A nickel-based heat-resistant alloy for high-parameter steam turbine fasteners and a preparation method thereof

Abstract

The present application relates to a kind of high parameter steam turbine fastener nickel-based heat-resistant alloy and its preparation method, belong to heat-resistant alloy technical field, solve the technical problem that the strength and toughness of existing 700 DEG C grade nickel-based heat-resistant alloy are difficult to consider under long time high temperature service condition.The alloy of the present application is composed of C, Cr, Co, Al, Ti, Mo, Zr, B and Ni element by mass percentage, and satisfy 5.50%≤Al+Ti≤6.00%, 0.35≤Ti / Al≤0.90;Its preparation method includes vacuum induction melting, control ingot cooling rate is 130~160 ℃ / min, forging and heat treatment.The room temperature impact toughness KU2 of the nickel-based heat-resistant alloy provided by the present application is greater than or equal to 73J, the fracture time under the condition of 700 DEG C / 585MPa is greater than or equal to 850 hours, and high strength and high toughness good match in 700 DEG C advanced ultra-supercritical unit steam turbine fastener are realized.

Description

Technical Field

[0001] This invention relates to the field of heat-resistant alloy technology, and in particular to a nickel-based heat-resistant alloy for high-parameter steam turbine fasteners and its preparation method. Background Technology

[0002] As the steam parameters of advanced ultra-supercritical units increase from 600℃ to 700℃, the service environment of components such as high-temperature fasteners in steam turbines becomes increasingly demanding. The upper limit of the service temperature of traditional ferritic heat-resistant steels (such as 1Cr11MoV, 20Cr1Mo1VTiB, etc.) is 600℃, which cannot meet the service requirements of 700℃-class units. Nickel-based high-temperature alloys must be used.

[0003] Currently, nickel-based superalloys for 700℃ fasteners mainly include grades such as Nimonic 105, Inconel 740H, and Haynes 282. Among them, Nimonic 105, as an age-hardening forged nickel-based alloy, uses the γ′ phase (Ni3(Al,Ti)) as the main strengthening phase. It exhibits a creep rupture strength exceeding 100 MPa at 700℃ and 100,000 hours, demonstrating good high-temperature strength. The modified GH4145 alloy, through the addition of trace amounts of Mg, Zr, and B elements and optimized heat treatment processes, successfully eliminated the notch sensitivity of the Inconel X-750 alloy in the mid-temperature range of 500~700℃.

[0004] However, the Nimonic 105 alloy has an impact absorption energy of less than 30 Joules under standard heat treatment conditions, resulting in insufficient toughness and increasing the possibility of brittle fracture in practical use. Although the modified GH4145 alloy improves its mid-temperature plasticity, its γ′ phase strengthening level is limited, making it difficult to meet the strength requirements for long-term service at 700℃. Summary of the Invention

[0005] Based on the above analysis, the present invention aims to provide a nickel-based heat-resistant alloy for high-parameter steam turbine fasteners and its preparation method, in order to solve the technical problem that existing 700℃-grade nickel-based heat-resistant alloys are difficult to balance in terms of strength and toughness under long-term high-temperature service conditions.

[0006] On one hand, embodiments of the present invention provide a nickel-based heat-resistant alloy for high-parameter steam turbine fasteners. The chemical composition of the alloy, by mass percentage, is as follows: C: 0.05~0.25%, Si: ≤0.10%, Mn: ≤0.10%, P: ≤0.010%, S: ≤0.010%, Cr: 13.0~17.0%, Co: 17.0~23.0%, Al: 3.0~5.0%, Ti: 1.60~3.00%, Mo: 3.0~6.0%, Zr: 0.03~0.12%, B: 0.005~0.01%, and satisfies: 5.50%≤Al+Ti≤6.00%, 0.35≤Ti / Al≤0.90, with the balance being Ni and unavoidable impurities.

[0007] Furthermore, the Al+Ti content is 5.65%–5.85%.

[0008] Furthermore, the γ′ strengthening phase in the alloy microstructure is spherical with an average grain size of 10–50 nm.

[0009] Furthermore, the room temperature impact toughness of the alloy is KU2≥73J.

[0010] Furthermore, the alloy exhibits a fracture time of ≥850 hours under conditions of 700℃ / 585MPa.

[0011] On the other hand, the present invention also proposes a method for preparing the above-mentioned high-parameter steam turbine fastener nickel-based heat-resistant alloy, the method comprising the following steps: S1: Vacuum induction melting is used to melt the proportioned raw materials and then cast them into steel ingots; S2: Forced cooling is adopted after casting, and the cooling rate of the steel ingot is controlled to be ≥10℃ / minute; S3: After demolding, the steel ingot is annealed and then forged. The forging opening temperature is 1150~1170℃ and the final forging temperature is ≥1000℃ to obtain the billet. S4: Heat treatment is performed on the forged billet.

[0012] Furthermore, in step S1, during the refining stage of vacuum induction melting, the melting power is controlled at 600-800W.

[0013] Furthermore, in step S1, Al is added by repeated immersion: the Al ingot is immersed 5-20 mm below the liquid surface to melt, then lifted out, and the operation is repeated until it is completely melted.

[0014] Furthermore, in step S2, the ingot is demolded when its surface cools to 300-350°C.

[0015] Furthermore, in step S4, the heat treatment includes: Solution treatment: Hold at 1100-1180℃ for 2-6 hours, then cool. Intermediate treatment: Hold at 1000~1060℃ for 12~20 hours, then cool. Aging treatment: Keep warm at 700-850℃ for 12-20 hours, then cool.

[0016] Compared with the prior art, the present invention can achieve at least one of the following beneficial effects: 1) This invention, through the synergistic proportioning of Al and Ti elements, controls the total Al+Ti content to 5.50%–6.00% and the Ti / Al ratio to 0.35–0.90, causing the γ′ phase to disperse and precipitate in a fine spherical rather than cubic shape. The spherical γ′ phase is geometrically symmetrical and has a low stress concentration factor, making it less likely to become a crack initiation point under stress, thus maintaining high strength (R at room temperature). m ≥1245MPa, R p0.2 ≥850MPa, 700℃R m ≥1050MPa, R p0.2 While achieving an impact strength of ≥760MPa, the room temperature impact toughness KU2 is increased to ≥73J. At the same time, this microstructure exhibits excellent microstructure stability during long-term service at 700℃, with creep rupture time reaching more than 850 hours and more than 2435 hours under 700℃ / 585MPa and 700℃ / 500MPa conditions, respectively, overcoming the shortcomings of Nimonic 105 alloy with impact toughness below 30J and insufficient high-temperature creep performance.

[0017] 2) This invention employs water cooling after casting. The steel ingot is rapidly cooled by adding a water-cooling jacket to the outside of the ingot mold, controlling the cooling rate to ≥10℃ / min. This rapid cooling reduces the segregation of alloying elements and ensures that primary carbides are uniformly dispersed in the matrix as fine particles, preventing carbide agglomeration and continuous precipitation along grain boundaries. Simultaneously, rapid cooling also suppresses the precipitation of the γ′ phase during the ingot cooling process, preventing its coarsening during subsequent aging. Combined with a three-stage heat treatment process—solution treatment, intermediate treatment, and aging treatment—the intermediate treatment further enhances the M… 23 C6-type carbides precipitate in a chain-like pattern at grain boundaries, effectively pinning the grain boundaries and improving the alloy's high-temperature creep rupture performance. The γ′ phase is distributed in fine, dispersed spherical shapes, precipitating within the grains with an average grain size of 10–50 nm, improving the alloy's impact toughness and fracture performance. The alloy exhibits fracture times exceeding 850 hours at 700℃ / 585 MPa and 2435 hours at 700℃ / 500 MPa.

[0018] 3) This invention does not add W or Nb elements, thus avoiding the problems of W and Nb increasing the resistance to hot deformation, promoting the precipitation of harmful TCP phases and causing solidification segregation. It also reduces the difficulty of hot working of the alloy. It can be formed at a forging opening temperature of 1150-1170℃ and a final forging temperature of ≥1000℃, avoiding the problem of complex hot working processes required for nickel-based alloys containing W and Nb, which is conducive to engineering mass production.

[0019] In this invention, the above-described technical solutions can be combined with each other to achieve more preferred combinations. Other features and advantages of this invention will be set forth in the following description, and some advantages may become apparent from the description or be learned by practicing the invention. The objects and other advantages of this invention can be realized and obtained from what is particularly pointed out in the description and drawings. Attached Figure Description

[0020] The accompanying drawings are for illustrative purposes only and are not intended to limit the invention. Throughout the drawings, the same reference numerals denote the same parts.

[0021] Figure 1 This is a transmission electron microscope (TEM) image of the γ′ strengthening phase in the alloy of Example 2 of the present invention; Figure 2 This is a transmission electron microscope (TEM) image of the γ′ strengthening phase in the alloy of Example 4 of the present invention. Figure 3 This is a scanning electron microscope image of the primary carbides in the alloy of Example 2 of the present invention; Figure 4 This is a scanning electron microscope image of the primary carbides in the alloy of Example 4 of the present invention; Figure 5 This is a scanning electron microscope image of the primary carbides in the alloy of Comparative Example 1 of the present invention. Detailed Implementation

[0022] The preferred embodiments of the present invention will now be described in detail with reference to the accompanying drawings, which constitute a part of the present invention and are used together with the embodiments of the present invention to illustrate the principles of the present invention, but are not intended to limit the scope of the present invention.

[0023] Existing 700℃-grade nickel-based heat-resistant alloys suffer from the problem of balancing strength and toughness. To address this issue, this invention proposes a high-parameter nickel-based heat-resistant alloy for steam turbine fasteners and its preparation method.

[0024] On one hand, a specific embodiment of the present invention discloses a nickel-based heat-resistant alloy for high-parameter steam turbine fasteners. The chemical composition of the alloy, by mass percentage, is as follows: C: 0.05~0.25%, Si: ≤0.10%, Mn: ≤0.10%, P: ≤0.010%, S: ≤0.010%, Cr: 13.0~17.0%, Co: 17.0~23.0%, Al: 3.0~5.0%, Ti: 1.60~3.00%, Mo: 3.0~6.0%, Zr: 0.03~0.12%, B: 0.005~0.01%, and satisfies: 5.50%≤Al+Ti≤6.00%, 0.35≤Ti / Al≤0.90, with the balance being Ni and unavoidable impurities.

[0025] The functions and roles of each element in the above alloy composition are as follows: Aluminum (Al): Al is the most important and essential element for the formation of the γ′ phase (Ni3Al). The Al content directly affects the volume fraction of the γ′ phase and the high-temperature strength of the alloy. When the Al content is too low, the volume fraction of the γ′ phase is insufficient, and the high-temperature strength of the alloy cannot meet the service requirements of 700℃; when the Al content is too high, the volume fraction of the γ′ phase is too large, the plastic deformation capacity of the alloy decreases, and the toughness deteriorates. This invention controls the Al content at 3.0% to 5.0% (e.g., 3.0%, 3.2%, 3.5%, 3.8%, 4.0%, 4.2%, 4.5%, 4.8%, 5.0%), which ensures a sufficient volume fraction of the γ′ phase to obtain high strength, while avoiding the decrease in toughness caused by excessive Al content.

[0026] Titanium (Ti): Ti can partially replace Al in the γ′ phase to form Ni3(Al,Ti), significantly increasing the ordering temperature and dissolution temperature of the γ′ phase, thereby raising the maximum service temperature of the alloy. However, excessive Ti content promotes the precipitation of the harmful η phase (Ni3Ti), consuming Ti and impairing the alloy's plasticity and toughness. This invention controls the Ti content to 1.60%–3.00% (e.g., 1.60%, 1.80%, 2.00%, 2.20%, 2.40%, 2.60%, 2.80%, 3.00%), improving the high-temperature stability of the γ′ phase while preventing the precipitation of the η phase.

[0027] Total Al+Ti content: Al and Ti are the main elements forming the γ′ phase (Ni3(Al,Ti)). The γ′ phase is the most important strengthening phase in the nickel-based heat-resistant alloy of this invention, and its volume fraction directly determines the high-temperature strength of the alloy. When the total Al+Ti content is too low, the volume fraction of the γ′ phase is insufficient, and the high-temperature strength of the alloy cannot meet the service requirements of 700℃; when the total Al+Ti content is too high, the volume fraction of the γ′ phase is too large, the plastic deformation capacity of the alloy decreases, and the toughness deteriorates. This invention controls the total Al+Ti content at 5.50% to 6.00% (e.g., 5.50%, 5.55%, 5.60%, 5.70%, 5.80%, 5.90%, 5.95%, 6.00%), ensuring a sufficient volume fraction of the γ′ phase to obtain high strength while retaining good plastic deformation capacity of the alloy.

[0028] Preferably, an Al+Ti content of 5.65% to 5.85% yields the best balance between strength and toughness.

[0029] The titanium to aluminum ratio (Ti / Al): The Ti / Al ratio determines the morphology and stability of the γ′ phase. When the Ti / Al ratio is low, the γ′ phase tends to be cubic in shape. The sharp corners of the cubic γ′ phase are prone to stress concentration, becoming crack initiation points and reducing the impact toughness of the alloy. When the Ti / Al ratio is moderate, the γ′ phase is spherical. The spherical γ′ phase is geometrically symmetrical and has a low stress concentration factor, making it less likely to become a crack initiation point under stress. This invention controls the Ti / Al ratio within the range of 0.35 to 0.90 (e.g., 0.35, 0.40, 0.45, 0.50, 0.55, 0.60, 0.70, 0.80, 0.90) to make the γ′ phase precipitate in a spherical shape, thereby improving impact toughness while maintaining high strength.

[0030] Furthermore, this invention does not add W and Nb elements. The addition of W and Nb increases the alloy's resistance to hot deformation, promotes the precipitation of harmful TCP phases (such as μ and σ phases), and causes solidification segregation, leading to a narrower hot working window and an increased risk of forging cracks, often requiring multi-stage forging or powder metallurgy processes. This invention employs a W- and Nb-free composition design, allowing for forging at a starting temperature of 1150–1170℃ and a final forging temperature ≥1000℃, which is beneficial for engineered mass production.

[0031] Zirconium (Zr) and Boron (B): Zr and B are grain boundary strengthening elements. Zr atoms tend to segregate at grain boundaries, filling grain boundary vacancies, inhibiting the formation and expansion of grain boundary voids, and forming high-melting-point compounds with harmful impurities such as sulfur and oxygen, thus improving grain boundary bonding. B atoms also tend to segregate at grain boundaries, which can delay grain boundary slip and the nucleation and connection of voids. In this invention, the Zr content is controlled at 0.03% to 0.12% (e.g., 0.03%, 0.04%, 0.05%, 0.06%, 0.08%, 0.10%, 0.12%), and the B content is controlled at 0.005% to 0.01% (e.g., 0.005%, 0.006%, 0.007%, 0.008%, 0.009%, 0.01%). Through the combined addition of Zr and B, the grain boundaries are synergistically strengthened, thereby improving the creep resistance and fracture life of the alloy.

[0032] Chromium (Cr): The main function of Cr is to form a protective oxide film of Cr₂O₃, improving the alloy's resistance to oxidation and hot corrosion. Simultaneously, Cr atoms dissolve in the γ matrix, causing lattice distortion and playing a solid solution strengthening role, and forming M... 23 C6-type carbides precipitate at grain boundaries, pinning them. This invention controls the Cr content to 13.0%–17.0% (e.g., 13.0%, 14.0%, 14.5%, 15.0%, 15.5%, 16.0%, 17.0%), ensuring antioxidant properties while preventing the precipitation of harmful phases.

[0033] Cobalt (Co): Co can reduce the stacking fault energy of the austenitic matrix, promote the formation of deformation twins during creep, delay microcrack propagation, and improve the creep resistance and fracture life of the alloy. Simultaneously, Co dissolves in the γ matrix, increasing the dissolution temperature of the γ′ phase and enhancing the stability of the high-temperature microstructure. This invention controls the Co content to 17.0%–23.0% (e.g., 17.0%, 18.0%, 19.0%, 20.0%, 21.0%, 22.0%, 23.0%) to achieve optimal creep performance within a reasonable cost.

[0034] Molybdenum (Mo): Mo has a large atomic radius, and when dissolved in an austenitic matrix, it causes lattice distortion, resulting in a strong solid solution strengthening effect. Simultaneously, Mo can improve the bonding force between matrix atoms and inhibit the coarsening of the γ′ strengthening phase. This invention controls the Mo content at 3.0%–6.0% (e.g., 3.0%, 3.5%, 4.0%, 4.5%, 5.0%, 5.5%, 6.0%) to achieve sufficient solid solution strengthening while avoiding the precipitation of harmful phases.

[0035] Carbon (C): C reacts with elements such as Cr, Ti, and Mo in the alloy to form stable carbides (such as MC-type, M-type). 23C6 type carbides precipitate at grain boundaries and within grains, pinning grain boundaries, hindering dislocation movement, and improving high-temperature strength and creep resistance. This invention controls the C content to 0.05%–0.25% (e.g., 0.05%, 0.08%, 0.10%, 0.15%, 0.20%, 0.25%) to obtain an appropriate amount of dispersed carbides, strengthening grain boundaries.

[0036] Silicon (Si) and manganese (Mn): As impurity elements, excessive Si and Mn content will reduce the alloy's oxidation resistance and hot workability. In this invention, both Si and Mn are controlled to ≤0.10%.

[0037] Phosphorus (P) and sulfur (S): P and S are harmful impurities that easily segregate at grain boundaries, reducing grain boundary bonding and causing thermal brittleness. In this invention, both P and S are controlled to ≤0.010%.

[0038] Furthermore, the nickel-based heat-resistant alloy for high-parameter turbine fasteners of the present invention has a microstructure in which the γ′ strengthening phase is spherical with an average grain size of 10–50 nm. Mg precipitates at the grain boundaries. 23 C6-type carbides (where M is mainly Cr, with solid solutions of elements such as Mo, Ni, and Co) are distributed in a discontinuous chain pattern, with a precipitation amount of 0.9% to 1.1% (mass fraction). They effectively pin grain boundaries, inhibit grain boundary slip and crack propagation at high temperatures, and synergistically enhance the high-temperature creep fracture performance of the alloy with spherical γ′ phases.

[0039] Furthermore, the nickel-based heat-resistant alloy for high-parameter turbine fasteners of the present invention exhibits excellent comprehensive mechanical properties. In terms of strength, the alloy's room temperature tensile strength R... m ≥1245MPa, yield strength R p0.2 ≥850MPa, tensile strength at 700℃ m ≥1050MPa, yield strength R p0.2 With a strength of ≥760MPa, it is comparable to Nimonic 105 alloy, meeting the basic requirements of fasteners for high strength.

[0040] Based on this, the present invention optimizes the microstructure, enabling the alloy to maintain good plasticity at both room temperature and 700℃ (room temperature elongation after fracture A is 21.0%~25.0%, reduction of area Z is 23.5%~32.5%, and at 700℃, the elongation after fracture A is 18.5%~28.5%, and the reduction of area Z is 24.0%~31.0%). In particular, the room temperature impact toughness KU2≥73J is much higher than that of Nimonic105 alloy (usually <30J), effectively solving the industry pain point of mismatch between high strength and toughness in existing similar alloys.

[0041] Meanwhile, the primary carbides in the alloy are controlled at level 4-5, with significant carbide refinement, ensuring the uniformity of the microstructure. In terms of creep resistance, the alloy of this invention exhibits fracture times ≥850 hours and ≥2435 hours under stress conditions of 700℃ / 585MPa and 700℃ / 500MPa, respectively, demonstrating excellent high-temperature creep resistance.

[0042] In summary, the alloy of this invention, while maintaining high-temperature strength comparable to Nimonic 105, significantly improves plasticity and impact toughness, and also possesses excellent high-temperature creep resistance, making it more suitable for the long-life and reliable service requirements of high-parameter turbine fasteners.

[0043] On the other hand, the present invention also proposes a method for preparing the above-mentioned high-parameter steam turbine fastener nickel-based heat-resistant alloy, the method comprising the following steps: S1: Vacuum induction melting is used to melt the proportioned raw materials and then cast them into steel ingots; S2: Forced cooling is used after casting, and the cooling rate of the steel ingot is controlled at ≥10℃ / minute. The ingot is demolded when the surface of the steel ingot cools to 300~350℃. S3: After demolding, the steel ingot is annealed and then forged. The forging opening temperature is 1150~1170℃ and the final forging temperature is ≥1000℃ to obtain the billet. S4: Heat treatment is performed on the forged billet.

[0044] Specifically, in S1, vacuum induction smelting is employed, and all raw materials used in the smelting process are pure metals, including metallic Cr, Al, sponge Ti, Ni plates, metallic Mo strips, metallic Co, sponge Zr, and Ni-B alloys. Before smelting, the content of impurity elements in each raw material is checked to ensure that P: ≤0.010%, S: ≤0.010%, Pb: ≤0.001%, As: ≤0.0025%, Sn: ≤0.0015%, Sb: ≤0.0025%, and Bi: ≤0.0001%.

[0045] During charging, a small layer of pure Ni plate is first laid at the bottom of the crucible. This protects the bottom of the crucible from direct impact from the high-density charge and also provides good thermal conductivity, which is beneficial for initial melting. The middle layer contains high-melting-point, non-oxidizing main elements such as Co, Cr, and Mo. The top layer of the crucible contains the remaining Ni and C (deoxidizer). This allows the ferroalloy to quickly form a molten pool in the early stages of melting, melting the high-melting-point elements in the middle and improving efficiency. To prevent oxidation, sponge Ti, sponge Zr, and Ni-B alloy are added in the later stages of smelting. Al is very reactive and easily oxygenates the molten steel. It is added using a repeated immersion method: after the surface of the molten steel cools and forms a film during smelting, the bundled Al ingots are slowly immersed 5–20 mm (e.g., 5 mm, 8 mm, 10 mm, 12 mm, 15 mm, 18 mm, 20 mm) below the liquid surface to melt, then lifted out, and the operation is repeated until all the ingots are melted. This method avoids the mechanical impact of directly adding Al ingots to the furnace bottom. At the same time, by controlling the temperature during melting, it effectively mitigates the localized and violent splashing caused by the concentrated melting of Al elements, thereby achieving temperature-controlled melting, precisely controlling the Al composition, and avoiding the oxidation and burning loss of Al elements.

[0046] During the initial preheating stage of the smelting process, the furnace power is controlled at 200-300W to preheat the furnace charge and begin degassing. As the temperature rises, the power is gradually increased to 800-1000W to rapidly melt the furnace charge and form a molten pool. During the melting stage, the melting temperature is controlled at 1500-1520℃, and the vacuum level is controlled at 1.0-1.5 Pa to facilitate gas escape. The melting stage lasts for 15-20 minutes.

[0047] Furthermore, in step S1, during the refining stage of vacuum induction melting, the power is reduced to 600-800W (e.g., 600W, 650W, 700W, 720W, 750W, 780W, 800W), sufficient only to maintain the smelting temperature and appropriate electromagnetic stirring. Conventional refining processes typically use high power (e.g., above 1000W) to enhance electromagnetic stirring and promote composition homogenization. However, high-power stirring intensifies the tumbling of molten steel, causing suspended inclusions to be re-entered into the molten steel, reducing alloy purity. This invention employs low-power refining, ensuring sufficient degassing and composition homogenization of the molten steel while avoiding the re-entry of inclusions due to strong stirring. This improves alloy purity, reduces the risk of inclusions becoming crack initiations, and indirectly enhances impact toughness.

[0048] The refining stage requires heating to 1520-1550℃ to ensure sufficient fluidity of the molten steel, promoting the aggregation and flotation of inclusions. The vacuum level should be below 1.2 Pa; a higher vacuum level is more conducive to the forward degassing reaction. The refining period should be controlled within 40-45 minutes.

[0049] The active elements are added after refining and degassing, at which point the molten steel is cooled to 1480-1510°C to minimize the loss of these active elements. Argon gas is introduced for protection during the addition of these elements to further improve the yield.

[0050] Specifically, in step S1, when casting the steel ingot, the casting temperature is 1460-1500℃. If the temperature is too high, it will aggravate macro segregation; if the temperature is too low, it may lead to poor fluidity of the molten steel.

[0051] Specifically, in step S2, after casting, forced cooling is achieved by adding a water-cooling jacket to the outside of the ingot mold. The cooling jacket is made of 45# steel with high thermal conductivity and filled with cooling water. The cooling rate is adjusted by controlling the water flow rate and volume. The cooling rate of the ingot is controlled to be ≥10℃ / min (e.g., 10℃ / min, 15℃ / min, 20℃ / min, 25℃ / min, 30℃ / min, 35℃ / min, 40℃ / min). Conventional ingot cooling often uses air cooling, with a cooling rate usually below 1℃ / min. This invention uses rapid cooling, which can reduce the segregation of alloying elements and the distribution of primary carbides, and suppress the coarsening of the γ′ phase during the ingot cooling process, laying the microstructure foundation for obtaining fine and dispersed spherical γ′ phases in subsequent heat treatment. The ingot is demolded when the surface cools to 300-350℃. Too high a demolding temperature can easily lead to ingot deformation, while too low a temperature increases the risk of cracking.

[0052] Specifically, in step S3, the demolded steel ingot is annealed and then forged. A 2000-ton high-speed forging mill is used for billet preparation, with a billet preparation temperature of 1150-1170℃ and a final forging temperature ≥1000℃, to obtain a billet (e.g., a 90mm×90mm square billet). Because this invention does not add W or Nb elements, it avoids the problems of W and Nb increasing hot deformation resistance, promoting the precipitation of harmful TCP phases, and causing solidification segregation. The alloy has good hot workability, and can be formed at a forging preparation temperature of 1150-1170℃ and a final forging temperature ≥1000℃, which is beneficial for engineering-scale mass production.

[0053] Furthermore, in step S4, the forged billet is subjected to heat treatment, which includes solution treatment, intermediate treatment, and aging treatment. (1) Solution treatment: Heat to 1100-1180℃ (e.g., 1100℃, 1120℃, 1130℃, 1150℃, 1160℃, 1170℃, 1180℃), hold for 2-6 hours (e.g., 2h, 3h, 4h, 5h, 6h), and then cool (e.g., air cooling); (2) Intermediate treatment: Heat to 1000-1060℃ (e.g., 1000℃, 1010℃, 1020℃, 1030℃, 1040℃, 1050℃, 1060℃), keep warm for 12-20 hours (e.g., 12h, 14h, 16h, 18h, 20h), and then cool (e.g., air cooling); (3) Aging treatment: Heat to 700-850℃ (e.g., 700℃, 730℃, 760℃, 790℃, 820℃, 850℃), keep warm for 12-20 hours (e.g., 12h, 14h, 16h, 18h, 20h), and then cool (e.g., air cooling).

[0054] The above three-stage heat treatment process, combined with the ingot cooling rate (≥10℃ / min), has the following synergistic effects: rapid ingot cooling reduces alloy element segregation, improves primary carbide distribution and inhibits the coarsening of the primary γ′ phase; solution treatment dissolves the residual coarse γ′ phase; and intermediate treatment promotes the growth of M... 23 The C6 phase nucleates at grain boundaries and is distributed in a chain-like pattern, hindering grain boundary movement at high temperatures and thus improving the alloy's creep retardation properties. Intermediate treatment at 1000–1060℃ for 12–20 hours can ensure the presence of M in the alloy. 23 The C6 phase content is 0.13~0.37% (mass percentage). Aging treatment at 700~850℃ for 12~20 hours can control the growth of the γ′ phase to the target size (10~50nm) with a content of 22.30~31.25% (mass percentage), ultimately obtaining a fine and dispersed spherical γ′ phase structure, achieving a good match between high strength, high toughness and high durability.

[0055] The high-parameter nickel-based heat-resistant alloy for steam turbine fasteners of this invention can be used to manufacture high-temperature structural components above 700°C, including but not limited to: fasteners (bolts, studs) for advanced ultra-supercritical power plant steam turbines, valve forgings, and turbine and combustion chamber components for gas turbines. Furthermore, it can also be applied to key load-bearing components in spacecraft, rocket engines, nuclear reactors, petrochemical equipment, and energy conversion devices such as coal conversion plants.

[0056] In summary, the high-parameter nickel-based heat-resistant alloy for steam turbine fasteners and its preparation method provided by this invention, through optimization of chemical composition and matching heat treatment process, ensures that the γ′ strengthening phase in its microstructure is spherically and uniformly distributed with an average particle size controllable within the range of 10–50 nm, thereby effectively balancing the alloy's high-temperature creep strength and room-temperature impact toughness. Regarding mechanical properties, the alloy of this invention exhibits a room-temperature tensile strength R0. m ≥1245MPa, yield strength R p0.2 ≥850MPa, tensile strength at 700℃ m ≥1050MPa, yield strength R p0.2With an impact toughness of ≥760MPa and room temperature impact toughness of KU2≥73J, and fracture times of ≥2435 hours and ≥850 hours under stress conditions of 700℃ / 500MPa and 700℃ / 585MPa respectively, it is significantly better than existing commercial nickel-based heat-resistant alloys and can fully meet the long-term stable operation requirements of high-parameter turbine fasteners under extreme service conditions.

[0057] The present invention will be described in more detail below through specific embodiments. These embodiments are merely descriptions of the best implementation of the invention and do not limit the scope of the invention in any way.

[0058] Examples 1-7 and Comparative Example 3 Seven furnaces of the high-parameter steam turbine fastener nickel-based heat-resistant alloy of the present invention were smelted according to the chemical composition shown in Table 2, respectively as Examples 1-7; at the same time, commercial Nimonic 105 alloy was used as Comparative Example 3 (its chemical composition is also shown in Table 2).

[0059] Examples 1-7 were all prepared according to the following process route: S1: Vacuum induction melting is used to melt the proportioned raw materials and then cast them into steel ingots.

[0060] The smelting raw materials are pure metals, including metallic Cr, Al, sponge Ti, Ni plates, metallic Mo strips, metallic Co, sponge Zr, and Ni-B alloys. The content of impurity elements is controlled as follows: P≤0.010%, S≤0.010%, Pb≤0.001%, As≤0.0025%, Sn≤0.0015%, Sb≤0.0025%, Bi≤0.0001%.

[0061] During the loading process, a small layer of pure Ni plate is first laid at the bottom of the crucible, then Co, Cr and Mo are loaded in the middle layer, and the remaining Ni and C are loaded in the top layer. Sponge Ti, sponge Zr and Ni-B alloy are added in the later stage of smelting.

[0062] The initial preheating power was 250W, which was gradually increased to 900W (heating rate was 60℃ / min, and power increased at a rate of 120W / min) for rapid melting. The melting temperature was 1510℃, the vacuum degree was below 1.2Pa, and the time was 18 minutes. The refining stage uses a power of 650W, a temperature of 1530℃, a vacuum of less than 1.2Pa, and takes 42 minutes.

[0063] After refining, the temperature is lowered to 1500℃, protected by argon gas, and sponge Ti, sponge Zr and Ni-B alloy are added. Al is added to a depth of 10mm below the liquid surface by repeated immersion until it is completely melted; the casting temperature is 1480℃.

[0064] S2: After casting, water cooling jacket is used for forced cooling, and the cooling rate of the steel ingot is controlled at 11℃ / minute. The ingot is demolded when the surface of the steel ingot cools to 310℃.

[0065] S3: After demolding, the steel ingot is annealed and then forged. The forging opening temperature is 1150℃ and the final forging temperature is 1010℃ to obtain the billet.

[0066] S4: Heat treatment of the forged billet, including solution treatment (1100℃ for 4 hours), intermediate treatment (1030℃ for 16 hours) and aging treatment (750℃ for 16 hours).

[0067] Comparative Example 3 is a commercial Nimonic 105 alloy, which was subjected to the standard heat treatment process of the alloy: solution treatment at 1100℃ for 4 hours, intermediate treatment at 1030℃ for 16 hours, and aging treatment at 750℃ for 16 hours.

[0068] Table 2 Alloy Chemical Composition (wt%)

[0069] As can be seen from Table 2, the Ti content (1.60-3.00%) of the alloys in Examples 1-7 of the present invention is significantly higher than that in Comparative Example 3 (0.9-1.5%), and the Ti / Al ratio (0.35-0.90) is also significantly higher than that in Comparative Example 3 (0.18-0.33).

[0070] Examples 8-9 and Comparative Examples 1-2 Examples 8-9 use alloys with the same chemical composition as Example 1. The only difference between Examples 8-9 and Example 1 is that some process parameters in the preparation method are different, as shown in Table 1.

[0071] Comparative Examples 1-2 used alloys with the same chemical composition as Example 1. The only difference between them and Example 1 was the preparation method: in step S2, the cooling rate of the steel ingot was controlled. Comparative Example 1 was prepared by air cooling (cooling rate of 0.55℃ / min), and Comparative Example 2 was prepared by controlled cooling (cooling rate of 5.0℃ / min). Both were lower than the ≥10℃ / min range specified in this invention. The specific process parameters are shown in Table 1.

[0072] Table 1

[0073] The test results of the above-mentioned embodiments and comparative alloys are shown in Table 3: Table 3

[0074] As can be seen from Table 3, the alloys of Examples 1-9 of this invention exhibit excellent comprehensive mechanical properties. Room temperature tensile strength R mThe yield strength is 1245–1277 MPa (all ≥1245 MPa), and the yield strength R is... p0.2 The tensile strength is 853–887 MPa (all ≥850 MPa); the tensile strength at 700℃ is R m The yield strength is 1050–1095 MPa (all ≥1050 MPa), and the yield strength R is... p0.2 The strength ranges from 763 to 817 MPa (all ≥760 MPa); the strength index is basically equivalent to that of comparative examples 1-3. The elongation at room temperature (A) was 21.0%–25.0%, and the reduction of area (Z) was 23.5%–32.5%. At 700℃, the elongation at room temperature (A) was 18.5%–28.5%, and the reduction of area (Z) was 24.0%–31.0%, indicating good overall plasticity. In particular, the room temperature impact toughness (KU2) reached 73–89 J (all ≥73 J), significantly higher than the 22–46 J of comparative examples 1-3.

[0075] Meanwhile, the primary carbide level of Examples 1-9 is controlled at level 4-5, which is far superior to level 7-8 of Comparative Examples 1-2. Compared with Comparative Example 3, the alloy of the present invention has better impact toughness and plasticity matching while maintaining a comparable degree of carbide refinement, and its comprehensive performance advantages are obvious.

[0076] The excellent balance between strength and toughness described above is attributed to the rational control of carbide morphology in this invention. Examples 2 and 4 are provided (see...). Figure 3 , Figure 4 The primary carbides are fine granular and dispersed throughout the matrix, without obvious agglomeration or continuous precipitation along grain boundaries; the morphology of the carbides in other embodiments is basically consistent with this (figures omitted). In contrast, the carbides in Comparative Examples 1-3 are irregularly blocky and aggregated along grain boundaries (see Figure 1). Figure 5 (Taking Comparative Example 1 as an example), coarse carbides become the source of crack initiation, leading to a sharp deterioration in toughness.

[0077] The above results show that the present invention successfully achieves the refinement and dispersion of carbides, and significantly improves room temperature impact toughness without sacrificing high temperature strength, thus solving the technical problem of achieving both high strength and high toughness in nickel-based heat-resistant alloys.

[0078] Based on the above room temperature and high temperature tensile and impact tests, the present invention further evaluated the high temperature creep performance of alloys of Examples 1-9 and Comparative Examples 1-3, and the results are shown in Table 4.

[0079] Table 4

[0080] As shown in Table 4, under conditions of 700℃ / 585MPa, the fracture times of alloys in Examples 1-9 of this invention ranged from 856.7h to 1208.3h (all ≥850h), while that of Comparative Example 3 (Nimonic 105) was only 678.3h. Under conditions of 700℃ / 500MPa, the fracture times of alloys in Examples 1-9 ranged from 2436.5h to 3121.1h (all ≥2435h), while that of Comparative Example 3 was only 2205.5h. This indicates that the alloys of this invention exhibit excellent creep rupture performance under high temperature and high stress conditions. Combined with the tensile and impact data in Table 3, it can be seen that the alloys of this invention achieve a significant improvement in toughness and creep rupture life without reducing strength, and their overall mechanical properties are significantly superior to those of the comparative examples.

[0081] Figure 1 and Figure 2 The images show transmission electron microscopy (TEM) images of the γ′ strengthening phase after the heat treatment of the alloys in Examples 2 and 4, respectively. The observations of other examples are basically consistent with these images (figures omitted). As can be seen from the figures, the γ′ strengthening phase exhibits a fine, dispersed spherical distribution with an average particle size of 10–50 nm. This fine, uniformly distributed spherical precipitate is geometrically symmetrical and has a low stress concentration factor, making it less likely to become a crack initiation point under high-temperature stress. This ensures that the alloy maintains good toughness and resistance to creep damage while achieving high strength.

[0082] In summary, this invention, through rational control of the proportions of key elements in the alloy and combined with optimized heat treatment processes, yielded a microstructure in which the γ′ strengthening phase is spherically and uniformly distributed (average particle size 10–50 nm) and the primary carbides are finely dispersed. This microstructure enables the alloy to exhibit excellent tensile strength and impact toughness at both room temperature and 700℃. Furthermore, it demonstrates creep rupture lifespans of ≥850 h at 700℃ / 585 MPa and ≥2435 h at 700℃ / 500 MPa, respectively. The overall mechanical properties are significantly superior to existing Nimonic 105 alloys, fully meeting the requirements for high-parameter turbine fasteners under high-temperature and high-pressure long-term service conditions, and possesses broad industrial application prospects.

[0083] The above description is only a preferred embodiment of the present invention, but the scope of protection of the present invention is not limited thereto. Any changes or substitutions that can be easily conceived by those skilled in the art within the scope of the technology disclosed in the present invention should be included within the scope of protection of the present invention.

Claims

1. A nickel-based heat-resistant alloy for high-parameter steam turbine fasteners, characterized in that, The chemical composition by mass percentage is as follows: C: 0.05~0.25%, Si: ≤0.10%, Mn: ≤0.10%, P: ≤0.010%, S: ≤0.010%, Cr: 13.0~17.0%, Co: 17.0~23.0%, Al: 3.0~5.0%, Ti: 1.60~3.00%, Mo: 3.0~6.0%, Zr: 0.03~0.12%, B: 0.005~0.01%, and satisfies the following conditions: 5.50%≤Al+Ti≤6.00%, 0.35≤Ti / Al≤0.90, with the balance being Ni and unavoidable impurities.

2. The nickel-based heat-resistant alloy according to claim 1, characterized in that, The Al+Ti content is 5.65%–5.85%.

3. The nickel-based heat-resistant alloy according to claim 1, characterized in that, The γ′ strengthening phase in the alloy microstructure is spherical with an average grain size of 10–50 nm.

4. The nickel-based heat-resistant alloy according to any one of claims 1 to 3, characterized in that, The room temperature impact toughness of the alloy is KU2≥73J.

5. The nickel-based heat-resistant alloy according to any one of claims 1 to 3, characterized in that, The fracture time of the alloy under 700℃ / 585MPa conditions is ≥850 hours.

6. A method for preparing a nickel-based heat-resistant alloy for high-parameter steam turbine fasteners according to any one of claims 1 to 5, characterized in that, Includes the following steps: S1: Vacuum induction melting is used to melt the proportioned raw materials and then cast them into steel ingots; S2: Forced cooling is adopted after casting, and the cooling rate of the steel ingot is controlled to be ≥10℃ / minute; S3: After demolding, the steel ingot is annealed and then forged. The forging opening temperature is 1150~1170℃ and the final forging temperature is ≥1000℃ to obtain the billet. S4: Heat treatment is performed on the forged billet.

7. The preparation method according to claim 6, characterized in that, In step S1, during the refining stage of vacuum induction melting, the melting power is controlled at 600-800W.

8. The preparation method according to claim 6, characterized in that, In step S1, Al is added by repeated immersion: the Al ingot is immersed 5-20 mm below the liquid surface to melt, then lifted out, and the operation is repeated until it is completely melted.

9. The preparation method according to claim 6, characterized in that, In step S2, the ingot is demolded when its surface cools to 300-350°C.

10. The preparation method according to claim 6, characterized in that, In step S4, the heat treatment includes: Solution treatment: Hold at 1100-1180℃ for 2-6 hours, then cool. Intermediate treatment: Hold at 1000~1060℃ for 12~20 hours, then cool. Aging treatment: Keep warm at 700-850℃ for 12-20 hours, then cool.

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