A nickel-based superalloy block and a continuous solution annealing method thereof

Abstract

The present application relates to nickel-based superalloy heat treatment technical field, specifically disclose a kind of nickel-based superalloy block and its continuous solution annealing method, the nickel-based superalloy is Nb-containing γ″ precipitation strengthening type nickel-based superalloy, alloy block thickness H is 80~150mm.This method uses step beam type continuous annealing furnace, including: S1 preheating uniform temperature, with ≤5 ℃ / min to 600~650 ℃ after heat preservation, heat preservation time is determined according to t_h=k×H, k takes 0.8~1.2min / mm;S2 transition heating, with 8~15 ℃ / min to solution temperature;S3 sub-solution limit solution, at less than the alloy δ phase dissolution temperature 5~30 ℃ heat preservation 30~90min;S4 staged cooling, first with 10~20 ℃ / min slow cooling to 600~650 ℃, then with ≥50 ℃ / min fast cooling to 200 ℃ below furnace air cooling, the whole process is carried out under ammonia decomposition gas or high-purity nitrogen protection.This method solves the problem of thermal stress cracking, structure inhomogeneity and precipitated phase control in the continuous solution annealing of large cross-section block.

Description

Technical Field

[0001] This invention relates to the field of heat treatment technology for nickel-based superalloys, and in particular to a nickel-based superalloy bulk material and its continuous solution annealing method. Background Technology

[0002] Against the backdrop of rapid upgrades and replacements in global aero-engines, gas turbines, nuclear power equipment, and advanced spacecraft, high-temperature alloy structural components have become core foundational components determining the lifespan, reliability, and safety of equipment. Nickel-based high-temperature alloys, with their excellent high-temperature strength, creep resistance, fatigue resistance, oxidation resistance, and corrosion resistance, are widely used in critical components such as engine hot-end parts and gas turbine blades and discs. After plastic forming processes such as forging and rolling, high-temperature alloys may exhibit residual stress, uneven microstructure, localized mixed grains, or work hardening. Without subsequent heat treatment to relieve stress and regulate microstructure, these components are highly susceptible to deformation, cracking, and fatigue failure under extreme conditions such as high temperature and high pressure. Therefore, solution annealing is an indispensable and crucial process in the transformation of high-temperature alloys from raw materials to finished products.

[0003] Ni-based superalloys with Nb precipitation strengthening, exemplified by GH4169 (corresponding to Inconel 718), primarily have γ″ phase (Ni3Nb, body-centered tetragonal DO) as their strengthening phase. 22 The alloy contains a γ″ phase (Ni3(Al,Ti), face-centered cubic L12 structure) and a γ′ phase (Ni3Nb, orthorhombic structure) at grain boundaries, and a δ phase (Ni3Nb, orthorhombic structure) at grain boundaries. Solution annealing not only requires the elimination of residual stress and work hardening, but also the precise control of the dissolution degree of the γ″ and γ′ strengthening phases, as well as the morphology and distribution of the δ phase. Sufficient dissolution of the γ″ phase is a prerequisite for subsequent age hardening to achieve high strength; while the δ phase, when present in appropriate amounts as isolated granules at grain boundaries, can effectively pin the grain boundaries and inhibit grain growth, playing a crucial role in obtaining a uniform, fine-grained microstructure. However, the dissolution temperature of the δ phase is significantly affected by the alloy composition (especially the Nb content) and the hot working history, typically fluctuating within the range of 980–1030 °C. Once the solution temperature exceeds the actual dissolution limit temperature of the δ phase in that batch of alloy, the δ phase completely dissolves, the grain boundaries lose pinning, leading to rapid grain coarsening and severe mixed crystals, causing a sharp decrease in the material's plasticity. Furthermore, if the δ phase remains in the rapid precipitation sensitive region (approximately 850–950 °C, with a nose temperature of approximately 927 °C on the TTT curve) for too long, the δ phase is prone to precipitate semi-continuously or in a network pattern along the grain boundaries, leading to grain boundary embrittlement and severely impairing the material's toughness. Therefore, the solution annealing process window for this type of alloy is extremely narrow, imposing extremely stringent requirements on temperature uniformity and time control.

[0004] Currently, there are still significant shortcomings in the domestic high-temperature alloy annealing technology and equipment. Solution annealing of nickel-based high-temperature alloy billets, square blocks, and discs generally adopts intermittent operation methods such as box furnaces or bogie hearth furnaces, which have the following prominent problems: (1) It is difficult to accurately control the heating rate and temperature uniformity. During the heating process of large cross-section blocks (thickness ≥ 80 mm), the temperature difference between the surface and the core is large, the thermal stress is concentrated, and cracks are easily generated; (2) Intermittent production includes the entire cycle of loading the furnace, heating, holding, and cooling, which takes several hours to tens of hours. The production efficiency is low, and the consistency of the microstructure and properties between batches is poor, which makes it difficult to meet the requirements of high reliability and traceability for aerospace-grade parts; (3) The waste heat of traditional furnaces cannot be recovered, resulting in high energy consumption and large footprint, which is significantly different from the requirements of green manufacturing trends. There are already publicly available continuous annealing technologies, such as CN112813369B, which discloses a high-strength, high-elasticity, and high-plasticity nickel-based high-temperature alloy strip and its preparation process. The processing object is strip rather than large cross-section blocks, which cannot be directly transferred to the processing of thick parts. Other literature on continuous annealing furnaces mainly focuses on furnace structure or load-bearing methods, and has not yet provided repeatable temperature gradient control and phase structure regulation parameters for GH4169 blocks with a thickness of 80~150mm. Therefore, how to achieve crack-free solution annealing, precise control of microstructure uniformity, and efficient production of large-section blocks under continuous conveying conditions has become an urgent technical problem to be solved. Summary of the Invention

[0005] In view of the above-mentioned deficiencies of the prior art, the technical problem to be solved by the present invention is: for Nb-containing γ″ precipitation-strengthened nickel-based superalloy large cross-section blocks (thickness H=80~150mm), how to systematically solve the contradiction between thermal stress cracking, precipitate control and production efficiency in a walking beam continuous annealing furnace, and provide a continuous solution annealing method that can achieve continuous production, effectively control the temperature gradient during heating and cooling to prevent thermal stress cracking, and at the same time ensure that the γ″ phase is fully dissolved and the δ phase is moderately retained to pin grain boundaries and suppress grain coarsening and grain boundary embrittlement.

[0006] To achieve the above objectives, this invention provides a nickel-based superalloy block and its continuous solution annealing method. The nickel-based superalloy is a Nb-containing γ″ precipitation-strengthened nickel-based superalloy, and the thickness H of the alloy block is 80–150 mm. The method employs a walking beam continuous annealing furnace, which sequentially includes a preheating zone, a transition heating zone, a solution annealing zone, a slow cooling zone, and a rapid cooling zone along the conveying direction, and includes the following steps: S1, Preheating and Uniform Temperature Measurement: The alloy block is placed in the preheating zone and heated from room temperature to 600-650℃ at a heating rate of ≤5℃ / min. Then, uniform temperature measurement and holding are performed within this temperature range. The uniform temperature measurement and holding time is t_h=k×H, where k is the holding time coefficient, which is 0.8~1.2min / mm, and H is the thickness of the alloy block in mm. At the end of preheating, the core temperature of the alloy block is not lower than 580℃, and the temperature difference between the surface and the core is not greater than 80℃. S2, Transition heating: The preheated and homogenized alloy block is sent into the transition heating zone and heated to the solution temperature at a heating rate of 8-15℃ / min. S3, Subsolution Limit: The alloy block is placed into the solution zone and held at the solution temperature for 30 to 90 minutes; the solution temperature is 5 to 30°C lower than the δ phase dissolution temperature of the alloy. S4, graded cooling: After the solution treatment and heat preservation are completed, the alloy block is sent to the cooling zone. First, it is cooled to 600-650°C in the slow cooling zone at a cooling rate of 10-20°C / min, and then cooled to below 200°C in the fast cooling zone at an average surface cooling rate of ≥50°C / min. After being taken out of the furnace, it is air-cooled to room temperature. The above steps are carried out in a protective atmosphere, which is ammonia decomposition gas or nitrogen with a purity ≥99.995%, the furnace pressure is 50-200 Pa, the dew point is ≤-60℃, and the oxygen content is ≤20 ppm.

[0007] In a preferred embodiment of the present invention, the nickel-based superalloy is GH4169 or Inconel 718.

[0008] In another preferred embodiment of the present invention, the heating rate in step S1 is 3°C / min, the temperature for uniform heat preservation is 650°C, and the heat preservation time coefficient k is 1.0 to 1.1 min / mm.

[0009] In another preferred embodiment of the present invention, the solution temperature in steps S2 and S3 is 985–1010 °C.

[0010] In another preferred embodiment of the present invention, the heating rate in step S2 is 10°C / min.

[0011] In another preferred embodiment of the present invention, the heat preservation time in step S3 is determined according to the thickness: when H≤100mm, the heat preservation time is 30-60min; when 100<H≤130mm, the heat preservation time is 45-75min; when 130<H≤150mm, the heat preservation time is 60-90min.

[0012] In another preferred embodiment of the present invention, the cooling rate of the slow cooling zone in step S4 is 15°C / min, cooling to 620°C; the average surface cooling rate of the fast cooling zone is 60°C / min, cooling to 140°C.

[0013] In another preferred embodiment of the present invention, the stepping cycle of the walking beam continuous annealing furnace is 60 to 120 seconds per step, which is adjusted according to the thickness H of the alloy block, and the dwell time accuracy of each temperature zone is ±10%.

[0014] In another preferred embodiment of the present invention, the protective atmosphere is ammonia decomposition gas, and the furnace is equipped with a hydrogen concentration monitoring and explosion-proof pressure relief device; the rapid cooling section adopts a forced circulation nitrogen injection or aerosol cooling device.

[0015] In another preferred embodiment of the present invention, the nickel-based superalloy bulk material treated by the continuous solution annealing method has a grain size of ASTM 7 to 9, an average grain size of 18 to 25 μm, and the δ phase (i.e., Ni3Nb) is distributed in the form of isolated particles at the grain boundaries, with an average grain boundary coverage of no more than 25%.

[0016] Compared with the prior art, the present invention has the following beneficial effects: 1. Effectively reduces the risk of thermal stress cracking. This invention breaks down the preheating and heating processes into multiple stages. First, a slow heating rate combined with a holding time linked to the thickness ensures uniform and thorough heating of the large-section alloy block, keeping the temperature difference between the core and surface within a small range and reducing the tendency for thermal stress cracking. The transition stage moderately increases the heating rate after the material enters the plastic temperature range, balancing safety and efficiency. The slow cooling stage in the staged cooling prioritizes the release of thermal stress from the initial cooling phase, further reducing the cracking risk of thick-section materials during the cooling process. Compared to traditional intermittent box furnaces, the crack incidence rate of the alloy block after annealing is significantly reduced.

[0017] 2. Improved microstructure uniformity. This invention sets the solution temperature within the sub-solution limit range of the δ phase, allowing the γ″ strengthening phase to fully dissolve while retaining a suitable amount of isolated granular δ phase at the grain boundaries. These δ phases pin the grain boundaries during high-temperature holding, inhibiting abnormal grain growth. The linkage between the stepper beam cycle and the alloy block thickness ensures that alloy blocks of different thicknesses achieve matched heat penetration and holding times in each temperature zone, controlling the grain size difference between different specifications within a narrow range after annealing.

[0018] 3. Strength and plasticity are well-matched. The staged cooling system, through a controlled slow cooling rate, allows the material to pass through the δ-phase precipitation-sensitive region, reducing embrittlement caused by semi-continuous or network precipitation at grain boundaries. The subsequent rapid cooling stage fixes the supersaturated solid solution, providing a driving force for age-hardening. After standard aging, the alloy block maintains high strength and elongation at both room temperature and high temperature, with minimal performance differences between the core and surface layers in large cross-sections. The following will further explain the concept, specific structure, and technical effects of this invention in conjunction with the accompanying drawings, to fully understand the purpose, features, and effects of this invention. Attached Figure Description

[0019] Figure 1 The image shows the metallographic structure of the GH4169 alloy block obtained in Example 2 of this invention after continuous solution annealing and aging treatment.

[0020] Figure 2 The image shows the SEM-BSE morphology of the δ phase of the GH4169 alloy block obtained in Example 2 of this invention after continuous solution annealing and aging treatment. Detailed Implementation

[0021] The following specific examples illustrate the implementation of the present invention. Those skilled in the art can easily understand other advantages and effects of the present invention from the content disclosed in this specification. The present invention can also be implemented or applied through other different specific embodiments, and various details in this specification can also be modified or changed based on different viewpoints and applications without departing from the spirit of the present invention. It should be noted that, unless otherwise specified, the following embodiments and features described therein can be combined with each other.

[0022] Some exemplary embodiments of the invention have been described for illustrative purposes. It should be understood that the invention may be implemented in other ways not specifically shown in the accompanying drawings.

[0023] This invention provides a nickel-based superalloy bulk material and its continuous solution annealing method. The nickel-based superalloy is a Nb-containing γ″ precipitation-strengthened nickel-based superalloy, typically represented by GH4169 or Inconel 718. The main strengthening phases of this type of alloy are the γ″ phase (Ni3Nb) and the γ′ phase (Ni3(Al,Ti)), with a δ phase (i.e., Ni3Nb) present at the grain boundaries. The goal of solution annealing is to fully dissolve the γ″ phase while retaining an appropriate amount of isolated granular δ phase to pin the grain boundaries and prevent abnormal grain growth at high temperatures. The thickness H of the alloy bulk material is 80–150 mm, constituting a large cross-section bulk material, and the temperature following behavior between the core and surface differs significantly during heat treatment.

[0024] This method employs a walking beam continuous annealing furnace, with preheating, transition heating, solution treatment, slow cooling, and rapid cooling zones sequentially arranged along the conveying direction. The entire process is carried out under a protective atmosphere. The protective atmosphere is ammonia decomposition gas (75% H2 + 25% N2) or nitrogen with a purity of not less than 99.995%. The furnace is maintained at a slightly positive pressure of 50–200 Pa, with a dew point not exceeding -60°C and an oxygen content not exceeding 20 ppm. When using ammonia decomposition gas, the furnace is equipped with hydrogen concentration monitoring and explosion-proof pressure relief devices.

[0025] During the preheating and homogenization stage, after the alloy block enters the preheating zone, it is heated from room temperature to 600–650°C at a rate not exceeding 5°C / min. In the initial heating phase of large cross-section blocks, heat conduction from the surface to the core takes time. Excessive heating can cause a significant temperature difference between the surface and the core, resulting in thermal stress sufficient to cause material cracking. The purpose of using a low heating rate is to allow the core sufficient time to follow the surface temperature change, controlling the core-surface temperature difference within a safe range. After reaching the target temperature, the alloy block undergoes homogenization and holding within this temperature range. The holding time t_h is determined by t_h = k × H, where H is the alloy block thickness (mm), and k is the holding time coefficient, taken as 0.8–1.2 min / mm. The purpose of holding is to allow the core temperature to continue approaching the surface temperature, ultimately ensuring that the core temperature is not lower than 580°C, and the temperature difference between the surface and the core is not greater than 80°C. Once this state is achieved, the material has entered the plastic temperature range, enabling plastic deformation to coordinate the thermal stress generated during subsequent heating. The total residence time in the preheating zone is the sum of the heating sub-stage time and the temperature equalization and heat preservation time.

[0026] During the transition heating stage, the alloy block, after preheating and homogenization, enters the transition heating zone and is heated to the solution temperature at a rate of 8–15 °C / min. At this point, the core of the material has sufficient temperature, and its plastic deformation capacity is significantly improved, allowing it to withstand a faster heating rate than in the preheating stage without cracking. Controlling the heating rate within this moderate range aims to shorten the transition time from the preheating temperature to the solution temperature while ensuring safety and balancing efficiency.

[0027] During the sub-solution stage, the alloy block is held at a temperature within the solution zone. The solution temperature is 5–30°C lower than the dissolution temperature of the δ phase in this batch of alloys. For GH4169 alloy, the preferred solution temperature is 985–1010°C. The reason for setting the solution temperature within the sub-solution zone below the δ phase dissolution temperature is that exceeding the dissolution temperature completely would cause the δ phase at the grain boundaries to dissolve entirely into the matrix. Grains that have lost their pinning effect would grow unrestrained at high temperatures, forming a mixed-grain structure and severely reducing plasticity. On the other hand, if the solution temperature is too low, the γ″ strengthening phase cannot dissolve sufficiently, and subsequent aging cannot fully elute, resulting in insufficient strength. Precisely controlling the temperature within the sub-solution zone achieves two goals simultaneously—the γ″ phase fully dissolves into the matrix, while isolated granular δ phase remains at the grain boundaries. The total holding time ranges from 30 to 90 minutes, determined in segments based on the thickness of the alloy block: 30–60 minutes for thicknesses not exceeding 100 mm, 45–75 minutes for thicknesses between 100 mm and 130 mm, and 60–90 minutes for thicknesses between 130 mm and 150 mm. This segmentation is because the thicker the cross-section, the longer the time required for the solution temperature to conduct from the surface to the core and for the γ″ phase to fully dissolve.

[0028] In the staged cooling phase, after solution treatment and holding, the alloy block enters the cooling zone. First, it is cooled to 600–650°C in the slow cooling zone at a rate of 10–20°C / min, then cooled to below 200°C in the rapid cooling zone at an average surface cooling rate of no less than 50°C / min, and finally air-cooled to room temperature. The staged cooling design aims to simultaneously avoid two failure modes. If rapid cooling occurs immediately after solution treatment, the core and surface of the large cross-section block will shrink asynchronously, generating large tensile stress internally and easily leading to quenching cracks. The slower cooling rate in the slow cooling section allows the material to shrink uniformly throughout, effectively releasing the thermal stress of this stage. On the other hand, if the material remains in the sensitive temperature range for rapid δ-phase precipitation (approximately 850–950°C) for too long, the δ-phase will connect along the grain boundaries to form a semi-continuous or even network-like brittle film, leading to grain boundary embrittlement. The cooling rate in the slow cooling section is designed to pass through this sensitive zone at a controlled rate, with insufficient dwell time to form a continuous network precipitation. Once the temperature drops to 600–650℃, it has left the temperature window for the large-scale precipitation of the δ phase. At this point, increasing the cooling rate to an average of no less than 50℃ / min on the surface can rapidly cool the material to below 200℃, fixing the supersaturated solid solution state formed at high temperatures and providing a driving force for subsequent aging strengthening. The rapid cooling section can be achieved using forced circulation nitrogen injection or aerosol cooling devices, with the cooling rate measured by thermocouples on the alloy block surface.

[0029] The walking beam's stepping cycle is 60–120 seconds per step, adjusted according to the alloy block thickness H, with the dwell time accuracy of each temperature zone controlled within ±10%. This linkage aims to ensure that each alloy block receives a uniform temperature holding time matching its thickness in the preheating zone, and sufficient dissolution time for the strengthening phase in the solution treatment zone. Furthermore, when alloy blocks of different thicknesses share the same continuous furnace production line, their heat treatment states remain highly consistent.

[0030] Example 1 This embodiment provides a nickel-based superalloy bulk material and its continuous solution annealing method.

[0031] GH4169 alloy forgings produced in the same batch using vacuum induction melting + vacuum arc remelting (VIM+VAR) have the following nominal chemical composition (mass percentage): Ni 52.8%, Cr 19.0%, Nb 5.3%, Mo 3.0%, Ti 0.9%, Al 0.5%, Fe balance. Differential scanning calorimetry (DSC) determined that the δ-phase dissolution termination temperature of this batch of alloy is approximately 1023℃. The forging dimensions are 80mm × 200mm × 500mm, with a thickness H = 80mm.

[0032] A walking beam continuous annealing furnace is employed, with preheating, transition heating, solution treatment, slow cooling, and rapid cooling zones sequentially arranged along the conveying direction. The temperature and atmosphere of each zone are independently controlled. The entire furnace is purged with ammonia decomposition gas (75% H2 + 25% N2), with a dew point of -62℃. The furnace is maintained at a slightly positive pressure of 100 Pa, with an oxygen content ≤15 ppm. The furnace is equipped with hydrogen concentration monitoring and explosion-proof pressure relief devices.

[0033] S1, Preheating and Isostatication: Place the alloy block on the walking beam and send it into the preheating zone. Increase the temperature from room temperature to 650℃ at a rate of 3℃ / min, taking approximately 208 minutes. After reaching 650℃, maintain the temperature at this isothermal level. The holding time coefficient k is set to 1.0, and the holding time t_h = 1.0 × 80 = 80 minutes. The total residence time in the preheating zone is the sum of the heating time and the isothermal holding time, approximately 288 minutes. Drill a 2mm diameter, 40mm deep blind hole in the center of the alloy block and insert a type K armored thermocouple to monitor the core temperature. The surface temperature is monitored by spot-welded thermocouples. At the end of preheating, the core temperature should not be lower than 580℃, and the temperature difference between the surface and the core should not exceed 80℃.

[0034] S2, Transition Heating: The preheated alloy block is fed into the transition heating zone by the walking beam and heated from 650℃ to 1000℃ at a heating rate of 10℃ / min for about 35 minutes.

[0035] S3, Sub-solution limit: The alloy block is placed in the solution zone and held at 1000℃ for 45 minutes. This solution temperature is approximately 23℃ lower than the final dissolution temperature of the δ phase of this batch of alloys, falling within the sub-solution limit range.

[0036] S4, Staged Cooling: After solution treatment and holding, the alloy block is placed in the cooling zone. First, it is cooled from 1000℃ to 620℃ in the slow cooling zone at a cooling rate of 15℃ / min, for approximately 25 minutes. Then, it enters the rapid cooling zone, where a forced-circulation nitrogen injection device is used to cool it from 620℃ to 140℃ at an average surface cooling rate of 60℃ / min, for approximately 8 minutes. The cooling rate is measured by thermocouples on the alloy block surface. Finally, it is removed from the furnace and air-cooled to room temperature.

[0037] The stepping cycle of the walking beam is set to 60 seconds per step based on the thickness of the alloy block, and the dwell time accuracy of each temperature zone is controlled within ±10%.

[0038] Example 2 This embodiment provides a nickel-based superalloy bulk material and its continuous solution annealing method.

[0039] GH4169 alloy forgings produced in the same batch using vacuum induction melting + vacuum arc remelting (VIM+VAR) have the following nominal chemical composition (mass percentage): Ni 52.8%, Cr 19.0%, Nb 5.3%, Mo 3.0%, Ti 0.9%, Al 0.5%, Fe balance. Differential scanning calorimetry (DSC) determined that the δ-phase dissolution termination temperature of this batch of alloy is approximately 1023℃. The forging dimensions are 120mm × 200mm × 500mm, with a thickness H = 120mm.

[0040] A walking beam continuous annealing furnace is employed, with preheating, transition heating, solution treatment, slow cooling, and rapid cooling zones sequentially arranged along the conveying direction. The temperature and atmosphere of each zone are independently controlled. The entire furnace is purged with ammonia decomposition gas (75% H2 + 25% N2), with a dew point of -62℃. The furnace is maintained at a slightly positive pressure of 100 Pa, with an oxygen content ≤15 ppm. The furnace is equipped with hydrogen concentration monitoring and explosion-proof pressure relief devices.

[0041] S1, Preheating and Isostatication: The alloy block is placed on the walking beam and sent into the preheating zone. The temperature is increased from room temperature to 650℃ at a rate of 3℃ / min, taking approximately 208 minutes. After reaching 650℃, isostatic holding is performed at this temperature, with a holding time coefficient k of 1.0. The isostatic holding time t_h = 1.0 × 120 = 120 minutes. The total residence time in the preheating zone is the sum of the heating time and the isostatic holding time, approximately 328 minutes. A blind hole with a diameter of 2mm and a depth of 60mm is drilled in the center of the alloy block, and a type K-armored thermocouple is inserted to monitor the core temperature. The surface temperature is monitored by spot-welded thermocouples. At the end of preheating, the core temperature should not be lower than 580℃, and the temperature difference between the surface and the core should not exceed 80℃. The monitored core temperature reached 593℃ at the end of S1, with a maximum core-to-surface temperature difference of 73℃.

[0042] S2, transitional heating: exactly the same as in Example 1.

[0043] S3, Sub-solution limit: The alloy block was placed in the solution treatment zone and held at 1000℃ for 60 minutes. This solution temperature is approximately 23℃ lower than the final dissolution temperature of the δ phase of this batch of alloys, falling within the sub-solution limit range. The lowest core temperature during the solution treatment stage was monitored to reach 986℃.

[0044] S4, graded cooling: exactly the same as in Example 1.

[0045] The stepping cycle of the walking beam is set to 80 seconds per step based on the thickness of the alloy block, and the dwell time accuracy of each temperature zone is controlled within ±10%.

[0046] Example 3 This embodiment provides a nickel-based superalloy bulk material and its continuous solution annealing method.

[0047] GH4169 alloy forgings produced in the same batch using vacuum induction melting + vacuum arc remelting (VIM+VAR) have the following nominal chemical composition (mass percentage): Ni 52.8%, Cr 19.0%, Nb 5.3%, Mo 3.0%, Ti 0.9%, Al 0.5%, Fe balance. Differential scanning calorimetry (DSC) determined that the δ-phase dissolution termination temperature of this batch of alloy is approximately 1023℃. The forging dimensions are 150mm × 200mm × 500mm, with a thickness H = 150mm.

[0048] A walking beam continuous annealing furnace is employed, with preheating, transition heating, solution treatment, slow cooling, and rapid cooling zones sequentially arranged along the conveying direction. The temperature and atmosphere of each zone are independently controlled. The entire furnace is purged with ammonia decomposition gas (75% H2 + 25% N2), with a dew point of -62℃. The furnace is maintained at a slightly positive pressure of 100 Pa, with an oxygen content ≤15 ppm. The furnace is equipped with hydrogen concentration monitoring and explosion-proof pressure relief devices.

[0049] S1, Preheating and Isostatication: The alloy block is placed on the walking beam and sent into the preheating zone. The temperature is increased from room temperature to 650℃ at a rate of 3℃ / min, taking approximately 208 minutes. After reaching 650℃, isostatic holding is performed at this temperature, with a holding time coefficient k of 1.1. The isostatic holding time t_h = 1.1 × 150 = 165 minutes. The total residence time in the preheating zone is the sum of the heating time and the isostatic holding time, approximately 373 minutes. A blind hole with a diameter of 2mm and a depth of 75mm is drilled in the center of the alloy block, and a type K-armored thermocouple is inserted to monitor the core temperature. The surface temperature is monitored by spot-welded thermocouples. At the end of preheating, the core temperature should not be lower than 580℃, and the temperature difference between the surface and the core should not exceed 80℃.

[0050] S2, transitional heating: exactly the same as in Example 1.

[0051] S3, Sub-solution limit: The alloy block is placed in the solution zone and held at 1000℃ for 75 minutes. This solution temperature is about 23℃ lower than the final dissolution temperature of the δ phase of this batch of alloys, and is within the sub-solution limit range.

[0052] S4, graded cooling: exactly the same as in Example 1.

[0053] The stepping cycle of the walking beam is set to 100 seconds per step based on the thickness of the alloy block, and the dwell time accuracy of each temperature zone is controlled within ±10%.

[0054] Comparative Example 1 This comparative example provides a nickel-based superalloy bulk material and its continuous solution annealing method.

[0055] GH4169 alloy forgings produced in the same batch using vacuum induction melting + vacuum arc remelting (VIM+VAR) have the following nominal chemical composition (mass percentage): Ni 52.8%, Cr 19.0%, Nb 5.3%, Mo 3.0%, Ti 0.9%, Al 0.5%, Fe balance. Differential scanning calorimetry (DSC) determined that the δ-phase dissolution termination temperature of this batch of alloy is approximately 1023℃. The forging dimensions are 120mm × 200mm × 500mm, with a thickness H = 120mm.

[0056] A walking beam continuous annealing furnace is employed, with preheating, transition heating, solution treatment, slow cooling, and rapid cooling zones sequentially arranged along the conveying direction. The temperature and atmosphere of each zone are independently controlled. The entire furnace is purged with ammonia decomposition gas (75% H2 + 25% N2), with a dew point of -62℃. The furnace is maintained at a slightly positive pressure of 100 Pa, with an oxygen content ≤15 ppm. The furnace is equipped with hydrogen concentration monitoring and explosion-proof pressure relief devices.

[0057] The alloy block is placed on a walking beam and fed into the furnace. The temperature is continuously increased from room temperature to 1000℃ at a heating rate of 12℃ / min, with a heating time of approximately 81 minutes. A blind hole with a diameter of 2mm and a depth of 60mm is drilled in the center of the alloy block, and a type K armored thermocouple is inserted to monitor the core temperature. The surface temperature is monitored by spot-welded thermocouples.

[0058] S3, Sub-solution limit: The alloy block is placed in the solution zone and held at 1000℃ for 60 minutes. This solution temperature is about 23℃ lower than the final dissolution temperature of the δ phase of this batch of alloys, and is within the sub-solution limit range.

[0059] S4, graded cooling: exactly the same as in Example 1.

[0060] The stepping cycle of the walking beam is set to 80 seconds per step based on the thickness of the alloy block, and the dwell time accuracy of each temperature zone is controlled within ±10%.

[0061] Comparative Example 2 This comparative example provides a nickel-based superalloy bulk material and its continuous solution annealing method.

[0062] GH4169 alloy forgings produced in the same batch using vacuum induction melting + vacuum arc remelting (VIM+VAR) have the following nominal chemical composition (mass percentage): Ni 52.8%, Cr 19.0%, Nb 5.3%, Mo 3.0%, Ti 0.9%, Al 0.5%, Fe balance. Differential scanning calorimetry (DSC) determined that the δ-phase dissolution termination temperature of this batch of alloy is approximately 1023℃. The forging dimensions are 120mm × 200mm × 500mm, with a thickness H = 120mm.

[0063] A walking beam continuous annealing furnace is employed, with preheating, transition heating, solution treatment, slow cooling, and rapid cooling zones sequentially arranged along the conveying direction. The temperature and atmosphere of each zone are independently controlled. The entire furnace is purged with ammonia decomposition gas (75% H2 + 25% N2), with a dew point of -62℃. The furnace is maintained at a slightly positive pressure of 100 Pa, with an oxygen content ≤15 ppm. The furnace is equipped with hydrogen concentration monitoring and explosion-proof pressure relief devices.

[0064] S1, Preheating and Isostatication: The alloy block is placed on the walking beam and sent into the preheating zone. The temperature is increased from room temperature to 650℃ at a rate of 3℃ / min, taking approximately 208 minutes. After reaching 650℃, isostatic holding is performed at this temperature, with a holding time coefficient k of 0.4. The isostatic holding time t_h = 0.4 × 120 = 48 minutes. The total residence time in the preheating zone is the sum of the heating time and the isostatic holding time, approximately 256 minutes. A blind hole with a diameter of 2mm and a depth of 60mm is drilled in the center of the alloy block, and a type K-armored thermocouple is inserted to monitor the core temperature. The surface temperature is monitored by spot-welded thermocouples.

[0065] S2, transitional heating: exactly the same as in Example 1.

[0066] S3, Sub-solution limit: The alloy block is placed in the solution zone and held at 1000℃ for 60 minutes. This solution temperature is about 23℃ lower than the final dissolution temperature of the δ phase of this batch of alloys, and is within the sub-solution limit range.

[0067] S4, graded cooling: exactly the same as in Example 1.

[0068] The stepping cycle of the walking beam is set to 80 seconds per step based on the thickness of the alloy block, and the dwell time accuracy of each temperature zone is controlled within ±10%.

[0069] Comparative Example 3 This comparative example provides a nickel-based superalloy bulk material and its continuous solution annealing method.

[0070] GH4169 alloy forgings produced in the same batch using vacuum induction melting + vacuum arc remelting (VIM+VAR) have the following nominal chemical composition (mass percentage): Ni 52.8%, Cr 19.0%, Nb 5.3%, Mo 3.0%, Ti 0.9%, Al 0.5%, Fe balance. Differential scanning calorimetry (DSC) determined that the δ-phase dissolution termination temperature of this batch of alloy is approximately 1023℃. The forging dimensions are 120mm × 200mm × 500mm, with a thickness H = 120mm.

[0071] A walking beam continuous annealing furnace is employed, with preheating, transition heating, solution treatment, slow cooling, and rapid cooling zones sequentially arranged along the conveying direction. The temperature and atmosphere of each zone are independently controlled. The entire furnace is purged with ammonia decomposition gas (75% H2 + 25% N2), with a dew point of -62℃. The furnace is maintained at a slightly positive pressure of 100 Pa, with an oxygen content ≤15 ppm. The furnace is equipped with hydrogen concentration monitoring and explosion-proof pressure relief devices.

[0072] S1, Preheating and Uniform Temperature Measurement: Place the alloy block on the walking beam and send it into the preheating zone. Increase the temperature from room temperature to 650℃ at a rate of 3℃ / min, taking approximately 208 minutes. After reaching 650℃, maintain the temperature for uniform uniform temperature measurement. The holding time coefficient k is set to 1.0, and the holding time t_h = 1.0 × 120 = 120 minutes. The total residence time in the preheating zone is the sum of the heating time and the uniform uniform temperature measurement time, approximately 328 minutes. Drill a blind hole with a diameter of 2mm and a depth of 60mm in the center of the alloy block and insert a type K armored thermocouple to monitor the core temperature. The surface temperature is monitored by a spot-welded thermocouple. At the end of preheating, the core temperature should not be lower than 580℃, and the temperature difference between the surface and the core should not exceed 80℃.

[0073] S2, Transition heating: The preheated alloy block is fed into the transition heating zone by the walking beam and heated from 650°C to 1000°C at a heating rate of 30°C / min, with a heating time of about 12min.

[0074] S3, Sub-solution limit: The alloy block is placed in the solution zone and held at 1000℃ for 60 minutes. This solution temperature is about 23℃ lower than the final dissolution temperature of the δ phase of this batch of alloys, and is within the sub-solution limit range.

[0075] S4, graded cooling: exactly the same as in Example 1.

[0076] The stepping cycle of the walking beam is set to 80 seconds per step based on the thickness of the alloy block, and the dwell time accuracy of each temperature zone is controlled within ±10%.

[0077] Comparative Example 4 This comparative example provides a nickel-based superalloy bulk material and its continuous solution annealing method.

[0078] GH4169 alloy forgings produced in the same batch using vacuum induction melting + vacuum arc remelting (VIM+VAR) have the following nominal chemical composition (mass percentage): Ni 52.8%, Cr 19.0%, Nb 5.3%, Mo 3.0%, Ti 0.9%, Al 0.5%, Fe balance. Differential scanning calorimetry (DSC) determined that the δ-phase dissolution termination temperature of this batch of alloy is approximately 1023℃. The forging dimensions are 120mm × 200mm × 500mm, with a thickness H = 120mm.

[0079] A walking beam continuous annealing furnace is employed, with preheating, transition heating, solution treatment, slow cooling, and rapid cooling zones sequentially arranged along the conveying direction. The temperature and atmosphere of each zone are independently controlled. The entire furnace is purged with ammonia decomposition gas (75% H2 + 25% N2), with a dew point of -62℃. The furnace is maintained at a slightly positive pressure of 100 Pa, with an oxygen content ≤15 ppm. The furnace is equipped with hydrogen concentration monitoring and explosion-proof pressure relief devices.

[0080] S1, Preheating and Uniform Temperature Measurement: Place the alloy block on the walking beam and send it into the preheating zone. Increase the temperature from room temperature to 650℃ at a rate of 3℃ / min, taking approximately 208 minutes. After reaching 650℃, maintain the temperature for uniform uniform temperature measurement. The holding time coefficient k is set to 1.0, and the holding time t_h = 1.0 × 120 = 120 minutes. The total residence time in the preheating zone is the sum of the heating time and the uniform uniform temperature measurement time, approximately 328 minutes. Drill a blind hole with a diameter of 2mm and a depth of 60mm in the center of the alloy block and insert a type K armored thermocouple to monitor the core temperature. The surface temperature is monitored by a spot-welded thermocouple. At the end of preheating, the core temperature should not be lower than 580℃, and the temperature difference between the surface and the core should not exceed 80℃.

[0081] S2, transitional heating: exactly the same as in Example 1.

[0082] S3, Sub-limit Solution Treatment: The alloy block is placed in the solution treatment zone and held at 1040℃ for 60 minutes. This solution treatment temperature is approximately 17℃ higher than the final dissolution temperature of the δ phase of this batch of alloys.

[0083] S4, graded cooling: exactly the same as in Example 1.

[0084] The stepping cycle of the walking beam is set to 80 seconds per step based on the thickness of the alloy block, and the dwell time accuracy of each temperature zone is controlled within ±10%.

[0085] Comparative Example 5 This comparative example provides a nickel-based superalloy bulk material and its continuous solution annealing method.

[0086] GH4169 alloy forgings produced in the same batch using vacuum induction melting + vacuum arc remelting (VIM+VAR) have the following nominal chemical composition (mass percentage): Ni 52.8%, Cr 19.0%, Nb 5.3%, Mo 3.0%, Ti 0.9%, Al 0.5%, Fe balance. Differential scanning calorimetry (DSC) determined that the δ-phase dissolution termination temperature of this batch of alloy is approximately 1023℃. The forging dimensions are 120mm × 200mm × 500mm, with a thickness H = 120mm.

[0087] A walking beam continuous annealing furnace is employed, with preheating, transition heating, solution treatment, slow cooling, and rapid cooling zones sequentially arranged along the conveying direction. The temperature and atmosphere of each zone are independently controlled. The entire furnace is purged with ammonia decomposition gas (75% H2 + 25% N2), with a dew point of -62℃. The furnace is maintained at a slightly positive pressure of 100 Pa, with an oxygen content ≤15 ppm. The furnace is equipped with hydrogen concentration monitoring and explosion-proof pressure relief devices.

[0088] S1, Preheating and Uniform Temperature Measurement: Place the alloy block on the walking beam and send it into the preheating zone. Increase the temperature from room temperature to 650℃ at a rate of 3℃ / min, taking approximately 208 minutes. After reaching 650℃, maintain the temperature for uniform uniform temperature measurement. The holding time coefficient k is set to 1.0, and the holding time t_h = 1.0 × 120 = 120 minutes. The total residence time in the preheating zone is the sum of the heating time and the uniform uniform temperature measurement time, approximately 328 minutes. Drill a blind hole with a diameter of 2mm and a depth of 60mm in the center of the alloy block and insert a type K armored thermocouple to monitor the core temperature. The surface temperature is monitored by a spot-welded thermocouple. At the end of preheating, the core temperature should not be lower than 580℃, and the temperature difference between the surface and the core should not exceed 80℃.

[0089] S2, transitional heating: exactly the same as in Example 1.

[0090] S3, Sub-solution limit: The alloy block is placed in the solution zone and held at 1000℃ for 60 minutes. This solution temperature is about 23℃ lower than the final dissolution temperature of the δ phase of this batch of alloys, and is within the sub-solution limit range.

[0091] S4, Staged Cooling: After solution treatment and heat preservation, the alloy block is placed in the cooling zone. A forced-circulation nitrogen injection device is used to directly cool the alloy from 1000°C to 140°C at an average surface cooling rate of 60°C / min, with a cooling time of approximately 14 minutes. The cooling rate is measured by thermocouples on the alloy block surface. The block is then removed from the furnace and air-cooled to room temperature.

[0092] The stepping cycle of the walking beam is set to 80 seconds per step based on the thickness of the alloy block, and the dwell time accuracy of each temperature zone is controlled within ±10%.

[0093] Comparative Example 6 This comparative example provides a nickel-based superalloy bulk material and its continuous solution annealing method.

[0094] GH4169 alloy forgings produced in the same batch using vacuum induction melting + vacuum arc remelting (VIM+VAR) have the following nominal chemical composition (mass percentage): Ni 52.8%, Cr 19.0%, Nb 5.3%, Mo 3.0%, Ti 0.9%, Al 0.5%, Fe balance. Differential scanning calorimetry (DSC) determined that the δ-phase dissolution termination temperature of this batch of alloy is approximately 1023℃. The forging dimensions are 120mm × 200mm × 500mm, with a thickness H = 120mm.

[0095] A walking beam continuous annealing furnace is employed, with preheating, transition heating, solution treatment, slow cooling, and rapid cooling zones sequentially arranged along the conveying direction. The temperature and atmosphere of each zone are independently controlled. The entire furnace is purged with ammonia decomposition gas (75% H2 + 25% N2), with a dew point of -62℃. The furnace is maintained at a slightly positive pressure of 100 Pa, with an oxygen content ≤15 ppm. The furnace is equipped with hydrogen concentration monitoring and explosion-proof pressure relief devices.

[0096] S1, Preheating and Uniform Temperature Measurement: Place the alloy block on the walking beam and send it into the preheating zone. Increase the temperature from room temperature to 650℃ at a rate of 3℃ / min, taking approximately 208 minutes. After reaching 650℃, maintain the temperature for uniform uniform temperature measurement. The holding time coefficient k is set to 1.0, and the holding time t_h = 1.0 × 120 = 120 minutes. The total residence time in the preheating zone is the sum of the heating time and the uniform uniform temperature measurement time, approximately 328 minutes. Drill a blind hole with a diameter of 2mm and a depth of 60mm in the center of the alloy block and insert a type K armored thermocouple to monitor the core temperature. The surface temperature is monitored by a spot-welded thermocouple. At the end of preheating, the core temperature should not be lower than 580℃, and the temperature difference between the surface and the core should not exceed 80℃.

[0097] S2, transitional heating: exactly the same as in Example 1.

[0098] S3, Sub-solution limit: The alloy block is placed in the solution zone and held at 1000℃ for 60 minutes. This solution temperature is about 23℃ lower than the final dissolution temperature of the δ phase of this batch of alloys, and is within the sub-solution limit range.

[0099] S4, Staged Cooling: After solution treatment and holding, the alloy block is placed in the cooling zone. It is continuously cooled from 1000℃ to 140℃ at a cooling rate of 12℃ / min, for approximately 72 minutes. The cooling rate is measured by thermocouples on the surface of the alloy block. It is then removed from the furnace and air-cooled to room temperature.

[0100] The stepping cycle of the walking beam is set to 80 seconds per step based on the thickness of the alloy block, and the dwell time accuracy of each temperature zone is controlled within ±10%.

[0101] To verify the effectiveness of this invention, the test samples were prepared as follows: GH4169 alloy forgings from the same batch as those used in the examples and comparative examples were taken and machined into standard test blocks of three sizes: 80mm×200mm×500mm, 120mm×200mm×500mm, and 150mm×200mm×500mm. These blocks were subjected to continuous solution annealing according to the complete processing steps of Examples 1, 2, 3, and Comparative Examples 1, 2, 3, 4, 5, and 6. Subsequently, all samples underwent a standard two-stage aging treatment to obtain test samples completely corresponding to each process scheme, which were sequentially labeled CRA001 to CRA009. Following this, the following performance tests were performed on all grouped samples.

[0102] 1. Macroscopic Defect and Crack Detection Fluorescent penetrant testing was performed according to GB / T 18851.1—2024 "Nondestructive Testing - Penetrant Testing - Part 1: General Rules", observation conditions were specified according to GB / T 5097—2020 "Nondestructive Testing - Penetrant Testing and Magnetic Particle Testing Observation Conditions", and ultrasonic testing was performed according to GB / T 6402—2024 "Ultrasonic Testing Methods for Steel Forgings". Twenty alloy blocks, after annealing and aging treatment, were taken from each group. Surface cracks were first checked using penetrant testing, followed by internal defects using ultrasonic testing. The crack detection rate was calculated as the proportion of samples with penetrant or ultrasonically detectable cracks to the total number of samples in the group. A lower crack detection rate indicates better thermal stress control during annealing, and the alloy block is less likely to crack due to excessive temperature difference between the core and surface during heating and cooling. Samples with detected through-cracks were not further subjected to mechanical property testing.

[0103] 2. Grain size test The method was performed according to the intercept method in GB / T 6394—2017 "Method for Determination of Average Grain Size of Metals". Samples were taken from three locations: the surface layer (5 mm from the surface), 1 / 4 thickness, and the core (H / 2). After polishing and chemical etching, the grain size was measured under a metallographic microscope. The average grain size difference ΔASTM between the core and the surface layer was calculated. The smaller the difference, the more uniform the grain size at different locations of the cross-section. The maximum grain size difference within the sample was also recorded; this value reflects the presence of local mixed-grain structures. Grain size and its uniformity directly determine the comprehensive mechanical properties and fatigue life of the material.

[0104] 3. δ phase observation The morphology, distribution characteristics, and grain boundary coverage of the δ phase were observed using scanning electron microscopy in backscattered electron mode. After electrolytic etching, the grain boundary regions were characterized at magnifications of 3000x to 5000x, and the chemical composition of the precipitated phases was confirmed by energy dispersive spectroscopy (EDS). The focus was on determining whether the δ phase was distributed as isolated granules at the grain boundaries, or as semi-continuous chains or networks. The morphology of the δ phase directly affects the plasticity of the material; isolated granular distribution helps suppress grain growth without compromising toughness, while continuous network distribution leads to grain boundary embrittlement. The average grain boundary coverage of the δ phase was calculated as the percentage of the sum of the projected lengths of the δ phase particles along the grain boundaries in the SEM images to the total length of the grain boundaries.

[0105] 4. Residual stress test According to GB / T 7704—2017 "Non-destructive Testing - X-ray Stress Measurement Method", the near-surface residual tensile stress of the alloy block was determined by X-ray diffraction. The test was performed on the original surface of the sample after annealing and aging treatment, without any additional processing to avoid introducing additional stress. A lower residual tensile stress value indicates more complete release of thermal stress in the slow cooling section of the graded cooling process, and a lower risk of failure of the alloy block due to residual stress superimposed on external loads during subsequent processing or service.

[0106] 5. Mechanical property testing Tensile specimens were prepared along the main deformation direction of forging, with three specimens per group. Room temperature tensile testing was performed according to GB / T 228.1-2021 "Metallic materials, tensile testing—Part 1: Test at room temperature" to determine yield strength, tensile strength, and elongation after fracture. High-temperature tensile testing at 650℃ was performed according to GB / T 228.2-2015 "Metallic materials, tensile testing—Part 2: Test at high temperature" to determine high-temperature yield strength and elongation. The results are expressed as the arithmetic mean of the three specimens. Room temperature strength and elongation reflect the material's load-bearing and plastic deformation capacity at room temperature, while high-temperature performance directly corresponds to the load-bearing capacity of the alloy block in service environments such as the hot end of an engine.

[0107] 6. Temperature Measurement A blind hole with a diameter of 2mm and a depth of H / 2 is drilled in the center of the alloy block. A type K-type armored thermocouple is inserted, and the hole is sealed with high-temperature resistant ceramic adhesive for real-time monitoring of the core temperature. The surface temperature is simultaneously measured by thermocouples spot-welded to the surface of the alloy block. All heating and cooling rates described in the process steps are based on the temperature changes measured by the thermocouples on the alloy block surface. The core temperature is used to verify whether the core-surface temperature difference at each stage meets the process requirements. This temperature measurement method does not reference any independent testing standards and is a conventional furnace temperature tracking and workpiece temperature measurement method in industrial heat treatment processes.

[0108] Table 1. Process parameters for each embodiment and comparative example Table 2 Results of tissue and physical property tests Table 3 Mechanical property test results (after standard aging, n=3) (a) The role of preheating and temperature equalization and accelerated transition in the plastic zone in preventing thermal stress cracking Comparative Example 1 omitted the preheating stage, directly heating the 120mm thick alloy block to 1000℃ at a rate of 12℃ / min. The excessive temperature difference between the surface and core resulted in significant thermal stress, leading to through-cracks in 65% of the samples and preventing the acquisition of a viable microstructure. Example 2, by heating at a low rate to 650℃ and maintaining a sufficiently uniform temperature (k=1.0), controlled the temperature difference between the surface and core within 80℃, reducing the risk of thermal stress cracking. No cracks were detected in any of the 20 samples. Comparative Example 2 used a relatively low holding time coefficient (k=0.4), resulting in insufficient heat penetration into the core. The actual core temperature during solution treatment was insufficient, leading to incomplete dissolution of the strengthening phase. After aging, the room temperature yield strength was 820MPa, and the core microstructure was uneven. In Comparative Example 3, the transition heating rate was increased to 30℃ / min. Despite preheating, the excessively rapid heating re-induced a significant temperature difference between the core and surface. Internal microcracks were detected in 20% of the samples, and the tensile strength decreased by approximately 200 MPa compared to Example 2, while the elongation decreased from 19.0% to 11.0%. These comparisons demonstrate that sufficient low-speed heating and thickness-coupled holding time during the preheating homogenization stage, combined with a heating rate of 8-15℃ / min in the transition section, are crucial conditions for preventing cracks in large-section bulk materials after annealing.

[0109] (II) The effect of sub-limit solid solution on retaining δ phase pinning and inhibiting grain coarsening In Comparative Example 4, the solution temperature was raised to 1040℃, which exceeded the dissolution end temperature of the δ phase in this batch of alloys (approximately 1023℃). The grain boundary granular δ phase was basically dissolved, the pinning effect was weakened, the austenite grains grew significantly, and the average grain size in the core increased from 22μm in Example 2 to 120μm. Mixed crystals appeared (local ASTM 3.5~8.0). The room temperature elongation decreased from 19.0% to 9.5%, and the high-temperature elongation at 650℃ was 5.5%.

[0110] In Example 2, the solution temperature was selected within the sub-solution limit range (1000℃). SEM confirmed that after solution treatment, discretely distributed granular δ-phase was retained, grain growth was suppressed, and the grain size stabilized at approximately ASTM grade 8, resulting in a good balance between strength and plasticity. Metallographic observation of the alloy block obtained in Example 2 yielded the following results: Figure 1 As shown. By Figure 1 As can be seen, the alloy microstructure consists of uniform equiaxed austenitic grains with clear and complete grain boundaries and a uniform grain size distribution. No abnormally large grains or mixed grain phenomena were observed. According to GB / T 6394—2017, the grain size, measured using the intercept method, is approximately ASTM grade 8, with an average grain size of approximately 22 μm. This indicates that controlling the solution temperature within the sub-solution limit range of the δ phase, preferably 985~1010℃ (the specific temperature depends on the batch Tδ), plays a crucial role in microstructure control.

[0111] (III) The dual effect of staged cooling on suppressing grain boundary embrittlement and eliminating quenching stress Comparative Example 5, by eliminating the slow cooling section and directly implementing rapid cooling, resulted in significant thermal stress in the 120mm thick block material due to the difference in shrinkage between the core and surface. The residual tensile stress reached 240MPa, and 30% of the samples exhibited quenching microcracks, leading to brittle fracture during tensile testing. Example 2, using a slow cooling section at 15℃ / min to preferentially release thermal stress, resulted in a residual tensile stress of 52MPa. Comparative Example 6, by eliminating the rapid cooling section and implementing slow cooling throughout, prolonged residence time in the δ-phase precipitation sensitive zone at 850~950℃ led to the δ-phase precipitating along grain boundaries in a semi-continuous chain-like to locally network-like manner. The room temperature elongation decreased from 19.0% to 8.0%, and the tendency for grain boundary embrittlement increased.

[0112] The alloy block obtained in Example 2 was further observed by scanning electron microscopy backscattered electron microscopy, and the results are as follows: Figure 2 As shown. By Figure 2As can be seen, bright white granular precipitates exist at the grain boundaries, which were confirmed by EDS analysis to be the δ phase, i.e., Ni3Nb. This δ phase is distributed as isolated granules along the grain boundaries in a discontinuous manner, without forming continuous chain-like, semi-continuous film-like, or network-like precipitate morphologies. This microstructure indicates that in the staged cooling process of this invention, the slow cooling section passes through the precipitation-sensitive zone at a cooling rate of 10~20℃ / min, reducing harmful continuous precipitation at grain boundaries; the rapid cooling section cools the alloy block to below 200℃ at an average surface cooling rate of not less than 50℃ / min, fixing the supersaturated solid solution and simultaneously addressing the requirements of preventing quenching cracks and suppressing grain boundary embrittlement.

[0113] In summary, by Figure 1 and Figure 2 It is known that the process of this invention enables GH4169 alloy blocks to obtain a uniform and fine equiaxed grain structure, and causes the δ phase to precipitate in discrete granular form at the grain boundaries, effectively avoiding abnormal grain growth, mixed grains, and continuous precipitation of the δ phase at grain boundaries. The preheating-temperature homogenization-transition heating-subsolution limit-stage cooling synergistic process of this invention solves the problems of thermal stress cracking, microstructure uniformity, and precipitate control in continuous solution annealing of large-section bulk Nb-containing γ″ precipitation-strengthened nickel-based superalloys, obtaining crack-free, uniformly grained, and well-matched strength and plasticity annealed alloy blocks, which are feasible for industrial application.

[0114] The above embodiments are merely illustrative of the principles and effects of the present invention and are not intended to limit the invention. Any person skilled in the art can modify or alter the above embodiments without departing from the spirit and scope of the present invention. Therefore, all equivalent modifications or alterations made by those skilled in the art without departing from the spirit and technical concept disclosed in the present invention should still be covered by the claims of the present invention.

Claims

1. A continuous solution annealing method for nickel-based superalloy bulk materials, characterized in that, The nickel-based superalloy is a Nb-containing γ″ precipitation-strengthened nickel-based superalloy, and the thickness H of the alloy block is 80-150 mm; the method uses a walking beam continuous annealing furnace, which sequentially includes a preheating zone, a transition heating zone, a solution treatment zone, a slow cooling zone, and a rapid cooling zone along the conveying direction, and includes the following steps: S1, Preheating and Uniform Temperature Measurement: The alloy block is placed in the preheating zone and heated from room temperature to 600-650℃ at a heating rate of ≤5℃ / min. Then, uniform temperature measurement and holding are performed within this temperature range. The uniform temperature measurement and holding time is t_h = k×H, where k is the holding time coefficient, which is 0.8-1.2min / mm, and H is the thickness of the alloy block in mm. At the end of preheating, the core temperature of the alloy block is not lower than 580℃, and the temperature difference between the surface and the core is not greater than 80℃. S2, Transition heating: The preheated and homogenized alloy block is sent into the transition heating zone and heated to the solution temperature at a heating rate of 8-15℃ / min. S3, Subsolution Limit: The alloy block is placed into the solution zone and held at the solution temperature for 30 to 90 minutes; the solution temperature is 5 to 30°C lower than the δ phase dissolution temperature of the alloy. S4, graded cooling: After the solution treatment and heat preservation are completed, the alloy block is sent to the cooling zone. First, it is cooled to 600-650°C in the slow cooling zone at a cooling rate of 10-20°C / min, and then cooled to below 200°C in the fast cooling zone at an average surface cooling rate of ≥50°C / min. After being taken out of the furnace, it is air-cooled to room temperature. The above steps are carried out in a protective atmosphere, which is ammonia decomposition gas or nitrogen with a purity ≥99.995%, the furnace pressure is 50-200 Pa, the dew point is ≤-60℃, and the oxygen content is ≤20 ppm.

2. The method according to claim 1, characterized in that, The nickel-based superalloy is GH4169 or Inconel718.

3. The method according to claim 1, characterized in that, The thickness H of the alloy block is 120 mm.

4. The method according to claim 1, characterized in that, The heating rate in step S1 is 3℃ / min, and the temperature for uniform temperature holding is 650℃.

5. The method according to claim 1, characterized in that, The heat preservation time coefficient k mentioned in step S1 is 1.0 to 1.1 min / mm.

6. The method according to claim 1, characterized in that, The solution temperature mentioned in steps S2 and S3 is 985–1010 °C.

7. The method according to claim 1, characterized in that, The heating rate in step S2 is 10℃ / min.

8. The method according to claim 6, characterized in that, The heat preservation time in step S3 is: When H≤100mm, keep warm for 30-60 minutes; When 100 < H ≤ 130 mm, keep warm for 45–75 min; When 130 < H ≤ 150 mm, keep warm for 60–90 minutes.

9. The method according to claim 1, characterized in that, In step S4, the cooling rate of the slow cooling zone is 15℃ / min, cooling to 620℃; the average surface cooling rate of the fast cooling zone is 60℃ / min, cooling to 140℃.

10. The method according to claim 1, characterized in that, The walking beam continuous annealing furnace has a walking cycle of 60-120 seconds per step, which is adjusted according to the thickness H of the alloy block, and the dwell time accuracy of each temperature zone is ±10%; the protective atmosphere is ammonia decomposition gas, and the furnace is equipped with hydrogen concentration monitoring and explosion-proof pressure relief devices.

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