WO2022100169A1 - 抗蠕变、长寿命镍基变形高温合金及其制备方法和应用 - Google Patents

抗蠕变、长寿命镍基变形高温合金及其制备方法和应用 Download PDF

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WO2022100169A1
WO2022100169A1 PCT/CN2021/111913 CN2021111913W WO2022100169A1 WO 2022100169 A1 WO2022100169 A1 WO 2022100169A1 CN 2021111913 W CN2021111913 W CN 2021111913W WO 2022100169 A1 WO2022100169 A1 WO 2022100169A1
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Prior art keywords
alloy
creep
long
nickel
resistant
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PCT/CN2021/111913
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English (en)
French (fr)
Inventor
束国刚
文新理
董建新
章清泉
江河
李振瑞
刘伟
余志勇
彭劼
柳海波
魏然
李国超
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中国联合重型燃气轮机技术有限公司
北京北冶功能材料有限公司
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Application filed by 中国联合重型燃气轮机技术有限公司, 北京北冶功能材料有限公司 filed Critical 中国联合重型燃气轮机技术有限公司
Priority to JP2023528611A priority Critical patent/JP7488423B2/ja
Priority to EP21890698.0A priority patent/EP4245873A4/en
Priority to KR1020237019601A priority patent/KR102658234B1/ko
Publication of WO2022100169A1 publication Critical patent/WO2022100169A1/zh

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    • 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
    • C22C19/055—Alloys based on nickel or cobalt based on nickel with chromium and Mo or W with the maximum Cr content being at least 20% but less than 30%
    • 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
    • F02—COMBUSTION ENGINES; HOT-GAS OR COMBUSTION-PRODUCT ENGINE PLANTS
    • F02C—GAS-TURBINE PLANTS; AIR INTAKES FOR JET-PROPULSION PLANTS; CONTROLLING FUEL SUPPLY IN AIR-BREATHING JET-PROPULSION PLANTS
    • F02C7/00—Features, components parts, details or accessories, not provided for in, or of interest apart form groups F02C1/00 - F02C6/00; Air intakes for jet-propulsion plants
    • F—MECHANICAL ENGINEERING; LIGHTING; HEATING; WEAPONS; BLASTING
    • F05—INDEXING SCHEMES RELATING TO ENGINES OR PUMPS IN VARIOUS SUBCLASSES OF CLASSES F01-F04
    • F05D—INDEXING SCHEME FOR ASPECTS RELATING TO NON-POSITIVE-DISPLACEMENT MACHINES OR ENGINES, GAS-TURBINES OR JET-PROPULSION PLANTS
    • F05D2220/00—Application
    • F05D2220/30—Application in turbines
    • F05D2220/32—Application in turbines in gas turbines
    • F—MECHANICAL ENGINEERING; LIGHTING; HEATING; WEAPONS; BLASTING
    • F05—INDEXING SCHEMES RELATING TO ENGINES OR PUMPS IN VARIOUS SUBCLASSES OF CLASSES F01-F04
    • F05D—INDEXING SCHEME FOR ASPECTS RELATING TO NON-POSITIVE-DISPLACEMENT MACHINES OR ENGINES, GAS-TURBINES OR JET-PROPULSION PLANTS
    • F05D2300/00—Materials; Properties thereof
    • F05D2300/10—Metals, alloys or intermetallic compounds
    • F05D2300/17—Alloys
    • F05D2300/175—Superalloys

Definitions

  • the invention belongs to the field of metal materials, in particular to a creep-resistant, long-life nickel-based deformed superalloy, particularly a preparation method of the creep-resistant, long-life nickel-based deformed superalloy, and further, to the anti-creep and long-life nickel-based deformed superalloy.
  • the initial gas temperature is getting higher and higher.
  • gas turbines as an example, the gas initial temperature of the most advanced G-class and H-class gas turbines has reached 1450-1500°C.
  • the initial gas temperature of the J-class gas turbine will reach 1600-1700°C, and the temperature of the alloy body of the aero-engine and gas turbine precision hot-end components will also reach 800-950°C. Therefore, the high-temperature mechanical properties of the alloy at 800-950°C are increasingly stringent. requirements.
  • Nimonic263, HastelloyX and Haynes230 alloys have good processing properties, these alloys can only serve for a long time below 800 °C, and their high temperature strength and creep resistance are seriously insufficient at temperatures above 800 °C; Inconel718 alloys have good processing properties, but only It can serve below 650 °C, and the organization will lose its stability and its performance will deteriorate at higher temperature.
  • R-41 and Waspaloy have high ⁇ ' content and fast precipitation rate, so it is difficult to hot work (forging, hot rolling), heat treatment or cold work of parts (cold bending, turning, welding, etc.) It is suitable for the manufacture of precision hot-end components that require complex machining processes in aero-engines and gas turbines.
  • the creep plastic elongation ( ⁇ p ) of R-41 at 816°C, 221MPa, and 100h is about 1%, and the value corresponding to Waspaloy is greater than 1%.
  • the two alloys are at 89MPa, 927°C.
  • the durable life ( ⁇ ) of both alloys is below 100h, which shows that the creep resistance and durable life of these two alloys do not meet the design requirements of advanced aero-engines and gas turbines.
  • Haynes282 alloy has good mechanical properties at room temperature and high temperature, and is easy to process and weld, but it also has three shortcomings: 1) low creep resistance, creep plastic elongation at 816°C, 221MPa, 100h About 1%; 2) Insufficient durable life, under 89MPa, 927 °C, the durable life is less than 200h; 3) The allowable service temperature is low, and the overhaul cycle of gas turbines is considered in practical engineering applications at about 8 years , In order to ensure the safety of components made of Haynes282 alloy through the 8-year maintenance cycle, the maximum operating temperature is strictly limited to about 800 °C.
  • the present invention is made based on the inventor's discovery and understanding of the following facts and problems:
  • advanced aero-engines and gas turbines not only require extremely high initial machining accuracy and assembly accuracy of precision hot-end components, but also require high temperatures of 800-950°C.
  • Excessive plastic deformation cannot occur during long-term service, that is, the alloy is required to have excellent creep resistance, so as to avoid failure of parts before the arrival of the maintenance cycle or difficult to disassemble and replace during maintenance.
  • the present invention aims to solve one of the technical problems in the related art at least to a certain extent.
  • the embodiment of the first aspect of the present invention proposes a creep-resistant, long-life nickel-based deformed superalloy, the alloy has excellent creep resistance and durable life, and the The creep plastic elongation is less than 0.5%, and the durable life under the condition of 89MPa and 927°C is more than 200h, so that the alloy can meet the design and use requirements of advanced aero-engines and gas turbines, and is suitable for the manufacture of advanced aero-engines and gas turbines. Precision hot-end components for long-term service.
  • a creep-resistant, long-life nickel-based deformed superalloy comprises: C: 0.04-0.08%, Cr: 18.50-21.50%, Co: 9.00-11.00%, Mo: 8.00-9.00%, Al: 2.00-3.00%, Ti: 1.10-1.49%, Nb: 0.81-2.00%, B: 0.003-0.009%, Sc: 0.001-0.10%, the balance is nickel and inevitable impurities,
  • the mass percentage contents of the elements Al, Ti and Mo in the alloy satisfy the relational formula: 11.59% ⁇ Al+Ti+Mo ⁇ 13.0%.
  • Sc element is added in the alloy of the embodiment of the present invention, and the addition of Sc introduces a new strengthening mechanism for the alloy of the embodiment of the present invention, forming a Ni 3 (Al, Al, Ti, Nb) composite strengthening phase, which is more high temperature resistant and more stable than the single traditional strengthening phase Ni 3 (Al, Ti), Ni 3 (Al, Ti, Nb), and significantly improves the creep resistance and lasting life of the alloy.
  • the addition of Sc in the alloy of the embodiment of the present invention has an effect on the pretreatment of molten steel, Sc refines the as-cast structure, significantly improves the dendrite segregation of the ingot, and improves the hot workability of the alloy on the one hand.
  • Sc added in the embodiment of the present invention has a post-processing effect on grain boundaries, that is, not only purifying molten steel, but also in liquid - During and after the solid transformation, Sc can still purify and strengthen the grain boundaries, making it difficult for S, P, five harmful elements and other unavoidable low-melting impurity elements to segregate at the grain boundaries, preventing the formation of grain boundaries at high temperatures Creep voids; 5.
  • the mass percentages of elements Al, Ti and Mo are limited in the alloys of the present invention to satisfy the relationship 11.59% ⁇ Al+Ti+Mo ⁇ 13.0%, so that the alloys not only have excellent creep resistance performance and long-lasting life, and has excellent welding performance.
  • the creep plastic elongation of the nickel-based deformed superalloy in the embodiment of the present invention is below 0.5%, and under the conditions of 89 MPa and 927 ° C.
  • the alloy has a lasting life of more than 200h, which can meet the requirements for the design and use of advanced aero-engines and gas turbines, and is suitable for manufacturing precision hot-end components for long-term service in equipment such as advanced aero-engines and gas turbines; 6.
  • the alloy of the embodiment of the present invention While having excellent creep resistance and long-lasting life, the alloy has a light weight because the density does not exceed 8.25g/cm 3 , which is conducive to reducing the fuel consumption of aero-engine and improving the maneuverability, and can meet the vibration requirements of the gas turbine during operation. As small as possible requirements, to prevent the formation of vibration damage.
  • the impurities are W ⁇ 0.50%, Fe ⁇ 1.50%, Si ⁇ 0.10%, Mn ⁇ 0.10%, P ⁇ 0.008%, S ⁇ 0.008%, Ta ⁇ 0.10%, Cu ⁇ 0.20%.
  • the mass percentage content of the elements Al, Sc and Ti in the alloy satisfies the relationship: 1.40% ⁇ (Al-1.8 Sc)/Ti ⁇ 2.6%.
  • the mass percentage content of the elements Al, Sc and Ti in the alloy satisfies the relationship: 2.22% ⁇ (Al-1.8 Sc)/Ti ⁇ 2.25%.
  • the creep-resistant, long-life nickel-based deformed superalloy according to the embodiment of the first aspect of the present invention includes: C: 0.04-0.08%, Cr: 18.50-21.50%, Co: 9.00-11.00%, Mo: 8.00 -9.00%, Al: 2.50-3.00%, Ti: 1.10-1.49%, Nb: 0.81-2.00%, B: 0.003-0.009%, Sc: 0.001-0.10%, the balance is nickel and inevitable impurities, with In terms of mass percentage content, the mass percentage contents of the elements Al, Ti and Mo in the alloy satisfy the relational formula: 11.59% ⁇ Al+Ti+Mo ⁇ 13.0%, and the mass percentage contents of the elements Al, Sc and Ti satisfy the relational formula : 2.22% ⁇ (Al-1.8Sc)/Ti ⁇ 2.25%.
  • Embodiments of the second aspect of the present invention provide for the use of creep-resistant, long-life nickel-based wrought superalloys in aeroengines.
  • the creep-resistant, long-life nickel-based deformation superalloy in the aircraft engine according to the embodiment of the second aspect of the present invention the creep-resistant, long-life nickel-based deformation of the embodiment of the first aspect of the present invention
  • Superalloys meet the requirements of advanced aero-engine design and use, and can be used in sophisticated equipment of advanced aero-engines.
  • Embodiments of the third aspect of the present invention provide for the use of creep-resistant, long-life nickel-based wrought superalloys in gas turbines.
  • the creep-resistant, long-life nickel-based deformed superalloy in the gas turbine according to the embodiment of the third aspect of the present invention the creep-resistant, long-life nickel-based deformed high temperature of the embodiment of the first aspect of the present invention
  • the alloy meets the requirements of gas turbine design and use, and can be used in the precision equipment of gas turbines.
  • the embodiment of the fourth aspect of the present invention provides a method for preparing a creep-resistant, long-life nickel-based deformed superalloy, which is characterized by comprising the following steps:
  • the alloy ingot obtained in the step a is directly forged and opened into an electrode rod without diffusion annealing, and remelted to obtain an alloy ingot, which is forged and opened into a desired billet, and heat-treated after processing.
  • the nickel-based deformed superalloy prepared by the method in the embodiment of the present invention has excellent creep resistance
  • the alloy ingots prepared by the preparation method of the embodiment of the present invention can be exempted from diffusion annealing treatment, or only need to adopt shorter alloy ingots. Diffusion annealing time can achieve the effect of homogenization of components, reduce energy consumption, shorten production cycle and improve production efficiency.
  • the step b when the diameter of the alloy ingot obtained by remelting is less than or equal to 200 mm, diffusion annealing is not performed, When the diameter of the alloy ingot obtained by remelting is greater than 200 mm, diffusion annealing is performed, the diffusion annealing temperature is 1150-1200° C., and the annealing time is 12-24 hours.
  • a creep-resistant, long-life nickel-based deformed superalloy comprises: C: 0.04-0.08%, Cr: 18.50-21.50%, Co: 9.00-11.00%, Mo: 8.00-9.00%, Al: 2.00-3.00%, Ti: 1.10-1.49%, Nb: 0.81-2.00%, B: 0.003-0.009%, Sc: 0.001-0.10%, the balance is nickel and inevitable impurities,
  • the mass percentage contents of the elements Al, Ti and Mo in the alloy satisfy the relational formula: 11.59% ⁇ Al+Ti+Mo ⁇ 13.0%.
  • the impurities are W ⁇ 0.50%, Fe ⁇ 1.50%, Si ⁇ 0.10%, Mn ⁇ 0.10%, P ⁇ 0.008%, S ⁇ 0.008%, Ta ⁇ 0.10%, Cu ⁇ 0.20%.
  • the nickel-based wrought superalloy according to the embodiment of the present invention may further include Zr of not more than 0.02%.
  • the strengthening element design scheme of high Al, low Ti, and high Nb is adopted, and the traditional Ni 3 (Al, Ti) strengthening phase is carried out Through modification, a Ni 3 (Al, Ti, Nb) strengthening phase with higher Al content and Nb is formed, which is more resistant to high temperature than the traditional Ni 3 (Al, Ti) strengthening phase; in the alloy of the embodiment of the present invention Sc element is added, and the addition of Sc introduces a new strengthening mechanism for the alloy of the embodiment of the present invention, forming a Ni 3 (Al, Ti, Nb) composite strengthening phase containing Sc, which is stronger than the single traditional strengthening phase Ni 3 (Al, Ti), Ni 3 (Al, Ti, Nb) are more resistant to high temperature and more stable, and significantly improve the creep resistance and lasting life of the alloy; the addition of Sc in the alloy of the embodiment of the present invention has a negative effect on the steel before the molten steel.
  • Sc refines the as-cast structure and significantly improves the dendrite segregation of the ingot.
  • This aspect improves the hot workability of the alloy, which can improve the cracking problem of the alloy in the direction of tensile stress, and prevent the alloy from forging, heat treatment.
  • Cracking during hot deformation such as rolling, on the other hand, the alloy ingots do not need to be subjected to high temperature and long time diffusion annealing or only need to use a short diffusion annealing time, which reduces energy consumption, reduces production costs, and at the same time shortens the production cycle.
  • the Sc added in the embodiment of the present invention has a post-processing effect on the grain boundary, that is, not only purifies the molten steel, but also in the liquid-solid transformation process and after the transformation is completed, Sc can still purify and strengthen the grain boundary, so that S, P, five harmful elements and other unavoidable low-melting impurity elements are difficult to segregate at the grain boundary, preventing the formation of creep voids at the grain boundary at high temperature;
  • the alloy of the embodiment of the present invention limits the elements Al, Ti and Mo The mass percentage content satisfies the relational formula 11.59% ⁇ Al+Ti+Mo ⁇ 13.0%, so that the alloy not only has excellent creep resistance and long-lasting life, but also has excellent welding performance.
  • the creep plastic elongation of the nickel-based deformed superalloy in the embodiment of the present invention is less than 0.5%, and under the conditions of 89MPa and 927°C, the durable life of the alloy reaches more than 200h, which can meet the design and use requirements of advanced aero-engines and gas turbines.
  • the alloy of the embodiment of the present invention has excellent creep resistance and long-lasting life, and at the same time, because the density does not exceed 8.25g/cm 3 , Therefore, the alloy has a light weight, which is conducive to reducing the fuel consumption of the aero-engine and improving the maneuverability. At the same time, it can meet the requirements of the gas turbine to have as little vibration as possible during the working process, and prevent the formation of vibration damage.
  • C mainly suppresses the growth of austenite grains during heating by forming MC-type carbides at the end of solidification in nickel-based superalloys, and forms M 23 C 6 and other types of carbides along the grain boundaries during heat treatment, which strengthens the crystallites.
  • the role of the boundary can delay the initiation, expansion and merging of creep voids, thereby increasing the high temperature durability of the alloy.
  • the C content is less than 0.04%, it is not enough to form a sufficient amount of MC and M 23 C 6 .
  • the C content is too high, the size of MC formed is larger, and it will consume too much Mo, Cr, Ti and Nb in the alloy.
  • the main function of Cr is to improve the oxidation resistance of the alloy, and has a certain solid solution strengthening effect. After aging treatment, it can also combine with C to form granular M 23 C 6 distributed along the grain, which can strengthen the grain boundary. effect.
  • the content of Cr is generally not more than 25%.
  • the The Cr content is controlled at 18.50-21.50%.
  • Co is both an important solid solution strengthening element and an important precipitation strengthening element.
  • Co element can be dissolved in the matrix to provide a good solid solution strengthening effect for the alloy, which can significantly reduce the stacking fault energy of the matrix, widen and expand the dislocation width, so that the dislocation is not easy to bunch up and cross-slip occurs, thereby improving the alloy. creep resistance and longevity.
  • Co can also partially replace the elements in the Ni 3 Al-type phase precipitation strengthening phase to improve the stability of the phase in long-term service; Co can also reduce the solid solubility of Al and Ti elements in the matrix, promote the precipitation of the ⁇ ' strengthening phase, and improve the stability of the phase. Increase its precipitation number and solution temperature. When the Co content is lower than 9%, the high temperature strength is low. When the Co content is higher than 11%, it is easy to form an ⁇ phase that affects its performance in long-term service, so the Co content is controlled at 9.00-11.00%.
  • Mo is one of the main solid-solution strengthening elements. It can be solid-dissolved in the alloy matrix as well as in the ⁇ ' strengthening phase. At the same time, it can improve the interatomic bonding force, increase the diffusion activation energy and recrystallization temperature, thereby effectively improving the High temperature strength. However, when Mo is too high, long-term high temperature aging is easy to generate ⁇ phase and reduce alloy toughness. Therefore, the Mo content is controlled at 8.00-9.00%.
  • W W and Mo have similar physical and chemical properties.
  • the role of W in nickel-based superalloys is mainly solid solution strengthening. Its atomic radius is relatively large, more than ten percent larger than the atomic radius of nickel, and the solid solution strengthening effect is obvious.
  • W is an element that accelerates high-temperature corrosion, and will form a harmful delta phase during long-term service, reducing the strength and toughness of the alloy.
  • the density of W is relatively high, and its density is 19.25 g/cm 3 . Adding a small amount of W to the nickel-based alloy can significantly increase the density of the alloy and increase the weight of the manufactured parts.
  • the alloy in the embodiment of the present invention is mainly used in aero-engines and gas turbines, the lighter the material is required, the better, so W is not added to the alloy in the embodiment of the present invention.
  • Al, Ti and Nb The three are the forming elements of the strengthening phase ⁇ ' in the aging-strengthened nickel-based alloy. It is generally believed that with the increase of the content of the three, the number of ⁇ ' increases, and the high-temperature creep and durability properties are improved, but too much ⁇ ' can deteriorate weldability and impair processability.
  • Ti and Nb will also combine with C to form MC-type carbides, which hinder the growth of grain boundaries and the sliding of grain boundaries at high temperatures, and play a role in improving high-temperature mechanical properties, but too much Ti and Nb will form large-grained MC-type carbides. Carbides are detrimental to the mechanical properties of the alloy.
  • the present invention finds through research that the high-temperature mechanical properties of the alloy depend not only on the amount of ⁇ ' phase, but also on its composition and characteristics. optimization, the best ⁇ ' strengthening effect can be obtained.
  • the strengthening element design scheme of high Al, low Ti and high Nb is adopted, and the traditional Ni 3 (Al, Ti) strengthening phase is modified to form a higher Al content and simultaneously contain Nb and
  • the Ni 3 (Al, Ti, Nb) of Sc is more resistant to high temperature than the traditional Ni 3 (Al, Ti) strengthening phase, thereby improving the creep resistance and long-lasting life of the alloy.
  • the specific control ranges of the three are: Al: 2.00-3.00%, preferably 2.50-3.00%, Ti: 1.10-1.49%, and Nb: 0.81-2.00%.
  • Sc Whether or not Sc is a rare earth element is still controversial in academia, and its role cannot be simply equated with rare earth elements.
  • the commonly used rare earth elements such as La, Ce, and Nd are rarely reported on the application of Sc in iron and steel materials, and the academic community has a relatively general understanding of the functions of rare earth elements.
  • Generalization the role of some rare earth elements is often generalized to the role of all rare earth elements, and there is a problem of generalization. It is generally believed that the role of rare earth elements is to remove inclusions, purify grain boundaries, improve oxidation resistance and There is little research and understanding on the role of each rare earth element in the adhesion of oxide films.
  • the appropriate amount of Sc added in the alloy of the embodiment of the present invention has three main effects: 1) The pretreatment effect of Sc on the molten steel, Sc improves the solidification nucleation rate, refines the as-cast grains, and significantly It improves the dendrite segregation of the ingot, which improves the hot workability of the alloy, improves the cracking problem of the alloy in the direction of tensile stress, and prevents the alloy from cracking during hot deformation such as forging and hot rolling.
  • the ingot does not need to undergo high-temperature and long-term diffusion annealing or only needs to use a short diffusion annealing time, which reduces energy consumption, reduces production costs, shortens production cycle, and improves production efficiency; 2)
  • the addition of Sc introduces a The new strengthening mechanism forms a Ni 3 (Al, Ti, Nb ) composite strengthening phase containing Sc, which is more resistant to high temperature and It is more stable and significantly improves the creep resistance and long-lasting life of the alloy;
  • Sc has a post-treatment effect on grain boundaries, that is, not only purifying molten steel, but also during the liquid-solid transformation and after the transformation is completed.
  • B The role of B is mainly manifested in two aspects. First, because the atomic radius of B is very small, only about 85 picometers, while the atomic radius of Ni is about 135 picometers, so B atoms are easily enriched at the grain boundaries, making harmful The low melting point elements cannot segregate at the grain boundary, which improves the bonding force of the grain boundary; the second is that the boride on the grain boundary can prevent the grain boundary from slipping, the initiation and expansion of voids, and improve the creep resistance and lasting life of the alloy. favorable. But too much B will deteriorate the hot workability and welding performance of the alloy, so the suitable B content of the alloy in the embodiment of the present invention is 0.003-0.009%.
  • Fe is a harmful element in nickel-based superalloys, but it is unavoidable in industrial production. In the alloys of the embodiments of the present invention, no more than 1.50% Fe is allowed, which makes it economical in industrial production. The production cost of the alloy is controlled at a reasonable level by using raw materials and return materials containing trace Fe.
  • Zr helps to purify the grain boundary and enhance the bonding force of the grain boundary.
  • the compound addition of Zr and B helps to maintain the high temperature strength and long-lasting life of the alloy, but excessive Zr easily causes hot working cracking and damages the welding performance.
  • the present invention The alloys in the examples control Zr to be ⁇ 0.02%.
  • Ni is the most important matrix element and the formation element of the precipitation strengthening phase ⁇ '. With Ni as the matrix, a large amount of alloying elements with different functions can be solid solution, such as Cr, Mo, Co, C, etc., Ni-based superalloys
  • the microstructure has strong stability, excellent high-temperature strength, toughness and processing and forming properties, and is suitable for the manufacture of creep-resistant, long-life precision hot-end components for advanced aero-engines and gas turbines.
  • the mass percentage content of the elements Al, Sc and Ti in the alloy satisfies the relationship: 1.40% ⁇ (Al-1.8 Sc)/Ti ⁇ 2.6%, preferably 2.22% ⁇ (Al-1.8Sc)/Ti ⁇ 2.25%.
  • the nickel-based deformed superalloy according to the embodiment of the present invention can further improve the precipitation phase when the elements Al, Sc and Ti satisfy the condition of 1.40% ⁇ (Al-1.8Sc)/Ti ⁇ 2.6%.
  • the thermal stability of the alloy significantly improves the creep resistance and long-lasting life of the nickel-based deformed superalloy.
  • the durable life under the condition of °C can reach more than 300h, especially when 2.22% ⁇ (Al-1.8Sc)/Ti ⁇ 2.25%, under the condition of 89MPa and 927°C, the durability of the nickel-based deformed superalloy of the embodiment of the present invention is The life can reach more than 330h.
  • Embodiments of the second aspect of the present invention provide for the use of creep-resistant, long-life nickel-based wrought superalloys in aeroengines.
  • the creep-resistant, long-life nickel-based deformed superalloy of the embodiment of the first aspect of the present invention satisfies the requirements for the design and use of advanced aero-engines, and can be applied to precision equipment of advanced aero-engines.
  • Embodiments of the third aspect of the present invention provide for the use of creep-resistant, long-life nickel-based wrought superalloys in gas turbines.
  • the creep-resistant, long-life nickel-based deformed superalloy according to the first aspect of the present invention satisfies the design and use requirements of gas turbines, and can be applied to precision equipment of gas turbines.
  • the embodiment of the fourth aspect of the present invention provides a method for preparing a creep-resistant, long-life nickel-based deformed superalloy, which is characterized by comprising the following steps:
  • the alloy ingot obtained in the step a is forged into an electrode rod, and remelted to obtain an alloy ingot.
  • diffusion annealing is not performed.
  • the diameter of the alloy ingot obtained by remelting is less than or equal to 200 mm >200mm, carry out diffusion annealing, diffusion annealing temperature is 1150-1200 °C, annealing time is 12-24 hours, forging and opening into required blanks for further processing into forgings, plates, strips, bars, pipes, Wire (including welding wire) or material for powder metallurgy, followed by heat treatment.
  • the obtained nickel-based deformed superalloy has excellent creep resistance, long service life and excellent welding performance, It can meet the requirements for the design and use of advanced aero-engines and gas turbines, and the alloy ingot prepared by the preparation method of the embodiment of the present invention can be free of diffusion annealing treatment, or only need to adopt a shorter diffusion annealing time to achieve the effect of uniform composition , reducing energy consumption, while shortening the production cycle and improving production efficiency.
  • the heat treatment includes one solution and two aging, first solution heat treatment, welding, cold bending After the forming process, two aging heat treatments are performed.
  • the solid solution system is: the solution temperature is 1100-1170°C, the temperature control accuracy is within ⁇ 10°C, the holding time is determined according to the size of the product, and it is cooled by air cooling or water cooling; the primary aging system is: 950-1010°C, The temperature control accuracy is within ⁇ 10°C, the temperature is kept for 1-3h, and the cooling rate is equivalent to air cooling; the secondary aging system is: 750-800°C, the temperature control accuracy is within ⁇ 10°C, and the holding time is 8- 10h, cool at a cooling rate equivalent to air cooling.
  • the temperature is 1180°C, and the time is 12h; the forging temperature is 1100°C, and it is forged into a 40mm thick slab after three fires, and then rolled into a 20mm thick plate and a 5mm thick plate; the 20mm thick plate is solid solution at 1150°C for 1h , water cooling, and then carry out aging treatment of 1010°C ⁇ 2h air cooling + 788°C ⁇ 8h air cooling to obtain creep-resistant, long-life nickel-based deformed superalloy, and test its mechanical properties; Using the welding wire of the same composition as the plate, two pieces of 5mm thick plates were welded together by TIG fusion welding method, and their welding performance was tested.
  • the composition of the alloy prepared in Example 1 is shown in Table 1, and the properties are shown in Table 2.
  • Example 3-10 The preparation methods of Examples 3-10 are the same as those of Example 1, except that the alloy components are different.
  • the alloy components prepared in Examples 3-10 are shown in Table 1, and the properties are shown in Table 2.
  • the alloy prepared in Comparative Example 1 The ingredients are shown in Table 1, and the properties are shown in Table 2.
  • the alloy prepared in Comparative Example 2 The ingredients are shown in Table 1, and the properties are shown in Table 2.
  • the alloy prepared in Comparative Example 3 The ingredients are shown in Table 1, and the properties are shown in Table 2.
  • the alloy obtained in Comparative Example 4 The ingredients are shown in Table 1, and the properties are shown in Table 2.
  • Comparative Example 5 The preparation method of Comparative Example 5 is the same as that of Example 1. The difference is that in the alloy composition, the content of Sc is 0.0005%.
  • Comparative Example 6 The preparation method of Comparative Example 6 is the same as that of Example 1, and the difference is that in the alloy composition, the content of Sc is 0.19%.
  • Comparative Example 7 The preparation method of Comparative Example 7 is the same as that of Example 1, except that the content of Nb in the alloy composition is 0.25%.
  • Comparative Example 8 The preparation method of Comparative Example 8 is the same as that of Example 1. The difference is that in the alloy composition, the content of Nb is 2.23%.
  • Comparative Example 9 The preparation method of Comparative Example 9 is the same as that of Example 1. The difference is that in the alloy composition, the content of Al is 1.45%.
  • Comparative Example 10 The preparation method of Comparative Example 10 is the same as that of Example 1. The difference is that in the alloy composition, the content of Al is 3.15%.
  • Comparative Example 11 The preparation method of Comparative Example 11 is the same as that of Example 1, except that in the alloy composition, the content of Ti is 0.95%.
  • Comparative Example 12 The preparation method of Comparative Example 12 is the same as that of Example 1. The difference is that in the alloy composition, the content of Ti is 1.55%.
  • Comparative Example 13 is a commercial Haynes 282 alloy, that is, the alloy composition meets the alloy disclosed in the patent application No. 201210057737.8.
  • the alloy composition in Test Example H prepared according to the method of the patent is: C: 0.088%, Cr: 19.3% , Co: 10.8%, Mo: 4.6%, Al: 1.63%, Ti: 1.85%, B: 0.003%, Nb: 0.04%, Fe: 0.2%, W: 6.1%, the balance is Ni, the comparative example 13
  • the properties of the obtained alloys are shown in Table 2.
  • Comparative Example 14 is the alloy disclosed in the patent application No. 201910811805.7.
  • the alloy composition of Example 2 prepared according to the method in the patent is: C: 0.04%, Cr: 18.3%, Co: 9.5%, Mo: 8.5 %, Al: 1.5%, Ti: 1.9%, Zr: 0.02%, B: 0.005%, Nb: 0.2%, V: 0.05%, Fe: 1.2%, Si: 0.1%, Mn: 0.2%, P: 0.006 %, S: 0.001%, Nd: 0.02%, and the balance is Ni.
  • the properties of the alloy prepared in Comparative Example 14 are shown in Table 2.
  • ⁇ p is the creep plastic elongation of aged alloy at 816°C, 221MPa , 100h;
  • ⁇ is the permanent life of the aged alloy at 89MPa and 927°C, and ⁇ is the permanent elongation after fracture of the aged alloy at 89MPa and 927°C;
  • R p0.2 is the room temperature tensile yield strength of the aged alloy
  • R m is the room temperature tensile strength of the aged alloy
  • A is the room temperature tensile elongation after fracture of the aged alloy
  • the detection conditions for welding cracks are: according to the national energy industry standard NB/T 47013.5-2015, the surface quality of 5mm sheet welds shall be detected by two methods of fluorescence penetration and color penetration, and 5mm shall be detected by X-ray according to the national standard GB/T 3323.1-2019. Internal quality of sheet welds.
  • the mass percentage contents of the elements Al, Ti and Mo satisfy the relational formula: 11.59 ⁇ Al+Ti+Mo ⁇ 13.0, and the mass percentage contents of the elements Al, Sc and Ti satisfy the relational formula: 1.40 ⁇ (Al-1.8Sc)/Ti ⁇ 2.6, under these two conditions, the nickel-based deformed superalloys prepared in Examples 1-6 have very excellent creep resistance and high-temperature durability, at 816°C, 221MPa The creep plastic elongation of the alloy is below 0.2% under the conditions of 100h and 100h, and the durable life under the conditions of 89MPa and 927°C is more than 300h.
  • Comparative Examples 1-2 the sum of the mass percentages of elements Al, Ti and Mo, that is, Al+Ti+Mo, were 11.49% and 11.40%, respectively, although the alloys in Comparative Examples 1 and 2 were at 89MPa and 927°C.
  • the lasting life of the alloy has basically reached more than 200h, but the creep plastic elongation of the alloy is more than 0.5% under the conditions of 816°C, 221MPa and 100h, and the room temperature tensile strength Rm does not meet the requirements, which are lower than The expected value is 1035MPa.
  • Comparative Examples 3-4 the sum of the mass percentages of elements Al, Ti and Mo, that is, Al+Ti+Mo, was 13.34% and 13.25%, respectively, although the alloys in The creep plastic elongation under the condition is below 0.2%, and the permanent life under the condition of 89MPa and 927°C basically reaches 300h, but welding cracks appear during the welding process.
  • the sum of the mass percentages of elements Al, Ti and Mo in the alloys of Comparative Example 3 and Comparative Example 4 exceeds 13%, the precipitation strengthening and solid solution strengthening are too strong, and the aging yield strength and tensile strength are too large. Although it has excellent creep resistance and long life, it has poor weldability and weld cracks.
  • the Sc content was 0.0005% and 0.19%, respectively.
  • the Sc content in the alloy of Comparative Example 5 was 0.0005%, due to the low Sc content, the alloy's performance at 816 ° C, 221 MPa, and 100 h was reduced.
  • the creep plastic elongation is 0.878%, and the durable life at 89MPa and 927°C is only 103.5 hours; the Sc content in the alloy of Comparative Example 6 is 0.19%.
  • the creep plastic elongation of the alloy under the condition is 1.389%, and the permanent life under the condition of 89MPa and 927°C is only 45 hours.
  • the amount of Sc added is very important to the creep resistance and longevity of the alloy.
  • the content of Sc is 0.001-0.1%.
  • the Nb content was 0.25% and 2.23%, respectively.
  • low Nb was used, and the Nb content was 0.25%.
  • the length is 0.808%, and the durable life under the condition of 89MPa and 927°C is 180 hours, which does not reach the expected value.
  • the content of Nb is 2.23%, and the creep plasticity of the alloy under the conditions of 816°C, 221MPa and 100h The elongation is 0.126%, and the permanent life at 89MPa and 927°C is 352.4 hours.
  • the alloy of Comparative Example 8 has excellent creep resistance and permanent life, its permanent fracture at 89MPa and 927°C The elongation was only 5.8%, which did not reach the expected value of ⁇ 10.0%.
  • increasing the Nb content in the alloy can improve the creep resistance and longevity, if the Nb content is too high, the elongation after fracture will be reduced. Therefore, the Nb content in the alloy of the embodiment of the present invention is controlled to be 0.81-2.00%.
  • Comparative Example 9 and Comparative Example 10 the Al content was 1.45% and 3.15%, respectively.
  • low Al was used, and the Al content was 1.45%.
  • the plastic elongation is 0.488%, and the durable life at 89MPa and 927°C is only 105 hours.
  • the creep plastic elongation of the alloy in Comparative Example 9 reaches the expected value, the durable life is only 105 hours; Comparative Example 10 Among them, the Al content is 3.15%, the creep plastic elongation of the alloy is 0.178% under the conditions of 816 °C, 221 MPa, and 100 h, and the permanent life under the conditions of 89 MPa and 927 °C reaches 352.3 hours, although the alloy of Comparative Example 10 has Excellent creep resistance and durable life, but its elongation after fracture at 89MPa and 927°C is only 8.6%, which does not reach the expected value of ⁇ 10.0%. Although increasing the Al content in the alloy can improve the creep resistance and longevity, if the Al content is too high, the elongation after fracture will be reduced. Therefore, the Al content in the alloy of the embodiment of the present invention is controlled to be 2.00-3.00%.
  • Comparative Examples 11 and 12 the Ti content was 0.95% and 1.55%, respectively.
  • the Ti content was 0.95%, and the creep plastic elongation of the alloy was 0.187 at 816 °C, 221 MPa, and 100 h. %, the durable life at 89MPa and 927°C is 212.5 hours.
  • the creep resistance and durable life of the alloy in Comparative Example 11 can meet the requirements, its room temperature tensile strength Rm is 1032MPa, which is lower than the expected value of 1035MPa , did not meet the requirements; in Comparative Example 12, the Ti content was 1.55%, the creep plastic elongation of the alloy was 0.14% under the conditions of 816 °C, 221 MPa, and 100 h, and the permanent life under the conditions of 89 MPa and 927 °C was 345.6 hours. , although the alloy of Comparative Example 12 has excellent creep resistance and longevity, its elongation after fracture at 89MPa and 927°C is only 9.5%, which does not reach the expected value of ⁇ 10.0%. In the alloy of the embodiment of the present invention, the Ti content is preferably 1.10 to 1.49%.
  • Comparative example 13 is a commercial 282 alloy
  • comparative example 14 is an alloy disclosed in application number 201910811805.7, the alloy density of these two comparative examples is higher than that of the embodiment, and Sc is not added to the alloys in comparative example 13 and comparative example 14, and
  • Comparative Example 13 low Al, high Ti, and low Nb were used, and in Comparative Example 14, low Al and low Nb were used.
  • the alloy in Comparative Example 13 was tested at 816 ° C, 221 MPa, and 100 h under the conditions of alloy creep The plastic elongation is 1.1%, and the permanent life under the conditions of 89MPa and 927°C is 102 hours.
  • the alloy in Comparative Example 14 has a creep plastic elongation of 2.201% under the conditions of 816°C, 221MPa and 100h.
  • the durable life under the conditions of 89MPa and 927°C is 40.3 hours, which has not reached the expected value, and cannot meet the requirements of precision hot-end components for long-term service in equipment such as advanced aero-engines and gas turbines.
  • the terms “one embodiment,” “some embodiments,” “example,” “specific example,” or “some examples” and the like mean a specific feature, structure, material, or description described in connection with the embodiment or example. Features are included in at least one embodiment or example of the invention. In this specification, schematic representations of the above terms are not necessarily directed to the same embodiment or example. Furthermore, the particular features, structures, materials or characteristics described may be combined in any suitable manner in any one or more embodiments or examples. Furthermore, those skilled in the art may combine and combine the different embodiments or examples described in this specification, as well as the features of the different embodiments or examples, without conflicting each other.

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Abstract

本发明公开了一种抗蠕变、长寿命镍基变形高温合金，其包括：C：0.04-0.08％、Cr：18.50-21.50％、Co：9.00-11.00％、Mo：8.00-9.00％、Al：2.00-3.00％、Ti：1.10-1.49％、Nb：0.81-2.00％、B：0.003-0.009％、Sc：0.001-0.10％，余量为镍和不可避免的杂质，以质量百分含量计，该合金中元素Al、Ti和Mo的质量百分含量满足关系式：11.59％≤Al+Ti+Mo≤13.0％。本发明的镍基变形高温合金，具有优异的抗蠕变性能和持久寿命，能够满足先进航空发动机和燃气轮机设计和使用的要求，适用于制造先进航空发动机和燃气轮机等装备中长期服役的精密热端部件。

Description

抗蠕变、长寿命镍基变形高温合金及其制备方法和应用
相关申请的交叉引用
本申请要求申请号为202011262753.1、申请日为2020年11月12日的中国专利申请的优先权和权益，上述中国专利申请的全部内容在此通过引用并入本申请。
技术领域
本发明属于金属材料领域，具体涉及一种抗蠕变、长寿命镍基变形高温合金，特别地还涉及该抗蠕变、长寿命镍基变形高温合金的制备方法，进一步地，还涉及该抗蠕变、长寿命镍基变形高温合金的应用。
背景技术
航空发动机和燃气轮机等装备中有大量精密热端部件，其特点是在高温和大载荷条件下具有良好尺寸精度以保证功能稳定性，一般采用高温合金制造。这些部件中包含筒件、管件和紧固件等，形状复杂、承受温度高(一般为600-800℃)，是由板材、管材或棒材经冷弯、焊接、车、铣等多工序加工而成，对合金加工性能、焊接性能、抗蠕变性能、持久寿命等要求比较高，目前使用的合金有Nimonic263、R-41、Waspaloy、Haynes282、HastelloyX、Haynes230、Inconel718等。
随着航空发动机和燃气轮机设计水平和应用技术的不断发展，燃气初温越来越高，以燃气轮机为例，目前最先进的G级和H级燃气轮机燃气初温达到了1450～1500℃，未来的J级燃气轮机燃气初温将达到1600～1700℃，航空发动机和燃气轮机精密热端部件合金本体承受的温度也将达到800～950℃，因此对合金的800～950℃高温力学性能提出了日益严苛的要求。
Nimonic263、HastelloyX、Haynes230合金虽然具有良好的加工性能，但这些合金只能在800℃以下长期服役，在800℃以上高温强度和抗蠕变性能严重不足；Inconel718合金虽然具有良好的加工性能，但只能在650℃以下服役，在更高温度下组织将失去稳定性而发生性能劣化。R-41和Waspaloy的强化相γ′含量高、析出速度快，无论是坯料的热加工(锻造、热轧)、热处理还是零件的冷加工(冷弯、车削、焊接等)难度都比较大，不适用于制造航空发动机和燃气轮机中需要经历复杂加工工艺的精密热端部件。而且R-41在816℃、221MPa、100h条件下的蠕变塑性伸长率(ε p)约为1％，Waspaloy对应的这一数值大于1％，这两种合金在89MPa、927℃条件下的持久寿命(τ)都在100h以下，可见这两种合金的抗蠕变性能和持久寿命都不满足先进航空发动机和燃气轮机的设计要求。Haynes282合金具 有较好的室温和高温力学性能，同时易于加工和焊接，但也存在三点不足：1)抗蠕变性能偏低，在816℃、221MPa、100h条件下的蠕变塑性伸长率约为1％；2)持久寿命不足，在89MPa、927℃条件下的持久寿命在200h以下；3)允许的服役温度较低，目前在实际工程应用中考虑到燃气轮机的大修周期在8年左右，为了确保使用Haynes282合金制造的部件安全度过8年检修周期，其最高使用温度被严格限定在800℃左右。
因此，需要研制一种具有优异的抗蠕变性能和持久寿命的镍基变形高温合金。
发明内容
本发明是基于发明人对以下事实和问题的发现和认识做出的：目前先进航空发动机和燃气轮机不仅对精密热端部件的初始加工精度和装配精度要求极高，而且要求在800-950℃高温长期服役过程中不能发生过量塑性变形，即要求合金具有优异的抗蠕变性能，以免在检修周期到来前发生零件失效或在检修时难以拆卸和更换。此外，希望进一步延长精密热端部件的检修周期，因此对合金的高温持久寿命提出了更严格的要求。在具体性能指标上，希望合金在816℃、221MPa、100h条件下的蠕变塑性伸长率(ε p)在0.5％以下，在89MPa、927℃条件下的持久寿命(τ)达到200h以上，同时其他主要力学性能指标不低于现有合金。目前精密热端部件使用的合金，如Nimonic263、R-41、Waspaloy、Haynes282、HastelloyX、Haynes230、Inconel718等均不能满足上述抗蠕变、长寿命的要求。
本发明旨在至少在一定程度上解决相关技术中的技术问题之一。
为此，本发明第一个方面的实施例提出一种抗蠕变、长寿命镍基变形高温合金，该合金具有优异的抗蠕变性能和持久寿命，在816℃、221MPa、100h条件下的蠕变塑性伸长率在0.5％以下，89MPa、927℃条件下的持久寿命在200h以上，使合金能够满足先进航空发动机和燃气轮机设计和使用的要求，适用于制造先进航空发动机和燃气轮机等装备中长期服役的精密热端部件。
根据本发明第一个方面实施例的一种抗蠕变、长寿命镍基变形高温合金，其包括：C：0.04-0.08％、Cr：18.50-21.50％、Co：9.00-11.00％、Mo：8.00-9.00％、Al：2.00-3.00％、Ti：1.10-1.49％、Nb：0.81-2.00％、B：0.003-0.009％、Sc：0.001-0.10％，余量为镍和不可避免的杂质，以质量百分含量计，该合金中元素Al、Ti和Mo的质量百分含量满足关系式：11.59％≤Al+Ti+Mo≤13.0％。
根据本发明第一个方面实施例的抗蠕变、长寿命镍基变形高温合金带来的优点和技术效果：1、本发明实施例中采用高Al、低Ti、高Nb的强化元素设计方案，对传统的Ni 3(Al、Ti)强化相进行了改性，形成了含Al更高并同时含有Nb的Ni 3(Al、Ti、Nb)强化相，比传统的Ni 3(Al、Ti)强化相更耐高温；2、本发明实施例的合金中加入了Sc元素，Sc的添加为本发明实施例的合金引入了一种新的强化机制，形成了含Sc的Ni 3(Al、Ti、Nb)复合强化相，比单一的传统强化相Ni 3(Al、Ti)、Ni 3(Al、Ti、Nb)更耐高温、更加稳定， 显著提高了合金的抗蠕变性能和持久寿命；3、本发明实施例的合金中Sc的添加具有对钢液前处理的作用，Sc细化了铸态组织，显著改善了铸锭的枝晶偏析，这一方面提高了合金的热加工性能，可改善合金在拉应力方向的开裂问题，防止合金在锻造、热轧等热变形过程中开裂，另一方面使得合金锭不必进行高温长时间的扩散退火或仅需要采用较短的扩散退火时间，降低了能源消耗，可降低生产成本，同时缩短了生产周期，提高了生产效率；4、本发明实施例中加入的Sc具有对晶界的后处理作用，即不仅净化钢液，而且在液-固转变过程中和转变完成后，Sc仍能净化并强化晶界，使得S、P、五害元素和其他不可避免的低熔点杂质元素难以在晶界偏聚，防止晶界在高温下形成蠕变空洞；5、本发明实施例的合金中限定了元素Al、Ti和Mo的质量百分含量满足关系式11.59％≤Al+Ti+Mo≤13.0％，使合金不仅具有优异的抗蠕变性能和持久寿命，而且具有优异的焊接性能，在816℃、221MPa、100h条件下，本发明实施例的镍基变形高温合金的蠕变塑性伸长率在0.5％以下，在89MPa、927℃条件下，合金的持久寿命达到200h以上，能够满足先进航空发动机和燃气轮机设计和使用的要求，适用于制造先进航空发动机和燃气轮机等装备中长期服役的精密热端部件；6、本发明实施例的合金在具有优异的抗蠕变性能和持久寿命的同时，因密度不超过8.25g/cm 3，所以合金自重轻，有利于降低航空发动机燃料消耗和提高机动性能，同时能够满足燃气轮机在工作过程中振动尽可能小的要求，防止形成振动破坏。
根据本发明第一个方面实施例的抗蠕变、长寿命镍基变形高温合金，其中，所述杂质为W≤0.50％，Fe≤1.50％，Si≤0.10％，Mn≤0.10％，P≤0.008％，S≤0.008％，Ta≤0.10％，Cu≤0.20％。
根据本发明第一个方面实施例的抗蠕变、长寿命镍基变形高温合金，其中，所述合金中元素Al、Sc和Ti的质量百分含量满足关系式：1.40％≤(Al-1.8Sc)/Ti≤2.6％。
根据本发明第一个方面实施例的抗蠕变、长寿命镍基变形高温合金，其中，所述合金中元素Al、Sc和Ti的质量百分含量满足关系式：2.22％≤(Al-1.8Sc)/Ti≤2.25％。
根据本发明第一个方面实施例的抗蠕变、长寿命镍基变形高温合金，其中，所述合金还包括不大于0.02％的Zr。
根据本发明第一个方面实施例的抗蠕变、长寿命镍基变形高温合金，其中，包括：C：0.04-0.08％、Cr：18.50-21.50％、Co：9.00-11.00％、Mo：8.00-9.00％、Al：2.50-3.00％、Ti：1.10-1.49％、Nb：0.81-2.00％、B：0.003-0.009％、Sc：0.001-0.10％，余量为镍和不可避免的杂质，以质量百分含量计，该合金中元素Al、Ti和Mo的质量百分含量满足关系式：11.59％≤Al+Ti+Mo≤13.0％，元素Al、Sc和Ti的质量百分含量满足关系式：2.22％≤(Al-1.8Sc)/Ti≤2.25％。
本发明第二个方面的实施例提供了抗蠕变、长寿命镍基变形高温合金在航空发动机中的应用。
根据本发明第二个方面实施例的抗蠕变、长寿命镍基变形高温合金在航空发动机中的应用的优点和技术效果：本发明第一方面实施例的抗蠕变、长寿命镍基变形高温合金满足 了先进航空发动机设计和使用的要求，能够应用在先进航空发动机的精密设备中。
本发明第三个方面的实施例提供了抗蠕变、长寿命镍基变形高温合金在燃气轮机中的应用。
根据本发明第三个方面实施例的抗蠕变、长寿命镍基变形高温合金在燃气轮机中的应用的优点和技术效果：本发明第一方面实施例的抗蠕变、长寿命镍基变形高温合金满足了燃气轮机设计和使用的要求，能够应用在燃气轮机的精密设备中。
本发明第四个方面的实施例提供了一种抗蠕变、长寿命镍基变形高温合金的制备方法，其特征在于，包括如下步骤：
a、取本发明第一个方面实施例中合金设计用量的原料在真空下熔炼，精炼去除气体，真空下浇铸成合金锭；
b、所述步骤a得到的合金锭不必进行扩散退火直接锻造开坯成电极棒，重熔，得到合金锭，锻造开坯成所需坯料，加工后进行热处理。
根据本发明第四个方面实施例的抗蠕变、长寿命镍基变形高温合金的制备方法的优点和技术效果：本发明实施例的方法制得的镍基变形高温合金，具有优异的抗蠕变性能、持久寿命和优异的焊接性能，能够满足先进航空发动机和燃气轮机设计和使用的要求，而且本发明实施例的制备方法制备的合金锭可以免去扩散退火处理，或者仅需要采用较短的扩散退火时间即可达到成分均匀化的效果，降低了能源消耗，同时缩短了生产周期，提高了生产效率。
根据本发明第四个方面实施例的抗蠕变、长寿命镍基变形高温合金的制备方法，其中，所述步骤b中，当重熔得到的合金锭直径≤200mm时，不进行扩散退火，当重熔得到的合金锭直径＞200mm时，进行扩散退火，扩散退火温度为1150-1200℃，退火时间为12-24小时。
具体实施方式
下面详细描述本发明的实施例，实施例是示例性的，旨在用于解释本发明，而不能理解为对本发明的限制。
根据本发明第一个方面实施例的一种抗蠕变、长寿命镍基变形高温合金，其包括：C：0.04-0.08％、Cr：18.50-21.50％、Co：9.00-11.00％、Mo：8.00-9.00％、Al：2.00-3.00％、Ti：1.10-1.49％、Nb：0.81-2.00％、B：0.003-0.009％、Sc：0.001-0.10％，余量为镍和不可避免的杂质，以质量百分含量计，该合金中元素Al、Ti和Mo的质量百分含量满足关系式：11.59％≤Al+Ti+Mo≤13.0％。优选地，所述杂质为W≤0.50％，Fe≤1.50％，Si≤0.10％，Mn≤0.10％，P≤0.008％，S≤0.008％，Ta≤0.10％，Cu≤0.20％。进一步地，本发明实施例的镍基变形高温合金还可以包括不大于0.02％的Zr。
根据本发明第一个方面实施例的抗蠕变、长寿命镍基变形高温合金：采用高Al、低Ti、 高Nb的强化元素设计方案，对传统的Ni 3(Al、Ti)强化相进行了改性，形成了含Al更高并同时含有Nb的Ni 3(Al、Ti、Nb)强化相，比传统的Ni 3(Al、Ti)强化相更耐高温；本发明实施例的合金中加入了Sc元素，Sc的添加为本发明实施例的合金引入了一种新的强化机制，形成了含Sc的Ni 3(Al、Ti、Nb)复合强化相，比单一的传统强化相Ni 3(Al、Ti)、Ni 3(Al、Ti、Nb)更耐高温、更加稳定，显著提高了合金的抗蠕变性能和持久寿命；本发明实施例的合金中Sc的添加具有对钢液前处理的作用，Sc细化了铸态组织，显著改善了铸锭的枝晶偏析，这一方面提高了合金的热加工性能，可改善合金在拉应力方向的开裂问题，防止合金在锻造、热轧等热变形过程中开裂，另一方面使得合金锭不必进行高温长时间的扩散退火或仅需要采用较短的扩散退火时间，降低了能源消耗，可降低生产成本，同时缩短了生产周期，提高了生产效率；本发明实施例中加入的Sc具有对晶界的后处理作用，即不仅净化钢液，而且在液-固转变过程中和转变完成后，Sc仍能净化并强化晶界，使得S、P、五害元素和其他不可避免的低熔点杂质元素难以在晶界偏聚，防止晶界在高温下形成蠕变空洞；本发明实施例的合金中限定了元素Al、Ti和Mo的质量百分含量满足关系式11.59％≤Al+Ti+Mo≤13.0％，使合金不仅具有优异的抗蠕变性能和持久寿命，而且具有优异的焊接性能，在816℃、221MPa、100h条件下，本发明实施例的镍基变形高温合金的蠕变塑性伸长率在0.5％以下，在89MPa、927℃条件下，合金的持久寿命达到200h以上，能够满足先进航空发动机和燃气轮机设计和使用的要求，适用于制造先进航空发动机和燃气轮机等装备中长期服役的精密热端部件；本发明实施例的合金在具有优异的抗蠕变性能和持久寿命的同时，因密度不超过8.25g/cm 3，所以合金自重轻，有利于降低航空发动机燃料消耗和提高机动性能，同时能够满足燃气轮机在工作过程中振动尽可能小的要求，防止形成振动破坏。
本发明第一个方面实施例中抗蠕变、长寿命镍基变形高温合金中各成分的作用如下：
C：C在镍基高温合金中主要通过在凝固末期形成MC型碳化物抑制加热时奥氏体晶粒长大，在热处理时沿晶界形成M 23C 6等类型碳化物，起到强化晶界的作用，能够延缓蠕变空洞萌生、扩展和合并，从而提高合金的高温持久寿命，当C含量小于0.04％时，不足以形成足够数量MC和M 23C 6。当C含量过高时形成的MC尺寸较大，并且会过多的消耗合金中的Mo、Cr、Ti和Nb，一方面不仅减少了Mo、Cr的固溶强化作用，另一方面用于形成Ni 3(Al、Ti)和Ni 3(Al、Ti、Nb)复合强化相的Ti和Nb将会减少，对合金的高温抗蠕变性能和持久性能产生不利影响，因此C应控制在不超0.08％。
Cr：Cr最主要的作用是提高合金的抗氧化性能，并具有一定的固溶强化效果，在时效处理后还可以与C结合形成沿晶分布的粒状M 23C 6，起到强化晶界的作用。但Cr含量过高时易于形成拓普密排相，降低合金长期组织性能稳定性，因此其含量一般不超过25％，本发明实施例中考虑兼顾抗氧化性能和长期组织性能的稳定性，将Cr含量控制在18.50-21.50％。
Co：Co既是重要的固溶强化元素，也是重要的析出强化元素。Co元素可固溶于基体中为合金提供良好的固溶强化效果，可显著降低基体堆垛层错能，拉宽扩展位错宽度，使位错不易束集而发生交滑移，从而提高合金的抗蠕变性能和持久寿命。Co也可部分替代Ni 3Al型相析出强化相中的元素，改善相长期服役中的稳定性；Co还可以降低Al、Ti元素在基体中的固溶度，促进γ′强化相的析出并提高其析出数量和固溶温度。当Co含量低于9％时，高温强度偏低，当Co含量高于11％时，在长期服役中易形成影响其性能的η相，因此将Co含量控制在9.00-11.00％。
Mo：Mo是主要的固溶强化元素之一，既可固溶于合金基体又可固溶于γ′强化相，同时可提高原子间结合力，提高扩散激活能和再结晶温度，从而有效提高高温强度。但Mo过高时长期高温时效易于生成μ相而降低合金韧性。因此，将Mo含量控制在8.00-9.00％。
W：W与Mo具有相似的物理化学性质，W在镍基高温合金中的作用主要是固溶强化，其原子半径比较大，比镍原子半径大百分之十几，固溶强化作用明显。但W是加速高温腐蚀的元素，而且在长期服役时会形成有害的δ相，降低合金强度和韧性。此外，W的密度较大，其密度为19.25g/cm 3，镍基合金中添加少量W就会明显提高合金密度，使所制造的部件重量增加。考虑到本发明实施例中的合金主要用在航空发动机和燃气轮机上，要求材料越轻越好，因此本发明实施例的合金中不添加W。
Al、Ti和Nb：三者是时效强化型镍基合金中强化相γ′的形成元素，一般认为随三者含量的增加，γ′数量增加，高温蠕变和持久性能提高，但过多的γ′会恶化焊接性能和损害加工性能。另外，Ti、Nb还会与C结合形成MC型碳化物，在高温时阻碍晶界长大和晶界滑动，起到提高高温力学性能的作用，但过多的Ti、Nb会形成大颗粒MC型碳化物，对合金的力学性能反而不利。本发明通过研究发现，合金的高温力学性能不仅取决于γ′相的多少，而且取决于其成分组成和特性，在Al、Ti、Nb总量不变的前提下，通过Al、Ti、Nb比例的优化，可以得到最佳的γ′强化效果。本发明实施例的合金中采用高Al、低Ti、高Nb的强化元素设计方案，对传统的Ni 3(Al、Ti)强化相进行了改性，形成了含Al更高并同时含有Nb和Sc的Ni 3(Al、Ti、Nb)，比传统的Ni 3(Al、Ti)强化相更耐高温，从而提高合金的抗蠕变性能和持久寿命。三者的具体控制范围为：Al：2.00-3.00％，优选为2.50-3.00％，Ti：1.10-1.49％，Nb：0.81-2.00％。
Sc：对于Sc是否属于稀土元素，在学术界还有争议，对其作用不能简单与稀土元素划等号。虽然稀土在钢铁材料中的添加已被广泛采用，但常用的是La、Ce、Nd等稀土元素，对于Sc在钢铁材料中的应用鲜有报道，学术界对稀土元素的功能认识也比较笼统和泛化，往往将某几种稀土元素的作用泛化为所有稀土元素的作用，存在着以偏概全的问题，一般认为稀土元素的作用是去除夹杂物、净化晶界、提高抗氧化性和氧化膜附着力，对于具体每一种稀土元素的作用研究和认识很少。目前Sc在金属材料中的应用主要集中在铝合金方面，一般认为在铝合金冶炼过程中添加Sc可形成Al 3Sc变质剂，提高凝固形核率，细化铸 态组织，减轻偏析，显著提高合金的强度和韧性。而对于Sc在高温合金中的添加及其作用机理尚未见详细科学报道。本发明研究发现，Sc在本发明实施例的合金中的适量添加主要有三方面作用：1)Sc对钢液的前处理作用，Sc提高了凝固形核率，细化了铸态晶粒，显著改善了铸锭的枝晶偏析，这一方面提高了合金的热加工性能，可改善合金在拉应力方向的开裂问题，防止合金在锻造、热轧等热变形过程中开裂，另一方面使得合金锭不必进行高温长时间的扩散退火或仅需要采用较短的扩散退火时间，降低了能源消耗，可降低生产成本，同时缩短了生产周期，提高了生产效率；2)Sc的添加引入了一种新的强化机制，形成了含Sc的Ni 3(Al、Ti、Nb)复合强化相，比单一的传统强化相Ni 3(Al、Ti)、Ni 3(Al、Ti、Nb)更耐高温、更加稳定，显著提高了合金的抗蠕变性能和持久寿命；3)Sc具有对晶界的后处理作用，即不仅净化钢液，而且在液-固转变过程中和转变完成后，Sc仍能净化并强化晶界，使杂质元素S、P、五害元素和不可避免的低熔点有害元素难以在晶界偏聚，从而防止晶界在高温下形成蠕变空洞，因此能够提高合金蠕变、持久等高温力学性能。Sc的添加是本发明实施例中合金具有抗蠕变和长寿命优点的主要原因之一，Sc添加过少时作用不明显，添加过多时将过量消耗合金中的Al元素，形成大尺寸Al 3Sc，减少细小Ni 3(Al、Ti、Nb)强化相的形成，反而对合金强度和塑性不利，对于本发明实施例中的合金，适宜的Sc含量应控制在0.001-0.1％。
B：B的作用主要表现为两方面，一是由于B的原子半径很小，只有约85皮米，而Ni原子半径约135皮米，因此B原子很容易在晶界富集，使得有害的低熔点元素不能在晶界偏聚，这样就提高了晶界结合力；二是晶界上的硼化物可以阻止晶界滑移、空洞萌生和扩展，对提高合金的抗蠕变性能和持久寿命有利。但过多的B却会恶化合金热加工性能和焊接性能，因此本发明实施例的合金选取适宜的B含量为0.003-0.009％。
Fe：Fe在镍基高温合金中是有害元素，但在工业生产中不可避免，在本发明实施例的合金中，不超过1.50％的Fe是允许的，这就使得在工业生产中可经济地利用含微量Fe的原材料和返回料，从而将合金的生产成本控制在合理水平。
Zr：Zr有助于净化晶界，增强晶界结合力，Zr与B的复合添加有助于保持合金的高温强度和持久寿命，但过量的Zr易引起热加工开裂并损害焊接性能，本发明实施例中的合金将Zr控制在≤0.02％。
Ni：Ni是最重要的基体元素和析出强化相γ′的形成元素，以Ni作为基体，可以大量的固溶具有不同作用的合金元素，如Cr、Mo、Co、C等，Ni基高温合金的组织具有很强的稳定性，具有优异的高温强度、韧性和加工成型性能，适用于制造先进航空发动机和燃气轮机的抗蠕变、长寿命精密热端部件。
根据本发明第一个方面实施例的抗蠕变、长寿命镍基变形高温合金，其中，所述合金中元素Al、Sc和Ti的质量百分含量满足关系式：1.40％≤(Al-1.8Sc)/Ti≤2.6％，优选为2.22％≤(Al-1.8Sc)/Ti≤2.25％。发明人在试验过程中发现，本发明实施例的镍基变形高温合金， 在元素Al、Sc和Ti满足1.40％≤(Al-1.8Sc)/Ti≤2.6％的条件时，能够进一步提高析出相的热稳定性，显著提升了镍基变形高温合金的抗蠕变性能和持久寿命，在816℃、221MPa、100h条件下合金的蠕变塑性伸长率能够降至0.2％以下，在89MPa、927℃条件下的持久寿命可以达到300h以上，特别是当2.22％≤(Al-1.8Sc)/Ti≤2.25％时，在89MPa、927℃条件下，本发明实施例的镍基变形高温合金的持久寿命可以达到330h以上。
本发明第二个方面的实施例提供了抗蠕变、长寿命镍基变形高温合金在航空发动机中的应用。本发明第一方面实施例的抗蠕变、长寿命镍基变形高温合金满足了先进航空发动机设计和使用的要求，能够应用在先进航空发动机的精密设备中。
本发明第三个方面的实施例提供了抗蠕变、长寿命镍基变形高温合金在燃气轮机中的应用。本发明第一方面实施例的抗蠕变、长寿命镍基变形高温合金满足了燃气轮机设计和使用的要求，能够应用在燃气轮机的精密设备中。
本发明第四个方面的实施例提供了抗蠕变、长寿命镍基变形高温合金的制备方法，其特征在于，包括如下步骤：
a、取本发明第一个方面实施例中合金设计用量的原料在真空下熔炼，原料全部融化后进行精炼去除气体，真空下浇铸成合金锭；
b、将所述步骤a得到的合金锭锻造开坯成电极棒，重熔，得到合金锭，当重熔得到的合金锭直径≤200mm时，不进行扩散退火，当重熔得到的合金锭直径＞200mm时，进行扩散退火，扩散退火温度为1150-1200℃，退火时间为12-24小时，锻造开坯成所需坯料，用于进一步加工成锻件、板材、带材、棒材、管材、丝材(包括焊丝)或粉末冶金用材料，之后进行热处理。
根据本发明第四个方面实施例的抗蠕变、长寿命镍基变形高温合金的制备方法，制得的镍基变形高温合金，具有优异的抗蠕变性能、持久寿命和优异的焊接性能，能够满足先进航空发动机和燃气轮机设计和使用的要求，而且本发明实施例的制备方法制备的合金锭可以免去扩散退火处理，或者仅需要采用较短的扩散退火时间即可达到成分均匀化的效果，降低了能源消耗，同时缩短了生产周期，提高了生产效率。
根据本发明第四个方面实施例的抗蠕变、长寿命镍基变形高温合金的制备方法，其中，所述热处理包括一次固溶和两次时效，先进行固溶热处理，经焊接、冷弯等成型加工后再进行两次时效热处理。固溶制度为：固溶温度1100-1170℃，温控精度在±10℃以内，根据产品尺寸确定保温时间，以风冷或水冷方式冷却；所述的一次时效制度为：950-1010℃，温控精度在±10℃以内，保温1-3h，以相当于空冷的冷却速度冷却；所述的二次时效制度为：750-800℃，温控精度在±10℃以内，保温时间8-10h，以相当于空冷的冷却速度冷却。
下面结合实施例详细描述本发明。
实施例1
按设计配比称取精选纯度满足要求的原材料，装入真空感应熔炼炉内，在真空条件下 熔炼；原料全部融化后保持0.1-0.5Pa的真空条件，进行30-35min的精炼以去除气体；精炼结束后，在真空条件下浇铸成合金锭；将合金锭锻造开坯成电极棒，经保护气氛电渣炉重熔后得到直径为230mm的合金锭，对合金锭进行扩散退火，扩散退火温度为1180℃，时间为12h；锻造开坯温度为1100℃，经三火锻造成40mm厚板坯，再一火轧制成20mm厚板材和5mm厚板材；20mm厚板材经1150℃固溶1h，水冷，再进行1010℃×2h空冷+788℃×8h空冷的时效处理，得到抗蠕变、长寿命镍基变形高温合金，检测其力学性能；5mm厚板材经1150℃固溶1h，水冷，利用与板材同成分的焊丝，采用TIG熔化焊方法将两块5mm厚板材拼焊在一起，检测其焊接性能。实施例1制得的合金成分见表1，性能见表2。
实施例2
按设计配比称取精选纯度满足要求的原材料，装入真空感应熔炼炉内，在真空条件下熔炼；原料全部融化后保持0.1-0.5Pa的真空条件，进行30-35min的精炼以去除气体；精炼结束后，在真空条件下浇铸成合金锭；将合金锭锻造开坯成电极棒，经真空自耗炉重熔后得到直径为200mm的合金锭，不进行扩散退火；锻造开坯温度为1050℃，经三火锻造成40mm厚板坯，再一火轧制成20mm厚板材和5mm厚板材；20mm厚板材经1150℃固溶1h，水冷，再进行1010℃×2h空冷+788℃×8h空冷的时效处理，得到抗蠕变、长寿命镍基变形高温合金，检测其力学性能；5mm厚板材经1150℃固溶1h，水冷，利用与板材同成分的焊丝，采用MIG熔化焊方法将两块5mm厚板材拼焊在一起，检测其焊接性能。实施例2制得的合金成分见表1，性能见表2。
实施例3-10与实施例1的制备方法相同，不同在于合金成分不同，实施例3-10制得的合金成分见表1，性能见表2。
对比例1
对比例1与实施例1的制备方法相同，不同之处在合金成分中，元素Al、Ti和Mo的质量百分含量之和即Al+Ti+Mo＝11.49％，对比例1制得的合金成分见表1，性能见表2。
对比例2
对比例2与实施例1的制备方法相同，不同之处在合金成分中，元素Al、Ti和Mo的质量百分含量之和即Al+Ti+Mo＝11.40％，对比例2制得的合金成分见表1，性能见表2。
对比例3
对比例3与实施例1的制备方法相同，不同之处在合金成分中，元素Al、Ti和Mo的质量百分含量之和即Al+Ti+Mo＝13.34％，对比例3制得的合金成分见表1，性能见表2。
对比例4
对比例4与实施例1的制备方法相同，不同之处在合金成分中，元素Al、Ti和Mo的质量百分含量之和即Al+Ti+Mo＝13.25％，对比例4制得的合金成分见表1，性能见表2。
对比例5
对比例5与实施例1的制备方法相同，不同之处在合金成分中，Sc的含量为0.0005％，对比例5制得的合金成分见表1，性能见表2。
对比例6
对比例6与实施例1的制备方法相同，不同之处在合金成分中，Sc的含量为0.19％，对比例6制得的合金成分见表1，性能见表2。
对比例7
对比例7与实施例1的制备方法相同，不同之处在合金成分中，Nb的含量为0.25％，对比例7制得的合金成分见表1，性能见表2。
对比例8
对比例8与实施例1的制备方法相同，不同之处在合金成分中，Nb的含量为2.23％，对比例8制得的合金成分见表1，性能见表2。
对比例9
对比例9与实施例1的制备方法相同，不同之处在合金成分中，Al的含量为1.45％，对比例9制得的合金成分见表1，性能见表2。
对比例10
对比例10与实施例1的制备方法相同，不同之处在合金成分中，Al的含量为3.15％，对比例10制得的合金成分见表1，性能见表2。
对比例11
对比例11与实施例1的制备方法相同，不同之处在合金成分中，Ti的含量为0.95％，对比例11制得的合金成分见表1，性能见表2。
对比例12
对比例12与实施例1的制备方法相同，不同之处在合金成分中，Ti的含量为1.55％，对比例12制得的合金成分见表1，性能见表2。
对比例13
对比例13为商用Haynes282合金，也即合金成分满足申请号201210057737.8的专利中所公开的合金，按照该专利的方法制得的试验例H中的合金成分为：C：0.088％、Cr：19.3％、Co：10.8％、Mo：4.6％、Al：1.63％、Ti：1.85％、B：0.003％、Nb：0.04％、Fe：0.2％，W：6.1％，余量为Ni，对比例13制得的合金性能见表2。
对比例14
对比例14为申请号201910811805.7的专利中所公开的合金，按照该专利中的方法制得的实施例2的合金成分为：C：0.04％、Cr：18.3％、Co：9.5％、Mo：8.5％、Al：1.5％、Ti：1.9％、Zr：0.02％、B：0.005％、Nb：0.2％、V：0.05％，Fe：1.2％，Si：0.1％、Mn：0.2％，P：0.006％、S：0.001％、Nd：0.02％，余量为Ni，对比例14制得的合金性能见表2。
表1
Figure PCTCN2021111913-appb-000001
注：表中各元素的含量均以wt％计。
表2
Figure PCTCN2021111913-appb-000002
注：1、ε p为时效态合金在816℃、221MPa、100h条件下的蠕变塑性伸长率；
2、τ为时效态合金在89MPa、927℃条件下的持久寿命，δ为时效态合金在89MPa、927℃条件下的持久断后伸长率；
3、R p0.2为时效态合金的室温拉伸屈服强度、R m为时效态合金的室温拉伸抗拉强度，A为时效态合金的室温拉伸断后伸长率；
4、焊接裂纹的检测条件为：按国家能源行业标准NB/T 47013.5-2015利用荧光渗透和着色渗透两种方法检测5mm板材焊缝表面质量，按国标GB/T 3323.1-2019利用X射线检测5mm板材焊缝内部质量。
5、表中(Al-1.8Sc)/Ti、Al+Ti+Mo的含量均以wt％计。
通过表1和表2各实施例和对比例的合金成分和性能数据可以看出，实施例1-10中，元素Al、Ti和Mo的质量百分含量满足关系式：11.59≤Al+Ti+Mo≤13.0，在816℃、221MPa、 100h条件下，蠕变塑性伸长率均在0.5％以下，在89MPa、927℃条件下，合金的持久寿命均达到了200h以上，均满足了先进航空发动机和燃气轮机设计和使用的要求。特别是实施例1-6，元素Al、Ti和Mo的质量百分含量满足关系式：11.59≤Al+Ti+Mo≤13.0，并且元素Al、Sc和Ti的质量百分含量满足关系式：1.40≤(Al-1.8Sc)/Ti≤2.6，在满足这两种条件下实施例1-6制得的镍基变形高温合金具有非常优异的抗蠕变性能和高温持久寿命，在816℃、221MPa、100h条件下合金的蠕变塑性伸长率均在0.2％以下，在89MPa、927℃条件下的持久寿命均达到了300h以上。
对比例1-2中，元素Al、Ti和Mo的质量百分含量之和即Al+Ti+Mo分别为11.49％和11.40％，虽然对比例1和2中的合金在89MPa、927℃条件下的持久寿命基本达到了200h以上，但在816℃、221MPa、100h条件下合金的蠕变塑性伸长率均在0.5％以上，并且室温拉伸抗拉强度R m也没有满足要求，均低于期望值1035MPa。
对比例3-4中，元素Al、Ti和Mo的质量百分含量之和即Al+Ti+Mo分别为13.34％和13.25％，虽然对比例3和4中的合金在816℃、221MPa、100h条件下的蠕变塑性伸长率在0.2％以下，在89MPa、927℃条件下的持久寿命基本达到了300h，但在焊接过程中出现了焊接裂纹。对比例3和对比例4的合金中元素Al、Ti和Mo的质量百分含量之和超出13％，析出强化和固溶强化作用过强，时效态屈服强度和抗拉强度富余量偏大，虽然具有优异的抗蠕变性能和持久寿命，但焊接性能较差，出现焊接裂纹。
对比例5和对比例6中，Sc的含量分别为0.0005％和0.19％，对比例5的合金中Sc含量为0.0005％时，由于Sc含量过低，在816℃、221MPa、100h条件下合金的蠕变塑性伸长率为0.878％，在89MPa、927℃条件下的持久寿命仅为103.5小时；对比例6的合金中Sc含量为0.19％，由于Sc含量过高，在816℃、221MPa、100h条件下合金的蠕变塑性伸长率为1.389％，在89MPa、927℃条件下的持久寿命仅为45小时，对比例5和对比例6的抗蠕变和持久寿命均未达到期望值，可见，本发明实施例中Sc的加入量对合金抗蠕变和持久寿命的影响十分重要，本发明实施例的合金中，Sc的含量为0.001-0.1％。
对比例7和8中，Nb的含量分别为0.25％和2.23％，对比例7的合金中采用了低Nb，Nb含量为0.25％，在816℃、221MPa、100h条件下合金的蠕变塑性伸长率为0.808％，在89MPa、927℃条件下的持久寿命为180小时，没有达到期望值；对比例8中，Nb的含量为2.23％，在816℃、221MPa、100h条件下合金的蠕变塑性伸长率为0.126％，在89MPa、927℃条件下的持久寿命为352.4小时，虽然对比例8的合金具有优异的抗蠕变性能和持久寿命，但其在89MPa、927℃条件下的持久断后伸长率仅为5.8％，没有达到≥10.0％的期望值。虽然提高合金中Nb含量能够提高抗蠕变性能和持久寿命，但如果Nb含量过高将会降低断后伸长率。因此，本发明实施例的合金中Nb含量控制为0.81-2.00％。
对比例9和对比例10中，Al的含量分别为1.45％和3.15％，对比例9的合金中采用了低Al，Al含量为1.45％，在816℃、221MPa、100h条件下合金的蠕变塑性伸长率为0.488％， 在89MPa、927℃条件下的持久寿命仅为105小时，对比例9中合金虽然蠕变塑性伸长率达到了期望值，但持久寿命仅为105小时；对比例10中，Al含量为3.15％，在816℃、221MPa、100h条件下合金的蠕变塑性伸长率为0.178％，在89MPa、927℃条件下的持久寿命达到352.3小时，虽然对比例10的合金具有优异的抗蠕变性能和持久寿命，但其在89MPa、927℃条件下的断后伸长率仅为8.6％，没有达到≥10.0％的期望值。虽然提高合金中Al含量能够提高抗蠕变性能和持久寿命，但如果Al含量过高将会降低断后伸长率。因此，本发明实施例的合金中Al含量控制为2.00-3.00％。
对比例11和12中，Ti的含量分别为0.95％和1.55％，对比例11的合金中，Ti含量为0.95％，在816℃、221MPa、100h条件下合金的蠕变塑性伸长率为0.187％，在89MPa、927℃条件下的持久寿命为212.5小时，虽然对比例11的合金抗蠕变性能和持久寿命能够满足要求，但其室温拉伸抗拉强度R m为1032MPa，低于期望值1035MPa，没有满足要求；对比例12中，Ti含量为1.55％，在816℃、221MPa、100h条件下合金的蠕变塑性伸长率为0.14％，在89MPa、927℃条件下的持久寿命为345.6小时，虽然对比例12的合金具有优异的抗蠕变性能和持久寿命，但其在89MPa、927℃条件下的断后伸长率仅为9.5％，没有达到≥10.0％的期望值。本发明实施例的合金中，优选Ti含量为1.10～1.49％。
对比例13为商用282合金，对比例14为申请号201910811805.7所公开的合金，这两个对比例的合金密度均高于实施例，对比例13和对比例14中的合金均没有加入Sc，并且对比例13中采用了低Al、高Ti和低Nb，对比例14中采用了低Al和低Nb，经过测试，对比例13中的合金，在816℃、221MPa、100h条件下合金的蠕变塑性伸长率为1.1％，在89MPa、927℃条件下的持久寿命为102小时，对比例14中的合金，在816℃、221MPa、100h条件下合金的蠕变塑性伸长率为2.201％，在89MPa、927℃条件下的持久寿命为40.3小时，均没有达到期望值，无法满足先进航空发动机和燃气轮机等装备中长期服役的精密热端部件的要求。
在本发明中，术语“一个实施例”、“一些实施例”、“示例”、“具体示例”、或“一些示例”等意指结合该实施例或示例描述的具体特征、结构、材料或者特点包含于本发明的至少一个实施例或示例中。在本说明书中，对上述术语的示意性表述不必须针对的是相同的实施例或示例。而且，描述的具体特征、结构、材料或者特点可以在任一个或多个实施例或示例中以合适的方式结合。此外，在不相互矛盾的情况下，本领域的技术人员可以将本说明书中描述的不同实施例或示例以及不同实施例或示例的特征进行结合和组合。
尽管上面已经示出和描述了本发明的实施例，可以理解的是，上述实施例是示例性的，不能理解为对本发明的限制，本领域的普通技术人员在本发明的范围内可以对上述实施例进行变化、修改、替换和变型。

Claims (10)

  1. 一种抗蠕变、长寿命镍基变形高温合金，其特征在于，包括：C：0.04-0.08％、Cr：18.50-21.50％、Co：9.00-11.00％、Mo：8.00-9.00％、Al：2.00-3.00％、Ti：1.10-1.49％、Nb：0.81-2.00％、B：0.003-0.009％、Sc：0.001-0.10％，余量为镍和不可避免的杂质，以质量百分含量计，该合金中元素Al、Ti和Mo的质量百分含量满足关系式：11.59％≤Al+Ti+Mo≤13.0％。
  2. 根据权利要求1所述的抗蠕变、长寿命镍基变形高温合金，其特征在于，所述杂质为W≤0.50％，Fe≤1.50％，Si≤0.10％，Mn≤0.10％，P≤0.008％，S≤0.008％，Ta≤0.10％，Cu≤0.20％。
  3. 根据权利要求1或2所述的抗蠕变、长寿命镍基变形高温合金，其特征在于，所述合金中元素Al、Sc和Ti的质量百分含量满足关系式：1.40％≤(Al-1.8Sc)/Ti≤2.6％。
  4. 根据权利要求3所述的抗蠕变、长寿命镍基变形高温合金，其特征在于，所述合金中元素Al、Sc和Ti的质量百分含量满足关系式：2.22％≤(Al-1.8Sc)/Ti≤2.25％。
  5. 根据权利要求1-4中任一项所述的抗蠕变、长寿命镍基变形高温合金，其特征在于，所述合金还包括不大于0.02％的Zr。
  6. 根据权利要求1-4中任一项所述的抗蠕变、长寿命镍基变形高温合金，其特征在于，包括：C：0.04-0.08％、Cr：18.50-21.50％、Co：9.00-11.00％、Mo：8.00-9.00％、Al：2.50-3.00％、Ti：1.10-1.49％、Nb：0.81-2.00％、B：0.003-0.009％、Sc：0.001-0.10％，余量为镍和不可避免的杂质，以质量百分含量计，该合金中元素Al、Ti和Mo的质量百分含量满足关系式：11.59％≤Al+Ti+Mo≤13.0％，元素Al、Sc和Ti的质量百分含量满足关系式：2.22％≤(Al-1.8Sc)/Ti≤2.25％。
  7. 权利要求1-6中任一项所述的抗蠕变、长寿命镍基变形高温合金在航空发动机中的应用。
  8. 权利要求1-6中任一项所述的抗蠕变、长寿命镍基变形高温合金在燃气轮机中的应用。
  9. 一种权利要求1-6中任一项所述的抗蠕变、长寿命镍基变形高温合金的制备方法，其特征在于，包括如下步骤：
    a、取设计用量的原料在真空下熔炼，精炼去除气体，真空下浇铸成合金锭；
    b、将所述步骤a得到的合金锭锻造开坯成电极棒，重熔，得到合金锭，锻造开坯成所需坯料，加工后进行热处理。
  10. 根据权利要求9所述的抗蠕变、长寿命镍基变形高温合金的制备方法，其特征在于，所述步骤b中，当重熔得到的合金锭直径≤200mm时，不进行扩散退火，当重熔得到的合金锭直径＞200mm时，进行扩散退火，扩散退火温度为1150-1200℃，退火时间为12-24小时。
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