A nickel-based wrought superalloy GH4780 alloy and a preparation method thereof

Abstract

The present application relates to a kind of nickel-based wrought superalloy GH4780 alloy and its preparation method, belong to nickel-based wrought superalloy technical field, solve the lower 760 ℃ long-time creep life of existing GH4780 alloy and / or the problem that strength and plasticity balance cannot be realized.The nickel-based wrought superalloy GH4780 alloy, by weight percentage, the content of B is 0.008%~0.030%, the content of Zr is 0.08%~0.30%, and the mass ratio of Zr and B is 8~12:1.The GH4780 alloy has higher 760 ℃ long-time creep life, and has better strength and plasticity.

Description

Technical Field

[0001] This invention relates to the field of nickel-based wrought superalloy technology, and in particular to a nickel-based wrought superalloy GH4780 and its preparation method. Background Technology

[0002] Nickel-based wrought superalloy GH4780 has become a core candidate material for hot-end components of aero-engines and gas turbines due to its excellent hot working properties and the combination of strength and plasticity at medium and high temperatures.

[0003] In existing GH4780 alloys, the B content is usually controlled between 0.002% and 0.007%, and the Zr content is controlled between 0.005% and 0.07%. This GH4780 alloy has a low long-term creep life at 760℃ and cannot achieve a balance between strength and plasticity. Summary of the Invention

[0004] Based on the above analysis, the present invention aims to provide a nickel-based wrought superalloy GH4780 and its preparation method to solve the problems of low creep life at 760℃ and / or inability to achieve a balance between strength and plasticity in existing GH4780 alloys.

[0005] On one hand, the present invention provides a nickel-based wrought superalloy GH4780 alloy, wherein the chemical composition of the nickel-based wrought superalloy GH4780 alloy, by weight percentage, comprises: C: 0.04%–0.06%, Cr: 21.5%–22.0%, Co: 18.0%–18.8%, W: 1.7%–1.85%, Nb: 0.60%–0.70%, Al: 1.0%–1.2%, Ti: 2.0%–2.2%, Ta: 0.9%–1.5%, B: 0.008%–0.030%, Zr: 0.08%–0.30%, Si: ≤0.08%, S: ≤10ppm, P: ≤40ppm, O+N+S: ≤50ppm, and Ni as the balance; The mass ratio of Zr to B is 8 to 12:1.

[0006] Preferably, the nickel-based wrought superalloy GH4780 has the following chemical composition by weight percentage: C: 0.045%–0.058%, Cr: 21.8%–21.95%, Co: 18.2%–18.7%, W: 1.74%–1.81%, Nb: 0.63%–0.68%, Al: 1.05%–1.15%, Ti: 2.04%–2.12%, Ta: 0.95%–1.21%, B: 0.012%–0.025%, Zr: 0.12%–0.25%, Si: ≤0.05%, S: ≤5ppm, P: ≤30ppm, O+N+S: ≤30ppm, with Ni as the balance; And / or, the mass ratio of Zr to B is 9 to 11:1.

[0007] Secondly, the present invention provides a method for preparing a nickel-based wrought superalloy GH4780 alloy, used to prepare the aforementioned nickel-based wrought superalloy GH4780 alloy, the preparation method comprising: Step (1): First, nickel, cobalt, tungsten, niobium, chromium, tantalum and nickel-carbon master alloy are loaded into a vacuum melting furnace for vacuum induction melting. After refining, aluminum and titanium are added in batches. Finally, at the end of the vacuum induction melting and at a superheat of 80-120°C, nickel-boron master alloy, zirconium and nickel-carbon master alloy for carbon adjustment are added and electromagnetically stirred for 3-5 minutes. After pouring, the ingot is demolded to obtain a vacuum induction casting. Step (2) involves electroslag remelting the vacuum induction casting ingot, wherein the electroslag remelting uses an alkaline pre-melted slag system; Step (3) involves vacuum consumable arc remelting of the ingot; Step (4) involves subjecting the ingot to a three-stage gradient homogenization annealing process; Step (5): The ingot is subjected to controlled heat treatment to obtain a billet; Step (6): The billet is subjected to multiple temperature-controlled rolling passes; Step (7) involves subjecting the billet to solution heat treatment and aging heat treatment; The three-stage gradient homogenization annealing process includes: Step (4-1): Increase the temperature from room temperature to 1100-1130℃ at a rate of 8-12℃ / min and hold for 20-30 hours; Step (4-2): Increase the temperature to 1150-1170℃ at a rate of 0.5-1℃ / min and hold for 30-40 hours; Step (4-3): Heat to 1180-1195℃ at a rate of 0.5-1℃ / min and hold for 45-55 hours; after holding, cool slowly to 900-1000℃ at a rate of 0.5-1℃ / min and remove from the furnace.

[0008] Preferably, in step (1), the vacuum degree of the vacuum induction melting is ≤5×10⁻⁶. -2 Pa, the refining temperature is 1500~1540℃, and the refining time is 10~15min.

[0009] Preferably, in step (1), the pouring temperature is 1460-1500℃, and the mold is kept warm for 2-3 hours after pouring.

[0010] Preferably, in step (2), the alkaline pre-melted slag comprises, by mass percentage: CaF2: 55%–60%, Al2O3: 20%–25%, CaO: 15%–20%, and MgO: 3%–5%.

[0011] Preferably, in step (2), the melting rate of the electroslag remelting is 3.5 to 4.0 kg / min, the water temperature is 28 to 32°C, and high-purity argon gas is introduced for protection throughout the process, with an argon gas flow rate ≥ 30 L / min.

[0012] Preferably, in step (3), during the vacuum self-consuming arc remelting, when the vacuum degree is <1×10 -3 Arc initiation occurs when Pa and leakage rate is <0.3 Pa·L / s, and melting rate is 3.4~3.8 kg / min.

[0013] Preferably, in step (5), the controllable hot working blanking includes sequentially performing rapid forging and radial forging.

[0014] Preferably, in step (7), the solution heat treatment includes: heating to 1070-1090°C at a rate of 5-8°C / min, holding at that temperature for 1-2 hours, and then air cooling; And / or, the aging heat treatment includes: heating to 790-810°C at a rate of 3-5°C / min, holding at that temperature for 6-8 hours, and then air cooling.

[0015] Compared with the prior art, the present invention can achieve at least one of the following beneficial effects: 1. In the nickel-based wrought superalloy GH4780 of the present invention, the B content is 0.008% to 0.030%, and the Zr content is 0.08% to 0.30%. B preferentially agglomerates at grain boundaries, reduces the grain boundary diffusion rate, and forms dispersed boron carbides with C and Cr to pin the grain boundaries. The large atomic radius Zr fills the grain boundary distortion region, reduces the grain boundary energy, inhibits the growth of harmful phase nuclei, and stabilizes the grain boundary MC carbides. Under the above-mentioned high addition amount, the agglomeration of B and Zr at the grain boundaries is sufficient to form continuous and stable M. 23 The (C,B)6 boron carbide pinning network and MC stabilizing network improve the long-term creep life at 760℃, solving the problem that the existing GH4780 alloy has insufficient grain boundary segregation and discontinuous pinning network due to low B and Zr content, which leads to a low long-term creep life at 760℃.

[0016] 2. In the nickel-based wrought high-temperature alloy GH4780 of the present invention, the mass ratio of Zr to B is 8-12:1 (preferably 10:1). Under this ratio, the M formed by B... 23(C,B)6 and Zr-stabilized MC form a synergistic dispersion structure of "boron carbide pinning + MC stabilization" at the grain boundaries, which ensures pinning strength and avoids continuous embrittlement. This avoids the problems of excessive coarsening and embrittlement of grain boundaries caused by boride when the B content is relatively high (Zr to B mass ratio is less than 8:1), and continuous embrittlement and hot working cracking caused by the enrichment of Zr-rich low-melting-point eutectic phase when the Zr content is relatively high (Zr to B mass ratio is greater than 12:1). Thus, it maintains excellent plasticity even at high addition levels, achieving the strengthening effect of "grain boundary purification + grain boundary pinning", suppressing grain boundary sliding during deformation, and achieving a balance between alloy strength and plasticity.

[0017] 3. In the nickel-based wrought high-temperature alloy GH4780 of the present invention, the B content is 0.008% to 0.030%, the Zr content is 0.08% to 0.30%, and the mass ratio of Zr to B is 8 to 12:1. Both B and Zr contents are relatively high, and the Zr content is greater than the B content. On the one hand, the grain boundary sliding and crack initiation during high-temperature creep can be suppressed by the synergistic segregation and dispersed precipitate pinning of B-Zr grain boundaries. On the other hand, the excess Zr after meeting the grain boundary segregation requirements enters the γ′ phase and MC carbides, and together with Al, Ti, Ta, and Nb, ensures and strengthens the precipitation strengthening effect of the matrix γ′ phase, realizing the dual synergy of grain boundary strengthening and γ′ phase precipitation strengthening, which greatly improves the creep life of the alloy at 760℃ / 460MPa, while maintaining excellent plasticity and meeting the long-term service requirements of hot-end components of aero-engines.

[0018] 4. In the preparation method of this invention, to address the problems of increased Zr-rich low-melting-point eutectic phase and decreased solidus line caused by high Zr content, a three-stage gradient homogenization annealing process is adopted to eliminate high Zr dendrite segregation without overheating. Combined with the timing of B and Zr addition during the melting process (adding B and Zr at the end of vacuum induction melting) and low superheat control, the problems of easy burn-off, segregation, and formation of low-melting-point eutectic phases with high B and Zr content are solved. This preparation method is compatible with existing nickel-based superalloy industrial production lines, requires no additional specialized equipment, and produces alloy bars with uniform grain size (grade difference < 3), excellent batch stability, and can be directly used for mass production of key components for aero-engines.

[0019] 5. The GH4780 alloy of the present invention combines excellent grain boundary stability, high-temperature creep performance and strong plasticity. At room temperature (23℃), the tensile strength is ≥1238MPa, the room temperature yield strength is ≥783MPa and the room temperature elongation is ≥26.5%; at 760℃, the tensile strength is ≥854MPa, the yield strength is ≥703MPa and the elongation is ≥13.8%; at 760℃ / 460MPa, the creep life is ≥70.0h and the creep elongation is ≥12.5%. It can be used not only for aero-engine turbine disks and casings, but also for aero-engine fasteners, industrial gas turbine combustion chamber components and other applications with high requirements for grain boundary performance.

[0020] 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 objectives and other advantages of this invention can be realized and obtained through the specific points highlighted in the description. Attached Figure Description

[0021] Figure 1 The grain boundary M of the nickel-based wrought high-temperature alloy GH4780 in Example 3 of this invention. 23 Scanning electron microscope (SEM) images of the dispersed distribution morphology of (C,B)6 boron carbides; Figure 2 For the grain boundary M of the alloy in Comparative Example 5 23 (C,B) Scanning electron microscope (SEM) image of (C,B) boron carbides with excessive coarsening and a continuous film-like distribution morphology; Figure 3 This is a scanning electron microscope (SEM) image of the γ′ phase morphology in the matrix of the nickel-based wrought high-temperature alloy GH4780 after electrolytic corrosion, as described in Example 3 of the present invention. Figure 4 The image shows the metallographic structure of the nickel-based wrought high-temperature alloy GH4780 after solution treatment and aging treatment in Example 3 of this invention. Detailed Implementation

[0022] The preferred embodiments of the present invention are described in detail below. These embodiments are used to illustrate the principles of the present invention and are not intended to limit the scope of the present invention.

[0023] Boron (B) and zinc (Zr) are grain boundary strengthening elements in nickel-based superalloys, achieving grain boundary strengthening through grain boundary segregation and dispersed precipitate pinning. Studies have found that the addition amounts of B and Zr in existing GH4780 alloy systems are generally too low, and the ratio of the two is not constrained, resulting in the following technical defects: (1) The existing GH4780 alloy usually has a B content of 0.002% to 0.007% and a Zr content of 0.005% to 0.07%. Within this low content range, the grain boundary segregation of B and Zr is insufficient, making it difficult to form continuous and stable boride pinning at the grain boundaries. The grain boundary strengthening and purification effects are limited, and the long-term creep life of the alloy at 760℃ is relatively low. (2) Existing technologies only specify the content range of B and Zr in isolation, without recognizing that there is an optimal ratio for the synergistic segregation of B and Zr. When the ratio of B to Zr is unbalanced, either excessive precipitation of borides at grain boundaries leads to embrittlement (when B is relatively too high), or enrichment of low-melting-point eutectic phases rich in Zr leads to continuous embrittlement of grain boundaries and hot working cracking (when Zr is relatively too high). Neither of these can achieve a balance between strength and plasticity.

[0024] (3) The existing preparation method of GH4780 alloy is only applicable to alloys with low B and Zr content. It has not been optimized for the metallurgical characteristics of B and Zr elements being easy to burn off, segregate, and form low melting point eutectic phases under high B and high Zr content. Using the existing preparation method of GH4780 alloy to prepare alloys with high B and high Zr will result in large fluctuations in the composition of alloy components, severe dendrite segregation, and insufficient batch performance stability, which cannot meet the requirements of industrial mass production of key components of aero-engines.

[0025] Based on this, in a first aspect, the present invention provides a nickel-based wrought superalloy GH4780 alloy, wherein the chemical composition of the nickel-based wrought superalloy GH4780 alloy, by weight percentage, comprises: C: 0.04%–0.06%, Cr: 21.5%–22.0%, Co: 18.0%–18.8%, W: 1.7%–1.85%, Nb: 0.60%–0.70%, Al: 1.0%–1.2%, Ti: 2.0%–2.2%, Ta: 0.9%–1.5%, B: 0.008%–0.030%, Zr: 0.08%–0.30%, Si: ≤0.08%, S: ≤10ppm, P: ≤40ppm, O+N+S: ≤50ppm, and Ni as the balance; The mass ratio of Zr to B is 8 to 12:1.

[0026] Compared with the prior art, the nickel-based wrought superalloy GH4780 of the present invention, by limiting the B and Zr content to a relatively high level, ensures that the segregation of B and Zr at the grain boundaries is sufficient to form a continuous and stable M. 23 (C,B)6 boron carbide pinning network and MC stabilizing network improve long-term creep life at 760℃; by limiting the mass ratio of Zr and B synergistic segregation, this invention avoids grain boundary embrittlement caused by relative excess of a single element, and achieves a balance between grain boundary strengthening and plasticity.

[0027] The design basis of each component and the role of key elements in the nickel-based wrought superalloy GH4780 of this invention are explained as follows: Element B (core grain boundary strengthening element): Element B preferentially segregates at grain boundaries, fills grain boundary vacancies, reduces grain boundary diffusion rate, and inhibits continuous precipitation and coarsening of grain boundary carbides; it combines with grain boundary C and Cr to form dispersed M. 23 (C,B)6 boron carbides pin grain boundaries, preventing grain boundary slippage during high-temperature deformation and significantly improving high-temperature creep life. The content of boron is 0.008% to 0.030%, for example, 0.008%, 0.010%, 0.015%, 0.020%, 0.025%, and 0.030%. The content control is based on the following: when B < 0.008%, the grain boundary segregation is insufficient, the pinning network is discontinuous, and the grain boundary strengthening effect is limited; when B > 0.030%, continuous coarse boron carbides form at the grain boundaries, leading to embrittlement and a sharp decrease in high-temperature plasticity.

[0028] Zr (a core grain boundary purification and strengthening element): It synergistically agglomerates with boron at grain boundaries, filling grain boundary distortion regions with its large atomic radius, reducing grain boundary energy, inhibiting the nucleation and growth of harmful phases, improving grain boundary bonding strength, and suppressing creep crack initiation. Excess Zr enters the γ′ phase and carbides, acting as a precipitation strengthening agent. The Zr content is 0.08%–0.30%, for example, 0.08%, 0.10%, 0.15%, 0.20%, 0.25%, and 0.30%. The content control is based on the following: when Zr < 0.08%, grain boundary purification and strengthening are insufficient; when Zr > 0.30%, coarse, Zr-rich, low-melting-point eutectic phases and other brittle phases form at grain boundaries, causing continuous embrittlement, a sharp decrease in plasticity and creep life, and an increased risk of hot working cracking.

[0029] C element: synergistically regulates grain boundary precipitation with B and Zr, with a content of 0.04%–0.06%, for example 0.04%, 0.05%, and 0.06%. C element forms an appropriate amount of MC carbide to pin grain boundaries and provides M for grain boundaries. 23 (C,B)6 precipitation provides C, avoiding excessive C which leads to continuous precipitation and embrittlement of grain boundary carbides.

[0030] Al+Ti elements: The main forming elements of the γ′ phase, with Al content of 1.0% to 1.2% (e.g., 1.0%, 1.1%, 1.2%) and Ti content of 2.0% to 2.2% (e.g., 2.0%, 2.1%, 2.2%). The total Al+Ti content is controlled at 3.0% to 3.4% (3.0%, 3.1%, 3.2%, 3.3%, 3.4%) to ensure the precipitation amount and thermal stability of the γ′ phase, and to work synergistically with B-Zr grain boundary strengthening to avoid excessive Al+Ti content leading to η phase precipitation.

[0031] Ta and Nb elements: Ta content is 0.9% to 1.5% (e.g., 0.9%, 1.0%, 1.2%, 1.4%, 1.5%), and Nb content is 0.60% to 0.70% (e.g., 0.60%, 0.65%, 0.70%). Ta and Nb assist in solid solution strengthening and stabilization of the γ′ phase, improve the thermal stability of the γ′ phase, and inhibit the coarsening of the γ′ phase under long-term heat exposure.

[0032] Cr, Co, and W elements: Cr ensures high-temperature oxidation resistance and corrosion resistance, with a content of 21.5% to 22.0% (e.g., 21.5%, 21.7%, 22.0%); Co reduces stacking fault energy and enhances solid solution strengthening, with a content of 18.0% to 18.8% (e.g., 18.0%, 18.2%, 18.4%, 18.6%, 18.8%); W assists in solid solution strengthening and optimizes high-temperature performance, with a content of 1.7% to 1.85% (e.g., 1.7%, 1.75%, 1.8%, 1.85%).

[0033] Impurity element control: Strictly control S≤10ppm, O+N+S≤50ppm, Si≤0.08%, P≤40ppm. S is a grain boundary weakening element. Strict control of impurities can maximize the grain boundary strengthening effect of B-Zr and avoid grain boundary embrittlement caused by impurities. Furthermore, for the alloys of this invention that rely on B and Zr for grain boundary strengthening, the control of Si and P is particularly crucial: Si and B exhibit the same tendency to segregate at grain boundaries, competing for grain boundary positions. First-principles calculations show that Si occupying grain boundary substitution sites reduces grain boundary dissociation energy and weakens grain boundary bonding, thereby offsetting the grain boundary strengthening effect brought about by B segregation. Experiments have also confirmed that even when the Si content is below the conventional commercial upper limit, its segregation towards grain boundaries still significantly reduces the creep plasticity improvement of B-containing alloys, and local Si enrichment lowers the local solidus temperature and induces initial melting. P also strongly segregates at grain boundaries and competes with B for grain boundary positions. When the P content is too high, it significantly weakens the grain boundary strengthening effect of B, while controlling P to a very low level allows the beneficial effect of B to extend to a higher content range. In addition, P segregation also lowers the initial melting temperature and promotes the precipitation of Laves phase and σ phase at grain boundaries. Therefore, to maximize the grain boundary strengthening effect of B and Zr, this invention controls Si at ≤0.08% and P at ≤40ppm.

[0034] More preferably, the chemical composition of the nickel-based wrought superalloy GH4780 alloy, by weight percentage, includes: C: 0.045%–0.058%, Cr: 21.8%–21.95%, Co: 18.2%–18.7%, W: 1.74%–1.81%, Nb: 0.63%–0.68%, Al: 1.05%–1.15%, Ti: 2.04%–2.12%, Ta: 0.95%–1.21%, B: 0.012%–0.025%, Zr: 0.12%–0.25%, Si: ≤0.05%, S: ≤5ppm, P: ≤30ppm, O+N+S: ≤30ppm, with Ni as the balance.

[0035] More preferably, the mass ratio of Zr to B is 9-11:1, and most preferably 10:1.

[0036] Secondly, in view of the metallurgical characteristics of B and Zr elements being prone to burn-off, segregation, and formation of Zr-rich low-melting-point eutectic phases under high B and high Zr compositions, this invention also provides a method for preparing a nickel-based wrought high-temperature alloy GH4780, used to prepare the above-mentioned nickel-based wrought high-temperature alloy GH4780, the preparation method comprising: Step (1): First, nickel, cobalt, tungsten, niobium, chromium, tantalum and nickel-carbon master alloy are loaded into a vacuum melting furnace for vacuum induction melting. After refining, aluminum and titanium are added in batches. Finally, at the end of the vacuum induction melting and at a superheat of 80-120℃ (e.g. 80℃, 90℃, 100℃, 110℃, 120℃), nickel-boron master alloy, zirconium and nickel-carbon master alloy for carbon adjustment are added. Electromagnetically stir for 3-5 minutes to ensure uniform composition and avoid burning and volatilization of B and Zr. After casting, the ingot is demolded to obtain a vacuum induction casting. Step (2) involves electroslag remelting the vacuum induction casting ingot, wherein the electroslag remelting uses an alkaline pre-melted slag system; Step (3) involves vacuum consumable arc remelting of the ingot; Step (4) involves subjecting the ingot to a three-stage gradient homogenization annealing process; Step (5): The ingot is subjected to controlled heat treatment to obtain a billet; Step (6): The billet is subjected to multiple temperature-controlled rolling passes; Step (7) involves subjecting the billet to solution heat treatment and aging heat treatment; The three-stage gradient homogenization annealing process includes: Step (4-1): Increase the temperature from room temperature to 1100-1130℃ at a rate of 8-12℃ / min and hold for 20-30 hours; Step (4-2): Increase the temperature to 1150-1170℃ at a rate of 0.5-1℃ / min and hold for 30-40 hours; Step (4-3): Heat to 1180-1195℃ at a rate of 0.5-1℃ / min and hold for 45-55 hours; after holding, cool slowly to 900-1000℃ at a rate of 0.5-1℃ / min and remove from the furnace.

[0037] Compared with existing technologies, this preparation method solves the problems of easy burn-off, easy segregation, and easy formation of low-melting-point eutectic phases of high-content B and Zr by limiting the timing of B and Zr addition during the melting process (B and Zr are added at the end of vacuum induction melting) and controlling the low superheat. The annealing process adopts a three-stage gradient homogenization annealing treatment to eliminate high Zr dendrite segregation without burning.

[0038] In step (1) of the present invention, since the Zr content is increased to 0.08% to 0.30%, the risk of Zr burn-off and volatilization increases. Therefore, nickel-boron master alloy and zirconium are added at the end of the smelting process and at low superheat, and the electromagnetic stirring time is extended to 3 to 5 minutes to ensure that the high content of B and Zr is evenly distributed and the burn-off is controllable.

[0039] For example, in step (1), after the furnace charge is fully melted and the surface of the melt is calm, the refining period begins. The refining process is carried out using intermittent electromagnetic stirring. The refining includes: electromagnetic stirring for 5 to 10 minutes after the furnace charge is completely melted, and stirring under a vacuum degree ≤ 5 × 10⁻⁶. - Refining at 1500–1540℃ for 10–15 minutes under 2 Pa.

[0040] For example, in step (1), the vacuum degree of the vacuum induction melting is ≤5×10 - ²Pa (e.g., 1×10) - ²Pa, 2×10 - ²Pa, 3×10 - ²Pa, 4×10 - ²Pa, 5×10 - The refining temperature is 1500–1540℃ (e.g., 1500℃, 1510℃, 1520℃, 1530℃, 1540℃), and the refining time is 10–15 min (10 min, 12 min, 14 min, 15 min). If the temperature is too low or the time is too short, the high-melting-point refractory elements will not dissolve completely, resulting in defects and excessive oxygen and sulfur impurities; if the temperature is too high or the time is too long, active and volatile elements will be easily burned off, and the solidification structure will deteriorate.

[0041] For example, in step (1), aluminum and titanium are added in batches after refining, and finally, at the end of vacuum induction melting and at a superheat of 80-120°C, nickel-boron master alloy, zirconium, and nickel-carbon master alloy for carbon adjustment are added. Specifically, this includes: after refining, cooling to 1450-1480°C, adding aluminum in batches and stirring for 1-2 minutes, then adding titanium in batches and stirring for 2-3 minutes; then cooling to 1440-1470°C, i.e., a superheat of 80-120°C, and refilling with high-purity argon gas to 5×10 3 ~2×10 4 After Pa, nickel-boron master alloy, zirconium, and nickel-carbon master alloy for carbon adjustment are added. The total amount of aluminum and titanium added satisfies the requirement that the total mass content of Al and Ti is 3.0% to 3.4% (e.g., 3.0%, 3.1%, 3.2%, 3.3%, 3.4%).

[0042] For example, in step (1), the pouring temperature is 1460-1500℃ (1460℃, 1470℃, 1480℃, 1490℃, 1500℃), and the mold is kept warm for 2-3 hours after pouring (e.g., 2 hours, 2.5 hours, 3 hours).

[0043] In step (2) of this invention, the electroslag remelting uses an alkaline pre-melted slag system to avoid the loss of B and Zr due to acidic slag. For example, the alkaline pre-melted slag system comprises, by mass percentage: CaF2: 55%–60% (e.g., 55%, 57%, 59%, 60%), Al2O3: 20%–25% (e.g., 20%, 22%, 24%, 25%), CaO: 15%–20% (e.g., 15%, 17%, 19%, 20%), and MgO: 3%–5% (e.g., 3%, 4%, 5%).

[0044] For example, in step (2), the electroslag remelting melting rate is 3.5–4.0 kg / min (e.g., 3.5 kg / min, 3.7 kg / min, 4.0 kg / min), the water temperature is 28–32°C (e.g., 28°C, 30°C, 31°C, 32°C), and high-purity argon gas is introduced throughout the process for protection, with an argon gas flow rate ≥30 L / min. This electroslag remelting can effectively remove sulfur and reduce the oxidation loss of boron and zirconium.

[0045] For example, step (2) further includes: turning the ingot obtained by electroslag remelting with a single-sided turning amount of 5-8 mm to remove surface defects and composition segregation layer.

[0046] For example, in step (3), during the vacuum self-consuming arc remelting, when the vacuum degree is <1×10 -Arc initiation occurs when the pressure is ³Pa and the leakage rate is <0.3Pa·L / s, with a melting rate of 3.4~3.8kg / min (e.g., 3.4kg / min, 3.5kg / min, 3.6kg / min, 3.7kg / min, 3.8kg / min).

[0047] For example, in step (3), the vacuum self-consuming arc remelting process uses a low-current arc stabilization coil with a current of 9-10A to avoid local segregation of high Zr content.

[0048] For example, in step (3), the vacuum consumable arc remelting process employs a hot capping process and helium gradient cooling at the end of the melting process to control the segregation of B and Zr dendrites during solidification, resulting in a high-purity, low-segregation ingot. Specifically, the hot capping process refers to gradually reducing the melting current and melting rate at the end of the vacuum consumable arc remelting process and extending the arc-ending and shrinkage compensation time, so that the molten pool smoothly transitions from steady-state solidification to transient solidification, and the top of the ingot solidifies slowly and sequentially and obtains sufficient shrinkage compensation, reducing end-tube shrinkage, porosity, and top macroscopic segregation; the helium gradient cooling process refers to filling the gap between the ingot and the crystallizer with high-purity helium during the remelting process, utilizing the characteristic that the thermal conductivity of helium is much higher than that of the vacuum gap to enhance lateral heat dissipation. Because this invention has high B and Zr content and a strong tendency for solidification segregation, the use of hot capping combined with helium gradient cooling can shorten the enrichment path and time of the solute in the pasty region, thereby suppressing the dendritic segregation of B and Zr and obtaining a high-purity, low-segregation ingot. The specific operating steps of the hot capping process and the helium gradient cooling are conventional operations in the field and will not be described in detail here.

[0049] In step (4) of the present invention, since the Zr content of the GH4780 alloy of the present invention is 0.08% to 0.30% compared with the existing low Zr content GH4780 alloy, the number of Zr-rich low-melting-point eutectic phases between dendrites increases significantly during solidification, and the solidus line (initial melting temperature) of the alloy decreases accordingly. If conventional two-stage homogenization annealing treatment is adopted (heating at 10 to 15℃ / min and holding at 1200 to 1220℃), there are two major risks: (1) the melting point of the low-melting-point eutectic phase is exceeded too quickly, and the grain boundary remelting occurs before it can dissolve, forming voids; (2) the temperature of the high-temperature section is too high, close to the reduced solidus line, and it is easy to overheat and crack. Therefore, this invention employs a three-stage gradient homogenization annealing process, adding a low-temperature eutectic step, lowering the peak temperature of the high-temperature stage, and extending the high-temperature diffusion time. Simultaneously, given that the B content is increased to 0.008%–0.030% in this invention, the aforementioned three-stage gradient homogenization, while extending the high-temperature diffusion time, also promotes the full diffusion and uniform distribution of B elements segregated at grain boundaries and between dendrites, avoiding localized boride aggregation and embrittlement. Specifically, this includes: Step (4-1): Increase the temperature from room temperature to 1100–1130°C (e.g., 1100°C, 1120°C, 1130°C) at a rate of 8–12°C / min (e.g., 8°C / min, 9°C / min, 10°C / min, 11°C / min, 12°C / min), and hold at this temperature for 20–30 h (e.g., 20 h, 22 h, 24 h, 26 h, 28 h, 30 h). This allows the Zr-rich low-melting-point eutectic phase between dendrites to fully dissolve below its melting point, preventing the eutectic phase from liquefying and forming voids during subsequent heating, and eliminating the η phase and the low-melting-point phase. Step (4-2): Increase the temperature to 1150–1170℃ (e.g., 1150℃, 1160℃, 1170℃) at a rate of 0.5–1℃ / min (e.g., 0.5℃ / min, 0.6℃ / min, 0.7℃ / min, 0.8℃ / min, 0.9℃ / min, 1℃ / min), and hold for 30–40 h (e.g., 30 h, 32 h, 34 h, 36 h, 38 h, 40 h) to further eliminate the eutectic phase and allow the B and Zr enriched at the grain boundaries to diffuse initially. Step (4-3): Increase the temperature to 1180–1195℃ (e.g., 1180℃, 1190℃, 1195℃) at a rate of 0.5–1℃ / min (e.g., 0.5℃ / min, 0.6℃ / min, 0.7℃ / min, 0.8℃ / min, 0.9℃ / min, 1℃ / min), and hold for 45–55h (e.g., 45h, 47h, 49h, 51h, 53h, 55h) to eliminate component segregation and control the residual elemental segregation coefficient below 0.2. After holding, slowly cool to 900–1000℃ at a rate of 0.5–1℃ / min (e.g., 0.5℃ / min, 0.6℃ / min, 0.7℃ / min, 0.8℃ / min, 0.9℃ / min, 1℃ / min) to avoid rapid cooling that could cause thermal stress cracks, and also to prevent premature segregation of B and Zr grain boundaries during the cooling process.

[0050] For example, in step (5), the controllable hot working blanking includes sequentially performing rapid forging and radial forging.

[0051] For example, the controllable hot working forging includes: heating the ingot to 1150-1190℃ (e.g., 1150℃, 1160℃, 1170℃, 1180℃, 1190℃) and holding it at that temperature for 5-8 hours (e.g., 5 hours, 6 hours, 7 hours, 8 hours); performing three consecutive upsetting and drawing operations on a high-speed forging mill, with each upsetting deformation amounting to 35%-45% (e.g., 35%, 37%, 40%, 43%, 45%) and each drawing deformation amounting to 40%-50% (40%, 42%, 44%, 46%, 48%, 50%), each heating temperature decreasing by 40℃ compared to the previous operation, and the final forging temperature not lower than 1050℃, and drawing the ingot to Φ110-Φ16 on the high-speed forging mill. 0mm intermediate billet; then heat the intermediate billet to 1080~1120℃ (e.g. 1080℃, 1090℃, 1100℃, 1110℃, 1120℃), hold for 1~2h, and perform 3~5 passes of hot radial forging on a radial forging machine, with a forging frequency of 80~120 times / min, a single pass area reduction rate of 12%~20% in the main deformation pass, and the billet rotates 15°~30° after each forging. The last pass is the shaping pass, with a single pass area reduction rate of 5%~8%, a rotation angle of 15°~20°, a total radial forging deformation of 35%~50%, and a final forging temperature of not less than 1000℃. Prepare Φ80~Φ120mm homogeneous bar billets with a grain size difference of <3 levels by multiple passes. It should be noted that the upsetting deformation and drawing deformation are process control values ​​for each heat treatment, and the intermediate billet size is the actual measured size after rapid forging; the total radial forging deformation is the cumulative reduction rate of the cross-sectional area of ​​the billet before and after radial forging, and is calculated based on the actual measured cross-sectional area of ​​the billet before and after radial forging. The total radial forging deformation η is calculated as η = (1 - A) / ( ... n Calculated by multiplying A0 by 100%, where A0 is the measured cross-sectional area of ​​the intermediate billet before radial forging, and A n The measured cross-sectional area of ​​the forged bar billet is given; for a circular cross-section, η = [1 - (d...]. n / d0) 2 ]×100%, where d0 is the measured cross-sectional diameter of the intermediate billet before radial forging, d n This represents the measured cross-sectional diameter of the forged bar billet.

[0052] For example, in step (6), the multi-pass temperature-controlled rolling includes: heating to 1130-1170℃ (e.g., 1130℃, 1140℃, 1150℃, 1160℃, 1170℃) at a rate of 5-10℃ / min (e.g., 5℃ / min, 6℃ / min, 7℃ / min, 8℃ / min, 9℃ / min, 10℃ / min), and holding at that temperature for 2-3 hours (e.g., 2 hours). h, 2.5h, 3h); perform 8 to 10 passes (e.g., 8 passes, 9 passes, 10 passes) of continuous temperature-controlled rolling, with an initial rolling temperature of 1090 to 1140℃ (e.g., 1090℃, 1100℃, 1110℃, 1120℃, 1130℃, 1140℃) and a final rolling temperature of 960 to 1040℃ (e.g., 960℃, 980℃, 1000℃, 1020℃, 1040℃).

[0053] For example, in step (6), the area reduction rate per pass is 15% to 25% (e.g., 15%, 17%, 20%, 23%, 25%). Based on the area reduction rates for 8 to 10 passes and the following stages, the total rolling deformation is 79% to 92%; the total rolling deformation is the cumulative reduction rate η of the cross-sectional area of ​​the billet before and after rolling, calculated as η = (1... A n Calculated as (A0)×100%, where A0 is the cross-sectional area of ​​the billet before rolling, A n This refers to the cross-sectional area of ​​the billet after the final rolling pass; when the area reduction ratios for each pass are r1, r2, ..., r... n When, η=[1 (1 r1)(1 r2)…(1 r n )]×100%. Immediately after rolling, recrystallize and anneal at 960~1000℃ (e.g. 960℃, 970℃, 980℃, 990℃, 1000℃) for 30~50min (e.g. 30min, 35min, 40min, 45min, 50min), and air cool to room temperature.

[0054] For example, in step (6), the area reduction rate of the first three passes is 20% to 25% (e.g., 20%, 22%, 24%, 25%) to break up coarse grains, the area reduction rate of the middle passes is 18% to 22% (e.g., 18%, 19%, 20%, 22%), and the area reduction rate of the last two passes is 15% to 18% (e.g., 15%, 16%, 17%, 18%).

[0055] For example, in step (6), a 280-type transverse rolling mill is used for multi-pass temperature-controlled rolling. Before the rolling temperature rises, the billet is sawn to a fixed length and the surface is ground around the whole circumference. The single-sided turning amount is 0.3 to 0.6 mm.

[0056] For example, in step (7), the solution heat treatment includes: heating at 5-8℃ / min (e.g., 5℃ / min, 6℃ / min, 7℃ / min, 8℃ / min) to 1070-1090℃ (e.g., 1070℃, 1080℃, 1090℃), holding at that temperature for 1-2h (e.g., 1h, 1.5h, 2h), and air cooling to achieve partial re-dissolution of MC carbides, controlling the matrix grain size to 100-200μm, so that B and Zr are uniformly dissolved.

[0057] For example, the aging heat treatment includes: heating to 790–810°C (e.g., 790°C, 800°C, 810°C) at a rate of 3–5°C / min (e.g., 3°C / min, 4°C / min, 5°C / min), holding at that temperature for 6–8 hours (e.g., 6 hours, 7 hours, 8 hours), and air cooling, to promote the segregation of B and Zr towards the grain boundaries to form dispersed M. 23 (C,B)6 boron carbide pins grain boundaries while simultaneously precipitating nanoscale γ′ phase, achieving a synergistic effect of grain boundary strengthening and precipitation strengthening.

[0058] The present invention will be further illustrated below with specific embodiments of the nickel-based wrought high-temperature alloy GH4780 and its preparation method.

[0059] Example 1 This embodiment provides a nickel-based wrought superalloy GH4780 and its preparation method.

[0060] The chemical composition of the nickel-based wrought superalloy GH4780, by weight percentage, includes: C: 0.054%, Si: 0.041%, B: 0.012%, Zr: 0.12% (Zr to B mass ratio is 10:1), Al: 1.08%, Ta: 0.95%, Ti: 2.04%, W: 1.74%, Nb: 0.63%, Co: 18.56%, Cr: 21.85%, N: 0.0004%, O: 0.0011%, S: 0.0004%, P: 0.0023%, with the balance being Ni.

[0061] The preparation method is as follows: Step (1): Nickel, cobalt, tungsten, niobium, chromium, tantalum, and nickel-carbon master alloy are charged into a vacuum induction furnace for vacuum induction melting. After the furnace charge is completely melted, it is electromagnetically stirred for 5 minutes at a vacuum degree of 3×10⁻⁶. -Refining at 1500℃ and 2Pa for 10 minutes; after refining, cooling to 1470℃, adding aluminum in batches and stirring for 1 minute, then adding titanium in batches and stirring for 2 minutes; then cooling to 1450℃, i.e., superheat of 90℃, and refilling with high-purity argon to 5×10³Pa, adding nickel-boron master alloy, zirconium, and nickel-carbon master alloy for carbon adjustment, electromagnetic stirring for 3 minutes to ensure uniform composition, and starting pouring at 1460℃ within 5 minutes after stirring, keeping the mold at the temperature for 2 hours after pouring, and demolding to obtain vacuum induction casting ingot; Step (2): The vacuum induction casting ingot is electroslag remelted using an alkaline pre-melted slag system. The alkaline pre-melted slag system comprises, by mass percentage: CaF2: 60%, Al2O3: 20%, CaO: 16%, MgO: 4%. The electroslag remelting rate is 3.5 kg / min, the water temperature is 28°C, and high-purity argon gas is introduced throughout the process for protection. The argon gas flow rate is 30 L / min. The ingot obtained by electroslag remelting is machined with a single-sided machining allowance of 5 mm to remove surface defects and compositional segregation layers. Step (3) involves vacuum arc remelting of the ingot. When the vacuum degree is <1×10 - Arc initiation occurs when the pressure is 3 Pa and the leakage rate is <0.3 Pa·L / s, the melting rate is 3.4 kg / min, the arc is stabilized by low current throughout the process, and the current of the arc stabilizing coil is 9 A; the vacuum self-consuming arc remelting adopts a hot capping process and helium gradient cooling at the end of the melting process. Step (4) involves subjecting the ingot to a three-stage gradient homogenization annealing process; the three-stage gradient homogenization annealing process includes: Step (4-1): Increase the temperature from room temperature to 1120℃ at a rate of 10℃ / min and hold for 25 hours; Step (4-2): Increase the temperature to 1160℃ at a rate of 0.8℃ / min and hold for 35 hours; Step (4-3): Heat to 1190℃ at a rate of 0.8℃ / min and hold for 50 hours; after holding, cool slowly to 1000℃ at a rate of 0.8℃ / min and remove from the furnace. Step (5): Controllable hot working of the ingot: Heat the ingot to 1180℃ and hold for 6 hours. Then, perform three upsetting and drawing operations on a high-speed forging mill. Each upsetting deformation is 40% and the drawing deformation is 45%. The temperature is reduced by 40℃ each time. The final forging temperature is 1060℃. The ingot is drawn to a Φ125mm intermediate billet on the high-speed forging mill. The intermediate billet is then heated to 1080℃ and held for 1 hour. Then, it is hot-forged in three passes on a radial forging mill. The forging frequency is 80 times / min. The single-pass area reduction rate of the main deformation pass is 17%. The billet rotates 15° after each forging. The last pass is the shaping pass. The single-pass area reduction rate is 7.1%. The rotation angle is 15°. The total radial forging deformation is 36.00%. The final forging temperature is 1000℃. The billet is radially forged to a Φ100mm bar billet with a grain size difference of <3 (grain size rating follows ASTM E112 standard). Step (6): The billet is subjected to multi-pass temperature-controlled rolling: a 280-type transverse rolling mill is used, with a single-sided grinding amount of 0.5 mm. It is heated to 1150℃ at a rate of 8℃ / min and held for 2.5 h. Ten passes are rolled, with an initial rolling temperature of 1120℃ and a final rolling temperature of 980℃. The area reduction rate of the first three passes is 23% to break up coarse grains, the area reduction rate of the middle passes is 20%, and the area reduction rate of the last two passes is 18%. According to the above formula, the total rolling deformation is 89.94%. After rolling, it is held at 980℃ for 40 min for recrystallization annealing and then air-cooled. Step (7): The billet is subjected to solution heat treatment and aging heat treatment. Solution heat treatment is performed by heating to 1080℃ at a rate of 6℃ / min, holding for 1h, and then air cooling. Aging heat treatment is performed by heating to 800℃ at a rate of 4℃ / min, holding for 8h, and then air cooling.

[0062] The nickel-based wrought superalloy GH4780 of Example 1 has an average grain size of approximately 200 μm, with dispersed M precipitates at the grain boundaries. 23 (C,B)6 phase, with no continuous brittle phase; γ′ phase area fraction is about 23.7%, average size is about 33.5 nm.

[0063] Example 2 This embodiment provides a nickel-based wrought superalloy GH4780 and its preparation method.

[0064] The chemical composition of the nickel-based wrought superalloy GH4780, by weight percentage, includes: C: 0.057%, Si: 0.042%, B: 0.025%, Zr: 0.25% (Zr to B mass ratio is 10:1), Al: 1.14%, Ta: 0.97%, Ti: 2.09%, W: 1.81%, Nb: 0.68%, Co: 18.26%, Cr: 21.88%, N: 0.0002%, O: 0.001%, S: 0.0004%, P: 0.0021%, with the balance being Ni.

[0065] The preparation method is as follows: Step (1): Nickel, cobalt, tungsten, niobium, chromium, tantalum, and nickel-carbon master alloy are charged into a vacuum induction furnace for vacuum induction melting. After the furnace charge is completely melted, it is electromagnetically stirred for 10 minutes at a vacuum degree of 2×10⁻⁶. - Refining at 1540℃ and 2Pa for 15 minutes; after refining, cooling to 1480℃, adding aluminum in batches and stirring for 2 minutes, then adding titanium in batches and stirring for 2 minutes; then cooling to 1470℃, i.e., superheat of 120℃, and refilling with high-purity argon to 5×10³Pa, adding nickel-boron master alloy, zirconium, and nickel-carbon master alloy for carbon adjustment, electromagnetic stirring for 3 minutes to ensure uniform composition, and starting pouring at 1500℃ within 5 minutes after stirring, keeping the mold at the temperature for 3 hours after pouring, and demolding to obtain vacuum induction casting ingot; Step (2): The vacuum induction casting ingot is electroslag remelted using an alkaline pre-melted slag system. The alkaline pre-melted slag system comprises, by mass percentage: CaF2: 56%, Al2O3: 24%, CaO: 16%, MgO: 4%. The electroslag remelting rate is 4 kg / min, the water temperature is 32°C, and high-purity argon gas is introduced throughout the process for protection. The argon gas flow rate is 40 L / min. The ingot obtained by electroslag remelting is machined with a single-sided machining allowance of 8 mm to remove surface defects and compositional segregation layers. Step (3) involves vacuum arc remelting of the ingot. When the vacuum degree is <1×10 - Arc initiation occurs when the pressure is 3 Pa and the leakage rate is <0.3 Pa·L / s, the melting rate is 3.8 kg / min, the arc is stabilized by low current throughout the process, and the current of the arc stabilizing coil is 10 A; the vacuum self-consuming arc remelting adopts a hot capping process and helium gradient cooling at the end of the melting process. Step (4) involves subjecting the ingot to a three-stage gradient homogenization annealing process; the three-stage gradient homogenization annealing process includes: Step (4-1): Increase the temperature from room temperature to 1100℃ at a rate of 8℃ / min and hold for 20 hours; Step (4-2): Increase the temperature to 1150℃ at a rate of 0.5℃ / min and hold for 30 hours; Step (4-3): Heat to 1180℃ at a rate of 0.5℃ / min and hold for 45 hours; after holding, cool slowly to 900℃ at a rate of 0.5℃ / min and remove from the furnace. Step (5): Controllable hot working of the ingot: Heat the ingot to 1150℃ and hold for 5 hours. Then, perform upsetting and drawing on a high-speed forging machine in 3 passes. The upsetting deformation is 35% and the drawing deformation is 40%. The temperature drops by 40℃ for each pass. The final forging temperature is 1050℃. Draw the ingot to Φ110mm on the high-speed forging machine. Then, heat the intermediate billet to 1120℃ and hold for 2 hours. Perform 5 passes of hot radial forging on a radial forging machine. The forging frequency is 120 times / min. The single-pass area reduction rate of the main deformation pass is 13%. The billet rotates 30° for each forging. The last pass is the shaping pass. The single-pass area reduction rate is 7.7%. The rotation angle is 20°. The total radial forging deformation is 47.11%. The final forging temperature is 1000℃. The radial forging is completed to Φ80mm billet with a grain size difference of <3. Step (6): The billet is subjected to multi-pass temperature-controlled rolling: a 280-type transverse rolling mill is used, with a single-sided grinding amount of 0.3 mm. It is heated to 1130℃ at 5℃ / min and held for 2 hours. Eight passes are rolled, with an initial rolling temperature of 1090℃ and a final rolling temperature of 960℃. The area reduction rate of the first three passes is 20% to break up coarse grains, the area reduction rate of the middle passes is 18%, and the area reduction rate of the last two passes is 15%. According to the above formula, the total deformation of the rolling is 79.60%. After rolling, it is held at 960℃ for 30 minutes for recrystallization annealing and then air-cooled. Step (7): The billet is subjected to solution heat treatment and aging heat treatment. Solution heat treatment is performed by heating to 1070℃ at a rate of 5℃ / min, holding for 1h, and then air cooling. Aging heat treatment is performed by heating to 790℃ at a rate of 3℃ / min, holding for 6h, and then air cooling.

[0066] The nickel-based wrought superalloy GH4780 in Example 2 has an average grain size of approximately 200 μm, with a small amount of dispersed M precipitates at the grain boundaries. 23 (C,B)6 boron carbide, no continuous brittle phase formed; γ′ phase area fraction of about 24.5%, average size of about 32.5 nm.

[0067] Example 3 This embodiment provides a nickel-based wrought superalloy GH4780 and its preparation method.

[0068] The chemical composition of the nickel-based wrought superalloy GH4780, by weight percentage, includes: C: 0.052%, Si: 0.040%, B: 0.018%, Zr: 0.18% (Zr to B mass ratio is 10:1), Al: 1.10%, Ta: 1.05%, Ti: 2.08%, W: 1.78%, Nb: 0.65%, Co: 18.45%, Cr: 21.90%, N: 0.0001%, O: 0.001%, S: 0.0003%, P: 0.0022%, with the balance being Ni.

[0069] The preparation method is as follows: Step (1): Nickel, cobalt, tungsten, niobium, chromium, tantalum, and nickel-carbon master alloy are charged into a vacuum induction furnace for vacuum induction melting. After the furnace charge is completely melted, it is electromagnetically stirred for 8 minutes at a vacuum degree of 3×10⁻⁶. - Refining at 1530℃ and 2Pa for 13 minutes; after refining, cooling to 1450℃, adding aluminum in batches and stirring for 2 minutes, then adding titanium in batches and stirring for 1 minute; then cooling to 1440℃, i.e., superheat of 80℃, and refilling with high-purity argon to 7×10³Pa, adding nickel-boron master alloy, zirconium, and nickel-carbon master alloy for carbon adjustment, electromagnetic stirring for 5 minutes to ensure uniform composition, and starting pouring at 1480℃ within 5 minutes after stirring, holding in the mold for 2.5 hours after pouring, and demolding to obtain vacuum induction casting; Step (2): The vacuum induction casting ingot is electroslag remelted using an alkaline pre-melted slag system. The alkaline pre-melted slag system comprises, by mass percentage: CaF2: 58%, Al2O3: 22%, CaO: 16%, MgO: 4%. The electroslag remelting rate is 3.8 kg / min, the water temperature is 30°C, and high-purity argon gas is introduced throughout the process for protection. The argon gas flow rate is 40 L / min. The ingot obtained by electroslag remelting is machined with a single-sided machining allowance of 7 mm to remove surface defects and compositional segregation layers. Step (3) involves vacuum arc remelting of the ingot. When the vacuum degree is <1×10 - Arc initiation occurs when the pressure is 3 Pa and the leakage rate is <0.3 Pa·L / s, the melting rate is 3.6 kg / min, the arc is stabilized by low current throughout the process, and the current of the arc stabilizing coil is 10 A; the vacuum self-consuming arc remelting adopts a hot capping process and helium gradient cooling at the end of the melting process. Step (4) involves subjecting the ingot to a three-stage gradient homogenization annealing process; the three-stage gradient homogenization annealing process includes: Step (4-1): Increase the temperature from room temperature to 1130℃ at a rate of 12℃ / min and hold for 30 hours; Step (4-2): Increase the temperature to 1170℃ at a rate of 1℃ / min and hold for 35 hours; Step (4-3): Heat to 1195℃ at a rate of 1℃ / min and hold for 55h; after holding, cool slowly to 1000℃ at a rate of 1℃ / min and remove from the furnace. Step (5): Controllable hot working of the ingot: Heat the ingot to 1190℃ and hold for 8 hours. Then, perform upsetting and drawing on a high-speed forging machine in 3 passes. The upsetting deformation is 45% and the drawing deformation is 50% each time. The temperature drops by 40℃ each time and the final forging temperature is 1060℃. Draw the ingot to Φ160mm on the high-speed forging machine. Then, heat the intermediate billet to 1100℃ and hold for 2 hours. Perform 5 passes of hot radial forging on a radial forging machine. The forging frequency is 100 times / min. The single-pass area reduction rate of the main deformation pass is 12%. The billet rotates 20° each time it is forged. The last pass is the shaping pass. The single-pass area reduction rate is 6.2%. The rotation angle is 17°. The total radial forging deformation is 43.75%. The final forging temperature is 1000℃. Radial forging is performed to a Φ120mm bar billet with a grain size difference of <3 levels. Step (6): The billet is subjected to multi-pass temperature-controlled rolling: a 280-type tandem rolling mill is used, with a single-sided grinding amount of 0.5 mm. It is heated to 1170℃ at a rate of 10℃ / min and held for 2.5 h. Nine passes are rolled, with an initial rolling temperature of 1140℃ and a final rolling temperature of 1040℃. The area reduction rate of the first three passes is 25% to break up coarse grains, the area reduction rate of the middle passes is 22%, and the area reduction rate of the last two passes is 15%. According to the above formula, the total deformation of the rolling is 88.72%. After rolling, it is held at 1000℃ for 50 min for recrystallization annealing and then air-cooled. Step (7): The billet is subjected to solution heat treatment and aging heat treatment. Solution heat treatment is performed by heating to 1090℃ at a rate of 8℃ / min, holding for 2 hours, and then air cooling. Aging heat treatment is performed by heating to 810℃ at a rate of 5℃ / min, holding for 7 hours, and then air cooling.

[0070] The average grain size of the nickel-based wrought superalloy GH4780 in Example 3 is approximately 200 μm (e.g., Figure 4 As shown), M precipitates are dispersedly distributed at grain boundaries. 23 (C,B)6 phase, without a continuous brittle phase (e.g.) Figure 1 As shown); the γ′ phase area fraction is approximately 24.0%, and the average size is approximately 33.0 nm (as shown). Figure 3 (As shown).

[0071] Example 4 This embodiment provides a nickel-based wrought superalloy GH4780 and its preparation method.

[0072] The chemical composition of the nickel-based wrought superalloy GH4780, by weight percentage, includes: C: 0.052%, Si: 0.040%, B: 0.015%, Zr: 0.18% (the mass ratio of Zr to B is 12:1, which is the upper limit of the mass ratio range described in this invention), Al: 1.10%, Ta: 1.05%, Ti: 2.08%, W: 1.78%, Nb: 0.65%, Co: 18.45%, Cr: 21.90%, N: 0.0002%, O: 0.001%, S: 0.0003%, P: 0.0023%, with the balance being Ni.

[0073] The preparation method is as follows: Step (1): Nickel, cobalt, tungsten, niobium, chromium, tantalum, and nickel-carbon master alloy are charged into a vacuum induction furnace for vacuum induction melting. After the furnace charge is completely melted, it is electromagnetically stirred for 6 minutes at a vacuum degree of 3×10⁻⁶. - Refining at 1520℃ and 2Pa for 12 minutes; after refining, cooling to 1450℃, adding aluminum in batches and stirring for 2 minutes, then adding titanium in batches and stirring for 1 minute; then cooling to 1440℃, i.e., superheat of 80℃, and refilling with high-purity argon to 7×10³Pa, adding nickel-boron master alloy, zirconium, and nickel-carbon master alloy for carbon adjustment, electromagnetic stirring for 4 minutes to ensure uniform composition, and starting pouring at 1480℃ within 5 minutes after stirring, holding in the mold for 2.5 hours after pouring, and demolding to obtain vacuum induction casting ingot; Step (2): The vacuum induction casting ingot is electroslag remelted using an alkaline pre-melted slag system. The alkaline pre-melted slag system comprises, by mass percentage: CaF2: 58%, Al2O3: 22%, CaO: 16%, MgO: 4%. The electroslag remelting rate is 3.7 kg / min, the water temperature is 30°C, and high-purity argon gas is introduced throughout the process for protection. The argon gas flow rate is 35 L / min. The ingot obtained by electroslag remelting is machined with a single-sided machining allowance of 6 mm to remove surface defects and compositional segregation layers. Step (3) involves vacuum arc remelting of the ingot. When the vacuum degree is <1×10 - Arc initiation occurs when the pressure is 3 Pa and the leakage rate is <0.3 Pa·L / s, the melting rate is 3.6 kg / min, the arc is stabilized by low current throughout the process, and the current of the arc stabilizing coil is 10 A; the vacuum self-consuming arc remelting adopts a hot capping process and helium gradient cooling at the end of the melting process. Step (4) involves subjecting the ingot to a three-stage gradient homogenization annealing process; the three-stage gradient homogenization annealing process includes: Step (4-1): Increase the temperature from room temperature to 1120℃ at a rate of 11℃ / min and hold for 28 hours; Step (4-2): Increase the temperature to 1165℃ at a rate of 0.8℃ / min and hold for 36 hours; Step (4-3): Heat to 1190℃ at a rate of 0.8℃ / min and hold for 52 hours; after holding, cool slowly to 980℃ at a rate of 0.8℃ / min and remove from the furnace. Step (5): Controllable hot working of the ingot: Heat the ingot to 1180℃ and hold for 7 hours. Then, perform upsetting and drawing in three passes on a high-speed forging machine. The upsetting deformation is 42% and the drawing deformation is 46% each time. The temperature drops by 40℃ each time, and the final forging temperature is 1060℃. Draw the ingot into a Φ145mm intermediate billet on the high-speed forging machine. Then, heat the intermediate billet to 1110℃ and hold for 2 hours. Perform four passes of hot radial forging on a radial forging machine. The forging frequency is 90 times / min. The single-pass area reduction rate of the main deformation pass is 15%. The billet rotates 20° each time it is forged. The last pass is the shaping pass. The single-pass area reduction rate is 6.3%. The rotation angle is 15°. The total radial forging deformation is 42.45%. The final forging temperature is 1000℃. Radial forging is performed into a Φ110mm bar billet with a grain size difference of <3 levels. Step (6): The billet is subjected to multi-pass temperature-controlled rolling: a 280-type tandem rolling mill is used, with a single-sided grinding amount of 0.5 mm. It is heated to 1160℃ at 9℃ / min and held for 2.5 h, and then subjected to 9 passes of continuous temperature-controlled rolling. The initial rolling temperature is 1130℃ and the final rolling temperature is 1020℃. The area reduction rate of the first 3 passes is 24% to break up coarse grains, the area reduction rate of the middle passes is 20%, and the area reduction rate of the last 2 passes is 16%. According to the above formula, the total rolling deformation is 87.31%. After rolling, it is held at 990℃ for 45 min for recrystallization annealing and then air-cooled. Step (7): The billet is subjected to solution heat treatment and aging heat treatment. The solution heat treatment is: heating to 1085℃ at 7℃ / min, holding for 1.5h, and air cooling; the aging heat treatment is: heating to 805℃ at 4℃ / min, holding for 7h, and air cooling.

[0074] The alloy in Example 4 has an average grain size of approximately 200 μm, with grain boundaries consisting of finely dispersed M-type grains. 23 The precipitates are mainly composed of (C,B)6 boron carbides, with a small amount of Zr-stabilized MC-type carbides. The average size of the precipitates is about 0.25 μm, and they are discontinuously distributed. No continuous brittle phase film was formed, and no Zr-rich low-melting-point eutectic phase was observed. The area fraction of the γ′ phase is about 24.2%, and the average size is about 33.5 nm.

[0075] Example 5 This embodiment provides a nickel-based wrought superalloy GH4780 and its preparation method.

[0076] The chemical composition of the nickel-based wrought superalloy GH4780, by weight percentage, includes: C: 0.052%, Si: 0.040%, B: 0.025%, Zr: 0.20% (the mass ratio of Zr to B is 8:1, which is the lower limit of the mass ratio range described in this invention), Al: 1.10%, Ta: 1.05%, Ti: 2.08%, W: 1.78%, Nb: 0.65%, Co: 18.45%, Cr: 21.90%, N: 0.0001%, O: 0.002%, S: 0.0002%, P: 0.0018%, with the balance being Ni.

[0077] The preparation method is as follows: Step (1): Nickel, cobalt, tungsten, niobium, chromium, tantalum, and nickel-carbon master alloy are charged into a vacuum induction furnace for vacuum induction melting. After the furnace charge is completely melted, it is electromagnetically stirred for 7 minutes at a vacuum degree of 3×10⁻⁶. - Refining at 1535℃ and 2Pa for 14 minutes; after refining, cooling to 1450℃, adding aluminum in batches and stirring for 2 minutes, then adding titanium in batches and stirring for 1 minute; then cooling to 1440℃, i.e., superheat of 80℃, and refilling with high-purity argon to 7×10³Pa, adding nickel-boron master alloy, zirconium, and nickel-carbon master alloy for carbon adjustment, electromagnetic stirring for 5 minutes to ensure uniform composition, and starting pouring at 1490℃ within 5 minutes after stirring, holding in the mold for 2.8 hours after pouring, and demolding to obtain vacuum induction casting ingot; Step (2): The vacuum induction casting ingot is electroslag remelted using an alkaline pre-melted slag system. The alkaline pre-melted slag system comprises, by mass percentage: CaF2: 56%, Al2O3: 24%, CaO: 16%, MgO: 4%. The electroslag remelting rate is 3.9 kg / min, the water temperature is 32°C, and high-purity argon gas is introduced throughout the process for protection. The argon gas flow rate is 38 L / min. The ingot obtained by electroslag remelting is machined with a single-sided machining allowance of 6.5 mm to remove surface defects and compositional segregation layers. Step (3) involves vacuum arc remelting of the ingot. When the vacuum degree is <1×10 - Arc initiation occurs when the pressure is 3 Pa and the leakage rate is <0.3 Pa·L / s, the melting rate is 3.7 kg / min, the arc is stabilized by low current throughout the process, and the current of the arc stabilizing coil is 10 A; the vacuum self-consuming arc remelting adopts a hot capping process and helium gradient cooling at the end of the melting process. Step (4) involves subjecting the ingot to a three-stage gradient homogenization annealing process; the three-stage gradient homogenization annealing process includes: Step (4-1): Increase the temperature from room temperature to 1130℃ at a rate of 12℃ / min and hold for 30 hours; Step (4-2): Increase the temperature to 1170℃ at a rate of 1℃ / min and hold for 38 hours; Step (4-3): Heat to 1195℃ at a rate of 1℃ / min and hold for 54 hours; after holding, cool slowly to 990℃ at a rate of 1℃ / min and remove from the furnace. Step (5): Controllable hot working of the ingot: Heat the ingot to 1185℃ and hold for 7.5h. Then, perform three upsetting and drawing operations on a high-speed forging machine. The upsetting deformation is 44% and the drawing deformation is 48% each time. The temperature drops by 40℃ each time. The final forging temperature is 1055℃. The ingot is drawn to Φ155mm on the high-speed forging machine. Then, heat the ingot to 1090℃ and hold for 2h. Then, perform four hot radial forging operations on a radial forging machine. The forging frequency is 80 times / min. The single-pass area reduction rate of the main deformation pass is 16%. The billet rotates 25° each time it is forged. The last pass is the shaping pass. The single-pass area reduction rate is 7.1%. The rotation angle is 15°. The total radial forging deformation is 44.95%. The final forging temperature is 1000℃. The ingot is radially forged to Φ115mm and the grain size difference is <3. Step (6): The billet is subjected to multi-pass temperature-controlled rolling: a 280-type transverse rolling mill is used, with a single-sided grinding amount of 0.6 mm. It is heated to 1170℃ at a rate of 10℃ / min and held for 3 hours. Ten passes of continuous temperature-controlled rolling are performed, with an initial rolling temperature of 1140℃ and a final rolling temperature of 1030℃. The area reduction rate of the first three passes is 25% to break up coarse grains, the area reduction rate of the middle passes is 22%, and the area reduction rate of the last two passes is 17%. According to the above formula, the total deformation of the rolling is 91.61%. After rolling, it is held at 1000℃ for 50 minutes for recrystallization annealing and then air-cooled. Step (7): The billet is subjected to solution heat treatment and aging heat treatment. Solution heat treatment is performed by heating to 1090℃ at a rate of 8℃ / min, holding for 2 hours, and then air cooling. Aging heat treatment is performed by heating to 810℃ at a rate of 5℃ / min, holding for 8 hours, and then air cooling.

[0078] The alloy in Example 5 has an average grain size of approximately 200 μm. The B content in this example is 0.025%, and the grain boundary M... 23 The amount of (C,B)6 boron carbide precipitates increased compared to Example 3, with an average size of about 0.30 μm. It still maintained a diffuse and discontinuous distribution, without coarsening into a continuous film, and no Zr-rich low-melting-point eutectic phase was observed. The matrix γ′ phase had an area fraction of about 24.8% and an average size of about 34.5 nm.

[0079] Example 6 This embodiment provides a nickel-based wrought superalloy GH4780 and its preparation method.

[0080] The chemical composition of the nickel-based wrought superalloy GH4780, by weight percentage, includes: C: 0.042%, Si: 0.065%, B: 0.010%, Zr: 0.11% (Zr to B mass ratio is 11:1), Al: 1.02%, Ta: 1.35%, Ti: 2.02%, W: 1.72%, Nb: 0.61%, Co: 18.10%, Cr: 21.60%, N: 0.0002%, O: 0.002%, S: 0.0002%, P: 0.0022%, with the balance being Ni.

[0081] The preparation method is as follows: Step (1): Nickel, cobalt, tungsten, niobium, chromium, tantalum, and nickel-carbon master alloy are charged into a vacuum induction furnace for vacuum induction melting. After the furnace charge is completely melted, it is electromagnetically stirred for 5 minutes at a vacuum degree of 3×10⁻⁶. - Refining at 1505℃ and 2Pa for 10 minutes; after refining, cooling to 1450℃, adding aluminum in batches and stirring for 2 minutes, then adding titanium in batches and stirring for 1 minute; then cooling to 1440℃, i.e., superheat of 80℃, and refilling with high-purity argon to 7×10³Pa, adding nickel-boron master alloy, zirconium, and nickel-carbon master alloy for carbon adjustment, electromagnetic stirring for 3 minutes to ensure uniform composition, and starting pouring at 1465℃ within 5 minutes after stirring, keeping the mold at 2 hours after pouring, and demolding to obtain vacuum induction casting ingot; Step (2): The vacuum induction casting ingot is electroslag remelted using an alkaline pre-melted slag system. The alkaline pre-melted slag system comprises, by mass percentage: CaF2: 60%, Al2O3: 20%, CaO: 16%, MgO: 4%. The electroslag remelting rate is 3.5 kg / min, the water temperature is 28°C, and high-purity argon gas is introduced throughout the process for protection. The argon gas flow rate is 30 L / min. The ingot obtained by electroslag remelting is machined with a single-sided machining allowance of 5 mm to remove surface defects and compositional segregation layers. Step (3) involves vacuum arc remelting of the ingot. When the vacuum degree is <1×10 - Arc initiation occurs when the pressure is 3 Pa and the leakage rate is <0.3 Pa·L / s, the melting rate is 3.4 kg / min, the arc is stabilized by low current throughout the process, and the current of the arc stabilizing coil is 10 A; the vacuum self-consuming arc remelting adopts a hot capping process and helium gradient cooling at the end of the melting process. Step (4) involves subjecting the ingot to a three-stage gradient homogenization annealing process; the three-stage gradient homogenization annealing process includes: Step (4-1): Increase the temperature from room temperature to 1105℃ at a rate of 8℃ / min and hold for 22 hours; Step (4-2): Increase the temperature to 1155℃ at a rate of 0.5℃ / min and hold for 32 hours; Step (4-3): Heat to 1185℃ at a rate of 0.5℃ / min and hold for 46 hours; after holding, cool slowly to 920℃ at a rate of 0.5℃ / min and remove from the furnace. Step (5): Controllable hot working of the ingot: Heat the ingot to 1155℃ and hold for 5 hours. Then, perform upsetting and drawing on a high-speed forging machine in 3 passes. The upsetting deformation is 36% and the drawing deformation is 41% each time. The temperature drops by 40℃ each time. The final forging temperature is 1050℃. The ingot is drawn to Φ120mm intermediate billet on the high-speed forging machine. Then, heat the intermediate billet to 1090℃ and hold for 2 hours. Perform 4 passes of hot radial forging on a radial forging machine. The forging frequency is 80 times / min. The single-pass area reduction rate of the main deformation pass is 19%. The billet rotates 25° each time it is forged. The last pass is the shaping pass. The single-pass area reduction rate is 5.6%. The rotation angle is 15°. The total radial forging deformation is 49.83%. The final forging temperature is 1000℃. The ingot is radially forged to Φ85mm bar billet with a grain size difference of <3. Step (6): The billet is subjected to multi-pass temperature-controlled rolling: a 280-type transverse rolling mill is used, with a single-sided grinding amount of 0.3 mm. It is heated to 1135℃ at a rate of 5℃ / min and held for 2 hours. It is subjected to 8 passes of continuous temperature-controlled rolling, with an initial rolling temperature of 1095℃ and a final rolling temperature of 965℃. The area reduction rate of the first 3 passes is 21% to break up coarse grains, the area reduction rate of the middle passes is 19%, and the area reduction rate of the last 2 passes is 15%. According to the above formula, the total deformation of the rolling is 81.07%. After rolling, it is held at 965℃ for 35 minutes for recrystallization annealing and then air-cooled. Step (7): The billet is subjected to solution heat treatment and aging heat treatment. Solution heat treatment is performed by heating to 1070℃ at a rate of 5℃ / min, holding for 1h, and then air cooling. Aging heat treatment is performed by heating to 790℃ at a rate of 3℃ / min, holding for 6h, and then air cooling.

[0082] The alloy in Example 6 has an average grain size of approximately 170 μm and grain boundary M 23 The precipitation of (C,B)6 boron carbides was relatively small, with an average size of about 0.20 μm. No continuous brittle phase or Zr-rich low-melting-point eutectic phase was observed. The γ′ phase had an area fraction of about 22.5% and an average size of about 30.0 nm.

[0083] Comparative Example 1 This comparative example provides a nickel-based wrought superalloy GH4780 alloy and its preparation method similar to Example 3. The main difference is that the contents of B and Zr are lower. The chemical composition of the nickel-based wrought superalloy GH4780 alloy by weight percentage includes: C: 0.058%, Si: 0.021%, B: 0.003%, Zr: 0.03% (Zr to B mass ratio is 10:1), Al: 1.11%, Ta: 0.96%, Ti: 2.12%, W: 1.79%, Nb: 0.67%, Co: 18.27%, Cr: 21.87%, N: 0.0001%, O: 0.001%, S: 0.0003%, P: 0.0022%, with the balance being Ni.

[0084] The average grain size of the nickel-based wrought superalloy GH4780 in Comparative Example 1 is approximately 180 μm, and the grain boundary M 23 (C,B)6 pinned network is discontinuous; γ′ phase area fraction is approximately 24.2%, average size is approximately 32.8 nm.

[0085] Comparative Example 2 This comparative example provides a nickel-based wrought superalloy GH4780 alloy and its preparation method, similar to Example 3. The main difference is that the contents of B and Zr are slightly lower than the lower limits defined in this invention. The chemical composition of the nickel-based wrought superalloy GH4780 alloy by weight percentage includes: C: 0.057%, Si: 0.049%, B: 0.007%, Zr: 0.07% (Zr to B mass ratio is 10:1), Al: 1.08%, Ta: 0.96%, Ti: 2.05%, W: 1.75%, Nb: 0.63%, Co: 18.6%, Cr: 21.91%, with the balance being Ni.

[0086] The nickel-based wrought superalloy GH4780 of Comparative Example 2 has an average grain size of about 190 μm; the γ′ phase area fraction is about 23.8%, and the average size is about 33.1 nm.

[0087] Comparative Example 3 This comparative example provides a nickel-based wrought superalloy GH4780 alloy and its preparation method similar to Example 3. The main difference is that the contents of B and Zr are too high. The chemical composition of the nickel-based wrought superalloy GH4780 alloy by weight percentage includes: C: 0.057%, Si: 0.021%, B: 0.035%, Zr: 0.35% (Zr to B mass ratio is 10:1), Al: 1.08%, Ta: 0.96%, Ti: 2.05%, W: 1.75%, Nb: 0.63%, Co: 18.6%, Cr: 21.91%, and the balance is Ni.

[0088] The nickel-based wrought superalloy GH4780 of Comparative Example 3 has an average grain size of about 200 μm, with continuous coarse borides and Zr-rich low-melting-point eutectic phases forming at the grain boundaries; the γ′ phase has an area fraction of about 23.9% and an average size of about 33.8 nm.

[0089] Comparative Example 4 This comparative example provides a nickel-based wrought superalloy GH4780 alloy and its preparation method, similar to Example 3. The main difference is that the mass ratio of Zr to B is relatively large. The chemical composition of the nickel-based wrought superalloy GH4780 alloy by weight percentage includes: C: 0.054%, Si: 0.041%, B: 0.012%, Zr: 0.24% (Zr to B mass ratio is 20:1), Al: 1.08%, Ta: 0.95%, Ti: 2.04%, W: 1.74%, Nb: 0.63%, Co: 18.56%, Cr: 21.85%, with the balance being Ni.

[0090] The average grain size of the nickel-based wrought superalloy GH4780 in Comparative Example 4 is about 200 μm. Due to the relative excess of Zr, the grain boundaries are enriched with coarse Zr-rich low-melting-point eutectic phases, and locally there is a continuous brittle phase distribution; the area fraction of the γ′ phase is about 24.0%, and the average size is about 33.0 nm.

[0091] Comparative Example 5 This comparative example provides a nickel-based wrought superalloy GH4780 alloy and its preparation method similar to Example 3. The main difference is that the mass ratio of Zr to B is smaller. The chemical composition of the nickel-based wrought superalloy GH4780 alloy by weight percentage includes: C: 0.052%, Si: 0.040%, B: 0.030%, Zr: 0.20% (Zr to B mass ratio is 6.7:1), Al: 1.10%, Ta: 1.05%, Ti: 2.08%, W: 1.78%, Nb: 0.65%, Co: 18.45%, Cr: 21.90%, with the balance being Ni.

[0092] The average grain size of the nickel-based wrought superalloy GH4780 in Comparative Example 5 is approximately 200 μm; due to the relative excess of boron, continuous coarse argon crystals precipitate at the grain boundaries. 23 (C,B)6 boron carbides, with an average size of approximately 0.45 μm, are locally distributed in a film-like manner (e.g. Figure 2 (As shown); the γ′ phase area fraction is approximately 24.3%, and the average size is approximately 33.7 nm.

[0093] Performance test examples The tensile and creep properties of the nickel-based wrought superalloy GH4780 in the examples and comparative examples were tested. The tensile properties were tested according to ASTM E8 (room temperature 23℃) and ASTM E21 (high temperature 760℃) standards; the creep properties were tested according to ASTM E139 standard. The test results are shown in Table 1.

[0094] Table 1. Tensile and creep properties of GH4780 alloys in the examples and comparative examples.

[0095] As can be seen from Table 1, the nickel-based wrought superalloy GH4780 from Examples 1 to 6 all exhibit good high-temperature creep performance. The alloy combines high strength and plasticity. At room temperature (23℃), the tensile strength is ≥1238MPa, the room temperature yield strength is ≥783MPa, and the room temperature elongation is ≥26.5%. At 760℃, the tensile strength is ≥854MPa, the yield strength is ≥703MPa, and the elongation is ≥13.8%. At 760℃ / 460MPa, the creep life is ≥70.0h and the creep elongation is ≥12.5%.

[0096] Comparing the parameters of each comparative example with Example 3 (B: 0.018%, Zr: 0.18%, Zr to B mass ratio of 10:1) as a benchmark, it can be seen that: the B and Zr contents of Comparative Example 1 (B: 0.003%, Zr: 0.03%) are lower than the content range of the present invention, and the grain boundary M 23(C,B)6 The stapled network is discontinuous, and its creep life at 760℃ / 460MPa decreased from 79.2h in Example 3 to 52.6h, a decrease of approximately 33.6%. Although the room temperature elongation is slightly higher than that of Example 3 (33.5% vs. 29.5%), its high-temperature long-term load-bearing capacity is significantly insufficient. The B and Zr content (B: 0.007%, Zr: 0.07%) in Comparative Example 2 is still lower than the content range of the present invention, and the creep life decreased to 61.8h, a decrease of approximately 22.0%, indicating that B and Zr content are lower than that of the present invention. When the Zr content is below the lower limit, the grain boundary segregation is insufficient to form a continuous and stable pinning network. In Comparative Example 3, the B and Zr contents (B: 0.035%, Zr: 0.35%) are higher than the range of this invention. Continuous coarse borides and a Zr-rich, low-melting-point eutectic phase form at the grain boundaries. The room temperature elongation decreases from 29.5% to 10.5%, the 760°C elongation decreases from 15.2% to 3.8%, the creep life decreases to 10.2 hours (a decrease of approximately 87.1%), and the creep elongation decreases to 2.1%. Plasticity and long-term high-temperature performance deteriorated sharply. Although the B and Zr contents of Comparative Example 4 were within the range of this invention, the Zr to B mass ratio was 20:1, higher than the upper limit of 8-12:1, indicating a relative excess of Zr. The Zr-rich low-melting-point eutectic phase was enriched at the grain boundaries and locally distributed continuously. The room temperature elongation decreased to 14.5%, the 760℃ elongation decreased to 5.5%, and the creep life decreased to 28.5h, a decrease of approximately 64.0%. Both plasticity and creep life decreased significantly, increasing the risk of hot working cracking. In Example 5, the contents of B and Zr are also within the scope of this invention, but the mass ratio of Zr to B is 6.7:1, which is lower than the lower limit of 8 to 12:1. B is relatively excessive, and the boron carbides at the grain boundaries are excessively coarsened and distributed in a continuous film, which breaks the grain boundary bonding force. The room temperature elongation decreases from 29.5% to 13.5%, the 760°C elongation decreases from 15.2% to 5.0%, the creep life decreases to 26.0h, a decrease of about 67.2%, and the creep elongation decreases to 4.5%. Both plasticity and creep life are significantly reduced. Comparison of Comparative Examples 1 to 5 with Example 3 shows that only when the B and Zr content ranges defined by this invention (B: 0.008% to 0.030%, Zr: 0.08% to 0.30%) and the Zr to B mass ratio are both 8 to 12:1, can a synergistic dispersion structure of "boron carbide pinning + MC stabilization" be formed at the grain boundaries, thus achieving both high creep life at 760℃ and excellent plasticity. Deviations in content from the range or ratios from 8 to 12:1 result in significant deterioration of creep life or plasticity, and the two cannot be achieved simultaneously.

[0097] 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 wrought superalloy GH4780, characterized in that, The chemical composition of the nickel-based wrought superalloy GH4780, by weight percentage, includes: C: 0.04%–0.06%, Cr: 21.5%–22.0%, Co: 18.0%–18.8%, W: 1.7%–1.85%, Nb: 0.60%–0.70%, Al: 1.0%–1.2%, Ti: 2.0%–2.2%, Ta: 0.9%–1.5%, B: 0.008%–0.030%, Zr: 0.08%–0.30%, Si: ≤0.08%, S: ≤10ppm, P: ≤40ppm, O+N+S: ≤50ppm, with Ni as the balance. The mass ratio of Zr to B is 8 to 12:

1.

2. The nickel-based wrought superalloy GH4780 according to claim 1, characterized in that, The nickel-based wrought superalloy GH4780 has the following chemical composition by weight percentage: C: 0.045%–0.058%, Cr: 21.8%–21.95%, Co: 18.2%–18.7%, W: 1.74%–1.81%, Nb: 0.63%–0.68%, Al: 1.05%–1.15%, Ti: 2.04%–2.12%, Ta: 0.95%–1.21%, B: 0.012%–0.025%, Zr: 0.12%–0.25%, Si: ≤0.05%, S: ≤5ppm, P: ≤30ppm, O+N+S: ≤30ppm, Ni is the balance; And / or, the mass ratio of Zr to B is 9~11:

1.

3. A method for preparing a nickel-based wrought superalloy GH4780, characterized in that, The preparation method for the nickel-based wrought superalloy GH4780 alloy according to claim 1 or 2 includes: Step (1): First, nickel, cobalt, tungsten, niobium, chromium, tantalum and nickel-carbon master alloy are loaded into a vacuum melting furnace for vacuum induction melting. After refining, aluminum and titanium are added in batches. Finally, at the end of the vacuum induction melting and at a superheat of 80-120°C, nickel-boron master alloy, zirconium and nickel-carbon master alloy for carbon adjustment are added and electromagnetically stirred for 3-5 minutes. After pouring, the ingot is demolded to obtain a vacuum induction casting. Step (2) involves electroslag remelting the vacuum induction casting ingot, wherein the electroslag remelting uses an alkaline pre-melted slag system; Step (3) involves vacuum consumable arc remelting of the ingot; Step (4) involves subjecting the ingot to a three-stage gradient homogenization annealing process; Step (5): The ingot is subjected to controlled heat treatment to obtain a billet; Step (6): The billet is subjected to multiple temperature-controlled rolling passes; Step (7) involves subjecting the billet to solution heat treatment and aging heat treatment; The three-stage gradient homogenization annealing process includes: Step (4-1): Increase the temperature from room temperature to 1100-1130℃ at a rate of 8-12℃ / min and hold for 20-30 hours; Step (4-2): Increase the temperature to 1150-1170℃ at a rate of 0.5-1℃ / min and hold for 30-40 hours; Step (4-3): Heat to 1180-1195℃ at a rate of 0.5-1℃ / min and hold for 45-55 hours; after holding, cool slowly to 900-1000℃ at a rate of 0.5-1℃ / min and remove from the furnace.

4. The preparation method according to claim 3, characterized in that, In step (1), the vacuum degree of the vacuum induction melting is ≤5×10 -2 Pa, the refining temperature is 1500~1540℃, and the refining time is 10~15min.

5. The preparation method according to claim 3, characterized in that, In step (1), the pouring temperature is 1460-1500℃, and the mold is kept warm for 2-3 hours after pouring.

6. The preparation method according to claim 3, characterized in that, In step (2), the alkaline premelted slag system comprises, by mass percentage: CaF2: 55%–60%, Al2O3: 20%–25%, CaO: 15%–20%, MgO: 3%–5%.

7. The preparation method according to claim 3, characterized in that, In step (2), the melting rate of the electroslag remelting is 3.5 to 4.0 kg / min, the water temperature is 28 to 32°C, and high-purity argon gas is introduced for protection throughout the process, with an argon gas flow rate ≥ 30 L / min.

8. The preparation method according to claim 3, characterized in that, In step (3), during the vacuum self-consuming arc remelting, when the vacuum degree is <1×10 -3 Arc initiation occurs when Pa and leakage rate is <0.3 Pa·L / s, and melting rate is 3.4~3.8 kg / min.

9. The preparation method according to claim 3, characterized in that, In step (5), the controllable hot working blanking includes sequentially performing rapid forging and radial forging.

10. The preparation method according to claim 3, characterized in that, In step (7), the solution heat treatment includes: heating to 1070-1090°C at a rate of 5-8°C / min, holding at that temperature for 1-2 hours, and then air cooling; And / or, the aging heat treatment includes: heating to 790-810°C at a rate of 3-5°C / min, holding at that temperature for 6-8 hours, and then air cooling.

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