A Fe-based superalloy, its preparation method and application

Abstract

A Fe-based superalloy, its preparation method and application belong to the field of alloy technology, overcoming the defects of low hot strength and poor thermal stability of existing Fe-based superalloys and heat-resistant steels. The Fe-based superalloy of the present invention, in terms of mass percentage, includes: C: 0.05 - 0.08%, Ni: 23 - 25%, Cr: 13.5 - 14.8%, W: 0.1 - 0.3%, Mo: 0.6 - 1.0%, Ti: 2.0 - 2.5%, Al: 1.0 - 1.5%, Co: 1.0 - 1.6%, Mn: ≤0.5%, Si: ≤0.3%, B: ≤0.005%, Zr: ≤0.03%, and the balance is Fe. Through the optimization of the alloy composition, the present invention significantly reduces the alloy cost while maintaining the high-temperature strength of the alloy, and has good organizational structure and thermal stability of strength.

Description

Technical Field

[0001] The present invention belongs to the technical field of alloys, and particularly relates to a ferrous-based superalloy, a preparation method thereof, and an application thereof. Background Art

[0002] The high-efficiency, clean and high-parameter ultra-supercritical coal-fired power generation technology is one of the main development trends of clean coal-fired power generation technology in the world today. To achieve steam parameters of 650 °C, a series of technical problems need to be solved. Among them, developing a heat-resistant alloy for the last-stage superheater / reheater of boilers with high thermal strength, resistance to high-temperature flue gas oxidation corrosion and high-temperature steam-water medium corrosion, good weldability and processability, and low price is the key to realizing high-parameter ultra-supercritical coal-fired power generation technology.

[0003] At present, a series of solid-solution strengthened and precipitation-strengthened nickel-based and nickel-iron-based superalloy materials have been developed at home and abroad, such as Inconel 740H, Haynes 282, CCA 617, Nimonic 263, USC41. These materials have excellent high-temperature creep strength and oxidation resistance, but their high prices, poor weldability, and high technical requirements for smelting and hot working limit their rapid popularization and application. As a cheap element, Fe can significantly reduce the cost of the alloy and improve its processability. Therefore, a series of ferrous-based heat-resistant alloys have also been developed at home and abroad, such as Inconel 718, GH2984, HR6W, HR35, Sanicro 25, Super304H, HR3C, G115, etc. Compared with nickel-based and nickel-iron-based superalloys, although the above-mentioned ferrous-based superalloys and heat-resistant steels have the advantage of raw material cost, they have low thermal strength and poor tissue stability, and are difficult to meet the service performance requirements. Summary of the Invention

[0004] Therefore, the technical problem to be solved by the present invention is to overcome the defects of low thermal strength and poor thermal stability of ferrous-based superalloys and heat-resistant steels in the prior art, and thus provide a ferrous-based superalloy, a preparation method thereof, and an application thereof.

[0005] For this purpose, the present invention provides the following technical solutions.

[0006] A ferrous-based superalloy, the components are calculated by mass percentage, including: C: 0.05 - 0.08%, Ni: 23 - 25%, Cr: 13.5 - 14.8%, W: 0.1 - 0.3%, Mo: 0.6 - 1.0%, Ti: 2.0 - 2.5%, Al: 1.0 - 1.5%, Co: 1.0 - 1.6%, Mn: ≤0.5%, Si: ≤0.3%, B: ≤0.005%, Zr: ≤0.03%, and the balance is Fe.

[0007] A preparation method of a ferrous-based superalloy, comprising the following steps:

[0008] Step 1: Melting the raw materials into a master alloy ingot under a vacuum environment;

[0009] Step 2: Performing homogenization annealing on the master alloy ingot;

[0010] Step 3: High-temperature rolling;

[0011] Step 4: Heat treatment.

[0012] Furthermore, Step 4 includes solution treatment and aging treatment;

[0013] The aging treatment includes first high-temperature aging, low-temperature aging, and second high-temperature aging.

[0014] Furthermore, the first high-temperature aging includes: heating the heat treatment furnace to 130 - 170 °C above the temperature at which the γ′ phase starts to precipitate, placing the alloy after solution treatment into the heat treatment furnace for 1 - 2 h, and then water-cooling to room temperature.

[0015] Furthermore, the low-temperature aging includes: heating the heat treatment furnace to 150 - 200 °C below the temperature at which the γ′ phase starts to precipitate, placing the alloy after the first high-temperature aging into the heat treatment furnace for 8 - 16 h, and then water-cooling to room temperature.

[0016] Furthermore, the second high-temperature aging includes: heating the heat treatment furnace to 30 - 150 °C below the temperature at which the γ′ phase starts to precipitate, placing the alloy after low-temperature aging into the heat treatment furnace for 4 - 24 h, and then water-cooling to room temperature.

[0017] Furthermore, Step 2 includes: heating the master alloy ingot from room temperature to 300 - 350 °C above the temperature at which the γ′ phase starts to precipitate at a heating rate of 10 - 30 °C / min, performing homogenization annealing for 20 - 35 h, and then air-cooling to room temperature.

[0018] Furthermore, Step 3 includes: performing high-temperature rolling on the master alloy ingot after homogenization annealing at a temperature above 200 - 250 °C at which the γ′ phase starts to precipitate, with a deformation amount of 15 - 25% per pass, and a final total deformation amount of not less than 60%.

[0019] Furthermore, the solution treatment includes: heating the heat treatment furnace to 250 - 300 °C above the temperature at which the γ′ phase starts to precipitate, placing the alloy after high-temperature rolling into the heat treatment furnace for solution treatment for 0.5 - 2 h, and then water-cooling to room temperature.

[0020] Application of the prepared alloy in thermal power units.

[0021] The technical solution of the present invention has the following advantages:

[0022] 1. The iron-based superalloy provided by the present invention, in terms of mass percentage, comprises: C: 0.05 - 0.08%, Ni: 23 - 25%, Cr: 13.5 - 14.8%, W: 0.1 - 0.3%, Mo: 0.6 - 1.0%, Ti: 2.0 - 2.5%, Al: 1.0 - 1.5%, Co: 1.0 - 1.6%, Mn: ≤0.5%, Si: ≤0.3%, B: ≤0.005%, Zr: ≤0.03%, and the balance is Fe.

[0023] The functions of each element in the alloy and the reasons for the selection of their content ranges are as follows:

[0024] Ni is an austenite stabilizing element. The addition of Ni can improve the stability of Fe-based austenite under low-temperature conditions. At the same time, Ni can also form an ordered phase γ' phase with elements such as Al and Ti, improving the strength of the alloy.

[0025] Fe is inexpensive. A large amount of addition can reduce the cost of the alloy and improve the workability of the alloy. A large amount of addition of Fe element can reduce the temperature at which the γ' phase starts to precipitate, thereby reducing the coarsening rate of the γ' phase during the aging process.

[0026] The main function of Cr is to improve the oxidation resistance, steam oxidation resistance and corrosion resistance of the alloy, especially the sulfur corrosion resistance. With the increase of Cr content, the oxidation resistance, steam oxidation resistance and corrosion resistance are improved. However, a large amount of Fe is added in the present invention, which will hinder the precipitation of the γ' phase. The reduction of the volume fraction of the γ' phase leads to a decrease in strength. At the same time, the interaction of a large amount of Fe with Cr, Co, and Mo usually leads to the massive precipitation of acicular or lamellar α-Cr phase, σ phase and η phase. In the present invention, by controlling the Cr content within a lower range (13.5 - 14.8%), the volume fraction of the γ' phase rebounds, avoiding a large reduction in strength and avoiding the precipitation of acicular or lamellar α-Cr phase, σ phase and η phase.

[0027] Both Ti and Al are γ' phase forming elements. The precipitated γ' phase can significantly improve the high-temperature strength of the alloy. In addition, Al has the ability to improve the resistance of the alloy to internal oxidation, and Ti has the function of resisting hot corrosion and sulfidation corrosion.

[0028] Mo, W and Co are important solid solution strengthening elements, which have a great promoting effect on the high-temperature strength and creep properties of the alloy.

[0029] Si is an important deoxidizing element. Appropriate addition can promote the oxidation resistance of the alloy.

[0030] Mn can replace Ni to form and stabilize austenite and fix the S in the alloy.

[0031] C, B, and Zr are grain boundary segregation elements. Adding a small amount can increase the binding force of the grain boundary, strengthen and purify the grain boundary, and change the morphology of the grain boundary precipitate phase, thereby improving the strength of the alloy. In addition, C and B can form carbides and borides with elements such as Ti, Mo, and Fe in the alloy. The precipitation of a small amount can pin the grain boundary and improve the creep properties of the alloy.

[0032] The alloy of the present invention has a relatively high Fe element content (about 54 - 57% Fe), and relatively low Ni, Cr, W, and Mo contents, significantly reducing the raw material cost of the alloy.

[0033] Through the composition optimization of the alloy, the present invention significantly reduces the alloy cost while maintaining the high-temperature strength of the alloy, and has good organizational structure and thermal stability of strength. No acicular phase precipitation is found after the alloy is thermally exposed, and the growth rate of the main strengthening phase γ' is very slow, indicating good thermal stability of the alloy.

[0034] 2. The preparation method of the iron-based superalloy provided by the present invention includes the following steps: Step 1, melting the raw materials into a master alloy ingot in a vacuum environment; Step 2, performing homogenization annealing on the master alloy ingot; Step 3, hot rolling; Step 4, solution treatment; Step 5, aging treatment.

[0035] Specifically, the melting of the present invention uses arc melting, abandoning the traditional triple melting process of superalloys, thereby reducing the preparation cost of the alloy.

[0036] 3. The preparation method of the iron-based superalloy provided by the present invention, the Step 3 includes: hot rolling the master alloy ingot after homogenization annealing at 200 - 250°C above the γ' phase precipitation temperature, with the deformation amount per pass being 15 - 25%, and the final total deformation amount not less than 60%. The present invention processes the alloy with large deformation amounts in multiple passes at a specific temperature, ensuring sufficient strain energy storage during the alloy processing, which is convenient for the occurrence of the subsequent recrystallization process.

[0037] 4. The preparation method of the iron-based superalloy provided by the present invention, the Step 4 includes solution treatment and aging treatment; the aging treatment includes the first high-temperature aging, low-temperature aging, and the second high-temperature aging.

[0038] The first high-temperature aging controls the precipitation of carbide at grain boundaries. The low-temperature aging promotes the nucleation of the main strengthening phase γ'-phase particles. The second high-temperature aging promotes the precipitation and growth of γ'-phase particles, ensuring that the alloy has excellent tensile strength. After heat treatment, the alloy has excellent high-temperature strength performance. Its tensile strength at 700 °C is not less than 680 MPa, and the elongation is higher than 25%. Moreover, the alloy has excellent microstructure stability. After 10,000 h of thermal exposure at 700 °C, the average size of the main strengthening phase γ'-phase particles is 45 - 55 nm; no acicular and lamellar phases are found to precipitate in the alloy; the alloy still has a high compressive yield strength after long-term thermal exposure, indicating that the alloy of the present invention has high thermal stability. Description of the Drawings

[0039] In order to more clearly illustrate the specific embodiments of the present invention or the technical solutions in the prior art, the following will briefly introduce the drawings required for use in the description of the specific embodiments or the prior art. Obviously, the following drawings are some embodiments of the present invention. For those of ordinary skill in the art, without creative efforts, other drawings can also be obtained based on these drawings.

[0040] Figure 1 It is a low-magnification image of the original aged microstructure of the alloy obtained in Example 1;

[0041] Figure 2 It is a high-magnification image of the original aged microstructure of the alloy obtained in Example 1;

[0042] Figure 3 It is the microstructure of the alloy obtained in Example 1 after thermal exposure aging at 700 °C for 1000 h;

[0043] Figure 4 It is the microstructure of the alloy obtained in Example 1 after thermal exposure at 700 °C for 5000 h;

[0044] Figure 5 It is the microstructure of the alloy obtained in Example 1 after thermal exposure at 700 °C for 10000 h;

[0045] Figure 6 It is a comparison diagram of the growth of γ'-phase at different thermal exposure times at 700 °C for the alloy of Example 1 and the samples of Comparative Examples 1 - 3;

[0046] Figure 7 It is a high-magnification image of the original aged microstructure of the alloy of Comparative Example 3;

[0047] Figure 8 It is the microstructure of the alloy of Comparative Example 3 after thermal exposure at 700 °C for 5000 h;

[0048] Figure 9 This is a high-magnification image of the microstructure of the alloy in the original aged state of Comparative Example 4.

[0049] Figure 10 This is the microstructure of the alloy in Comparative Example 4 after thermal exposure at 700 °C for 1000 h. Specific Embodiments

[0050] The following embodiments are provided to better understand the present invention further. They are not limited to the described optimal embodiments, and do not limit the content and protection scope of the present invention. Any product that is the same as or similar to the present invention obtained by anyone under the inspiration of the present invention or by combining the features of the present invention with other prior art features falls within the protection scope of the present invention.

[0051] For those embodiments where specific experimental steps or conditions are not indicated, the operations or conditions of the conventional experimental steps described in the literature in this field can be followed. For reagents or instruments whose manufacturers are not indicated, they are all conventional reagent products that can be obtained through commercial purchase.

[0052] Example 1

[0053] This example provides a ferrous superalloy, which includes, by mass percentage: C: 0.07%, Ni: 24%, Cr: 14.8%, Co: 1.0%, W: 0.2%, Mo: 0.8%, Ti: 2.1%, Al: 1.3%, Si: 0.015%, B: 0.003%, and the balance is Fe.

[0054] The preparation method of the ferrous superalloy in this example includes the following steps:

[0055] Step 1: After considering the burn-off, complete the batching. Introduce argon as the protective gas into the vacuum arc melting furnace, and obtain a 150 g master alloy ingot with uniform alloy composition through electromagnetic stirring and flipping.

[0056] Step 2: Heat the master alloy ingot in the furnace at a rate of 10 °C / min to 1160 °C for homogenization treatment for 24 h, and then air-cool to room temperature.

[0057] Step 3: Perform hot rolling on the homogenized master alloy ingot at a temperature 230 °C above the γ′ phase precipitation temperature, with a deformation amount of 20% per pass and a final total deformation amount of 60%. After each pass of rolling, heat the ingot in the furnace for heat preservation for 30 min.

[0058] Step 4: After the alloy rolling is completed, heat it up to 1100 °C for solution treatment for 0.5 h, and then water-cool it to room temperature. Subsequently, directly put the alloy into a heat treatment furnace at 1000 °C for high-temperature aging for 1 h, and then air-cool it to room temperature. Next, directly put the alloy into a heat treatment furnace at 650 °C for low-temperature aging for 8 h, and then water-cool it to room temperature. Finally, directly put the alloy into a heat treatment furnace at 750 °C for high-temperature aging for 4 h, and then water-cool it to room temperature.

[0059] Example 2

[0060] This example provides a Fe-based superalloy, which includes by mass percentage: C: 0.05%, Ni: 25%, Cr: 13.5%, W: 0.3%, Mo: 0.6%, Ti: 2.5%, Al: 1.0%, Co: 1.2%, Mn: 0.5; Si: 0.015%, B: 0.003%, Zr: 0.02, and the balance is Fe.

[0061] The preparation method of the Fe-based superalloy in this example includes the following steps:

[0062] Step 1: After considering the burn-off, complete the batching. Pass argon gas into the vacuum arc melting furnace as the protective gas, and obtain a 150 g master alloy ingot with uniform alloy composition through electromagnetic stirring and flipping.

[0063] Step 2: Heat the master alloy ingot in the furnace at a rate of 20 °C / min to 1160 °C for homogenization treatment for 24 h, and then air-cool it to room temperature.

[0064] Step 3: Perform hot rolling on the homogenized master alloy ingot at 230 °C above the γ′ phase precipitation temperature, with a deformation amount of 20% per pass and a final total deformation amount of 60%. After each pass of rolling is completed, heat it back in the furnace and keep it warm for 30 min.

[0065] Step 4: After the alloy rolling is completed, heat it up to 1100 °C for solution treatment for 0.5 h, and then water-cool it to room temperature. Subsequently, directly put the alloy into a heat treatment furnace at 1000 °C for high-temperature aging for 1 h, and then air-cool it to room temperature. Next, directly put the alloy into a heat treatment furnace at 650 °C for low-temperature aging for 8 h, and then water-cool it to room temperature. Finally, directly put the alloy into a heat treatment furnace at 750 °C for high-temperature aging for 4 h, and then water-cool it to room temperature.

[0066] Example 3

[0067] This example provides a Fe-based superalloy, which includes by mass percentage: C: 0.08%, Ni: 23%, Cr: 14%, W: 0.1%, Mo: 1.0%, Ti: 2.0%, Al: 1.5%, Co: 1.0%, Mn: 0.2; Si: 0.1%, B: 0.004%, Zr: 0.01, and the balance is Fe.

[0068] The preparation method of the iron-based superalloy in this embodiment includes the following steps:

[0069] Step 1: After considering the burn-off, complete the batching. Pass argon gas as the protective gas into the vacuum arc melting furnace, and obtain a 150g master alloy ingot with uniform alloy composition through electromagnetic stirring and flipping.

[0070] Step 2: Heat the master alloy ingot in the furnace at a rate of 30°C / min to 1160°C for homogenization treatment for 24h, and then air-cool to room temperature.

[0071] Step 3: Perform hot rolling on the homogenized master alloy ingot at a temperature 230°C above the γ′ phase precipitation temperature. The deformation per pass is 20%, and the final total deformation is 60%. After each pass of rolling, heat it back in the furnace and keep it warm for 30min.

[0072] Step 4: After the alloy rolling is completed, heat it up to 1100°C for solution treatment for 0.5h, and then water-cool to room temperature. Subsequently, directly put the alloy into a heat treatment furnace at 1000°C for high-temperature aging for 1h, and then air-cool to room temperature. Next, directly put the alloy into a heat treatment furnace at 650°C for low-temperature aging for 8h, and then water-cool to room temperature. Finally, directly put the alloy into a heat treatment furnace at 750°C for high-temperature aging for 4h, and then water-cool to room temperature.

[0073] Comparative Example 1

[0074] The alloy composition adopts the composition in the patent with the application number: 201911296733.3, and the name: A high-strength, high-toughness and oxidation-resistant iron-nickel-based superalloy and its preparation method (the content of Fe element in the alloy is 40%).

[0075] The alloy in this comparative example, by mass percentage, includes: the alloy composition of this comparative example, by mass percentage, is C: 0.05%, Fe: 40%, Cr: 16%, Co: 1.0%, W: 0.3%, Mo: 0.6%, Ti: 1.8%, Al: 1.6%, Si: 0.025%, B: 0.002%, Nb: 0.05%, Mn: 0.1%, and the balance is Ni.

[0076] The preparation method of the alloy in this comparative example is the same as that of Example 1.

[0077] Comparative Example 2

[0078] This comparative example is the nickel-iron-based alloy GH2984 (33% Fe).

[0079] Comparative Example 3

[0080] The alloy of this comparative example, by mass percentage, includes: C: 0.02%, Fe: 25%, Cr: 21%, W: 0.7%, Mo: 0.6%, Ti: 1.7%, Al: 1.6%, Si: 0.002%, B: 0.003%, Nb: 1.4%, Mn: 0.3%, Zr: 0.02%, and the balance is Ni.

[0081] The preparation method of the alloy of this comparative example includes:

[0082] Step 1: Prepare a formed alloy tube by vacuum induction melting;

[0083] Step 2: Put the alloy tube into a heat treatment furnace at 1100°C for solution treatment for 0.5 h and then water-cool to room temperature. Subsequently, directly put the alloy tube into a heat treatment furnace at 1000°C for high-temperature aging for 0.5 h, and then water-cool to room temperature. Next, directly put the alloy tube into a heat treatment furnace at 650°C for low-temperature aging for 8 h and then water-cool to room temperature. Finally, directly put the alloy ingot into a heat treatment furnace at 750°C for high-temperature aging for 4 h and then water-cool to room temperature.

[0084] Comparative Example 4

[0085] The alloy composition of this comparative example, by mass percentage, is C: 0.05%, Fe: 55%, Cr: 16%, Co: 1.0%, W: 0.3%, Mo: 0.6%, Ti: 1.8%, Al: 1.6%, Si: 0.025%, B: 0.002%, Nb: 0.05%, Mn: 0.1%, and the balance is Ni.

[0086] The preparation method of the alloy of this comparative example is the same as that of Example 1.

[0087] Test Example

[0088] Figure 1 、 Figure 2 They are the low-magnification image and high-magnification image of the microstructure of the alloy prepared in Example 1 respectively. Fine blocky and discontinuously distributed MC carbides precipitate at grain boundaries and twin boundaries, and a large number of fine γ' phases precipitate inside the grains. These phases are the key to improving the high-temperature performance of the material.

[0089] The tensile yield strength of the alloys of the present invention and the comparative examples at 700°C is shown in Table 1. It can be seen from Table 1 that although a large amount of Fe element is added to the alloy of the present invention, through the synergistic cooperation of various components, the prepared iron-based alloy still has good high-temperature strength, which is equivalent to that of Comparative Example 1 and is significantly improved compared with Comparative Example 2. At the same time, the high-temperature alloy of the present application also has excellent tensile plasticity, and the elongation at 700°C is 27%. The alloy of the present invention has excellent strength and plasticity.

[0090] Table 1 is the tensile yield strength of the alloy at 700°C

[0091] alloy yield strength (MPa) Example 1 562 Example 2 572 Example 3 551 Comparative Example 1 573 Comparative Example 2 539

[0092] The tissue stability test includes: thermally exposing the alloy at 700 °C and measuring the compressive yield strength after a certain period of exposure. The test results are shown in Table 2.

[0093] Table 2 Compressive yield strength of the alloy at 700 °C after different thermal exposure times

[0094]

[0095] As can be seen from Table 2, Comparative Example 1

[0096] As can be seen from Table 2, although a large amount of Fe element is added to the alloy of the present invention, through the synergistic cooperation of each component, the reduction amplitude of the compressive yield strength of the prepared iron-based alloy after thermal exposure at 700 °C for 5000 h is very small, and it still has a very high strength, which can meet the service requirements. The thermal stability of the present invention is equivalent to that of Comparative Example 1 and is significantly improved compared with Comparative Example 2. The present invention greatly reduces the cost while having good high-temperature strength and thermal stability.

[0097] The microstructure and the types of the second phase of the alloy prepared in Example 1 after thermal exposure at 700 °C for different times were determined by using a scanning electron microscope (SEM), a transmission electron microscope (TEM) and an energy dispersive spectrometer (EDS), as Figures 3 - 5 shown. After aging at 700 °C / 5000 h and 700 °C / 10000 h, the sizes of the main strengthening phase γ′ phase in the alloys of Example 1 and Comparative Examples 1-3 are as Figure 6 shown. It can be seen from the figure that after aging at 700 °C / 10000 h, the average size of the γ′ phase particles in the alloy of Example 1 is 51 nm. Moreover, the coarsening rate of the γ′ phase particles in the alloy of Example 1 is less than that of the main strengthening phase γ′ phase in the alloys of Comparative Examples 1-3. After thermal exposure at 700 °C / 5000 h, as Figures 7 - 10 shown, a large amount of plate / needle-shaped σ phase or η phase has precipitated in the alloys of Comparative Examples 2-4, and the microstructure is unstable, while no plate / needle-shaped σ phase and η phase are found in the alloy of Example 1 after thermal exposure at 700 °C for 10000 h. This is because the reasonable ratio of elements such as Cr, Fe, Ti, Mo, and W in the alloy content designed in the present invention not only endows the material with excellent tissue thermal stability, but also obtains a more dispersed and finer strengthening phase. The compression test was carried out on the specimen by using an MTS universal testing machine. The compressive yield strength of the alloy of Example 1 at 700 °C after thermal exposure at 700 °C / 5000 h is 497 MPa. Combining the microstructure evolution and the mechanical property test results shows that the iron-based superalloy exhibits good thermal stability at 700 °C.

[0098] The microstructure of the alloy of Comparative Example 4 is as follows Figure 9 shown. The alloy structure is mainly composed of γ′ phase and carbides. After long-term thermal exposure at 700 °C for 1000 h, the alloy microstructure is as follows Figure 10 shown. It can be seen from Figure 10 that a large number of needle-shaped / plate-shaped topologically close-packed phases precipitate in the alloy, and these topologically close-packed phases will significantly deteriorate the mechanical properties of the alloy, especially the long-term creep properties. It can be seen from this that by simply increasing the content of Fe element in the alloy and reducing the content of Ni, the precipitation of needle-shaped and plate-shaped harmful phases in the alloy cannot be avoided during long-term service, and the obtained alloy microstructure is unstable.

[0099] Obviously, the above embodiments are merely examples given for clear illustration and are not limitations on the implementation manners. For those of ordinary skill in the art, other different forms of changes or modifications can be made on the basis of the above description. It is not necessary and impossible to list all the implementation manners here. And the obvious changes or modifications derived therefrom are still within the protection scope of the present invention.

Claims

1. A ferrous superalloy, characterized in that by mass percentage, the composition includes: C: 0.05 - 0.08%, Ni: 23 - 25%, Cr: 13.5 - 14.8%, W: 0.1 - 0.3%, Mo: 0.6 - 1.0%, Ti: 2.0 - 2.5%, Al: 1.0 - 1.5%, Co: 1.0 - 1.6%, Mn: ≤0.5%, Si: ≤0.3%, B: ≤0.005%, Zr: ≤0.03%, and the balance is Fe; The preparation method of the ferrous superalloy includes the following steps: Step 1: Melt the raw materials into a master alloy ingot in a vacuum environment; Step 2: Perform homogenization annealing on the master alloy ingot; Step 3: High-temperature rolling; Step 4: Heat treatment; The said Step 3 includes: Perform high-temperature rolling on the master alloy ingot after homogenization annealing at 200 - 250 °C above the γ′ phase precipitation start temperature, with the deformation per pass being 15 - 25%, and the final total deformation being not less than 60%; The said Step 4 includes solution treatment and aging treatment; The said aging treatment includes the first high-temperature aging, low-temperature aging, and the second high-temperature aging; The said first high-temperature aging includes: Heat the heat treatment furnace to 130 - 170 °C above the γ′ phase precipitation start temperature, put the alloy after solution treatment into the heat treatment furnace for 1 - 2 h, and cool it to room temperature with water; The said low-temperature aging includes: Heat the heat treatment furnace to 150 - 200 °C below the γ′ phase precipitation start temperature, put the alloy after the first high-temperature aging into the heat treatment furnace for 8 - 16 h, and cool it to room temperature with water; The said second high-temperature aging includes: Heat the heat treatment furnace to 30 - 150 °C below the γ′ phase precipitation start temperature, put the alloy after low-temperature aging into the heat treatment furnace for 4 - 24 h, and cool it to room temperature with water; The said solution treatment includes: Heat the heat treatment furnace to 250 - 300 °C above the γ′ phase precipitation start temperature, put the alloy after high-temperature rolling into the heat treatment furnace for solution treatment for 0.5 - 2 h, and then cool it to room temperature with water.

2. The ferrous superalloy according to claim 1, characterized in that the said Step 2 includes: Heat the master alloy ingot from room temperature to 300 - 350 °C above the γ′ phase precipitation start temperature at a heating rate of 10 - 30 °C / min, perform homogenization annealing for 20 - 35 h, and then air-cool it to room temperature.

3. Application of the ferrous superalloy according to claim 1 or 2 in a thermal power unit.

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