Material suitable for fast reactor fuel assembly and method for preparing tube blank suitable for fast reactor fuel assembly

Abstract

The embodiment of the invention relates to the field of iron-based alloys, in particular to a material suitable for a fast reactor fuel assembly and a method for preparing a tube blank suitable for the fast reactor fuel assembly. According to the material, the content of Ni and Mn is reduced, the diffusion speed of M23C6 is reduced, and the coarsening speed of M23C6 is inhibited; n is added, a VN phase is separated out, pinning dislocation is achieved, subcrystal formation is delayed, and the tissue coarsening speed is slowed down; the content of B is reduced, and the content of Cr is increased, so that the generated M23C6 and VN can jointly pinning the lath boundary; nb, Zr, Ta and V are introduced, MX phase precipitation is promoted, MX and M23C6 strengthening and M23C6 and VN pinning lath boundaries are achieved, the structure stability is improved, the high-temperature durability of the material is improved, and the safety requirements for high temperature, high irradiation dose and long-time use of a fast reactor core assembly are met. According to the method, the pipe blank with high-temperature durability and irradiation resistance can be prepared, it is ensured that the cladding obtained through subsequent machining can keep a high-temperature-resistant and irradiation-resistant structure, and the economical efficiency and safety of a reactor are improved.

Description

Technical Field

[0001] The embodiments of this application relate to the field of iron-based alloys, and in particular to a material suitable for fast reactor fuel assemblies and a method for preparing tube blanks suitable for fast reactor fuel assemblies. Background Technology

[0002] The statements herein are provided merely as background information in connection with this application and do not necessarily constitute prior art.

[0003] Fast reactors typically use metallic fuels to improve fast neutron economy and heat transfer efficiency. These metallic fuels require cladding to isolate the coolant, prevent corrosion and radioactive leakage, and maintain structural integrity. During reactor operation, under sustained high temperatures and intense radiation, the cladding material undergoes degradation such as swelling, creep, and embrittlement, leading to decreased mechanical properties and dimensional instability. Therefore, it is necessary to improve the material properties of the metallic fuel cladding to ensure the long-term safe operation of the reactor. Summary of the Invention

[0004] A brief overview of this application is provided below to offer a basic understanding of certain aspects thereof. It should be understood that this overview is not an exhaustive summary of the application. It is not intended to identify key or essential parts of the application, nor is it intended to limit its scope. Its purpose is merely to present certain concepts in a simplified form as a prelude to the more detailed description that follows.

[0005] This application provides a material suitable for fast reactor fuel assemblies, the composition of which is: C: 0.18%-0.23%; Si: ≤0.30%; Mn: 0.20%-0.70%; N: 0.06%-0.15%; Ni: 0.20%-0.60%; Cr: 10.50%-12.50%; Mo: 0.40%-0.60%; V: 0.20%-0.40%; W: 1.0%-2.5%; Nb: 0.10%-0.45%; B: 0.002%-0.008%; Ta: 0.05%-0.15%; Al: ≤0.020%; Ti: ≤0.10%; Zr: ≤0.015%; Cu: ≤0.02%; Co: ≤0.015%; O: ≤0.010%; P: ≤0.015%; S: ≤0.010%, with the remainder being Fe.

[0006] The materials for fast reactor fuel assemblies provided in the embodiments of this application reduce diffusion and suppress coarsening by lowering the Ni and Mn content; by adding N, the VN phase can be precipitated, pinning dislocations, delaying subgrain formation, and slowing down the coarsening rate; by lowering the B content and increasing the Cr content, the generated M... 23C6 and VN can jointly anchor the lath boundary; by introducing Nb, Zr, Ta, and V elements into composite reinforcing materials, the precipitation of MX phase is promoted, achieving the desired MX and M phase precipitation. 23 C6 Enhancement, M 23 C6 and VN pinning strips at the interface enhance microstructure stability, thereby improving the material's high-temperature durability and enabling it to meet the safety and reliability requirements of fast reactor core assemblies under high temperature, high radiation dose, and long-term use. The materials provided in the embodiments of this application, due to their low B content, high microstructure stability Cr content, and even lower Ni and Mn content, can reduce M... 23 C6 diffusion rate, slowing down M 23 C6 coarsening speed.

[0007] This application also provides a method for preparing a tube blank suitable for fast reactor fuel assemblies, wherein the material composition of the prepared tube blank is the same as that described above.

[0008] The method for preparing tube blanks suitable for fast reactor fuel assemblies provided by the embodiments of this application can prepare tube blanks with high-temperature durability and radiation resistance, ensuring that the cladding obtained by subsequent processing can maintain a high-temperature and radiation-resistant structure, thereby improving the economics and safety of the reactor. Attached Figure Description

[0009] To further illustrate the above and other advantages and features of this application, the specific embodiments of this application will be described in more detail below with reference to the accompanying drawings. The drawings, together with the following detailed description, are included in and form a part of this specification. Elements having the same function and structure are indicated by the same reference numerals. It should be understood that these drawings only depict typical examples of this application and should not be considered as limiting the scope of this application.

[0010] Figure 1 This is a schematic diagram of the grain boundary carbide distribution of a material. Figure 2 This is a schematic diagram of the distribution of grain boundary carbides and nitrides in the materials provided in the embodiments of this application; Figure 3 This is a schematic diagram illustrating the phase volume fraction change of the tube blank material obtained by the preparation method provided in the embodiments of this application; Figure 4 Yes Figure 3 A magnified schematic diagram of the precipitated phase in the sample; Figure 5 This is a schematic diagram illustrating the durability level of the tube blank material obtained by the preparation method provided in the embodiments of this application. Detailed Implementation

[0011] Exemplary embodiments of this application will be described below with reference to the accompanying drawings. For clarity and brevity, not all features of actual implementations are described in the specification. However, it should be understood that many implementation-specific decisions must be made in the development of any such actual embodiment to achieve the developer's specific goals, such as complying with constraints related to the system and business, and these constraints may vary depending on the implementation. Furthermore, it should be understood that while development work can be very complex and time-consuming, such development work is merely a routine task for those skilled in the art who benefit from the content of this application.

[0012] It should also be noted that, in order to avoid obscuring this application with unnecessary details, only the equipment structure and / or processing steps closely related to the solution according to this application are shown in the accompanying drawings, while other details that are not closely related to this application are omitted.

[0013] The following disclosure provides several different implementations or examples for carrying out this application. To simplify the disclosure of this application, specific examples of components and methods are described below. Of course, these are merely examples and are not intended to limit this application. In the description of the embodiments of this application, "multiple" means at least two, such as two, three, etc., unless otherwise explicitly specified.

[0014] Currently, research on fast reactor cladding materials mainly focuses on 15-15Ti stainless steel, ferritic / martensitic heat-resistant steel, and ODS steel. 15-15Ti stainless steel meets the requirements for high-temperature applications; however, it exhibits a high tendency for radiation swelling under high radiation doses. While ODS steel combines high-temperature strength and resistance to radiation swelling, it is not yet practically applicable. 9-12%Cr ferritic / martensitic heat-resistant steel possesses excellent resistance to neutron radiation swelling and is readily applicable, but some grades suitable for nuclear energy applications have relatively low high-temperature strength, requiring further improvement in heat resistance. Therefore, it is necessary to strengthen ferritic / martensitic steel to achieve both high-temperature resistance and resistance to radiation hardening or embrittlement.

[0015] Conventional 12% Cr ferritic / martensitic heat-resistant steel exhibits poor creep performance under the same working conditions, as shown by: M 23 C6 is prone to coarsening and has difficulty sustaining lath boundaries. After prolonged exposure to high temperatures, it forms subgrains, resulting in lath coarsening, poor microstructure stability, and reduced durability.

[0016] Conventional techniques for suppressing martensitic lath coarsening mainly involve adding Co and higher levels of B, but this can negatively impact the material's radiation resistance.

[0017] To address the aforementioned problems, embodiments of this application provide a material suitable for fast reactor fuel assemblies, with the following composition: C: 0.18%-0.23%; Si: ≤0.30%; Mn: 0.20%-0.70%; N: 0.06%-0.15%; Ni: 0.20%-0.60%; Cr: 10.50%-12.50%; Mo: 0.40%-0.60%; V: 0.20%-0.40%. %; W: 1.0%-2.5%; Nb: 0.10%-0.45%; B: 0.002%-0.008%; Ta: 0.05%-0.15%; Al: ≤0.020%; Ti: ≤0.10%; Zr: ≤0.015%; Cu: ≤0.02%; Co: ≤0.015%; O: ≤0.010%; P: ≤0.015%; S: ≤0.010%, with the remainder being Fe.

[0018] The materials for fast reactor fuel assemblies provided in the embodiments of this application reduce diffusion and suppress coarsening by lowering the Ni and Mn content; by adding N, the VN phase can be precipitated, pinning dislocations, delaying subgrain formation, and slowing down the coarsening rate; by lowering the B content and increasing the Cr content, the generated M... 23 C6 and VN can jointly anchor the lath boundary; by introducing Nb, Zr, Ta, and V elements into composite reinforcing materials, the precipitation of MX phase is promoted, achieving the desired MX and M phase precipitation. 23 C6 Enhancement, M 23 C6 and VN pinning strips at the interface enhance microstructure stability, thereby improving the material's high-temperature durability and enabling it to meet the safety and reliability requirements of fast reactor core assemblies under high temperature, high radiation dose, and long-term use. The materials provided in the embodiments of this application, due to their low B content, high microstructure stability Cr content, and even lower Ni and Mn content, can reduce M... 23 C6 diffusion rate, slowing down M 23 C6 coarsening speed.

[0019] Figure 1 This is a schematic diagram of the grain boundary carbide distribution of a material, such as... Figure 1 As shown, the black lines represent austenite grain boundaries, and the orange lines in the magnified area represent lath boundaries. Black dots are distributed within the lath boundaries, and these black dots represent M... 23 C6, Figure 1 The precipitated phases are mainly M-type phases distributed along the original austenite grain boundaries and lath boundaries. 23 C6.

[0020] Figure 2 This is a schematic diagram of the grain boundary carbide and nitride distribution of the material provided in the embodiments of this application, as shown below. Figure 2As shown, the black lines represent austenite grain boundaries, and the orange lines in the magnified area represent lath boundaries. Black and red dots are distributed within the lath boundaries; the black dots represent M... 23 C6, the red dot represents the MX phase. Figure 2 The precipitated phase in retains M 23 Along with C6, the MX phase was added. The MX phase is mainly distributed inside the slats, while the M... 23 C6 is still mainly distributed at all levels of grain boundaries, and the two complement each other in spatial distribution, thus enabling composite reinforcement of materials.

[0021] This application also provides a method for preparing a tube blank suitable for fast reactor fuel assemblies, wherein the material composition of the prepared tube blank is the same as that described above.

[0022] The method for preparing tube blanks suitable for fast reactor fuel assemblies provided by the embodiments of this application can prepare tube blanks with high-temperature durability and radiation resistance, ensuring that the cladding obtained by subsequent processing can maintain a high-temperature and radiation-resistant structure, thereby improving the economics and safety of the reactor.

[0023] In some embodiments, a method for preparing a tube blank suitable for fast reactor fuel assemblies includes the following steps: S10: vacuum induction smelting of raw materials; S20: electroslag remelting of the product obtained in step S10; S30: forging the product obtained in step S20 to obtain the tube blank. By combining the purification smelting processes of vacuum induction and electroslag remelting, precise control of the chemical composition of the tube blank can be achieved, and the requirements for non-metallic inclusions and the nitrogen / carbide levels of titanium in the material can be met.

[0024] Figure 3 This is a schematic diagram illustrating the phase volume fraction change of the tube blank material obtained by the preparation method provided in the embodiments of this application, as shown below. Figure 3 As shown, the yellow-green curve represents austenite (γ-Fe), and the purple curve represents the carbide precipitate (M). 23 C6), the red curve corresponding to the temperature below 1000℃ on the horizontal axis represents α-ferrite (α-Fe), and the red curve corresponding to the temperature above 1000℃ on the horizontal axis represents δ-ferrite (δ-Fe).

[0025] δ-ferrite typically impairs the low-temperature toughness and weld heat-affected zone toughness of materials. While δ-ferrite formation is usually suppressed by adding austenite-forming elements such as Ni, Mn, C, and N, Ni and Mn promote carbide coarsening, negatively impacting high-temperature creep performance. Excessive C accelerates carbide coarsening, and excessive N increases the risk of material cracking. Figure 3As shown, the liquidus temperature of the tube blank material obtained by the preparation method provided in the embodiments of this application is about 1487°C. δ-ferrite begins to transform into austenite from 1393°C and is completely transformed into austenite at 1138°C. The region between 914°C and 1138°C is the fully austenitic region. The austenitic phase transformation temperature is 826°C. This indicates that the composition of the material provided in the embodiments of this application achieves a balance between the ferrite-forming elements and the contents of C and N.

[0026] High W content can also promote the formation of δ-ferrite and the precipitation of Laves phase during long-term use, which has an adverse effect on the high-temperature strength of the material. Therefore, appropriate W and Mo contents are required to obtain a fully martensitic structure and reduce the tendency of Laves phase precipitation. Figure 4 Yes Figure 3 A magnified schematic diagram of the precipitated phase in the image, as shown below. Figure 4 As shown, the black curve represents the change in the volume fraction of the Laves phase. The trend of the black curve indicates that the growth and coarsening kinetics of the Laves phase during cooling are very slow. The total amount of W and Mo is moderate and effective, which can trigger the precipitation of the Laves phase in the high-temperature tempering range without excess, so as to avoid damaging the toughness.

[0027] Figure 5 This is a schematic diagram illustrating the creep life level of the tube blank material obtained by the preparation method provided in the embodiments of this application. By fitting the relationship between the Rasmiller parameter and the stress of each material, the creep life level of each material is predicted, such as... Figure 5 As shown, the green scatter plots and curves represent the stress variation with the Rasmusville parameter (LMP) of the billet material obtained by the preparation method provided in the embodiments of this application; the blue scatter plots and curves represent the stress variation with LMP of T92 steel; the black scatter plots and curves represent the stress variation with LMP of HT9 steel; and the red pentagrams indicate the positions of the material performance indicators required for fast reactor operation. Figure 5 It can be seen that neither HT9 steel nor T92 steel can meet the material performance requirements, and HT9 steel has a large gap with the requirements. However, the fitting curve of the tube blank material obtained by the preparation method provided in the embodiments of this application is above the material performance index required by the operating conditions, and can meet the requirements of the fast reactor operating conditions.

[0028] In some embodiments, in step S10, the furnace vacuum degree before smelting is set to ≤1.0 Pa; in the refining stage, the furnace vacuum degree is set to ≤5.0 Pa; in the alloying stage, the furnace vacuum degree is set to ≤5.0 Pa; and the composition of the raw materials is adjusted according to the elemental analysis results of N, Si, and Al. By controlling the furnace vacuum degree before smelting to ≤1.0 Pa, impurities carried by the raw materials can be removed, ensuring a pure melt environment in the subsequent smelting stages, which is beneficial to improving smelting quality. Furthermore, controlling the furnace vacuum degree in the refining and alloying stages to ≤5.0 Pa ensures a moderate furnace vacuum environment, which can ensure accurate and stable alloy composition. Elemental analysis of N, Si, and Al can reflect the degree of deoxidation and purity of the melt. If too much N, Si, and Al remains, it will affect the toughness and high-temperature strength of the material by forming specific oxides or nitrides. Fine-tuning the material composition according to the analysis results can calibrate the alloy composition, resulting in a more accurate composition and higher performance of the smelted alloy.

[0029] In some embodiments, after adjusting the composition of the raw materials to meet predetermined requirements based on the elemental analysis results of N, Si and Al, a protective gas (e.g., argon) can be introduced into the vacuum induction furnace to make the pressure of the protective atmosphere meet predetermined requirements (e.g., 5000-8000 Pa), and metallic manganese can be added under the protective atmosphere.

[0030] In some embodiments, if the vacuum level in the furnace during the refining stage does not meet the requirement of ≤1.0Pa, the power can be reduced for heat preservation to avoid the alloy quality from being reduced due to continued melting under insufficient vacuum conditions. Heat preservation can buy time to restore the required vacuum level in the furnace and stabilize the melt state.

[0031] In some embodiments, in step S10, the temperature of the refining stage is set to about 1560°C; the temperature of the early stage of the alloying stage is set to ≤1570°C; and the temperature of the later stage of the alloying stage is set to 1530°C-1600°C. With the above settings, it is possible to ensure that an alloy ingot with high purity, accurate composition and uniform structure is obtained.

[0032] In some embodiments, when the melt in the vacuum induction furnace reaches a state of calm surface and no bubbles, the temperature is adjusted to enter the refining stage, and the refining stage time can be set to more than 45 minutes.

[0033] In some embodiments, sampling and analysis can be performed in the early stages of alloying to allow for adjustments to the composition based on elemental analysis results of N, Si, and Al.

[0034] In some embodiments, casting can be carried out in the later stage of alloying, the casting temperature can be set to 1530℃-1600℃, and the casting method can be top casting. During the casting process, the crucible of the vacuum induction furnace should be moved smoothly to ensure that the melt flows out stably and to avoid interruption of flow and slag entrapment. After casting, vacuum solidification is performed, and the alloy ingot is obtained by demolding after complete solidification.

[0035] In some embodiments, in step S20, the alloy ingot obtained in step S20 can be electroslag remelted using a CaF2-CaO-Al2O3-MgO quaternary slag system.

[0036] Because boron (B) is chemically extremely reactive and easily reacts with oxygen or nitrogen, it is prone to burn-off and uneven distribution. The above-mentioned quaternary slag system can fix boron oxides and reduce boron activity, thereby avoiding burn-off caused by uneven distribution of boron. It can also optimize the electroslag remelting rate and remove large particle inclusions.

[0037] In some embodiments, the ratio of the CaF2-CaO-Al2O3-MgO quaternary slag system can be set to 67:15:16:2.

[0038] In some embodiments, in step S20, during the slag-forming stage, the slag-adding current is set to 3000A, the slag-adding time is set to 25-40min, the refining current is set to 4500A, and the refining time is set to 20-30min. Through the above settings, the slag melting process during the slag-forming stage can be made uniform and stable, and the input power during the refining stage can be increased, making the reaction more thorough.

[0039] In some embodiments, in step S20, the steady-state remelting time is controlled at 8.5-10.5h, the voltage is controlled at 55-68V, and the current is controlled at 9000-14000A. Through the above settings, the power input of steady-state remelting can be stable, which is beneficial to obtaining a pure and homogeneous alloy ingot.

[0040] In some embodiments, during step S20, the voltage and current of the smelting process are detected at predetermined intervals (e.g., 30 minutes) during the steady-state remelting stage. This helps to detect hidden process defects in a timely manner during the monitoring of steady-state remelting and avoid uneven quality of the obtained alloy ingots.

[0041] In some embodiments, in step S20, during the ingot casting stage after electroslag remelting, the cooling time is not less than 60 minutes; after demolding, the electroslag remelted ingot is slowly cooled in a protective cover for not less than 48 hours. By controlling the cooling rate and cooling time, the internal structure of the alloy ingot can be made uniform, and the thermal stress caused by the excessive internal and external temperature difference can be reduced.

[0042] In some embodiments, annealing can be performed within 3 days of demolding.

[0043] In some embodiments, in step S20, during the annealing stage, the annealing heating rate is ≤100°C. o C / h, and at a temperature of 680±20 o The holding time at temperature C is ≥20h, and after holding, the annealing cooling rate is ≤40. o C / h, and at temperatures ≤300 o At step C, the product is removed from the furnace and air-cooled to reduce the internal stress caused by the temperature difference in the material cross section and improve the uniformity and stability of the structure.

[0044] In some embodiments, in step S30, before forging, the alloy ingot obtained by electroslag remelting in step S20 can be heated to a predetermined temperature (e.g., 1170°C) and homogenized for a predetermined time (e.g., 5 hours). By subjecting the alloy ingot to long-term high-temperature homogenization treatment, elemental segregation can be reduced.

[0045] In some embodiments, during the rapid forging stage, the alloy ingot can be forged in multiple passes (e.g., 5 passes) to obtain an intermediate billet, with the final forging temperature of each pass being greater than a predetermined value (e.g., 850°C).

[0046] In some embodiments, the heating temperature of the intermediate billet can be 1120°C, the heat preservation time after forging can be 4 hours, and the head and tail of the billet forged in the final heat are removed (for example, the head removal amount is greater than 5%, and the tail removal amount is greater than 4%).

[0047] In some embodiments, during the finishing forging stage, the intermediate billet is returned to the heating furnace and held at a temperature of 1100°C. o C, heat treatment time 4 hours, and then forged to the finished bar size, with a forging ratio of 16.

[0048] By precisely controlling the forging temperature of the alloy using the above methods and employing a gradually decreasing process, the rapid forging temperature and the precision forging temperature can be reduced step by step, thereby reducing the precipitation of ferrite phase.

[0049] In some embodiments, a combination of upsetting and drawing processes can be used during forging to eliminate element segregation and obtain forging billets with uniform structure and composition.

[0050] Regarding the embodiments of this application, it should also be noted that, without conflict, the embodiments of this application and the features in the embodiments can be combined with each other to obtain new embodiments.

[0051] The above are merely specific embodiments of this application, but the scope of protection of this application is not limited thereto. The scope of protection of this application shall be determined by the scope of the claims.

Claims

1. A material suitable for fast reactor fuel assemblies, characterized in that, Its components are: C: 0.18%-0.23%; Si: ≤0.30%; Mn: 0.20%-0.70%; N: 0.06%-0.15%; Ni: 0.20%-0 .60%; Cr: 10.50%-12.50%; Mo: 0.40%-0.60%; V: 0.20%-0.40%; W: 1.0%-2.5%; Nb ：0.10%-0.45%； B：0.002%-0.008%； Ta: 0.05%-0.15%; Al: ≤0.020%; Ti: ≤0.10%; Zr: ≤0.015%; Cu: ≤0.02%; Co: ≤0.015%; O: ≤0.010%; P: ≤0.015%; S: ≤0.010%, the remainder is Fe.

2. A method for preparing a tube blank suitable for fast reactor fuel assemblies, characterized in that, The material composition of the prepared tube blank is the same as that described in claim 1.

3. The method according to claim 2, characterized in that, It includes the following steps: S10: Vacuum induction smelting of raw materials; S20: Electroslag remelting the product obtained in step S10; S30: The product obtained in step S20 is forged to obtain the tube blank.

4. The method according to claim 3, characterized in that, In step S10, the vacuum degree inside the furnace before smelting is set to ≤1.0 Pa; During the refining stage, the vacuum level inside the furnace is set to ≤5.0 Pa; During the alloying stage, the vacuum level inside the furnace is set to ≤5.0 Pa; Furthermore, the composition of the raw materials is adjusted based on the elemental analysis results of N, Si, and Al.

5. The method according to claim 3, characterized in that, In step S10, the temperature for the refining stage is set to approximately 1560°C; The temperature in the early stage of alloying is set to ≤1570℃; The temperature for the later stage of alloying was set to 1530℃-1600℃.

6. The method according to claim 3, characterized in that, In step S20, a quaternary slag system of CaF2-CaO-Al2O3-MgO is used.

7. The method according to claim 3, characterized in that, In step S20, during the slag-forming stage, the slag-adding current is set to 3000A and the slag-adding time is set to 25-40min. The refining current is set to 4500A and the refining time is set to 20-30min.

8. The method according to claim 3, characterized in that, In step S20, the steady-state remelting time is controlled at 8.5-10.5 hours, the voltage at 55-68V, and the current at 9000-14000A.

9. The method according to claim 8, characterized in that, In step S20, during the steady-state remelting stage, the smelting voltage and current are monitored at predetermined intervals.

10. The method according to claim 3, characterized in that, In step S20, during the ingot casting stage after electroslag remelting, the cooling time shall not be less than 60 minutes; After being removed from the furnace and demolded, the electroslag remelted ingot is slowly cooled inside a protective cover for no less than 48 hours.

11. The method according to claim 3, characterized in that, In step S20, during the annealing stage, the annealing temperature rise rate is ≤100°C. o C / h, and at a temperature of 680±20 o The heat preservation time at temperature C is ≥20 hours. After heat preservation, the annealing cooling rate is ≤40°C. o C / h, and at temperatures ≤300 o At step C, the product is removed from the furnace and air-cooled.

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