Oxide dispersion strengthened nickel-based superalloy and preparation method thereof

Abstract

The invention discloses a preparation method of an oxide dispersion strengthened nickel-based superalloy, which comprises the following steps: firstly, preparing powder by circulating gas atomization for many times, returning coarse powder, mixing with a new material, remelting, realizing controllable accumulation of the oxygen content of the powder, and then forming by utilizing optimized LPBF process parameters, thereby obtaining the oxide dispersion strengthened nickel-based superalloy. Oxygen and active elements are promoted to react in a molten pool by utilizing a rapid solidification effect, and oxide dispersion strengthening particles of 10-50 nm are generated in situ; and finally, through specific solid solution and aging heat treatment, the ODS nickel-based high-temperature alloy component with a refined structure and excellent performance is obtained. According to the method, efficient cyclic utilization of powder resources and synchronous improvement of material performance are achieved, and the method has the advantages of economy and high performance and is suitable for manufacturing of hot end components such as aero-engines and gas turbines.

Description

Technical Field

[0001] This invention belongs to the field of nickel-based superalloy preparation technology, and particularly relates to an oxide dispersion strengthened nickel-based superalloy and its preparation method. Background Technology

[0002] Nickel-based superalloys are widely used in hot-end components such as aero-engines and gas turbines due to their excellent high-temperature strength, creep resistance, and corrosion resistance. Laser selective melting (LPBF), as a mainstream metal additive manufacturing technology, has extremely high requirements for the sphericity, particle size distribution, and oxygen content of the powder raw materials. Currently, the commonly used vacuum induction gas atomization (VIGA) technology produces powders in which only about 30% of the fine powder (15-60 μm) can be directly used for LPBF. A large amount of coarse powder (>60 μm) needs to be remelted or scrapped, resulting in low material utilization and high costs.

[0003] On the other hand, oxide dispersion strengthening (ODS) technology can significantly improve the high-temperature performance of superalloys. However, traditional ODS alloy preparation processes (such as mechanical alloying) suffer from problems such as long processes, high energy consumption, uneven particle distribution, and difficulty in compatibility with additive manufacturing, making it difficult to meet the integrated manufacturing requirements of complex components. Although some studies have attempted to utilize rapid solidification during LPBF to promote the reaction of oxygen with elements such as Al and Ti to form oxide particles and achieve a certain degree of in-situ strengthening, these studies typically lack systematic control over oxygen content and integrated design of the entire "powdering-forming-heat treatment" process, making it difficult to achieve controllable adjustment of oxygen content and stable generation of ODS particles. Therefore, existing technologies still cannot ensure the controllability of the ODS strengthening mechanism and the stability of mechanical properties while achieving efficient recycling of powder. Summary of the Invention

[0004] The main objective of this invention is to provide an oxide dispersion strengthened nickel-based superalloy and its preparation method, so as to solve the problems of low powder utilization, difficulty in controlling the ODS strengthening process, and inability of traditional processes to balance economy and high performance in the prior art.

[0005] Therefore, the present invention provides a method for preparing oxide dispersion strengthened nickel-based superalloys, comprising the following steps: S1. Circulating atomization powder production: Using nickel-based superalloy ingots as initial raw materials, vacuum induction gas atomization powder production is performed to obtain fine powder that can be used for LPBF and coarse powder that cannot be used directly; the obtained coarse powder is mixed with new nickel-based superalloy ingots in a predetermined mass ratio and used as raw material for vacuum induction gas atomization powder production again; this cyclic process is repeated, and by controlling the number of cycles and the return ratio of coarse powder, the oxygen content of the finally obtained fine powder is controlled within the range of 0.015 wt% to 0.060 wt%; the reason for controlling the O content within the above range is that when the O content exceeds 0.5 wt%, large particle inclusions are easily formed, leading to a decrease in alloy performance; S2, Selective Laser Melting: The circulating atomized fine powder obtained in step S1 is used as raw material and selective laser melting is performed under a protective atmosphere. Nano-oxide dispersion reinforced particles are generated in situ in the molten pool through optimized laser process parameters. S3. Heat treatment for microstructure control: The formed part obtained in step S2 is subjected to standard solution treatment and aging treatment to obtain a nickel-based high-temperature alloy component with fine grain structure and dispersed nano-oxide.

[0006] Specifically, the return ratio of the coarse powder is 30%-50%, preferably 40%; this is because the yield of fine powder from atomization is commonly 50-60% in engineering, which complements the coarse powder ratio, achieving efficient material utilization and continuous production. Furthermore, if the return ratio is too low, the O content may be introduced too little.

[0007] Specifically, the number of cycles is 3 to 6, preferably 5.

[0008] Specifically, in step S1, the process parameters for vacuum induction gas atomization powder production are: vacuum degree not higher than 10. -2 Pa, melting temperature 1550-1600°C, atomizing gas is Ar, gas pressure 2.5-3.5 MPa.

[0009] Specifically, in step S2, the process parameters for laser selective melting forming are: laser power 200-300 W, scanning speed 800-1000 mm / s, layer thickness 20-40 μm, and scanning spacing 80-120 μm; preferably, the laser power is 250 W, the scanning speed is 900 mm / s, the layer thickness is 30 μm, and the scanning spacing is 100 μm.

[0010] Specifically, in step S2, an interlayer rotation scanning strategy is adopted, with a rotation angle of 60° to 90°.

[0011] Specifically, in step S2, the particle size of the in-situ generated nano-oxide dispersion-reinforced particles is 10-50 nm, and their composition is at least one of Cr2O3 and Al2O3-TiN composite structures.

[0012] Specifically, the nickel-based superalloy is GH4169, IN625, or IN718.

[0013] Specifically, for GH4169 alloy and IN718 alloy, the standard solution treatment regime is 1100°C for 1 hour; the standard aging treatment is a two-stage aging process, which is 720°C for 4 hours and then 620°C for 8 hours.

[0014] Specifically, for IN625 alloy, the standard solution treatment regime is 1150°C for 1 hour, and the standard aging treatment regime is 700°C for 6 hours.

[0015] This application also provides an oxide dispersion-strengthened nickel-based superalloy component, which is prepared by the above method. The component has a grain size of 2-4 μm and nano-oxide particles with a particle size of 10-50 nm are dispersed inside and outside its matrix.

[0016] Compared with the prior art, the present invention has the following beneficial effects: 1. Resource recycling and cost advantages: By establishing a closed-loop material flow of "coarse powder return - circulating atomization", the waste coarse powder in the traditional process is transformed into high-value LPBF raw materials, and the comprehensive utilization rate of powder is increased from about 30% to more than 80%, which significantly reduces the cost of raw materials.

[0017] 2. Controllable and reliable enhancement effect: It is the first to achieve gradient design and precise control of powder oxygen content through cyclic atomization process, which provides a stable and controllable oxygen source for the in-situ reaction in the subsequent LPBF process, fundamentally solving the technical problem of high randomness and difficulty in repeating the generation of in-situ ODS particles.

[0018] 3. Superior microstructure and properties: Utilizing optimized LPBF process parameters, densification and in-situ reaction are simultaneously achieved in the molten pool, resulting in clean, firmly bonded, and uniformly distributed nano-oxide particles at the matrix interface. Combined with subsequent heat treatment, a dual-phase composite reinforced microstructure with fine grains (2-4 μm) and dispersed ODS particles is obtained, enabling the material to exhibit excellent strength, plasticity, and creep properties at both room temperature and high temperature.

[0019] 4. Strong process integration and universality: This invention innovatively integrates the three links of material preparation, component forming and performance improvement into one, forming a short-process, high-efficiency and high-performance manufacturing technology path, which can be widely applied to various grades of nickel-based high-temperature alloy systems and has great engineering promotion value. Attached Figure Description

[0020] To more clearly illustrate the technical solutions in the embodiments of the present invention, the accompanying drawings used in the description of the embodiments will be briefly introduced below. Obviously, the accompanying drawings described below are only some embodiments of the present invention. For those skilled in the art, other drawings can be obtained based on these drawings without creative effort.

[0021] Figure 1 This is a flowchart of the cyclic atomization powder production and LPBF experiment of the present invention; Figure 2 This is a microstructure image of the heat-treated state after five cycles of powder LPBF printing according to the present invention; Figure 3 This is a comparison chart of the mechanical properties of GH4169 prepared by LPBF after powder recycling. Detailed Implementation

[0022] The technical solutions of the embodiments of the present invention will be clearly and completely described below with reference to the accompanying drawings. Obviously, the described embodiments are only some embodiments of the present invention, and not all embodiments. Based on the embodiments of the present invention, all other embodiments obtained by those skilled in the art without creative effort are within the scope of protection of the present invention.

[0023] In the description of this invention, it should be understood that the terms "center," "longitudinal," "lateral," "length," "width," "thickness," "upper," "lower," "front," "rear," "left," "right," "vertical," "horizontal," "top," "bottom," "inner," "outer," "clockwise," "counterclockwise," "axial," "radial," and "circumferential" indicate the orientation or positional relationship based on the orientation or positional relationship shown in the accompanying drawings. They are used only for the convenience of describing this invention and simplifying the description, and are not intended to indicate or imply that the device or element referred to must have a specific orientation, or be constructed and operated in a specific orientation. Therefore, they should not be construed as limitations on this invention.

[0024] Furthermore, the terms "first" and "second" are used for descriptive purposes only and should not be construed as indicating or implying relative importance or implicitly specifying the number of technical features indicated. Thus, a feature defined as "first" or "second" may explicitly or implicitly include one or more of that feature. In the description of this invention, "a plurality of" means two or more, unless otherwise explicitly specified.

[0025] This invention aims to propose an integrated preparation method of "circulating atomized powder + LPBF in-situ ODS synthesis". The method first involves multiple cycles of gas atomization to produce powder, then returning the coarse powder to be mixed and remelted with virgin material, achieving a controllable accumulation of oxygen content in the powder (0.015 wt% to 0.060 wt%). Then, optimized LPBF process parameters are used for shaping, utilizing the rapid solidification effect in the molten pool to promote the reaction of oxygen with active elements, generating 10-50 nm oxide dispersion-strengthened particles in situ. Finally, standard solution treatment followed by aging heat treatment yields ODS nickel-based superalloy components with refined microstructure and excellent performance. This invention achieves efficient recycling of powder resources and simultaneous improvement of material properties, combining economic and high-performance advantages, and is suitable for the manufacture of hot-end components such as aero-engines and gas turbines. Figure 1 This demonstrates the entire process flow from coarse powder return, mixing and smelting, gas atomization powder preparation to LPBF printing and heat treatment. This process achieves controlled accumulation of oxygen content and stable powder properties. The specific steps are as follows:

[0026] Circulating atomization powder production is the fundamental step in this invention's technical solution. It aims to achieve efficient resource recycling through multiple returns of coarse powder, while simultaneously gradually controlling the oxygen content of the powder to provide a raw material basis for subsequent ODS in-situ strengthening. A vacuum induction gas atomization (VIGA) device is used, initially using brand-new GH4169 ingots (composition: Ni-based, Cr 18-21 wt%, Fe balance, Nb 4.75-5.5 wt%, Mo 2.8-3.3 wt%, Ti 0.65-1.15 wt%, Al 0.2-0.8 wt%, oxygen content <0.01 wt%) as the starting material. Of course, other nickel-based superalloys such as IN625, IN718, or Hastelloy X can also be used. The total melting amount is controlled at 20 kg. The vacuum degree is not higher than 10. -2 Under induction heating power of Pa, the ingot is melted to 1550-1600°C, forming a liquid metal stream. This stream falls through a ceramic nozzle (0.4-0.6 mm in diameter) and is atomized into powder by a high-pressure inert gas (Ar, pressure 2.5-3.5 MPa). The atomizing gas flow rate is controlled at 5-10 m / s to ensure rapid cooling of the droplets (cooling rate 10 m / s). 5-6 K / s), forming a powder with high sphericity (sphericity > 92%).

[0027] After atomization, fine powder (15-60 μm, used for LPBF) and coarse powder (>60 μm, accounting for approximately 70%) are separated by vibrating sieve. The coarse powder is not discarded directly but returned according to a set mass ratio, mixed with new ingots, and remelted for atomization. Based on thermodynamic simulation optimization, the return ratio was determined to be 40%. The oxygen content in the returned powder had accumulated to 0.015 wt%, and the total oxygen content after mixing could be controllably increased to 0.017 wt%, avoiding embrittlement due to over-oxidation. The cycle was repeated 5 times, with composition analysis (ICP-OES) and oxygen content detection (inert gas melting method) performed before and after each melting process, ultimately resulting in the powder oxygen content gradually increasing from 0.017 wt% to 0.038 wt%. Of course, the return ratio and number of cycles can be selected according to the specific oxygen content requirements. It is only necessary to ensure that the oxygen content of the final fine powder is controlled between 0.015 wt% and 0.060 wt%. However, too many cycles (>6 times) may cause Nb and Ti segregation, so the upper limit is set to 5 times to balance economy and performance.

[0028] The oxygen accumulation mechanism in the above steps stems from the partial dissolution of the oxide layer (Al2O3) on the surface of the returned powder during remelting, providing a source of active oxygen, while the VIGA vacuum environment inhibits additional oxidation. The powder particle size distribution follows the Rosin-Rammler model, D 50 With a particle size of approximately 35 μm and a span value <1.0, excellent flowability is ensured (Hall flow rate <13 s / 50 g). To further improve powder uniformity, auxiliary measures can be introduced: adding 0.1-0.5 wt% deoxidizer (such as Ca or Mg) before atomization to control inclusions <10 ppm; preheating the gas to 200-300°C to reduce the droplet cooling rate and promote a 15-20% increase in fine powder yield. Compared with traditional single-cycle VIGA, this cyclic process achieves a powder utilization rate of 82%, reduces energy consumption by 30%, and lowers overall cost by approximately 37%, making it suitable for nickel-based alloys such as IN718 (GH4169) and IN625.

[0029] Powder property control and characterization The resulting recycled powder must meet stringent LPBF requirements: sphericity, flowability, and oxygen content are key parameters. Sphericity is measured using an optical microscope and image analysis software (ImageJ), with a target of >95% to ensure uniform powder spreading without bridging. Flowability and bulk density are tested according to ASTM B213 standards, with a recycled powder flowability of 12-14 s / 50 g and a density of 2.8-3.0 g / cm³. 3It is superior to commercial pre-alloyed powders (15-18 s / 50 g). The oxygen content gradient design (0.017-0.038 wt%) is based on the ODS strengthening requirements: low oxygen (<0.02 wt%) is beneficial to the initial printing density, and high oxygen (0.03-0.04 wt%) promotes the subsequent oxide formation, but does not exceed 0.05 wt% to prevent porosity defects.

[0030] Chemical stability was verified by XRF and GDMS: Cr and Nb fluctuations after cycling were <0.5 wt%, with no significant volatilization. SEM microstructure observation showed a smooth powder surface, a small amount of satellite powder (<5%), and no hollow internal structure (CT scan confirmed porosity <1%). For optimized properties, post-processing can be performed: plasma spheroidization or mechanical sieving to further narrow the particle size distribution and improve the uniformity of the powder layer thickness (30-50 μm). This control ensures the stability of the powder in the LPBF melt pool and avoids crack initiation due to oxygen heterogeneity.

[0031] LPBF forming is the core of the strengthening mechanism of this invention, utilizing laser-induced rapid solidification (10 6 The oxygen-alloying element reaction is driven by a laser (K / s). A commercial LPBF device (such as the EOS M290) is used, with the substrate preheated to 100-150°C and a protective atmosphere of Ar (O2 < 100 ppm). Process parameters: laser power 200-300 W (preferably 250 W), scanning speed 800-1000 mm / s (preferably 900 mm / s), layer thickness 20-40 μm (preferably 30 μm), scanning spacing 80-120 μm (preferably 100 μm), energy density E = P / (vdt) ≈ 60-80 J / mm². 3 These parameters were optimized through orthogonal experiments to ensure a relative density >99.5% and minimize keyhole (<0.5%) and incomplete fusion defects.

[0032] The scanning strategy employs an alternating 60° rotation pattern to promote isotropic microstructure and suppress columnar crystal preferential orientation. Marangoni convection intensity is modulated by the surface tension gradient (dγ / dT ≈ -0.3 mN / m·K), with increased oxygen content enhancing convection (velocities reaching 1-5 m / s), driving Al and Ti to diffuse towards the melt pool edge, forming oxide nuclei. Printed samples are rectangular (30×30 ×100 mm) or turbine blade models, with minimized support structures. During the process, melt pool fluctuations are monitored in situ (high-speed camera, 10 kHz), and power is adjusted to avoid splashing (<10 particles / s). Compared to traditional parameters, this optimization reduces crack sensitivity by 20% due to oxygen-induced liquid phase segregation buffering residual stress.

[0033] Variant processes: For large components, segmented printing (single-layer exposure time <5 s) or powder mixing (recycled powder + commercial powder 1:1) can reduce costs. This LPBF path is compatible with alloys such as IN718, with a printing efficiency >10 cm. 3 / h.

[0034] ODS in-situ formation mechanism In-situ formation of ODS depends on non-equilibrium conditions of LPBF: the peak temperature of the molten pool is 2500-3000 K, and the oxygen solubility soars (Henry's Law), reacting with Al (solubility <1 wt%), Ti, and Cr to form nano-oxides. The mechanism consists of three stages: (1) Oxygen diffusion: circulating powder oxygen (0.017-0.038 wt%) desorbs under laser heating, with a diffusion coefficient D0 ≈10 -9 m 2 / s; (2) Nucleation: Local supersaturation promotes the formation of Al2O3 or Cr2O3 crystal nuclei, and the Gibbs free energy ΔG<0; (3) Growth and composite: TiN coats the oxide nuclei to form a core-shell structure (total diameter 10-50 nm), which is semi-coherent with the γ matrix (mismatch <5%).

[0035] TEM / EDS confirmed uniform particle distribution (spacing 50-100 nm), grain boundary pinning (Zener force > 10 MPa), and inhibition of columnar crystal growth. Strengthening effects include: Orowan mechanism (τ = Gb / λ, where λ is the spacing, increasing strength by 200-300 MPa); grain boundary pinning (Hall-Petch); and synergistic precipitation of γ′ / γ″ (volume fraction 50-60%). Compared to externally added ODS, in-situ method resulted in no particle agglomeration, cleaner interfaces, and a 3-5 times increase in creep lifetime.

[0036] Heat treatment system The heat treatment regime varies depending on the alloy grade, and a standard solution treatment plus aging regime can be used accordingly. For example, for GH4169 and IN718 alloys, a standard solution treatment plus double aging regime is used. Specifically, the printed sample is solution treated in a vacuum furnace at 1100°C for 1 h to dissolve the δ phase through a martensitic process, with a cooling rate of 10 K / min to avoid stress. Double aging is then performed at 720°C for 4 h (γ″ precipitation), followed by water quenching; and at 620°C for 8 h (γ′ stabilization), followed by air cooling. The microstructure of the heat-treated state is as follows: Figure 2 As shown, the grain size decreased from 8.4 μm to 2–4 μm (EBSD), with a twinning ratio >30%. Backscattered electron microscopy (BSE) and EBSD orientation imaging revealed that nano-ODS particles (10–50 nm) were uniformly distributed within the grains and at the grain boundaries, and the ODS particles were stable without coarsening. Figure 3As shown, the room temperature tensile strength increased from approximately 1250 MPa to 1477 MPa with increasing cycle number. After five cycles, the alloy still retained 943.8 MPa at 650°C, and the elongation remained above 24%. This is because the nano-ODS particles dispersed inside and outside the matrix have a significant grain boundary pinning and microstructure refinement effect. To meet the high density requirements (density > 99.9%) of high-load components, a further heat treatment can be performed after HIP (1150°C, 100 MPa, 4 h).

[0037] This invention effectively solves the problems of powder waste, performance degradation, and difficulty in controlling ODS in the traditional nickel-based high-temperature alloy powder preparation and forming process. It proposes a sustainable manufacturing path that balances economy and performance, providing new technological support for the high-performance and low-cost manufacturing of high-temperature components for aero-engines and gas turbines.

[0038] The present invention will be further described in detail below through specific embodiments, but the scope of protection of the present invention is not limited to the following embodiments.

[0039] Example 1: Preparation of GH4169 alloy turbine blade blanks Taking GH4169 alloy as an example, the implementation steps of the method of this invention are described in detail. This implementation focuses on the preparation of LPBF and ODS strengthening of atomized powder in 5 cycles, verifying the feasibility and performance improvement of the technical solution. The entire process is completed on a laboratory scale (20kg batch), with equipment including a VIGA atomization furnace, an LPBF printer (SLM Solutions), SEM / TEM, and a universal testing machine.

[0040] Step 1: Initial Powder Preparation. Select a high-purity GH4169 ingot (100 mm in diameter, 500 mm in length), with a composition conforming to AMS 5662 standard. Place the ingot in a VIGA equipment crucible (Al2O3, capacity 30 kg) and evacuate it to 10°C. -2 Pa, Ar protection. Induction heating for melting, temperature monitoring (infrared thermometer), atomization valve opened when reaching 1570°C. Nozzle gap 0.5 mm, gas pressure 3.0 MPa, liquid flow rate 0.8 kg / s. Atomization for 10 min, collecting approximately 15 kg of powder. After cooling and sieving: fine powder yield 30% (4.5 kg, 15-60 μm), coarse powder 70% (10.5 kg, >60 μm). Oxygen content determined to be 0.012 wt%, sphericity 96.2%, D 50 =38 μm.

[0041] Step 2: First cycle atomization. Coarse powder was pulverized (ball milling, <2 h, to avoid contamination), and 40% (4.2 kg) was mixed with 60% new ingots (7.8 kg), for a total of 12 kg. Remelted to 1580°C, oxygen accumulation was reduced to 0.017 wt%. Atomization parameters remained the same, producing 3.6 kg of fine powder and 8.4 kg of coarse powder. SEM observation: Micro-oxide film on powder surface (thickness <5 nm), flowability 13 s / 50 g. Composition analysis: Nb 5.2 wt%, Ti 1.0 wt%, no inclusions.

[0042] Step 3: Subsequent cycles (2nd-5th). Repeat Step 2, returning 40% coarse powder + 60% new material each time, with a fixed melting amount of 12 kg. Oxygen content gradient: 0.022 wt% for the 2nd cycle, 0.028 wt% for the 3rd cycle, 0.033 wt% for the 4th cycle, and 0.038 wt% for the 5th cycle. Total cycle time: 8 hours, cumulative fine powder yield: 18 kg, utilization rate: 82%. Powder characteristics for the 5th cycle: sphericity 95.8%, flowability 12.5 s / 50 g, oxygen 0.038 wt%, suitable for LPBF (uniform spreading thickness <5 μm deviation). If oxygen exceeds the standard, add 0.2 wt% Zr for stabilization.

[0043] Step 4: LPBF Printing Implementation. Fifth-cycle powder was used and dried (80°C, 2 h). Equipment parameters: laser Yb fiber, wavelength 1070 nm, spot diameter 80 μm. A rectangular sample (30 × 30 × 100 mm) was printed on a 316L stainless steel substrate, preheated to 120°C. Single-layer strategy: power 250 W, speed 900 mm / s, layer thickness 30 μm, spacing 100 μm, contour scanning power 80%. Printing direction Z-axis, interlayer time 10 s. Process monitoring: melt pool depth 0.1-0.15 mm, no keyhole (X-ray CT <0.2%). Total printing time 60 min, sample density 99.7%. Splash control: gas flow rate 20 L / min, particle escape <5%.

[0044] Step 5: Heat treatment is performed. The printed sample is placed in a vacuum tube furnace (10...). -5 Solution treatment at 1100°C for 1 h followed by Ar quenching (rate 50 K / min). Then, double aging was performed: 720°C for 4 h followed by air cooling; and 620°C for 8 h followed by air cooling. After heat treatment, microstructure analysis showed: OM revealed equiaxed crystals of 2.5 μm; BSE / TEM showed uniform distribution of ODS particles (Al₂O₃-TiN, diameter 20 nm), with a γ′ volume fraction of 58%, and semi-coherent interfaces (HRTEM). EBSD showed twin boundaries Σ₃ > 25%, and texture intensity < 2 times randomness.

[0045] Step 6: Performance Verification. Room temperature tensile test (GB / T 228.1): Strength 1477 MPa, Yield 1250 MPa, Elongation 24%. 650°C tensile test: 943 MPa, Creep life (700°C, 200 MPa) >200 h (40% improvement). Oxidation test (1000°C, 100 h): Weight gain 0.3 mg / cm³. 2 It is lower than that of traditional IN718 (0.5 mg / cm³). 2 Fracture surface SEM: The dimples are fine and dense, with no ODS particle shedding. Its performance is comprehensively superior to similar materials prepared using traditional processes. Example

[0046] Unlike Example 1, the material was changed to IN625 alloy, and the cyclic atomization process was adjusted to control the final powder oxygen content at 0.025 wt%. The LPBF laser power was adjusted to 220 W, and the scanning speed was 850 mm / s. The heat treatment regime was 1150°C × 1h solution treatment + 700°C × 6h single-stage aging. The resulting component had a uniform microstructure, and its room temperature strength was approximately 15% higher than that of commercial IN625 powder LPBF parts, while maintaining good plasticity. Example

[0047] Unlike Example 1, the material was changed to IN718 alloy, the number of cycles was adjusted to 4, and the final oxygen content of the powder was controlled at 0.030 wt%. The LPBF printing parameters were set at a slightly lower energy density to verify the reliability of the optimal parameter range: laser power 230 W, scanning speed 880 mm / s, layer thickness 30 μm, and scanning spacing 100 μm.

[0048] The powder obtained by cyclic atomization had a sphericity of 95.5%, a flowability of 13.2 s / 50 g, and a relative density >99.6%. After LPBF forming, the sample was dense with no obvious unfusion defects. After heat treatment of 1100°C×1 h solution treatment + two-stage aging (720°C×4 h + 620°C×8 h), the microstructure showed equiaxed crystals of 2–4 μm, with 20–40 nm Al2O3-TiN composite particles dispersed inside.

[0049] The room temperature tensile strength was 1420 MPa, and the high temperature (650°C) strength was 910 MPa, both slightly lower than the best performance of Example 1, but still significantly higher than that of conventional IN718 prints. This example demonstrates that the ODS strengthening effect remains stable and reliable when the key parameters vary within the range defined in this invention. Example

[0050] Unlike Example 1, this example uses a six-cycle atomization route to verify the stability of in-situ ODS formation near the upper limit of oxygen content. The coarse powder return ratio is maintained at 40%, resulting in a final powder oxygen content of 0.042 wt%. LPBF printing parameters are adjusted to a slightly higher energy density: laser power 280 W, scanning speed 950 mm / s, layer thickness 30 μm, and scanning spacing 90 μm. After printing, the relative density reaches 99.7%, the molten pool morphology is stable, and there is no keyhole effect. After heat treatment, the microstructure shows a slightly finer grain size (2–3 μm), and TEM observation reveals a uniform distribution of ODS particles of approximately 15–35 nm. Using a higher energy density laser printing process, the laser energy is concentrated, which can easily lead to slight local overheating. Mechanical property tests show a room temperature tensile strength of 1460 MPa and a high-temperature strength of 935 MPa at 650°C.

[0051] Comparative Example 1 Using non-circulating atomized powder, the resulting oxygen content was less than 0.015 wt%. Due to insufficient oxygen source, the reaction between oxygen and Al / Ti in the LPBF melt pool was inadequate. TEM observation revealed an extremely small number of nano-oxide particles (average spacing >300 nm) and unstable particle size (a small number of coarse particles >80 nm), failing to form an effective ODS pinning mechanism. Grain growth was significant, with an average grain size reaching 8–10 μm. The tested room temperature tensile strength was only 1310 MPa, and the high-temperature strength at 650°C was 820 MPa, a decrease of nearly 120 MPa compared to Example 1.

[0052] Comparative Example 2 Unlike Example 1, this comparative example uses powder obtained from three cycles of atomization as the raw material for LPBF forming. The oxygen content of the powder is approximately 0.022 wt%, which is within the controllable oxygen content range set by this invention. The LPBF process parameters are consistent with those of Example 1, but the formed parts were not subjected to solution treatment and aging treatment after obtaining them.

[0053] The microstructure of the printed components showed that the grain morphology was mainly composed of incompletely eliminated slender columnar crystals, with an average grain size remaining at 6–8 μm, failing to form a uniform equiaxed crystal structure as after solution treatment. Although a small number of nano-oxide particles were formed during the printing process, their quantity was limited, their distribution was uneven, and their stability was insufficient due to the lack of subsequent heat treatment. Furthermore, the γ′ / γ″ precipitate phase was not formed, and the matrix remained in a non-equilibrium state. Mechanical property tests showed a room temperature tensile strength of 1160 MPa, a yield strength of 960 MPa, a high-temperature tensile strength of 760 MPa at 650°C, and an elongation of 10–12%. Compared to Example 1, the strength and plasticity of this comparative example were significantly reduced.

[0054] The above embodiments fully verify the effectiveness, reliability, and universality of the technical solutions of the present invention. Unless otherwise stated, if any technical solution disclosed in the present invention specifies a numerical range, then the disclosed numerical range is a preferred numerical range. Any person skilled in the art should understand that the preferred numerical range is merely a range of numerical values ​​among many implementable values ​​that have a relatively obvious or representative technical effect. Because there are many numerical values, it is impossible to exhaustively list them all. Therefore, the present invention discloses only some numerical values ​​to illustrate the technical solutions of the present invention. Furthermore, the numerical values ​​listed above should not constitute a limitation on the scope of protection of the present invention.

[0055] Furthermore, if the present invention discloses or relates to mutually fixedly connected components or structural parts, then unless otherwise stated, a fixed connection can be understood as: a detachable fixed connection (e.g., using bolts or screws), or a non-detachable fixed connection (e.g., riveting, welding). Of course, mutually fixed connections can also be replaced by an integral structure (e.g., manufactured using a casting process) (except where it is obviously impossible to use an integral molding process).

[0056] Furthermore, unless otherwise stated, the terms used to indicate positional relationships or shapes in any of the technical solutions disclosed in this invention include states or shapes that are similar to, analogous to, or close to those states or shapes. Any component provided by this invention can be assembled from multiple individual components or can be a single component manufactured using a one-piece molding process.

[0057] The above embodiments are merely illustrative examples to clearly illustrate the present invention and are not intended to limit the implementation. Those skilled in the art will recognize that other variations or modifications can be made based on the above description. It is neither necessary nor possible to exhaustively list all embodiments here. However, obvious variations or modifications derived therefrom are still within the scope of protection of this invention.

Claims

1. A method for preparing an oxide dispersion-strengthened nickel-based superalloy, characterized in that, Includes the following steps: S1. Circulating atomization powder production: Using nickel-based superalloy ingots as initial raw materials, vacuum induction gas atomization powder production is carried out to obtain fine powder that can be used for LPBF and coarse powder that cannot be used directly; the obtained coarse powder is mixed with new nickel-based superalloy ingots in a predetermined mass ratio and used as raw materials for vacuum induction gas atomization powder production again. This cyclic process is repeated. By controlling the number of cycles and the return ratio of coarse powder, the oxygen content of the final fine powder is controlled within the range of 0.015 wt% to 0.060 wt%. S2, Selective Laser Melting: The circulating atomized fine powder obtained in step S1 is used as raw material and selective laser melting is performed under a protective atmosphere. Nano-oxide dispersion reinforced particles are generated in situ in the molten pool through optimized laser process parameters. S3. Heat treatment for microstructure control: The formed part obtained in step S2 is subjected to standard solution treatment and aging treatment to obtain a nickel-based high-temperature alloy component with fine grain structure and dispersed nano-oxide.

2. The method for preparing oxide dispersion strengthened nickel-based superalloys according to claim 1, characterized in that: In step S2, the process parameters for laser selective melting forming are: laser power 200-300 W, scanning speed 800-1000 mm / s, layer thickness 20-40 μm, and scanning spacing 80-120 μm.

3. The method for preparing oxide dispersion strengthened nickel-based superalloys according to claim 1 or 2, characterized in that: In step S1, the process parameters for vacuum induction gas atomization powder production are: vacuum degree not higher than 10. -2 Pa, melting temperature 1550-1600°C, atomizing gas is Ar, gas pressure 2.5-3.5 MPa.

4. The method for preparing oxide dispersion strengthened nickel-based superalloys according to claim 1 or 2, characterized in that, The return rate of the coarse powder is 30%-50%, and the number of cycles is 3 to 6.

5. The method for preparing oxide dispersion strengthened nickel-based superalloys according to claim 1 or 2, characterized in that, In step S2, an interlayer rotation scanning strategy is adopted, with a rotation angle of 60° to 90°.

6. The method for preparing oxide dispersion strengthened nickel-based superalloys according to claim 1 or 2, characterized in that, In step S2, the in-situ generated nano-oxide dispersion-reinforced particles have a particle size of 10-50 nm and their composition is at least one of Cr2O3 and Al2O3-TiN composite structures.

7. The method for preparing oxide dispersion strengthened nickel-based superalloys according to claim 1 or 2, characterized in that, The nickel-based superalloy is GH4169, IN625, or IN718.

8. The method for preparing oxide dispersion strengthened nickel-based superalloys according to claim 7, characterized in that, For GH4169 and IN718 alloys, the standard solution treatment regime is 1100°C for 1 hour; the standard aging treatment is a two-stage aging process, which involves holding at 720°C for 4 hours and then holding at 620°C for 8 hours.

9. The method for preparing oxide dispersion strengthened nickel-based superalloys according to claim 7, characterized in that, For IN625 alloy, the standard solution treatment regime is 1150°C for 1 hour and the standard aging treatment regime is 700°C for 6 hours.

10. An oxide dispersion-strengthened nickel-based superalloy component, characterized in that, The component is prepared by any one of claims 1 to 9, wherein the grain size is 2-4 μm and nano-oxide particles with a particle size of 10-50 nm are dispersedly distributed inside and outside its matrix.

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