High-temperature-resistant nickel-based alloy for high-temperature reactor and preparation method thereof

Abstract

The invention discloses a high-temperature-resistant nickel-based alloy for a high-temperature reactor. The high-temperature-resistant nickel-based alloy comprises the following chemical components in percentage by weight: 0.05-0.1% of C; 19 to 23% of Cr; 30 to 35 percent of Ni; 0.15 to 0.6 percent of Al; 0.15 to 0.6 percent of Ti; 0.02 to 0.3 percent of Nb; 0.02 to 0.3 percent of Ta; 0.001 to 0.04 percent of N; 0.001 to 0.006 percent of Mg; the content of V is 0.1 to 0.6 percent; si < 1%; mn is less than 1.5%; and the balance of Fe. The invention further discloses a preparation method of the high-temperature-resistant nickel-based alloy for the high-temperature reactor. Compared with the prior art, the high-temperature-resistant nickel-based alloy for the high-temperature reactor has excellent high-temperature instantaneous mechanical property and long-time lasting creep property; and the design requirements of metal parts in high-temperature reactors such as a high-temperature gas cooled reactor and a gas cooled fast reactor on high-temperature instantaneous mechanical property, long-time lasting creep property and long-time service thermal aging mechanical property can be met.

Description

Technical Field

[0001] The invention belongs to the technical field of nuclear engineering, and in particular relates to a high temperature resistant nickel-based alloy for high temperature reactors and a preparation method thereof. Background Art

[0002] At present, high-temperature reactors have become the focus of research and development due to their inherent safety and reliability. However, as the operating temperature in the reactor increases, higher requirements are also placed on the strength of metal components in the core. In gas-cooled micro-reactors, gas-cooled fast reactors and other reactor types, metal components such as control rod cladding, control rod drive mechanism, neutron source cladding, and pressure vessels in the core must be used. Under the expected peak operating temperature of 750°C and high temperature conditions of up to 900°C in transient accident conditions, high requirements are placed on the high temperature resistance of these metal component materials. In the ASME III-5 "High Temperature Reactor" volume, UNSN08810 material is the material with the highest allowable temperature.

[0003] UNS N08810 alloy is a solid solution strengthened single-phase austenitic heat-resistant steel with high Cr and Ni content. It has become the recommended material in the American "ASTM" high-temperature pressure vessel design criteria due to its good high-temperature performance and structural stability. In the "ASTM" criteria, although the allowable temperature range of UNS N08810 alloy basically meets the service temperature requirements of metal parts in high-temperature reactors, its high-temperature endurance and creep strength do not meet the design requirements in the actual design conditions of many reactor types. Summary of the invention

[0004] The technical problem to be solved by the present invention is to provide a high-temperature resistant nickel-based alloy for high-temperature reactors and a preparation method thereof in view of the above-mentioned deficiencies in the prior art. Compared with the prior art, the high-temperature resistant nickel-based alloy has excellent high-temperature transient mechanical properties and long-term sustained creep properties, and can meet the design requirements of metal parts in high-temperature reactors such as high-temperature gas-cooled reactors and gas-cooled fast reactors for high-temperature transient mechanical properties, long-term sustained creep properties and long-term service thermal aging mechanical properties.

[0005] The technical solution of the present invention to solve the above technical problems is:

[0006] According to a first aspect of the present invention, there is provided a high temperature resistant nickel-based alloy for high temperature reactors, comprising the following chemical components in weight percentage:

[0007] C 0.05-0.1%; Cr 19-23%; Ni 30-35%; Al 0.15-0.6%; Ti0.15-0.6%; Nb0.02-0.3%; Ta 0.02-0.3%; N 0.001-0.04%; Mg0.001-0.006%; V 0.1-0.6%; Si<1%; Mn<1.5%; the balance is Fe.

[0008] Preferably, the high temperature resistant nickel-based alloy comprises the following chemical components in weight percentage:

[0009] C 0.09-0.1%; Cr 19-21%; Ni 33-35%; Al 0.5-0.6%; Ti 0.2-0.45%; Nb0.06-0.1%; Ta 0.06-0.1%; N 0.007-0.03%; Mg 0.001-0.003%; V 0.1-0.15%; Si 0.1-0.7%; Mn 0.5-1%; the balance is Fe.

[0010] Preferably, the high temperature resistant nickel-based alloy comprises a fully austenite phase matrix, granular MC-type carbides finely dispersedly distributed in the crystals of the fully austenite phase matrix, and M-type carbides finely and discretely distributed at the grain boundaries of the fully austenite phase matrix. 23 C6 type carbide.

[0011] Preferably, the grain size of the full austenite phase matrix is ​​coarser than grade 5.0.

[0012] Preferably, the particle size of the MC type carbide is ≤0.5 μm, 23 C6 carbide is granular, M 23 The number density of C6 carbides is not less than 6000 / mm, M 23 The average size of C6 carbides is ≤0.1μm;

[0013] And / or, the intragranular B-type fine TiN inclusions of the fully austenite phase matrix are ≤2.0 level, the D-type fine TiN inclusions are ≤2.5 level, and the size of M (C, N) type carbonitrides is ≤2 μm;

[0014] and / or, heat aging at 650°C for 3000 h in atmospheric conditions 23 C6 carbides are granular or short rod-shaped, with a number density of 300-400 / mm and an average size of ≤2μm.

[0015] Preferably, 4.5*(C content)≤(Nb+Ta+Ti+V)≤9*(C content);

[0016] And / or, the content ratio of Cr / V is 120≤≤230.

[0017] According to a second aspect of the present invention, a method for preparing a high temperature resistant nickel-based alloy comprises:

[0018] According to the chemical composition weight percentage of C 0.05-0.1%, Cr 19-23%, Ni 30-35%, Al 0.15-0.6%, Ti 0.15-0.6%, Nb 0.02-0.3%, Ta 0.02-0.3%, N 0.001-0.04%, Mg 0.001-0.006%, V 0.1-0.6%, Si<1%, Mn<1.5%, and the balance is Fe, raw materials are taken and smelted to obtain alloy ingots;

[0019] The alloy ingot is subjected to high temperature homogenization treatment, forging, and hot rolling in sequence to form a bar;

[0020] The rod is subjected to solution treatment and aging treatment in sequence to obtain a high temperature resistant nickel-based alloy.

[0021] Preferably, the smelting process adopts a double vacuum melting method of vacuum induction melting and vacuum consumable remelting, wherein the rising rate of the melt level during the vacuum consumable remelting process is controlled at 5.5-6.5 mm / min.

[0022] Preferably, the temperature of the high temperature homogenization treatment is 1220-1240° C., and the holding time is 15-25 hours.

[0023] Preferably, the holding temperature of the forging process is 1220-1240°C, the time interval between the first 5-10 forgings after each furnace is not shorter than 1-5s, and the final forging temperature is not lower than 1000°C.

[0024] Preferably, the holding temperature of the hot rolling process is 1180-1220°C, the deformation of the first pass after each furnace is not more than 20%; and / or, the deformation of the second pass is not more than 35%; and / or, the deformation of subsequent passes is 30-60%; and / or, the final rolling temperature is not less than 1000°C; and / or, water cooling is used after rolling is completed.

[0025] Preferably, the temperature of the solution treatment is 1121-1220° C.; and / or, the time of the solution treatment is 2-4 h; and / or, water cooling is adopted after the solution treatment.

[0026] Preferably, the temperature of the aging treatment is 600-630° C.; and / or, the time of the aging treatment is 2-4 h; and / or, water cooling is used after the aging treatment is completed.

[0027] Beneficial effects:

[0028] The invention discloses a high-temperature resistant nickel-based alloy for a high-temperature reactor and a preparation method thereof. The room temperature yield strength Rp0.2 of the high-temperature resistant nickel-based alloy in a standard heat-treated state is ≥210MPa, and the tensile strength Rm is ≥543MPa; the yield strength Rp0.2 of high-temperature tensile test at 650°C is ≥243MPa, and the tensile strength Rm is ≥488MPa; the yield strength Rp0.2 of high-temperature tensile test at 900°C is ≥98MPa, and the tensile strength Rm is ≥138MPa; the 30000h-650°C creep 1% total strain stress value is 107MPa, and the performance can be improved by more than 20% compared with the mature UNS N08810 material on the market, and the alloy has excellent high-temperature transient mechanical properties and long-term sustained creep properties, and can meet the design requirements of metal parts in high-temperature reactors such as high-temperature gas-cooled reactors and gas-cooled fast reactors for high-temperature transient mechanical properties, long-term sustained creep properties, and long-term service thermal aging mechanical properties. BRIEF DESCRIPTION OF THE DRAWINGS

[0029] Figure 1 This is a standard heat-treated microstructure of the high-temperature resistant nickel-based alloy for high-temperature reactors in Preparation Example 1 of the present invention;

[0030] Figure 2 This is a standard heat-treated microstructure of the high-temperature resistant nickel-based alloy for high-temperature reactors in Preparation Example 2 of the present invention;

[0031] Figure 3 This is a standard heat-treated microstructure of the high-temperature resistant nickel-based alloy for high-temperature reactors in Preparation Example 3 of the present invention;

[0032] Figure 4 This is a standard heat-treated microstructure of the high-temperature-resistant nickel-based alloy for high-temperature reactors in Preparation Example 4 of the present invention;

[0033] Figure 5 This is a standard heat-treated microstructure morphology of the UNS N08810 alloy in Comparative Example 1 of the present invention. DETAILED DESCRIPTION

[0034] In order to enable those skilled in the art to better understand the technical solution of the present invention, the technical solution of the present invention will be clearly and completely described below in conjunction with the drawings in the present invention. Obviously, the described embodiments are part of the embodiments of the present invention, not all embodiments. Based on the embodiments of the present invention, all other embodiments obtained by ordinary technicians in this field without making creative work belong to the protection scope of the present invention.

[0035] In view of the problem that the materials that can be used for metal parts in high-temperature reactors in the prior art have good high-temperature performance and structural stability and can meet the high-temperature mechanical properties, their high-temperature durability, creep and other strengths do not meet the design requirements in the actual design conditions of many reactor types. The present invention discloses a high-temperature resistant nickel-based alloy, comprising the following chemical components in weight percentage: C 0.05-0.1%; Cr 19-23%; Ni 30-35%; Al 0.15-0.6%; Ti 0.15-0.6%; Nb 0.02-0.3%; Ta 0.02-0.3%; N 0.001-0.04%; Mg 0.001-0.006%; V 0.1-0.6%; Si <1%; Mn <1.5%; the balance is Fe.

[0036] Correspondingly, the present invention also discloses a method for preparing a high temperature resistant nickel-based alloy for a high temperature reactor, comprising: taking raw materials according to the chemical composition weight percentage of C 0.05-0.1%, Cr 19-23%, Ni 30-35%, Al 0.15-0.6%, Ti 0.15-0.6%, Nb 0.02-0.3%, Ta 0.02-0.3%, N 0.001-0.04%, Mg 0.001-0.006%, V 0.1-0.6%, Si <1%, Mn <1.5%, and the balance is Fe, and smelting to obtain an alloy ingot;

[0037] The alloy ingot is subjected to high temperature homogenization treatment, forging, and hot rolling in sequence to form a bar;

[0038] The rod is subjected to solution treatment and aging treatment in sequence to obtain a high temperature resistant nickel-based alloy.

[0039] Example 1

[0040] like Figure 1 As shown, this embodiment discloses a high temperature resistant nickel-based alloy for high temperature reactors, comprising the following components in weight percentage:

[0041] C 0.05-0.1%; Cr 19-23%; Ni 30-35%; Al 0.15-0.6%; Ti0.15-0.6%; Nb0.02-0.3%; Ta 0.02-0.3%; N 0.001-0.04%; Mg0.001-0.006%; V 0.1-0.6%; Si<1%; Mn<1.5%; balance Fe.

[0042] In this embodiment, the high temperature resistant nickel-based alloy for high temperature reactor (hereinafter referred to as the present alloy) is specifically a high temperature resistant austenitic steel, which includes a fully austenitic phase matrix. In addition, the present alloy contains MC type carbide forming elements such as C and Ti, which can form MC type carbides during the smelting solidification process and the long-term thermal aging process. In addition, the present alloy can form fine granular or short rod-shaped semi-continuously distributed MC carbides at the grain boundary when it is in a high temperature (650-750°C) environment for a long time. 23 C6 type carbides, that is, the alloy also includes granular MC type carbides finely dispersed in the crystals of the full austenite phase matrix, and M type carbides finely and discretely distributed at the grain boundaries of the full austenite phase matrix. 23 C6 type carbide.

[0043] Compared with the prior art, the alloy contains multiple alloying elements such as C, Cr, Ni, and Ti. Through the coordinated optimization of multiple alloying elements, the solid solution strengthening effect can be improved, the alloy strength can be improved, and the dispersion strengthening effect can be improved by increasing the number of finely dispersed granular MC-type carbides in the crystal, thereby obtaining good high-temperature mechanical properties. 23 C6 carbide can improve the grain boundary strength, ultimately obtaining excellent high temperature transient mechanical properties and long-term creep properties.

[0044] Specifically, the C element can achieve the purpose of solid solution strengthening by forming interstitial solid solution, and increasing the C content can significantly improve the strength of the alloy. At the same time, the supersaturated C element can be combined with MC-type carbide-forming elements such as Ti to form MC-type carbides during the smelting solidification process and the long-term thermal aging process. The growth rate of MC-type carbides is low, and they can be dispersed and precipitated in a relatively small size during the long-term thermal aging process. MC-type carbides can hinder dislocation movement and have a strong precipitation strengthening effect, thereby improving the long-term high-temperature mechanical properties of the alloy. The present invention has been verified by a large number of experiments, and the C content is preferably controlled at 0.09-0.1%.

[0045] Ti is the main element forming MC carbide. If the Ti content is too low, the precipitation of MC carbide is insufficient, and the long-term high-temperature mechanical properties of the alloy cannot be guaranteed; while if the Ti content is too high, the MC carbide will coarsen, which will damage the long-term thermal aging performance and welding performance of the alloy. The present invention has been verified by a large number of experiments, and the Ti content is preferably not less than 0.2% and not higher than 0.45%.

[0046] At the same time, an appropriate amount of Nb and Ta elements are also added to the alloy. Nb and Ta elements have a strong affinity with C elements, which can rob some C atoms in the crystal to generate more stable MC-type carbides containing Nb and Ta, so that the carbides in the crystal are refined. In addition, Ti, Nb and Ta are all strong MC-type carbide-forming elements. The simultaneous addition of the three elements can produce a synergistic effect, which can more effectively promote the fine dispersion precipitation of MC-type carbides and better improve the long-term high-temperature mechanical properties of the alloy. The present invention has been verified by a large number of experiments. The Nb content is preferably controlled at 0.06-0.1%, and the Ta content is preferably controlled at 0.06-0.1%. Otherwise, the content is too low to achieve sufficient MC refinement effect, and the content is too high to easily lead to the precipitation of harmful phases such as Laves phase.

[0047] Cr element can reduce M 23 The precipitation tendency of C6 type carbides reduces the grain boundary M 23 Coarsening of C6 carbides during long-term aging. The present invention has been verified by a large number of experiments, and the Cr content is preferably controlled at 19-21%.

[0048] In addition, considering that the number and size of carbides at the grain boundaries of the alloy increase after the C content is increased, the carbides at the grain boundaries are in the shape of long rods with large continuous distribution, which are prone to stress concentration when subjected to force, inducing the formation of holes and cracks, and leading to problems such as reduced service performance of the alloy. In order to solve the above problems, on the one hand, by adding Ti, Nb, and Ta elements to fix the C element to a certain extent, MC type carbides are formed, and the M at the grain boundaries is reduced. 23 On the other hand, in order to improve the morphology and distribution of grain boundary carbides, appropriate amounts of Mg and V are added to this alloy.

[0049] Mg is a grain boundary strengthening element, which tends to be concentrated at the grain boundary, occupying some of the atomic sites at the grain boundary, reducing the C atom content at the grain boundary, and thus reducing M 23 The purpose of C6 particle size. V is a strong carbide-forming element. The atomic radius of V (0.135nm) is close to the atomic radius of Cr (0.127nm). It is easy to interact with Cr and replace M 23 The Cr site in C6 forms a Cr-V bond with short bond length and strong bond energy, which reduces the carbide lattice constant and promotes M 23 C6 carbide undergoes lattice transformation during long-term service. 23 The surface of C6 carbide particles forms vanadium carbide particles with better stability. 23 C6 carbides are encapsulated to limit the M 23 The size of C6 carbides grows excessively. The combined addition of Mg and V can effectively extend the M during the long-term service of this alloy. 23The precipitation incubation period of C6 carbides, the refinement of carbide precipitation particle size, and the inhibition of carbide coarsening and growth significantly reduce the performance degradation tendency of the alloy during long-term service. Considering that the content of Mg and V is too low to play a sufficient role in M 23 The Mg content is preferably controlled at 0.001-0.003%, and the V content is preferably controlled at 0.1-0.15%.

[0050] The N element is a strong austenite forming and stabilizing element, which can expand the range of the austenite phase region in the phase diagram. The atomic radius of N is relatively small, and it is easy to enter the crystal interstitial position. As an interstitial atom, it plays a role in solid solution strengthening, effectively improving the high-temperature instantaneous and high-temperature long-term mechanical properties of the alloy. Considering that too little N content cannot play an effective solid solution strengthening role, and at the same time, due to the strong affinity between N and Ti elements, it is easy to form larger TiN or Ti (C, N), although nitrides can pin the grain boundaries to limit grain growth and play a role in fine grain strengthening, but too much nitrides will cause the alloy to have excessive inclusion levels, mixed crystals, and increase the volatility of the alloy's high-temperature long-term performance. Therefore, the N content is not easy to be too low, and it is preferably controlled at 0.007wt.%-0.03wt.%.

[0051] The present alloy contains high Ni content and Al content. Ni element is the main solid solution strengthening element in the present alloy, and increasing Ni content helps to improve the strength of the alloy. In the present embodiment, the Ni content is preferably controlled at 33-35%, otherwise, if the Ni content is too low, the solid solution strengthening effect will be insufficient and the stability of the alloy structure will be reduced, while if the Ni content is too high, other second phases may be precipitated.

[0052] Al element has the effect of improving oxidation resistance and solid solution strengthening. In this embodiment, in order to reduce M 23 The precipitation tendency of C6 type carbides reduces the Cr content, and the reduction of Cr content may lead to problems such as reduced oxidation resistance of the alloy. Al and Cr are both ferrite forming elements. Al can increase the diffusion rate of Cr. By increasing the Al content, the oxidation resistance of the alloy can be guaranteed. In this embodiment, the Al content is preferably controlled at Al 0.5-0.6%.

[0053] Si and Mn act as deoxidizers, which are helpful for precipitation deoxidation of the alloy during the welding process. In order to ensure the deoxidation effect, the Si and Mn contents should not be too low. The Si content is preferably greater than 0.1%, and the Mn content is preferably greater than 0.5%. However, too much Si element can easily lead to the formation of eutectic phase in the weld joint structure, which promotes the formation of liquefaction cracks in the heat-affected zone. In order to ensure the weldability of the alloy, the Si content needs to be strictly controlled. The Si content in the alloy is preferably not higher than 0.7%. At the same time, Mn is an easily oxidized element, and since the alloy is used to control Mn, the Si content in the alloy should be less than 0.7%. 23 The precipitation tendency of C6 type carbides reduces the Cr content. In order to ensure the integrity and continuity of the oxide film on the surface of the material during oxidation or corrosion, the Mn content needs to be strictly controlled to avoid external oxidation of the Mn element. Therefore, the Mn content in the alloy is preferably controlled below 1%.

[0054] In some embodiments, the grain size of the fully austenite phase matrix is ​​preferably coarser than 5.0 to ensure the long-term creep performance of the alloy.

[0055] In some embodiments, in order to obtain a high-temperature resistant nickel-based alloy for a high-temperature reactor with more excellent high-temperature mechanical properties and long-lasting creep properties, and to meet the service mechanical properties requirements of metal parts in a high-temperature reactor, the particle size of the MC-type carbide is preferably ≤0.5 μm. In the standard heat treatment state, M 23 C6 carbide is granular, M 23 The number density of C6 type carbides is preferably not less than 6000 / mm, M 23 The average size of C6 carbides is preferably ≤0.1 μm; and / or, the intragranular B-type fine TiN inclusions of the full austenite phase matrix are ≤2.0 level, the D-type fine TiN inclusions are ≤2.5 level, and the size of M(C,N)-type carbonitrides is ≤2 μm; and / or, by refining the grain boundary M 23 The size of C6 carbides is controlled, and the amount of precipitation is controlled, so that the alloy obtains better thermal aging mechanical properties. Specifically, the M 23 C6 carbides are in the form of granules or short rods, with a number density of 300-400 / mm and an average size of ≤2μm. This can not only precipitate small discrete granular M 23 C6 carbide can improve the grain boundary strength and avoid the M 23 C6 carbides coarsen.

[0056] In some embodiments, 4.5*(C content)≤(Nb+Ta+Ti+V) content≤9*(C content); and / or, 120≤Cr / V content ratio≤230.

[0057] Specifically, adjusting the relationship between the content of (Nb+Ta+Ti+V) and the content of C to 4.5*(content of C)≤(content of Nb+Ta+Ti+V)≤9*(content of C) can promote the formation of MC-type carbides and control the formation of M 23 The precipitation amount of C6 type carbides. Nb, Ta and Ti are all strong MC type carbide forming elements. While increasing the precipitation amount of MC type carbides, they consume C element appropriately to control the M in this alloy. 23 The grain boundary precipitation state of C6 type carbides can make the grain boundary M of this alloy in the service process 23 The coarsening tendency of C6 carbides is reduced, thereby controlling M 23 The number density of C6 type carbides is to avoid stress concentration at the grain boundary and reduce the mechanical properties. If the addition amount of Nb, Ta and Ti is too small, it will not only lead to insufficient formation of MC type carbides, but also cause M in the service process. 23 The coarsening of C6 carbides reduces the mechanical properties of the alloy. If the addition of Nb, Ta and Ti is too much, it will not only lead to the aggregation and coarsening of MC carbides and the precipitation of harmful intermetallic compounds, but also cause the grain boundary M 23 The amount of C6 carbides is insufficient, which reduces the grain boundary strength of the alloy and deteriorates the mechanical properties of the alloy.

[0058] Adjusting the relationship between Cr content and V content to 120≤Cr / V content ratio≤230 can control the grain boundary M 23 The size and precipitation morphology of C6 carbides make M 23 The precipitation form of C6 carbide is fine particles or short rods. Too high Cr / V content ratio is not conducive to the 23 The surface of C6 carbide forms vanadium carbide particles with better stability, which has a great influence on the M 23 The coarsening restriction effect of C6 carbides is limited, and it is difficult to effectively inhibit the M 23 C6 carbides grow up, resulting in long rods of large size. 23 The formation of C6 type carbides deteriorates the long-term service performance of the alloy; while the Cr / V content ratio is too low, which may lead to a low Cr content in the alloy and a decrease in corrosion resistance, or a high V content, where the V element preferentially forms a large amount of VC type carbides with C and cannot enter M 23 C6 carbide controls its precipitation morphology.

[0059] It should be noted that in order to distribute a large number of fine and dispersed granular MC-type carbides in the crystal of the alloy, with the particle size of the MC-type carbides being ≤0.5 μm, the present invention adopts at least the following three measures:

[0060] First, increase the C content to 0.09-0.1%, and add strong MC-type carbide-forming elements (Ti 0.2-0.45%, Nb 0.06-0.1%, Ta 0.06-0.1%) at the same time. The simultaneous addition of multiple strong MC-type carbide-forming elements can produce a synergistic effect and promote the fine dispersion and precipitation of MC-type carbides;

[0061] Second, the total content of Nb, Ta, Ti, and V in the alloy is controlled: 4.5*(C content)≤(Nb+Ta+Ti+V)≤9*(C content) to promote the formation of MC-type carbides and regulate the precipitation size;

[0062] Third, the temperature of the high-temperature homogenization treatment during the preparation of the alloy is strictly controlled to be 1220-1240°C and the holding time is 15-25h to dissolve the primary large-sized MC carbides and ensure the secondary fine and uniform precipitation of MC carbides.

[0063] It should be noted that in order to precipitate fine and discrete granular M 23 C6 carbide improves the grain boundary strength and avoids the formation of M during the thermal aging process of heat-resistant steel. 23 The C6 carbides coarsen, making the M 23 C6 type grain boundary carbides are in the form of particles or short rods, with a number density of 300 / mm and an average size of ≤2μm. The present invention mainly adopts the following three measures:

[0064] First, the addition of Mg 0.02-0.06% as a grain boundary segregation element and V 0.1-0.15% as a strong carbide forming element jointly refines the M 23 C6 particle size;

[0065] Second, control the Cr / V content ratio in the alloy: 120≤Cr / V content ratio≤230 to control the grain boundary M 23 Size and precipitation morphology of C6 type carbides;

[0066] Third, aging treatment is performed with an aging temperature of 600-630°C and an aging time of 2-4 hours. After the aging treatment, water cooling is used to obtain a higher amount of grain boundary carbides to avoid service coarsening.

[0067] Example 2

[0068] This embodiment discloses a high temperature resistant nickel-based alloy for high temperature reactors. Compared with the embodiment 1, the difference is that the high temperature resistant nickel-based alloy for high temperature reactors described in the embodiment 2 includes the following chemical components in weight percentage:

[0069] C 0.09-0.1%; Cr 19-21%; Ni 33-35%; Al 0.5-0.6%; Ti 0.2-0.45%; Nb0.06-0.1%; Ta 0.06-0.1%; N 0.007-0.03%; Mg 0.001-0.003%; V 0.1-0.15%; Si 0.1-0.7%; Mn 0.5-1%; the balance is Fe.

[0070] During the service of metal parts in high-temperature reactors, the coarsening of grain boundary carbides is the main reason for the reduction of thermal aging mechanical properties of metal parts in high-temperature reactors. The high-temperature resistant nickel-based alloy for high-temperature reactors described in Examples 1 and 2 has a room temperature yield strength Rp0.2≥210MPa and a tensile strength Rm≥543MPa in the standard heat-treated state; a yield strength Rp0.2≥243MPa and a tensile strength Rm≥488MPa at 650℃; a yield strength Rp0.2≥98MPa and a tensile strength Rm≥138MPa at 900℃. In addition, the total strain stress value of 1% creep at 30000h-650℃ is 107MPa, which can improve the performance of the mature UNSN08810 material on the market by more than 20%, and can suppress M during long-term service at high temperatures (650-750℃). 23 The excessive coarsening of C6 carbides enables the alloy to maintain excellent thermal aging mechanical properties for a long time, and has excellent high-temperature transient mechanical properties and long-term creep performance, which can meet the design requirements of metal parts in high-temperature reactors such as high-temperature gas-cooled reactors and gas-cooled fast reactors, including high-temperature transient mechanical properties, long-term creep properties and long-term service thermal aging mechanical properties.

[0071] Example 3

[0072] This embodiment discloses a method for preparing the high temperature resistant nickel-based alloy for high temperature reactors described in Embodiments 1 and 2, comprising:

[0073] S1, according to the weight percentage of each chemical component in the high temperature resistant nickel-based alloy: C0.05-0.1%, Cr 19-23%, Ni 30-35%, Al 0.15-0.6%, Ti 0.15-0.6%, Nb 0.02-0.3%, Ta 0.02-0.3%, N0.001-0.04%, Mg 0.001-0.006%, V 0.1-0.6%, Si<1%, Mn<1.5%, and the balance is Fe, take raw materials, and smelt them to obtain alloy ingots.

[0074] In some embodiments, the smelting process adopts a double vacuum melting method of vacuum induction melting plus vacuum consumable remelting, and the rising rate of the melt level during the vacuum consumable remelting process is controlled at 5.5-6.5 mm / min to control the remelting rate.

[0075] Specifically, since the high-temperature resistant nickel-based alloy for high-temperature reactors (hereinafter referred to as the present alloy) prepared by the method of this embodiment contains a variety of volatile and easily burnt elements such as N and Mg, a slow melting rate will lead to serious element burnout, and a fast melting rate will cause the elements to easily segregate, affecting the uniformity of the structure. When the melt level rising rate is controlled at 5.5-6.5 mm / min, an appropriate remelting melting rate can be obtained to ensure that the element content is controlled within the composition range.

[0076] S2, the alloy ingot is then subjected to high temperature homogenization treatment, forging, and hot rolling in sequence to form a rod;

[0077] S3, performing solution treatment and aging treatment on the rod in sequence to obtain a high temperature resistant nickel-based alloy for high temperature reactors.

[0078] In some embodiments, the temperature of the high temperature homogenization treatment is 1220-1240°C, and the holding time is 15-25h. By performing the high temperature homogenization treatment at 1220-1240°C, not only can the elements such as Nb, Ta and Ti in the alloy with a large segregation tendency and a slow diffusion rate be fully diffused and uniformly diffused in the ingot, but also the coarse primary MC carbides precipitated by the above elements can be dissolved, ensuring the secondary fine dispersion and uniform precipitation of MC carbides, thereby eliminating the element segregation and coarse precipitation phase in the alloy ingot.

[0079] In some embodiments, the holding temperature of the forging process is 1220-1240°C, the time interval between the first 5-10 forgings after each furnace is not shorter than 1-5s, and the final forging temperature is not lower than 1000°C. This can not only ensure the precipitation of MC-type carbides during the deformation process, but also avoid a significant increase in the internal temperature of the ingot, and avoid deformation heat causing over-melting of the grain boundaries in the first few forgings.

[0080] In some embodiments, the holding temperature of the hot rolling process is 1180-1220°C, the deformation of the first pass after each furnace is not more than 20%; and / or, the deformation of the second pass is not more than 35%; and / or, the deformation of subsequent passes is 30-60%; and / or, the final rolling temperature is not less than 1000°C; and / or, water cooling is used after rolling is completed.

[0081] In some embodiments, considering that the selection of the solution treatment system is related to the smelting composition (especially the nitrogen content) and the microstructure after thermal deformation, the solution treatment temperature is 1121-1220°C, preferably 1160-1220°C, according to the element composition and its ratio in the alloy (for example, the N content and the carbide-forming element (Ti, Nb, Ta) content); and / or, the solution treatment time is 2-4h; and / or, water cooling is adopted after completing the solution treatment.

[0082] Compared with the conventional technology, the method of this embodiment adopts a higher solution temperature to dissolve the precipitated VC carbide, so that V enters M during the service thermal aging process. 23 C6 type carbide to control M 23 The size and shape of C6 carbides ensure that the grain size of the obtained alloy is coarser than grade 5.0.

[0083] In some embodiments, the temperature of the aging treatment is 600-630° C.; and / or, the time of the aging treatment is 2-4 hours; and / or, water cooling is used after the aging treatment is completed.

[0084] Compared with the traditional method, the method of this embodiment adopts a lower aging temperature of 580-630°C, which can not only promote the pre-formation of a large number of small-sized carbide core particles at the grain boundary of the prepared alloy before service, but also reduce the individual coarsening tendency of carbides during service; in addition, it can also avoid M 23 The growth of tiny cores of C6 carbides causes Cr depletion near the grain boundaries.

[0085] Several groups of preparation examples are provided below to describe in detail the preparation method of the high temperature resistant nickel-based alloy for high temperature reactors of this embodiment, as follows:

[0086] Preparation Example 1

[0087] This preparation example discloses a method for preparing a high temperature resistant nickel-based alloy for high temperature reactors, and the steps are as follows:

[0088] First, raw materials were taken according to the alloy chemical composition shown in Preparation Example 1 in Table 1, and double vacuum melting of vacuum induction melting and vacuum consumable remelting was adopted, wherein the rising rate of the melt level during the vacuum consumable remelting was controlled at 6.0 mm / min, to obtain an alloy ingot.

[0089] Then, the alloy ingot is subjected to high-temperature homogenization treatment, forging and hot rolling, wherein the temperature of the high-temperature homogenization treatment is 1230°C and the holding time is 20 hours; the forging temperature is 1235°C, the time interval between the first 8 forgings after each furnace is not less than 3s, and the final forging temperature is not lower than 1100°C; the holding temperature of the hot rolling process is 1200°C, the first deformation amount after each furnace is not more than 20%, the second deformation amount is not more than 35%, the subsequent deformation amount is 30-60%, the final rolling temperature is not lower than 1000°C, and water cooling is used after rolling to obtain rods.

[0090] Then, the rod is subjected to solution treatment and aging treatment in sequence, wherein the temperature of the solution treatment is 1170°C, the time of the solution treatment is 3 hours, the grain size is controlled at level 4-5, and water cooling is adopted after the solution treatment; the temperature of the aging treatment is 615°C, the time of the aging treatment is 3 hours; and / or, water cooling is adopted after the aging treatment to obtain a high temperature resistant nickel-based alloy for high temperature reactors.

[0091] Table 1 Alloy chemical composition (wt.%)

[0092]

[0093]

[0094] Preparation Example 2

[0095] This preparation example discloses a method for preparing a high temperature resistant nickel-based alloy for high temperature reactors. The steps are basically the same as those of Preparation Example 1, except that the chemical compositions of the alloys are different, as shown in Preparation Example 2 in Table 2.

[0096] Preparation Example 3

[0097] This preparation example discloses a method for preparing a high temperature resistant nickel-based alloy for high temperature reactors. The steps are basically the same as those of Preparation Example 1, except that the chemical compositions of the alloys are different, as shown in Preparation Example 3 in Table 2.

[0098] Preparation Example 4

[0099] This preparation example discloses a method for preparing a high temperature resistant nickel-based alloy for high temperature reactors. The steps are basically the same as those of Preparation Example 1, except that the chemical compositions of the alloys are different, as shown in Preparation Example 4 in Table 2.

[0100] Comparative Example 1

[0101] Comparative Example 1 is a mature UNS N08810 alloy on the market, and its composition is shown in Comparative Example 1 in Table 1.

[0102] The standard heat treatment microstructures of the high temperature resistant nickel-based alloys for high temperature reactors prepared in Preparation Examples 1-4 and the comparative example UNS N08810 alloy were tested, and the results are as follows: Figure 1-5 shown.

[0103] Depend on Figure 1-Figure 5 It can be seen that the standard heat-treated microstructures of the high-temperature resistant nickel-based alloys for high-temperature reactors prepared in Preparation Examples 1 and 2 are similar, and the standard heat-treated microstructures of Preparation Examples 3 and 4 are similar. In addition, the standard heat-treated microstructures of the high-temperature resistant nickel-based alloys for high-temperature reactors prepared in Preparation Examples 1-4 all have obvious granular M 23 C6 type grain boundary carbides, calculated M 23The number density of C6 carbides is not less than 6000 / mm, the average size is ≤0.1μm, the size of intragranular MC carbides is less than 0.5μm, and the intragranular and grain boundary strengthening effects are strong. However, in comparative example 1, the size of intragranular MC carbides is greater than 1μm, and no grain boundary carbides are precipitated, and the intragranular and grain boundary strengthening effects are weak.

[0104] In addition, the high temperature resistant nickel-based alloys for high temperature reactors prepared in Preparation Examples 1-4 and the comparative example UNS N08810 alloy were subjected to the following performance tests under the same conditions:

[0105] (1) Tensile property test: The tensile property tests were performed on the comparative example 1 and preparation examples 1-4 at room temperature, 650°C and 900°C, respectively. The test results are shown in Table 2.

[0106] As shown in Table 2, compared with the mature UNS N08810 alloy (Comparative Example 1) on the market, the room temperature and high temperature mechanical properties of the high temperature reactor resistant nickel-based alloys prepared in Preparation Examples 1-4 are better, especially Preparation Examples 1 and Preparation Examples 2, which can be used as preferred solutions. The reason why the room temperature and high temperature mechanical properties of the alloys in Preparation Examples 1-4 are better is mainly attributed to the synergistic optimization of multiple alloying elements in Preparation Examples 1-4, which can not only improve the solid solution strengthening effect, but also improve the dispersion precipitation strengthening and grain boundary strengthening effects.

[0107] Table 2 Tensile properties test results

[0108]

[0109] (2) Creep performance: The creep performance at 10,000 h was fitted and extrapolated to the creep performance at 30,000 h. The stress value of the high-temperature resistant nickel-based alloy for high-temperature reactors prepared in Preparation Examples 1-4 at 30,000 h-650°C creep 1% total strain was greater than or equal to 107 MPa, which was 20% higher than that of Comparative Example 1.

[0110] (3) Thermal aging mechanical properties: After testing, the high temperature resistant nickel-based alloy for high temperature reactor prepared in Preparation Examples 1-4 has a 650°C tensile yield strength Rp0.2≥405MPa and a tensile strength Rm≥531MPa after thermal aging for 3000 hours under 650°C atmospheric conditions.

[0111] In summary, the high-temperature resistant nickel-based alloy for high-temperature reactors prepared by the preparation method of this embodiment has significantly improved mechanical properties compared with the prior art, and has excellent high-temperature transient mechanical properties and long-term sustained creep properties, which can meet the design requirements of metal parts in high-temperature reactors such as high-temperature gas-cooled reactors and gas-cooled fast reactors currently under development, including high-temperature transient mechanical properties, long-term sustained creep properties and long-term service thermal aging mechanical properties.

[0112] It is to be understood that the above embodiments are merely exemplary embodiments used to illustrate the principles of the present invention, but the present invention is not limited thereto. For those of ordinary skill in the art, various modifications and improvements can be made without departing from the spirit and essence of the present invention, and these modifications and improvements are also considered to be within the scope of protection of the present invention.

Claims

1. A high temperature resistant nickel-based alloy for high temperature reactors, characterized in that: The chemical composition includes the following weight percentages: C 0.05-0.1%; Cr 19-23%; Ni 30-35%; Al 0.15-0.6%; Ti0.15-0.6%; Nb 0.02-0.3%; Ta 0.02-0.3%; N 0.001-0.04%; Mg0.001-0.006%; V 0.1-0.6%; Si<1%; Mn<1.5%; the balance is Fe.

2. The high temperature resistant nickel-based alloy for high temperature reactor according to claim 1, characterized in that: The chemical composition includes the following weight percentages: C 0.09-0.1%; Cr 19-21%; Ni 33-35%; Al 0.5-0.6%; Ti 0.2-0.45%; Nb 0.06-0.1%; Ta 0.06-0.1%; N 0.007-0.03%; Mg 0.001-0.003%; V 0.1-0.15%; Si 0.1-0.7%; Mn 0.5-1%; the balance is Fe.

3. The high temperature resistant nickel-based alloy for high temperature reactor according to claim 1 or 2, characterized in that: The high temperature resistant nickel-based alloy comprises a full austenite phase matrix, granular MC type carbides finely dispersedly distributed in the crystals of the full austenite phase matrix, and M 23 C6 type carbide.

4. The high temperature resistant nickel-based alloy for high temperature reactor according to claim 3, characterized in that: The grain size of the full austenite phase matrix is ​​coarser than grade 5.

0.

5. The high temperature resistant nickel-based alloy for high temperature reactor according to claim 4, characterized in that: The particle size of the MC type carbide is ≤0.5 μm. 23 C6 carbide is granular, M 23 The number density of C6 carbides is not less than 6000 / mm, M 23 The average size of C6 carbides is ≤0.1μm; And / or, the intragranular B-type fine TiN inclusions of the fully austenite phase matrix are ≤2.0 level, the D-type fine TiN inclusions are ≤2.5 level, and the size of M (C, N) type carbonitrides is ≤2 μm; and / or, heat aging at 650°C for 3000 h in atmospheric conditions 23 C6 carbides are granular or short rod-shaped, with a number density of 300-400 / mm and an average size of ≤2μm.

6. The high temperature resistant nickel-based alloy for high temperature reactor according to claim 1 or 2, characterized in that: 4.5*(C content)≤(Nb+Ta+Ti+V)≤9*(C content); And / or, the content ratio of Cr / V is 120≤≤230.

7. A method for preparing a high temperature resistant nickel-based alloy for a high temperature reactor, characterized in that: include: According to the chemical composition weight percentage of C 0.05-0.1%, Cr 19-23%, Ni 30-35%, Al 0.15-0.6%, Ti0.15-0.6%, Nb 0.02-0.3%, Ta 0.02-0.3%, N 0.001-0.04%, Mg 0.001-0.006%, V 0.1-0.6%, Si<1%, Mn<1.5%, and the balance is Fe, raw materials are taken and smelted to obtain alloy ingots; The alloy ingot is subjected to high temperature homogenization treatment, forging, and hot rolling in sequence to form a bar; The rod is subjected to solution treatment and aging treatment in sequence to obtain a high temperature resistant nickel-based alloy.

8. The method for preparing a high temperature resistant nickel-based alloy for a high temperature reactor according to claim 7, characterized in that: The smelting process adopts a double vacuum melting method of vacuum induction melting and vacuum consumable remelting, wherein the rising rate of the melt level during the vacuum consumable remelting process is controlled at 5.5-6.5mm / min.

9. The method for preparing a high temperature resistant nickel-based alloy for a high temperature reactor according to claim 7, characterized in that: The temperature of high temperature homogenization treatment is 1220-1240℃, and the insulation time is 15-25h.

10. The method for preparing a high temperature resistant nickel-based alloy for a high temperature reactor according to claim 7, characterized in that: The holding temperature of the forging process is 1220-1240℃, the time interval between the first 5-10 forgings after each furnace is not shorter than 1-5s, and the final forging temperature is not lower than 1000℃.

11. The method for preparing a high temperature resistant nickel-based alloy for a high temperature reactor according to claim 7, characterized in that: The holding temperature of the hot rolling process is 1180-1220°C, the deformation of the first pass after each furnace is not more than 20%; and / or, the deformation of the second pass is not more than 35%; and / or, the deformation of subsequent passes is 30-60%; and / or, the final rolling temperature is not less than 1000°C; and / or, water cooling is used after rolling is completed.

12. The method for preparing a high temperature resistant nickel-based alloy for a high temperature reactor according to claim 7, characterized in that: The temperature of the solution treatment is 1121-1220° C.; and / or, the time of the solution treatment is 2-4 hours; and / or, water cooling is adopted after the solution treatment.

13. The method for preparing a high temperature resistant nickel-based alloy for a high temperature reactor according to claim 7, characterized in that: The temperature of the aging treatment is 600-630° C.; and / or, the time of the aging treatment is 2-4 hours; and / or, water cooling is used after the aging treatment is completed.

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