High-strength oxygen-enriched-ablation-resistant nickel-based superalloy and preparation method thereof

Abstract

A high-strength, oxygen-rich ablation-resistant nickel-based superalloy and its preparation method are disclosed, belonging to the technical field of nickel-based superalloys. The alloy comprises the following components: Cr: 16.0–17.0%, Mo: 2.5–3.5%, W: 0–1.0%, Cu: 0.5–3.0%, Ta: 0–3.0%, Nb: 4.5–5.5%, Al: 0.5–1.8%, Ti: 0.5–1.0%, Co: 0–9.0%, Fe: 0–12.0%, C: 0.02–0.06%, V: 0.3–0.6%, with the remainder being Ni. The preparation process steps are as follows: Step 1: Vacuum induction melting to prepare an ingot; Step 2: Homogenization treatment of the alloy ingot; Step 3: Forging of the homogenized alloy ingot; Step 4: Heat treatment of the forged billet. The high-strength, oxygen-enriched ablation-resistant nickel-based superalloy is characterized by containing 0.5–3.0% Cu, with a calorific value <1430 cal / g, giving the alloy excellent resistance to oxygen-enriched ablation. After forging, the alloy exhibits excellent room-temperature and high-temperature tensile fracture strength. This invention is primarily intended for designing a high-temperature alloy material for use in high-temperature, high-pressure, oxygen-enriched environments, such as a turbopump material for liquid oxygen / kerosene rocket engines.

Description

Technical Field

[0001] This invention relates to the field of high-temperature alloy preparation, and in particular to a high-strength nickel-based high-temperature alloy resistant to oxygen-rich ablation and its preparation method. Background Technology

[0002] Nickel-based superalloys are widely used as hot-end components of advanced aerospace engines due to their excellent high-temperature mechanical properties. Their research, development, manufacturing and application have become one of the important indicators for measuring a country's level of materials development.

[0003] As a core hot-end component of liquid oxygen / kerosene rocket engines, the turbopump must withstand the extremely harsh operating environment of high temperature, high pressure, and oxygen-rich conditions. Under these conditions, metallic materials are prone to undergoing a combustion phenomenon that differs from both ordinary oxidation and melting, but is similar to that of wood, nylon, etc., known as metal combustion. Metal combustion is often accompanied by a rapid increase in temperature, flame formation, volume reduction, and violent heat release, severely limiting the operating temperature of high-temperature alloy materials.

[0004] As engine thrust increases, the operating temperature and pressure of turbopumps gradually increase, placing higher demands on the mechanical properties and resistance to oxygen-rich ablation of turbopump materials. Therefore, to meet the requirements of higher-thrust liquid oxygen / kerosene rocket engines, it is urgently necessary to develop high-temperature alloy materials with higher ignition points. Summary of the Invention

[0005] This invention provides a high-strength nickel-based superalloy resistant to oxygen-rich ablation and its preparation method. The main purpose is to enable the alloy to possess both excellent tensile strength and resistance to oxygen-rich ablation.

[0006] To achieve the above objectives, the present invention mainly provides the following technical solutions:

[0007] A high-strength, oxygen-rich ablation-resistant nickel-based superalloy (GH4022), wherein the deformed superalloy comprises the following components by weight percentage:

[0008] Cr: 16.0–17.0%, Mo: 2.5–3.5%, W: 0–1.0%, Cu: 0.5–3.0%, Ta: 0–3.0%, Nb: 4.5–5.5%, Al: 0.5–1.8%, Ti: 0.5–1.0%, Co: 0–9.0%, Fe: 0–12.0%, C: 0.02–0.06%, V: 0.3–0.6%, with the remainder being Ni.

[0009] The preferred alloy composition (wt.%) is as follows: Cr: 16.0–17.0%, Mo: 2.5–3.5%, W: 0.8–1.2%, Cu: 1.0–2.0%, Ta: 1.0–2.0%, Nb: 4.5–5.5%, Al: 0.5–1.0%, Ti: 0.5–1.0%, Co: 8.5–9.5%, Fe: ≤0.5%, C: 0.02–0.06%, V: 0.3–0.6%, with the remainder being Ni.

[0010] The chemical composition design of alloy GH4022 in this invention is mainly based on the following reasons:

[0011] The combustion characteristics of metallic materials in high-temperature, high-pressure, oxygen-rich environments are related to the heat of combustion of the elements. Elements with higher heat of combustion release more heat, which promotes sustained combustion of the material. The heats of combustion of commonly used strengthening alloying elements in high-temperature alloys and Cu are shown in Table 1.

[0012] As can be seen from the table, Cu has a very low heat of combustion, and adding Cu to alloys can improve their resistance to oxygen-enriched ablation. However, copper is an impurity element in high-temperature alloys, and its content must be controlled. Early studies suggested that Cu segregates at grain boundaries, reducing grain boundary bonding and thus negatively impacting the alloy's durability. However, rocket engines operate for very short periods, so the requirements for material durability are not high; instead, the instantaneous mechanical properties of the alloy are more important. This provides a design window for the composition design of oxygen-enriched ablation resistant alloys—an alloy design approach that combines excellent resistance to high-pressure oxygen-enriched ablation with tensile properties through the addition of Cu.

[0013] Table 1. Heat of combustion of metallic elements (cal / g)

[0014] Heat of combustion 980 7425 4710 1090 1460 585 1560 2490 2595 1345 955

[0015] Based on this, embodiments of the present invention provide a deformed high-temperature alloy, in which Co and Cu replace Fe, W replaces Mo, and Ta replaces Al, giving the alloy both excellent mechanical properties and resistance to oxygen-rich ablation. However, adding too much Cu will lower the solidus and liquidus temperatures of the alloy, which is detrimental to its combustion performance. Therefore, the preferred alloy contains 1.0 to 2.0 wt.% Cu.

[0016] Co has an inhibitory effect on the TCP phase, but excessively high Co content will lower the solution temperature, leading to a decrease in the high-temperature performance of the alloy. More importantly, compared with Fe, Co has a lower calorific value, which is beneficial to the alloy's combustion performance. Therefore, to ensure the alloy's high-temperature performance while also considering combustion performance, the Co content is controlled between 8.5 and 9.5 wt.%.

[0017] Cr is a key element for improving the hot corrosion resistance of alloys. An appropriate amount of Cr must be added to the alloy, with 16.0 to 17.0 wt.% of Cr being preferred.

[0018] W and Mo are solid solution strengthening elements that can produce lattice distortion, hindering dislocation movement within the matrix and thus improving the alloy's strength and creep resistance. However, excessive W and Mo can also lead to the precipitation of harmful phases, negatively impacting the alloy's hot corrosion resistance. Compared to Mo, W has a lower calorific value. Therefore, in the original alloy, W is used to partially replace Mo, controlling the Mo content between 2.5 and 3.5 wt.% and the W content between 0.8 and 1.2 wt.%.

[0019] Nitrogen (Nb) is a common solid solution strengthening element, as well as a precipitation strengthening element. Nb has a larger atomic radius than W and Mo, therefore its solid solution strengthening effect is more pronounced. Simultaneously, Nb is also the forming element of the important strengthening phase γ” (Ni3Nb) in this alloy. The γ” phase in the alloy is disk-shaped and has a significant coherent stress strengthening effect, thus enabling the alloy to achieve relatively high yield strength. However, excessive Nb content can also cause the precipitation of the Laves phase; therefore, 4.5–5.5% Nb is added to the alloy.

[0020] Al, Ti, and Ta are γ' phase forming elements, which have an age-hardening precipitation strengthening effect on the alloy, ensuring high high-temperature strength and creep resistance. High Ti content and Ti / Al ratio in the alloy easily lead to the formation of the harmful η phase, affecting the alloy's hot working properties. More importantly, Al and Ti have very high heats of combustion; excessively high Al and Ti contents are detrimental to the alloy's resistance to oxygen-rich ablation. Therefore, the Al and Ti content in this alloy should ideally be 0.5–1.0 wt.%. Compared to Al and Ti, Ta has a lower heat of combustion; substituting Ta for Al can balance mechanical and combustion properties, but excessively high Ta content significantly reduces the alloy's combustion performance. Therefore, 1.0–2.0% Ta is added to the alloy.

[0021] C and V are the most widely used microalloying elements in high-temperature alloys. Adding C is for purifying the alloy melt (deoxidation), and the precipitated carbides can strengthen grain boundaries and improve the alloy's mechanical properties. Vanadium in high-temperature alloys is mainly distributed in the γ phase, accounting for approximately 70%–87% of the total vanadium content, followed by the γ' phase, accounting for approximately 14%–29%. The atomic radius of V distributed in the austenitic matrix is ​​5% larger than that of Ni, causing lattice distortion and playing a significant role in solid solution strengthening. In addition, V also has a grain-refining effect on the alloy. Therefore, the alloy requires C and V contents of 0.02–0.06 wt.% and 0.3–0.6 wt.%, respectively.

[0022] The proper proportions of the above elements are the guarantee that the alloy of this invention obtains good mechanical properties and resistance to oxygen-rich ablation.

[0023] The preparation method of the above-mentioned deformed high-temperature alloy includes the following steps:

[0024] 1) Vacuum induction melting is used to prepare ingots. The refining temperature is 1500℃±10℃, the time is 20 minutes, the vacuum degree is 1~20Pa, and then the ingots are cast into ingots (Φ80mm) at a casting temperature of 1450℃±10℃.

[0025] 2) The alloy ingot is subjected to homogenization treatment at a temperature of 1150℃±10℃ for a time of not less than 20 hours.

[0026] 3) The homogenized alloy ingot is forged using a two-upsetting, two-drawing forging process to obtain a deformed high-temperature alloy (deformation between 50% and 60%). The forging temperature is 970℃-1100℃.

[0027] 4) After forging, the billet is subjected to heat treatment. The solution treatment regime is: 960±20℃, hold for 1h, and air cool; the aging treatment regime is: 780℃±10℃, hold for 8h, furnace cool to 700℃±10℃, hold for 8h, and air cool.

[0028] The advantages and beneficial effects of this invention are explained below:

[0029] Compared with the wrought high-temperature alloys commonly used in my country's existing liquid oxygen / kerosene rocket engines, the alloy of this invention simultaneously improves the material's room temperature and high-temperature tensile strength, as well as its resistance to oxygen-enriched ablation. The alloy of this invention has a room temperature tensile strength ≥1250 MPa and a tensile strength at 750℃ ≥780 MPa. The oxygen-enriched ablation pressure threshold has been increased from the original 5 MPa to 8 MPa, significantly improving the material's safety in high-temperature, high-pressure, and oxygen-enriched environments. Attached Figure Description

[0030] The present invention will be further described in detail below with reference to the accompanying drawings and embodiments:

[0031] Figure 1 The grain structure of the deformed high-temperature alloy prepared in the embodiments of the present invention;

[0032] Figure 2 The morphology of the precipitated phases in the deformed high-temperature alloy prepared in the embodiments of the present invention;

[0033] Figure 3 The morphology of the precipitated phase of the deformed high-temperature alloy prepared in the embodiments of the present invention;

[0034] Figure 4 The room temperature tensile strength of the deformed high-temperature alloy prepared in the embodiments of the present invention;

[0035] Figure 5 The tensile strength at 750°C of the deformed high-temperature alloy prepared in the embodiments of the present invention;

[0036] Figure 6 The oxygen-enriched combustion performance of the deformed high-temperature alloy prepared in the embodiments of the present invention. Detailed Implementation

[0037] The present invention will be further described in detail below with reference to the accompanying drawings and embodiments. The alloy compositions of embodiments 1-5 of the present invention are specifically shown in Table 2. For ease of comparison, Table 2 also lists the chemical compositions of two high-temperature alloy materials for liquid oxygen / kerosene rocket engines, comparing alloys 1 and 2. In Table 2, "remaining" in the Ni content column means "balance". The preparation steps of this deformed high-temperature alloy are as follows:

[0038] 1) Add the alloy raw materials into a vacuum induction furnace for melting treatment, and pour the melted alloy liquid into an alloy ingot (Φ80mm) without obvious casting defects.

[0039] 2) Treat the alloy ingot at a temperature range of 1140-1160℃ for 20 hours.

[0040] 3) The homogenized alloy ingot is hot forged at a temperature between 970-1100℃ and then prepared into a 40*40mm square billet through a two-forging process (final forging temperature ≥970℃, deformation controlled between 50% and 60%).

[0041] 4) The forged billet is subjected to heat treatment. The heat treatment regime is as follows: 960±20℃, hold for 1 hour, air cool, 780℃±10℃, hold for 8 hours, furnace cool to 700℃±10℃, hold for 8 hours, air cool.

[0042] Grain structure of forged square billet after heat treatment as follows Figure 1 As shown, the average grain size is approximately 20 μm. Massive carbides and short rod-shaped σ phases precipitate at the grain boundaries, such as... Figure 2 As shown. After heat treatment, spherical γ' strengthening phase precipitated, with a volume fraction of approximately 45%, as shown. Figure 3 As shown in the figure. After forging, the square billet was machined into standard tensile specimens (M10*Φ5mm) and oxy-fuel combustion specimens (Φ3.2*110mm), and tensile and oxy-fuel combustion performance tests were performed. The test results are shown in the figure. Figure 4-6 As shown.

[0043] Table 2 Alloy composition (wt.%) of embodiments of the present invention

[0044]

[0045]

[0046] Based on Comparative Alloy 1, Cu was added to obtain Examples 1 and 2, which significantly improved the combustion performance of the alloys. However, a Cu content exceeding 3% was detrimental to the tensile properties of the alloys. Adding Al to Example 1 resulted in Example 3, which significantly improved the tensile properties of the alloy, but had the worst combustion performance. Adding Ta to Example 1 resulted in Example 4, which exhibited the best tensile properties and slightly worse combustion performance, but was comparable to that of Comparative Alloy 1. Based on Example 4, replacing Mo with 1% W and Fe with Co and Ni resulted in Example 5, which showed the best combustion performance, significantly better than Comparative Alloy 1 and Comparative Alloy 2, while its room temperature tensile strength was comparable to that of Comparative Alloy 1 and Comparative Alloy 2, and its high-temperature tensile strength was superior to that of Comparative Alloy 1 and Comparative Alloy 2.

[0047] In summary, the wrought superalloy and its preparation method provided by this invention can simultaneously improve the tensile strength and oxygen-enriched combustion performance of the alloy, thereby solving the problem of the contradiction between material mechanical properties and resistance to oxygen-enriched ablation. The wrought superalloy of this invention has a very low calorific value (<1430 cal / g), which is far lower than that of the comparative alloys (1520 cal / g and 1580 cal / g). Therefore, the wrought superalloy of this invention has excellent resistance to oxygen-enriched combustion.

[0048] Matters not covered in this invention are common knowledge.

[0049] Although embodiments of the invention have been shown and described, it will be understood by those skilled in the art that various changes, modifications, substitutions and alterations can be made to these embodiments without departing from the principles and spirit of the invention, the scope of which is defined by the appended claims and their equivalents.

Claims

1. A high-strength nickel-based superalloy resistant to oxygen-rich ablation, characterized in that: The chemical composition of the nickel-based superalloy, by weight percentage, is as follows: Cr: 16.34%, Mo: 3.03%, W: 0.99%, Cu: 2.05%, Ta: 1.96%, Nb: 4.64%, Al: 0.91%, Ti: 0.61%, Co: 8.97%, Fe: 0.45%, C: 0.041%, V: 0.44%, the remainder being Ni; The high-strength, oxygen-rich ablation-resistant nickel-based superalloy has a tensile strength of ≥1250 MPa at room temperature. The high-strength, oxygen-rich ablation-resistant nickel-based superalloy has a tensile strength ≥780 MPa at 750 ℃. The high-strength, oxygen-rich ablation-resistant nickel-based superalloy exhibits a combustion length of less than 30 mm in a pure oxygen environment at room temperature below 8 MPa, with an oxygen concentration ≥ 99.5%.

2. A method for preparing a high-strength, oxygen-rich ablation-resistant nickel-based superalloy as described in claim 1, characterized in that: The preparation process steps are as follows: Step 1: Prepare ingots by vacuum induction melting, refining temperature 1500 ℃ ± 10 ℃, time 20 minutes, vacuum degree 1~20Pa, and then pour into ingots at a pouring temperature of 1450 ℃ ± 10 ℃; Step 2: The ingot is subjected to homogenization treatment at a temperature of 1150 ℃ ± 10 ℃ for a time of not less than 20 hours. Step 3: The homogenized alloy ingot is forged using a two-upsetting and two-drawing forging process to obtain a deformed high-temperature alloy with a deformation of 50% to 60%; the forging temperature is 970 ℃ to 1100 ℃. Step 4: After forging, the billet is subjected to heat treatment. The solution treatment regime is: 960 ± 20 ℃, hold for 1 h, and air cool; the aging treatment regime is: 780 ℃ ± 10 ℃, hold for 8 h, furnace cool to 700 ℃ ± 10 ℃, hold for 8 h, and air cool.

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