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WO2024149404A1 - 一种盘轴一体涡轮盘及其制备方法 - Google Patents

一种盘轴一体涡轮盘及其制备方法 Download PDF

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Publication number
WO2024149404A1
WO2024149404A1 PCT/CN2024/073096 CN2024073096W WO2024149404A1 WO 2024149404 A1 WO2024149404 A1 WO 2024149404A1 CN 2024073096 W CN2024073096 W CN 2024073096W WO 2024149404 A1 WO2024149404 A1 WO 2024149404A1
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WO
WIPO (PCT)
Prior art keywords
disk
furnace
integrated turbine
shaft integrated
turbine disk
Prior art date
Application number
PCT/CN2024/073096
Other languages
English (en)
French (fr)
Inventor
曲敬龙
孙少斌
侯为学
杨成斌
荣义
杨姗洁
唐超
王民庆
杜金辉
张北江
付锐
谢兴飞
吕少敏
张筱萌
罗志强
王广磊
Original Assignee
北京钢研高纳科技股份有限公司
四川钢研高纳锻造有限责任公司
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Publication of WO2024149404A1 publication Critical patent/WO2024149404A1/zh

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Classifications

    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C19/00—Alloys based on nickel or cobalt
    • C22C19/03—Alloys based on nickel or cobalt based on nickel
    • C22C19/05—Alloys based on nickel or cobalt based on nickel with chromium
    • C22C19/051—Alloys based on nickel or cobalt based on nickel with chromium and Mo or W
    • C22C19/056—Alloys based on nickel or cobalt based on nickel with chromium and Mo or W with the maximum Cr content being at least 10% but less than 20%
    • B—PERFORMING OPERATIONS; TRANSPORTING
    • B21—MECHANICAL METAL-WORKING WITHOUT ESSENTIALLY REMOVING MATERIAL; PUNCHING METAL
    • B21J—FORGING; HAMMERING; PRESSING METAL; RIVETING; FORGE FURNACES
    • B21J5/00—Methods for forging, hammering, or pressing; Special equipment or accessories therefor
    • B21J5/02—Die forging; Trimming by making use of special dies ; Punching during forging
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C1/00—Making non-ferrous alloys
    • C22C1/02—Making non-ferrous alloys by melting
    • C22C1/023—Alloys based on nickel
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22F—CHANGING THE PHYSICAL STRUCTURE OF NON-FERROUS METALS AND NON-FERROUS ALLOYS
    • C22F1/00—Changing the physical structure of non-ferrous metals or alloys by heat treatment or by hot or cold working
    • C22F1/10—Changing the physical structure of non-ferrous metals or alloys by heat treatment or by hot or cold working of nickel or cobalt or alloys based thereon
    • F—MECHANICAL ENGINEERING; LIGHTING; HEATING; WEAPONS; BLASTING
    • F01—MACHINES OR ENGINES IN GENERAL; ENGINE PLANTS IN GENERAL; STEAM ENGINES
    • F01D—NON-POSITIVE DISPLACEMENT MACHINES OR ENGINES, e.g. STEAM TURBINES
    • F01D5/00—Blades; Blade-carrying members; Heating, heat-insulating, cooling or antivibration means on the blades or the members
    • F01D5/02—Blade-carrying members, e.g. rotors
    • Y—GENERAL TAGGING OF NEW TECHNOLOGICAL DEVELOPMENTS; GENERAL TAGGING OF CROSS-SECTIONAL TECHNOLOGIES SPANNING OVER SEVERAL SECTIONS OF THE IPC; TECHNICAL SUBJECTS COVERED BY FORMER USPC CROSS-REFERENCE ART COLLECTIONS [XRACs] AND DIGESTS
    • Y02—TECHNOLOGIES OR APPLICATIONS FOR MITIGATION OR ADAPTATION AGAINST CLIMATE CHANGE
    • Y02P—CLIMATE CHANGE MITIGATION TECHNOLOGIES IN THE PRODUCTION OR PROCESSING OF GOODS
    • Y02P10/00—Technologies related to metal processing
    • Y02P10/25—Process efficiency

Definitions

  • the present application relates to the technical field of turbine disks, and in particular to a disk-shaft integrated turbine disk and a preparation method thereof.
  • Turbine disks are key rotating parts used in aerospace. Compared with ground-based gas turbine turbine disk forgings, aircraft engine turbine disks are smaller in size and rotate faster, requiring materials to have high durability and fatigue performance, and the uniformity of grain size is closely related to this. Therefore, aircraft engine turbine disks have higher requirements for homogeneous and fine grains.
  • Advanced aircraft engine turbine disk forgings are often in the shape of a disk-shaft integrated shape. Controlling the coordinated deformation of the disk body and the shaft and obtaining homogeneous and fine grains are crucial to improving the service life of aircraft engine turbine disk forgings.
  • Existing disk-shaft integrated turbine disk forgings often have mixed crystals, resulting in the high-temperature durability and low-cycle fatigue performance of the disk-shaft integrated turbine disk not being able to meet the requirements of high-temperature use.
  • the present application aims to provide a disk-shaft integrated turbine disk and a preparation method thereof, so as to solve the problem that the existing disk-shaft integrated turbine disk has mixed crystals and is difficult to simultaneously meet the comprehensive requirements of high temperature durability and low cycle fatigue performance.
  • the present application provides a disk-shaft integrated turbine disk, which includes a disk portion and a shaft portion passing through the disk portion, and the disk portion and the shaft portion are integrally formed;
  • the components of the disk-shaft integrated turbine disk, in percentage by mass, include: C: 0.02% ⁇ 0.04%, Cr: 18.5% ⁇ 20.0%, Co: 13.0% ⁇ 14.0%, Mo: 4.0% ⁇ 4.90%, Al: 1.3% ⁇ 1.6%, Ti: 2.80% ⁇ 3.25%, Ti/Al: 2.25 ⁇ 2.38, (Al+Ti): 4.35% ⁇ 4.58%, O: ⁇ 20PPm, N: ⁇ 20PPm, S ⁇ 10PPm, P ⁇ 80PPm, Nickel: balance.
  • the grain size of the integrated turbine disk reaches above level 6.5, and the grain size difference is below level 2.
  • the preparation method comprises:
  • Step 1 Design a closed mold and select the appropriate billet according to the blanking weight
  • Step 2 Apply thermal insulation coating to the sides and ends of the rod and then soft-wrap them;
  • Step 3 forging the rod blank at 1030-1060° C. to obtain a cake blank
  • Step 4 Remove the wrapping from the dough and let it cool
  • Step 5 After the cake is soft-wrapped, it is closed-die forged at 1030-1060° C. to obtain a forging;
  • Step 6 Remove the cover and cool
  • Step 8 Stabilization and aging treatment are performed to obtain a disk-shaft integrated turbine disk.
  • step 1 the rod blank is obtained by vacuum induction melting, electroslag remelting, vacuum arc remelting and homogenization treatment.
  • vacuum induction melting includes:
  • step 1 the vacuum arc remelting includes:
  • helium is introduced for cooling during smelting.
  • the helium flow rate is increased from 0 ml/min to 110 ml/min in the first 0.5 h, and then decreased from 110 ml/min to 20 ml/min within 0.5 h in the hot capping stage.
  • step 3 the upsetting cake is pressed by a single fire, with a deformation amount of 50% to 60% and a pressing rate of 5 to 10 mm/s.
  • step 7 the solution treatment process is: heating to 995-1050° C., keeping the temperature for 3.5-4.5 hours, and oil cooling.
  • the microstructure of the obtained disk-shaft integrated turbine disk mainly includes equiaxed austenite grains and uniformly distributed carbides, as well as dispersed ⁇ 'phase; the carbides mainly include M23C6 and MC; M23C6 is in the shape of short rods and is discontinuously distributed at the grain boundaries; MC is in the shape of blocks and is discontinuously distributed within the grains.
  • the present invention has the following beneficial effects:
  • the disk-shaft integrated turbine disk of the present application improves the solid solution strengthening effect and the grain boundary strength of the alloy by precisely controlling the contents of individual elements such as C, Cr, Co, Al, and Ti in the alloy; and ensures the optimal matching of the content and size of the ⁇ ' phase in the disk-shaft integrated turbine disk by synergistically controlling the contents of O, N, S, and P, reduces the content of inclusions, improves the purity, plasticity, and fatigue properties of the alloy of the disk-shaft integrated turbine disk, ensures the uniformity of the grains and the precipitation and distribution of the grain boundary phase, and ensures the comprehensive performance of the disk-shaft integrated turbine disk.
  • the preparation method of the disk-shaft integrated turbine disk of the present application controls the content of key components, and through the triple smelting and blanking, die forging, and heat treatment processes, a disk-shaft integrated turbine disk that meets the requirements of 700°C service aircraft engines is prepared.
  • the disk-shaft integrated turbine disk has uniform and fine grains (the grain size reaches above level 6.5, for example, level 7 to 8, the grains in different parts are uniform and consistent, and the grain size difference is less than 2 levels), the crack propagation rate is low, and the comprehensive performance for long-term service is excellent.
  • the performance of the disk-shaft integrated turbine disk of the present application is as follows: room temperature tensile performance: tensile strength ⁇ b ⁇ 1300MPa (e.g. 1330-1420MPa); yield strength ⁇ 0.2 ⁇ 1000MPa (e.g. 1010-1050MPa); elongation after fracture ⁇ 5 ⁇ 20.0% (e.g. 21%-24%); cross-sectional shrinkage ⁇ ⁇ 24.0% (e.g. 25%-36%); 535°C performance: tensile strength ⁇ b ⁇ 1200MPa (e.g. 1210-1320MPa); yield strength ⁇ 0.2 ⁇ 875MPa (e.g.
  • the room temperature tensile strength and yield strength of the disk-shaft integrated turbine disk of the present application are less dispersed, and the coefficient of variation Cv value is 2.63% to 4.75%; the elongation at break and the area reduction Cv value are less than 15%, which is within the normal range and has a low dispersion.
  • the performance fluctuation of the disk-shaft integrated turbine disk is small.
  • FIG1 is a schematic diagram of the overall structure of a disk-shaft integrated turbine disk of the present application.
  • FIG2 is a schematic diagram of the grain size detection position of the disk-shaft integrated turbine disk of the present application.
  • FIG3 is a grain map of the disk-shaft integrated turbine disk of Example 1 of the present application.
  • FIG4 is a microstructure diagram of a disk-shaft integrated turbine disk according to Example 1 of the present application.
  • FIG. 5 is a grain map of Comparative Example 4.
  • the inventors found in their in-depth research that the current disc-shaft integrated turbine disc has too wide a control over its composition and a wide manufacturing process window.
  • the C content is generally high, and the content and distribution dispersion of the grain boundary phase are large, which makes it difficult to stably control the disc-shaft integrated turbine disc during thermal deformation, and easily causes mixed crystals, resulting in large performance fluctuations of the disc-shaft integrated turbine disc, making it difficult to simultaneously meet the comprehensive requirements of the rotating parts for high-temperature durability and low-cycle fatigue performance. Therefore, after in-depth research, the inventors have made more precise control over the composition and process of the disc-shaft integrated turbine disc, and provided a disc-shaft integrated turbine disc that can serve for a long time at 700°C.
  • the present application provides a disk-shaft integrated turbine disk, as shown in Figures 1 and 2, the disk-shaft integrated turbine disk includes a disk portion 1 and a shaft portion 2 penetrating the disk portion 1, and the disk portion 1 and the shaft portion 2 are integrally formed;
  • the material of the disk-shaft integrated turbine disk is a nickel-based high-temperature alloy, specifically, the components of the nickel-based high-temperature alloy include, by mass percentage: C: 0.02% ⁇ 0.04%, Cr: 18.5% ⁇ 20.0%, Co: 13.0% ⁇ 14.0%, Mo: 4.0% ⁇ 4.90%, Al: 1.3% ⁇ 1.6%, Ti: 2.80% ⁇ 3.25%, Ti/Al: 2.25 ⁇ 2.38, (Al+Ti): 4.35% ⁇ 4.58%, O: ⁇ 20PPm, N: ⁇ 20PPm, S ⁇ 10PPm, P ⁇ 80PPm, nickel: balance.
  • C in nickel-based alloys mainly affects the mechanical properties of materials by forming MC carbides during solidification and precipitating M 23 C 6 and M 6 C during heat treatment.
  • the granular discontinuous carbides M 23 C 6 precipitated at the grain boundary can prevent grain boundary sliding and crack propagation, thereby improving the durability.
  • Improve the lasting plasticity and toughness; higher C content is easy to increase the precipitation of MC and M 23 C 6 , and increase the uneven distribution of M 23 C 6 at the grain boundary.
  • the inventors have conducted in-depth research and found that when the C content is 0.01%, 0.03%, 0.04%, 0.05%, and 0.1%, the content of M 23 C 6 is 0.19%, 0.58%, 0.78%, 0.97%, and 1.95% respectively.
  • the C content is less than 0.04%, the carbides are distributed discontinuously at the grain boundary; when the C content is greater than 0.04%, the carbides begin to be distributed continuously at the grain boundary; when the C content is 0.06%, there are not only small-sized continuous carbides at the grain boundary, but also large-sized discrete carbides, with an average size of 3 ⁇ m, and the lasting life of the alloy is also reduced. Therefore, in this application, the C content is controlled at 0.02% to 0.04%.
  • Cr Cr is an indispensable alloying element in nickel-based high-temperature alloys. Its main functions are as follows: (1) Solid solution strengthening: Cr in the ⁇ matrix of the high-temperature alloy causes lattice distortion, generates elastic stress field strengthening, and improves the strength of the ⁇ solid solution; (2) Precipitation strengthening: Cr dissolved in the ⁇ solid solution can also form a series of carbides with C, mainly M23C6 type carbides, which are mainly distributed at the grain boundaries. The granular discontinuous carbides evenly distributed at the grain boundaries can effectively prevent grain boundary sliding and migration, and improve the endurance strength; (3) Anti-oxidation: One of the most important functions of Cr in the ⁇ matrix is to form a Cr2O3 type oxide film , which has good anti-oxidation properties. The higher the Cr content, the better the anti-oxidation properties. Considering the cost of the alloy, the Cr content is controlled at 18.5% to 20.0%.
  • Co is one of the main solid solution strengthening elements of nickel-based alloys. Adding Co to the ⁇ matrix can reduce the stacking fault energy of the matrix. When the stacking fault energy is reduced, the probability of stacking faults increases, making the cross-slip of dislocations more difficult. In this way, deformation requires greater external force, which is manifested as an increase in strength. Moreover, the stacking fault energy is reduced, the creep rate is reduced, and the creep resistance is increased. In addition, the Co element can also reduce the solubility of ⁇ ' forming elements Ti and Al in the matrix, thereby increasing the number of ⁇ ' precipitation phases in the alloy and increasing the service temperature of the alloy. However, Co is a scarce resource in my country. Considering the cost factor of the alloy, this application controls the Co element content to 13.0% to 14.0%.
  • Mo enters the nickel-based alloy matrix and mainly plays a role in solid solution strengthening.
  • the solid solution strengthening of the Mo element is also one of the characteristics of the present application.
  • the Mo element content is controlled at 4.0% to 4.9%.
  • Al is the main element that forms the ⁇ ′ phase. About 20% of the Al added to the alloy enters the ⁇ solid solution, which plays a role in solid solution strengthening; 80% of the Al forms Ni 3 Al with Ni, which plays a role in precipitation strengthening. In order to ensure that the alloy has the ⁇ ′ phase necessary to maintain the high temperature strength of 700°C, the Al content of the alloy is limited to 1.3% to 1.6%.
  • Ti About 10% of Ti added to the nickel-based alloy enters the ⁇ solid solution, which plays a certain role in solid solution strengthening, and about 90% enters the ⁇ ′ phase. Under the condition of a certain Al content, as the Ti content increases, the amount of the ⁇ ′ phase increases, and the high-temperature strength of the alloy increases. In order to ensure that the alloy has the ⁇ ′ phase necessary to maintain the high-temperature strength of 700°C, the present application limits the Ti content to 2.80% to 3.25%, and controls Ti/Al: 2.25 to 2.38, (Al+Ti): 4.35% to 4.58%.
  • O Reducing the content of O and N elements can reduce the number of inclusions in the material, which is beneficial to improving the plasticity and toughness of the alloy.
  • N easily forms Ti(C, N) with Ti.
  • the increase of Ti(C, N) increases the possibility of fatigue source formation and also reduces the Ti element content required for the strengthening phase ⁇ ′ phase. Therefore, it is necessary to control O: ⁇ 20PPm, N: ⁇ 20PPm in this application.
  • S Higher S content affects the plasticity and long-term performance of the alloy. At the end of alloy smelting and solidification, the higher the S content, the easier it is for sulfides to precipitate.
  • the S element has a significant effect on nickel-based alloys above 800°C, which is more obvious during the smelting and billeting process of steel ingots.
  • the inventors found in their research that the 100ppm S experimental group failed in smelting, and the 56ppm S group had severe cracking during the billeting process of the steel ingot. When the S content exceeds 10ppm, the endurance life and plasticity at 730°C decrease to varying degrees. Therefore, it is necessary to control S: ⁇ 10PPm in this application.
  • the components of the above-mentioned nickel-based high-temperature alloy may include, by mass percentage: C: 0.026% ⁇ 0.037%, Cr: 18.5% ⁇ 19.7%, Co: 13.0% ⁇ 13.98%, Mo: 4.10% ⁇ 4.70%, Al: 1.3% ⁇ 1.45%, Ti: 2.95% ⁇ 3.25%, Ti/Al: 2.26 ⁇ 2.38, (Al+Ti): 4.36% ⁇ 4.58%, O: ⁇ 10PPm, N: ⁇ 20PPm, S ⁇ 8PPm, P ⁇ 40PPm, Nickel: balance.
  • the microstructure of the above-mentioned disk-shaft integrated turbine disk mainly includes equiaxed austenite grains and uniformly distributed carbides, as well as dispersed ⁇ 'phase; wherein the carbides mainly include M23C6 and MC; M23C6 is in the shape of short rods, discontinuously distributed at the grain boundaries, and has a content of about 0.5% to 0.75% (mass fraction); MC is in the shape of blocks, discontinuously distributed in the crystals, and has a small content of about 0.1 to 0.2% (mass fraction); the ⁇ ' phase is spherical and dispersed inside the grains, the particle size of the ⁇ ' phase is about 60 to 200 nm, and the content is 24% to 26%, wherein the content of the ⁇ ' phase with a particle size of about 60 to 100 nm is about 12 to 16%, and the content of the ⁇ ' phase with a size greater than 100 nm is about 8% to 14%.
  • the carbides mainly include M23C6
  • the grain size of the above-mentioned integrated turbine disk reaches above level 6.5, for example, level 7-8, the grains in different parts are uniform, and the grain size difference is below level 2 (for example, the grain size difference is below level 1.5).
  • the present application also provides a method for preparing the above-mentioned disk-shaft integrated turbine disk, comprising:
  • Step 1 Design a closed mold according to the product shape and select a suitable billet according to the blanking weight
  • Step 2 Apply thermal insulation coating to the sides and end faces of the rod blank and then soft-wrap it (i.e. wrap it with thermal insulation cotton);
  • Step 3 forging the rod blank at 1030-1060° C. to obtain a cake blank
  • Step 4 Remove the wrapping from the dough and let it cool
  • Step 5 After the cake is soft-wrapped, it is closed-die forged at 1030-1060° C. to obtain a forging;
  • Step 6 Remove the cover and cool
  • Step 8 Stabilization and aging treatment are performed to obtain a disk-shaft integrated turbine disk.
  • the rod blank is obtained by vacuum induction melting, electroslag remelting, vacuum arc remelting and homogenization treatment.
  • the method for preparing the rod blank includes:
  • Step 1.1 Weigh the raw materials according to the alloy composition ratio
  • Step 1.2 Vacuum induction melting, including:
  • S1.2.2 Melting and refining Maintain low power baking at 50-100kw, evacuate the air, and use a gradient (e.g. 200kw, 400kw, 600kw) to increase the power to 1000-1500KW. After the materials in the furnace are fully melted, add the remaining nickel plates, remaining carbon powder, and chromium blocks until the materials are completely melted. After melting, electromagnetic stirring is performed on the alloy liquid to promote the rapid reduction of O and N elements; the temperature during the melting period is controlled at 1500°C ⁇ 1560°C, and the temperature during the refining period is controlled at 1500°C ⁇ 1560°C;
  • a gradient e.g. 200kw, 400kw, 600kw
  • Step 1.3 Electroslag remelting, including:
  • the ingot is surface turned, with a single-side turning amount of 5 to 10 mm and a diameter controlled within the range of 400 to 420 mm for later vacuum arc remelting;
  • the electroslag remelting treatment using the above method can effectively reduce the S content in the alloy
  • Step 1.4 Vacuum Arc Remelting:
  • Step 1.5 homogenize to obtain a rod blank.
  • the use of a crucible of mid-term heats (>10 heats) for smelting can effectively reduce the content of gas elements, because the crucible wall has severe outgassing in the first few heats of the crucible.
  • the first part of nickel plates and the first part of carbon powder are loaded into the furnace; then, the second part of nickel plates, the second part of carbon powder, the first part of cobalt blocks and molybdenum bars are loaded into the furnace; and then, the second part of cobalt blocks and the third part of carbon powder are loaded into the furnace.
  • the first part of nickel plates accounts for 1/2 to 2/3 of the total nickel plate mass; the first part of carbon powder accounts for about 1/4 of the total carbon powder mass, the second part of nickel plates accounts for about 1/6 of the total nickel plate mass, the second part of carbon powder accounts for about 1/4 of the total carbon powder mass, the first part of cobalt blocks accounts for about 1/2 of the total cobalt blocks mass, the second part of cobalt blocks accounts for about 1/2 of the total cobalt blocks mass, and the third part of carbon powder accounts for about 1/4 of the total carbon powder mass.
  • the specific steps of adding aluminum blocks and sponge titanium in three batches include: first adding the first part of sponge titanium and the first part of aluminum blocks; then adding the second part of sponge titanium and the second part of aluminum blocks; and adding the third part of aluminum blocks after an interval of 8 to 12 minutes.
  • the mass ratio of the first part of sponge titanium and the second part of sponge titanium is about 1:1; the mass ratio of the first part of aluminum blocks, the second part of aluminum blocks and the third part of aluminum blocks is about 1:1:1.
  • adding Ti to the above S1.2.3 can reduce Ti inclusions, such as the formation of Ti(C,N).
  • Al is added in batches to reduce exothermic temperature rise
  • Al and Ti are added in batches to control the content of Al and Ti elements at the same time.
  • the bottom of the molten pool can be moved upward, changing from the original inverted cone to a flat disk; at the same time, the two-phase area of the mushy zone is reduced, the diffusion distance of the metal elements during the solidification process is reduced, and the segregation tendency of the alloy is reduced.
  • the melting rate of vacuum consumable remelting affects the microstructure of the alloy.
  • the melting rate is controlled within the range of (3.4-4.0) kg/min and the cooling water temperature is controlled between 18°C and 28°C.
  • the purpose of homogenization annealing is to eliminate the low melting point phase in the steel ingot and reduce element segregation.
  • Homogenization annealing is set up with two stages of heat preservation. In the first stage of heat preservation, the temperature is low, and the main function is to eliminate the low melting point phase in the alloy; the second stage of heat preservation can promote the uniform diffusion of segregated elements.
  • the process steps of homogenization annealing include:
  • the temperature is slowly increased from a furnace temperature of ⁇ 400°C to 1150°C to 1165°C, and the heating time is 10 to 15 hours.
  • the temperature is controlled to 1150-1165°C, and the holding time and the size of the ingot are controlled to meet the requirements of complete remelting of the low-melting-point phase.
  • the temperature is controlled to rise to 1180-1195°C, and the holding time and the size of the ingot are controlled to meet the requirements of completing the residual segregation coefficient to reach 0.2.
  • the composition of the rod blank is: C: 0.02-0.04%, Cr: 18.5%- 20.0%, Co: 13.0 ⁇ 14.0%, Mo: 4.0 ⁇ 4.90%, Al: 1.3% ⁇ 1.6%, Ti: 2.80% ⁇ 3.25%, Ti/Al: 2.25 ⁇ 2.38, (Al+Ti): 4.35% ⁇ 4.58%, O: ⁇ 20PPm, N: ⁇ 20PPm, S ⁇ 10PPm, P ⁇ 80PPm, Nickel: balance.
  • the method for preparing the rod blank further includes:
  • Step 1.6 the alloy billet kept at 1150°C to 1200°C is subjected to multiple upsetting, drawing and radial forging in sequence.
  • the deformation amount of each upsetting is 30% to 50%, and the deformation amount of each drawing is 30% to 60%; after each upsetting and drawing, the temperature is reduced by 40 to 50°C until the holding temperature drops to 1050 to 1120°C, and then the temperature is no longer reduced. If the deformation amount of each fire is too large, the alloy is at risk of cracking, and if the deformation amount is too large, the alloy grain structure is at risk of mixed crystals in the large deformation zone. If the deformation amount is too small, the alloy is not deformed sufficiently, and the purpose of grain crushing and recrystallization cannot be achieved.
  • the thickness of the thermal insulation cotton is 10 to 15 mm.
  • the upsetting cake is pressed by one fire, the deformation is 50% to 60%, and the pressing rate is 5 to 10 mm/s; the upsetting cake is pressed at 1030 to 1060°C because at this temperature, the alloy is in a single-phase region as a whole, and the plastic deformation can be more sufficient.
  • the temperature drops, and the ⁇ ' phase precipitates, which can provide pinning effect for grain refinement and slow down grain growth.
  • the die forging adopts a fire pressing
  • the deformation amount is 50% to 60%
  • the pressing rate is 5 to 10 mm/s.
  • the design of closed die forging can not only reduce the weight of the blank, but also design the mold shape according to the deformation characteristics of the disk and shaft, and control the deformation amount of different parts.
  • the solution treatment process is: heating to 995-1050°C, keeping warm for 3.5-4.5h, and oil cooling; considering that the heating rate is too fast, the core of the alloy will not reach the solution temperature and the holding time will be insufficient; therefore, the temperature is controlled to rise slowly, starting from a furnace temperature of ⁇ 400°C, and the heating rate is controlled to be 4-6°C/min. Too high a temperature or too long a holding time will lead to grain growth, which is not conducive to the performance of the alloy.
  • the temperature is controlled to rise to 995-1050°C. Keep warm for 3.5 to 4.5 hours.
  • the stabilization and aging treatment process is: heating to 840-850°C and keeping it for 3-4.5 hours, and air cooling; heating to 755-765°C and keeping it for 15-17 hours, and then air cooling to room temperature.
  • the temperature is slowly raised, starting from a furnace temperature of ⁇ 400°C, and the heating rate is controlled to be 4-6°C/min.
  • the performance of the integrated turbine disk with a disk and a shaft obtained is as follows: room temperature performance: tensile strength ⁇ b ⁇ 1300MPa (e.g. 1330-1420MPa); yield strength ⁇ 0.2 ⁇ 1000MPa (e.g. 1010-1050MPa); elongation after fracture ⁇ 5 ⁇ 20.0% (e.g. 21%-24%); cross-sectional shrinkage ⁇ ⁇ 24.0% (e.g. 25%-36%); 535°C performance: tensile strength ⁇ b ⁇ 1200MPa (e.g. 1210-1320MPa); yield strength ⁇ 0.2 ⁇ 875MPa (e.g.
  • the disk-shaft integrated turbine disk of the present application improves the solid solution strengthening effect and the grain boundary strength of the alloy by precisely controlling the contents of individual elements such as C, Cr, Co, Al, and Ti in the alloy; and ensures the optimal matching of the content and size of the ⁇ ' phase in the disk-shaft integrated turbine disk by synergistically controlling the contents of O, N, S, and P, reduces the content of inclusions, improves the purity, plasticity, and fatigue properties of the alloy of the disk-shaft integrated turbine disk, ensures the uniformity of the grains and the precipitation and distribution of the grain boundary phase, and ensures The comprehensive performance of the integrated turbine disc.
  • the preparation method of the disk-shaft integrated turbine disk of the present application controls the content of key components, and through the triple smelting and blanking, die forging, and heat treatment processes, a disk-shaft integrated turbine disk that meets the requirements of 700°C service aircraft engines is prepared.
  • the disk-shaft integrated turbine disk has uniform and fine grains (the grain size reaches above level 6.5, for example, level 7 to 8, the grains in different parts are uniform and consistent, and the grain size difference is less than 2 levels), the crack propagation rate is low, and the comprehensive performance for long-term service is excellent.
  • Examples 1-5 of the present application provide a disk-shaft integrated turbine disk and a preparation method thereof.
  • the method for preparing the disk-shaft integrated turbine disk of Examples 1-5 comprises:
  • Embodiment 1 is a diagrammatic representation of Embodiment 1:
  • Step 1 Design a closed mold according to the product shape, and select a suitable billet according to the blanking weight; the blanking weight is 90 ⁇ 1kg, the billet specification is ⁇ 180 ⁇ 430mm, and the specification tolerance is ⁇ 1mm;
  • Step 2 Apply thermal insulation coating to the sides and ends of the rod blank and then cover them with thermal insulation cotton;
  • Step 3 Place the rod blank into a heating furnace and heat it to 1030°C for 6 hours; then press it into a cake; the pressing speed is 10mm/s, and the cake height is precisely controlled to 190 ⁇ 1mm;
  • Step 4 Remove the wrapping from the dough and let it cool
  • Step 5 After the cake is soft-wrapped, it is placed in a heating furnace and heated to 1030°C for 6 hours; then the cake is placed in a closed die forging die forging to obtain a disc forging; the deformation of the upsetting cake and die forging is 50% respectively;
  • Step 6 Remove the cover and cool
  • Step 7 Solution treatment: start heating from 400°C in the furnace at a rate of 5°C/min to 1025°C, keep warm for 4.5h, and cool with oil;
  • Step 8 Stabilization and aging treatment to obtain a disk-shaft integrated turbine disk: start heating from a furnace temperature of 400°C at a heating rate of 5°C/min, heat to 850°C, keep warm for 4.5 hours, and air cool; start heating from a furnace temperature of 400°C at a heating rate of 5°C/min, heat to 765°C, keep warm for 17 hours, and then air cool to room temperature.
  • the method for preparing the rod blank in step 1 above includes:
  • Step 1.1 Weigh the raw materials according to the alloy composition ratio
  • Step 1.2 Vacuum Induction Melting:
  • S1.2.1 Loading First, load the first part of nickel plate and the first part of carbon powder into the furnace; then load the second part of nickel plate, the second part of carbon powder, the first part of cobalt block and molybdenum bar into the furnace; then load the second part of cobalt block and the third part of carbon powder into the furnace; after the furnace charge is completely loaded into the furnace, evacuate to a vacuum degree of ⁇ 0.1Pa and prepare to turn on the power;
  • the first part of nickel plates accounts for about 2/3 of the total nickel plate mass;
  • the first part of carbon powder accounts for about 1/4 of the total carbon powder mass,
  • the second part of nickel plates accounts for about 1/6 of the total nickel plate mass,
  • the second part of carbon powder accounts for about 1/4 of the total carbon powder mass,
  • the first part of cobalt blocks accounts for about 1/2 of the total cobalt blocks mass,
  • the second part of cobalt blocks accounts for about 1/2 of the total cobalt blocks mass, and the third part of carbon powder accounts for about 1/4 of the total carbon powder mass;
  • Step 1.3 Electroslag remelting:
  • the ingot is surface turned, with a single-side turning amount of 5 to 10 mm and a diameter controlled within the range of 400 to 420 mm for later vacuum arc remelting;
  • Step 1.4 Vacuum Arc Remelting:
  • Step 1.5 Homogenize the steel ingot by annealing it from 400°C to 1150°C over 11h. Keep warm for 47 hours; continue to heat up to 1180°C, keep warm for 65 hours and then air cool;
  • Step 1.6 Prepare the rod blank by fast forging + radial forging:
  • the alloy rod blank kept at 1200°C is subjected to three upsetting, drawing and radial forging in sequence; wherein, the deformation amount of each upsetting is 30%, and the deformation amount of each drawing is 30%; after each upsetting and drawing, the temperature is reduced by 40°C until the holding temperature drops to 1080°C, and then the blank is opened on the radial forging machine for multiple passes to the finished rod blank.
  • the pre-melted slag composition is CaF 2 65%, Al 2 O 3 25%, CaO 25%, MgO 8%, TiO 2 5%;
  • Step 1.5 homogenizing annealing the steel ingot: heating from 400°C to 1165°C over 15h, keeping it at that temperature for 49h; further heating to 1195°C, keeping it at that temperature for 67h, and then air cooling;
  • Step 1.6 The alloy billet kept at 1200°C is subjected to three upsetting and drawing processes and radial forging processes in sequence; wherein the deformation amount of each upsetting process is 50%, and the deformation amount of each drawing process is 60%; after each upsetting and drawing process, the temperature is lowered by 50°C until the holding temperature is reduced to 1050°C, and then the billet is opened on a radial forging machine in multiple passes to a finished billet;
  • step 3 and step 5 the temperature is 1060°C;
  • Step 7 Solution treatment: start heating from 400°C in the furnace at a rate of 5°C/min to 1025°C, keep warm for 4.5h, and cool with oil;
  • Step 8 Start heating from a furnace temperature of 400°C at a rate of 5°C/min, heat to 850°C, keep warm for 4.5 hours, and air cool; Start heating from a furnace temperature of 400°C at a rate of 5°C/min, heat to 765°C, keep warm for 17 hours, and then air cool to room temperature.
  • Embodiment 3 is a diagrammatic representation of Embodiment 3
  • the pre-melted slag composition is CaF 2 55%, Al 2 O 3 20%, CaO 20%, MgO 5%, TiO 2 3%;
  • Step 1.5 Perform homogenization annealing on the steel ingot: heat from 400°C to 1160°C over 13h, keep warm for 48h; continue heating to 1190°C, keep warm for 66h and then air cool.
  • Step 1.6 The alloy bar billet kept at 1180°C is subjected to three upsetting and drawing processes and radial forging processes in sequence; wherein the deformation amount of each upsetting process is 50%, and the deformation amount of each drawing process is 50%; after each upsetting and drawing process, the temperature is lowered by 40°C until the holding temperature is reduced to 1060°C, and then the billet is opened on a radial forging machine in multiple passes to a finished bar;
  • step 3 and step 5 the temperature is 1050°C; the deformation of the upsetting cake and the die forging is 60% respectively;
  • Step 7 Solution treatment: start heating from 400°C in the furnace at a rate of 5°C/min to 1020°C, keep warm for 4h, and cool with oil;
  • Step 8 Start heating from a furnace temperature of 400°C at a rate of 5°C/min, heat to 845°C, keep warm for 4 hours, and air cool; Start heating from a furnace temperature of 400°C at a rate of 5°C/min, heat to 760°C, keep warm for 16 hours, and then air cool to room temperature.
  • Example 4-5 The difference between Example 4-5 and Example 3 is that the chemical composition of the steel is different, see Table 1, and the other parameters are the same.
  • the present application also provides 5 comparative examples.
  • the chemical compositions of the disc-shaft integrated turbine discs of Examples 1-5 and Comparative Examples 1-5 are shown in Table 1.
  • Comparative Example 4 The components of Comparative Example 4 are the same as those of Example 1.
  • the die forging temperature of Comparative Example 4 is 1090° C.
  • the deformation amounts of the upsetting cake and die forging are 30% and 40%, respectively.
  • Comparative Examples 1-3 and 5 have different components from Example 3, but the preparation methods are the same.
  • the metallographic structures of the embodiments and comparative examples are shown in Table 2. Specifically, as shown in FIG2 , the grain sizes at different positions are checked. It can be seen that the grains of the embodiments of the present application are uniform. The grain size reaches 6.5 or above, for example, 7 to 8, and the grains at different positions are uniform and consistent, and the grain size difference is less than 2 levels, for example, the grain size difference is 1 level.
  • Table 3 shows the room temperature mechanical properties of the embodiments and comparative examples of the present application
  • Table 4 shows the room temperature mechanical properties of the present application.
  • Table 5 is the 535°C mechanical properties of the embodiments and comparative examples
  • Table 5 is the 730°C/550MPa endurance performance of the embodiments and comparative examples of the present application
  • Table 6 is the 500°C low cycle fatigue performance of the embodiments and comparative examples of the present application.
  • FIG. 3 is a grain map of Example 1 of the present application
  • FIG. 4 is a microstructure map of Example 1 of the present application
  • FIG. 5 is a grain map of Comparative Example 4.
  • the average grain size of the integrated turbine disk of the present application is 7-8, which is finer than 5 and the difference is less than 2 (for example, the grain size difference is 1), and the coefficient of variation Cv ranges from 2.33% to 4.33%, which is less than 15%.
  • the data is normal and the degree of dispersion is low.
  • the room temperature tensile strength and yield strength of 535°C tensile strength have a small dispersion, and the coefficient of variation Cv value is 2.63% to 4.75%; the elongation at break and the surface shrinkage Cv value are less than 15%, which are within the normal range and have a low dispersion.
  • the content of elements in the alloy of the integrated disk-shaft turbine disk is precisely controlled; and by synergistically controlling the values of Ti/Al and (Al+Ti), combined with the precise control of the preparation method, the uniform grains of the microstructure of the integrated disk-shaft turbine disk are ensured, the grain size reaches above level 6.5, and the content and size of the ⁇ ' phase in the alloy are optimally matched; by precisely controlling the content of O, N, S, and P, the content of inclusions in the alloy is reduced, the purity, plasticity, and fatigue properties of the alloy are improved, the uniformity of the grains, the precipitation and distribution of the grain boundary phase are ensured, and the comprehensive performance of the alloy is ensured.
  • the room temperature tensile and 535°C tensile tensile strength and yield strength of the integrated disk-shaft turbine disk of this application have a small dispersion, and the coefficient of variation Cv value is 2.63% to 4.75%; the elongation at break and the surface reduction Cv value are less than 15%, which is within the normal range and has a low dispersion.
  • the performance fluctuation of the integrated disk-shaft turbine disk is small.

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Abstract

本申请公开了一种盘轴一体涡轮盘及其制备方法，属于涡轮盘技术领域，解决了现有技术中盘轴一体涡轮盘存在混晶，难以同时满足对高温持久及低周疲劳性能的综合要求的问题。盘轴一体涡轮盘包括盘部和贯穿盘部的轴部，盘部和轴部一体成型；盘轴一体涡轮盘的组分以质量百分比计，包括：C：0.02％～0.04％，Cr：18.5％～20.0％，Co：13.0％～14.0％，Mo：4.0％～4.90％，Al：1.3％～1.6％，Ti：2.80％～3.25％，Ti/Al：2.25～2.38，(Al+Ti)：4.35％～4.58％，O：≤20PPm，N：≤20PPm，S≤10PPm，P≤80PPm,镍：余量。本申请的盘轴一体涡轮盘的综合性能优异。

Description

一种盘轴一体涡轮盘及其制备方法 技术领域
本申请涉及涡轮盘技术领域，特别涉及一种盘轴一体涡轮盘及其制备方法。
背景技术
涡轮盘是航空航天用的关键转动部件，相比地面燃气轮机涡轮盘锻件，航空发动机涡轮盘尺寸更小且转速更快，要求材料有高持久和疲劳性能，而晶粒尺寸的均匀性与此密切相关，因此航空发动机涡轮盘对材料均质细晶的要求更高。而先进航空发动机涡轮盘锻件常为盘轴一体形状，控制盘体与轴部协同变形、获得均质细晶，对提高航空发动机涡轮盘锻件服役寿命至关重要。现有的盘轴一体涡轮盘锻件经常存在混晶，造成盘轴一体涡轮盘的高温持久及低周疲劳性能不能满足高温使用需求。
发明内容
鉴于上述情况，本申请旨在提供一种盘轴一体涡轮盘及其制备方法，用于解决现有的盘轴一体涡轮盘存在混晶，难以同时满足对高温持久及低周疲劳性能的综合要求的问题。
本申请的目的主要是通过以下技术方案实现的：
一方面，本申请提供了一种盘轴一体涡轮盘，盘轴一体涡轮盘包括盘部和贯穿盘部的轴部，盘部和轴部一体成型；盘轴一体涡轮盘的组分以质量百分比计，包括：C：0.02％～0.04％，Cr：18.5％～20.0％，Co：13.0％～14.0％，Mo：4.0％～4.90％，Al：1.3％～1.6％，Ti：2.80％～3.25％，Ti/Al：2.25～2.38，(Al+Ti)：4.35％～4.58％，O：≤20PPm，N：≤20PPm，S≤10PPm，P≤80PPm,镍：余量。
进一步的，盘轴一体涡轮盘的晶粒度达到6.5级以上，晶粒度级差2级以下。
进一步的，制备方法包括：
步骤1：设计闭式模具，根据下料重量选择合适的棒坯；
步骤2：将棒坯侧面及端面涂抹保温涂料后进行软包套；
步骤3：将棒坯在1030～1060℃镦饼，得到饼坯；
步骤4：将饼坯去除包套，冷却；
步骤5：将饼坯软包套后在1030～1060℃闭式模锻，得到锻件；
步骤6：去除包套，冷却；
步骤7：固溶处理；
步骤8：稳定化及时效处理得到盘轴一体涡轮盘。
进一步的，步骤1中，棒坯采用真空感应熔炼、电渣重熔、真空电弧重熔和均匀化处理得到。
进一步的，步骤1中，真空感应熔炼包括：
S1.2.1布料：分批次将镍板、碳粉、钴块、钼条装入炉中；炉料完全入炉后，抽真空，准备通电；
S1.2.2熔化精炼：维持低功率烘烤，抽空排气，采用梯度上升功率至1000～1500KW，炉内料全熔后加入剩余镍板与剩余碳粉、铬块直至炉料完全熔清，熔清后对合金液进行电磁搅拌；
S1.2.3出炉浇铸：降功率保温，分批加入铝块、海绵钛，炉料熔化后，向炉内充Ar气后加入镍合金并施加电磁搅拌，调节功率至浇注温度出钢，炉冷一段时间后，破空处理得到真空感应铸锭。
进一步的，步骤1中，真空电弧重熔包括：
S1.4.1将电极调入结晶器内，对中处理后，在炉内完成电极焊接工作；
S1.4.2真空度小于1Pa，漏气率小于0.3Pa/min时送电冶炼；
S1.4.3熔炼时充入氦气冷却；
S1.4.4铸锭在炉内冷却后，对真空自耗电弧进行破空处理，得到钢锭。
进一步的，S1.4.3中，熔炼时充入氦气冷却，按照流量控制，前0.5h氦气流量从0ml/min升到110ml/min，热封顶阶段在0.5h内从110ml/min降低到20ml/min。
进一步的，步骤3中，镦饼采用一火压制，变形量为50％～60％，压下速率为5～10mm/s。
进一步的，步骤7中，固溶处理工艺为：升温至995～1050℃保温3.5～4.5h，油冷。
进一步的，步骤8中，得到的盘轴一体涡轮盘的微观组织主要包括等轴奥氏体晶粒和均匀分布的碳化物，以及弥散分布的γ’相；碳化物主要包括M23C6和MC；M23C6呈短棒状，断续分布在晶界；MC呈块状，断续分布在晶内。
与现有技术相比，本申请有益效果如下：
a)本申请的盘轴一体涡轮盘通过精确控制合金中的C、Cr、Co、Al、Ti等单个元素的含量，提高了合金固溶强化的作用和合金的晶界强度；并通过协同控制Ti/Al、(Al+Ti)的值，保证了盘轴一体涡轮盘中γ’相的含量与尺寸最佳匹配，通过精确控制O、N、S、P的含量，减少了夹杂物的含量，提高了盘轴一体涡轮盘的合金的纯净度和塑性及疲劳性能，保证了晶粒的均匀性及晶界相的析出及分布，保证了盘轴一体涡轮盘的综合性能。
b)本申请的盘轴一体涡轮盘的制备方法通控制关键成分含量，并通过三联冶炼及开坯、模锻、热处理过程制备出了满足700℃服役航空发动机需求的盘轴一体涡轮盘，其晶粒均匀细小(晶粒度达到6.5级以上，例如7～8级，不同部位的晶粒均匀一致，晶粒度级差2级以下)、裂纹扩展速率低，长期服役综合性能优越。
c)本申请的盘轴一体涡轮盘性能如下：室温拉伸性能：抗拉强度σb≥1300MPa(例如1330～1420MPa)；屈服强度σ0.2≥1000MPa(例如1010～1050MPa)；断后伸长率δ5≥20.0％(例如21％～24％)；断面收缩率ψ≥24.0％(例如25％～36％)；535℃性能：抗拉强度σb≥1200MPa(例如1210～1320MPa)；屈服强度σ0.2≥875MPa(例如885～950MPa)；断后伸长率δ5≥19％(例如19％～22％)；断面收缩率ψ≥23％(例如24％～29％)；730℃/550MPa持久性能：持久时间τ≥35h(例如41～47h)；815℃/295MPa持久性能：持久时间τ≥52h(例如53～70h)；断后伸长率δ5≥12％(例如13％～18％)；低周疲劳性能：500℃/应变控制0～0.7％/0.33Hz，＞3×104周(例如33884～47323)。
d)本申请的盘轴一体涡轮盘的室温拉伸和535℃拉伸抗拉强度和屈服强度离散度较小，变异系数Cv值为2.63％～4.75％；断裂伸长率和面缩率Cv值小于15％，在正常范围内，离散度低。该盘轴一体涡轮盘的性能波动小。
本申请的其他特征和优点将在随后的说明书中阐述，并且，部分的从说明书中变得显而易见，或者通过实施本申请而了解。本申请的目的和其他优点可通过在所写的说明书、权利要求书、以及附图中所特别指出的结构来实现和获得。
附图说明
附图仅用于示出具体实施例的目的，而并不认为是对本申请的限制，在整个附图中，相同的参考符号表示相同的部件。
图1为本申请的盘轴一体涡轮盘的整体结构示意图；
图2为本申请的盘轴一体涡轮盘的晶粒度检测位置示意图；
图3为本申请实施例1的盘轴一体涡轮盘的晶粒图；
图4为本申请实施例1的盘轴一体涡轮盘的微观组织图；
图5是对比例4的晶粒图。
附图说明
1-盘部，2-轴部。
具体实施方式
下面结合附图来具体描述本申请的优选实施例，其中，附图构成本申请一部分，并与本申请的实施例一起用于阐释本申请的原理。
发明人在深入研究中发现，目前的盘轴一体涡轮盘的成分控制过于宽泛，制备工艺窗口也较宽，并且现有工艺中，C含量一般较高，晶界相的含量及分布分散度较大，在热变形过程中难以稳定控制，容易造成混晶，导致盘轴一体涡轮盘的性能波动大，难以同时满足转动件对高温持久及低周疲劳性能的综合要求。因此，发明人经过深入研究，对盘轴一体涡轮盘的成分和工艺进行了更加精确的控制，提供了一种能在700℃长期服役的盘轴一体涡轮盘。
本申请提供了一种盘轴一体涡轮盘，如图1和图2所示，盘轴一体涡轮盘包括盘部1和贯穿盘部1的轴部2，盘部1和轴部2一体成型；盘轴一体涡轮盘的材料采用镍基高温合金，具体的，镍基高温合金的组分以质量百分比计包括：C：0.02％～0.04％，Cr：18.5％～20.0％，Co：13.0％～14.0％，Mo：4.0％～4.90％，Al：1.3％～1.6％，Ti：2.80％～3.25％，Ti/Al：2.25～2.38，(Al+Ti)：4.35％～4.58％，O：≤20PPm，N：≤20PPm，S≤10PPm，P≤80PPm,镍：余量。
具体的，O+N+S≤50PPm。
以下对本申请中盘轴一体涡轮盘所含组分的作用及用量选择作具体说明：
C：镍基合金中C主要通过在凝固过程形成MC碳化物以及在热处理过程中析出的M23C6、M6C等影响材料的力学性能。在晶界析出的颗粒状不连续碳化物M23C6，可以阻止晶界滑移和裂纹扩展，提高持久寿命， 改善持久塑性和韧性；较高的C含量容易增加MC、M23C6的析出，并增加晶界M23C6分布的不均匀性，发明人深入研究发现：C含量分别为0.01％、0.03％、0.04％、0.05％、0.1％时，M23C6的含量分别为0.19％、0.58％、0.78％、0.97％、1.95％。C含量小于0.04％时，碳化物在晶界呈断续状分布；当C含量为大于0.04％时，碳化物开始在晶界呈连续状分布；当C含量为0.06％时，晶界不仅有小尺寸的连续状碳化物，还有大尺寸离散状的碳化物存在，其平均尺寸为3μm，合金的持久寿命也随之降低。因此，本申请中将C含量控制在0.02％～0.04％。
Cr：Cr是镍基高温合金中不可缺少的合金化元素，其主要作用有如下几点：(1)固溶强化：高温合金γ基体中的Cr引起晶格畸变，产生弹性应力场强化，而使γ固溶体强度提高；(2)析出强化：溶解于γ固溶体的Cr还能与C形成一系列碳化物，主要以M 23C 6型碳化物为主，该碳化物主要分布在晶界处，均匀的分布于晶界的颗粒状不连续碳化物，可有效地阻止晶界滑移和迁移，提高持久强度；(3)抗氧化：Cr在γ基体中一种十分重要的作用就是形成Cr2O3型氧化膜，具有良好的抗氧化性能，且Cr含量越高，抗氧化性能越好。考虑合金的成本问题，Cr含量控制在18.5％～20.0％。
Co：Co是镍基合金主要固溶强化元素之一，Co加入γ基体可降低基体堆垛层错能，层错能降低，层错出现的几率增大，使位错的交滑移更加困难，这样变形就需要更大的外力，表现为强度的提高；而且层错能降低，蠕变速率降低，蠕变抗力增加。此外，Co元素还可以降低γ′形成元素Ti、Al在基体中的溶解度，从而提高合金中的γ′析出相的数量，提高合金的服役温度。但Co在我国属于稀缺资源，考虑合金的成本因素，本申请将Co元素含量控制在13.0％～14.0％。
Mo：Mo进入镍基合金基体，主要起固溶强化作用。特别是在降低 成本而降低Co含量减弱固溶强化作用时，辅以Mo元素的固溶强化也是本申请的特点之一。本申请将Mo元素含量控制在4.0％～4.9％。
Al：Al是形成γ′相的主要元素，加入合金中的Al约有20％进入γ固溶体，起固溶强化作用；80％的Al与Ni形成Ni3Al起沉淀强化作用。为了确保合金中有维持700℃高温强度所必须的γ′相，本申请将合金Al含量限制在1.3％～1.6％。
Ti：加入镍基合金的Ti，约有10％进入γ固溶体，起一定固溶强化作用，约90％进入γ′相。在一定Al含量的条件下，随着Ti含量增加，γ′相数量增加，合金高温强度增加，为了确保合金中有维持700℃高温强度所必须的γ′相，本申请将Ti含量限制在2.80％～3.25％，同时控制Ti/Al：2.25～2.38，(Al+Ti)：4.35％～4.58％。
O：降低O、N元素含量可以减少材料中的夹杂数量，有利于提高合金的塑性和韧性。在合金熔化期间，N容易与Ti形成Ti(C，N)，Ti(C，N)的增多增加了疲劳源形成的可能性，同时也会减少强化相γ′相所需的Ti元素含量。因此，本申请中需要控制O：≤20PPm，N：≤20PPm。
S：较高的S含量对合金的塑性和长时性能有影响。在合金冶炼凝固末期，S含量越高，硫化物越容易析出。S元素对镍基合金在800℃以上产生明显影响，在钢锭冶炼和开坯过程中表现较为明显，发明人在研究中发现：100ppm含量S实验组冶炼失败，而S含量为56ppm组钢锭开坯过程中开裂严重，S含量超过10ppm后，730℃下持久寿命和塑性均有不同程度的下降。因此，本申请中需要控制S：≤10PPm。
P：随着P含量的增高，镍基合金的持久寿命急剧降低，发明人在研究中发现：当P含量超过80ppm时，730℃持久寿命小于25h，当超过100ppm时，持久寿命只有2h左右；且当P含量超过80ppm之后，镍基合金730℃持久有显著的缺口敏感性。因此，本申请中需要控制P： ≤80PPm。
为了进一步改善上述盘轴一体涡轮盘的综合性能，上述镍基高温合金的组分以质量百分比计可以包括：C：0.026％～0.037％，Cr：18.5％～19.7％，Co：13.0％～13.98％，Mo：4.10％～4.70％，Al：1.3％～1.45％，Ti：2.95％～3.25％，Ti/Al：2.26～2.38，(Al+Ti)：4.36％～4.58％，O：≤10PPm，N：≤20PPm，S≤8PPm，P≤40PPm,镍：余量。
具体的，O+N+S≤40PPm。
具体的，上述盘轴一体涡轮盘的微观组织主要包括等轴奥氏体晶粒和均匀分布的碳化物，以及弥散分布的γ’相；其中碳化物主要包括M23C6和MC；M23C6呈短棒状，断续分布在晶界，含量约为0.5％～0.75％(质量分数)；MC呈块状，断续分布在晶内，含量较少，约为0.1～0.2％(质量分数)左右；γ’相呈圆球状，弥散分布在晶粒内部，γ’相的粒径约60～200nm，含量为24％～26％，其中粒径约60～100nm的γ’相的含量约为12～16％，尺寸大于100nm的γ’相的含量约为8％～14％。
具体的，上述盘轴一体涡轮盘的晶粒度达到6.5级以上，例如7～8级，不同部位的晶粒均匀一致，晶粒度级差2级以下(例如晶粒度级差1.5级以下)。
本申请还提供了上述盘轴一体涡轮盘的制备方法，包括：
步骤1：根据产品形状设计闭式模具，根据下料重量选择合适的棒坯；
步骤2：将棒坯侧面及端面涂抹保温涂料后进行软包套(即包覆保温棉)；
步骤3：将棒坯在1030～1060℃镦饼，得到饼坯；
步骤4：将饼坯去除包套，冷却；
步骤5：将饼坯软包套后在1030～1060℃闭式模锻，得到锻件；
步骤6：去除包套，冷却；
步骤7：固溶处理；
步骤8：稳定化及时效处理得到盘轴一体涡轮盘。
具体的，上述步骤1中，为了保证棒坯的组分合格，晶粒均匀，需要精确控制棒坯的制备方法，具体的，棒坯采用真空感应熔炼、电渣重熔、真空电弧重熔和均匀化处理得到。
具体的，上述步骤1中，棒坯的制备方法包括：
步骤1.1：按照合金成分配比称取原料；
步骤1.2：真空感应熔炼，包括：
S1.2.1布料：分批次将镍板、碳粉、钴块、钼条装入炉中；炉料完全入炉后，抽真空到真空度≤0.1Pa，准备通电；
S1.2.2熔化精炼：维持50～100kw低功率烘烤，抽空排气，采用梯度(例如200kw、400kw、600kw)上升功率至1000～1500KW，炉内料全熔后加入剩余镍板与剩余碳粉、铬块直至炉料完全熔清，熔清后对合金液进行电磁搅拌，促使O元素和N元素快速降低；熔化期温度控制在1500℃～1560℃，精炼期温度控制在1500℃～1560℃；
S1.2.3出炉浇铸：降功率保温，分批加入铝块、海绵钛，炉料熔化后，向炉内充Ar气至20000Pa～30000Pa后加入镍合金并施加电磁搅拌促进熔化和均匀成分，调节功率至浇注温度出钢，出钢温度控制在1450℃～1510℃，炉冷一段时间后，破空处理得到真空感应铸锭；
S1.2.4表面处理：铸锭冷却后，切除冒口，并进行表面车光处理，以便后期进行电渣重熔处理；
步骤1.3：电渣重熔，包括：
S1.3.1选取CaF2 45％～65％，Al2O3 15％～25％，CaO 15％～25％，MgO 2％～8％，TiO2 0～5％作为预熔渣；
S1.3.2电极经过焊接打磨入炉后，通入氩气，氩气流量不低于 30L/min，然后进行化渣、起弧、重熔、补缩工作，待铸锭炉冷2h后，脱模得到ESR铸锭，重熔阶段控制熔速在3.7Kg/min～4.2Kg/min之间，重熔过程水温控制在28℃～35℃；
S1.3.3铸锭进行表面车削处理，单边车削量5～10mm，直径控制在400～420mm范围内，以便后期进行真空电弧重熔处理；
采用上述方法进行电渣重熔处理，可以有效降低合金中的S元素含量；
步骤1.4：真空电弧重熔：
S1.4.1对真空电弧炉结晶器进行清理，然后将电极调入结晶器内，对中处理后，在炉内完成电极焊接工作；
S1.4.2真空度小于1Pa，漏气率小于0.3Pa/min时送电冶炼；
S1.4.3熔炼时充入氦气冷却，按照流量控制，前0.5h氦气流量从0ml/min升到110ml/min，热封顶阶段在0.5h内从110ml/min降低到20ml/min；
S1.4.4铸锭在炉内冷却一定时长后，对真空自耗电弧进行破空处理，得到钢锭；
步骤1.5：均匀化处理得到棒坯。
具体的，上述S1.2.1中，大尺寸炉料尽量铺设在底部，以防止熔化过程中搭桥现象的发生。
具体的，上述S1.2.1中，采用中期炉次(＞10炉次)的坩埚进行冶炼，可有效降低气体元素的含量，因为坩埚前期几个炉次，坩埚壁放气严重。
具体的，上述S1.2.1中，先将第一部分镍板和第一部分碳粉装入炉中；然后将第二部分镍板、第二部分碳粉、第一部分钴块、钼条装入炉中；再将第二部分钴块和第三部分碳粉装入炉中。
具体的，上述S1.2.1中，第一部分镍板占总镍板质量的1/2～2/3；第一部分碳粉占总碳粉质量的1/4左右，第二部分镍板占总镍板质量的1/6左右，第二部分碳粉占总碳粉质量的1/4左右，第一部分钴块占总钴块质量的1/2左右，第二部分钴块占总钴块质量的1/2左右，第三部分碳粉占总碳粉质量的1/4左右。
具体的，上述S1.2.1中，考虑到精炼时真空度提高将促进碳氧反应，随着CO气泡的上浮排出，有利于H和N的析出、非金属夹杂的上浮、氮化物的分解、微量有害元素的挥发；但过高的真空度会加剧坩埚与金属间的反应、增加合金元素的挥发损失。因此，控制抽真空到真空度≤0.1Pa，准备通电。
具体的，上述S1.2.3中，分3批加铝块、海绵钛的具体步骤包括：先加入第一部分海绵钛和第一部分铝块；然后加入第二部分海绵钛和第二部分铝块；间隔8～12min后加入第三部分铝块。其中第一部分海绵钛和第二部分海绵钛的质量比约为1：1；第一部分铝块、第二部分铝块和第三部分铝块的质量比约为1：1：1。
具体的，在上述S1.2.3中加入Ti，能够减少Ti的夹杂，例如Ti(C，N)形成。
具体的，在上述S1.2.3中分批次加Al，减少放热升温，分批次投料Al、Ti可同时控制Al、Ti元素含量。
具体的，在上述S1.4.3中，通过前0.5h氦气流量从0ml/min升到110ml/min，热封顶阶段在0.5h内从110ml/min降低到20ml/min的控制，可促使熔池的底部上移，由原来的倒锥形变成扁平的盘状；同时缩小糊状区两相区，减少金属元素在凝固过程中的扩散距离，降低合金的偏析倾向。
具体的，在上述S1.4.3中，真空自耗重熔的熔速影响合金的显微疏 松数量，为减少该缺陷，熔速控制范围：(3.4～4.0)kg/min，冷却水温控制在18℃～28℃之间。
具体的，上述步骤1.5中，均匀化退火的目的是消除钢锭中的低熔点相和减轻元素偏析。均匀化退火设置了2段保温，在第1阶段保温中，温度较低，主要作用是消除合金中的低熔点相；第2阶段的保温可促进偏析元素扩散均匀。具体的，均匀化退火的工艺步骤包括：
S1.5.1、升温至1150～1165℃，保温47～49h；
S1.5.2、继续升温至1180～1195℃，保温65～67h后空冷。
具体的，S1.5.1中，缓慢升温，由炉温≤400℃缓慢升温至1150℃～1165℃，升温时间10～15h。
具体的，S1.5.1中，考虑到温度过高与保温时间过长对低熔点相的回熔并不起到积极作用且高温下的长时间保温会导致铸锭表面氧化层厚度加深、晶粒长大，不利于后续锻造加工；而温度过低与保温时间过短则无法保证低熔点相完全回熔，残留的低熔点相容易成为锻造的裂纹源。因此，控制升温至1150～1165℃，控制保温时间与钢锭的尺寸符合低熔点相完全回熔即可。
具体的，S1.5.2中，温度越高且保温时间越长，偏析元素的回熔越充分，但是当温度与时间达到匹配平衡点后，偏析元素的回熔将保持稳定。发明人深入研究认为残余偏析系数达到0.2时即可完成均匀化热处理。因此温度过高与保温时间过长，对偏析元素回熔积极作用较小，反而会导致晶粒粗大，能源浪费，生产效率下降；而温度较低和保温时间过短，则无法保证偏析元素大部分回熔，枝晶偏析将在锻造过程中降低热加工塑性。因此控制升温至1180～1195℃，控制保温时间与钢锭的尺寸符合完成残余偏析系数达到0.2时即可。
具体的，上述步骤1中，棒坯的成分为：C：0.02～0.04％，Cr：18.5％～ 20.0％，Co：13.0～14.0％，Mo：4.0～4.90％，Al：1.3％～1.6％，Ti：2.80％～3.25％，Ti/Al：2.25～2.38，(Al+Ti)：4.35％～4.58％，O：≤20PPm，N：≤20PPm，S≤10PPm，P≤80PPm,镍：余量。
具体的，上述棒坯的制备方法还包括：
步骤1.6、将保温在1150℃～1200℃的合金棒坯依次进行多次镦粗、拔长、径向锻造。
具体的，上述步骤1.6中，每次镦粗的变形量为30％～50％，每次拔长的变形量为30％～60％；每火次镦粗、拔长后降温40～50℃，直至保温温度降低至1050～1120℃之后，不再降温。若每火次变形量太大，合金有开裂的风险，且变形量过大，合金晶粒组织在大变形区有混晶的风险，若变形量过小，则合金变形不充分，无法完成晶粒破碎与再结晶的目的。
具体的，上述步骤2中，保温棉厚度为10～15mm。
具体的，上述步骤3中，镦饼采用一火压制，变形量为50％～60％，压下速率为5～10mm/s；在1030～1060℃镦饼是因为该温度下，合金整体处于单相区，塑性变形可更充分，锻造结束后，温度下降，γ’相析出，可为晶粒细化提供钉扎作用，减缓晶粒长大。
具体的，上述步骤5中，模锻采用一火压制，变形量为50％～60％，压下速率为5～10mm/s，闭式模锻的设计，既可减小下料重量，又可以针对盘部和轴部的变形特性设计模具形状，控制不同部位的变形量。
具体的，上述步骤7中，固溶处理工艺为：升温至995～1050℃保温3.5～4.5h，油冷；考虑到升温速度过快会导致合金心部未达到固溶温度而导致保温时间不足；因此，控制缓慢升温，由炉温≤400℃开始升温，控制升温速度为4～6℃/min。温度过高或保温时间过长会导致晶粒长大不利于合金性能。而温度较低或保温时间过短，又无法溶解或部分溶解γ’相，为后续时效过程获得合适的γ’相做好准备；因此，控制升温至995～1050℃ 保温3.5～4.5h。
具体的，上述步骤8中，稳定化及时效处理工艺：升温至840～850℃保温3～4.5h，空冷；升温至755～765℃保温15～17h后空冷到室温。
具体的，上述步骤8中，考虑到升温速度过快会导致合金心部未达到稳定化和时效处理温度而导致保温时间不足，导致析出的γ’相不充分，数量和尺寸未达到最佳匹配值；温度过高或保温时间过长会导致γ’相尺寸较大，偏离最佳尺寸数量比，影响合金强度。因此，缓慢升温，由炉温≤400℃开始升温，控制升温速度为4～6℃/min。
具体的，上述步骤8中，得到的盘轴一体涡轮盘的性能如下：室温性能：抗拉强度σb≥1300MPa(例如1330～1420MPa)；屈服强度σ0.2≥1000MPa(例如1010～1050MPa)；断后伸长率δ5≥20.0％(例如21％～24％)；断面收缩率ψ≥24.0％(例如25％～36％)；535℃性能：抗拉强度σb≥1200MPa(例如1210～1320MPa)；屈服强度σ0.2≥875MPa(例如885～950MPa)；断后伸长率δ5≥19％(例如19％～22％)；断面收缩率ψ≥23％(例如24％～29％)；730℃/550MPa持久性能：持久时间τ≥35h(例如41～47h)；断后伸长率δ5≥24％(例如25％～38％)；815℃/295MPa持久性能：持久时间τ≥52h(例如53～70h)；断后伸长率δ5≥12％(例如13％～18％)；低周疲劳性能：500℃/应变控制0～0.7％/0.33Hz，＞3×104周(例如33884～47323)。
与现有技术相比，本申请的盘轴一体涡轮盘通过精确控制合金中的C、Cr、Co、Al、Ti等单个元素的含量，提高了合金固溶强化的作用和合金的晶界强度；并通过协同控制Ti/Al、(Al+Ti)的值，保证了盘轴一体涡轮盘中γ’相的含量与尺寸最佳匹配，通过精确控制O、N、S、P的含量，减少了夹杂物的含量，提高了盘轴一体涡轮盘的合金的纯净度和塑性及疲劳性能，保证了晶粒的均匀性及晶界相的析出及分布，保证了 盘轴一体涡轮盘的综合性能。
本申请的盘轴一体涡轮盘的制备方法通控制关键成分含量，并通过三联冶炼及开坯、模锻、热处理过程制备出了满足700℃服役航空发动机需求的盘轴一体涡轮盘，其晶粒均匀细小(晶粒度达到6.5级以上，例如7～8级，不同部位的晶粒均匀一致，晶粒度级差2级以下)、裂纹扩展速率低，长期服役综合性能优越。
实施例1-5
下面以具体的实施例与对比例来展示本申请的盘轴一体涡轮盘的成分和工艺参数精确控制的优势。本申请的实施例1-5提供了一种盘轴一体涡轮盘及其制备方法。
实施例1-5的盘轴一体涡轮盘的组分见下表1。
实施例1-5的盘轴一体涡轮盘的制备方法包括：
实施例1:
步骤1：根据产品形状设计闭式模具，根据下料重量选择合适的棒坯；其下料重量90±1kg，棒坯规格φ180×430mm，规格公差±1mm；
步骤2：将棒坯侧面及端面涂抹保温涂料后包覆保温棉；
步骤3：将棒坯放入加热炉后升温至1030℃，保温6h；然后进行镦饼，得到饼坯；压制速度10mm/s，精确控制饼坯高度190±1mm；
步骤4：将饼坯去除包套，冷却；
步骤5：将饼坯软包套后放入加热炉后升温至1030℃，保温6h；然后将饼坯放入闭式模锻模具进行模锻，得到盘锻件；镦饼和模锻的变形量分别为50％；
步骤6：去除包套，冷却；
步骤7：固溶处理：由炉温400℃开始升温，升温速度为5℃/min，升温至1025℃，保温4.5h，油冷；
步骤8：稳定化及时效处理得到盘轴一体涡轮盘：由炉温400℃开始升温，升温速度为5℃/min，升温至850℃保温4.5h，空冷；由炉温400℃开始升温，升温速度为5℃/min，升温至765℃保温17h后空冷到室温。
具体的，上述步骤1中的棒坯的制备方法包括：
步骤1.1：按照合金成分配比称取原料；
步骤1.2：真空感应熔炼：
S1.2.1布料：先将第一部分镍板和第一部分碳粉装入炉中；然后将第二部分镍板、第二部分碳粉、第一部分钴块、钼条装入炉中；再将第二部分钴块和第三部分碳粉装入炉中；炉料完全入炉后，抽真空到真空度≤0.1Pa，准备通电；
其中，第一部分镍板占总镍板质量的2/3左右；第一部分碳粉占总碳粉质量的1/4左右，第二部分镍板占总镍板质量的1/6左右，第二部分碳粉占总碳粉质量的1/4左右，第一部分钴块占总钴块质量的1/2左右，第二部分钴块占总钴块质量的1/2左右，第三部分碳粉占总碳粉质量的1/4左右；
S1.2.2熔化精炼：维持50～100kw低功率烘烤，抽空排气，采用200kw、400kw、600kw)上升功率至1000～1500KW，炉内料全熔后加入剩余镍板与剩余碳粉、铬块直至炉料完全熔清，熔清后对合金液进行电磁搅拌，促使O元素和N元素快速降低；熔化期温度控制在1500℃～1560℃，精炼期温度控制在1500℃～1560℃；
S1.2.3出炉浇铸：降功率保温，分批加入铝块、海绵钛(先加入第一部分海绵钛和第一部分铝块；然后加入第二部分海绵钛和第二部分铝块；间隔8～12min后加入第三部分铝块。其中第一部分海绵钛和第二部分海绵钛的质量比约为1：1；第一部分铝块、第二部分铝块和第三部分铝块的质量比约为1：1：1)，炉料熔化后，向炉内充Ar气至20000Pa～30000Pa 后加入镍合金并施加电磁搅拌促进熔化和均匀成分，调节功率至浇注温度出钢，出钢温度控制在1450℃～1510℃，炉冷一段时间后，破空处理得到真空感应铸锭；
S1.2.4表面处理：铸锭冷却后，切除冒口，并进行表面车光处理，以便后期进行电渣重熔处理；
步骤1.3：电渣重熔：
S1.3.1选取CaF2 45％～65％，Al2O3 15％～25％，CaO 15％～25％，MgO 2％～8％，TiO2 0～5％作为预熔渣；
S1.3.2电极经过焊接打磨入炉后，通入氩气，氩气流量不低于30L/min，然后进行化渣、起弧、重熔、补缩工作，待铸锭炉冷2h后，脱模得到ESR铸锭，重熔阶段控制熔速在3.7Kg/min～4.2Kg/min之间，重熔过程水温控制在28℃～35℃；
S1.3.3铸锭进行表面车削处理，单边车削量5～10mm，直径控制在400～420mm范围内，以便后期进行真空电弧重熔处理；
步骤1.4：真空电弧重熔：
S1.4.1对真空电弧炉结晶器进行清理，然后将电极调入结晶器内，对中处理后，在炉内完成电极焊接工作；
S1.4.2真空度小于1Pa，漏气率小于0.3Pa/min时送电冶炼；
S1.4.3熔炼时充入氦气冷却，按照流量控制，前0.5h氦气流量从0ml/min升到110ml/min，热封顶阶段在0.5h内从110ml/min降低到20ml/min。熔速控制范围：(3.6～4.0)kg/min，冷却水温控制在18℃～24℃之间；
S1.4.4铸锭在炉内冷却一定时长后，对真空自耗电弧进行破空处理，得到Φ508mm的VAR铸锭；
步骤1.5：将钢锭进行均匀化退火：由400℃经过11h升温至1150℃， 保温47h；继续升温至1180℃，保温65h后空冷；
步骤1.6：采用快锻+径向锻造制备棒坯：将保温在1200℃的合金棒坯依次进行三次镦粗和拔长、径向锻造；其中，每次镦粗的变形量为30％，每次拔长的变形量为30％；每火次镦粗、拔长后降温40℃，直至保温温度降低至1080℃之后在径锻机上多道次开坯至成品棒坯。
实施例2
本实施例的制备方法与实施例1大致相同，不同之处在于：
S1.3.1中，预熔渣组分为CaF2 65％，Al2O3 25％，CaO 25％，MgO 8％，TiO25％；
步骤1.5：将钢锭进行均匀化退火：由400℃经过15h升温至1165℃，保温49h；继续升温至1195℃，保温67h后空冷；
步骤1.6：将保温在1200℃的合金棒坯依次进行三次镦粗和拔长、径向锻造；其中，每次镦粗的变形量为50％，每次拔长的变形量为60％；每火次镦粗、拔长后降温50℃，直至保温温度降低至1050℃之后在径锻机上多道次开坯至成品棒坯；
步骤3和步骤5中，温度为1060℃；
步骤7：固溶处理：由炉温400℃开始升温，升温速度为5℃/min，升温至1025℃，保温4.5h，油冷；
步骤8：由炉温400℃开始升温，升温速度为5℃/min，升温至850℃保温4.5h，空冷；由炉温400℃开始升温，升温速度为5℃/min，升温至765℃保温17h后空冷到室温。
实施例3:
本实施例的制备方法与实施例1大致相同，不同之处在于：
S1.3.1中，预熔渣组分为CaF2 55％，Al2O3 20％，CaO 20％，MgO 5％，TiO2 3％；
步骤1.5：将钢锭进行均匀化退火：由400℃经过13h升温至1160℃，保温48h；继续升温至1190℃，保温66h后空冷。
步骤1.6：将保温在1180℃的合金棒坯依次进行三次镦粗和拔长、径向锻造；其中，每次镦粗的变形量为50％，每次拔长的变形量为50％；每火次镦粗、拔长后降温40℃，直至保温温度降低至1060℃之后在径锻机上多道次开坯至成品棒材；
步骤3和步骤5中，温度为1050℃；镦饼和模锻的变形量分别为60％；
步骤7：固溶处理：由炉温400℃开始升温，升温速度为5℃/min，升温至1020℃，保温4h，油冷；
步骤8：由炉温400℃开始升温，升温速度为5℃/min，升温至845℃保温4h，空冷；由炉温400℃开始升温，升温速度为5℃/min，升温至760℃保温16h后空冷到室温。
实施例4-5与实施例3的区别在于钢的化学成分不同，见表1，其余参数相同。
本申请同时提供了5种对比例，实施例1-5和对比例1-5的盘轴一体涡轮盘的化学成分见表1。
对比例4与实施例1的组分相同，制备方法中，对比例4的模锻温度为1090℃，镦饼和模锻的变形量分别为30％和40％。
对比例1-3和5与实施例3的组分不同，制备方法相同。
表1 化学成分wt％

实施例和对比例的金相组织见表2。具体的，如图2所示，晶粒度检查了不同位置的。可见，本申请的实施例的晶粒均匀。晶粒度达到6.5级以上，例如7～8级，不同部位的晶粒均匀一致，晶粒度级差2级以下，例如晶粒度级差1级。
表2 金相组织
表3是本申请的实施例和对比例的室温力学性能，表4是本申请的 实施例和对比例的535℃力学性能，表5是本申请的实施例和对比例的730℃/550MPa持久性能，表6是本申请的实施例和对比例的500℃低周疲劳性能。
表3 室温拉伸力学性能
表4 535℃拉伸力学性能
表5 持久性能
表6 500℃/应变控制0～0.7％/0.33Hz低周疲劳性能
图3是本申请实施例1的晶粒图；图4是本申请实施例1的微观组织图，图5是对比例4的晶粒图。
本申请的盘轴一体涡轮盘的平均晶粒度为7～8级，细于5级且级差小于2级(例如晶粒度级差1级)，变异系数Cv范围2.33％～4.33％，小于15％，数据正常，离散程度低。室温拉伸和535℃拉伸抗拉强度和屈服强度离散度较小，变异系数Cv值为2.63％～4.75％；断裂伸长率和面缩率Cv值小于15％，在正常范围内，离散度低。
综上分析可见，本申请中通过精确控制盘轴一体涡轮盘的合金中的元素的含量；并通过协同控制Ti/Al、(Al+Ti)的值，结合制备方法的精确控制，保证了盘轴一体涡轮盘的微观组织的晶粒均匀，晶粒度达到6.5级以上，合金中γ’相的含量与尺寸最佳匹配；通过精确控制O、N、S、P的含量，减少了合金中夹杂物的含量，提高了合金的纯净度和塑性及疲劳性能，保证了晶粒的均匀性及晶界相的析出及分布，保证了合金的综合性能。本申请的盘轴一体涡轮盘的室温拉伸和535℃拉伸抗拉强度和屈服强度离散度较小，变异系数Cv值为2.63％～4.75％；断裂伸长率和面缩率Cv值小于15％，在正常范围内，离散度低。该盘轴一体涡轮盘的性能波动小。
以上所述，仅为本申请较佳的具体实施方式，但本申请的保护范围并不局限于此，任何熟悉本技术领域的技术人员在本申请揭露的技术范围内，可轻易想到的变化或替换，都应涵盖在本申请的保护范围之内。

Claims (10)

  1. 一种盘轴一体涡轮盘，其特征在于，所述盘轴一体涡轮盘包括盘部(1)和贯穿所述盘部(1)的轴部(2)，所述盘部(1)和所述轴部(2)一体成型；所述盘轴一体涡轮盘的组分以质量百分比计，包括：C：0.02％～0.04％，Cr：18.5％～20.0％，Co：13.0％～14.0％，Mo：4.0％～4.90％，Al：1.3％～1.6％，Ti：2.80％～3.25％，Ti/Al：2.25～2.38，(Al+Ti)：4.35％～4.58％，O：≤20PPm，N：≤20PPm，S≤10PPm，P≤80PPm,镍：余量。
  2. 根据权利要求1所述的盘轴一体涡轮盘，其特征在于，所述盘轴一体涡轮盘的晶粒度达到6.5级以上，晶粒度级差2级以下。
  3. 一种权利要求1或2所述的盘轴一体涡轮盘的制备方法，其特征在于，所述制备方法包括：
    步骤1：设计闭式模具，根据下料重量选择合适的棒坯；
    步骤2：将棒坯侧面及端面涂抹保温涂料后进行软包套；
    步骤3：将棒坯在1030～1060℃镦饼，得到饼坯；
    步骤4：将饼坯去除包套，冷却；
    步骤5：将饼坯软包套后在1030～1060℃闭式模锻，得到锻件；
    步骤6：去除包套，冷却；
    步骤7：固溶处理；
    步骤8：稳定化及时效处理得到盘轴一体涡轮盘。
  4. 根据权利要求3所述的制备方法，其特征在于，所述步骤1中，棒坯采用真空感应熔炼、电渣重熔、真空电弧重熔和均匀化处理得到。
  5. 根据权利要求4所述的制备方法，其特征在于，所述步骤1中，真空感应熔炼包括：
    S1.2.1布料：分批次将镍板、碳粉、钴块、钼条装入炉中；炉料完全入炉后，抽真空，准备通电；
    S1.2.2熔化精炼：维持低功率烘烤，抽空排气，采用梯度上升功率至1000～1500KW，炉内料全熔后加入剩余镍板与剩余碳粉、铬块直至炉料完全熔清，熔清后对合金液进行电磁搅拌；
    S1.2.3出炉浇铸：降功率保温，分批加入铝块、海绵钛，炉料熔化后，向炉内充Ar气后加入镍合金并施加电磁搅拌，调节功率至浇注温度出钢，炉冷一段时间后，破空处理得到真空感应铸锭。
  6. 根据权利要求4所述的制备方法，其特征在于，所述步骤1中，真空电弧重熔包括：
    S1.4.1将电极调入结晶器内，对中处理后，在炉内完成电极焊接工作；
    S1.4.2真空度小于1Pa，漏气率小于0.3Pa/min时送电冶炼；
    S1.4.3熔炼时充入氦气冷却；
    S1.4.4铸锭在炉内冷却后，对真空自耗电弧进行破空处理，得到钢锭。
  7. 根据权利要求6所述的制备方法，其特征在于，所述S1.4.3中，熔炼时充入氦气冷却，按照流量控制，前0.5h氦气流量从0ml/min升到110ml/min，热封顶阶段在0.5h内从110ml/min降低到20ml/min。
  8. 根据权利要求3所述的制备方法，其特征在于，所述步骤3中，镦饼采用一火压制，变形量为50％～60％，压下速率为5～10mm/s。
  9. 根据权利要求3所述的制备方法，其特征在于，所述步骤7中，固溶处理工艺为：升温至995～1050℃保温3.5～4.5h，油冷。
  10. 根据权利要求3至9任一项所述的制备方法，其特征在于，所述步骤8中，盘轴一体涡轮盘的微观组织主要包括等轴奥氏体晶粒和均匀分布的碳化物，弥散分布的γ’相；碳化物主要包括M23C6和MC；M23C6呈短棒状，断续分布在晶界；MC呈块状，断续分布在晶内。
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