevail.

[0038] Various modifications and variations can be made to the specific embodiments described in this specification without departing from the scope or spirit of the invention, as will be apparent to those skilled in the art. Other embodiments derived from this specification will also be readily apparent to those skilled in the art. This specification and embodiments are merely exemplary.

[0039] The terms “include,” “including,” “have,” “contain,” etc., used in this article are all open-ended terms, meaning that they include but are not limited to.

[0040] This invention provides a method for suppressing high-temperature brittle fracture of high-strength nickel-based alloys manufactured by laser additive manufacturing, specifically including the following steps:

[0041] S1. Adding element B to a high-strength nickel-based alloy yields a pre-alloyed powder;

[0042] S2. The pre-alloyed powder is prepared into alloy specimens using a laser powder bed melting process;

[0043] S3. The alloy specimen is subjected to heat treatment.

[0044] In some optional embodiments of the present invention, in step S1, the amount of element B added is based on the total B content in the obtained pre-alloyed powder being 0.1-0.8 wt.% (e.g., 0.4 wt.% or 0.5 wt.%). The element B is in the micrometer range.

[0045] In this invention, the high-strength nickel-based alloy is a precipitation-strengthened nickel-based alloy with a total mass fraction of γ′ phase-forming elements Al+Ti+Ta greater than 7wt.%. That is, the total proportion of Al, Ti and Ta in the high-strength nickel-based alloy used must be greater than 7wt.% in order to meet the cracking requirements.

[0046] In some optional embodiments of the present invention, the method of adding element B includes the following two:

[0047] Method 1: Added in the form of alloying elements during the smelting process of the base material; the composition of the base material includes Cr: 12-15wt.%, Mo: 2.6-3wt.%, Al: 3-4wt.%, Ti: 4-5wt.%, Nb: 1-1.5wt.%, Fe: 0.8-1wt.%, W: 1-1.5wt.%, Co: 20-25wt.%, Ni: balance.

[0048] Method 2: Directly mix high-strength nickel-based alloy with element B, specifically by mixing through mechanical ball milling.

[0049] In some optional embodiments of the present invention, the heat treatment method is selected from one of atmospheric pressure heat treatment, hot isostatic pressing, and heat treatment under an argon protective atmosphere.

[0050] In some optional embodiments of the present invention, the specific operation steps of the atmospheric pressure heat treatment include: holding at 1050-1200℃ for 1-3 hours, and then aging heat treatment at 800-900℃ for 8-24 hours.

[0051] In some optional embodiments of the present invention, the specific operation steps of the hot isostatic pressing treatment include: applying an isostatic pressure of 100-200 MPa at 1050-1170°C and holding it at that temperature for 1-3 hours, and then performing an aging heat treatment at 800-900°C under normal pressure for 8-24 hours.

[0052] In some optional embodiments of the present invention, the specific operation steps of the heat treatment under the argon protective atmosphere include: holding at 1050-1200℃ for 1-3 hours in a flowing argon atmosphere, and then aging heat treatment at 800-900℃ for 8-24 hours.

[0053] This invention introduces boron (B) to induce a high proportion of Cr5B3 particles to precipitate along grain boundaries in high-strength nickel-based alloys after heat treatment. This strengthens the alloy grain boundaries and effectively suppresses high-temperature brittle fracture in LPBF-formed high-strength nickel-based alloys. The result is a laser powder bed fusion formed component with no forming cracks and excellent high-temperature mechanical properties.

[0054] IN738 is a typical representative of high-strength nickel-based alloys. To avoid cracking during LPBF forming, existing research generally removes Zr and B elements from IN738 alloys. Therefore, in the following embodiments of the present invention, the IN738 used is a commercially available product that does not contain Zr and B elements, and satisfies the requirement that the total proportion of Al, Ti, and Ta is greater than 7 wt.%.

[0055] The technical solution of the present invention will be further illustrated by the following embodiments.

[0056] Example 1

[0057] A method for suppressing high-temperature brittle fracture of high-strength nickel-based alloys manufactured by laser additive manufacturing specifically includes the following steps:

[0058] S1. Add element B to IN738 alloy and ball mill to mix so that the total content of B in the obtained pre-alloyed powder is 0.5 wt.%.

[0059] S2. The pre-alloyed powder obtained in step S1 is prepared into alloy specimens using laser powder bed melting (LPBF) process;

[0060] S3. The alloy specimen obtained in step S2 is subjected to atmospheric pressure heat treatment in a muffle furnace. The specific process is as follows: hold at 1050℃ for 2 hours, then furnace cool to 800℃ and hold for 12 hours to obtain the sample (denoted as: IN738+0.5B).

[0061] Comparative Example 1

[0062] Same as Example 1, except that element B is not added, and the specific steps include:

[0063] S1. The IN738 alloy was prepared into alloy specimens using laser powder bed melting (LPBF) process;

[0064] S2. The alloy specimen obtained in step S1 is subjected to atmospheric pressure heat treatment in a muffle furnace. The specific process is as follows: hold at 1050℃ for 2 hours, then cool in the furnace to 800℃ and hold for 12 hours to obtain the sample (denoted as: IN738).

[0065] Example 2

[0066] A method for suppressing high-temperature brittle fracture of high-strength nickel-based alloys manufactured by laser additive manufacturing specifically includes the following steps:

[0067] S1. Cr: 13.5 wt.%, Mo: 2.7 wt.%, Al: 3.8 wt.%, Ti: 4.5 wt.%, Ta: 2 wt.%, Nb: 1.2 wt.%, Fe: 0.8 wt.%, W: 1.3 wt.%, Co: 24 wt.%, B: 0.4 wt.%, and the balance Ni are mixed and smelted to obtain a novel nickel-based alloy;

[0068] S2. The novel nickel-based alloy obtained in step S1 is prepared into alloy specimens using laser powder bed melting (LPBF) process;

[0069] S3. The alloy specimen obtained in step S2 is subjected to atmospheric pressure heat treatment in a muffle furnace. The specific process is as follows: hold at 1100℃ for 2 hours, then furnace cool to 900℃ and hold for 12 hours to obtain the sample (denoted as the new nickel-based alloy +0.4B).

[0070] Comparative Example 2

[0071] Same as Example 2, except that element B is not added, and the specific steps include:

[0072] S1. Cr: 13.5 wt.%, Mo: 2.7 wt.%, Al: 3.8 wt.%, Ti: 4.5 wt.%, Ta: 2 wt.%, Nb: 1.2 wt.%, Fe: 0.8 wt.%, W: 1.3 wt.%, Co: 24 wt.%, and the balance Ni are mixed and smelted to obtain a high-strength nickel-based alloy;

[0073] S2. The novel nickel-based alloy obtained in step S1 is prepared into alloy specimens using laser powder bed melting (LPBF) process;

[0074] S3. The alloy specimen obtained in step S2 is subjected to atmospheric pressure heat treatment in a muffle furnace. The specific process is as follows: hold at 1100℃ for 2 hours, then cool in the furnace to 900℃ and hold for 12 hours to obtain the sample (denoted as the new nickel-based alloy).

[0075] Figure 1 The images are optical microscope images, where (a) is an optical microscope image of IN738 alloy formed by LPBF in Comparative Example 1; and (b) is an optical microscope image of IN738 alloy formed by LPBF with 0.5 wt.% B element in Example 1. Figure 1 As can be seen in (a) of the image, the sample is almost completely dense, and no cracks were observed in the image; from Figure 1 As can be seen in (b) of the image, the sample is almost completely dense, and no cracks were observed in the image.

[0076] Figure 2These are transmission electron microscope (TEM) images, where (a) is a TEM image of the IN738 alloy in Comparative Example 1, with the inset showing the energy dispersive spectral density (EDS) spectrum of Cr; (b) is a TEM image of the IN738 alloy with 0.5 wt.% B added in Example 1, with the inset showing the EDS spectrum of Cr (since B is a light element, its distribution is difficult to detect using existing EDS methods). Figure 2 As can be seen in (a), the dendrite size in the microstructure is about 1 μm, and no Cr segregation is observed in the microstructure; Figure 2 As can be seen in (b), the dendrite size in the microstructure is about 1 μm, and obvious Cr element segregation is observed at the dendrite boundaries.

[0077] Figure 3 Images are obtained using scanning electron microscopy (SEM) and transmission electron microscopy (TEM), where (a) is a SEM image of the IN738 alloy after heat treatment in Comparative Example 1; (b) is a SEM image of the IN738 alloy with 0.5 wt.% B added after heat treatment in Example 1; and (c) is a TEM image of the second phase in the IN738 alloy with 0.5 wt.% B added after heat treatment in Example 1, with the inset showing the selected area electron diffraction pattern of the second phase particles. Figure 3 As can be seen in (a) of the sample, granular second phase of about 0.2 μm precipitated along the grain boundaries in the heat-treated sample; from Figure 3 As can be seen in (b) of the figure, a granular second phase of about 0.3 μm precipitated along the grain boundaries in the heat-treated sample. Compared with the IN738 sample in Figure (a), the size and content of the second phase at the grain boundaries are significantly increased. Figure 3 (c) indicates that the second phase particles precipitated in the IN738+0.5B sample are Cr5B3 phase.

[0078] Figure 4 A comparison of the tensile curves of samples prepared in Comparative Example 1 and Example 1 at 900°C; from Figure 4 The tensile test results at 900℃ show that the IN738 sample fractured before reaching its yield strength, with a plastic elongation of 0, exhibiting brittle fracture; the IN738+0.5B sample showed significant plastic deformation, with a tensile strength of 570 MPa and a plastic elongation of 5.5%.

[0079] Figure 5 The images are optical microscope images, where (a) is an optical microscope image of the novel nickel-based alloy formed by LPBF in Comparative Example 2; (b) is an optical microscope image of the novel nickel-based alloy formed by LPBF in Example 2 after adding 0.4 wt.% of B element; from Figure 5 As can be seen in (a), the sample is almost completely dense, and no cracks were observed in the image; from Figure 5As can be seen from (b) in the sample, no large number of cracks were observed.

[0080] Figure 6 The images are scanning electron microscope (SEM) images: (a) is an SEM image of the novel nickel-based alloy after heat treatment in Comparative Example 2; (b) is an optical micrograph of the novel nickel-based alloy after heat treatment in Example 2 with the addition of 0.4 wt.% B; from Figure 6 As can be seen in (a) of the sample, a small amount of granular second phase, approximately 0.1 μm in size, precipitated along the grain boundaries in the heat-treated sample; from Figure 6 As can be seen in (b), a large number of granular Cr5B3 second phases with a size of about 2 μm precipitated along the grain boundaries in the heat-treated sample.

[0081] Figure 7 This is a comparison of the tensile curves of the samples prepared in Comparative Example 2 and Example 2 at 900°C; from Figure 7 The tensile test results at 900℃ show that the novel nickel-based alloy obtained in Comparative Example 2 fractured before reaching the yield strength, with a plastic elongation of 0, exhibiting brittle fracture. The novel nickel-based alloy +0.4B sample obtained in Example 2 showed significant plastic deformation, with a tensile strength of 630 MPa and a plastic elongation of 5%.

[0082] Comparative Example 3

[0083] Same as Example 1, except that step S1 specifically involves adding element B to the IN738 alloy, ball milling and mixing to obtain a total B content of 1 wt.% in the pre-alloyed powder. Step S3 specifically involves holding the powder at 1200°C for 2 hours and then air cooling it to room temperature.

[0084] Figure 8 The images are optical microscope images, where (a) is an optical microscope image of the IN738 alloy with 1 wt.% B element added during LPBF forming in Comparative Example 3; (b) is an optical microscope image of the IN738 alloy with 0.5 wt.% B element added after heat treatment in Example 1; and (c) is an optical microscope image of the IN738 alloy with 1 wt.% B element added after heat treatment in Comparative Example 3. The results showed that after adding 1 wt.% B element to the IN738 alloy, the sample was almost completely dense, and no cracks were observed in the images. Figure 8 As shown in (a) above, the IN738 alloy with 0.5 wt.% B added after heat treatment at 1200℃ still exhibits columnar grains that have undergone epitaxial growth along the forming direction, as shown in (a). Figure 8As shown in (b). However, the IN738 alloy with 1 wt.% B element added underwent grain boundary overmelting during heat treatment at 1200℃, resulting in spheroidization of large-sized spherical grains and the presence of large-sized segregated phases at the grain boundaries. This is mainly because the excessive addition of B element leads to an excessive number of low-melting-point segregated phases at the grain boundaries, which can melt at relatively low temperatures. The spheroidized grains and the large-sized segregated phases at the grain boundaries are detrimental to the high-temperature service life of the alloy, such as... Figure 8 As shown in (c) in the figure.

[0085] The above are merely preferred embodiments of the present invention, but the scope of protection of the present invention is not limited thereto. Any variations or substitutions that can be easily conceived by those skilled in the art within the scope of the technology disclosed in the present invention should be included within the scope of protection of the present invention. Therefore, the scope of protection of the present invention should be determined by the scope of the claims.

Claims

1. A method for suppressing high-temperature brittle fracture of high-strength nickel-based alloys manufactured by laser additive manufacturing, characterized in that, Boron (B) was added to a high-strength nickel-based alloy to achieve a total B content of 0.4-0.8 wt.% in the resulting pre-alloyed powder. Specifically, the following steps are included: Adding boron (B) to a high-strength nickel-based alloy yields a pre-alloyed powder. The pre-alloyed powder was prepared into alloy specimens using a laser powder bed melting process; The alloy specimen was subjected to heat treatment; The high-strength nickel-based alloy is a precipitation-strengthened nickel-based alloy with a total mass fraction of γ′ phase-forming elements Al+Ti+Ta greater than 7wt.%.

2. The method for suppressing high-temperature brittle fracture of high-strength nickel-based alloys manufactured by laser additive manufacturing according to claim 1, characterized in that, The method for adding element B is selected from either of the following two methods: Method 1: Add as an alloying element during the base material smelting process; Method 2: Directly mix high-strength nickel-based alloy powder with element B.

3. The method for suppressing high-temperature brittle fracture of high-strength nickel-based alloys manufactured by laser additive manufacturing according to claim 2, characterized in that, In Method 2, the direct mixing method is ball milling.

4. The method for suppressing high-temperature brittle fracture of high-strength nickel-based alloys manufactured by laser additive manufacturing according to claim 1, characterized in that, The heat treatment method is selected from one of atmospheric pressure heat treatment, hot isostatic pressing, and heat treatment under an argon protective atmosphere.

5. The method for suppressing high-temperature brittle fracture of high-strength nickel-based alloys manufactured by laser additive manufacturing according to claim 4, characterized in that, The specific steps of the atmospheric pressure heat treatment include: holding at 1050-1200℃ for 1-3 hours, and then aging heat treatment at 800-900℃ for 8-24 hours.

6. The method for suppressing high-temperature brittle fracture of high-strength nickel-based alloys manufactured by laser additive manufacturing according to claim 4, characterized in that, The specific steps of the hot isostatic pressing treatment include: holding at 1050-1170℃ and 100-200MPa for 1-3 hours, and then holding at normal pressure and 800-900℃ for 8-24 hours.

7. The method for suppressing high-temperature brittle fracture of high-strength nickel-based alloys manufactured by laser additive manufacturing according to claim 4, characterized in that, The specific operation steps of the heat treatment under the argon protective atmosphere include: holding at 1050-1200℃ for 1-3 hours in a flowing argon atmosphere, and then aging heat treatment at 800-900℃ for 8-24 hours.

Images