Nuclear shielding material and preparation method thereof

Abstract

The invention discloses a nuclear shielding material and a preparation method thereof, and relates to the technical field of nuclear shielding. The preparation raw materials of the nuclear shielding material comprise tungsten carbide and a high-entropy alloy; the high-entropy alloy comprises the following components in percentage by mass: 20 to 40 percent of Fe, 20 to 40 percent of Ni, 12 to 30 percent of Cr, 14 to 25 percent of Cu and 0 to 15 percent of Co. The nuclear shielding material is mainly formed by compounding the high-entropy alloy and the WC, the hardness of the material is improved, meanwhile, the wear resistance and the high-temperature oxidation resistance are improved, and the yield strength and the nuclear shielding performance of a ceramic composite nuclear shielding material are considered. Therefore, the nuclear shielding material obtained by the invention can also realize mechanical strength, wear resistance, oxidation resistance and nuclear shielding performance at the same time.

Description

Technical Field

[0001] The present invention relates to the technical field of nuclear shielding, and in particular to a nuclear shielding material and a preparation method thereof. Background Art

[0002] With the continuous development of marine equipment, the demand for new reactor types is growing stronger. To simplify shielding designs for new reactor types and reduce the volume of shielding materials, there is an urgent need to conduct comprehensive shielding material design, preparation, and performance research, providing a shielding material database for the design and construction of new reactor types. WC-based metal-ceramic composites, with their advantages of high hardness, wear resistance, acid and alkali corrosion resistance, and high temperature resistance, are currently one of the most widely used cemented carbide materials and have the potential to be used as nuclear shielding materials.

[0003] High entropy alloy is a commonly used cemented carbide bonding phase in WC-based metal-ceramic composites. Its unique four major effects can make WC-based metal-ceramic composites form a simple phase structure and have excellent comprehensive properties, such as good mechanical properties, wear resistance, corrosion resistance, thermal stability and oxidation resistance, and resistance to radiation damage. However, the commonly used high entropy alloy components are generally based on transition metal elements, among which high entropy alloys with Co, Cr, Cu, Fe, Ni, Al, and Mn as components are the most widely studied. For example, the CoCrNiCuFeAl / (W,Ti)(C,N) metal ceramic material disclosed in the prior art is based on CoCrNiCuFeAl, TiC, WC, TiN x (x=0.3~0.9) and TiC as the main components, as well as WC-(Tix,Cr 1-x )3C2-(Fe,Cr)-(Ni,Cu)-Co multi-component cemented carbide. However, ceramic composites formed by these high-entropy alloys are not suitable for the requirements of nuclear radiation field for mechanical strength, wear resistance, oxidation resistance and other properties of materials. They cannot well balance mechanical properties and nuclear shielding performance. As a result, WC-based metal-ceramic composites containing these traditional high-entropy alloys are difficult to be applied as nuclear shielding materials. Summary of the Invention

[0004] The present invention aims to solve at least one of the above-mentioned technical problems existing in the prior art. To this end, the present invention provides a nuclear shielding material, which is a ceramic composite material. By adjusting the composition and content of the high-entropy alloy, the mechanical strength, wear resistance, and oxidation resistance are improved, thereby achieving good mechanical strength and wear resistance.

[0005] The second aspect of the present invention is to provide a method for preparing a nuclear shielding material.

[0006] In order to achieve the above object, the technical solution adopted by the present invention is:

[0007] A first aspect of the present invention provides a nuclear shielding material, wherein the raw materials for preparing the material include tungsten carbide and a high entropy alloy; the high entropy alloy includes the following components by mass percentage:

[0008] Fe: 20 to 40%, Ni: 20 to 40%, Cr: 12 to 30%, Cu: 14 to 25%, Co: 0 to 15%.

[0009] In some embodiments of the present invention, in the nuclear shielding material, the mass content of unavoidable impurities in the high entropy alloy is ≤0.35%.

[0010] In some embodiments of the present invention, in the nuclear shielding material, the high entropy alloy is composed of the following components by mass percentage:

[0011] Fe: 25-35%, Ni: 25-35%, Cr: 12-30%, Cu: 14-25%, Co: 0-15%, inevitable impurities ≤ 0.35%.

[0012] In some embodiments of the present invention, the mass percentage of Co in the high entropy alloy in the nuclear shielding material is 0.

[0013] In some specific embodiments of the present invention, in the nuclear shielding material, the high entropy alloy is composed of the following components by mass percentage:

[0014] Fe: 25-35%, Ni: 25-35%, Cr: 15-30%, Cu: 15-25%, inevitable impurities ≤ 0.35%.

[0015] In some embodiments of the present invention, in the nuclear shielding material, the crystal phase of the high entropy alloy is FCC phase.

[0016] In some embodiments of the present invention, in the nuclear shielding material, the high entropy alloy is in powder form, and the particle size D50 is 7 to 9 μm.

[0017] In some specific embodiments of the present invention, the high entropy alloy is in powder form, and the particle size D90 is 13-15 μm.

[0018] In some embodiments of the present invention, in the nuclear shielding material, the method for preparing the high entropy alloy comprises the following steps:

[0019] The raw material components of the high entropy alloy are mixed and smelted, and then powdered to obtain the high entropy alloy.

[0020] In some specific embodiments of the present invention, the raw materials include pure metals and / or alloys of various elements; the pure metals or alloys are in bulk; and before the mixed smelting, impurities and / or oxide films on the metal surface are removed.

[0021] In some specific embodiments of the present invention, the smelting pressure is (3-7)×10 -2 MPa.

[0022] In some specific embodiments of the present invention, the smelting pressure is (4-6)×10 -2 MPa.

[0023] In some specific embodiments of the present invention, the smelting process is first vacuumed to 3×10 -3 MPa or less, for example, (2 to 3) × 10 -3 MPa, and then fill with inert gas, such as argon, to reach the required melting pressure.

[0024] In some specific embodiments of the present invention, the smelting time is 20 to 40 minutes.

[0025] In some specific embodiments of the present invention, the smelting time is 25 to 35 minutes.

[0026] In some specific embodiments of the present invention, the powder making method includes vacuum air atomization; the conditions of the vacuum air atomization method meet at least one of the following a) to d):

[0027] a) Air pressure is 2-6 MPa;

[0028] b) The air flow velocity is 100-150 mm / s;

[0029] c) The liquid flow diameter is 5 to 10 mm;

[0030] d) The air flow injection angle is 30° to 60°.

[0031] In some specific embodiments of the present invention, the air pressure of the vacuum air atomization is 3-5 MPa.

[0032] In some specific embodiments of the present invention, the air flow velocity of the vacuum air atomization is 120-140 mm / s.

[0033] In some specific embodiments of the present invention, the diameter of the liquid flow of the vacuum air atomization is 7 to 9 mm.

[0034] In some specific embodiments of the present invention, the air flow spray angle of the vacuum air atomization is 40° to 50°.

[0035] In some specific embodiments of the present invention, secondary smelting is performed before the vacuum gas atomization; the conditions of the secondary smelting are the same as the conditions of the smelting before powder making.

[0036] In some embodiments of the present invention, in the nuclear shielding material, the mass ratio of the tungsten carbide to the high entropy alloy is (9-14):1.

[0037] In some specific embodiments of the present invention, the mass ratio of the tungsten carbide to the high entropy alloy is (11-12):1.

[0038] In some embodiments of the present invention, the raw materials for preparing the nuclear shielding material further include at least one of a grain inhibitor, an antioxidant, and a filler; the grain inhibitor includes vanadium carbide and / or chromium carbide; the antioxidant includes carbohydrazide; and the filler includes carbon black.

[0039] In some embodiments of the present invention, the mass of the grain inhibitor, antioxidant, and filler accounts for 0.5-0.7%, 0.2-0.4%, and 0.2-0.3% of the total mass of tungsten carbide and high entropy alloy, respectively.

[0040] In some embodiments of the present invention, the porosity of the nuclear shielding material is 0.9-3.5%.

[0041] In some specific embodiments of the present invention, the porosity of the nuclear shielding material is 0.9-2%.

[0042] It should be noted that the porosity is obtained by calculating the ratio of the cross-sectional area to the pore area using an imaging method.

[0043] The second aspect of the present invention provides a method for preparing the nuclear shielding material according to the first aspect of the present invention, comprising the following steps:

[0044] The nuclear shielding material is obtained by mixing tungsten carbide and high entropy alloy and ball milling, followed by cold pressing and sintering.

[0045] In some embodiments of the present invention, the ball-to-material ratio of the ball milling is (7-9):1.

[0046] In some specific embodiments of the present invention, the ball milling method is: every 1.5 to 2.5 hours of rotation and 15 to 25 minutes of rest.

[0047] In some specific embodiments of the present invention, the rotation speed of the ball mill is 250-350 rpm.

[0048] In some specific embodiments of the present invention, the ball milling time is 20 to 30 hours.

[0049] In some embodiments of the present invention, a grain inhibitor, an antioxidant and carbon black are further added during the ball milling process.

[0050] In some embodiments of the present invention, a binder and a process control agent are further added during the ball milling process; the binder comprises polyvinyl alcohol; and the process control agent comprises stearic acid.

[0051] In some specific embodiments of the present invention, the mass of the adhesive and the process control agent accounts for 2-3% and 0.1-0.3% of the total mass of the tungsten carbide and the high entropy alloy respectively.

[0052] In some embodiments of the present invention, a drying step is further included before the cold pressing; the drying conditions are: temperature 70-90° C., time 7-9 hours.

[0053] In some embodiments of the present invention, the cold pressing step comprises: placing the powder obtained after ball milling into a mold with an inner wall coated with a lubricant, and pressing it into blocks using a press.

[0054] In some embodiments of the present invention, the lubricant comprises a boron nitride lubricant.

[0055] In some embodiments of the present invention, the sintering temperature is 1000-1300°C.

[0056] In some specific embodiments of the present invention, the sintering temperature is 1000-1200°C.

[0057] In some specific embodiments of the present invention, the sintering temperature is 1050-1150°C.

[0058] In some specific embodiments of the present invention, the sintering temperature is 1095-1105°C.

[0059] In some embodiments of the present invention, the sintering pressure is 20-30 MPa.

[0060] In some specific embodiments of the present invention, the holding time of the sintering is 50 to 70 minutes.

[0061] In some specific embodiments of the present invention, the holding time of the sintering is 55 to 65 minutes.

[0062] Compared with the prior art, the present invention has at least the following beneficial technical effects:

[0063] (1) The nuclear shielding material provided by the present invention contains a high-entropy alloy of specific composition. The mechanical properties, wear resistance, and high-temperature oxidation resistance of the high-entropy alloy are improved by adjusting the composition and content of each element, and can achieve good mechanical strength and wear resistance. At the same time, the content of Co element in the high-entropy alloy is relatively low, which reduces the Co-60 isotope formed by cobalt under long-term radiation environment. When the Co element content is as low as 0, the formation of Co-60 can be avoided, thereby avoiding the problem of significantly extending the nuclear radiation half-life, which causes an increase in maintenance time and the cost of nuclear fuel shielding, and significantly improving the nuclear shielding performance.

[0064] (2) In the nuclear shielding material of the present invention, by controlling the content of each element in the high entropy alloy, especially the content of Cr and Cu, the high entropy alloy can present a single FCC phase. The slip plane of the FCC phase is highly densely packed, and the lattice resistance to dislocation slip is low, which can effectively improve the plastic toughness of the material and help to improve the yield strength and microhardness of the high entropy alloy, thereby comprehensively improving the mechanical properties of the nuclear shielding material.

[0065] (3) The nuclear shielding material of the present invention contains a high-entropy alloy prepared by smelting and vacuum gas atomization, which has a relatively low particle size, with a D50 particle size of only 7 to 9 μm. It is easier to be refined and adhere to the surface of WC particles during the ball milling process, which is beneficial to the binding of the high-entropy alloy to the WC particles, thereby improving the performance of the nuclear shielding material.

[0066] (4) The nuclear shielding material of the present invention is mainly composed of a high entropy alloy and WC composite, wherein the high entropy alloy as a bonding agent meets the WC phase to form an M3W3C phase (M = Fe, Ni, Cr, Cu or Co), which improves the hardness of the nuclear shielding material, improves the wear resistance and high temperature oxidation resistance, and also retains part of the FCC phase, thereby taking into account the yield strength of the nuclear shielding material. Therefore, the nuclear shielding material obtained by the present invention can achieve the effect of combining mechanical strength, wear resistance, oxidation resistance and nuclear shielding performance.

[0067] (5) The present invention can also reduce the porosity by adjusting the sintering conditions of the nuclear shielding material, avoid Cu overflow loss, and further improve the mechanical strength and wear resistance of the nuclear shielding material. BRIEF DESCRIPTION OF THE DRAWINGS

[0068] Figure 1 The SEM images of the high entropy alloy powders of Preparation Examples 1 to 3 are shown; Figure 1 a to c represent the SEM images of the high entropy alloys of Preparation Examples 1 to 3, respectively.

[0069] Figure 2 1 to 3 are XRD patterns of the high entropy alloy powders prepared in Examples 1 to 3.

[0070] Figure 3 is the particle size distribution diagram of the high entropy alloy powder of Preparation Examples 1 to 3; wherein, Figure 3 Figures a to c in the figure respectively represent the particle size distribution diagrams of the high entropy alloys of Preparation Examples 1 to 3.

[0071] Figure 4 The SEM image and EDS analysis image of the nuclear shielding material of Example 1 are shown; wherein, Figure 4 Figures a to b represent SEM images and EDS spectra respectively.

[0072] Figure 5 The SEM image and EDS analysis image of the nuclear shielding material of Example 2; wherein, Figure 5 Figures a to b represent SEM images and EDS spectra respectively.

[0073] Figure 6 The SEM image and EDS analysis image of the nuclear shielding material of Example 3; wherein, Figure 6 Figures a to b represent SEM images and EDS spectra respectively.

[0074] Figure 7 The SEM image and EDS analysis image of the nuclear shielding material of Example 4 are shown; wherein, Figure 7 Figures a to b represent SEM images and EDS spectra respectively.

[0075] Figure 8 The diagrams of porosity testing of nuclear shielding materials of Examples 1 to 4 by image method and the porosity testing results are shown; wherein, Figure 8 Figures a to b in the figure respectively represent the schematic diagram of the porosity calculation of the nuclear shielding material of Example 1 using the image method and the porosity test results of the nuclear shielding materials of Examples 1 to 4.

[0076] Figure 9 1 is the XRD pattern of the nuclear shielding material in Examples 1 to 4.

[0077] Figure 10 is a friction coefficient curve diagram of the nuclear shielding material of Examples 1 to 4; wherein, Figure 10 Figures a to d in the figure represent the friction coefficient curves of Examples 1 to 4 respectively.

[0078] Figure 11 is a graph showing the average friction coefficient and wear rate of the nuclear shielding materials of Examples 1 to 4; wherein, Figure 11 Figures a to b in the figure represent the average friction coefficient diagram and the wear rate diagram respectively.

[0079] Figure 12 The figure shows the results of the oxidation resistance test of the nuclear shielding materials of Examples 1 to 4; wherein, Figure 12 Figures a to b in the figure represent the optical photographs and oxidation weight gain curves after the antioxidant test.

[0080] Figure 13 This is the EDS surface scanning result of the oxide layer peeling area of ​​the nuclear shielding material in Example 1.

[0081] Figure 14 The SEM image and EDS analysis image of the nuclear shielding material of Example 5; wherein, Figure 14 Figures a to b represent SEM images and EDS spectra respectively.

[0082] Figure 15 The SEM image and EDS analysis image of the nuclear shielding material of Example 6; wherein, Figure 15 Figures a to b represent SEM images and EDS spectra respectively.

[0083] Figure 16 The SEM image and EDS analysis image of the nuclear shielding material of Example 7; wherein, Figure 16 Figures a to b represent SEM images and EDS spectra respectively.

[0084] Figure 17 1 is the XRD pattern of the nuclear shielding materials of Examples 5 to 7.

[0085] Figure 18 Graph showing the hardness test results of the nuclear shielding materials of Examples 5 to 7.

[0086] Figure 19 Graph showing the bending strength test results of the nuclear shielding materials of Examples 5 to 7.

[0087] Figure 20 is a friction curve diagram of the nuclear shielding material of Examples 5 to 6; wherein, Figure 21 Figures a to c in the figure respectively represent the friction curves of Examples 5 to 6 and 316 stainless steel.

[0088] Figure 21 is a graph showing the average friction coefficient and wear rate of the nuclear shielding materials of Examples 5 to 7; wherein, Figure 21 Figures a to b in the figure represent the average friction coefficient diagram and wear rate diagram respectively.

[0089] Figure 22 The following are oxidation weight gain curves of the nuclear shielding materials of Examples 5 to 7 at different oxidation times. DETAILED DESCRIPTION

[0090] The present invention is further described in detail below using specific examples. Unless otherwise specified, the raw materials, reagents, and apparatus used in the examples and comparative examples are available from conventional commercial sources or by conventional methods. Unless otherwise specified, all experimental or testing methods are conventional in the art.

[0091] The following is a detailed description with reference to different preparation examples, embodiments and comparative examples.

[0092] Preparation Example 1

[0093] This preparation example provides a high entropy alloy. The proportions of the various elements in the high entropy alloy are shown in Table 1. The preparation method includes the following steps:

[0094] S1. Batching: Pure metal blocks of Fe, Ni, Cr, and Cu (purity ≥ 99.99%) were batched in a mass ratio of Fe:Ni:Cr:Cu = 1:1:1:0.6, with a total weight of 30 kg; before batching, a grinding wheel was used to remove impurities and oxide films from the metal surface;

[0095] S2. Melting: Place the metal raw materials in the order of melting point from low to high into the vacuum arc melting furnace (HVAM-10) and evacuate to 2.5×10 -3 MPa, and then filled with argon gas to a pressure of 5×10 -2 MPa, under argon protection, the metal block is melted by induction coil heating and electromagnetic stirring. The heating temperature is sufficient to melt the metal block. The melting time is 30 minutes to obtain the alloy billet;

[0096] S3. Powdering: The alloy billet is placed in the smelting device of the vacuum air atomization powder making equipment (PS-ZC-SC-ZF-02) for secondary smelting. The smelting conditions are the same as those in step S2. Subsequently, the molten metal droplets are ejected using a tightly coupled annular gap nozzle to form spherical high-entropy alloy powder; the parameters of the injection process are: air pressure of 4 MPa, air flow velocity of 130 mm / s, liquid flow diameter of 8 mm, and air flow injection angle controlled at 45°.

[0097] The high entropy alloy prepared in this example is denoted as HEA-1.

[0098] Preparation Example 2

[0099] This preparation example provides a high entropy alloy. The differences from Preparation Example 1 are shown in Table 1. The preparation method differs from Example 1 only in that the mass ratio of Fe:Ni:Cr:Cu is 1:1:0.5:0.6. The high entropy alloy of this preparation example is designated HEA-2.

[0100] Preparation Example 3

[0101] This preparation example provides a high-entropy alloy. The differences from Preparation Example 1 are shown in Table 1. The preparation method differs from Example 1 only in that the mass ratio of Fe:Ni:Cr:Cu:Co is 1:1:0.5:0.6:0.5. The high-entropy alloy of this preparation example is designated HEA-3.

[0102] Table 1 Proportions of elements in the high entropy alloys of Preparation Examples 1 to 3 (mass percentage)

[0103]

[0104] It should be noted that the proportions of the elements in the high entropy alloy shown in Table 1 are obtained by analyzing the high entropy alloy using EDS, and positions A and B represent different test positions of the high entropy alloy powder, such as Figure 1 It can also be seen from Table 1 that the mass percentage of each element in the high entropy alloy obtained by the present invention is relatively close to the mass percentage of the components added during the preparation process.

[0105] Example 1

[0106] This embodiment provides a nuclear shielding material, including tungsten carbide (WC) and the high entropy alloy of Preparation Example 1; the preparation method includes the following steps:

[0107] S1. Ball milling of ingredients: Micron-sized WC (particle size range: 1-3 μm), high-entropy alloy, grain inhibitor (vanadium carbide: chromium carbide mass ratio of 1:1), antioxidant carbohydrazide, and carbon black were weighed in proportion and poured into a ball mill jar. Polyvinyl alcohol and stearic acid were added. After sealing, the jar was ball milled in a planetary high-energy ball mill (PM2L). The relevant ball milling parameters included wet milling, a ball milling speed of 300 r / min -1 The ball milling time is 24 hours, with a 20-minute rest period after every 2 hours, and a ball-to-material ratio of 8:1;

[0108] Among them, the mass ratio of micron-scale WC: high entropy alloy: grain inhibitor: antioxidant: carbon black: polyvinyl alcohol: stearic acid is 92:8:0.6:0.3:0.28:2.5:0.2;

[0109] S2 cold pressing: The ball-milled powder was dried at 80°C for 8h and loaded into a graphite mold coated with boron nitride lubricant, and the powder was pressed into a block using a press;

[0110] S3. Sintering: Place the cold pressed powder and mold in a vacuum hot pressing sintering furnace (ZT-45-20) for sintering at 1050°C, with a sintering pressure of 25 MPa, a holding time of 60 min, and a vacuum degree of <10 -2 MPa; finally, a nuclear shielding material block is obtained.

[0111] Example 2

[0112] This embodiment provides a nuclear shielding material, which is different from the embodiment 1 in that the sintering temperature is 1100° C.; the rest is the same as the embodiment 1.

[0113] Example 3

[0114] This embodiment provides a nuclear shielding material, which is different from the embodiment 1 in that the sintering temperature is 1200° C.; the rest is the same as the embodiment 1.

[0115] Example 4

[0116] This embodiment provides a nuclear shielding material, which is different from the embodiment 1 in that the sintering temperature is 1300° C.; the rest is the same as the embodiment 1.

[0117] Example 5

[0118] This embodiment provides a nuclear shielding material, which is different from the embodiment 1 in that the sintering temperature is 1100° C. and the sintering pressure is 30 MPa; the rest is the same as the embodiment 1; the obtained nuclear shielding material is recorded as WC-HEA1.

[0119] Example 6

[0120] This embodiment provides a nuclear shielding material, including tungsten carbide (WC) and the high entropy alloy of Preparation Example 2; the preparation method is different from that of Example 1 in that: the sintering temperature is 1100°C, the sintering pressure is 30 MPa, and the rest is the same as Example 1; the obtained nuclear shielding material is recorded as WC-HEA2.

[0121] Example 7

[0122] This embodiment provides a nuclear shielding material, including tungsten carbide (WC) and the high entropy alloy of Preparation Example 3; the preparation method is different from that of Example 1 in that: the sintering temperature is 1100°C, the sintering pressure is 30 MPa, and the rest is the same as Example 1; the obtained nuclear shielding material is recorded as WC-HEA3.

[0123] Result detection

[0124] 1. Performance Characterization of High Entropy Alloys of Preparation Examples 1 to 3

[0125] 1) Morphology characterization of high entropy alloy powder

[0126] Figure 1 The SEM images of the high entropy alloy powders of Preparation Examples 1 to 3 are all 10 μm. Figure 1 Figures a through c represent SEM images of the high-entropy alloys from Preparation Examples 1 through 3, respectively. As can be seen from the figures, the high-entropy alloy powders are primarily spherical in shape, with a smooth surface free of pores and impurities, demonstrating that the powders prepared by the vacuum atomization method of the present invention are of excellent quality and achieve the desired effect.

[0127] 2) Crystal structure characterization of high entropy alloys

[0128] Figure 2The following are XRD patterns of the high-entropy alloy powders prepared in Examples 1-3. The XRD patterns show that while the elemental composition and content of the three high-entropy alloy powders vary, the diffraction peaks of each powder are essentially identical, with the strongest peak appearing at a diffraction angle of 43.6°. Using JADE analysis software, these peaks can be identified as FCC phases. The formation of solid solution phases in high-entropy alloys is related to the combined atomic size difference δ between the components. Because the atomic radii of the five elements, Cr, Cu, Fe, and Ni, are nearly identical and lack lattice distortion, the FCC phase, with its high atomic packing density, is easily formed. Compared to BCC and HCP phases, the FCC phase has a more closely packed slip plane and lower lattice resistance to dislocation slip, effectively improving the material's plasticity and toughness.

[0129] In addition, combined with the composition ratio of each element in the high entropy alloy in Table 1, it can be further concluded that HEA-2 with a lower Cr content belongs to the Fe-Ni-rich FCC phase (JCPDS: 15-0806), while HEA-1 belongs to the Fe-Ni-Cr-rich FCC phase (JCPDS: 33-0397).

[0130] 3) Characterization of high entropy alloy powder particle size

[0131] Figure 3 is the particle size distribution diagram of the high entropy alloy powder of Preparation Examples 1 to 3, wherein: Figure 3 Figures a through c in the figure represent the particle size distributions of the high-entropy alloys from Preparation Examples 1 through 3, respectively. As can be seen from these figures, the particle sizes of the three high-entropy alloy powders range from 7.51 to 8.06 μm, with D50 values ​​ranging from 13.62 to 14.41 μm, indicating a relatively small overall particle size. Small-sized high-entropy alloy powders are more easily refined and adhere to the WC particle surface during ball milling, enhancing the binding force of individual droplets on the WC particles during spray granulation, thereby improving the microstructure, sphericity, and deposition efficiency of the granulated powders.

[0132] 2. Performance Characterization of Nuclear Shielding Materials of Examples 1 to 7

[0133] 1) Effects of different sintering temperatures on nuclear shielding materials in Examples 1 to 4

[0134] ① The influence of sintering temperature on the microstructure of nuclear shielding materials

[0135] Figures 4 to 7 The SEM images and EDS spectra of the nuclear shielding materials of Examples 1 to 4 are shown in order, wherein: Figures 4 to 7 Figure a in the figure shows SEM images, and Figure b in the figure shows EDS spectra. Figure 8 a in the figure is a schematic diagram of the porosity of the nuclear shielding material calculated by the image method in Example 1. Examples 2 to 4 are calculated in the same way. Figure 8 b in the table is the porosity test result of Examples 1 to 4. Figures 4 to 8 It can be seen that a small number of pores exist within the nuclear shielding material blocks in each example. Using an image-based porosity analysis, it was found that when the overall porosity of the nuclear shielding material exceeds 5.0%, the density is low. The existence of pores is due to the fact that during the continuous dissolution and precipitation of WC grains, the dissolved WC grains produce some vacancies, and the precipitated WC grains also agglomerate with surrounding WC grains, resulting in a large number of voids, which ultimately affects the density of the nuclear shielding material.

[0136] In addition, when the sintering temperature reaches 1300℃, the dissolution and precipitation rate of WC grains into the high-entropy alloy will be promoted, exacerbating the segregation growth behavior of WC grains, resulting in a porosity of 7.52% and reduced density. At the same time, since the high-entropy alloy acts as a binder phase, its delayed diffusion effect will affect the flow of the liquid binder phase, causing the binder phase to fail to fully wrap the WC grains, and also leading to an increase in the internal porosity of the bulk material. When the sintering temperature exceeds 1200℃, the Cu content in the nuclear shielding material decreases from 1.1wt% to 0.3wt%. During the high-temperature sintering process, the presence of low-melting-point Cu makes it easier for the high-entropy alloy to convert into a Cu-rich liquid phase. Higher sintering temperatures can further increase the liquid phase content. At this time, part of the liquid phase overflows and is lost along the edge of the sample under high pressure stress, resulting in a decrease in the Cu content in the nuclear shielding material, affecting the strength and hardness of the nuclear shielding material.

[0137] Figure 9 It is the XRD pattern of the nuclear shielding material in Examples 1 to 4. As can be seen from the figure, when the sintering temperature is 1050°C, the nuclear shielding material is still mainly composed of WC and FCC phases of high entropy alloys, but a small amount of low-melting-point copper atoms in the FCC phase will react with iron atoms to form a Cu-Fe alloy phase. As the sintering temperature reaches 1100°C, atoms such as copper, iron, and chromium in the high-entropy alloy react with tungsten atoms and carbon atoms dissolved in the high-entropy alloy to form a M3W3C (M is a metal) metastable carbide phase. As the sintering temperature increases further, the reaction behavior of low-melting-point copper atoms and iron atoms in the high-entropy alloy intensifies, and the diffraction peak intensity of the Cu-Fe alloy phase is significantly enhanced. The content of dissolved carbon and tungsten atoms does not increase significantly with the increase in sintering temperature, and the reaction rate of related elements in the high-entropy alloy is stable, so the diffraction peak of the M3W3C phase does not change significantly.

[0138] ②The influence of sintering temperature on the hardness of nuclear shielding materials

[0139] The hardness of the nuclear shielding materials of Examples 1 to 4 was tested in accordance with GB / T 230.1-2018. The results showed that the hardness of the nuclear shielding materials prepared at different sintering temperatures remained above 90 HRA. The formation of high-hardness nuclear shielding materials is primarily related to material density and grain size. On the one hand, under high temperature and high pressure, the high-entropy alloy phase containing the low-melting-point metal (Cu) fully contacts the surface of the WC particles and forms a microscopic metallurgical bond, ensuring the high density of the samples. On the other hand, the addition of vanadium carbide and chromium carbide grain inhibitors prevents excessive growth of WC particles during high-temperature sintering. This increases the contact area between small grains, enhances grain growth resistance, and increases the number of grain boundaries, effectively preventing dislocation migration and improving the hardness of the nuclear shielding materials. Furthermore, the high-entropy alloy phase also inhibits WC grain growth. This is primarily due to the delayed diffusion effect of the high-entropy alloy as a binder phase, which reduces the diffusion rate of W and C atoms in the binder phase, hindering the dissolution-precipitation growth of the WC grains. Compared with other sintering temperatures, the hardness of the nuclear shielding material block sintered at 1300°C is relatively the lowest (90 HRA), while the hardness of the nuclear shielding material sintered at 1100°C is relatively the highest (91.5 HRA). This is mainly because higher sintering temperatures (such as 1300°C) cause the high-entropy alloy phase to liquefy and overflow from the edge of the hot-pressed sample, resulting in a significant decrease in the high-entropy binder phase content. The loss of the high-entropy alloy phase not only reduces the interfacial bonding effect between WC and the high-entropy alloy phase, but also weakens its own inhibitory effect on the WC grain growth behavior.

[0140] ③The influence of sintering temperature on the wear resistance of nuclear shielding materials

[0141] The wear resistance of nuclear shielding materials was tested using a reciprocating friction and wear tester. The test conditions were: dry friction, load 10N, friction distance 5mm, total friction distance 500m, frequency 5Hz. The results are as follows Figure 10 As shown, Figure 10 Figures a through d in the figure represent the friction coefficient curves of the nuclear shielding materials of Examples 1 through 4, respectively. It can be seen that at the beginning of friction, the friction coefficient gradually increases and fluctuates greatly. As the friction time increases, the friction coefficient curve gradually stabilizes and the fluctuation amplitude decreases, indicating that the friction specimen has entered the stable wear stage from the pre-wear stage. Compared with other sintering temperatures, the friction coefficient value of the nuclear shielding material sintered at 1100°C is generally lower, fluctuating within the range of 0.3 to 0.5. The friction coefficient value of the nuclear shielding material sintered at 1300°C is relatively high, with the largest fluctuation amplitude (0.8 to 1.2), indicating poor wear resistance.

[0142] Figure 11 is a graph showing the average friction coefficient and wear rate of the nuclear shielding materials of Examples 1 to 4, wherein: Figure 11Figures a and b in the figure represent the average friction coefficient and wear rate diagrams, respectively. As can be seen from the figure, the average friction coefficient of the nuclear shielding material sintered at 1300℃ is the highest (0.98), while the average friction coefficient of the nuclear shielding material sintered at 1050℃ is the lowest, only 0.34. Further calculation of the wear rate of the nuclear shielding material shows that compared with other sintering temperatures, the wear rate of the nuclear shielding material sintered at 1100℃ is relatively low, only 0.96×10 -6 mm 3 / N·m, showing better wear resistance. It is worth noting that although the average friction coefficient of the nuclear shielding material sintered at 1300℃ is higher, the wear rate (1.18×10 -6 mm 3 / N·m) and nuclear shielding materials sintered at 1200℃ (1.17×10 -6 mm 3 / N·m) are basically consistent, indicating that the wear resistance is not only affected by the friction coefficient.

[0143] ④Effect of sintering temperature on the oxidation resistance of nuclear shielding materials

[0144] The oxidation resistance of the nuclear shielding materials of Examples 1 to 4 was tested in an atmospheric high-temperature heating furnace. During the test, the nuclear shielding materials were kept at 700°C for different times (30min, 90min and 150min). After the test, the weight change of the samples was weighed and the oxidation weight gain was calculated. At the same time, the same test was carried out using conventional commercially available cemented carbide YG-10 as a control group. Oxidation weight gain can be defined as the increase in weight per unit area of ​​the sample surface before and after oxidation, in units of mg·cm -2 .

[0145] Figure 12 The figure shows the results of the oxidation resistance test of the nuclear shielding materials of Examples 1 to 4, wherein: Figure 12 a~b in the figure represent the optical photograph and oxidation weight gain curve respectively. As can be seen from the figure, with the increase of oxidation time, the oxidation weight gain trend of the nuclear shielding material slows down, and all of them show good high-temperature oxidation resistance. The main reasons for this result are as follows: a. The formation of the oxide film with a passivating effect inhibits the continuous diffusion and oxidation of oxygen into the interior of the nuclear shielding material; b. The hysteresis diffusion effect of the high-entropy alloy greatly reduces the diffusion rate of the metal atoms inside it, effectively slowing down the oxidation rate of the high-entropy bonding phase. Compared with other sintering temperatures, the oxidation weight gain of the nuclear shielding material sintered at 1100℃ after oxidation for 30min, 90min and 150min is 3.95mg·cm respectively. -2 , 5.13 mg·cm -2 and 6.42 mg·cm -2The growth trend is relatively low, which is mainly due to the lower porosity of the core shielding material and the fewer oxidation channels, which slows down the oxidation rate of the sample surface. In addition, the traditional cemented carbide (YG-10) as the control group has an oxidation weight gain (9.03 mg cm -2 ) is significantly higher than that of the sample sintered at 1100°C, which shows that by controlling the sintering temperature, the nuclear shielding material obtained by the invention has more application advantages in the high-temperature field.

[0146] Figure 13 This is the EDS surface scanning result of the oxide layer peeling area of ​​the nuclear shielding material of Example 1. As shown in the figure, the surface oxide layer of the nuclear shielding material is mainly distributed with Fe, Cu, Cr and O elements, while the oxide layer peeling area is mainly distributed with W and O elements. This shows that the surface of the oxide layer is mainly composed of oxides containing Fe, Cu and Cr, while the interior of the oxide layer is mainly composed of oxides containing W. Through further analysis, it can be determined that the surface of the oxide layer is mainly composed of MWO4 (M = Ni, Cr, Cu or Fe), and there is also a small amount of WO3, while the interior of the oxide layer is mainly composed of WO3. MWO4 is a metal oxide M formed by the oxidation of high entropy alloy elements. x O y It further reacts with WO3, which is formed by the high-temperature reaction of WC and O and is mainly loose and porous. In addition, the WO3 reaction produces oxidizing gas (CO2). The release of CO2 creates a large pressure difference between the inside and outside of the pores in the oxide layer, causing crack propagation and localized spalling of the oxide layer surface.

[0147] 2) Performance Characterization of Nuclear Shielding Materials of Examples 5 to 7

[0148] ① Characterization of tissue morphology

[0149] Figures 14 to 16 The SEM images and EDS spectra of the nuclear shielding materials of Examples 5 to 7 are shown in sequence, wherein: Figures 14 to 16 Figures a to b in the figure represent SEM images and EDS spectra respectively. Figures 14 to 16 It can be seen that the overall structure of the nuclear shielding materials of each embodiment is dense, with only a small amount of pores. Calculated by the image method, the porosity of WC-HEA1 (Example 5) is the lowest, only 1.13±0.21%. In addition, comparing this result with the porosity of the nuclear shielding material obtained in Example 2 (5.15±0.6%), it is found that increasing the sintering pressure can significantly reduce the porosity. This is mainly because the increase in sintering pressure promotes close bonding between the particle units, reduces the diffusion resistance between the particles, and thus makes the structure of the nuclear shielding material denser. Further EDS analysis of the nuclear shielding materials of Examples 5 to 7 found that the nuclear shielding materials are mainly composed of elements such as W, C, Fe, Ni, Cr, Cu and Co.

[0150] ② Characterization of tissue composition

[0151] Figure 17 The XRD patterns of the nuclear shielding materials of Examples 5 to 7 are shown. It can be seen that the diffraction peaks of Examples 5 to 7 are basically the same, and the main phase structures include WC, FCC, M3W3C (M = Fe, Ni, Cr, Cu or Co) and Cu-Fe phases. During the high-temperature sintering process, the W and C atoms dissolved in the high-entropy alloy react with the atoms of the high-entropy alloy to form the M3W3C phase. The formation of M3W3C consumes the WC content, affecting the hardness and wear resistance of the nuclear shielding material. In addition, while the M3W3C reaction is being generated, the size of the W atoms dissolved in the high-entropy alloy is quite different from that of the Fe, Ni, Cr, Co or Cu atoms, resulting in severe lattice distortion, and thus the diffraction peak intensity of the FCC phase is reduced.

[0152] ③Hardness characterization

[0153] Figure 18 The following graph shows the hardness test results for the nuclear shielding materials of Examples 5-7. The hardness tests were conducted in accordance with GB / T 230.1-2018. As can be seen from the graph, the hardness of the nuclear shielding materials of Examples 5-7 all exceeded 89 HRA. Compared with the nuclear shielding materials of the other examples, WC-HEA1 had the highest hardness, reaching 91.9±0.6 HRA. WC-HEA2 had a slightly lower hardness than WC-HEA1, at 91.3±0.5 HRA. The hardness of WC-HEA3 was lower than that of the other two examples. Combined with the aforementioned analysis of the microstructure of the nuclear shielding materials of Examples 5-7, it can be found that low porosity is the primary factor contributing to the higher hardness of WC-HEA1. On the one hand, the carbide particles are more closely connected, effectively preventing dislocation migration during deformation and improving the material's resistance to plastic deformation. On the other hand, the high-entropy alloy binder phase fully contacts the WC particle surface, forming a microscopic metallurgical bond, which can inhibit the initiation and propagation of interfacial cracks.

[0154] ④ Characterization of flexural strength

[0155] Figure 19The figure shows the test results of the flexural strength of the nuclear shielding materials of Examples 5 to 7. In addition to hardness, flexural strength is also an important indicator for measuring the mechanical properties of nuclear shielding materials. The present invention uses a three-point bending test to characterize the flexural strength of the nuclear shielding material block, and the specific test process is carried out in accordance with the YBT 5349-2014 standard. The test obtains no less than 3 groups of valid data, and the average value is taken for flexural strength evaluation. It can be seen from the test results that the flexural strength of the nuclear shielding materials of Examples 5 to 7 exceeds 940MPa, and the flexural strength is high. In addition, the flexural strength of WC-HEA3 reaches 1286.25±161.48MPa, which is higher than that of WC-HEA1 and WC-HEA2. This is because the presence of Co element in WC-HEA3 can improve the wettability between the high-entropy alloy binder phase and the WC hard phase, which is beneficial to further improve the grain boundary bonding state. In addition, its porosity is better, and more micropores can consume more crack propagation energy through crack deflection, inhibiting crack propagation behavior, thereby improving the flexural strength. Therefore, the flexural strength is higher.

[0156] ⑤ Wear performance characterization

[0157] Figure 20 is a friction curve diagram of the nuclear shielding material of Examples 5 to 6, wherein Figure 20 a to b in the figure represent the friction curves of the nuclear shielding materials of Examples 5 to 7, respectively. Figure 20 Figure c shows the friction curve using commercially available 316 stainless steel as a control. It can be seen that at the beginning of friction, the friction coefficient of the nuclear shielding materials of each example gradually increases and fluctuates significantly. As the friction time increases, the friction coefficient curve gradually stabilizes and the fluctuation amplitude decreases (between 0.4 and 0.7), indicating that the friction specimen has entered the stable wear stage after the pre-wear stage.

[0158] Figure 21 The average friction coefficient and wear rate of the nuclear shielding materials of Examples 5 to 7 are shown in FIG. Figure 21 a~b in the figure represent the average friction coefficient and wear rate diagrams respectively, and the conventional commercial 316 stainless steel is used as a control. As can be seen from the figure, the average friction coefficient of WC-HEA2 is relatively the lowest, only 0.42±0.12, while the friction coefficients of WC-HEA1, WC-HEA3 and 316L stainless steel are similar, about 0.52. In terms of wear rate, the wear rate of 316L stainless steel reaches 5.11×10 -6 mm 3 / N·m, which is 3.10 to 5.2 times the wear rate of the nuclear shielding materials in Examples 5 to 7. It can be seen that the nuclear shielding material of the present invention has better wear resistance. The main reason for this result is that the presence of hard phases such as WC and M3W3C (M = Fe, Ni, Cr, Cu or Co) can effectively improve the hardness and elastic modulus of the nuclear shielding material, enhance the surface resistance to plowing and spalling under friction shear stress, and thus significantly reduce the wear rate of the sample. Compared with WC-HEA2 and WC-HEA3, WC-HEA1 has the lowest wear rate, about 0.98×10 -6 mm 3 / Nm, which can be attributed to the relatively higher hardness and lower porosity of WC-HEA1. High hardness effectively reduces the depth of the friction pair's penetration into the material under load, preventing the formation of hard spots during friction and reducing surface plowing and spalling. Furthermore, low porosity means a lower content of micro-defects such as pores and cracks within the material, effectively suppressing crack initiation and propagation in the friction contact area under alternating loads, further mitigating surface wear of nuclear shielding materials.

[0159] ⑥ Characterization of antioxidant properties

[0160] The oxidation resistance of the nuclear shielding materials of Examples 5 to 7 was tested using a high-temperature heating furnace in an atmospheric atmosphere. During the test, the nuclear shielding materials were kept at 700°C for different times (4h, 6h, 8h and 10h). After the test, the weight change of the samples was measured and the oxidation weight gain was calculated. The oxidation weight gain can be defined as the increase in weight per unit area of ​​the sample surface before and after oxidation, with the unit being mg·cm -2 The test results are as follows. Figure 22 shown.

[0161] Figure 22 The weight gain curves of the nuclear shielding materials of Examples 5 to 7 at different oxidation times are shown in FIG. As can be seen from the figure, the weight gain of WC-HEA1 after oxidation for 2.5 h is only 6.42±0.42 mg·cm -2 This is due to the optimized oxidation of the Cr element in the high-entropy alloy to form Cr2O3. As oxidation continues, this oxide grows laterally to form a dense oxide film with a passivating effect. Furthermore, the delayed diffusion effect of the high-entropy alloy can slow the outward diffusion of metal atoms in the alloy, further improving the oxidation resistance of the nuclear shielding material and improving its high-temperature resistance.

[0162] When the oxidation time increases to 4 h, the Cr2O3 passivation film cannot completely cover the surface of the nuclear shielding material. Oxygen contacts and oxidizes WC to form WO3, accompanied by the release of oxidizing gas (CO2). This leads to an increase in defects such as pores inside the sample and an increase in the number of oxidation channels, resulting in an increase in the oxidation weight gain of WC-HEA1 to 9.56±2.34 mg·cm-2 , and as the oxidation time extended to 10 h, the oxidation weight gain reached 61.31±5.88 mg·cm -2 After 10h of oxidation, the weight gain of WC-HEA3 reached 109.12~113.86mg·cm -2 It can be seen that when the oxidation time increases from 4h to 10h, the oxidation weight gain trend of WC-HEA1 is relatively the lowest.

[0163] In summary, in the nuclear shielding material provided by the present invention, a relatively dense structure is formed by adjusting the composition and content of each element in the high-entropy alloy, so that the high-entropy alloy presents a single FCC phase, its slip plane is more closely packed, and the lattice resistance to dislocation slip is lower, which can effectively improve the plasticity and toughness of the material, and improve the mechanical properties, wear resistance, and high-temperature oxidation resistance. At the same time, the high-entropy alloy particle size D50 prepared by the present invention is in the range of 7 to 9 μm, so it is easier to be refined and adhere to the surface of WC particles during the ball milling process, which is beneficial to the binding of the high-entropy alloy to the WC particles, thereby enhancing the performance of the nuclear shielding material.

[0164] In addition, in the nuclear shielding material of the present invention, the Co element content in the high entropy alloy is low and can even be free of Co element, which reduces the Co-60 isotope formed by the stimulation of cobalt under a long-term radiation environment, ensures shielding performance, and the high entropy alloy formed not only has excellent mechanical strength, wear resistance and high temperature oxidation resistance, but also takes into account the nuclear shielding effect. The nuclear shielding material formed by the high entropy alloy and WC composite also has mechanical properties, wear resistance and oxidation resistance, and can further enhance the mechanical strength and hardness of the nuclear shielding material by adjusting the sintering temperature, reduce porosity, and avoid Cu overflow loss. Therefore, the nuclear shielding material of the present invention can meet the requirements of the nuclear radiation field for the performances such as material mechanical strength, wear resistance, oxidation resistance, and nuclear shielding ability.

[0165] The above embodiments are preferred implementation modes of the present invention, but the implementation modes of the present invention are not limited to the above embodiments. Any other changes, modifications, substitutions, combinations, and simplifications that do not deviate from the spirit and principles of the present invention should be considered as equivalent replacement methods and are included in the scope of protection of the present invention.

Claims

1. A nuclear shielding material, characterized in that: The raw materials for preparation include tungsten carbide and high entropy alloy; the high entropy alloy includes the following components by mass percentage: Fe: 20 to 40%, Ni: 20 to 40%, Cr: 12 to 30%, Cu: 14 to 25%, Co: 0 to 15%.

2. The nuclear shielding material according to claim 1, characterized in that In the nuclear shielding material, the mass content of inevitable impurities in the high entropy alloy is ≤0.35%.

3. The nuclear shielding material according to claim 1 or 2, characterized in that: In the nuclear shielding material, the crystal phase of the high entropy alloy is FCC phase.

4. The nuclear shielding material according to claim 1 or 2, characterized in that: In the nuclear shielding material, the high entropy alloy is in powder form with a particle size of D 50 7~9μm.

5. The nuclear shielding material according to claim 1, characterized in that In the nuclear shielding material, the preparation method of the high entropy alloy comprises the following steps: The raw material components of the high entropy alloy are mixed and smelted, and then powdered to obtain the high entropy alloy.

6. The nuclear shielding material according to claim 5, characterized in that The smelting pressure is (3-7)×10 - 2 MPa; And / or, the smelting time is 20 to 40 minutes.

7. The nuclear shielding material according to claim 5, characterized in that The powder making method includes a vacuum air atomization method; the conditions of the vacuum air atomization method meet at least one of the following a) to d): a) Air pressure is 2-6 MPa; b) The air flow velocity is 100-150 mm / s; c) The liquid flow diameter is 5 to 10 mm; d) The air flow injection angle is 30° to 60°.

8. The nuclear shielding material according to claim 1, characterized in that In the nuclear shielding material, the mass ratio of the tungsten carbide to the high entropy alloy is (9-14):

1.

9. A method for preparing the nuclear shielding material according to any one of claims 1 to 8, characterized in that: The following steps are involved: The nuclear shielding material is obtained by mixing tungsten carbide and high entropy alloy and ball milling, followed by cold pressing and sintering.

10. The preparation method according to claim 9, characterized in that The ball-to-material ratio of the ball mill is (7-9):1; And / or, the sintering temperature is 1000-1300°C.

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