WO2022233283A1 - 长时稳定性好的高温合金及其制备方法 - Google Patents

长时稳定性好的高温合金及其制备方法 Download PDF

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WO2022233283A1
WO2022233283A1 PCT/CN2022/090690 CN2022090690W WO2022233283A1 WO 2022233283 A1 WO2022233283 A1 WO 2022233283A1 CN 2022090690 W CN2022090690 W CN 2022090690W WO 2022233283 A1 WO2022233283 A1 WO 2022233283A1
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superalloy
crucible
alloy
preparing
long
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PCT/CN2022/090690
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English (en)
French (fr)
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束国刚
高杨
刘伟
安宁
徐超
李振瑞
段方苗
彭劼
余志勇
刘海稳
孙健
安杨
李崇巍
李占青
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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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Publication of WO2022233283A1 publication Critical patent/WO2022233283A1/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
    • C22C19/057—Alloys based on nickel or cobalt based on nickel with chromium and Mo or W with the maximum Cr content being less 10%
    • 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

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  • the application belongs to the technical field of superalloys, and in particular relates to a superalloy with good long-term stability and a preparation method of a superalloy with good long-term stability.
  • the present application aims to solve one of the technical problems in the related art at least to a certain extent.
  • the embodiments of the present application propose a superalloy with good long-term stability and a preparation method thereof.
  • This superalloy with good long-term stability not only has excellent mechanical properties, but also has long-term stability and long-lasting life. It fully meets the requirements for the design and use of advanced aero-engines and gas turbines, and is suitable for medium- and long-term service parts such as turbine blades, hot-end components of aero-engines and gas turbines.
  • the superalloy with good long-term stability includes: C: 0.05-0.16%; Cr: 8.0-9.5%; Co: 9-10.5%; W: 9.0-10.5%; Mo: 0.2-1.0% ; Ta: 2.5-3.5%; Al: 5.0-6.0%; Ti: 0.5-1.5%; B: 0.01-0.02%; Hf: 1.0-2.0%; Zr: 0.004-0.06%; Mg: 0.001-0.005%; Si ⁇ 0.15%; Mn ⁇ 0.05%; the balance is Ni and inevitable impurities, by mass;
  • the superalloys according to the embodiments of the present application have good long-term stability.
  • the strengthening element design scheme of high Al, low Ti and high Ta is adopted, and the traditional Ni 3 (Al, Ti) strengthening phase is modified to form a Ni 3 containing higher Al and Ta at the same time.
  • (Al, Ti, Ta) strengthening phase compared with the traditional Ni 3 (Al, Ti) strengthening phase, the superalloy in the embodiments of the present application has more excellent high temperature resistance performance.
  • the mass percentage content of the elements B and Mg in the alloys of the embodiments of the present application is limited to satisfy the relational formula 0.032% ⁇ B+12.6Mg ⁇ 0.068%, so that the alloys not only have excellent tensile properties and long-lasting life, but also After being treated at 950°C for 3000 hours, no TCP phase is precipitated, and it has excellent long-term service stability. Therefore, the alloys in the examples of the present application can meet the requirements for the design and use of advanced aero-engines and gas turbines.
  • the mass percentage contents of the Al, Ti and Ta satisfy the relational formula 9.1% ⁇ Al+Ti+Ta ⁇ 9.9%.
  • the density of the superalloy does not exceed 8.25 g/cm 3 .
  • the superalloy has a density in the range of 8.11 to 8.25 g/cm 3 .
  • Ni, Cr, Co, W, Mo, Ta, B, Hf, Zr, Mg, Si, Mn and some C raw materials are added to the crucible, and the crucible is heated under vacuum conditions, melting the raw material in the crucible, followed by holding the temperature;
  • an alloy material with good long-term stability can be obtained.
  • the element C in the nickel-based superalloy belongs to the element that is easy to burn out. Therefore, the element C is added in steps and in different proportions in the present application, and part of C is added for the first time. High, the alloying material is faster, and the C element reacts with O in the raw material in the crucible to generate CO2 and other gases that overflow, which is beneficial to remove the gas in the alloy. The second addition of the remaining C can effectively control the C content in the alloy and improve the mechanical properties of the alloy.
  • the method of the embodiment of the present application uses raw materials with specific element ratios, and the obtained alloy not only has excellent tensile properties and long-lasting life, but also does not precipitate TCP phase after the alloy is treated at 950 ° C for 3000 hours, and has excellent long-term service. Stability, able to meet the requirements of advanced aero-engine and gas turbine design and use.
  • part of the C raw material added in the step a is 10-20% of the designed amount of the C raw material.
  • the vacuum condition is a degree of vacuum ⁇ 0.1Pa.
  • the heating time of the crucible is 10-30 min.
  • the holding temperature is 1600°C to 1650°C.
  • the incubation time is 10-30 min.
  • the temperature of the natural cooling is 1200°C to 1400°C.
  • the natural cooling time is 5-15 min.
  • argon gas is passed into the crucible, and the pressure in the crucible is -0.02--0.1 MPa.
  • the vacuum degree in the crucible is less than 0.1 Pa after the vacuuming.
  • the casting temperature is greater than or equal to 1560°C.
  • Figure 1 is the SEM image of the alloy prepared in Example 1 after being treated at 950°C for 3000h;
  • FIG. 2 is the SEM image of the alloy prepared in Comparative Example 1 after being treated at 950° C. for 3000 h.
  • the superalloy according to the embodiment of the present application includes: C: 0.05-0.16%; Cr: 8.0-9.5%; Co: 9-10.5%; W: 9.0-10.5%; Mo: 0.2-1.0%; Ta: 2.5-3.5 %; Al: 5.0-6.0%; Ti: 0.5-1.5%; B: 0.01-0.02%; Hf: 1.0-2.0%; Zr: 0.004-0.06%; Mg: 0.001-0.005%; Si ⁇ 0.15%; Mn ⁇ 0.05%; the balance is Ni and inevitable impurities, by mass;
  • the superalloys according to the embodiments of the present application have good long-term stability.
  • the strengthening element design scheme of high Al, low Ti and high Ta is adopted, and the traditional Ni 3 (Al, Ti) strengthening phase is modified to form a Ni 3 containing higher Al and Ta at the same time.
  • the alloy in the embodiment of the present application defines The mass percentage content of elements B and Mg satisfies the relationship of 0.032% ⁇ B+12.6Mg ⁇ 0.068%, so that the alloy not only has excellent tensile properties and long-lasting life, but also has tensile properties at room temperature Rm ⁇ 966MPa, Rp0.2 ⁇ 725MPa , A ⁇ 6%; high temperature stretching 900°CRm ⁇ 700MPa, Rp0.2 ⁇ 420MPa, A ⁇ 6%, 980°C, lasting time under 200MPa>65h; the alloy does not precipitate TCP phase after 3000h treatment at 950°C, with excellent long-term service stability. It can meet the requirements of advanced aero-engine and gas turbine design and use.
  • 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, propagation and merging of microcracks, thereby improving the high temperature durability of the alloy.
  • the C content is less than 0.05%, 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 Ta in the alloy.
  • Ni 3 (Al, Ti) and Ni 3 (Al, Ti, Ta) composite strengthening phases will be reduced, which will adversely affect the high temperature properties and lasting properties of the alloy, so C should be controlled within 0.16%.
  • Cr 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 Cr content is too high, it is easy to form the TCP phase, which reduces the long-term microstructure and performance stability of the alloy. Therefore, its content generally does not exceed 25%.
  • the Cr content is controlled at 8.0-9.5%.
  • 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 9.0-10.5%.
  • W and Mo are one of the main solid-solution strengthening elements, which can be solid-dissolved in the alloy matrix and in the ⁇ ' strengthening phase, and at the same time can improve the bonding force between atoms, increase the diffusion activation energy and recrystallization temperature , so as to effectively improve the high temperature strength.
  • 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 0.2-1.0%.
  • the atomic radius of W 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. Therefore, the W content is controlled at 9.0-10.5%.
  • Al, Ti and Ta The three are the forming elements of the strengthening phase ⁇ ' in nickel-based alloys. It is generally believed that with the increase of the content of the three, the amount of ⁇ ' increases, and the high-temperature creep and durability improve, but too much ⁇ ' will cause Deterioration of processability.
  • Ti and Ta 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 Ta will form large-grained MC-type carbides. Carbides are detrimental to the mechanical properties of the alloy.
  • 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 Ta is adopted, and the traditional Ni 3 (Al, Ti) strengthening phase is modified, and the alloy with higher Al content and Ta at the same time is formed.
  • Ni 3 (Al, Ti, Ta) which is more resistant to high temperatures than conventional Ni 3 (Al, Ti) strengthening phases, thereby improving the tensile properties and longevity of the alloy.
  • the specific control range of the three is: Al: 5.0-6.0%, Ti: 0.5-1.5%, Ta: 2.5-3.5%.
  • 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. However, 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 application is 0.01-0.02%.
  • 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 reduces the processing performance.
  • the alloys in the examples of this application Control Zr at 0.004-0.06%.
  • Mg Superalloys are microalloyed with Mg, and Mg atoms are segregated at grain boundaries, and this segregation is a balanced segregation.
  • the segregation of Mg at the grain boundary improves the bonding force of the grain boundary and increases the strength of the grain boundary.
  • Mg atoms are not only segregated at grain boundaries, but also at carbide phase boundaries, ⁇ phase boundaries. Mg atoms also enter into ⁇ and carbides, which have a favorable effect on the mechanical properties.
  • the segregation of a small amount of Mg at the grain boundary reduces the grain boundary energy and phase boundary energy, and improves and refines the morphology of other grain boundary precipitation phases at the grain boundary carbide level.
  • carbides can be lumped or spheroidized, effectively inhibiting grain boundary sliding, reducing grain boundary stress concentration, and eliminating notch sensitivity.
  • Mg and sulfur and other harmful impurities form high melting point compounds such as MgS, which purifies the grain boundary, so that the concentration of S, O, P and other impurity elements in the grain boundary is significantly reduced, and the harmful effects of S, O, P and other impurities are reduced.
  • a small amount of Mg can improve durability and plasticity, improve creep performance and high temperature tensile plasticity, increase impact toughness and fatigue strength, and also improve hot workability and yield for some alloys.
  • the content should not be too high, too high will deteriorate the performance, for example, Ni-Ni 2 Mg low melting point (1050°C) eutectic can be formed, which will deteriorate the hot working performance.
  • the Mg content is controlled at 0.001 to 0.005%.
  • the mass percentage content of the Al, Ti and Ta satisfies the relational formula 9.1% ⁇ Al+Ti+Ta ⁇ 9.9%.
  • the strengthening element design scheme of high Al, low Ti and high Ta is adopted to modify the traditional Ni 3 (Al, Ti) strengthening phase to form Ni with higher Al and Ta at the same time.
  • 3 (Al, Ti, Ta) strengthening phase compared with the traditional Ni 3 (Al, Ti) strengthening phase, the superalloy in the embodiments of the present application has better high temperature resistance performance.
  • the mass percentage contents of the elements Al, Ti and Ta are limited to satisfy the relational formula 9.1% ⁇ Al+Ti+Ta ⁇ 9.9%, and the durability of the alloy can be further improved by adjusting Al, Ti and Ta, The alloy has more excellent tensile properties and long-lasting life.
  • the density of the superalloy is less than or equal to 8.25 g/cm 3 .
  • the superalloy has excellent tensile properties and long-lasting life, and because the density does not exceed 8.25g/cm 3 , the alloy has a light weight, which is conducive to reducing the fuel consumption of aero-engine and improving the maneuverability, and can meet the requirements of gas turbines. It is required that the vibration is as small as possible during the working process to prevent the formation of vibration damage.
  • impurities are generally unavoidable.
  • the unavoidable impurities are very small, usually below 1%, such as 0.1%, 0.01% or 0.01% or below, or even undetectable.
  • Ni, Cr, Co, W, Mo, Ta, B, Hf, Zr, Mg, Si, Mn and some C raw materials are added to the crucible, and the crucible is heated under vacuum conditions, melting the raw material in the crucible, followed by holding the temperature;
  • the method for preparing a superalloy according to the embodiment of the present application can prepare an alloy material with good long-term stability.
  • the C element in the nickel-based superalloy is an element that is easy to burn out. Therefore, in this application, C element is added in different proportions, and part of C is added for the first time.
  • the refining temperature is higher, the alloying material is faster, and the C element reacts with O in the raw material in the crucible to generate CO 2 It is beneficial to remove the gas in the alloy when the gas overflows.
  • the second addition of the remaining C can effectively control the C content in the alloy and improve the mechanical properties of the alloy.
  • the method of the embodiment of the present application uses raw materials with specific element ratios, and the obtained alloy not only has excellent tensile properties and long service life, but also does not precipitate TCP phase after the alloy is treated at 950 ° C for 3000 hours, and has excellent long-term service. Stability, able to meet the requirements of advanced aero-engine and gas turbine design and use.
  • Cr comes from at least one of ferrochromium, metallic chromium, bromine ferrochromium, low nitrogen bromine ferrochromium, and high-purity low-oxygen chromium;
  • Co comes from at least one of electrolytic cobalt, metal cobalt, and Jinchuan cobalt plate;
  • W comes from at least one of ferrotungsten, tungsten rod and high-purity tungsten block;
  • Mo comes from at least one of metal molybdenum, ferromolybdenum, and molybdenum-chromium rods;
  • Ta comes from at least one of metal Ta and smelting Ta
  • Al comes from at least one of electrolytic aluminum, aluminum rods, and pure aluminum strips;
  • Ti comes from at least one of pure titanium rods, metal titanium, and sponge titanium;
  • B is from ferroboron
  • Hf comes from at least one of metal hafnium and high-purity hafnium rod;
  • Zr comes from sponge zirconium
  • Ni comes from at least one of electrolytic nickel, metallic nickel, Jinchuan nickel, and high-purity nickel;
  • Si comes from at least one of high-carbon silicon, high-carbon ferrosilicon, high-purity polysilicon, and metal silicon particles;
  • Mn comes from at least one of ferromanganese, metal manganese, and electrolytic manganese flakes.
  • the part of the C raw material added in the step a is 10-20% of the designed amount of the C raw material.
  • the vacuum condition is that the degree of vacuum is less than 0.1Pa.
  • the degree of vacuum in the crucible needs to be controlled. Under higher vacuum degree, vacuuming while heating is conducive to the overflow of gas in the alloy, accelerating the reaction of C and O in the alloy, and improving the purity of the alloy.
  • the crucible heating time in the step a is 15-30 min.
  • rapid heating and rapid smelting of the alloy under a relatively high degree of vacuum can reduce energy consumption on the one hand, and on the other hand, the raw materials in the crucible are rapidly melted to form a liquid-phase molten pool, so as to provide conditions for the reaction of C and O .
  • the holding temperature is 1600°C to 1650°C, and/or the holding time is 10-30 minutes.
  • taking heat preservation measures can make all elements of the alloy mix evenly, and on the other hand, can fully remove the gas in the alloy at this heat preservation temperature.
  • the natural cooling time is 5-15 minutes, and the natural cooling temperature is 1200°C to 1400°C.
  • C, Al and Ti are all elements that are easy to burn out, and are easily reacted with gases to cause deviation of alloy elements. Therefore, it is necessary to reduce the temperature of the alloy liquid in the crucible to 1200°C to 1400°C, which not only ensures that the alloy liquid is liquid, but also avoids the temperature range of maximum burning loss.
  • argon gas is introduced into the crucible, and the pressure in the crucible is -0.02 to -0.1 MPa.
  • the pressure in the crucible is -0.02 to -0.1 MPa.
  • the alloy to be liquid is evacuated after being calm, so that the degree of vacuum in the crucible is less than 0.1Pa.
  • the overflow of gas in the alloy is effectively suppressed, and the splash of the alloy liquid is prevented. Gas extraction improves the purity of the alloy.
  • the temperature of the crucible in the step b is adjusted so that the temperature is ⁇ 1560°C, and the temperature is adjusted to be ⁇ 1560° C.
  • the mold is demolded and the surface is sandblasted and ground to remove the oxide scale to obtain a superalloy.
  • the pouring temperature is lower than 1560° C.
  • the fluidity of the alloy liquid is poor, which is not conducive to the flow of the alloy liquid during the pouring process. This will cause the temperature of the alloy liquid to be too low when it is poured into the mold, and even start to solidify. In this case, the alloy liquid is prone to cracks when it is cooled in the mold, and fractures occur during demolding.
  • the alloy composition ratio is C: 0.11%; Cr: 8.7%; Co: 9.7%; W: 9.8%; Mo: 0.6%; Ta: 3.1%; Al: 5.6%; Ti: 0.8%; B: 0.016%; Hf: 1.5%; Zr: 0.03%; Mg: 0.003%; Si: 0.01%; Mn: 0.01%; the balance is Ni and inevitable impurities.
  • Cr is metal chromium
  • Co electrolytic cobalt
  • W is tungsten rod
  • Mo is metal molybdenum
  • Ta is metal Ta
  • Al is electrolytic aluminum and aluminum rod
  • Ti is pure titanium rod
  • B ferroboron
  • Hf is metal hafnium
  • Zr is sponge zirconium
  • Ni is electrolytic nickel
  • Si is high carbon silicon
  • Mn is metal manganese.
  • the long-term service stability can meet the requirements of advanced aero-engine and gas turbine design and use.
  • the alloy composition ratio is C: 0.16%; Cr: 8.1%; Co: 10.5%; W: 10.2%; Mo: 0.8%; Ta: 2.8%; Al: 5.9%; Ti: 1.2%; B: 0.015%; Hf: 1.8%; Zr: 0.06%; Mg: 0.002%; Si: 0.02%; Mn: 0.01%; the balance is Ni and inevitable impurities.
  • Cr is high-purity low-oxygen chromium
  • Co is Jinchuan cobalt plate
  • W is ferrotungsten and tungsten rod
  • Mo is metal molybdenum and ferromolybdenum
  • Ta is smelting Ta
  • Al is electrolytic aluminum and aluminum rod
  • Ti is pure titanium rod
  • B is ferroboron
  • Hf is metal hafnium and high-purity hafnium rod
  • Zr is sponge zirconium
  • Ni is electrolytic nickel and metal nickel
  • Si is high carbon silicon and high carbon ferrosilicon
  • Mn It is ferromanganese, metal manganese and electrolytic manganese sheet
  • C, Cr, Ni raw materials are mixed into the bottom of the crucible, and Mo and Ta are installed in the upper part of the crucible in proportion.
  • the other ingredients are placed in layers. Pumping low vacuum and heating to complete melting within 30min, then controlling the temperature at 1650°C, and controlling the degree of vacuum to be less than 0.1Pa for 30min, then stop heating, keep it for 15min, and add C, Al, Ti; while adding, argon is flushed into the furnace To -0.1MPa, vacuumize the alloy liquid after it is calm, and the vacuum degree is less than 0.1Pa. Adjust the temperature to tap at 1570°C for casting, cool to room temperature and demould, and then sandblast the surface to remove the oxide scale.
  • the alloy composition ratio is C: 0.08%; Cr: 9.5%; Co: 9%; W: 9.3%; Mo: 0.4%; Ta: 3.4%; Al: 5.1%; Ti: 0.6%; B: 0.019%; Hf: 1.1%; Zr: 0.01%; Mg: 0.002%; Si: 0.009%; Mn: 0.02%; the balance is Ni and inevitable impurities.
  • Cr is ferrochromium, metallic chromium, low nitrogen bromine iron chromium
  • Co is one or more of electrolytic cobalt and Jinchuan cobalt plate
  • W is ferrotungsten, tungsten rod, high One or several kinds of pure tungsten blocks
  • Mo is metal molybdenum, ferromolybdenum, molybdenum-chromium rod
  • Ta is metal Ta and smelting Ta
  • Al is electrolytic aluminum and pure aluminum bar
  • Ti is pure titanium rod, metal titanium and sponge titanium ;
  • B is ferroboron;
  • Hf is metal hafnium and high-purity hafnium rod;
  • Zr is sponge zirconium;
  • Ni is electrolytic nickel, metal nickel, and high-purity nickel;
  • Si is high-carbon ferrosilicon, high-purity polysilicon and metal silicon particles;
  • Mn is manganese Iron, metal manganese and electrolytic manganese sheets; C, Cr, Ni raw materials are mixed into the bottom
  • the other ingredients are placed in layers. Pumping low vacuum and heating to complete melting within 25min, then controlling the temperature at 1640°C, controlling the degree of vacuum to be less than 0.1Pa for 27min, then stopping heating, maintaining for 10min, adding C, Al, and Ti; while adding, argon was flushed into the furnace To -0.07MPa, vacuumize the alloy liquid after it is calm, and the vacuum degree is less than 0.1Pa. Adjust the temperature to tap at 1580°C for casting, cool to room temperature and demould, and then sandblast the surface to remove the oxide scale.
  • Example 4-8 The method of Example 4-8 is the same as that of Example 1, and the difference lies in the alloy composition.
  • the alloy composition of Example 4-8 is shown in Table 1, and the performance data is shown in Table 2.
  • Example 1 Example 2 Example 3 Example 4 Example 5 Example 6
  • Example 8 C 0.11 0.16 0.08 0.06 0.08 0.01 0.012 0.013 Cr 8.7 8.1 9.5 8.2 9.4 8.7 9.2 8.3 Co 9.7 10.5 9 10.2 9.9 10.4 9.7 9.2 W 9.8 10.2 9.3 10.3 9.9 10.2 9.1 Mo 0.6 0.8 0.4 0.3 0.9 0.6 0.5 1 Ta 3.1 2.8 3.4 3.5 3.1 2.7 2.5 3 Al 5.6 5.9 5.1 6 5.1 5.5 5.8 5.2 Ti 0.8 1.2 0.6 1.2 0.6 1.5 0.5 0.9 B 0.016 0.015 0.019 0.015 0.012 0.014 0.014 0.018 hf 1.5 1.8 1.1 1.9 1.2 1.3 1.8 1.5 Zr 0.03 0.06 0.01 0.006 0.03 0.05 0.01 0.06 Mg 0.003 0.002 0.002 0.004 0.004 0.002 0.003 0.003 Si 0.01 0.02 0.009 0.01 0.04 0.02 0.01 0.04 Mn
  • the terms “one embodiment,” “some embodiments,” “example,” “specific example,” or “some examples,” etc. mean the specific features, structures, materials, or characteristics described in connection with the embodiment or example. Features are included in at least one embodiment or example of the present application. 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.

Abstract

本申请公开了一种长时稳定性好的高温合金，所述高温合金包括：C：0.05～0.16％；Cr：8.0～9.5％；Co：9～10.5％；W：9.0～10.5％；Mo：0.2～1.0％；Ta：2.5～3.5％；Al：5.0～6.0％；Ti：0.5～1.5％；B：0.01～0.02％；Hf：1.0～2.0％；Zr：0.004～0.06％；Mg：0.001～0.005％；Si≤0.15％；Mn≤0.05％；余量为Ni和不可避免的杂质，以质量计；其中，B、Mg的质量百分含量满足关系式0.032％≤B+12.6Mg≤0.068％。本申请的长时稳定性好的高温合金不仅具有优良的力学性能，还具有长时稳定性和持久寿命，完全满足先进航空发动机和燃气轮机设计和使用的要求，适用于航空发动机和燃气轮机热端部件透平叶片等中长期服役的零部件。

Description

长时稳定性好的高温合金及其制备方法
相关申请的交叉引用
本申请基于申请号为202110492321.8、申请日为2021年5月6日的中国专利申请提出，并要求该中国专利申请的优先权，该中国专利申请的全部内容在此引入本申请作为参考。
技术领域
本申请属于高温合金技术领域，具体涉及一种长时稳定性好的高温合金和一种长时稳定性好的高温合金的制备方法。
背景技术
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800-1200℃高温长期服役过程中不能发生失效断裂，即要求合金具有优异的高温长时稳定性能，以免在检修周期到来前发生零件失效。目前国内外现有的高温合金中能够完全满足上述要求的合金几乎没有，一般能够达到上述力学性能的合金长时稳定较差，在2000h处理后析出TCP相，导致性能下降。
发明内容
本申请是基于发明人对以下事实和问题的发现和认识做出的：目前在先进航空发动机、燃气轮机等领域使用的高温合金无法满足需求，所需求的高温合金在具体性能指标上，有以下需求：室温下拉伸性能Rm≥966MPa，Rp0.2≥725MPa，A≥6％；高温拉伸900℃Rm≥700MPa，Rp0.2≥420MPa，A≥6％，980℃，200MPa下持久时间＞65h；950℃处理3000h后不能析出TCP相，但是目前没有高温合金可以满足上述需求。
本申请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，本申请实施例提出一种长时稳定性好的高温合金及其制备方法，这种长时稳定性好的高温合金不仅具有优良的力学性能，还具有长时稳定性和持久寿命，完全满足先进航空发动机和燃气轮机设计和使用的要求，适用于航空发动机和燃气轮机热端部件透平叶片等中长期服役的零部件。
根据本申请实施例的长时稳定性好的高温合金包括：C：0.05～0.16％；Cr：8.0～9.5％；Co：9～10.5％；W：9.0～10.5％；Mo：0.2～1.0％；Ta：2.5～3.5％；Al：5.0～6.0％；Ti：0.5～1.5％； B：0.01～0.02％；Hf：1.0～2.0％；Zr：0.004～0.06％；Mg：0.001～0.005％；Si≤0.15％；Mn≤0.05％；余量为Ni和不可避免的杂质，以质量计；
其中，B、Mg的质量百分含量满足关系式0.032％≤B+12.6Mg≤0.068％。
根据本申请实施例的高温合金具有良好的长时稳定性。本申请实施例中采用高Al、低Ti、高Ta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Ta的Ni 3(Al、Ti、Ta)强化相，相比于传统的Ni 3(Al、Ti)强化相，本申请实施例中的高温合金具有更优异的耐高温性能。进一步地，本申请实施例的合金中限定了元素B、Mg的质量百分含量满足关系式0.032％≤B+12.6Mg≤0.068％，使合金不仅具有优异的拉伸性能和持久寿命，合金在950℃处理3000h后没有析出TCP相，具有优良的长时服役稳定性，因此本申请实施例的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。
在一些实施例中，所述Al、Ti和Ta的质量百分含量满足关系式9.1％≤Al+Ti+Ta≤9.9％。
在一些实施例中，所述高温合金的密度不超过8.25g/cm 3。
在一些实施例中，所述高温合金的密度在8.11至8.25g/cm 3的范围内。
根据本申请实施例的高温合金的制备方法包括：
a、根据所需合金配比，将Ni、Cr、Co、W、Mo、Ta、B、Hf、Zr、Mg、Si、Mn和部分C原料加于坩埚内，真空条件下加热所述坩埚，使所述坩埚中的原料熔化，随后保温；
b、停止加热所述坩埚，自然冷却，之后向坩埚中通入氩气，将Al、Ti和剩余C原料加入所述坩埚中，抽真空，使所述坩埚中的原料熔化，浇铸，得到高温合金。
根据本申请实施例的高温合金的制备方法可以获得具有良好的长时稳定性的合金材料。在本申请实施例中，在镍基高温合金中的C元素属于易烧损元素，因此，本申请将C元素分步、按照不同比例加入，第一次加入部分C，在冶炼初期精炼温度较高，合金化料较快，C元素和坩埚内原料中的O进行反应生成CO 2等气体溢出，有利于去除合金中的气体。第二次加入剩余的C，能够有效控制合金中C含量，提高合金的力学性能。此外，本申请实施例的方法，采用特定元素配比的原料，制得的合金不仅具有优异的拉伸性能和持久寿命，合金在950℃处理3000h后没有析出TCP相，具有优良的长时服役稳定性，能够满足先进航空发动机和燃气轮机设计和使用的要求。
在一些实施例中，所述步骤a中加入的部分C原料为C原料设计用量的10-20％。
在一些实施例中，所述步骤a中，所述真空条件为真空度＜0.1Pa。
在一些实施例中，所述步骤a中，所述坩埚的加热时间为10-30min。
在一些实施例中，所述步骤a中，所述保温温度为1600℃～1650℃。
在一些实施例中，所述步骤a中，所述保温时间为10-30min。
在一些实施例中，所述步骤b中，所述自然冷却的温度为1200℃～1400℃。
在一些实施例中，所述步骤b中，所述自然冷却的时间为5-15min。
在一些实施例中，所述步骤b中，向所述坩埚中通入氩气，至所述坩埚中气压为-0.02～-0.1MPa。
在一些实施例中，所述步骤b中，所述抽真空后使所述坩埚内真空度＜0.1Pa。
在一些实施例中，所述步骤b中，所述浇铸温度≥1560℃。
附图说明
图1是实施例1制得的合金经950℃处理3000h后的SEM图；
图2是对比例1制得的合金经950℃处理3000h后的SEM图。
具体实施方式
下面详细描述本申请的实施例，所述实施例的示例在附图中示出。下面通过参考附图描述的实施例是示例性的，旨在用于解释本申请，而不能理解为对本申请的限制。
根据本申请实施例的高温合金包括：C：0.05～0.16％；Cr：8.0～9.5％；Co：9～10.5％；W：9.0～10.5％；Mo：0.2～1.0％；Ta：2.5～3.5％；Al：5.0～6.0％；Ti：0.5～1.5％；B：0.01～0.02％；Hf：1.0～2.0％；Zr：0.004～0.06％；Mg：0.001～0.005％；Si≤0.15％；Mn≤0.05％；余量为Ni和不可避免的杂质，以质量计；
其中，B、Mg的质量百分含量满足关系式0.032％≤B+12.6Mg≤0.068％。
根据本申请实施例的高温合金具有良好的长时稳定性。本申请实施例中采用高Al、低Ti、高Ta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Ta的Ni 3(Al、Ti、Ta)强化相，相比于传统的Ni 3(Al、Ti)强化相，本申请实施例中的高温合金具有更优异的耐高温性能；本申请实施例的合金中限定了元素B、Mg的质量百分含量满足关系式0.032％≤B+12.6Mg≤0.068％，使合金不仅具有优异的拉伸性能和持久寿命，室温下拉伸性能Rm≥966MPa，Rp0.2≥725MPa，A≥6％；高温拉伸900℃Rm≥700MPa，Rp0.2≥420MPa，A≥6％，980℃，200MPa下持久时间＞65h；合金在950℃处理3000h后没有析出TCP相，具有优良的长时服役稳定性。能够满足先进航空发动机和燃气轮机设计和使用的要求。
其中，本申请实施例的合金中各个主要元素的作用如下：
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，能够延缓微裂纹萌生、扩展和合并，从而提高合金的高温持久寿命，当C含量小于0.05％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Ta，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、Ta)复合强化相的Ti和Ta将会减少，对合金的高温性能和持久性能产生不利影响，因此C应控制在不超0.16％。
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TCP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％，本申请实施例中考虑兼顾耐腐蚀性和长期组织性能的稳定性，将Cr含量控制在8.0-9.5％。
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9.0-10.5％。
W和Mo：W和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0.2-1.0％。W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。因此将W含量控制在9.0-10.5％。
Al、Ti和Ta：三者是镍基合金中强化相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′会恶化加工性能。另外，Ti、Ta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、Ta会形成大颗粒MC型碳化物，对合金的力学性能反而不利。本申请通过研究发现，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、Ta总量不变的前提下，通过Al、Ti、Ta比例的优化，可以得到最佳的γ′强化效果。本申请实施例的合金中采用高Al、低Ti、高Ta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Ta的Ni 3(Al、Ti、Ta)，比传统的Ni 3(Al、Ti)强化相更耐高温，从而提高合金的拉伸性能和持久寿命。三者的具体控制范围为：Al：5.0-6.0％，Ti：0.5-1.5％，Ta：2.5-3.5％。
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，因此本申请实施例的合金选取适宜的B含量为0.01-0.02％。
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易降低加工性能，本申请实施例中的合金将Zr控制在0.004-0.06％。
Mg：高温合金用Mg微合金化，Mg原子偏聚于晶界，这种偏聚属平衡偏聚。Mg偏聚于晶界提高晶界结合力，增加晶界强度。Mg原子不仅偏聚于晶界，而且还偏聚于碳化物相界，γ`相界。Mg原子还进入γ`和碳化物中，从而对力学性能产生有利影响。微量Mg在晶界偏聚降低晶界能和相界能，改善和细化晶界碳化物级其他晶界析出相的形态。例如，可以使碳化物块化或球化，有效抑制晶界滑动，降低晶界应力集中，消除缺口敏感性。Mg与硫等有害杂质形成高熔点的化合物MgS等，净化晶界，使晶界的S、O、P等杂质元素的浓度明显降低，减少S、O、P等杂质的有害作用。微量Mg提高持久时间和塑性，改善蠕变性能和高温拉伸塑性，增加冲击韧性和疲劳强度，对有些合金还可改善热加工性能，提高收得率。但是含量不能太高，太高使性能恶化，生成例如可以生成Ni-Ni 2Mg低熔点(1050℃)共晶，使热加工性能变坏。同理，含量太低则不能充分发挥其有利作用。因此将Mg含量控制在0.001～0.005％。
根据本申请实施例的长时稳定性好的高温合金，其中，所述Al、Ti和Ta的质量百分含量满足关系式9.1％≤Al+Ti+Ta≤9.9％。本申请实施例中，采用高Al、低Ti、高Ta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Ta的Ni 3(Al、Ti、Ta)强化相，相比于传统的Ni 3(Al、Ti)强化相，本申请实施例中的高温合金耐高温性能更好。
本申请实施例的合金中限定了元素Al、Ti和Ta的质量百分含量满足关系式 9.1％≤Al+Ti+Ta≤9.9％，通过调控Al、Ti和Ta能够进一步提高合金的持久时长，使合金具有更优异的拉伸性能和持久寿命。
本申请实施例中，所述高温合金的密度≤8.25g/cm 3。其中，所述高温合金在具有优异的拉伸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。
应当理解的是，在合金中，杂质一般是不可避免的。在本申请的Ni高温合金中，不可避免的杂质非常少量，通常在1％以下，例如0.1％、0.01％或0.01％以下，甚至是检测不到的。
根据本申请实施例的长时稳定性好的高温合金的制备方法包括：
a、根据所需合金配比，将Ni、Cr、Co、W、Mo、Ta、B、Hf、Zr、Mg、Si、Mn和部分C原料加于坩埚内，真空条件下加热所述坩埚，使所述坩埚中的原料熔化，随后保温；
b、停止加热所述坩埚，自然冷却，之后向坩埚中通入氩气，将Al、Ti和剩余C原料加入所述坩埚中，抽真空，使所述坩埚中的原料熔化，浇铸，得到高温合金。
根据本申请实施例的的高温合金的制备方法能够制备具有良好长时稳定性的合金材料。在本申请实施例中，在镍基高温合金中的C元素属于易烧损元素。因此本申请将C元素分步，按照不同比例进行加入，第一次加入部分C，在冶炼初期精炼温度较高，合金化料较快，C元素和坩埚内原料中的O进行反应生成CO 2等气体溢出，有利于去除合金中的气体。第二次加入剩余的C，能够有效控制合金中C含量，提高合金的力学性能。此外，本申请实施例的方法，采用特定元素配比的原料，制得的合金不仅具有优异的拉伸性能和持久寿命，合金在950℃处理3000h后没有析出TCP相，具有优良的长时服役稳定性，能够满足先进航空发动机和燃气轮机设计和使用的要求。
其中，本申请实施例中，所述高温合金中各个组分的原料如下：
Cr来自铬铁、金属铬、锦铁铬、低氮锦铁铬、高纯低氧铬中的至少一种；
Co来自电解钴、金属钴、金川钴板中的至少一种；
W来自钨铁、钨棒、高纯钨块中的至少一种；
Mo来自金属钼、钼铁、钼铬棒中的至少一种；
Ta来自金属Ta、熔炼Ta中的至少一种；
Al来自电解铝、铝棒、纯铝条中的至少一种；
Ti来自纯钛棒、金属钛、海绵钛中的至少一种；
B来自硼铁；
Hf来自金属铪、高纯铪棒中的至少一种；
Zr来自海绵锆；
Ni来自电解镍、金属镍、金川镍、高纯镍中的至少一种；
Si来自高碳硅、高碳硅铁、高纯多晶硅、金属硅颗粒中的至少一种；
Mn来自锰铁、金属锰、电解锰片中的至少一种。
根据本申请实施例的长时稳定性好的高温合金的制备方法，其中，所述步骤a中加入的部分C原料为C原料设计用量的10-20％。根据本申请实施例的长时稳定性好的高温合金的制备方法，其中，所述步骤a中，所述真空条件为真空度＜0.1Pa。在本申请实施例中，合金化料时，需要控制坩埚内的真空度。在较高真空度下，一边加热一边抽真空，有利于合金中气体的溢出，加速合金中C和O的反应，提高合金纯净度。
本申请实施例中，所述步骤a中坩埚加热时间为15-30min。在本申请实施例中，合金在较高真空度下，急速升温快速熔炼一方面能够降低能耗，另一方面，坩埚中原料快速熔化形成液相熔池，以提供C和O发生反应的条件。
根据本申请实施例的长时稳定性好的高温合金的制备方法，其中，所述步骤a中，所述保温温度为1600℃～1650℃，和/或，所述保温时间为10-30min。在本申请实施例中，在高温合金精炼期，采取保温措施一方面能够使合金各元素混合均匀，另一方面在此保温温度下能够充分去除合金中的气体。
根据本申请实施例的长时稳定性好的高温合金的制备方法，其中，所述步骤b中，所述自然冷却的时间为5-15min，所述自然冷却的温度为1200℃～1400℃。在本申请实施例中，C、Al和Ti元素都属于易烧损元素，极易和气体反应导致合金元素偏差。因此需要将坩埚内合金液温度降至1200℃～1400℃，既保证了合金液为液态，也可以避开最大烧损的温度区间。
根据本申请实施例的长时稳定性好的高温合金的制备方法，其中，所述步骤b中，向所述坩埚中通入氩气，至所述坩埚中气压为-0.02～-0.1MPa。在本申请实施例中，在加入C、Al、Ti原料后，合金液中的O会急剧与这些元素发生反应，即使对温度进行控制，也无法完全抑制其反应，而且在反应的过程中，液态合金液液面会随着反应而发生剧烈的翻涌，本申请实施例中，向坩埚中通入氩气能够有效抑制反应和防止合金液喷溅。
根据本申请实施例的长时稳定性好的高温合金的制备方法，其中，所述步骤b中，所述待液态合金平静后抽真空，使坩埚中真空度＜0.1Pa。在本申请实施例中，通入氩气后，有效抑制了合金中气体的溢出，防止合金液喷溅，在加入C、Al、Ti原料合金液平静后，采用 抽真空处理将合金液中的气体抽出，提高了合金的纯净度。
根据本申请实施例的长时稳定性好的高温合金的制备方法，其中，所述步骤b中，所述调整所述步骤b中坩埚的温度，调整使温度≥1560℃，浇铸进入模具中，待其自然冷却至室温后，脱模进行表面喷砂修磨，去除氧化皮，得到高温合金。在本申请实施例中，若倒出温度低于1560℃，合金液流动性差，在倒出过程中不利于合金液的流动。这将导致合金液倒入模具时的温度过低，甚至开始凝固，这种情况下合金液在模具中冷却时极易产生裂纹，脱模时发生断裂。
下面结合实施例详细描述本申请。
实施例1
合金成分配比为C：0.11％；Cr：8.7％；Co：9.7％；W：9.8％；Mo：0.6％；Ta：3.1％；Al：5.6％；Ti：0.8％；B：0.016％；Hf：1.5％；Zr：0.03％；Mg：0.003％；Si：0.01％；Mn：0.01％；余量为Ni和不可避免的杂质。按上述合金的化学成份质量百分比配料，其中，Cr为金属铬；Co为电解钴；W为钨棒；Mo为金属钼；Ta为金属Ta；Al为电解铝、铝棒；Ti为纯钛棒；B为硼铁；Hf为金属铪；Zr为海绵锆；Ni为电解镍；Si为高碳硅；Mn为金属锰。
将C、Cr、Ni原料混合放入坩埚底部，Mo、Ta按比例装在坩埚中的上部。其他原料分层放置。抽低真空同时15min内加热至完全熔化，随后控制温度在1600℃，真空度控制在小于0.1Pa保持10min后停止加热，保持5min，加入C、Al、Ti；加入的同时向炉内冲氩气至-0.02MPa，合金液平静后抽真空，真空度小于0.1Pa。调整温度在1560℃的条件下进行浇铸，冷却至室温脱模进行表面喷砂修磨去除氧化皮。
如图1所示，为实施例1制得的合金经950℃处理3000h后的SEM图，从图1可以看出，合金经过3000h处理后，组织中强化相γ`相界清晰，γ`相均匀，没有发生聚集重大。同时也没有析出TCP有害相。
按照本实施例的方法制备的高温合金B+12.6Mg＝0.0538，Al+Ti+Ta＝9.5％，合金密度为8.17g/cm 3，具有优异的拉伸性能和持久寿命，室温下拉伸性能Rm 1025MPa，Rp0.2 849MPa，A 6.5％；高温拉伸900℃Rm 802MPa，Rp0.2 455MPa，A 6.5％，980℃，200MPa下持久时间110h；950℃处理3000h后没有析出TCP相，具有优良的长时服役稳定性，能够满足先进航空发动机和燃气轮机设计和使用的要求。
实施例2
合金成分配比为C：0.16％；Cr：8.1％；Co：10.5％；W：10.2％；Mo：0.8％；Ta：2.8％； Al：5.9％；Ti：1.2％；B：0.015％；Hf：1.8％；Zr：0.06％；Mg：0.002％；Si：0.02％；Mn：0.01％；余量为Ni和不可避免的杂质。按上述合金的化学成份质量百分比配料，其中，Cr高纯低氧铬；Co为金川钴板；W为钨铁和钨棒；Mo为金属钼和钼铁；Ta为熔炼Ta；Al为电解铝和铝棒；Ti为纯钛棒；B为硼铁；Hf为金属铪和高纯铪棒；Zr为海绵锆；Ni为电解镍和金属镍；Si为高碳硅和高碳硅铁；Mn为锰铁、金属锰和电解锰片；将C、Cr、Ni原料混合放入坩埚底部，Mo、Ta按比例装在坩埚中的上部。其他原料分层放置。抽低真空同时30min内加热至完全熔化，随后控制温度在1650℃，真空度控制在小于0.1Pa保持30min后停止加热，保持15min，加入C、Al、Ti；加入的同时向炉内冲氩气至-0.1MPa，合金液平静后抽真空，真空度小于0.1Pa。调整温度在1570℃的条件下出钢进行浇铸，冷却至室温脱模进行表面喷砂修磨去除氧化皮。
按照本实施例制备的高温合金B+12.6Mg＝0.0402％，Al+Ti+Ta＝9.9％，合金密度为8.11g/cm 3，使合金具有优异的拉伸性能和持久寿命，室温下拉伸性能Rm 995MPa，Rp0.2 839MPa，A 7.5％；高温拉伸900℃Rm 810MPa，Rp0.2 455MPa，A 7.0％，980℃，200MPa下持久时间102h；950℃处理3000h后没有析出TCP相，具有优良的长时服役稳定性。能够满足先进航空发动机和燃气轮机设计和使用的要求。
实施例3
合金成分配比为C：0.08％；Cr：9.5％；Co：9％；W：9.3％；Mo：0.4％；Ta：3.4％；Al：5.1％；Ti：0.6％；B：0.019％；Hf：1.1％；Zr：0.01％；Mg：0.002％；Si：0.009％；Mn：0.02％；余量为Ni和不可避免的杂质。按上述合金的化学成份质量百分比配料，其中，Cr为铬铁、金属铬、低氮锦铁铬；Co为电解钴和金川钴板的一种或几种；W为钨铁、钨棒、高纯钨块的一种或几种；Mo为金属钼、钼铁、钼铬棒；Ta为金属Ta和熔炼Ta；Al为电解铝和纯铝条；Ti为纯钛棒、金属钛和海绵钛；B为硼铁；Hf为金属铪和高纯铪棒；Zr为海绵锆；Ni为电解镍、金属镍、和高纯镍；Si为高碳硅铁、高纯多晶硅和金属硅颗粒；Mn为锰铁、金属锰和电解锰片；C、Cr、Ni原料混合放入坩埚底部，Mo、Ta按比例装在坩埚中的上部。其他原料分层放置。抽低真空同时25min内加热至完全熔化，随后控制温度在1640℃，真空度控制在小于0.1Pa保持27min后停止加热，保持10min，加入C、Al、Ti；加入的同时向炉内冲氩气至-0.07MPa，合金液平静后抽真空，真空度小于0.1Pa。调整温度在1580℃的条件下出钢进行浇铸，冷却至室温脱模进行表面喷砂修磨去除氧化皮。
按照本实施例制备的高温合金B+12.6Mg＝0.0442％，Al+Ti+Ta＝9.1％，合金密度为8.25g/cm 3，使合金具有优异的拉伸性能和持久寿命，室温下拉伸性能Rm 975MPa，Rp0.2  739MPa，A 7％；高温拉伸900℃Rm 730MPa，Rp0.2 435MPa，A 6.5％，980℃，200MPa下持久时间72h；950℃处理3000h后没有析出TCP相，具有优良的长时服役稳定性。能够满足先进航空发动机和燃气轮机设计和使用的要求。
实施例4-8
实施例4-8的方法与实施例1相同，不同之处在于合金成分，实施例4-8的合金成分见表1，性能数据见表2。
表1
  实施例1 实施例2 实施例3 实施例4 实施例5 实施例6 实施例7 实施例8
C 0.11 0.16 0.08 0.06 0.08 0.01 0.012 0.013
Cr 8.7 8.1 9.5 8.2 9.4 8.7 9.2 8.3
Co 9.7 10.5 9 10.2 9.9 10.4 9.7 9.2
W 9.8 10.2 9.3 10.3 9.3 9.9 10.2 9.1
Mo 0.6 0.8 0.4 0.3 0.9 0.6 0.5 1
Ta 3.1 2.8 3.4 3.5 3.1 2.7 2.5 3
Al 5.6 5.9 5.1 6 5.1 5.5 5.8 5.2
Ti 0.8 1.2 0.6 1.2 0.6 1.5 0.5 0.9
B 0.016 0.015 0.019 0.015 0.012 0.014 0.014 0.018
Hf 1.5 1.8 1.1 1.9 1.2 1.3 1.8 1.5
Zr 0.03 0.06 0.01 0.006 0.03 0.05 0.01 0.06
Mg 0.003 0.002 0.002 0.004 0.004 0.002 0.003 0.003
Si 0.01 0.02 0.009 0.01 0.04 0.02 0.01 0.04
Mn 0.01 0.01 0.02 0.04 0.02 0.03 0.02 0.01
B+12.6Mg 0.0538 0.0402 0.0442 0.0654 0.0624 0.0392 0.0518 0.0558
Al+Ti+Ta 9.5 9.9 9.1 10.7 8.8 9.7 8.8 9.1
表2
Figure PCTCN2022090690-appb-000001
对比例1
与实施例1相同，其不同之处在于，合金中B、Mg的百分含量不同，其中B：0.01％；Mg：0.001％，B+12.6Mg＝0.0226％。
对比例1制得的合金性能数据见表3
如图2所示，为对比例1制得的合金经950℃处理3000h后的SEM图，图2可以看出，经过3000h处理后，合金组织中析出长条形的TCP相，数量密集。合金受力后微裂纹容易在此处进行生长并扩散，并且最终导致了合金失效。
对比例2
与实施例1相同，其不同之处在于，合金中B、Mg的百分含量不同，其中B：0.02％；Mg：0.005％，B+12.6Mg＝0.083％。
对比例2制得的合金性能数据见表3
对比例3
与实施例1相同，其不同之处在于，合金中不含有Ta。
对比例3制得的合金性能数据见表3。
表3
Figure PCTCN2022090690-appb-000002
在本申请中，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申请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申请的实施例，可以理解的是，上述实施例是示例性的，不能理解为对本申请的限制，本领域的普通技术人员在本申请的范围内可以对上述实施例进行变化、修改、替换和变型。

Claims (15)

  1. 一种高温合金，包括：C：0.05～0.16％；Cr：8.0～9.5％；Co：9～10.5％；W：9.0～10.5％；Mo：0.2～1.0％；Ta：2.5～3.5％；Al：5.0～6.0％；Ti：0.5～1.5％；B：0.01～0.02％；Hf：1.0～2.0％；Zr：0.004～0.06％；Mg：0.001～0.005％；Si≤0.15％；Mn≤0.05％；余量为Ni和不可避免的杂质，以质量计；
    其中，B、Mg的质量百分含量满足关系式0.032％≤B+12.6Mg≤0.068％。
  2. 根据权利要求1所述的高温合金，其中所述Al、Ti和Ta的质量百分含量满足关系式9.1％≤Al+Ti+Ta≤9.9％。
  3. 根据权利要求1所述的高温合金，其中所述高温合金的密度不超过8.25g/cm 3。
  4. 根据权利要求3所述的高温合金，其中所述高温合金的密度在8.11至8.25g/cm 3的范围内。
  5. 一种权利要求1至4中任一项所述的高温合金的制备方法，包括：
    a、根据合金设计配比，将Ni、Cr、Co、W、Mo、Ta、B、Hf、Zr、Mg、Si、Mn和部分C原料加于坩埚内，真空条件下加热所述坩埚，使所述坩埚中的原料熔化，保温；
    b、停止加热所述坩埚，自然冷却，之后向坩埚中通入氩气，将Al、Ti和剩余C原料加入所述坩埚中，抽真空，使所述坩埚中的原料熔化，浇铸，得到高温合金。
  6. 根据权利要求5所述的高温合金的制备方法，其中所述步骤a中加入的部分C原料为C原料设计用量的10-20％。
  7. 根据权利要求5或6所述的高温合金的制备方法，其中所述步骤a中，所述真空条件为真空度＜0.1Pa。
  8. 根据权利要求5至7中任一项所述的高温合金的制备方法，其中所述步骤a中，所述坩埚的加热时间为10-30min。
  9. 根据权利要求5至8中任一项所述的高温合金的制备方法，其中所述步骤a中，所述保温温度为1600℃～1650℃。
  10. 根据权利要求5至9中任一项所述的高温合金的制备方法，其中所述步骤a中，所述保温时间为10-30min。
  11. 根据权利要求5至10中任一项所述的高温合金的制备方法，其中所述步骤b中，所述自然冷却的温度为1200℃～1400℃。
  12. 根据权利要求5至11中任一项所述的高温合金的制备方法，其中所述步骤b中，所述自然冷却的时间为5-15min。
  13. 根据权利要求5至12中任一项所述的高温合金的制备方法，其中所述步骤b中，向所述坩埚中通入氩气，至所述坩埚中气压为-0.02～-0.1MPa。
  14. 根据权利要求5至13中任一项所述的高温合金的制备方法，其中所述步骤b中，所述抽真空后使所述坩埚内真空度＜0.1Pa。
  15. 根据权利要求5至14中任一项所述的高温合金的制备方法，其中所述步骤b中，所述浇铸温度≥1560℃。
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